U.S. patent application number 15/922596 was filed with the patent office on 2018-08-30 for high capacity electrode materials for batteries and process for their manufacture.
This patent application is currently assigned to UCHICAGO ARGONNE, LLC. The applicant listed for this patent is UCHICAGO ARGONNE, LLC. Invention is credited to Christopher S. JOHNSON, Tijana RAJH, Elena SHEVCHENKO, Sanja TEPAVCEVIC, Hui XIONG.
Application Number | 20180248185 15/922596 |
Document ID | / |
Family ID | 47996690 |
Filed Date | 2018-08-30 |
United States Patent
Application |
20180248185 |
Kind Code |
A1 |
JOHNSON; Christopher S. ; et
al. |
August 30, 2018 |
HIGH CAPACITY ELECTRODE MATERIALS FOR BATTERIES AND PROCESS FOR
THEIR MANUFACTURE
Abstract
The present invention provides a nanostructured metal oxide
material for use as a component of an electrode in a lithium-ion or
sodium-ion battery. The material comprises a nanostructured
titanium oxide or vanadium oxide film on a metal foil substrate,
produced by depositing or forming a nanostructured titanium dioxide
or vanadium oxide material on the substrate, and then, optionally,
charging and discharging the material in an electrochemical cell
from a high voltage in the range of about 2.8 to 3.8 V, to a low
voltage in the range of about 0.8 to 1.4 V over a period of about
1/30 of an hour or less. Lithium-ion and sodium-ion electrochemical
cells comprising electrodes formed from the nanostructured metal
oxide materials, as well as batteries formed from the cells, also
are provided.
Inventors: |
JOHNSON; Christopher S.;
(Naperville, IL) ; XIONG; Hui; (Woodridge, IL)
; RAJH; Tijana; (Naperville, IL) ; SHEVCHENKO;
Elena; (Riverside, IL) ; TEPAVCEVIC; Sanja;
(Chicago, IL) |
|
Applicant: |
Name |
City |
State |
Country |
Type |
UCHICAGO ARGONNE, LLC |
Chicago |
IL |
US |
|
|
Assignee: |
UCHICAGO ARGONNE, LLC
Chicago
IL
|
Family ID: |
47996690 |
Appl. No.: |
15/922596 |
Filed: |
March 15, 2018 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
|
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14227341 |
Mar 27, 2014 |
9935314 |
|
|
15922596 |
|
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|
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PCT/US2012/042230 |
Jun 13, 2012 |
|
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14227341 |
|
|
|
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61540604 |
Sep 29, 2011 |
|
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Current U.S.
Class: |
1/1 |
Current CPC
Class: |
H01M 4/0452 20130101;
H01M 4/0442 20130101; H01M 4/661 20130101; H01M 4/48 20130101; H01M
4/0445 20130101; Y02E 60/10 20130101; H01M 10/0525 20130101; H01M
4/483 20130101; H01M 10/054 20130101; H01M 4/139 20130101; C25D
9/10 20130101; C25D 9/08 20130101; H01M 4/13 20130101; C25D 11/26
20130101 |
International
Class: |
H01M 4/48 20060101
H01M004/48; H01M 4/04 20060101 H01M004/04; H01M 10/054 20060101
H01M010/054; C25D 9/08 20060101 C25D009/08; C25D 9/10 20060101
C25D009/10; C25D 11/26 20060101 C25D011/26; H01M 10/0525 20060101
H01M010/0525; H01M 4/13 20060101 H01M004/13; H01M 4/139 20060101
H01M004/139; H01M 4/66 20060101 H01M004/66 |
Goverment Interests
CONTRACTUAL ORIGIN OF THE INVENTION
[0002] The United States Government has rights in this invention
pursuant to Contract No. DE-AC02-06CH11357 between the United
States Government and UChicago Argonne, LLC representing Argonne
National Laboratory.
Claims
1. A nanostructured metal oxide electrode for use in a sodium-ion
battery, the electrode comprising a nanostructured film of
bilayered V.sub.2O.sub.5 on a metal foil substrate.
2. The electrode of claim 1, wherein the bilayered V.sub.2O.sub.5
is amorphous and the layers thereof have an interlayer spacing of
greater than about 12 .ANG..
3. The electrode of claim 2, wherein the bilayered V.sub.2O.sub.5
is produced by electrochemical deposition from a VO.sup.2+ salt
solution onto the metal foil substrate, and drying the deposited
film to remove water therefrom.
4. The electrode of claim 1, wherein the bilayered V.sub.2O.sub.5
comprises nanostructured crystalline bilayered V.sub.2O.sub.5
produced by electrochemical deposition from a VO.sup.2+ salt
solution onto the metal foil substrate to form an amorphous
bilayered V.sub.2O.sub.5 film, drying the film, and then charging
and discharging the dried film in an electrochemical cell from a
high voltage in the range of about 3.3 to 3.8 V, to a low voltage
in the range of about 1.4 to 1.6 V, over a period of about 1/30 of
an hour or less.
5. The electrode of claim 1, wherein the substrate comprises nickel
or stainless steel.
6. A sodium-ion electrochemical cell comprising a cathode, an anode
and a nonaqueous sodium-containing electrolyte therebetween,
wherein the anode comprises the electrode of claim 1.
7. A sodium-ion battery comprising a plurality of electrochemically
linked electrochemical cells of claim 6.
8. A sodium-ion electrochemical cell comprising a cathode, an anode
and a nonaqueous sodium-containing electrolyte therebetween,
wherein the anode comprises the nanostructured vanadium oxide
electrode of claim 2.
9. A sodium-ion battery comprising a plurality of electrochemically
linked electrochemical cells of claim 8.
10. A sodium-ion electrochemical cell comprising a cathode, an
anode and a nonaqueous sodium-containing electrolyte therebetween,
wherein the anode comprises the nanostructured vanadium oxide
electrode of claim 4.
11. A sodium-ion battery comprising a plurality of
electrochemically linked electrochemical cells of claim 10.
12. A sodium-ion electrochemical cell comprising a cathode, an
anode and a nonaqueous sodium-containing electrolyte therebetween,
wherein the anode comprises a film of densely packed amorphous
TiO.sub.2 nanotubes on a surface of a metal foil substrate; and
wherein the densely packed nanotubes have a wall thickness of about
18 to 25 nm and an outer tubular diameter of about 100 to 150
nm.
13. The sodium-ion electrochemical cell of claim 12, wherein the
substrate is a titanium foil.
14. The sodium-ion electrochemical cell of claim 12, wherein the
densely packed TiO.sub.2 nanotubes are oriented perpendicular to
the surface of the metal foil substrate.
15. The sodium-ion electrochemical cell of claim 12, wherein the
substrate is a titanium foil, and the densely packed amorphous
TiO.sub.2 nanotubes are produced by electrochemical anodization of
a surface of the titanium foil, and subsequently drying the film to
remove water therefrom.
16. The sodium-ion electrochemical cell of claim 15, wherein the
densely packed TiO.sub.2 nanotubes are oriented perpendicular to
the surface of the substrate.
17. A sodium-ion battery comprising a plurality of
electrochemically linked electrochemical cells of claim 12.
18. A sodium-ion battery comprising a plurality of
electrochemically linked electrochemical cells of claim 14.
19. A sodium-ion battery comprising a plurality of
electrochemically linked electrochemical cells of claim 15.
20. A sodium-ion battery comprising a plurality of
electrochemically linked electrochemical cells of claim 16.
Description
CROSS-REFERENCE TO RELATED APPLICATION
[0001] This application is a division of U.S. application Ser. No.
14/227,341, filed on Mar. 27, 2014, which is a continuation of
International Application No. PCT/US2012/042230, filed on Jun. 13,
2012, which claims the benefit of U.S. Provisional Application Ser.
No. 61/540,604 filed on Sep. 29, 2011, each of which is
incorporated herein by reference in its entirety.
FIELD OF THE INVENTION
[0003] This invention relates to high capacity materials for
batteries and processes for preparing such materials. More
particularly, this invention relates to amorphous titanium oxide
and vanadium oxide materials useful as components of electrodes for
lithium-ion and sodium-ion batteries.
BACKGROUND OF THE INVENTION
[0004] Limited energy resources and the growing demand to decrease
greenhouse gas emissions have intensified research of carbon-free
energy sources. Batteries that store high-energy densities will
play a large role in implementation of green energy technologies
and non-petroleum vehicular mobility. To date, rechargeable Li-ion
batteries offer the highest energy density of any battery
technology, and are expected to provide a solution for our future
energy-storage requirements. Unfortunately, Li-ion batteries have a
number of limitations, such as capacity loss over time during
long-term cycling due to phase transitions leading to detrimental
volume changes in the electrode materials. This can cause local
atom rearrangements that block the diffusion of Li ions, leading to
high over potentials and loss of capacity. In addition, Li-ion
batteries that have been charged quickly can form dendritic Li
deposition at the commercial graphite anode and can create a safety
problem that in the worst scenario could cause thermal runaway,
cell rupture and explosion. Use of chemically inert anode materials
such as metal oxides with lithiation voltages positive of
Li-deposition compared to carbon can address these issues and
improve safety of Li-ion battery operation.
[0005] High-performance battery materials are critical for the
development of new alternative energy storage systems. While Li-ion
batteries are a mature technology for energy storage, disadvantages
include cost, Li supply, safety, reliability and stability.
Moreover, the electrolyte stability over time is additional major
concern for long-term operation and advanced applications. Thus,
the discovery, research and development of new transporting ions
that can provide an alternative choice to Li batteries are
essential for further advancement of energy storage materials.
Sodium-based batteries are attractive due to the promise of low
cost associated with the abundance of sodium, and enhanced
stability of non-aqueous battery electrolytes due to the lower
operating voltages. However, lower voltage leads to insufficient
energy density, thus cathode materials for Na batteries must
possess high-capacities. Since the ionic volume of sodium is almost
twice that of lithium, unique crystalline structures have to be
used to accommodate incorporation of large ions.
[0006] Titanium dioxide is one of the few metal oxide materials
that intercalates Li ions at reasonably low voltage (approximately
1.2 V vs. Li/Li.sup.+) and is suitable as a battery anode material.
The first attempts of using TiO.sub.2 for a durable and safe
electrode material were focused on microcrystalline TiO.sub.2
materials such as rutile, anatase, and TiO.sub.2(B). These
electrodes materials showed moderate specific capacities (e.g.,
maximum Li uptake of 0.5 Li/Ti for anatase and TiO.sub.2(B), and no
activity for rutile) due to the limited room temperature reactivity
and conductivity at microscale. Recently, the idea of using
TiO.sub.2 electrodes has been revisited with the consideration that
nanosize morphologies provide enhanced intercalation kinetics and
large surface area associated with high accessibility of
transporting ions. Reversible capacities with stoichiometries up to
about approximately 0.5-0.7Li/Ti have been demonstrate; however,
repetitive cycling caused loss of capacity independent of the
crystalline modification. It was expected that all TiO.sub.2-based
electrodes will have an intrinsic capacity limitation, because of a
low number of crystallographic sites, and their electronically
insulating structure. Both experimental and theoretical studies of
intercalation of Li ions in crystalline polymorphs of TiO.sub.2
show that high lithium content can be obtained exclusively at
elevated temperatures raising concerns from the application point
of view.
[0007] Vanadium pentoxide (V.sub.2O.sub.5) has been intensively
studied as the positive electrode (anode) material for lithium ion
batteries. In previous studies, various fabrication methods were
used: sputtering, thermal evaporation, thermal decomposition,
electrophoretic deposition and many chemical routes, such as
hydrothermal synthesis, and sol-gel method. It has been reported
that chemical composition, crystal structure and crystallinity of
V.sub.2O.sub.5 may have pivotal roles in lithium ion intercalation
capacity and cycling stability. To date, the practical application
of vanadium oxides in lithium ion batteries was still limited due
to the poor cyclic stability and undesirable phase transitions to
inactive materials.
