Res. Rep. Fac. Eng. Mie Univ., Vol. 29, pp. 65-72 (2004)

９

Original Paper

Novel negative electrode materials with high capacity density forfurther rechargeable lithium ion batteries
Y. Liua, Ｔ. Matsumuraa, A. Hiranob, T. Ichikawab, N. Imanishib and Y. Takedab( Satellite Venture Business Laboratory, b Department of Chemistry for Materials)a(Received October 20, 2004)
Abstract
Solid-state rechargeable batteries, especially, lithium ion batteries, are principle and promising
power sources for a wide variety of electronics. Electrode material is a key for developing further
lithium ion batteries, which are likely to require good reliability and high energy density. However,graphitic carbon that is currently used as negative electrode material in the commercial Li-ion
batteries appears to be unsatisfied due to low theoretic capacity of 372 mAh g-1 and poor thermal
stability under lithiated state. Therefore, there is even-increasing research in the feasibility of thereplacement of graphitic anodes. By this motivation, a series of novel negative electrode materials
that demonstrate extremely high capacities and different Li reactive mechanisms have been
proposed in this group. In this paper, we present a brief introduction with respect to some recent
results. In particular, the characterizations of these anode materials are described and discussed
based on different electrolytes such as liquid and solid systems.
Key words: Power sources; Rechargeable lithium ion batteries; Negative electrode materials;
Energy density; Working voltage; Capacity; Li reactive mechanisms.
BackgroundSearch for new energy sources becomes an urgent challenge in our society and is greatly related to thetechnology development. Such energy requires high efficiency, convenient, pollution-free and safety. As
alternative energy sources and energy storage media, batteries have many advantages and are principle andpromising power sources for a wide variety of electronics, e.g., portable, entertainment, computing andtelecommunication equipments, as well as electric vehicles (EV) and hybrid electric vehicles (HEV) [1]. A typicalbattery is composed of many electrochemical cells in series in terms of exporting-energy requirement. Single cell
contains a positive electrode (cathode) and a negative electrode (anode), which are separated by anionic-conducting but electron-insulating electrolyte. The working voltage of the cell is due to the potential
difference between the cathode and the anode (both sources of chemical reactions). Furthermore, the chemical
reactions processed in both cathode and anode can generate
electrons and currents in case of an external circuit connection
between the two electrodes. The amount of electrical energy
produced by the battery is dependent on several key parameterssuch as working voltage (V) and specific capacity (Ah Kg-1), bothof which are linked directly to the chemistry of the system, and isexpressed according to either per unit of weight (Wh kg-1 ) or per
unit of volume (Wh L-1). At present, fuel cell, nickel metal
hydride and lithium ion batteries are the major research and attracteven-increasing attention. In particular, lithium ion batteries havea remarkable expansion and now supply over 63 % of the total
portable battery market in worldwide (90% in Japan, data citedfrom 2001) [2]. This is attributed to that lithium ion cell has rather
high working voltage and thus large energy density. There is along historical development over 30 years for Li-battery
technology. The initial motivation for using Li as anode in
Fig.1. Principle of the lithium ion cell.
batteries is due to the fact that metallic lithium possess the mostelectropositive character (-3.04 V vs. standard hydrogenY. Liua, T. Matsumura, A. Hirano, T. Ichikawa, N. Imanishi and Y. Takeda

10electrode) and the lightest weight (equivalent weight M= 6.94 g mol-1, specific gravity ρ=0.53 g cm-3), as well as
the largest capacity density (ca. 3.8 Ah Kg-1); thereby battery based on lithium anode shows very high discharge
voltage and correspondingly large energy density. However, lithium dendrites that occur on the surface of thelithium anode during the repeated charge and discharge tend to cause a short current circuit and thereby a potential
safety problem. Fortunately, this deterrent was tackled successfully by Sony Corporation, Japan, in June 1991, by
introducing a kind of Li-intercalation host, such as graphite, in place of metallic lithium. Correspondingly,
lithiated LiCoO2 (lithium cobaltate) is adopted as cathode that gives high potential and acts as Li-sources. The“rocking chair” concept of the lithium ion cell proposed and commercialized firstly by Sony Corporation, as
shown in Fig.1, make mobile Li+ become shuttle between the cathode (LiCoO2) and the anode (graphite) via theelectrolytes in a charge and discharge process accompanied with an electrochemical electrode reaction. During thecharging process under an applied external voltage, a certain amount of Li+ are extracted from the cathode into theelectrolyte, which is composed of lithium salt dissolved in a liquid solvent or a solid polymer/glass and is notconsumed by the electrochemical charge–discharge reaction. Simultaneously, an equivalent amount of lithium ion
are forced from the electrolytes intercalate into the graphite anode. The movement of the lithium ion during
discharge is reverse to that of charge stage and is spontaneous. Charge compensation takes place through theexternal circuit during charge and discharge. By this way, the electrical energy can be saved during charging andreleased during discharging reversibly. The “rocking chair” manner highly avoids Li-dendrites and thus
demonstrates high reliability. The currently commercial lithium ion cell consisting of LiCoO2 and graphite hashigh discharge voltage of ca. 3.7 V and large energy density of ca. 110 Wh Kg-1, which are apparently superior tothose of Ni-Cd and Ni-MeH cells. Obviously, electrode materials that can reversibly storage lithium ion are thekey for developing further lithium ion batteries. The research in the cathode that must possess high electrodepotential mostly focuses on how to improve the electrode capacity, which is below to 200 mAh g-1 in the current
system, indicating that the weight of cathode to anode must be over 2 times. The large efforts have resulted inseries new compounds such as layered mixed transition metal (Ni-Mn-Co) oxides (e.g., LiNi1/3Co1/3Mn1/3O2) [3]and olivine lithium iron phosphate (e.g., LiFePO4) [4], which appear to be promising cathode alternatives for
LiCoO2 and is in the stage of application. In the anode side, research remains more urgent, although various exotic
carbonaceous and non-carbonaceous materials, e.g., disordered carbon [5], Li-alloys [6], lithium transition metal
nitrides [7], transition metal oxides [8] and composite tin oxides [9], have been deeply studied with an improved
capacity. It means that the “standard” carbonaceous materials, such as graphite and coke, are still likely to be theanode of choice for the next few years. A considerable investigation in our group for developing further Li-ion
batteries has been under progress over 10 years and has lead to some remarkable results including of the electrodematerials, the electrolytes and the battery system. In this paper, we give a brief introduction of the anode researchof recent two years. Especially, the synthesis, characterization, electrochemical behavior and safety problem ofthese novel anode materials are discussed and presented in detail according to the types of electrolytes, such as
liquid organic electrolytes and solid polymer electrolytes.
Long time
Part I: Liquid electrolytes based system
1. Li-alloy based composites
1.1

Introduction

The replacement of the Li metal anode by Li-alloys in rechargeable Li-ion cell has been under investigation
since Dey demonstrated the possibility of electrochemical formation of LixAl in liquid organic electrolytes in 1971
[10]. Since that, it has been found that numerous metals, e.g., Mg, Ca, Al, Si, Ge, Sn, Pb, Sb, Bi, etc., canreversibly accommodate lithium under an electrochemical alloying process according to the mechanism: xLi+ + e+ M = LixM [6]. Such a reaction can generate larger theoretic specific and volumetric capacities than the currently
used lithiated graphite (LiC6). In addition, the reactive potential of the lithium alloying process is close to that of
the metallic lithium. Therefore, there has been a long time consideration for using Li-alloys as promising anodecandidates for Li-ion cell. In general, Li-alloys have several advantages over carbonaceous materials which favor
their use in a Li-ion cell, including of higher lithium storage capacities, in particular, when the charge densities arecomparable to metallic lithium, and a relatively higher operating voltage, retarding the deposition of metallic
lithium during charge, and thus improving safety and fast-charging capability. Besides, Li-alloys appear not tohave much sensitive to the solvent co-intercalation compared with the graphitic anode materials. However, there isa major problem that encumbers the application of Li-alloys. This is mainly due to that the volume mismatch of
the alloy hosts between Li intercalation and extraction are comparatively large (typically in the scale of over ca.
200 %), as shown in Fig.2, and such that the host matrix undergoes drastic mechanical stresses during charge anddischarge. Such a mechanical degradation results in the metal cracks and the host particle pulverizes, and thusNovel negative electrode materials