[0008] In contrast to lithium batteries, a report of Liu et al.
considered employment of vanadium oxide by insertion/deinsertion of
sodium-ion into NaV.sub.6O.sub.15 nanorods. Even though the
morphology of the reported material could be retained, the initial
discharge capacity of 142 mAh/g substantially decreased by cycling
at higher current densities, which led to poor overall performance.
In order to explore full potential in utilization of nanostructured
V.sub.2O.sub.5 electrodes for sodium batteries a detailed
fundamental insight is necessary.
[0009] There is an ongoing need for new, more efficient, electrode
materials for lithium-ion and sodium-ion batteries. The present
invention addresses this need.
SUMMARY OF THE INVENTION
[0010] The present invention provides a nanostructured metal oxide
material for use as a component of an electrode of a lithium-ion or
sodium-ion battery. The material comprises a nanostructured
titanium oxide or vanadium oxide film on a metal foil substrate.
The film is produced by depositing or forming a nanostructured
material selected from titanium dioxide and vanadium oxide on the
surface of the substrate, and then charging and discharging the
material in an electrochemical cell from a high voltage in the
range of about 2.8 to 3.8 V, to a low voltage in the range of about
0.8 to 1.4 V over a period of about 1/30 of an hour or less.
[0011] In one embodiment, the nanostructured metal oxide material
of the present invention comprises a film of titanium dioxide
nanotubes. Preferably, the titanium dioxide nanotubes comprise
densely packed TiO.sub.2 nanotubes oriented perpendicular to the
surface of the metal foil substrate. In one form of this
embodiment, the substrate is a titanium foil and the densely packed
TiO.sub.2 nanotubes are amorphous and are produced by
electrochemical anodization of a surface of the titanium foil, and
subsequently drying the film to remove water therefrom (e.g., by
heating at a temperature in excess of 100.degree. C.). The densely
packed TiO.sub.2 nanotubes can have a wall thickness of about 8 to
25 nm and an outer tubular diameter of about 50 to 150 nm. In one
preferred form, the nanotubes have a wall thickness of about 8 to
12 nm (e.g., about 10 nm) and an outer tubular diameter of about 50
to 70 nm (e.g., about 65 nm). In another form, the densely packed
TiO.sub.2 nanotubes have a wall thickness of about 18 to 25 nm and
an outer tubular diameter of about 100 to 150 nm.
[0012] In another preferred form of this embodiment the titanium
dioxide comprises nanostructured crystalline TiO.sub.2 produced by
electrochemical anodization of a surface of the titanium foil to
form amorphous nanotubes, subsequently drying the film to remove
water, and then charging and discharging the amorphous film in an
electrochemical cell from a high voltage to a low voltage over a
short period of time. For example, charging and discharging from a
high voltage in the range of about 2.6 to 3 V (e.g., about 2.8 to
2.9 V), to a low voltage in the range of about 0.7 to 1 V (e.g.,
about 0.8 to 0.9 V), over a period of about 1/30 of an hour or
less.
[0013] In another preferred embodiment, the nanostructured metal
oxide material comprises a film of bilayered V.sub.2O.sub.5. In one
preferred form, the bilayered V.sub.2O.sub.5 is amorphous and the
layers thereof have an interlayer spacing of greater than about 12
.ANG.. In another preferred form, the bilayered V.sub.2O.sub.5 is
amorphous and is produced by electrochemical deposition from a
VO.sup.2+ salt solution (e.g., aqueous VOSO.sub.4) onto the metal
foil substrate, and drying deposited film to remove water
therefrom. In another preferred form, the bilayered V.sub.2O.sub.5
comprises nanostructured crystalline bilayered V.sub.2O.sub.5
produced by electrochemical deposition from a VO.sup.2+ salt
solution onto the metal foil substrate to form an amorphous
bilayered V.sub.2O.sub.5 film, drying film, and then charging and
discharging the dried film in an electrochemical cell from a high
voltage in the range of about 3.3 to 3.8 V (e.g., about 3.5 V), to
a low voltage in the range of about 1.4 to 1.6 V (e.g., about 1.5
V), over a period of about 1/30 of an hour or less.
[0014] The metal foil substrate can comprise any metal suitable for
use as a current collector for an electrode in lithium-ion or
sodium-ion electrochemical cell or battery. Preferably, when the
metal oxide material comprises titanium oxide, the substrate
comprises titanium. When the metal oxide material comprises
vanadium oxide, the substrate preferably comprises nickel or
stainless steel.
[0015] In one aspect, the present invention provides a lithium-ion
electrochemical cell comprising a cathode, an anode, and a
nonaqueous lithium-containing electrolyte therebetween. In one
preferred form, the anode comprises a nanostructured titanium oxide
material of the invention as described herein. If desired, the
anode can comprise a carbon material or other suitable components
in addition to the titanium oxide material. The cathode in this
embodiment can comprise a vanadium oxide material as described
herein, or can comprise any other cathode material suitable for use
in electrochemical cells, such cathode materials being well known
in the battery art.
[0016] In another preferred form of the lithium-ion electrochemical
cell, the cathode comprises a nanostructured vanadium oxide
material of the invention as described herein. The anode in this
form can comprise a titanium oxide material as described herein, or
can comprise any other anode material suitable for use in
electrochemical cells, such anode materials being well known in the
battery art.
[0017] The electrolyte in the lithium-ion electrochemical cells and
batteries can comprise any lithium-ion containing electrolyte
material that is suitable for use in lithium-ion electrochemical
cells and batteries. Such electrolyte materials also are well known
in the battery art.
[0018] The present invention also provides a sodium-ion
electrochemical cell comprising a cathode, an anode, and a
nonaqueous sodium-containing electrolyte therebetween. In one
preferred form, the anode comprises a nanostructured titanium oxide
material of the invention as described herein. If desired, the
anode can comprise a carbon material or any other suitable
component in addition to the titanium oxide material. The cathode
in this embodiment can comprise a vanadium oxide material as
described herein, or can comprise any other cathode material
suitable for use in electrochemical cells, such cathode materials
being well known in the battery art.
[0019] In another preferred form of the sodium-ion electrochemical
cell, the cathode comprises a nanostructured vanadium oxide
material of the invention as described herein. The anode in this
form can comprise a titanium oxide material as described herein, or
can comprise any other anode material suitable for use in
electrochemical cells, such anode materials being well known in the
battery art.
[0020] The electrolyte in the sodium-ion electrochemical cells and
batteries can comprise any sodium-ion containing electrolyte
material that is suitable for use in sodium-ion electrochemical
cells and batteries. Such electrolyte materials also are well known
in the battery art.
[0021] The electrochemical cells of the present invention can be
assembled in to a battery by electrochemically linking the cells
(e.g., in series or parallel or both), if desired. In addition, the
batteries and electrochemical cells can include other components
besides the cathode, anode and electrolyte, such as a housing,
porous separators, binders, and the like, as is well known in the
battery art.
[0022] The present invention provides an electrochemically driven
transformation of amorphous TiO.sub.2 nanotubes for Li-ion battery
anodes into a face-centered-cubic Li.sub.2Ti.sub.2O.sub.4
crystalline phase that self-improves as the cycling proceeds. The
intercalation and deintercalation processes of Li ions in the
electrochemically-grown TiO.sub.2 nanotubes provide spontaneous
development of a long-range order in amorphous TiO.sub.2 in the
presence of high concentration of Li ions (>75%). Reversible
intercalation and preservation of the resultant cubic structure
during cycling is maintained, and the electrode shows superior
stability for more than 600 cycles. Moreover, the adopted cubic
structure supports very fast (26 s) charging with good retention of
capacity. A full lithium-ion cell with a
LiNi.sub.0.5Mn.sub.1.5O.sub.4 cathode provided an average 2.8 V
which shows near theoretical reversible capacity of 310 mAh/g,
corresponding to a Li.sub.1.84Ti.sub.2O.sub.4 stoichiometry. The
observed capacity is the highest among all reported values for any
form of TiO.sub.2 anodes to date. In addition, in the same manner
we obtained the first reversible metal oxide anode for sodium-ion
batteries that also improves with cycling. The simplicity of the
electrode design and the use of fast electrochemical cycling foster
the ability for these materials to maximize their capacity in
operando, opening a new avenue for synthesis of safe and durable
high power/high capacity batteries.
[0023] The vanadium oxide embodiments of this invention are
particularly suited for use in sodium-ion cells and batteries. The
approach to achieving sodium intercalation is to use materials that
have adjustable d-spacing and two-dimensional layered structure
that can accommodate large volume changes. Furthermore, to
facilitate reversible insertion of sodium ions, a host lattice that
has short range order would be preferred to crystalline structures
due to the conservation of low-entropic energy associated with
ordering of intercalated atoms. The present invention provides,
inter alia, electrochemically grown nanostructured bilayered
vanadium pentoxide as a highly efficient 3 V cathode material for
rechargeable sodium batteries. With capacity of 250 mAh/g,
excellent rate capability and life cycle, as well as high energy
density of 760 Wh/kg, this material can be used in advanced energy
storage applications.
BRIEF DESCRIPTION OF THE DRAWINGS
[0024] FIG. 1. Electrochemical characterization of TiO.sub.2NTs in
the Li system. (A) Charge/discharge galvanostatic curves of
amorphous 65 nm outer diameter (O.D.) TiO.sub.2NT and
microcrystalline TiO.sub.2 conventional electrodes mixed with
binder and carbon black cycled between 2.5 and 0.9 V vs.
Li/Li.sup.+ at 0.05 A/g (approximately C/5, discharge the electrode
in 5 hours). The process of Li-ion coupled electron diffusion in NT
during charge and discharge is presented schematically where Ti
octahedra are light gray and Li octahedra after intercalation are
dark gray or black. Charge/discharge profiles of amorphous
TiO.sub.2NT (B) and anatase TiO.sub.2NT (D) over a wide range of
current densities cycled between 2.5 and 0.9 V vs. Li/Li.sup.+.
Note in amorphous TiO.sub.2NT (B) 17 A/g corresponds to
approximately 140C (discharge the electrode in ( 1/140) h) and 0.05
A/g correspond to C/5 (discharge the electrode in 5 h). (C)
Self-improving of amorphous TiO.sub.2NT's capacity with cycling at
high rate (7 A/g, approximately 32C). Inset shows the crystalline
structure obtained after cycling.
[0025] FIG. 2. Structural characterization of electrodes from
amorphous TiO.sub.2NTs at different stages of cycling in Li cells.
(A) Ex situ Synchrotron XRD measurements; at the stages of
discharged to 1.25 V vs. Li/Li.sup.+, discharged to 0.9 V, followed
by charged to 2.5 V. Simulation of the XRD pattern with fcc
structure (227 space group) is in agreement with the experimental
pattern and demonstrates the change of the volume of the
crystalline unit cell upon cycling, showing that the volume
decreases for approximately 17 .ANG..sup.3 as Li ions are
reversibly removed from the lattice. All measurements were
performed on 5 .mu.m-long TiO.sub.2NT with 65 nm O.D. and 10-nm
wall thickness. Simulation of XRD assumed the same sites occupied
by Li and Ti atoms with 0.5 site occupancy Inset: pre-edge feature
in the Ti K-edge X-ray Near Edge Structure (XANES) for the same
samples. T.sub.2O.sub.3 XANES spectrum is shown for reference. SEM
and TEM images of charged TiO.sub.2NTs. (B) SEM image of charged
TiO.sub.2NTs with 65 O.D. and 10-nm wall thickness after cycling.
(C) individual nanotube after cycling having 10-nm wall thickness
shown on HRTEM image. (D) HRTEM image at a region of the same tube
shows fringes corresponding to a lattice spacing of 2.05 .ANG.. (E)
Selected area electron diffraction (SAED) confirms fcc structure
and high crystallinity of TiO.sub.2NTs after phase transformation
upon electrochemical cycling.