11

looses mechanical and electrical contact. Eventually, capacity fading is inevitable. By contrast, the dimensional
stability of the carbonaceous host matrix during cycling is associated with very low volume changes during
intercalation and removal of lithium (ca. 8%, for graphite) that results in good charge reversibility. There havebeen two major designs to overcome the pulverization dilemma of Li-alloys. One is based on the concept using
composite or multiphase structure, in which the active material/phase is supported by the less-active or evennon-active matrices, to keep the mechanical integrity [11].
The other approach is to suppress
the drastic volume changes bydesigning a proper morphology of thehosts, characterized by thin layers
comprising small particle size active
material [12]. The effort to conquer
the morphology instability of the
Li-alloys leads to a challenge in thematerial research. Recently, we pay
our attention on silicon, which has a
low reactive potential of ca. 0.2 V vs.
Li/Li-1 and has the highest capacity ofca. 4000 mAh g-1, which is 10 times as
Fig.2 Capacity density and correspondingly volume changes
that of the current graphite, among all
before/after Li intercalation of different Li hosts.
known lithium storage hosts. [13-17].
However, similar to other Li-alloys,
silicon possesses poor capacity retention during cycling due to the volume mismatch upon Li intercalation andextraction. As mentioned, a promising approach to overcome the volume effect is to create a composite
microstructure comprising the active silicon uniformly dispersed in an inert matrix. The Sn-Fe-C nano-composites
consisting of active FeSn2 and inactive SnFeC proposed by Dahn’ and co-workers suggest that high-energy
mechanical milling (HEMM) could be an appropriate process to built such a composite microstructure [18]. By
this way, some examples such as TiN-Si, TiB2-Si, FeSi-graphite, and SiC-Si composites have been reported with
the improved cyclability [19-22].
From an application standpoint, a considerable improvement in the first cycle efficiency and the cycling life
remains to be done for such anode systems. There has been another attention that focuses on composites
consisting of silicon and carbonaceous matrices. For instance, silicon covered with carbon by thermal vapor
deposition (TVD) showed a better cycling stability than the pure silicon [23-25]. Dahn et al. reported that bypyrolyzing siloxane polymer or blends of pitch and polysine, silicon could be incorporated into carbon at atomic
scale. These materials had a stable capacity of ca. 500 mAh g-1 but low initial faradaic yield [26, 27]. We also
reported that pyrolysis of pitch or poly (vinyl chloride) (PVC) embedded with silicon and graphite provided aneffective way to suppress the volume effects of silicon [28-32]. The enhanced cyclability is attributed to the small
volume expansion of carbon on lithium intercalation (ca. 9% for graphite) and the ability of the ductile carbon toaccommodate the volume change of silicon, reducing mechanical strain within the electrode and consequent
electrode disintegration. Although this type of pyrolyzed carbon has a large potential hysteresis between
Li-insertion and extraction, it may mostly function as an elastic network with electron/ion conductivity that
permits the silicon in the carbon matrix to operate while maintaining electrode integrity. It seems that to a great
extent the cycle performance of the pyrolyzed Si-C systems depends on the distribution of silicon and the contact
strength between the active host and the matrix. However, a single pyrolysis step appears to be insufficient. We
found that a combination of high-energy mechanical milling (HEMM) treatment with pyrolyzed process in the
preparation for the Si-C composites can overcome the mentioned detriment from pyrolysis reaction and HEMM
step alone. By this way, a series of Si-based composites have been proposed. The Si-C composite synthesizedfrom two PVC pyrolysis reactions, combined with an intervening HEMM step, shows good cycling and large
capacity of 900 mAh g-1, plus high initial coulombic efficiency [28,29]. The Si-M-C composites prepared fromballmilling silicon with the inert conducting components and followed by a PVC pyrolysis process also show both
good capacity and capacity retention [30-32]. In this part, characterization of the proposed Si based materials arepresented and discussed in detail.
1.2. Experimental
1.2.1 Material preparations.
The basic preparation of the Si-C (silicon/disordered carbon) composite was as follows: poly (vinyl
chloride) (PVC, Aldrich) and silicon particles (<1 µm, >99.8%) were homogeneously mixed and the weight ratio

12

Y. Liua, T. Matsumura, A. Hirano, T. Ichikawa, N. Imanishi and Y. Takeda

of silicon vs. PVC was 3:7. The mixture was heated at 900℃ in an Ar atmosphere for 1 h at a heating rate of 5 ℃
min-1 and allowed cool down to room temperature normally. The products were further treated by a high-energy
mechanical milling (HEMM) step in a sealed bowl in Ar at a rotational speed of 500 rpm for 2-10 h. The resulting
samples were mixed with PVC again (Milled product vs. PVC was 3:7, wt.%). The mixture was processed by apyrolysis reaction following the same procedures as the first heating process. The preparation of the Si-M-C (M, akind of co-ballmilling component; C, pyrolyzed PVC) composites was as follows: Si-M mixture consisting ofsilicon particles and TiN (ca. 1.5 µm), graphitic carbon (average particle size: 10-20 µm) with a weight ratio of
1:1 was treated by high-energy mechanical milling in a sealed bowl in Ar at a rotational speed of 500 rpm for 2-10
h. The resulting product was mixed with poly (vinyl chloride) (milled product vs. PVC was 3:7, wt. %) and heatedat 900℃ in an Ar atmosphere for 1 h at heating rate of 5 ℃ min-1 and allowed cool down to room temperaturenormally. The final samples were ground and sieved. The SiMg2 and NiSi2 (Aldrich) materials were prepared bytreating the commercially obtained products with HEMM and further sieved. Nano-size SiO1.1 (typical particle
size: 50 nm, prepared by CVD) was kindly grafted from Nagaku Ltd.co. Powder X-ray diffraction (XRD) patterns
were obtained using automated powder diffractometer with Cu Kα radiation (Rotaflex RU-200B, Rigaku-denki
Corporation). The morphology performance of the materials was characterized with scanning electron microscopy(SEM) and Electron Probe Micro-Analysis (EPMA).
1.2.3 Cell assembly and electrochemical tests.The electrode containing 8 wt% acetylene black (AB), 80 wt% active materials and 12 wt% Poly (vinylidene
fluoride) (PVDF) was prepared by a normal casting. The active powders and AB were homogeneously mixed in a
0.02 g mL-1 PVDF/1-methyl-2-pyrrolidone (NMP) solution, and the vicious mixture was cast onto a 300-µm thickNi foam, which served as a current collector. The electrode was further dried at 120℃ under vacuum for 2 h until
NMP solvent was entirely removed. After pressing, the geometric area of the electrodes was 1.0 cm2, and thetypical thickness was 190~200 µm. To evaluate the electrochemical properties of the electrodes, a half-cell
containing LiClO4 / EC+DMC (Ethylene carbonate plus diethyl carbonate as 1:1 in volume) electrolyte was used,
and lithium metal was utilized as the counter electrode. All the three layers, including test electrode, separator andlithium metal, were stacked in a 2025 coin type cell in a glove box. Unless stated elsewhere, cycling was carriedout at a constant current density of 0.18 mA mg-1 and a voltage cutoff at 1.5/0.02V vs.Li/Li+. Charge anddischarge of the cell refer, respectively, to lithium extraction from, and insertion into, the active hosts. Electrodecapacity was calculated according to the weight of active materials.
1.3

Results and discussion

1.3.1

Electrochemical behavior of pure Si insertion host
The electrochemical characterization of the silicon with different particle size was investigated andcompared in Fig.3. All the silicon is reactive to lithium that results in a large capacity. A decrease in the particle
size enhanced the solid electrolyte interface (SEI) film formation on the surface of the active hosts; thereby it
increased the capacity loss in the first cycle. There is an obvious shift in the discharge potential plateau from thefirst cycle to the second cycle that could relate to
10 µm, η =69.0%
1.6
the irreversible phase transformation of silicon
200 um, η = 74.5 %< 1 µm, η =66.2%
from crystalline to amorphous state in the first
1.4
electrochemical alloying process. Capacity
1.2
retention shows high similarity for the silicon with
Discharge
1.0
different particle size, indicating that theChargeLi intercalation
Li extraction
morphology deteriorating mechanism of Si upon
0.8
charge and discharge is obviously dissimilar from
0.6
normal alloy hosts, such as Sn, whose cyclability is
0.4
highly dependent on the particle size, as mentioned
η1(The first cycle efficiency)
Cycle 2in the introduction part. Fig.4 shows such an=Discharge capacity/Charge capacity
0.2
imaginable mechanism difference in the
0.0
Cycle 1morphology changes upon Li insertion and
0
500
1000
1500
2000
extraction for Si (a) and Sn (b) based hosts.
-1
Capacity / mAh g
Differing form that of Sn, the repeated Li insertion
into and extraction from the active Si do not cause
a particle collapse. By contrast, the volume
Fig.3. Charge and discharge profiles of the silicon
expansion and shrink of Si possess homogeneouswith different particle size.
change due to the electrochemical driven
11

Voltage / V

1Novel negative electrode materials

13

amorphous character. However, the volume effects inevitably result in electron-contacting loss in both Si and Sn,
and finally cause capacity fading. A further understanding with respect to the morphology failure mechanism of Si
is still in the progress.

Voltage / V

Dis.capacity / mAh g-1

1.3.2 High mechanical stable Si-C composites for Li intercalation
To suppress the volume
effects of silicon upon cycling,
we try to disperse ultrafine silicon
into a feasible carbonaceous
matrix by means of pyrolyzed
PVC reaction and HEMM
process. The inset in Fig.5 shows
charge and discharge profiles of
the Si-C electrode at different
cycles. Insertion potential ofsilicon in the first cycle mainly
occurs below 0.1V, yielding an
extremely flat plateau. The
averagepotentials
ofthecomposite in the discharge andcharge stage are 0.2 V and 0.4 V.
The electrochemical behavior of
the composite is dominated by
Fig.4. Mechanism of morphology changes of Si and Sn Li upon Li
silicon insertion host, suggesting
intercalation and extraction.
only small amount lithium
3000
storages in the carbonaceous matrix. The faradaic
1.6
1.4
yield of extraction to insertion is close to 100 % after
Cycle 1
1.2
Cycle 22500
the second cycle, versus 82.6% at the first cycle.
1.0
Cycle 29
0.8
The synthesis steps have obvious effects upon
0.6
the electrochemical behavior of the composites.
2000
0.4
0.2
Fig.5 illustrates the cycling performance of the Si-C
0.0
composite prepared from different procedures. One
1500
0
200
400
600
800
1000
-1
pyrolysis reaction has the worst capacity retention.
Capacity / mAh g
SEM analysis reveals that the Si-distribution, as
1000
shown in Fig.6a, is relatively poor. Some cracks andOne pyrolysis
lots of silicon over the surface of the composite
500
Two pyrolysis
suggest the unstable morphology and incompletely
HEMM after one pyrolysis
Two pyrolysis with HEMM
embedding can hardly prevent Si from deterioration.
0
0
6
12
18
24
30
The composite synthesized from two pyrolysis
Cycle number
processes without HEMM step has a slightly
improvement in the cycling performance. By
contrast, HEMM appears to be highly favorable forFig.5 Cycling performance of the Si-based composites
long cyclability. HEMM is well known as an
prepared from different procedure. The inset figure iseffective energy transfer way via the vigorously
charge and discharge profiles of the Si-C composite.
interfacial collisions among precursors and balls. It

Fig.6 SEM and corresponding EPMA images of the composites prepared from a) one pyrolysis reaction; b)after one pyrolysis reaction and HEMM; and c) two pyrolysis reaction with a an intervening HEMM step.

14

Y. Liua, T. Matsumura, A. Hirano, T. Ichikawa, N. Imanishi and Y. Takeda

Dis.capacity / mAh g-1

Current / mAenhances Si/C contact and generates many Si-C clusters with a multiphase structure formed under the strong
particle impacts, as shown in Fig.6b. Such a structure can achieve the reasonably improved interface affinity
between the silicon and the carbonaceous matrix. HEMM also increases the packing density by reducing theporosity of the pyrolyzed PVC from the first pyrolysis, and ensures the adherent integrity between the dispersed
silicon and the carbonaceous matrix. However, HEMM increase the initial irreversible capacity due to an
enhanced SEI formation on the surface of the particles. The composite made from HEMM step between two
pyrolysis reactions possesses the best cycling. The corresponding EPMA analysis reveal that no obvious siliconregions could be identified, implying that the silicon particles are completely embedded in the disordered carbon.
The large particle of the composite also leads to a reasonable initial coulombic efficiency, as shown in Fig.6c.
A favorable structure concerning this type of composite is that silicon is completely and uniformly trappedand embedded in the carbonaceous matrix. We believe that
3000
1.0
2
the key points determining the mechanical stability are
0.5
1
homogeneous Si-distribution and high adhesion strength
2500
0.0
between the silicon and the matrix. By this way, a possible
-0.5
-1.0
2000
aggregation of silicon can be highly avoided and a good
Cycle 2
-1.5
electron-contacting channel can always be maintained. A
Si/CCycle 1-2.0
1500
comparison with pure silicon in the cycling performance
1.5/0.02 V vs.Li-2.5-0.2 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6
reveals that the Si-C composite has a remarkably enhanced
Voltage / V
1000
cyclability, as shown in Fig.7. Controlling Li-insertion-1
level at 600 mAh g can effectively improve the500morphology stability of the composite electrode. This canPure silicon
Si/C: 600 mAh g / 1.5 V vs.Li
be explained by the alleviated electrode disintegration and
0
0
10
20
30
40
deactivation. Reducing the Li-insertion degree suppresses
Cycle
numberelectrode degradation on cycling.
The synthesis temperature has an obvious effect upon
the electrochemical behavior of the samples. The inset in
Fig.7. Cycling performance of the Si-M-C
Fig.7 illustrates the CV curves of the 1000℃ product.
composites based electrodes. The inset figure is
There are three cathodic peaks appear during theCV profiles of the Si-C composite prepared at
Li-insertion process at the first cycle: 0.8V, 0.15 V and1000 ℃.
0.01V. The latter two peaks mainly link to Li-intercalation
into disordered carbon and silicon respectively. The former
one, however, could be due to a new compound formed under the prolysis reaction. Also during the anodic scan,
three peaks can be found at 0.3 V, 0.55 V and 1.1 V, representing three Li-extraction stages which correspond tosilicon, disordered carbon and the new compound. The 0.8V cathodic peak and the 1.1V anodic peak display thehigh reversibility.
-1

1200

2000
1500
1000

500

6

12

18

24

30

Si (220)0

Si (111)

0

1000

800
600
400
200

0

Si (331)2500

Si (400)

Capacity / mAh g

3000

Si (311)-1

3500

GraphiteDis. capacity / mAh g -1

The introducing of relatively inert components in the Si-C composites
Fig.8 shows X-ray diffraction patterns of the(2)(1)
resulting powders of a) Si-C, b) Si-graphite-C and c)
Si-C
Si-TiN-C composites. All the composites indicate the
SiliconSi-M-C
presence of crystalline silicon. Especially, the peak
Si-C Si-graphite-C
reflection of silicon in the Si-C composite isremarkable, suggesting that HEMM may have a weak
Si-TiN-C
1
0.25 0.5
2influence upon the crystallite of silicon. By contrast, it
Cycle number
Charge rate / C
is characteristic that intensity decreasing and peakbroadening reflected from silicon in the Si-TiN-C
composites are noticeable compared with those in theaSi-graphite-C and Si-C composites. This is probablybattributed to that the relatively hard TiN tends to
TiN TiN
TiN
TiN
c
reduce the crystallite size of silicon and make the
20
30
40
50
60
70
80
active particles partly amorphous during the HEMM
2 θ / deg.
treatment.
Fig.8. XRD patterns of the resulting composites of a)
All the composites have a high similarity in the
Si-C, b) Si-graphite-C and c) Si-TiN-C. The inset
voltage trends upon charge and discharge. The cycling
figure (1) is cycling performance of different
performance of the Si-based composites was shown incomposites and inset figure (2) is Charging rate as a
the inset (1) in Fig.8. The composites have afunction upon the reversible capacity of the Si-C andremarkably enhancement in the morphology stability
Si-M-C composites.
for Li insertion. As mentioned, there are several
Intensity / a.u.

1.3.3

Novel negative electrode materials

15

possibilities for such a cycling improvement. First, HEMM increases the Si-distribution and decreases the active
particle size. Secondly, the improved adhesion strength between the silicon and the relatively inert components
treated under HEMM can hold a good electron-contacting. Finally, pyrolyzed PVC can cover the milled
Si-clusters and function as an elastic network with electron/ion conductivity that permits the silicon in the carbon
matrix to operate while maintaining electrode integrity. The pyrolyzed PVC also can function as a binder to holdthe mechanical integrity among different components. Due to the poor electron-conductivity of silicon in theorigin, the Si-C composite suffers from an obvious loss in the reversible capacity with increasing the charging rate,
as shown in the inset (2) in Fig.8. In the case of the Si-M-C composites, the reversible capacity can be remained tobe ca. 50%, when the charging rate was increased to 2C. A decreased particle size of the silicon active hosts may
take this consequence. It means that ballmilling silicon with the conducting TiN and graphite prior to the pyrolysis
reaction tends to decrease the active particles size and brings an obviously improvement in the charging rate over
the Si-C composite. However, silicon suffers from a loss in the capacity utilization due to an inactive phaseformed between the silicon and the co-milling components. Especially, the milling time is extended.

Dis.capacity / mAh g

Voltage / V-1

1.3.4 Comparison in the cycling stability of different Si hosts materials
Adopting intermetallic multiphase structure also can
2.5
apparently alleviate the morphology instability of silicon
for Li-insertion. Fig.9 shows the charge and discharge
Pure silicon
curves of the electrodes based on NiSi2, SiMg2 and Si-C
2.0
composite at the second cycle. Two obvious discharge
Si-C
NiSi2
potential plateaus at 0.2V and 0.1V of the SiMg2
1.5electrode is linked to Li-intercalation process to form aSiMg2
single crystalline LiMg2Si, following a reaction between
1.0
Li and Si, finally lithium may alloy with Mg. Also, twocycle number
charge potential plateaus appear at 0.3 V and 0.7 V,
representing the Li-extraction stage that reveres to the
0.5
S-C
Li-Si-Mg alloy process. In the case of NiSi2 and Si-C
Mg Si
composite, similar electrochemical behaviors suggest that
Si Ni
0.0
the insertion mechanism of both hosts is only related to
-1000
-500
0
500
1000
the electrochemical Li-Si alloying reaction. Although-1
Capacity / mAh gsilicon content in Si2Ni (ca.48 wt %) is almost equal tothat of the silicon/disordered carbon composite, the latter
Fig.9. Charge and discharge profiles of different
shows much larger reversible capacity, indicating thesilicon based hosts. The inset figure is cycling
higher electrochemical reactivity. The cycling
performance of different silicon based hosts.
performance of the electrodes based on pure silicon,
NiSi2, SiMg2 and silicon/disordered carbon composite isshown in the inset in Fig.9. All the composite shows the
1.6
2400
relatively improved cycling stability compared to pure
1.4
SiO1.1
2200
1.2
Si-C
silicon. However, due to the comparable “crash”
1.0
2000
0.8
characteristic of the metallic matrix, the mechanical
0.6
1800
0.4
stability of NiSi2 and SiMg2 is obviously inferior to the1600
0.2
silicon/disordered carbon composite. Especially in the0.0
1400
0 200 400 600 800 1000120014001600
Si-C, η = 82 %case of SiMg2, a close reactive potential scale of Mg and
1200-1
Capacity / mAh g
1000
Si will easily cause to a rapid capacity decline,
800
suggesting that the Li-Si-Mg process is only partially600
reversible. Oxide such as inert Li2O and SiO2 also play a
400
SiO , η =46%
noticeable effect to enhance the morphology stability of200silicon upon volume change. X-ray diffraction patterns
0
0
6
12
18
24
30
36
reveal that the single phase of intermetallic NiSi2 andCycle number
SiMg2 can be recognized. By contrast, the relatively
weak peak indicates that SiO1.1 contains lot of
amorphous SiO and SiO2 besides crystalline silicon. The
Fig.10. Cycling performance of different siliconinset in Fig.10 shows the comparison of the charge andbased hosts. The inset figure is charge anddischarge behaviors between SiO1.1 and Si-C. The
discharge profiles of different silicon based hosts.
discharge plateau of SiO1.1 is slightly lower than that ofsilicon/disordered carbon, indicating Si and SiO2 within
the SiO1.1 electrode may be obstructive to the initial Li-insertion. Cycling performance of the electrodes based on
SiO1.1 and Si-C composite is shown in the inset in Fig.10. A significantly initial capacity loss of the SiO1.1
3500
3000
2500
2000
1500
1000