[0026] FIG. 3. Comparison of capacities of TiO.sub.2NT anodes in
half (right) and full cell (left) using Li and Na transporting
ions. (A) Capacity measured in terms of total number of stored
electrons represented by Ti oxidation states in different
charged/discharged states obtained from Ti K-edge XANES. The edge
position of each sample is determined from the intercept of the
main edge and pre-edge contributions. (B) Charge/discharge voltage
profile of the TiO.sub.2NT-LiNi.sub.0.5Mn.sub.1.5O.sub.4 battery at
ambient temperature at a rate of C/15 (C) Charge/discharge
galvanostatic curves of amorphous 120 nm O.D. TiO.sub.2NT in Na
half cell (red for discharge and black for charge) cycled between
2.5 and 0.9 V vs. Na/Na.sup.+ at 0.05 A/g (C/3, discharge the
electrode in 3 h). (D) Charge/discharge voltage profile of the
TiO.sub.2NT-Na.sub.1.0Li.sub.0.2Ni.sub.0.25Mn.sub.0.75O.sub.6
battery at ambient temperature and various rates.
[0027] FIG. 4. SEM images of TiO.sub.2NT electrodes: amorphous
TiO.sub.2NT (A) before, and (B) after cycling; annealed TiO.sub.2NT
(C) before, and (D) after cycling. Inset in (A): cross-sectional
image of TiO.sub.2NT.
[0028] FIG. 5. SEM images of TiO.sub.2NT electrodes: amorphous
TiO.sub.2NT (A) before, and (B) after cycling in Na system.
[0029] FIG. 6. HRTEM images of TiO.sub.2NT electrodes: (A)
amorphous TiO.sub.2NT before cycling, inset: area showing area from
the same TiO.sub.2NT after electron beam irradiation for 10 s. (B)
SAED pattern from the area circled in the inset of (A), (C) HRTEM
image of amorphous TiO.sub.2NT cycled with Li. (D) SAED pattern
from the cycled TiO.sub.2NT.
[0030] FIG. 7. Cycling performance of amorphous-to-crystalline
TiO.sub.2NT electrode. The TiO.sub.2NT electrode had been cycled at
a wide range of current densities prior to the cycle life test.
[0031] FIG. 8. DFT based geometric optimization shows that the
self-improved cubic phase is thermodynamically stable with an
optimized lattice constant approximately 8.27 .ANG., which is in
good agreement with the experimental observation. (B) The radial
distribution function showing the onset of localized ordering with
increasing Li loading.
[0032] FIG. 9. Arrhenius plot for Li ion diffusion in cubic
structure constructed from the results of 1 ns MD simulations.
[0033] FIG. 10. Time evolution of the configurational energy of the
100% lithiated TiO.sub.2 system.
[0034] FIG. 11. b-values for TiO.sub.2NT electrodes as a function
of potential for cathodic sweep (transporting ion insertion): (A)
Li and (B) Na.
[0035] FIG. 12. EPR spectra: (A) Discharged sample has three
features with g values 2.004, 1.980 and the strongest signal at
1.927. The resonance response of the discharged sample is power
dependent. The sample was then exposed to air in order to remove
excess electrons and measured under the same conditions. The
surface under the integrated spectrum of oxidized sample is 27
times smaller than that of discharged sample. Weak pitch spectrum
is shown in blue for the reference. (B) The sample warmed at 77 K
shows very weak EPR spectrum with the resonance at 1.928 shifted
towards the g-tensor of free electron to g=1.970. Upon cooling the
resonance shifts back to its original value indicating localization
of charges at liquid helium temperatures.
[0036] FIG. 13. Synchrotron X-ray diffraction, Scanning Electron
Microscopy and molecular simulations of electrodeposited vanadium
oxide: (A) bilayered V.sub.2O.sub.5 annealed in vacuum at
120.degree. C., and (B) orthorhombic V.sub.2O.sub.5 annealed in
oxygen at 500.degree. C. In each case V.sub.2O.sub.5 was deposited
at an anodic current density of 5 mA/cm.sup.2. Bottom line-spectrum
represents standard XRD of orthorhombic V.sub.2O.sub.5, JCPDS card
#001-0359.
[0037] FIG. 14. (A) First five charge-discharge cycles of bilayered
V.sub.2O.sub.5 and orthorhombic V.sub.2O.sub.5 electrodes. Both
cells were cycled at 20 mA/g, within the potential window of
3.8-1.5 V (vs. Na/Na+) from 1 M NaClO.sub.4 in PC. (B) Ex situ
synchrotron x-ray diffraction patterns of orthorhombic
V.sub.2O.sub.5 before and after cycling with Na-ions: after 10 and
after 83 cycles. All films were deposited at an anodic current
density of 5 mA/cm.sup.2. (C) Capacity retention and Coulombic
efficiency of bilayered V.sub.2O.sub.5 electrodes after first 350
cycles. (D) Cycle performance comparison of same electrochemically
grown bilayered V.sub.2O.sub.5 film: deposited on Ni substrate and
pressed on stainless steel mesh current collector. Both cells were
cycled at 630 mA/g, within the potential window of 3.8-1.5 V (vs.
Na/Na+) from 1 M NaClO.sub.4 in propylene carbonate (PC).
[0038] FIG. 15. (A) Normalized V K-edge XANES for VO.sub.2
(V.sup.4+ standard), bilayered V.sub.2O.sub.5 and orthorhombic
V.sub.2O.sub.5(V.sup.5+ standard) electrodes. (B) Phase-uncorrected
Fourier transforms of V K-edge EXAFS (k.sup.3-weighted) for
bilayered V.sub.2O.sub.5 and orthorhombic V.sub.2O.sub.5
electrodes. (C) XPS spectra of bilayered V.sub.2O.sub.5, before
(top) and after 10 cycles of charging with Na+-ions (bottom) in the
V 2p.sub.3/2 and O 1s core level regions.
[0039] FIG. 16. SAXS and WAXS spectra for bilayered V.sub.2O.sub.5:
electrochemically deposited vacuum annealed sample, after
discharging with the current of 630 .mu.A (black), 120 .mu.A (gray)
or 20 .mu.A (light gray), as well as after cycling at 120 .mu.A in
charge state. Model structures and critical interlayer spacing
depicting transformations occurring upon Na+ intercalation and
deintercalation are also shown.
[0040] FIG. 17. Surface morphology of electrolytic vanadium oxide
film after annealing in vacuum at 120.degree. C. and oxygen at
500.degree. C.: before cycling (a, c) and after cycling with
Na.sup.+ ions (b, d) respectively. All films were deposited at an
anodic current density of 5 mA/cm.sup.2.
[0041] FIG. 18. Model of bilayered structure using monoclinic C 2/m
crystal symmetry with 11 nm domains of crystallinity with lattice
parameters a=11.68 .ANG., b=3.64 .ANG. and c=13.69 .ANG. and its
corresponding simulated XRD pattern.
[0042] FIG. 19. XPS spectra of bilayered V.sub.2O.sub.5, before
(top) and after 10 cycles of charging with Na.sup.+ ions (bottom)
in the V 2p and O 1s core level regions.
[0043] FIG. 20. Synchrotron X-ray diffraction spectra of bilayered
V.sub.2O.sub.5 annealed in vacuum at 120.degree. C., before and
after Na.sup.+ ion intercalation. In each case V.sub.2O.sub.5 was
deposited at an anodic current density of 5 mA/cm.sup.2.
DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS
[0044] The present invention provides new, more efficient electrode
materials for lithium-ion and sodium ion batteries. In one
embodiment, the present invention provides anodes for lithium-ion
or sodium-ion batteries comprising a nanostructured amorphous
titanium or vanadium oxide on a metallic substrate.
[0045] In one preferred embodiment the electrode comprises titanium
dioxide nanotubes on a titanium metal (e.g., titanium foil)
support. In another preferred embodiment, the electrode material
comprises a bilayered V2O5 film on a metal (e.g., nickel)
substrate.
[0046] Certain aspects and features of the present invention are
illustrated by the following discussion, which is not to be
considered as limiting.
Titanium Oxide Materials.
[0047] Herein, we explore a new approach of creating high
capacity/high power electrodes starting from an amorphous,
disordered nanoscale material. Recently, it has been shown that
increased concentration of interfacial regions in amorphous
materials may form percolation pathways for fast diffusion of ions.
Our approach to forming new improved anodes is to take advantage of
enhanced diffusion in amorphous materials. This allows for
accumulation of a high local concentration of intercalated Li-ions
at the reduced transition metal centers (i.e. Ti.sup.3+). Ordering
of densely packed materials henceforth would result in the
formation of a crystalline structure that accommodates the highest
concentration of Li ions, ensuring the highest possible electrode
capacity.
[0048] We synthesized densely packed, vertically oriented amorphous
TiO.sub.2 nanotubes (TiO.sub.2NTs) using electrochemical
anodization of Ti foil. Electrochemical synthesis is a method of
choice for preparation of nanoscale architectures that require
electronic conductivity, and it eliminates the need for conductive
carbon additives and binders typically used in electrodes that can
alter their long-term stability. As the potential drives the
electrons that are used to create reactive species, the chemical
reaction propagates to the extent that electrons can penetrate the
structure, fostering electrical connectivity. The ease of
controlling the kinetic parameters of synthesis by simple
manipulation of voltage and current density offers a great
advantage in controlling size and shape of formed nanostructures.
In addition, nanostructures have direct and electronic connection
to a metallic current collector (e.g. Ti foil). As-synthesized
TiO.sub.2NTs are amorphous, as determined by X-ray diffraction
(XRD).
[0049] When amorphous TiO.sub.2NTs were used as electrodes in a
lithium half-cell, a monotonic linear decreasing voltage is
observed during the first discharge (Li.sup.+ insertion), followed
by a characteristic plateau at approximately 1.1 V vs. Li/Li.sup.+.
The existence of the plateau indicates a phase transition, implying
that injected electrons do not contribute to decrease of the
potential difference between electrodes but instead are consumed in
the change of the structure of the materials with enhanced capacity
for Li-ions (FIG. 1A). The phase transition to a new structure
improves the capacity by an additional 125 mAh/g and occurs at a
slightly lower potential than the one when transformation began.
However, unlike a standard two-phase equilibrium phase
transformation, after the first discharge, the shoulder
characteristic of phase transition vanishes and a smooth sloping
voltage/capacity curve indicative of solid-solution type
intercalation develops. The capacity of the first charge (Li.sup.+
deinsertion) corresponds to the theoretical capacity of 335 mAh/g,
while an extraordinary high capacity of 520 mAh/g on the first
discharge, previously observed in nanoparticulate TiO.sub.2
electrodes, is a consequence of the consumption of Li atoms in
reduction of surface OH groups and residual surface H.sub.2O after
vacuum annealing at 110.degree. C.
[0050] Further charge/discharge profiles do not show evidence of
capacity loss or Coulombic inefficiency, and Li ions reversibly
intercalate/deintercalate into TiO.sub.2NTs (FIG. 7). The specific
reversible capacity following phase transformation becomes as high
as 271 mAh/g at a rate of C/5 (C/n, discharge the sample in n
hours) in a half cell (FIG. 1A), which is significantly larger than
the highest capacities (170 mAh/g at room temperature) observed in
previously investigated polymorphs of TiO.sub.2 such as anatase or
rutile. This value is also higher than the expected theoretical
value of 250 mAh/g for the highest capacity polymorphs
corresponding to the stoichiometry of Li.sub.0.75TiO.sub.2. Control
experiments with conventional powdered anatase TiO.sub.2 electrode
also confirms that the TiO.sub.2NT electrode shows a surprisingly
higher capacity (approximately 10 times) than the traditionally
used microcrystalline electrodes assembled using electroactive
binders and carbon black diluent (FIG. 1A) of comparable capacity
reported in the literature.
[0051] It is surprising that the amorphous metal oxide electrode
materials of the present invention exhibit high capacity on
cycling, when conventional wisdom indicates that their collective
electron orbitals are not expected to participate in conduction of
charge within the electrode material. It is generally accepted that
a large fraction of amorphous materials usually possess "dead"
volume that adds weight and lowers the energy density of the
battery.