500

0

0

6

12

18

24

30

36

2

Discharge capacity / mAh g-1

Voltage / V

2

1

1.1

1

16

Y. Liua, T. Matsumura, A. Hirano, T. Ichikawa, N. Imanishi and Y. Takeda

electrode can be noticed, relating to an irreversible reaction of amorphous SiO with lithium to form Li2O and Si at
the first cycle. The huge specific area of nanosize active hosts also takes part consequence. Due to buffering effectof Li2O and SiO2, the high capacity retention is also remarkable for SiO1.1. Bear in mind that a noticeably large
irreversible capacity loss of SiO1.1 in the first cycle is inevitable; the electrochemical performance of the Si-C
composite is obviously superior to that of SiO1.1.
1.3.5 Thermal stability of the Si-C based composites under lithiated state
It is well known that the thermal stability of the graphite anode in the inorganic electrolytes is controlled by
the SEI film formed on the surface of the lithiated active hosts [35]. Since the reactive potential of the composite
is close to that of graphite after the first cycle and the SEI film is formed on the surface of the carbonaceous
matrix, it is reasonable that the thermal behavior of the Si-C under lithiated state is similar to graphite, as shown in
Fig.11. There are two sharp exothermic peaks and a small one starting from 130℃ to 280℃ for the graphite
electrode, which correspond to a mild heat generation caused by the reaction (SEI formation) of the electrolytes
and the lithiated graphite. The exothermic peak at 280 ℃ is probably due to the reaction of lithiated graphite andelectrolytes by a breakdown of the SEI as reported by J.R.Dahn et al. and J.Yamaki et al [35, 36]. However, thematched exothermic peak at 280℃ of the Si-C composite significantly turns into weak, indicating a relatively low
heating generation due to the comparatively increased insertion potential. Moreover, the lithium storage indisordered carbon is extremely lower than that of silicon. The later, on the other hand, is embedded inside and a
direct contact with electrolytes is highly avoided. It also leads to the relatively improved thermal stability.

Volatge / V

Heat Flow / a.u.

1.3.6 High energy density Li-ion cells consisting of silicon based anodes
A comparison with mesocarbon microbead (MCMB,a kind of commercial graphitic carbon) in the charge
4.0 (a)3.5
and discharge profiles at the second cycle reveals that
3.0
the Si-based composite has a slight increase in the Li
LiCo Ni O / SiC
2.5
insertion and extraction potential plateau for about 0.12
2.0
1.5 (b)V, as shown in inset (b) in Fig.11. This, in turn, canGraphite / Li
1.0
prevent lithium dendrite formation at a high charging
0.5
SiC / Li metal
0.0
rate and leads to an enhanced operation safety. However,-100 0 100 200 300 400 500 600 700
the Si-based composite surfers from a loss in the-1
Capacity / mAh g-3volume capacity due to a low density (1.7 g cm , incase of Si-C) compared with that of MCMB (ca. 2.25 g
Si-Ccm-3). A full cell using the Si-C composite anode andthe LiNi0.8Co0.2O2 cathode was fabricated to study its
Graphiteelectrochemical characteristics, as shown in the inset (a)
100
200
300
400
500
o
in Fig.11. For full utilization of lithium storage capacity
Temperature/ Cof the Si-C composite anode, the weight of cathode has
to be much over that of the anode (7−8 times). The cellFig.11. The DSC curves of silicon/disordered carbon
has a large anode capacity of ca. 600 mAh g-1 withincomposite and graphite under lithiated state. Scan
the working potential of 2.3−3.9 V that might result in
rate: 5 ℃ min-1 under Ar flow of 100ml min-1. The
high energy density.
inset figure is the charge and discharge profiles ofdifferent cells.
1.4 Conclusion
0.2

0.8

2

The volume effects of silicon upon Li insertion and extraction can be effectively suppressed by designing acomposite microstructure containing that ultrafine silicon are uniformly dispersed in a ductile conducting
carbonaceous matrix with electron/ion conductivity. By this way, several types of Si-based composites have beendeveloped by means of pyrolysis process in accompany with high-energy mechanical milling step. The pyrolyzed
carbon can mostly function as an elastic network with electron/ion conductivity that permits the silicon in thematrix to operate while maintaining electrode integrity. The Si-based composites have large capacity of ca. 1000mAh g-1 and good cycling performance, as well as acceptable first cycle efficiency. We expect that further
optimization of these Si-composite based anodes might lead to practical lithium-ion batteries with high energy
density.
2. Lithium transition metal nitrides based composite electrodes
2.1 Introduction

Novel negative electrode materials

17

In view of the morphology stability, especially, under highLi-utilization, lithium transition metal nitrides with hexagonal
symmetry, P6/mmm, such as Li2.6M0.4N (M=Co, Cu, Ni), aresuperior to Li-alloy in the cyclability as promising anodecandidates. The ternary Li2.6Co0.4N is composed of Co
substituting lithium between the Li2+N3- layers of Li3N. Li3N
has two kinds of Li ion. That is, location (1): Li ion is locatedin the Li2N layer and (2): Li ion is located between the Li2N
layers. It is believed that Co substitutes Li ion in the location(2), as shown in Fig.12. The intercalation mechanism of thelayered lithium transition metal nitrides has been primarily
investigated by this group and T. Shodai et al. [7,37-40]. The
solid solution range of the Li3-xMxN is located in 0<x<0.5,
0<x<0.6 and 0<x<0.3 for M=Co, Ni and Cu correspondingly.
From an electrochemical standpoint, the most promising andmost studied of the compound in this family is Li2.6Co0.4N. It
Fig.12. Structure model of the Li2.6Co0.4N.
was found that the large capacity of 900 mAh g-1 of Li2.6Co0.4N
is associated with a structure change from a crystal to an
amorphous phase in the first Li-extraction process [41]. Such a transformation indicates the atom rearrangement
and lithium in the compounds can be electrochemically extracted and re-intercalated with high reversibility.
However, Li2.6Co0.4N does not show satisfactory capacity retention upon long cycles. The capacity fading is
attributed to the electrochemical instability under high Li-extraction and the interfacial incomparability on thesurface of the active hosts [42,43]. In the pervious work, we have developed a series co-doped lithium transition
metal nitrides by means of combination of solid-state reaction at high temperature and high-energy mechanicalmilling (HEMM) process [44]. These compounds are superior to ternary Li2.6Co0.4N in terms of capacity retention
and cost. On the other hand, lithium must be extracted from the nitrides in an initial anodic oxidation, indicatingthat these Li-rich type anodes cannot directly combine with the typical high potential cathodes such as LiCoO2and LiMn2O4 to make Li-ion cells. We have showed that the Li-rich compounds can be used as insertion hosts inseveral ways. For example, lithium in the compounds can be pre-extracted by a chemical way before assembling acell [45]. Lithium metal nitrides also can be introduced to make compensation for the irreversible capacities of thealloy, carbon and metal oxides based electrodes in the first cycle [46-48]. Furthermore, we designed a type ofSi-graphite-Li2.6Co0.4N composite via two HEMM steps [49]. Due to the capacity compensation from the lithiatedLi2.6Co0.4N in the first cycle, the composite material demonstrates an extremely high initial faradaic yield of 92%,
as well as a large capacity and an acceptable cyclability. Recently, we have reported a composite electrodeconsisting of spinel LiTi2O4 and lithium metal nitrides. LiTi2O4 can extract lithium from the lithium transition
metal nitrides and function as an inert-conducting material in the electrodes within a special potential range [50].
The composite electrode shows high volumetric capacity and excellent cycling stability. More recently, we foundthat the novel composite anodes, which consist of transition metal oxide, such as Co3O4, and hexagonal lithiumtransition metal nitrides, demonstrate high electrochemical capacities of about 500 mAh g-1 and 100 % capacity
retention at the first cycle, as well as good cyclability [51]. Transition oxides can thermodynamically extract
lithium from the lithium transition meal nitrides that leads to a delithated state of the nitrides and to form metallic
Co and Li2O. Thus, lithium transition metal nitrides can be directly treated as being in a delithiated state. Such acomposite anode also shows high interfacial compatibility and low reaction heating with solid PEO electrolytes.
In this part, some of the above obtained results based on the hexagonal lithium transition metal nitrides arepresented and discussed in detail.
2.2 Experimental
2.2.1 Materials preparations and XRD measurements
The hexagonal compound was prepared by a solid state reaction following a pervious report [7]. The basic
procedure was as follows: mixtures of Li3N and powders of transition metal were pressed into tablets with 8 mm
in diameter and 5-8 mm in thickness in an Ar atmosphere. The tablets were heated at 700℃ for 12 h under a N2
stream at a heating rate of 35 ℃ min-1. The reactions were allowed cool down to room temperature normally. For
the compounds containing Fe, heating temperature was increased to 800 ℃. The resulting products were groundin a glove box and further treated by high-energy mechanical milling (HEMM) at a rotational speed of 500 rpm
for 20 h. Powder X-ray diffraction (XRD) patterns were obtained using automated powder diffractometer with Cu
Kα radiation (Rotaflex RU-200B, Rigaku-denki Corporation). The samples assembling and measurements were18