[0052] Our observation of high reversible capacity in amorphous
nanotube electrodes that exceeds the theoretically predicted values
for any of the TiO.sub.2 polymorphs likely is the consequence of a
different intercalation mechanism, different structure, or the
presence of different intercalation sites in nanotube electrodes
compared to traditional bulk electrodes. The rate capability study
confirms superior diffusion of Li ions in these electrodes compared
to anatase TiO.sub.2NT electrode obtained by annealing the same
TiO.sub.2NT at 450.degree. C. (FIG. 1B, D). This was especially
apparent at high rate (17 A/g) when the high capacity of 135 mAh/g
was preserved in charging amorphous TiO.sub.2NT within 26 seconds
compared to the anatase NTs, having only 38 mAh/g. This capacity is
comparable to the theoretical capacity (175 mA/g) of
Li.sub.4Ti.sub.5O.sub.12, which is considered as a promising safe
alternative to graphite anode. Moreover, the capacity of the
amorphous sample self-improved with cycling at fast rate
(approximately 32C) and its Columbic efficiency rose >99% even
after fast charge-discharge cycles (FIG. 1C), while our similar
experiments on anatase NT electrodes show a slight decrease of
capacity in fast charge-discharge cycles. These results suggest
that cycling of Li.sup.+ into and out of the amorphous TiO.sub.2NT
material initiates ordering of TiO.sub.2 into a new type of
material that is much better suited for electron conduction and
intercalation of Li ions. This new phenomenon of self-improving
amorphous materials can be a very attractive approach for the
preparation of high performance electrode materials.
[0053] We used a combination of experimental and computational
techniques to understand the atomic-scale features of this
transformation. XRD synchrotron measurements were employed to
investigate phase transition of the amorphous TiO.sub.2NTs during
the process of cycling. Due to the high brilliance and the phase
contrast, synchrotron XRD offers exceptionally high resolution in
reciprocal space and sufficient sensitivity to probe even single
layer atoms and therefore is ideal for investigation of nanoscale
materials. We found that cycling amorphous TiO.sub.2NTs repeatedly
above approximately 1.1 V, (the potential before phase transition
occurs) produces materials with no significant improvement of
crystallinity (FIG. 2A). However, cycling at potentials below the
shoulder indicative of a phase transition results in the formation
of a new crystalline material with high degree of symmetry. The
spectrum can be fitted remarkably well with the face-centered-cubic
(fcc) Li.sub.2Ti.sub.2O.sub.4 crystalline phase (Fd3 m, space group
227) with lattice parameters that depend on the Li content. In this
phase Li and Ti are randomly distributed among all octahedral sites
in a nearly ideal cubic closed packed oxygen array that was
previously obtained in a bulk form upon chemical treatment with
n-butyl-lithium of the spinel LiTi.sub.2O.sub.4 phase. It should be
noted that while XRD indicates development of a long-range order
immediately after first discharge, a pre-edge feature in the Ti
K-edge X-ray Near Edge Structure (XANES) (symmetry forbidden Is-3d
transition in Ti that gains intensity with decreasing
centro-symmetry) shows that the structure does not develop
short-range order expected for fully ordered octahedral system
(FIG. 2A, inset). While the oxidation state of Ti in lithiated
samples approaches that of Ti.sub.2O.sub.3, the short range order
does not evolve. High-resolution transmission electron microscopy
(HRTEM), and selected-area electron diffraction (SAED) images on
the charged sample confirm long-range order. The HRTEM image (FIG.
2D) revealed a well-defined 2.05 .ANG. crystal lattice spacing,
which corresponds to the newly emerged (400) plane of the charged
sample shown in the XRD spectrum. A SAED pattern viewed down the
[001] zone axis on the tubes also indicated (400) and (440) crystal
planes (FIG. 2E) indicative of a cubic phase.
[0054] The cell parameter of the discharged unit cell of 8.27
.ANG., is somewhat smaller than the one previously reported for
chemically prepared bulk Li.sub.2Ti.sub.2O.sub.4 (a=8.376 .ANG.) of
the same symmetry, suggesting a good balance of electrostatic
forces. Our geometric optimization using density functional theory
(DFT) calculations gives an optimized lattice constant of 8.27
.ANG. (FIG. 8A), in agreement with the experimental observation.
The size of the unit cell of the charged sample is significantly
smaller than that of the spinel phase LiTi.sub.2O.sub.4 (8.403
.ANG.), suggesting that during de-intercalation, Li ions do not
occupy tetrahedral sites that are generally taken in the spinel
phase. Moreover, investigation of the samples at different stages
of discharge shows a linear decrease of the crystal unit cell with
decreasing content of Li, suggesting that there is no transition
through a spinel phase, whereby Ti and Li form a solid solution
between charged and discharged compositions. The small 3% volume
change between discharged and charged state does not seem to affect
morphology and long-term stability of TiO.sub.2NTs, and the sample
maintains the same capacity and unperturbed morphology (FIG. 4, 7)
for >600 cycles (the extent of our measurement). Indeed, the
continuous change of the lattice parameters and the sloping shape
of the voltage curves after phase transformation are typical
features of a solid-solution intercalation mechanism, leading to
stable and reversible rechargeable batteries.
[0055] Molecular dynamics (MD) simulations suggest that atomic
diffusivity in the lithiated amorphous TiO.sub.2 structure depends
strongly on the lithium content. The calculated diffusivity of Ti
and O ions for 25% Li loading was significantly lower than that for
75% Li loading in the case of amorphous TiO.sub.2. This kinetic
fluidity in amorphous TiO.sub.2 suggests that crystallization to
the thermodynamically stable cubic phase of similar composition
should be possible during initial discharge at longer timescales
than are accessible to MD simulations. Once the cubic structure
forms, atomic motion in the case of Ti and O is confined to
localized vibrations, but the lithium mobility is very high.
[0056] We found that diffusion coefficients of Ti and O in
amorphous TiO.sub.2 having low concentrations of Li are fairly
small, comparable to those in other titania polymorphs. This
suggests that introducing low amount of Li should not produce any
large-scale phase transformations. However, at high lithiation
level (>75%), atomic diffusion of Ti and O becomes significant,
suggesting facile rearrangement of atoms in the structure and
leading to phase transformation. This rapid kinetic diffusion, in
turn, suggests that crystallization to thermodynamically stable
phases of similar composition should be possible during initial
discharge of the TiO.sub.2 electrodes. Our analysis of the
formation energies also indicates that the cubic structure is
energetically more favorable than amorphous titania with the same
lithium loading (FIG. 10).
[0057] In order to further understand the evolution of crystalline
structure we evaluate the dynamics of the phase transformation from
amorphous to crystalline material occurring at high Li loadings, by
performing MD simulations for the 100% Li loaded amorphous
TiO.sub.2 composition. These results indicate evolution of
long-range order in highly lithiated TiO.sub.2, as revealed by the
formation of regularly ordered layers of oxygen. These well-defined
layers of oxygen are separated by layers of mixed metal (Ti and Li)
atoms that do not show short-range order but are randomly
distributed between ordered oxygen arrays. We noted that this
layered formation resembles the cubic phase observed in the
experiments. The entire structure has long-range order while within
the layers, disorder exists. Further simulations performed up to 6
nanoseconds (ns) do not show increase in short-range ordering,
suggesting that the structure should possess stable long-range
order but that short-range ordering is kinetically inhibited. It
should be noted that the MD simulations also indicated that for Li
concentrations much lower than 100%, the lithiated amorphous oxide
maintains its disordered structure for the same time and
temperature conditions (1400 and 1200 K), indicating that the
kinetics of this transformation is strongly correlated to the
Li:TiO.sub.2 stoichiometry.
[0058] We believe that existence of long-range order in the absence
of short-range order within layers (chemical disorder) is an
important factor in the thermodynamic stability of the de-lithiated
structure. The fact that, in this structure, all layers continue to
contain metal atoms, even in the charged state, provides stability
and prevents collapsing of the delithiated structure. Accordingly,
our DFT calculations also show that the optimized structure for the
fully de-lithiated titania preserves cubic symmetry.
[0059] Interestingly, previous reports of bulk fcc
Li.sub.2Ti.sub.2O.sub.4 suggest that this structure cannot be
employed for reversible cycling. Chemically synthesized bulk cubic
Li.sub.2Ti.sub.2O.sub.4 can be oxidized to Li.sub.0.1TiO.sub.2, but
removal of Li results in a defective disordered cubic structure
that is virtually unreactive for further Li intercalation. On the
other hand, nanoscale architectures obtained in this work by
self-assembly of amorphous samples show entirely reversible Li
insertion/extraction, demonstrating the existence of a solid
solution mechanism and good coulombic efficiencies, and confirming
facile diffusion of Li ions (FIG. 1A).
[0060] To understand how the diffusion characteristics of lithium
in the self-improved cubic structure compare with those in other
titania structures, MD simulations of Li ion diffusion in a variety
of highly lithiated crystalline and amorphous titania structures
were performed, including amorphous and cubic
Li.sub.2Ti.sub.2O.sub.4, as well as anatase TiO.sub.2. Comparison
of the trajectories suggested that Li ion transport in fully
lithiated samples follows cubic (100% Li)>amorphous (100%
Li)>anatase (50% Li), which is consistent with the high
reversible capacity observed for the self-improved structure as
compared to other titania polymorphs. In fact, calculation of Li
ion diffusion in the cubic Li.sub.2Ti.sub.2O.sub.4 reveals a
remarkably low diffusion activation barrier (257 meV--FIG. 9)
implying that rapid lithiation and de-lithiation of the cubic 227
structures should be possible during successive discharge/charge
cycles. This finding supports our XANES measurements that confirm
commendable discharging/charging even at fast cycling rate (FIG.
3A). We note that the diffusion mechanism for Li ion transport can
be significantly different for each of the above oxide structures.
The diffusion mechanism in anatase is vacancy driven and occurs via
zigzag hops between the octahedral sites, while Li ions bond with
both Ti and O, and create a Li--Ti--O network in the amorphous
TiO.sub.2. The Li diffusion mechanism in the cubic structure that
is identified from the MD simulations proceeds by simultaneous
crossing of Li ions between Li-rich and Ti-rich rows in the cubic
227 crystal; suggesting that this elementary diffusion pathway is
active during charge/discharge cycles. We note that these types of
Li exchange processes are associated with low diffusion barriers,
and are not possible in other crystalline titania polymorphs that
lack this local geometry.
[0061] To demonstrate the practical significance of this material
in batteries, the TiO.sub.2NT (cubic form) anode was coupled with a
5 V spinel (LiNi.sub.0.5Mn.sub.1.5O.sub.4) cathode for testing in a
Li-ion full cell configuration. The 5 V spinel was chosen as the
cathode in order to maximize the cell voltage, and the cell was
constructed such that the specific capacity is limited by the mass
of anode material. The TiO.sub.2NT/LiNi.sub.0.5Mn.sub.1.5O.sub.4
cell shows an average cell voltage of 2.8 V and exceptionally good
specific capacity that improves as the cycling proceeds. The
success of this Li-ion cell made by TiO.sub.2NT anode demonstrates
that it maintained the advantage of a Li-ion battery (large output
potential) while it moved the operating potential at the anode to a
higher window, thus avoiding possible safety hazards from Li
plating. The discharge capacity of the cell improves with cycles
and reaches approximately 310 mAh/g in less than 25 cycles (C/15
rate) (FIG. 3B), which is higher than any reported capacity based
on TiO.sub.2.
[0062] The success of the synthetic approach of this invention for
Li suggests that the use of amorphous electrochemically synthesized
1D nanostructures for creating rechargeable batteries that utilize
transporting ions other than Li, for example Na ions, thus
presenting an unique opportunity to generate new classes of battery
materials. Sodium is a cheap, nontoxic and abundant element that is
uniformly distributed around the world and therefore would be ideal
as transporting ion for alternative rechargeable batteries.
However, to date, no low-voltage metal oxide anodes capable of
operating with sodium ion at room temperature have been reported.
The reason for this is likely the prohibitively large ionic radius
of Na ion (1.02 .ANG.) compared to the size of Li ion (0.76 .ANG.),
and therefore insertion of Na ion requires large distortion of the
metal oxide lattice which would require unacceptably elevated
temperatures not realistic for operation of batteries.
[0063] We applied the same method as for Li, i.e., the in situ
electrochemically driven self-assembling approach at room
temperature, in the presence of large concentration of Na ion. We
find that, under electrochemical cycling conditions, narrow
amorphous TiO.sub.2NTs do not incorporate sodium (e.g. 65 nm O.D.
with 10 nm wall thickness). However, when larger tubes having 120
nm O.D. and 20 nm wall thickness are used as a host material, we
again observe self-improving phenomenon, but in this case upon
cycling at a slow rate (0.05 A/g, FIG. 3C). We found that in the
first charge the capacity reaches 80 mAh/g at 0.9 V vs.
Na/Na.sup.+. Each subsequent cycle improves the capacity, and the
average of the voltage/capacity curve and Columbic efficiency
increased upon each cycle. However, the shoulder indicative of
phase transformation during the first discharge cycle is not
observed even when the samples were cycled to more negative
potentials (<0.5 V). After 15 cycles the capacity reaches 150
mAh/g, corresponding to a stoichiometry of 0.6 Na/Ti.
[0064] The question that naturally arises is why Na ions do not
insert into the smaller (65 nm O.D.) tube while Li ions do, and why
the capacity in Na cells self-improves only under slow rate cycling
while Li cell improves upon fast cycling. One of the possible
reasons is relatively strong adsorption of Li ions on the TiO.sub.2
surface in aprotic electrolyte solutions, while no adsorption was
found for Na ions under the same conditions. This in turn enables a
high local concentration of Li ions even in a small NT diameter,
while Na ions remain mainly outside the NT in the bulk electrolyte.
Upon injecting electrons into NT, Li ions from the surface are able
to insert rapidly into the structure and induce phase
transformation to high capacity structures that improves the
performance of the cell even at high cycling rate (approximately 7
A/g). On the other hand, Na ions first need to diffuse through the
electrolyte, adsorb on the surface of TiO.sub.2NTs, and then insert
into the TiO.sub.2NTs host material. In the larger NTs, however, a
larger amount of electrolyte (and therefore Na ions) are trapped
within the NT electrode compartment, enabling larger flux of Na
ions upon discharge. As accumulation of Na ions at the electrode
surface is slower compared to Li ions, capacity self-improvement of
the TiO.sub.2NT should occur at a slower rate. Indeed we observe Na
insertion at approximately 140 times slower rate of cycling
(approximately 0.05 A/g) compared to Li ion. The same arguments
also account for the different mechanism of cycling in Li and Na
cells. The Li-ion cells operate via a diffusion mechanism at high
potentials (2.4-1.8 V) where a small driving force is sufficient to
insert adsorbed ions into the TiO.sub.2 matrix. On the other hand,
Na-ion electrodes at high potentials operate under a
pseudocapacitive mechanism. At relatively negative potentials, a
pseudocapacitative mechanism converts into diffusion mechanism
(FIG. 11), indicating that critical concentration of Na ions in the
vicinity of TiO.sub.2 surface has to be reached in order for the
insertion to begin.
[0065] Surprisingly, the TiO.sub.2NT materials also are suitable as
anode materials for a Na-ion battery that can operate at room
temperature. The Na-ion cell was made by a TiO.sub.2NT anode and a
Na.sub.1.0Li.sub.0.2Ni.sub.0.25Mn.sub.0.75O.sub.6 cathode. The cell
shows an operation voltage of approximately 1.8 V and a discharge
capacity of approximately 80 mAh/g (C/8) (FIG. 3D). The voltage is
smaller compared to a Li-ion battery, but it is still higher than
1.2 V, the operation voltage of the nickel-metal hydride battery
used as the power source for hybrid vehicles like Toyota Prius. In
addition, this Na-ion cell comprised of a TiO.sub.2NT anode shows
excellent rate capability with around 70% low-rate capacity
retained at 11C. This Na-ion oxide battery therefore indeed holds
promise for the further development of new ambient temperature
Na-ion battery systems that combine novel electrodes to form full
energy storage devices that are inexpensive with good
performance.
[0066] It is at the nanoscale that near theoretical capacity and
high power electrodes can be achieved using simple
self-organization processes. The electrostatic attraction of
electrochemically-altered materials provides a strong driving-force
for the diffusion of a large concentration of transporting ions
into amorphous metal oxide frameworks. This consequently leads to
ordering of the atomic building blocks, transporting ions and host
metal octahedra into a crystalline array. Inducing crystallization
of nanomaterials in operando (i.e., in situ in a battery during
charging and discharging) allows realization of the highest
possible electrode capacity by optimizing the balance of
electrostatic forces. The present invention demonstrates that much
higher capacities can be realized if the system is naturally
allowed to choose and optimize its crystalline structure through a
process of self-organization and self-improvement. The small
diffusion length and large surface area of nanostructures also
enable exceptionally fast charging leading to high power batteries.
Electrochemically induced structural evolution into high
capacity/high power electrodes provides a powerful modular approach
to the design of improved battery materials with programmable
physical and chemical properties.
Synthesis of TiO.sub.2 Nanotube Electrodes.
[0067] TiO.sub.2 nanotubes were synthesized by electrochemical
anodization. Pure titanium thin foil (0.0127 mm, 99.8%, Alfa Aesar)
was cleaned by acetone following an isopropyl alcohol rinse before
anodization. The back of the Ti foil was protected by nail polish
to ensure uniform current distribution. The anodization was carried
out in a two-electrode cell with Ti metal as the working electrode
and a Pt mesh as the counter electrode under constant potentials
(15-30 V) at room temperature using electrolytes of formamide with
0.8 wt % ammonium fluoride (Aldrich) and 5 vol % water. The
as-anodized samples were ultrasonically cleaned in deionized (D.I.)
water for 30 seconds. All amorphous TiO.sub.2NT samples were vacuum
baked at 110.degree. C. overnight before assembly in
electrochemical cells. The crystallized TiO.sub.2NT samples were
prepared by annealing as-prepared TiO.sub.2NT under O.sub.2 at
450.degree. C. for 4 hours.
Preparation of Positive Electrodes.
[0068] The LiNi.sub.0.5Mn.sub.1.5O.sub.4 spinel and
Na.sub.1.0Li.sub.0.2Ni.sub.0.25Mn.sub.0.75O.sub..delta.-layered
cathode materials were prepared by solid state reactions as
follows: quantities of Na.sub.2CO.sub.3 and/or Li.sub.2CO.sub.3,
and Ni.sub.0.25Mn.sub.0.75CO.sub.3 (prepared by co-precipitation)
to achieve the desired stoichiometry were ground together in an
agate mortar until visually homogenous. The mixture was initially
fired in a muffle furnace at 550.degree. C. for 12 h, reground, and
then calcined at 850.degree. C. for 12 h and allowed to slowly cool
to ambient temperature. Structures were verified by powder XRD and
the relative metals contents were measured by ICP-OES and found to
be the target compositions within the error of the determination.
Electrode films were cast from slurries in N-methyl-2-pyrrolidone
containing the active material at a concentration of about 82
percent by weight (wt %), carbon black (4 wt %, CABOT XC72),
graphite (4 wt %, TIMCAL SFG-6), and poly(vinylidene difluoride)
(10 wt %; KYNAR KF 1120) using doctor blade or spin casting
techniques.
Electrochemical Insertion/Extraction of Li.sup.+ and Na.sup.+.
[0069] Li half-cells were assembled in coin-type cells (HOHSEN
2032) with a Li metal foil as the negative electrode, microporous
polyolefin separators (CELGARD 2325), and 1.2 M LiPF.sub.6 in
ethylene carbonate/ethyl methyl carbonate (3:7 weight ratio)
electrolyte (Tomiyama). Na half-cells were assembled in similar
setup with a Na metal foil counter electrode, glass fiber separator
(WHATMAN GF/F), and 1 M NaClO.sub.4 (Aldrich) in propylene
carbonate electrolyte. For comparison, a traditional laminate
electrode was made by mixing 84 wt % active material (anatase
powder, Aldrich), 4 wt % graphite (TIMCAL SFG-6), 8 wt %
poly(vinylidene difluoride) binder (KYNAR), and 4 wt % carbon black
(Toka). Half-cells were cycled galvanostatically at varying
currents between 2.5 and 0.9 V vs. Li/Li.sup.+ or 2.5 and 0.5 V vs.
Na/Na.sup.+ respectively using an automated Maccor battery tester
at ambient temperature. Li-ion and Na-ion full-cells were assembled
in the same manner as half-cells with a
LiNi.sub.0.5Mn.sub.1.5O.sub.4 cathode and a
Na.sub.1.0Li.sub.0.2Ni.sub.0.25Mn.sub.0.75O.sub..delta. cathodes,
respectively. Li-ion full-cells were cycled galvanostatically
between 2 to 3.5 V. Na-ion full-cells were cycled galvanostatically
between 1 to 2.6 V. Cyclic voltammograms of the cells were recorded
in a SOLARTRON 1470E Potentiostat/Galvanostat. Electrodes removed
from cells for analysis were thoroughly washed with dry dimethyl
carbonate (Aldrich) and allowed to dry under inert atmosphere. All
cell assembly and disassembly operations were performed in a
He-filled dry glove box (oxygen level <2 ppm). The net weights
of the TiO.sub.2 nanotube films were obtained by peeling off the
nanotube film from Ti substrate using adhesives and checking by SEM
to make sure no residual TiO.sub.2 nanotube is left on the
substrate.
Synchrotron XRD.
[0070] X-ray diffraction measurements were performed at beamline
13-ID-D of GSECARS at the Advanced Photon Source (APS) at Argonne
National Laboratory. The X-ray beam (37 keV energy, corresponding
to X-ray wavelength of .lamda.=0.3344 .ANG.) was focused to a 2
.mu.m diameter spot with a Kirkpatrick-Baez mirror system. The
distance and tilting of the MAR165-CCD detector were calibrated
using a CeO.sub.2 standard. Charged TiO.sub.2NT samples were
prepared by stripping nanotube films from Ti support onto KAPTON
tape and sandwiching with additional KAPTON tape. Electrochemically
lithiated TiO.sub.2NT samples were scratched-off and sealed inside
a 3 mm-diameter hole in a piece of aluminum foil by sealing KAPTON
sheet to the foil using epoxy. All procedures were carried out in a
He atmosphere glove box. Simulations of XRD patterns were carried
out using CRYSTALMAKER software (CrystalMaker Software, Ltd.).
Electron Microscopy.
[0071] Scanning Electron Microscopy (SEM) images were recorded with
a JEOL JSM-7500F Field Emission SEM operating at 30 kV. High
resolution electron transmission microscopy (HRTEM) images and
electron diffraction (ED) were obtained using an FEI Tecnai F30
microscope equipped with a field emission gun operated at 300 kV.
TEM samples were prepared by scratching cycled TiO.sub.2NT samples
from Ti support onto a carbon-coated copper TEM grid (ultrathin
carbon film on a perforated carbon support film, 400 mesh, Ted
Pella Inc.).
Molecular Dynamics (MD) and DFT Simulations.
[0072] MD simulations were performed using the DLPOLY MD package
and electronic structure calculations were within Vienna ab-initio
simulation package (VASP). The MD simulations utilized the shell
potential model in which the polarizable ion consists of two
particles--core and shell--that share the ion's charge and are
connected via a spring constant. The atoms were treated as point
particles, and interact via long-range Coulomb forces and
short-range interactions. The short-range interactions were
represented by the Buckingham potential, which is described by
parameterized functions fitted to reproduce the structure and
energetics of Li--Ti--O system. In all MD simulations, a system
size of 4320 atoms was used. The calculations spanned a temperature
range of 850 K-1400 K, and statistics typically were collected from
1 ns simulations with a time step of 0.2 femtoseconds (fs). DFT
calculations were performed using GGA-PBE. The cubic structure was
made of 64 atoms, while the amorphous cell was constructed using 96
atoms. The initial structure of the amorphous cell for DFT
calculations was obtained from the equilibrated configurations
generated from the MD simulations. At least three different
amorphous configurations were simulated to ensure that the results
were not biased by the initial starting configuration. For all the
systems, the cell shape was allowed to change, and ions were
allowed to relax. The calculations were performed using 550 eV
energy cut-off and 5.times.5.times.5 k-point sampling.