Y. Liua, T. Matsumura, A. Hirano, T. Ichikawa, N. Imanishi and Y. Takeda

performed under air-sealing condition.
2.2.2 Electrode fabrications and electrochemical measurements
A given weight of the electrode components, including 20 wt% acetylene black (AB), 70 wt% active
materials and 10 wt% Poly (vinylidene fluoride) (PVDF), was homogeneously mixed in an agate mortar in a glove
box and further pressed onto a 300-mesh stainless steel grid, which served as a current collector. The geometricarea of the electrodes was 0.55 cm2, and the typical thickness was 100~160 µm. Electrode capacity was calculated
according to the weight of active materials. To evaluate the electrochemical properties of the electrode, a half-cell
containing LiPF6 / EC+DMC (Ethylene carbonate plus diethyl carbonate as 1:1 in volume) electrolytes was used.
Li metal was utilized as both the counter and reference electrode. The structure of such a cell is similar to that of
Si based cells. Basically, all the three layers, including test electrode, separator and Li metal, were stacked in a
2025 coin type cell in a glove box. Unless stated elsewhere, cycling was started from discharge and carried out at
a potential cutoff at 1.4/0.01 V and a constant current density of 0.15 mA cm-2. The rest time between charge anddischarge was 1 min. Charge and discharge of the cell refer, respectively, to lithium extraction from, and insertion
into, the active hosts. The impedance response was measured with a Solartron SI 1260 (impedance / gain-phaseanalyzer) in a frequency range from 0.1 to 106 Hz.
2.3 Results and discussion

Li2.6Co0.2Ni0.1Cu0.1N

Li2.6Co0.1Ni0.1Cu0.2N

Li2.6Ni0.4N

Li2.6Co0.1Ni0.2Cu0.1N

Li2.6Cu0.2Ni0.2N

Li2.6Cu0.1Ni0.3N

Li2.6Cu0.4N

Li2.6Cu0.3Ni0.1N

Li2.6Co0.2Cu0.2N

Li2.6Co0.1Cu0.3N

Li2.6Co0.4N

Li2.6Co0.3Cu0.1N

Li2.6Co0.2Ni0.2N

Li2.6Co0.3Ni0.1N

Li2.6Co0.1Ni0.3N

Lattice parameter / A

Li2.6Ni0.4N

Intensity / a.u.

2.3.1 The doped elementals upon the electrochemical behavior of the ternary Li2.6Co0.4N
3.9
The composition of the resulting compounds
is confirmed by Inductively Coupled Plasma
a-axis(ICP). Fig.13 shows X-ray diffraction patterns ofc-axis3.8
the prepared lithium transition metal nitrides. The
XRD responses are well in accordance with those
3.7
of Li2.6Co0.4N recorded on JCPDS card 5-605,
indicating a typical hexagonal structure,
3.6
isostructural with Li3N, in which part of lithiumin Li3N is substituted by Co, Cu, Ni or Fe. No
impurity phase, such as Li2O, residual Li3N andmetals, can be detected. The collected lattice
Cu,Ni,Co
Li3N Li2Oaparameters of the hexagonal compounds arebcdshown in the inset in Fig.13. Due to the atomefg
radius of the doped metals, e.g., Ni<Co<Cu, theha-axis gradually decreases from Li2.6Ni0.4N to
20
30
40
50
60
Li2.6Co0.4N, until Li2.6Cu0.4N. As a result, the2 θ / deg.
affinity between the Li-M (M=Cu, Co, Ni) layer
1+3+(2) and the Li -N -Li layer (1) is decreased with
Fig.13. XRD patterns of the hexagonal lithium metal
the increasing of the affinity within the
nitrides: a) Li2.6Co0.4N; b) Li2.6Co0.2Cu0.2N; c)
Li1+-N3-Li+ layers (1). The a-axis increases again
Li2.6Co0.2Ni0.2N; d) Li2.6Cu0.2Ni0.2N; e) Li2.6Co0.25Cu0.15N;in the range from Li2.6Cu0.4N to Li2.6Ni0.4N. Thef) Li2.6Co0.2Cu0.1Ni0.1N; g) Li2.6Co0.25Cu0.1Ni0.05N;
c-axis is changed along with the oppositeh)Li2.6Co0.2Cu0.15Fe0.05N. The inset figure is Latticedirection of a-axis. When Co content is 0.2, bothparameters of the lithium metal nitrides.
a and c axis show the largest value in the case of
Co, Cu and Ni coexist.
As found, the electrochemical behavior of the Li2.6Co0.4N is superior to those of Li2.6Ni0.4N and Li2.6Cu0.4N.
To keep the charge balance of Li2.6Co0.4N upon oxidation and reduction state, both cobalt and nitrogen play animportant role. Fig.14 shows charge and discharge curves of the selected 1) Li2.6Co0.4N, 2) Li2.6Co0.2Cu0.2N, 3)
Li2.6Co0.2Cu0.1Ni0.1N and 4) Li2.6Co0.2Cu0.15Fe0.05N electrodes at different cycle. Due to the vacant sites introducedby the existence of Co2+, the fully lithiated nitrides still permit small amount of lithium intercalation into in thefirst cycle, as shown in Fig.14a. The compounds gradually undergo an irreversible transformation from a crystal toan amorphous phase in the initial Li-extraction stage, resulting in the noticeable potential plateaus at 1.0, 1.2, 1.1
and 1.1 V for Li2.6Co0.4N, Li2.6Co0.2Cu0.2N, Li2.6Co0.2Cu0.1Ni0.1N and Li2.6Co0.2Cu0.15Fe0.05N, respectively.
Compared with Li2.6Co0.4N, the obvious increases of Li2.6Co0.2Cu0.2N, Li2.6Co0.2Cu0.1Ni0.1N andLi2.6Co0.2Cu0.15Fe0.05N in the charge potentials indicate that part Co substituted by Cu, Ni and Fe may be
obstructive to the lithium extraction. At the fifth cycle, the Li2.6Co0.4N and Li2.6Co0.2Cu0.15Fe0.05N electrodes asshown in Fig.14b have a small potential fluctuation at ca.0.9 V and 1.3 V in the Li-extraction stage. Thisconsequence is related to an atom rearrangement associated with the doped elementals, Li and N.

Novel negative electrode materials

1.6

2 4 3

a) Cycle 1

1.4

Dis. Capacity / mAh g

1

1.2
1.0
0.8

1) Li2.6Co0.4N2) Li2.6Co0.2Cu0.2N3) Li2.6Co0.2Cu0.1Ni0.1N4) Li2.6Co0.2Cu0.15Fe0.05N

0.6
0.4
0.2

Voltage / V

0.0
-0.2
0

200

400

600

800

1.6

b) Cycle 5

1.4

3

2

4

1000

1

1.2
1.0
0.8
0.6
0.4
0.2
0.0
-0.2
0

200

400

600

800

1000

-1

Capacity / mAh g

Fig.14. Charge and discharge profiles of thelithium metal nitrides electrodes at a) cycle 1 andb) cycle 5.
50-Z'' / ohm

40

Li2.6Co0.4N
Li2.6Co0.2Cu0.2N

30
20
10
0

-1

Cycling comparison reveals that the co-doped
compounds have a remarkably enhanced cyclability, as
shown in Fig.15. By optimizing the tradeoffs for capacity
and cycling in accordance with the ternary Li2.6M0.4N, wehave confirmed that the appropriate compositions in thisfamily should be in the range, which can be stated by aformula as Li2.6Co0.4-x-yCuxNiyN (0.1≤x≤0.2, 0≤y≤0.1,
0.1≤x+y≤0.2). Furthermore, a certain amount of doped Fe
was found to increase the capacity and is in favor for thecycling. Compared to Li2.6Co0.4N, the obviously improved
cyclability of the co-doped compound may be partly
attributed to the enhanced interfacial compatibility
between the active particles and the electrolytes. The inset
in Fig.15 shows Cole-Cole plots of the impedance
response of Li2.6Co0.4N and Li2.6Co0.2Cu0.2N electrodes
with the delithiated stage vs. Li after 30 cycles. With partCo substituted by Cu, the total resistances of the half cellsignificantly turn into lessen, resulting in not only an
improved interfacial compatibility, but also an enhanced
capacity retention. In addition, according to the capacity
utilization in Li2.6Co0.4N and Li2.6Co0.2Cu0.2N, thecomparatively low Li-extraction level in Li2.6Co0.2Cu0.2N isfavorable for the cycling stability. We have observed thatthe capacity fading of the ternary Li2.6-zCo0.4N can beeffectively suppressed if the Li-extraction is limited atz≤1.0.
To further unveil the reliance of the electrochemical
behavior upon the structure change along with
Li-extraction, the variation of the lattice parameters of
Li2.6-ZCo0.4N and Li2.6-ZCo0.2Cu0.2N with various
charge-discharge depths in the scale of0<Z<1.0
in
the
first
Li-insertion/extraction process was
1200
measured and shown in Fig.16. The
pointed sites of a~e in a1 and b1 indicate
1000
the special Li-insertion/extraction stage
while
the
corresponding
lattice800
parameters are shown in a2 and b2. In
the charge process, it is obvious that the600
a-axis decreases with the c-axis
increases, relating to the composition
400
range where the potential gradually
increases to 1.0 V for Li2.6Co0.4N.
200
Above z~0.3, no remarkable change inthe cell parameter is observed until0
z=1.0, but decline in the XRD
0
responses peak turn into noticeable. It

19

0 10 20 30 40 50 60 70 80 90 100

Z ' / ohm

20

40

Li2.6Cu0.4N
Li2.6Co0.4N
Li2.6Ni0.4N
Li2.6Co0.2Ni0.2N
Li2.6Co0.2Cu0.2N
Li2.6Cu0.2Ni0.2N
Li2.6Co0.3Cu0.1N
Li2.6Co0.3Ni0.1N
Li2.6Co0.25Cu0.15N
Li2.6Co0.2Cu0.1Ni0.1N
Li2.6Co0.25Cu0.125Ni0.075N
Li2.6Co0.25Cu0.1Ni0.05N
Li2.6Co0.25Cu0.05Ni0.1N
Li2.6Co0.15Cu0.1Ni0.1Fe0.05N
Li2.6Co0.2Ni0.15Fe0.05N
Li2.6Co0.2Cu0.15Fe0.05N
Li2.5Ni0.45Fe0.05N
Li2.5Ni0.4Cu0.05Fe0.05N
60

Cycle number

suggests that the Li-extraction is
Fig.15. Cycling performance of the lithium metal nitrides electrodes.
associated with the two-phase reaction
The inset figure is Cole-Cole plots of the impedance response of theand the crystalline structure finallya) Li2.6Co0.4N and b) Li2.6Co0.2Cu0.2N with the charged state after 30
becomes amorphous. In the case ofcycles.
Li2.6-ZCo0.2Cu0.2N, the slight decreasingof a-axis along with the increasing of c-axis occurs much early, and there is no change in the lattice parameter
above z~0.2. However, we notice that the crystalline response peaks, which are linked to the phase transformation,
apparently reduce. It indicates that the formation of an amorphous phase appears in the very early charge region.
As a result, lithium extraction from the Li2.6-ZCo0.2Cu0.2N becomes difficult, leading to a decreased capacity
compared to that of the Li2.6-zCo0.4N.

20

Y. Liua, T. Matsumura, A. Hirano, T. Ichikawa, N. Imanishi and Y. Takeda

Volatge / V

Lattice parameter / A

2.3.2 The effects of the morphology performance upon the electrochemical behavior
3.90
1.6
Morphology performance of the particles after was
a1
a2
found to show obvious effects upon the
1.4
3.85
electrochemical behavior. Fig.17 shows charge and1.2
discharge curves of Li2.6Co0.2Cu0.2N without and witha-axis3.80
ec-axisd
1.0
HEMM. The initial insertion capacities weresignificantly enlarged when the particles were treated
3.75
0.8
c
d
under HEMM for 20h. Compared with the un-milled
e
a b
c
0.6
3.70
samples, the milled nitrides had an obviouslyb0.4
decreased ratio of Li vs. doped elementals from theLi2.6Co0.4N
3.65
ICP measurements. It appears to be related to the0.2
Chargeincreased Li-free sites and the valance change of part
a
0.0
3.60
0
200
400
600
800
1000
-0.2 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6
Co. Thus, we assumed that HEMM enables the
Capacity / mAh g
Lithium extraction / Z in Li Co N
removal of part lithium in the Li-rich compounds. A
3.90
1.6
decrease in the first extraction potential plateau
b2
b
1.4
1
3.85
probably is attributed to that HEMM make theea-axis1.2
dcrystalline lithium transition metal nitrides partly
c-axis3.80
amorphous. The milled samples have an increased
1.0ccapacity and a decreased Li-extraction potential. This
e
d
3.75
0.8
a bc
can be explained by the reduced potential polarization
0.6barising from the improved electrochemical kinetics,
3.70
e.g., the shorten lithium diffusion path and the0.4
Li2.6Co0.2Cu0.2N
3.65
improved interfacial charge transfer that relates to theCharge0.2
enlarged reactive area of the insertion hosts. The
a
0.0
3.60
HEMM alleviates the obstruction to the lithium
0
200
400
600
800
1000
-0.2 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6
Capacity/mAhg
Lithium extraction / Z in Li Co Cu N
extraction; thereby increases the Li-extraction degree
at a certain extent. A comparable cycling stability ofthe milled lithium transition metal nitrides is available
Fig.16. Variation of lattice parameters of the a)in accompany with a deep Li-extraction, as shown in
Li2.6Co0.4N and b) Li2.6Co0.2Cu0.2N with various
the inset in Fig.17. However, the Li2.6Co0.2Cu0.2N
Li-insertion/extraction depths in the first cycle.
directly synthesized by HEMM does not bring the
3.0
same consequence. It is probably due to the1000
incomplete reaction and poor atom distribution.
800a
We have observed a lot of residual Li3N and
2.5
600
b
transition metals in the XRD patterns for thec400
sample prepared by HEMM. This means thata) by Heating/HEMM
2.0200b) by Heating
high-energy mechanical milling step alone isc) by HEMM
0
insufficient for achieving the good crystalline
0 10 20 30 40 50 60 70 80
1.5
Cycle number
formation for the multi-doped compounds.
1
-1

0.4

2.6-Z

0.2

Volatge / V

Lattice parameter / A

2.6-Z

-1

Voltage / V

Dis.capacity / mAh g

-1

0.2

1

1.0
2
2.3.3 Novel composite electrodes based2
Without milling
lithium transition metal nitrides
With milling
0.5
1. Carbonaceous materials-lithium metal
nitrides composites
2
2
MCMB is a kind of graphitic carbon and10.0
1
currently used as anode material in the
-100
0
100
200
300
400
500
600
700
800
900
commercial Li-ion batteries. The composite-1
Capacity / mAh g
consisting of MCMB and Li2.6Co0.4N wasdeveloped via HEMM. SEM observation on thecomposite reveal that after ballmilling theFig.17. Charge and discharge curves of the electrodes based
Li2.6Co0.4N powders are severely smashed intoon Li2.6Co0.2Cu0.2N without and with HEMM at the first andnano- or submicro- scale, and dispersed andthe second cycle. The inset figure is Cycling performance ofadhered compactly on the surface of the MCMB
the electrodes of the Li2.6Co0.2Cu0.2N under different
particles, as shown in Fig.18. Such morphology
preparations.
is favorable for the electrochemical kinetics of
the Li2.6Co0.4N particles by adopting MCMB as an appropriate electron-conductive matrix. MCMB has a low
reactive potential at ca.0.1 V vs. lithium and a reversible capacity of 310 mAh g-1. As compared, we canapparently observe the increased initial coulombic efficiency (100% vs.89%) and the increased reversible capacity

Novel negative electrode materials

21

Dis.capacity / mAh g-1

Voltage / V(450 mAh g-1 vs.310 mAh g-1) for
the Li2.6Co0.4N-MCMB composite
electrode, as shown in the inset in
Fig.19. The high faradic yield in
the first cycle is dependent on theextraction capacity compensationof Li2.6Co0.4N. Compared with
MCMB, the composite has arelatively enhanced safety due tothe improved reactivity potential.
The potential trends of thecomposite electrode reflect theFig.18. SEM photos of a) the mechanical mixture of MCMB-Li2.6Co0.4N
mixing potential characteristics ofandb) the HEMM MCMB-Li2.6Co0.4N composite prepared by HEMM.
these two types of active hosts.
During cycling the Li2.6Co0.4N
1600
yields charge potentials mainly in the realm of
2.0
composite, η1=100%
0.6-1.1V vs.Li/Li+ and discharge potentials of
1400
MCMB, η1=89%+
1.5
0.15-0.5 V vs.Li/Li . As a result, an obviously
1200
potential hystersis of the composite electrode1.0
between charge and discharge is still remarkable.
1000
0.5
Fig.19 shows the cycling performance of theelectrodes based on the MCMB-lithium metal
800
0.0
MCMBnitrides composite and the pure MCMB. The
0
100
200
300
400
500
Co
Cu
Fe
N
Li
600
2.6
0.2
0.15
0.05
-1
capacity-cycle curves deliver high stability,
Capacity / mAh g
indicating a high reversibility of Li-insertion and400
extraction.
200

Intensity / a.u.

Voltage / V



X



X

X

X

X

Dis. Capacity / mAh g-1

MCMB-Li2.6Co0.2Cu0.2N
MCMB-Li2.6Co0.4N
MCMB
2. Si-C-Li2.6Co0.4N composites
0
Fig.20 shows charge and discharge profiles of
0
6
12
18
24
30
36
the Si−graphite−Li2.6Co0.4N composite synthesized
Cycle number
from two HEMM processes. In the first discharge,
the electrochemical behavior of the composite is only
Fig.19. Cycling performance of the lithium metal
associated to silicon and graphite because fully
nitrides-MCMB composites and the MCMB electrodes.
lithiated Li2.6Co0.4N has no extra vacancy for the
The inset figure is charge and discharge curves of theinitial Li intercalation. The capacity above 0.8 V
Li2.6Co0.4N-MCMB composite and the MCMB
from Li2.6Co0.4N plays an important role for theelectrodes at the first cycle.
compensation for the irreversible capacity. As a4.0
result, a high first-cycle efficiency of ca. 90 %
(1) 
(2)
 silicon
1200
graphite
was obtained. From the second cycle to the
3.5
Li2.6Co0.4N 800
subsequent cycles, the voltage trends show high
3.0coincidence that indicates a good reversibility of400a
Li insertion and extraction. Moreover, graphite
0
2.5
b
0
2
4
6
8
10 12
in the composite is favorable for reducing the
Cycle numbercpotential hysteresis arising from silicon and
2.0
Li2.6Co0.4N in charge and discharge. An XRD
1.5
measurement was further conducted on the
20
30
40
50
60
2θ / deg
compositewithrespect
tovarious1.0
charge-discharge stages, as shown in the inset
Cycle 1
Cycle 2
(1) in Fig.20. Before cycling, the diffraction
0.5
Cycle 6peaks reflected from silicon, graphite and
0.0
Li2.6Co0.4N can be distinguished (see a). The
0
200
400
600
800
1000
1200
diffraction peaks from silicon and Li2.6Co0.4N-1disappeared with the first discharge and charge,
Capacity / mAh g

Fig.20. Charge and discharge curves of the Si−graphite−Li2.6Co0.4N composite electrode at different cycles.
The inset figure (1) is XRD patterns of the Si−graphite−Li2.6Co0.4N composite electrode: (a) before charge anddischarge, (b) after first discharge and (c) after first discharge. The inset figure (2) is cycling performance of
the Si−C−Li2.6Co0.4N composite.