X-Ray Absorption Near Edge Structure.
[0073] X-ray spectroscopy measurements were performed at the
PNC/XSD bending magnet beamline (20-BM-B) of the APS at Argonne
National Laboratory. Measurements at the Ti K-edge were performed
under transmission mode using gas ionization chambers to monitor
the incident and transmitted X-ray intensities. A third ionization
chamber was used in conjunction with a Ti-foil standard to provide
internal calibration for the alignment of the edge positions. The
incident beam was monochromatized using a Si (111) double crystal
fixed exit monochromator. Harmonic rejection was accomplished using
a Rhodium coated mirror. The charged samples were prepared by
stripping NT films onto KAPTON tape. The discharged samples were
stripped onto KAPTON tape in a dry glove box under a helium
atmosphere and attached to a He-filled, sealed holder. The
reference standards were prepared by spreading thin, uniform layers
of powders on KAPTON tape and stacking a few layers to attain the
desired absorption step height. Each spectrum was normalized using
data processing software package IFEFFIT. Alignment of each sample
reference spectrum with respect to Ti standard spectrum was within
the range of +0.03 eV.
Electron Paramagnetic Resonance.
[0074] EPR measurements were performed on a Bruker Biospin Elexsys
E 580 at ANL. The instrument was operated at a frequency of 9.3 GHz
with a modulation frequency of 100 kHz. The field was swept at 3.5
Gauss/s over 300 to 1500 G range. The calibration of the g-values
was done with DPPH. The Li content in charged samples was
calculated by integrating the area under the EPR curve.
Structural Parameters for Different TiO.sub.2NTs
[0075] The volume of the unit cell expands slightly (approximately
1%) with cycling from 1st to 582nd cycle. The extent of lattice
expansion between charged and discharged samples was less
pronounced in the larger NTs (120 O.D. with 20 nm wall thickness)
and volumes ranged between 562 to 557 .ANG..sup.3 (Table 1).
TABLE-US-00001 TABLE 1 Size Tube Wall Unit cell length O.D.
thickness a Cell volume (.mu.m) (nm) (nm) Cycle (state) (.ANG.)
(.ANG..sup.3) 2 65 10 1 (charged) 8.19 550 2 65 10 172 (charged)
8.22 555.6 2 65 10 582 (charged) 8.22 555.6 3 120 20 172 (charged)
8.25 562 3 120 20 .sup. 172 (discharged) 8.23 557
Size Effect of TiO.sub.2NT Array Electrodes.
[0076] TiO.sub.2NT electrodes were prepared with varying size and
their specific capacities were compared in Table 2. It has been
observed that capacity decreases with the increase of the tube
length. This can be understood as being due to an increase in ohmic
resistance along the tube length, leading to decreased capacity.
The capacity of larger NTs (120 O.D.) was enhanced indicating that
the diffusion of Li ions into larger tubes was more efficient
possibly due to the larger amount of ions in larger NTs.
TABLE-US-00002 TABLE 2 Size Tube Wall length O.D. thickness
Capacity (mAh/g) (.mu.m) (nm) (nm) Amorphous Annealed 2 65 10 271
210 3 65 10 170 145 5 65 10 150 105 3 120 20 191
Characterization of Morphology.
1) SEM Characterization.
[0077] a. Li System.
[0078] SEM images of both amorphous and annealed TiO.sub.2NT
electrodes before and after electrochemical cycling in Li system
are shown in FIG. 4. The nanotube morphology remained intact,
indicating high mechanical strength of this material due to the
unique nano-size effect.
[0079] b. Na System.
[0080] FIG. 5 shows SEM images of TiO.sub.2NT electrodes before (A)
and after (B) cycling with Na. The TiO.sub.2NT material maintained
nanotube morphology after cycling. It should be noted that the
debris on the cycled tubes is from residues of electrolyte.
2) TEM Characterization.
[0081] FIG. 6A, C compares the HRTEM images of amorphous
TiO.sub.2NT electrode before and after cycling. It can be observed
that the morphology of the nanotube preserved after the electrode
was cycled at a wide range of current densities. Interestingly,
when TEM e-beam was focused on the amorphous (before cycling)
TiO.sub.2NT with the highest magnification (1M) and the smallest
size (2 nm) for 10 seconds, the TEM image and the ED pattern (FIG.
6B) showed that the nanotubes were locally crystallized to form
anatase crystal structure, which is different from the structure
after electrochemical cycling (FIG. 6D). We also carried out
post-annealing of the cycled TiO.sub.2NT at 450.degree. C. for 4 h,
after which the XRD pattern shows a mixture of anatase, rutile and
the cubic phase. This indicates that the newly formed fcc structure
of TiO.sub.2NT upon cycling was also thermodynamically stable.
Cycle Life.
[0082] The cycle life of amorphous TiO.sub.2NT sample in Li system
was evaluated by electrochemically cycling for more than 500 cycles
at approximately 0.6 A/g (2C) in a potential range of 2.5-0.9 V vs.
Li/Li.sup.+. It has been found that this material has excellent
cycling stability with approximately 95% retention of capacity and
close to 100% columbic efficiency (FIG. 7).
Charge Storage Behavior for TiO.sub.2NT: Capacitive Vs. Diffusion
Processes.
[0083] The charge storage behavior of TiO.sub.2NT electrodes in Li
and Na systems were studied using cyclic voltammetry under
different scan rates (0.1-5 mV/s) with Li or Na metal serving as
both counter and reference electrodes. There can be three
contributions to charge storage: the faradaic contribution due to
the intercalation of transporting ions (Li.sup.+, Na.sup.+);
faradaic contribution from surface processes, referred to as
pseudocapacitance; and non-faradaic contribution due to the double
layer capacitance. Contributions to the charge storage mechanisms
can be characterized by analyzing the cyclic voltammograms with
varying scan rates according to:
i=a.nu..sup.b (1)
where measured current i obeys a power law relationship with scan
rate .nu.. Both a and b are adjustable parameters and b can be
determined from plotting log(i) versus log(.nu.). For a process
limited by diffusion, b would be 0.5 according to the following
equation:
i = nFAC * D 1 / 2 v 1 / 2 ( .alpha. nF RT ) 1 / 2 .pi. 1 / 2 .chi.
( bt ) ( 2 ) ##EQU00001##
where n is number of electrons involved in the electrode reaction,
F is faraday constant, A is the surface area of the electrode
material, C* is the surface concentration, D is the diffusion
coefficient, .nu. is the scan rate, R is the gas constant, .alpha.
is the transfer coefficient, T is the temperature, and .chi.(bt) is
the normalized current for a totally irreversible system in cyclic
voltammetry. For a purely capacitive process, b is normally 1
according to the following equation:
i.sub.c=.nu.C.sub.dA (3)
where C.sub.d is the capacitance.
[0084] For charge storage in Li system with TiO.sub.2NT (FIG. 11A),
the increase of b values in the potential range from 2.3 to 1.8 V
indicates the switch of dual contributions from diffusional and
capacitive mechanisms to a capacitive process dominated the
mechanism. On the other hand, the opposite phenomenon was observed
in Na system (FIG. 11B) with capacitive-limiting mechanism changing
to mixed contributions by diffusion and surface capacitive
processes when discharging proceeds.
EPR Study.
[0085] EPR measurements (FIG. 12) were performed on a Bruker
Biospin Elexsys E 580 instrument operated at a frequency of 9.3 GHz
with a modulation frequency of 100 kHz. The field was swept at 3.5
G/s over 300 to 1500 G range. The calibration of the g-values was
done with DPPH. The microwave power shown on the spectra was 2 mW
and modulation amplitude 5 G. The Li content was calculated by
integrate the area under the curve.
Vanadium Oxide Materials.
[0086] Vanadium pentoxide was synthesized by electrochemical
deposition from the aqueous vanadyl sulfate electrolyte on a metal
foil substrate, e.g., a Ni foil substrate, and then annealed in
vacuum at 120.degree. C. to remove intercalated water.
Electrochemically synthesized electrodes offer long-range
electronic conductivity, which improves responsiveness to applied
potential, and therefore, their intercalation properties.
Utilization of electrochemical deposition also brings high level of
control to the structure, morphology, and uniformity of electrodes
by adjusting the crucial parameters such as: applied current,
potential and electric pulses, as well as the temperature and
concentration of the electrolyte. In addition, vanadium oxide
electrode is deposited directly on a current collector without the
use of electronic conductor (such as carbon black or nickel powder)
and/or a polymer binder (such as polyvinylidene difluoride).
[0087] Based on the scanning electron micrograph (FIG. 13 top
right), such V.sub.2O.sub.5 electrode is found to be composed of
nanoribbons with highly porous structure. The electrochemically
grown interconnected ribbons allow excellent electron conductivity,
while the high porosity enables efficient penetration of the
electrolyte and ensures high utilization of electrode material. The
electrode with such morphology represents an efficient matrix for
ion transport in which high surface area of electrode diminishes
limitations caused by diffusion.
[0088] The result from synchrotron X-ray diffraction (XRD) (FIG.
13A) indicates that V.sub.2O.sub.5 structure is composed of 2D
bilayered stacks indicated by narrow features in the intermediate
and low d-spacing range of the difractogram. These bilayered stacks
are separated by large interlayer spacing, which is shown as an
intense broad peak at high d-spacing (approx. 13.5 .ANG.). The
pattern contains a small number of Bragg-like features, indicating
the presence of intermediate range ordering and a pronounced
diffused component. The bilayered V.sub.2O.sub.5 is reminiscent of
the V.sub.2O.sub.5 xerogels in which monoclinic bilayers of
V.sub.2O.sub.5 stack up with the spacing that expands (or
contracts) as the xerogel incorporates (releases) water molecules.
This structure is a stacking of V.sub.2O.sub.5 bilayers made of
base-faced square pyramidal VO.sub.5 units arranged in parallel at
equidistant positions (model in FIG. 13A). The distance of closest
approach between the bilayer stacks is approximately 13.5 .ANG..
This is the most noticeable period of repetition in the structure
as manifested by the strength and position of the low-angle peak in
the XRD pattern. The width of this feature suggests that stacking
sequence is imperfect, confirming disordering in this system. The
high-angle domain region is dominated by one peak at 3.44 .ANG.,
which is also the highest intensity peak in the simulated
diffraction of monoclinic bilayered model structure and corresponds
to combined diffraction of 201 and 111 directions composed of 2.85
.ANG. apart single layers (FIG. 18).
[0089] When electrochemically grown V.sub.2O.sub.5 is annealed in
oxygen atmosphere at 500.degree. C., SEM image revealed that
nanoribbons were converted to fine rod-like shape that exhibit
orthorhombic crystalline structure (FIG. 13B). Based on diffraction
patterns and simulated model structure the orthorhombic phase is
consisted of single layers of VO.sub.5 square pyramids, while the
interlayer spacing in z direction is almost completely diminished
compared to bilayered V.sub.2O.sub.5 reaching only 4.4 .ANG..
[0090] Considering these findings, the crucial question is which
one of these structures supports better the reversible
intercalation of sodium transporting ions. While bilayered
structure is atomically ordered only in a short range, the spacing
between bilayers is more random. However, there is a lot of void
space between randomly spaced bilayers after removal of
intercalated water, which, if flexible enough, can readjust the
spacing to enable intercalation of larger Na+ ions (model, FIG.
13A). On the other hand, orthorhombic rod-like crystals have long
range order that may enable unhindered diffusion of intercalated
ions throughout the entire structure that could lead to highly
reversible capacity.