22

Y. Liua, T. Matsumura, A. Hirano, T. Ichikawa, N. Imanishi and Y. Takeda

Vol.capacity / mAh cm-3

as shown in b and c, respectively. The changed crystalline reflections remained same upon increasing of the cycle
number. This result means that the electrochemical Li insertion and extraction make silicon and Li2.6Co0.4N turninto amorphous. A SEM photo analysis reveals that all the components in the composite after HEMM tend to form
large agglomerates, indicating that silicon could be highly dispersed and embedded within the ductile
graphite−Li2.6Co0.4N matrix. Obviously, such a granular structure is in favor for increasing the interfacial affinityof silicon with the matrix and leading to an ensured electrical contact. On the other hand, the low absolute volume
effect of the composite using Li2.6Co0.4N reduces the mechanical stress of the dispersed silicon. Thus, thecomposite demonstrates good cycling stability, as shown in the inset (2) in Fig.20. In comparison, siliconpossesses the highest capacity of ca. 3.7 Ah g-1 in the first Li intercalation; but it suffers from a remarkable loss inthe extraction capacity. By contrast, Li2.6Co0.4N has a very low insertion capacity and very high extraction
capacity. The Si−graphite−Li2.6Co0.4N composite shows a large reversible capacity of ca. 1 Ah g-1 as three times
as graphite and a high initial faradaic yield of 90 % that might lead to a high energy density of the lithium ion
batteries.
3.0
1200
1000

Voltage / V

3. LiTi2O4-lithium metal nitrides composites
2.5800
The spinel LiTi2O4 demonstrates a high600
conductivity
andanextremely
constant
2.0
400
Li-insertion
charge/discharge potential versus lithium at ca.1.55V,
Li-extraction200-1
as well as a reversible capacity of ca.170 mAh g .
1.5
0
0
10
20
30
40
50
60
70
80
The Li-intercalation potential of LiTi2O4 (ca.1.50V)
Cycle number
is slightly higher than the Li-extraction potential of
1.0
Cycle 1Li2.6Co0.4N (ca.1.0V); therefore LiTi2O4 can extract
Cycle 2lithium from Li2.6Co0.4N. Open potential of 0.95 V
Cycle 100.5
Cycle 20can be observed for the electrode when lithium in theLi2.6Co0.4N was completely extracted. The suitable
0.0
weight ratio of the nitrides to the LiTi2O4 in the
0
100
200
300
400
500
600
700
800
900
electrode is dependent on the effective capacity and-3Volumetric capacity / mAh cm
the molecular weight of the two active hosts. A
Fig.21.
Chargeand discharge profiles of thepossible chemical reaction between the two hosts is
Co
N-LiTiOLi
2.6
0.4
2 4 electrode at different cycles. The“Li2.6-xCo0.4N + z LiTi2O4 →Li2.6-x-zCo0.4N + z
inset
figureisthecycling performance.
Li2Ti2O4 (x + z ≤ 1.6)”, where x means the lithium
loss in the Li2.6Co0.4N compared to the theoretical
Li:Co ratio as 2.6:0.4, and z is the possibly extracted lithium related to LiTi2O4. Thus, in view of the fact that the
introducing of HEMM enables the removal of part lithium in Li2.6Co0.4N and LiTi2O4 has no capacity contribution
in the electrode within a special potential range, the capacity utilization of the milled Li2.6Co0.4N can be improvedby reducing the contents of LiTi2O4. In the case of 22 wt% milled Li2.6Co0.4N plus 70 wt% LiTi2O4, the first cycle
efficiency is close to 100%. Fig.21 shows the charge and discharge curves of the electrode based on
Li2.6Co0.2Cu0.2N and LiTi2O4 at different cycles. In general, a decreased particle size leads to an increased density.
For example, the density of the mixture consisting of the separately milled Li2.6Co0.2Cu0.2N and LiTi2O4 is 3.2 g
cm-3 according to the experimental data, versus 2.2 g cm-3 of the mixture based on the un-milled samples. Largevolumetric capacity is dependent on the material’s high specific capacity and density. In addition, LiTi2O4 can act
as an inert-conducting additive in place of AB in the electrode within a potential range of 1.4V~0.01V vs.Li/Li+.
Thus, the composite electrode shows a comparative volumetric capacity. Taking Li2.6Co0.2Cu0.2N and LiTi2O4 as
the active hosts, the volumetric capacity of the composite electrode is 798 mAh cm-3, which is attractive comparedto that of graphite (ca. 650-700 mAh cm-3). The Li2.6Co0.2Cu0.2N-LiTi2O4 composite electrode exhibits highcycling stability, as shown in the inset in Fig.21.
4. Co3O4-lithium metal nitrides composites
Fig.22 shows charge and discharge curves of the Li/Co3O4 (a), Li/Li2.6Co0.2Cu0.2N (b) and
Li/Co3O4-Li2.6Co0.2Cu0.2N (c) cells. The Co3O4 electrode had a large insertion capacity of 1100 mAh g-1 alongwith a potential plateau of 1.1 V in the first cycle. However, its capacity retention at the subsequent cycle was
below to 180 mAh g-1 within the potential range of 0.01-1.4V. The electrochemical behavior of Li2.6Co0.2Cu0.2N
was in contrast to that of Co3O4, which had a high extraction capacity of over 800 mAh g-1 with a potential plateau
of 1.1 V but a very low first insertion capacity of around 200 mAhg-1. In contrast, 100 % capacity retention at thefirst cycle was observed in the Co3O4-Li2.6Co0.2Cu0.2N composite electrode. XRD studies of the
Co3O4-Li2.6Co0.2Cu0.2N composite on different charged states, as shown in Fig.23, revealed that the diffraction
peaks from Li2.6Co0.2Cu0.2N and Co3O4 disappeared by soaking the composite in the electrolyte for 2h. The
diffraction intensities remained constant with increasing charge depth and number of cycles. These results indicate

Novel negative electrode materials

23

Voltage / V

that a thermodynamically spontaneous reaction between Co3O4 and Li2.6Co0.2Cu0.2N leads to an amorphous
delithiated state of Li2.6Co0.2Cu0.2N and probably to form metallic Co and Li2O. We propose that the reaction
mechanism of the composite anode as follows:(1)
Li2.6Co0.2Cu0.2N + xCo3O4 → Li2.6-8xCo0.2Cu0.2N + 4xLi2O + 3xCo
(2).
Li2.6-8xCo0.2Cu0.2N + zLi+ + ze- → Li2.6-8x+zCo0.2Cu0.2N
Reaction (1) is a chemical reaction and reaction (2) is an
2.5
electrochemical one. If Li2.6-8xCo0.2Cu0.2N is thea. Co3O4
Cycle 1
2.0
composition of the fully charged state, then reaction (2)Cycle 2Li extraction
1.5
becomes reversible during charge and discharge even at
1.0
the first cycle. We estimated x and z as 0.16 and 1.7,
respectively. These values correspond to 60/40 weight
0.5
Li insertion
ratio of Li2.6-8xCo0.2Cu0.2N/Co3O4 and the reversible
0.0
capacity of 477 mAh g-1. We observed a first charge and
0
200
400
600
800
1000 1200
2.5
discharge efficiency of 100% and a reversible capacity ofb. Li2.6Co0.2Cu0.2N
ca. 500 mAh g-1 for the composite anode of 62.5 wt %2.0
Li2.6-xCo0.2Cu0.2N plus 37.5 wt % Co3O4 as shown in Fig.
1.5
Li extraction
1c. The high first cycle efficiency proposed with this1.0
anode can greatly extend the capacity utilization for the
0.5
cathode. It is remarkable because the capacity of the
Li insertion
conventional cathode materials in lithium batteries is less
0.0
-1
0
200
400
600
800
than 200 mA g . The irreversible anode capacity in the
2.5
first cycle requires an extra capacity of cathode active
c. Li2.6Co0.2Cu0.2N-Co3O4
2.0
materials; thereby it inevitably depresses the total energy
density of the batteries. Due to lithium metal nitride as the
1.5
Li extraction
active hosts, such a composite demonstrate excellent
1.0
capacity retention with cycling, as shown in the inset in
0.5
23.
Li insertion

0.0

2.4

Conclusions

0

100

200

300

400

500

600

-1

Capacity / mAh g

Intensity / a.u.

Dis. capacity / mAh g-1

In summary, we have developed a series of lithium
Fig.22. Charge-discharge curves of the cells: a)mixed transition metal (co-doped) nitrides with high
Co3O4 / electrolyte / Li, b) Li2.6Co0.2Cu0.2N /electrochemical capacity and good cyclability. For
Electrolyte / Li and c) Co3O4 - Li2.6Co0.2Cu0.2N
producing these compounds, a combination of a solid-state(37.5 - 62.5 wt. %) / electrolyte/ Li at room
reaction under an appropriate temperature and atemperature.
high-energy mechanical milling step was800
involved. Research reveals that the granular
700
Li2.6Co0.2Cu0.2N-Co3O4
structure of the nitrides shows obvious effects
600
500
upon the electrochemical behavior. The
400
lithiated compounds can be used in several
300
ways to form new composite electrodes which
200 Li2.6Co0.2Cu0.15Fe0.05N-Co3O4d) Charge to 1.4 V
demonstrate high insertion capacity, 100 % first
100c) Discharge to 0.01 V
cycle efficiency and excellent capacityb) Soaking in electrolyte
0
10
20
30
40
50
retention ability. Therefore, they are promisinga) Before soaking
Cycle number
anode candidates for further Li-ion batteries.
Co O
3

d

c
b
a

Part II Solid-state polymer electrolytes
based system

Li2.6Co0.4N

1. Introduction

20

Co

4

Co3O4

Li2.6Co0.4N3040

2 θ / deg.