[0091] In order to evaluate functionality of these materials, each
type of structure is formed on the separate electrode and they were
submitted to cycling against sodium metal electrode (FIG. 14A).
Both electrodes exhibited observable capacities, however, the
bilayered V.sub.2O.sub.5 electrode demonstrates higher
electrochemical activity and stable reversible capacity on repeated
cycling than its orthorhombic counterpart. In the 3.8 to 1.5 V
range, we observed significantly large specific capacity of 250
mAh/g at 20 mA/g (C/8; C/n, discharge rate in n hours) for
bilayered V.sub.2O.sub.5 electrode compared to only 150 mAh/g for
orthorhombic--V.sub.2O.sub.5. The high capacity of bilayered
V.sub.2O.sub.5 is comparable to the theoretical limit of 236 mAh/g
for Na.sub.2V2O.sub.5, suggesting that this structure can
accommodate one Na+ ion for each V atom.
[0092] The cycling curve is composed of two distinct parts: the
first with a smooth slope indicative of solid solution
intercalation mechanism that reaches about 235 mAh/g (theoretical
capacity), and the second one at potentials >2.25 V when the
slope changes to a steep curve indicative of capacitative
mechanism. This small contribution of 15 mAh/g might be due to the
capacitance of the surface layer or due to the presence of small
amount of NiO formed at the interface between V.sub.2O.sub.5 and Ni
substrate during the electrochemical deposition. Moreover, the
shape of galvanostatic curves in FIG. 14A reveals entirely
different mechanism of intercalation of Na+ transporting ions in
two different V.sub.2O.sub.5 structures. Incorporation of Na+ ions
into orthorhombic electrode is accompanied by two phase
transitions, manifested by existence of two plateaus both in the
discharge and charge cycles suggesting that orthorhombic
crystalline structure changes twice to accommodate increasing Na+
concentration. On the other hand, the incorporation of Na+ ions
into bilayered structure shows smooth, solid states solution
intercalation with no phase transitions. Importantly, the capacity
of bilayered V.sub.2O.sub.5 does not change with cycling including
the first cycle. This behavior indicates that there is no side
reactions of injected electrons with either electrolyte or surface
of nanocrystalline bilayered V.sub.2O.sub.5.
[0093] In contrast, the capacity of orthorhombic electrode rapidly
decreases with cycling, as previously reported. The phase
transformation in orthorhombic crystalline electrode is less
pronounced as the cycling proceeds concomitantly with the decrease
of the electrode capacity. This behavior suggests that the change
in crystalline structure is associated with the fading of the
orthorhombic electrode performance. Indeed, examination of XRD
patterns with cycling reveals that orthorhombic V.sub.2O.sub.5
experienced deterioration in its crystallinity after first 10
cycles, which is followed by loss of crystallinity after prolonged
cycling (FIG. 14B). After 82 cycles the broadening and
disappearance of the peaks, especially in longer d-spacing range
(>4 .ANG.), shows significant reduction of the crystalline
domains from a few hundred to only few nm. Reduction of the domains
of crystallinity is accompanied with loss of the electronic
conductivity that directly causes loss of electrode capacity.
[0094] On the other hand, bilayered V.sub.2O.sub.5 electrodes were
stable to repeated cycling and during prolonged galvanostatic
cycling (up to 350 cycles) at high current density of 630 mA/g, and
the average capacity of bilayered V.sub.2O.sub.5 electrodes remains
at 85% of its initial value (FIG. 14C). Furthermore, nanostructured
bilayered V.sub.2O.sub.5 electrodes exhibited excellent discharge
capacity and cycle stability even at high-rate charge-discharge
processes, e.g., the discharge capacity at faster cycling rate of
60 mA/g decreases slightly to 200 mAh/g, while at rate of 630 mA/g
(6 min discharge) decreases to 150 mA/g.
[0095] FIG. 14D compares charge-discharge performance at fast
cycling (630 mA/g) of bilayered V.sub.2O.sub.5 prepared
electrochemically versus mechanically pressed on stainless steel
mash. Interestingly, bilayered electrode that is electrochemically
grown on Ni substrate shows improvement of the electrode capacity
with the cycling (FIG. 14D) most likely due to improved supply of
Na+ ions through nanoporous electrode. For comparison, when
bilayered V.sub.2O.sub.5 was pealed from Ni substrate, mixed with
conductive carbon additive and polymer binder and pressed on
stainless steel mash current collector initial capacity immediately
reaches maximal capacity at the cycling rate of 140 mAh/g, but
substantially decreases by prolonged cycling. This underlines
importance of superior electronic contact and excellent ionic
conductivity that is obtained by electrochemical deposition, which
are key factors in stable and reversible operation of batteries.
The monotonic slope of the bilayered V.sub.2O.sub.5 voltage
profiles confirms the absence of phase transitions during the
charge-discharge processes (FIG. 14). The results indicate that
bilayered V.sub.2O.sub.5 structure is capable of sustaining
single-phase intercalation of large Na+ ions in wide concentration
range and does not change morphology during repeated cycling. (FIG.
17) These findings suggest the advantage of non-3D crystalline
structures that are far away from the thermodynamic equilibrium for
applications in batteries operating on sodium ion exchange.
[0096] The oxidation state of the material is essential for redox
processes that are taking place in rechargeable batteries, and for
that reason we utilized X-ray absorption near edge spectra (XANES)
in the V K-edge range for the bilayered V.sub.2O.sub.5,
orthorhombic--V.sub.2O.sub.5(V.sup.5+ standard) and VO.sub.2
(V.sup.4+ standard). The obtained values were consistent with the
previous results of V.sub.2O.sub.5 compounds and showed identical
pre-edge shape and peak positions, implying that the V redox state
in the bilayered V.sub.2O.sub.5 sample is close to V.sup.5+, in a
common VO.sub.5 environment (FIG. 15A). However, the area of the
pre-edge peak for bilayered is smaller than that of orthorhombic
V.sub.2O.sub.5 electrode. These variations in the area of the
pre-edge peak indicate that the local structure of V in bilayered
V.sub.2O.sub.5 has a higher degree of local symmetry than that of V
in orthorhombic V.sub.2O.sub.5.
[0097] In addition, comparison of the V-K-edge k.sup.3 weighted
EXAFS spectra show important differences (FIG. 15B). The structure
of the first two peaks is dominated by single scattering
contributions from the first and second coordination spheres of V-O
and V-V correlations. The scattering contributions of V-O bond in
orthorhombic V.sub.2O.sub.5 is split into two components, one
shorter bond distance corresponding to distances within square
pyramidal environment, and longer V-O* distance with major
contribution from oxygen atoms from neighboring planes in
orthorhombic structure. This longer bond distance is absent in
Fourier transform spectra of bilayered structure due to the shorter
bond between V and O atoms in bisquare pyramidal arrangement.
[0098] The position and shape of the V 2p and O 1s XPS peaks before
and after cycling of bilayered V.sub.2O.sub.5 electrodes confirms
the reversibility of Na+ intercalation (FIG. 15C). The peak at
516.9 eV is assigned to the V.sup.5+ 2p.sub.3/2, while shoulder at
515.7 eV to the V.sup.4+ 2p.sub.3/2 orbital, which is additionally
confirmed by the spin-orbit splitting of about 7.5 eV between V
2p.sub.3/2 and V 2p.sub.1/2 (FIG. 19). Composition of the
electrochemically synthesized bilayered V.sub.2O.sub.5 electrode
before the intercalation of Na+ shows slightly reduced content of
V, indicated by 25% of total vanadium in V.sup.4+. However, after
82 cycles, V 2p.sub.3/2 peak shape for the sample in a charged
(oxidized) state constitutes of 95% of V.sup.5+ and 5% of V.sup.4+
state, confirming efficient reversible intercalation and
deintercalation of Na+ ions in bilayered V.sub.2O.sub.5 structure.
In the O 1s region, the main peak attributed to lattice oxygen
(O--V) is located at 531 eV, however, we observed additional
shoulder at the higher binding energy side in electrochemically
synthesized V.sub.2O.sub.5, which was previously assigned to
chemisorbed water or adsorbed carbon dioxide molecules. This side
peak shifted upon cycling suggesting surface adsorption of
electrolyte molecules that replace initially adsorbed carbon
dioxide during electrode cycling.
[0099] In order to understand mechanism of sodium
insertion/de-insertion as well as the limits of discharge capacity
and cyclability of bilayered V.sub.2O.sub.5 electrodes, it is
important to understand the response of the bilayered
V.sub.2O.sub.5 structure to the intercalation of sodium. For this
purpose, we have used X-ray scattering that is a technique of
choice for determining the changes in both short and long range
order in crystalline and non-crystalline materials. Small and Wide
angle X-ray scattering (SAXS and WAXS) measurements in situ
(non-operando) confirm the initial structure of vacuum annealed
bilayered V.sub.2O.sub.5 that was obtained from XRD measurements,
i.e., layered structure with bilayers spaced at average distance of
13.5 .ANG. apart, and the structural order within bilayers showing
characteristic 3.44 .ANG. spacing (FIG. 16). SAXS measurements
confirm the change of periodicity and the stacking order upon
intercalation and deintercalation of Na+ ions.
[0100] Upon initial discharging we observed that the interlayer
d-spacing of the (001) plane (13.5 .ANG.) and corresponding (002)
plane (6.7 .ANG.) dramatically change upon intercalation of sodium
to 16.1 .ANG. and 8.0 .ANG., respectively. Concomitantly, the width
of the layer spacing peak narrows, and several peaks in the
wide-angle region sharpen, suggesting 3D-like ordering of the
structure upon sodium intercalation. Moreover, superimposed to the
broader scattering pattern corresponding to newly developed sodium
ion assembled layered structure one can observe a set of very
narrow peaks, which originate from defined distances of
intercalated sodium atom with atoms constituting the V.sub.2O.sub.5
bilayer. The intensity of these sharp features depends on the
applied discharge currents (FIG. 16, black and gray curves). One
can observe exceptionally strong narrow peak at 4.9 .ANG. that
corresponds to the distance between intercalated Na+ ions and
neighboring oxygen atoms that terminate the bilayered structure,
which was also observed in XRD of discharged sample (FIG. 20).
[0101] Upon electrode oxidation (charging), and consequent Na+
deintercalation, this peak completely vanishes, confirming complete
removal of Na ions in agreement with XPS restoration of V.sup.5+
state. Also all of the peaks associated with layered structure
disappear and only those associated with short range order within
bilayered structure are preserved. This suggests that after
deintercalation of Na+ ions from the structure the stacking order
is removed (FIG. 16). These measurements strongly indicate that
bilayered stacking and large spacing of bilayered structure is
crucial for efficiency and stability of reversible Na+ ion
intercalation in V.sub.2O.sub.5 electrodes. These flexible
non-crystalline layers reassemble into organized structure each
time Na+ ions intercalate into electrode. Electrostatic interaction
between sodium ions and bilayer terminating oxygen atoms fixates
the stacking of bilayers. However, upon oxidation (charging) the
electron density of terminating oxygen atoms decreases, weakening
their interaction with Na+ ions that causes deintercalation of Na+,
which leaves random ordering between the layers.
[0102] The extent of ordering and corresponding intensity of
scattering features is dependent on the current density used for
intercalation of Na+ ions. Interestingly, the highest current leads
to more pronounced peaks, suggesting that higher concentration
(flux) of intercalated atoms produces better ordering of the
lattice. Moreover, the specific capacity of fast cycled electrode
improves with the cycling. The full capacity of this electrode is
established only after a number of cycles at high scanning rate,
emphasizing the importance of availability of Na+ ions on electrode
surface for obtaining theoretical capacities. Therefore, the
superior discharge capacity retention of bilayered V.sub.2O.sub.5
electrodes compared to the other nanostructured vanadium pentoxide
electrode reported in the literature may be attributed to the
combined effects of their structural and surface properties. Highly
accessible nanoscale architecture that combines conductivity with
high surface area plays crucial role for ion and electron transport
at the electrode/electrolyte interface.
[0103] It has been shown recently that the increased number of
electrode/electrolyte interfacial interactions in nanostructured
materials is critical to the formation of percolation pathways for
fast diffusion of ions. Extensive adsorption of large sodium ions
at the nanostructured electrochemical interface is important to
initiate the intercalation process. At the same time the layered
nature of the active material with low energy sites drives the Na+
diffusion into the material bulk structure. The combination of
these processes is necessary for maximizing the flow of Na+ ions
into the material. The net effect is enhanced pseudocapacitance
that maximizes the electrostatic attraction of Na+ cations into
V.sub.2O.sub.5 matrix. In addition, charge transfer ability of the
bilayered V.sub.2O.sub.5 electrode is improved with the presence of
more conductive surface defect species such as lower valence
vanadium atoms (FIG. 15C) and associated oxygen vacancies. Enhanced
charge transfer conductivity improves electron transport during
sodium ion intercalation/deintercalation at the
electrode/electrolyte interface. Undoubtedly, this is sufficient to
stimulate efficient diffusion of the large concentration of
transporting ions into metal oxide layered framework with
concomitant reduction/oxidation of V atoms.
Synthesis of nanostructured V.sub.2O.sub.5 electrodes.
[0104] Nanostructured V.sub.2O.sub.5 was synthesized by
electrochemical deposition on pure Ni foil (0.0127 mm, 99.8%),
which was cleaned in acetone and isopropyl. The electrochemical
deposition was carried out in a three-electrode cell with Ni foil
as the working electrode, Pt mesh as a counter electrode and
Ag/AgCl as a reference electrode in aqueous 0.1 M VOSO.sub.4
solution at a constant potential of 1.5 V. Bilayered V.sub.2O.sub.5
electrodes were synthesized by vacuum annealing at 120.degree. C.
for 20 hours. The crystallized orthorhombic V.sub.2O.sub.5
electrodes were obtained by annealing of as-prepared V.sub.2O.sub.5
under O.sub.2 atmosphere at 500.degree. C. for 4 hours.
Electrochemical Insertion/Extraction of Na.sup.+ and
Characterization.
[0105] Na-half cells were assembled in He-filled dry glove box into
coin-type cells with a Na foil as the negative electrode, an
electrolyte of 1 M NaClO.sub.4 in propylene carbonate (PC) and
glassy fiber separator. For comparison, traditional electrode was
made by mixing 84 wt % active material (V.sub.2O.sub.5 powder,
Aldrich), 4 wt % graphite (TIMCAL SFG-6), 8 wt % poly(vinylidene
difluoride) binder (KYNAR), and 4 wt % carbon black (Toka) and
pressed on stainless steel mesh current collector. All cells were
tested galvanostatically by automated Maccor battery tester at
ambient temperature. Cyclic voltammograms of the cells were
recorded in a Solartron 1470E Potentialstat/Galvanostat.
Synchrotron XRD Measurements.
[0106] X-ray diffractions were performed at the beamline 13-ID-D of
GSECARS sector at Advanced Photon Source (APS) at Argonne National
Laboratory. The X-ray beam (37 keV energy, corresponding to X-ray
wavelength of .lamda.=0.3344 .ANG.) was focused to a 2 .mu.m
diameter spot with a Kirkpatrick-Baez mirror system. The distance
and tilting of the MAR165-CCD detector were calibrated using a
CeO.sub.2 standard. V.sub.2O.sub.5 samples were prepared by
stripping the V.sub.2O.sub.5 film from Ni support onto KAPTON
tapes. Electrodes removed from cells for analysis were thoroughly
washed with dry dimethyl carbonate and dried under inert
atmosphere. The charged samples were also prepared by stripping
V.sub.2O.sub.5 films onto KAPTON tape. The discharged samples were
scratched-off and sealed inside a 3 mm-diameter hole in a piece of
aluminum foil by sealing KAPTON sheet to the foil using epoxy. All
cell operations were performed in a He-filled dry glove box (oxygen
level <2 ppm). Simulations of XRD patterns were carried out
using CRYSTALMAKER software (CrystalMaker Software, Ltd.) and are
shown in FIG. 18.
Synchrotron SAXS/WAXS Measurements.
[0107] SAXS/WAXS data were collected at Beamline 12ID-B of the
Advanced Photon Source (APS) at the Argonne National Laboratory
(ANL). The X-ray was focused, and the spot size on the sample was
about 50 m.times.50 .mu.m. SAXS and WAXS data were presented in
momentum transfer, q (q=4.pi. sin .theta./.lamda., where .theta. is
one-half of the scattering angle, and .lamda.=1.033 .ANG. is the
wavelength of the 12 keV energy probing X-ray), measured in the
range 0.01-2.3 .ANG..sup.-1. The charged samples were prepared by
stripping V.sub.2O.sub.5 films onto KAPTON tape. The discharged
samples were scratched-off and sealed inside a 3 mm-diameter hole
in a piece of aluminum foil by sealing KAPTON sheet to the foil
using epoxy. All cell assembly and disassembly operations were
performed in a He-filled dry glove box (oxygen level <2
ppm).
XPS.
[0108] A SCIENTA hemispherical electron analyzer (SES100) was used
to obtain the XPS/UPS measurements and total energy resolution of
spectra, including photon energy, was set to less than about 0.1
eV. The acceptance angle for incoming electrons is (+) 5 degree.
All experimental data were taken under the pressure of
2.times.10.sup.-10 torr or less.
Electron Microscopy.
[0109] Scanning Electron Microscopy (SEM) images were recorded with
a JEOL JSM-7500F Field Emission SEM operating at 30 kV.
XANES.
[0110] X-ray spectroscopy (XAS) and extended X-ray absorption fine
structure (EXAFS) measurements were performed at PNC-XOR bending
magnet beamline (20-BM-B) of APS in Argonne National Laboratory.
Measurements at V K-edge were performed under transmission mode
using gas ionization chambers to monitor the incident and
transmitted X-ray intensities. A third ionization chamber was used
in conjunction with a Ti-foil standard to provide internal
calibration for the alignment of the edge positions. The incident
beam was monochromatized using a Si (111) double crystal fixed exit
monochromator. Harmonic rejection was accomplished using a Rhodium
coated mirror. The charged samples were prepared by stripping
V.sub.2O.sub.5 films onto KAPTON tape. The discharged samples were
scratched-off and sealed inside a 3 mm-diameter hole in a piece of
aluminum foil by sealing KAPTON sheet to the foil using epoxy. All
cell assembly and disassembly operations were performed in a
He-filled dry glove box (oxygen level <2 ppm). The reference
standards (V.sup.5+ and V.sup.4+) were prepared by spreading thin,
uniform layers of the V.sub.2O.sub.5 and VO.sub.2 power in KAPTON
tape and stacking a few layers to attain the desired absorption
step height. Each spectrum was normalized using data processing
software package IFEFFIT. Alignment of each sample reference
spectrum with respect to V standard spectrum is within the range of
.+-.0.03 eV.
SEM Characterization of Morphology Before and after Cycling.
[0111] SEM images of both layered and crystalline V.sub.2O.sub.5
electrodes before and after electrochemical cycling in Na system
are shown in FIG. 17. As-prepared and 120.degree. C. vacuum
annealed V.sub.2O.sub.5 electrodes (a) are composed of nanoribbons
with highly porous structure and there is no noticeable change in
the surface morphology after cycling (c) indicating high mechanical
strength of this material. It should be noted that the debris on
the cycled nanoribbons is from residues of the electrolyte.
Crystalline structure of orthorhombic V.sub.2O.sub.5 particles (b)
deteriorated during cycling (d). Large irreversible capacity in the
1st cycle (FIG. 14A) due to series of irreversible phase
transformation during the first discharge process, badly affect the
overall battery performance. The breakdown in the crystal structure
(FIG. 15) did not result in the immediate failure of the cathode
material, only a continuous decline in electrode capacity (FIG.
14A).
XPS Characterization of V Valence State Before and after
Cycling.
[0112] The V2p and O1s peaks are shown in FIG. 19. The peak at
516.9 eV is assigned to the V.sup.5+ 2p.sub.3/2 orbital, peak at
524 eV to V.sup.5+ 2p.sub.1/2 orbital and that at 515.7 eV to the
V.sup.4+ 2p.sub.3/2 orbital. The spin-orbit split of V2p.sub.3/2
and V2p.sub.1/2 is about 7.5 eV, consistent with the previous
reports in the literature.
XRD Characterization of Bilayered V.sub.2O.sub.5 Structure Before
and after Na+ Intercalation.
[0113] Synchrotron X-ray diffraction spectra of bilayered
V.sub.2O.sub.5 vacuum annealed at 120.degree. C., before and after
cycling with Na.sup.+-ions are shown in FIG. 20. The bilayered
V.sub.2O.sub.5 structure is composed of 2D bilayered stacks. The
interlayer distance between the two single sheets of V.sub.2O.sub.5
making up the bilayer is close to 3.44 .ANG.. This structure is
reminiscent of the V.sub.2O.sub.5 xerogels in which monoclinic
bilayers of V.sub.2O.sub.5 stack up with the spacing that expands
(or contracts) as the xerogel incorporates (releases) water
molecules in its structure. Upon Na.sup.+ intercalation, one can
observe exceptionally strong narrow peak at 4.9 .ANG. that
corresponds to the distance between intercalated Na.sup.+ ions and
neighboring oxygen atoms that terminate the bilayered
structure.
[0114] In summary, nanostructured thin films of vanadium oxide
prepared on Ni metal substrates by electrochemical deposition have
been used as cathodes for Na-ion batteries without the need of
electro-conductive carbon additives and polymeric binder. The above
theoretical capacity of 250 mAh/g at room temperature, and
electrolyte stable redox potentials around 3 V translates into an
energy density of about 760 Wh/kg. Stable and reversible high
capacity on repeated cycling (C/8) makes bilayered V.sub.2O.sub.5 a
suitable cathode material for high-energy density rechargeable
sodium batteries.
[0115] All references, including publications, patent applications,
and patents, cited herein are hereby incorporated by reference to
the same extent as if each reference were individually and
specifically indicated to be incorporated by reference and were set
forth in its entirety herein.
[0116] The use of the terms "a" and "an" and "the" and similar
referents in the context of describing the invention (especially in
the context of the following claims) are to be construed to cover
both the singular and the plural, unless otherwise indicated herein
or clearly contradicted by context. The terms "comprising,"
"having," "including," and "containing" are to be construed as
open-ended terms (i.e., meaning "including, but not limited to,")
unless otherwise noted. Recitation of ranges of values herein are
merely intended to serve as a shorthand method of referring
individually to each separate value falling within the range,
unless otherwise indicated herein, and each separate value is
incorporated into the specification as if it were individually
recited herein. All numerical values obtained by measurement (e.g.,
weight, concentration, physical dimensions, removal rates, flow
rates, and the like) are not to be construed as absolutely precise
numbers, and should be considered to encompass values within the
known limits of the measurement techniques commonly used in the
art, regardless of whether or not the term "about" is explicitly
stated. All methods described herein can be performed in any
suitable order unless otherwise indicated herein or otherwise
clearly contradicted by context. The use of any and all examples,
or exemplary language (e.g., "such as") provided herein, is
intended merely to better illuminate certain aspects of the
invention and does not pose a limitation on the scope of the
invention unless otherwise claimed. No language in the
specification should be construed as indicating any non-claimed
element as essential to the practice of the invention.
[0117] Preferred embodiments of this invention are described
herein, including the best mode known to the inventors for carrying
out the invention. Variations of those preferred embodiments may
become apparent to those of ordinary skill in the art upon reading
the foregoing description. The inventors expect skilled artisans to
employ such variations as appropriate, and the inventors intend for
the invention to be practiced otherwise than as specifically
described herein. Accordingly, this invention includes all
modifications and equivalents of the subject matter recited in the
claims appended hereto as permitted by applicable law. Moreover,
any combination of the above-described elements in all possible
variations thereof is encompassed by the invention unless otherwise
indicated herein or otherwise clearly contradicted by context.
* * * * *