Co3O4 Li2.6Co0.4N
50

Since the discovery of ionic conductivity
Fig.23. XRD patterns of the Co3O4-Li2.6Co0.2Cu0.2N
in the Poly (ethylene oxide) (PEO)-lithium salt
composite electrode at different stages. The inset figure iscomposites, lithium-ion batteries based on such
cycling performance of the composites.
category of electrolytes has attracted
worldwide attentions. In view of the highenergy density and feasible design, all solid-state PEO rechargeable lithium batteries could be promising and ideal

24

Y. Liua, T. Matsumura, A. Hirano, T. Ichikawa, N. Imanishi and Y. Takeda

power source for EV&HEV [52]. So far, research has been extensively concentrated on two major points: theimprovement of ionic conductivity of PEO-LiX complex at the relatively low temperature and the modification ofinterfacial characteristics between metallic lithium anode and polymer electrolytes. It is generally accepted that
there is some instable issue between metallic lithium and PEO electrolytes caused by lithium dendrite grown[53,54]. Insertion hosts such as carbon, Li-alloy and Li-M-O, in place of lithium, present the promising alternative
anodes that can avoid the above problem [55,56]. Unfortunately, carbon still cannot be adopted successfully in the
PEO electrolytes due to the interfacial incompatibility, regardless of its wonderful electrochemical manners in theliquid system. On the other hand, some Li-M-O materials, e.g., spinel Li4Ti5O12, demonstrate zero-strain effects
for Li-insertion, resulting in long cycling life in the solid PEO electrolytes. Relatively low capacity and highLi-insertion potential of lithium titanates suffers loss of the comparable energy density. In contrast, for the large
lithium storage capacity and low operating voltage close to lithium, Li-alloy could be considered as theappropriate candidate. It is believed that the morphology instability, which caused by the volume change of active
hosts in association with lithium insertion, severely hinders the development of Li-alloy in both solid and liquid
electrolytes. Decreasing particle size and adopting of multi-phase structure is favor for greatly enhancing thecyclability. Some positive results have been obtained, such as Li-Al electrode and interfacial layers of Li-Bi and
Li-Pb alloys between lithium metal and solid polymer electrolytes [57,58]. However, the cycling performance ofLi-alloy as anode in the solid PEO electrolytes is still poor.
Recently, we have developed a novel lithium ion cell consisting of the composite anode, such as
SnSb-Li2.6Co0.4N, SiO1.1-Li2.6Co0.4N, Co3O4-Li2.6Co0.2Cu0.2N, and the LiNi0.8Co0.2O2 cathode based on the solid
PEO-LiN(CF3SO2)2 electrolytes [59-63]. Such cells can offer large reversible anode capacity at 450 mAh g-1 andhigh working voltage, as well as an ensured safety without metallic lithium. In this part, the electrochemical
performance of several types of novel anodes with the solid-state polymer PEO electrolytes are presented anddiscussed in detail.
2. Experimental
2.1 Active materials
LiCo0.2Ni0.8O2 (typical particle size is 10 µm) was prepared by the normal solid state reaction. Ultrafine
SnSb alloy powder (ca. 0.2 µm) was produced by chemical precipitation from aqueous solutions [12] and lithium
metal nitrides was synthesized as mentioned above. Nickel powder (99.9%, ca. 0.03 µm), PEO (MW=6 × 105),
LiN(CF3SO2)2, BaTiO3 (typical particle size, ca.0.1µm), PVDF (MW=534000) were commercially purchased. The
Co3O4 powder (< 2 µm) was obtained by a decomposition reaction of CoCO3 at 800 0C in air.
2.2 Polymer electrolyte and cathode film
All the procedure was carries out in the glove box protected by Ar. PEO electrolytes (Li/O ratio: 1/18) wereprepared under the casting technique. A given weight of PEO (MW = 6 × 105) and LiN(CF3SO2)2 was dissolved
completely in anhydrous acetonitrile (AN). BaTiO3 was dispersed homogeneously in the solution as filler. The
viscous was strongly stirred overnight and cast into a Teflon dish. After AN was slowly and completely
evaporated under N2 flow, the obtained film was further dried at 90 ℃ under vacuum at least 8 h. Cathode films
consisting of 52 wt% LiNi0.8Co0.2O2, 10 wt% acetylene black(AB) and 38 wt% PEO-LiN(CF3SO2)2 were preparedby the normal casting method similar to PEO film. The thickness of the cathode and PEO film was about 200 µm.
The conductivities of the composite polymer electrolyte were observed as high as 1.7x10-3 S cm-1 at 80 0C and
0.82x10-3 S cm-1 at 65 0C.
2.3 Anode preparations
Different preparations were adopted and nominated as (a) dry mixing, (b) wet mixing, (c) hot pressing and (d)casting, respectively. For the dry mixing way, the electrode components were mechanically mixed in an agate
mortar. In the case of method (b), the components were homogeneously mixed by the aid of Hexane as a disperser.
Polymer binder was added and mixed after Hexane was entirely evicted. Mixtures prepared by (a) and (b) werepressed onto a 280-mesh stainless steel grid, which served as a current collector. Normally, area and thickness ofthe composite anodes were 0.55 cm2 and 50-60 µm, correspondingly. Anode prepared by hot pressing was
described as follows: PEO powders were carefully sieved and only the smallest particles were passed and used.
Before hot pressing, the electrode components were homogeneously mixed and pressed directly onto Cu-foil (20µm) in a steel mold at a temperature of 70-90℃ and pressure of 10-60 MPa. The press time was 0.5 h-6 h. The
thickness of the obtained film was 60-90µm (including Cu-foil). The procedure of method (d) was similar to thatof the PEO film preparation. PEO and LiN(CF3SO2)2 (Li/O ratio: 1/18) were dissolved in AN solvent entirely.
SnSb and conducting Ni powders were added and followed a strong stirring. At last, Li2.6Co0.4N was added at 10

Novel negative electrode materials

25

min before the casting. The viscous solution was cast onto a 280-mesh stainless steel grid and further dried at
90 ℃ under vacuum for 2 h.
2.4 Cell assembling and electrochemical measurement
To examine the cycling performance of the composite anode, half-cell was used and metallic lithium wasutilized as counter electrode. The full cells were composed of the LiNi0.8Co0.2O2 cathode and the SnSb-Li2.6Co0.4N
composite anode. Basically all the three layers, including composite anode, PEO electrolyte and cathode film (orlithium counter electrode), were stacked in coin type cells. A small constant pressure was kept inside cells by themeans of the Ni foam as filler. Considering the corrosion instability of Al foil with LiN(CF3SO2)2, stainless plate
was used as cathode current collector. Thickness ratio of cathode film and current collector is about 1/3. Before
assembly, the cathode film was tightly press on the current collector under slightly heat and pressure. Unlessstated elsewhere, cycling properties tests were performed at a current density at 0.15 mA cm-2 and the working
temperature was 65 ℃. The voltage cutoff is 1.4-0.1V vsLi/Li+ and 2.1-3.6V for half cell and full cell respectively.
Before electrochemical test, the cells were preheated for 2 hrs at a temperature of 75℃. The discharge capacity
was adopted for all figs, corresponding to Li-insertion into the SnSb-Li2.6Co0.4N composite in half-cell andLi-extraction out from the SnSb-Li2.6Co0.4N composite in full cell, respectively.
3. Results and discussion

Dis.capacity / mAh g

Voltage / V-1

3.1 Solid-state PEO cells based on SnSb-Li2.6Co0.4N composite anodeDue to the drastic volume effects in association with Li intercalation into and extraction from SnSb hosts, theanode morphology stability, as well as the interfacial
4.0
700
compatibility between the composite anode and the
600
3.5
500
solid PEO electrolytes, dominate the cycling400
performance of the all solid-state cells. In general,
300
3.0
o
o
the increased area of electrode is favorable for the
60 C
65 C
200
o
o
70
C
75
C
100
electrochemical properties via decreasing the bulk
2.5
o
o
80 C
85 C
0
resistance of the cell. However, the non-uniform
o
-100
90 C
2.0
dispersion of anode components prepared by dry
-200
o
0
2
4
6
8 10 12 14 16
65 C, η1 = 117.5 %mixing method tends to be aggravationoCycle number
75 C, η1 = 107 %1.5
accompanied with the enlarged electrode area.
o
90 C, η1 = 98.1 %Non-uniform dispersion of lithium salt and nano-Ni
1.0
conducting powder will cause the relatively low firstcycle efficiency and the slightly increased voltage
0.5
polarization. In addition, ultrafine SnSb particles
0.0with non-uniform distribution will form the large
-100
0
100
200
300
400
500
600
700
800
aggregation region and easily result in the local-1
volume mismatch during charging and discharging.
Capacity / mAh g
At last, rapid capacity decline is inevitable.
Fig.24. The first charge and discharge profiles of
Research reveals that the composite anode prepared
composite electrodes at different operating temperature.
by hot pressing and casting way demonstrates
The inset figure is cycling performance of composite
severe deterioration compared to dry mixing way.
electrodes at different operating temperature.
Based on this, a solution stirring preparation by
the aid of a dispersant, e.g., hexane, was introduced to achieve the homogeneous distribution of the electrodecomponents. The cycling stability of the composite electrode prepared by the aid of solvent dispersion is
obviously superior to those from other preparations. The observed results suggest that the appropriate anodepreparation has an apparent effect upon the electrochemical behavior of the SnSb-Li2.6Co0.4N composite anode.
On the other hand, good electrode performance of the composite electrode requires a combination of suitable
polymer binders. Poly(vinylidene fluoride) (PVDF) or poly(vinylidene fluoride) -hexafluoropropylene(PVDF-HFP) as binder in lithium ion batteries demonstrates enough binding ability, high mechanical andchemical stability even at elevated temperature. The electrochemical stable window of fluorinated polymers is as
wide as 0 to 5 V vs. Li/Li+. Moreover, these polymers are easer to be dispersed and can improve the adhesion
between the active material and current collector. The mechanical strength of the composite electrode can begreatly reinforced by the use of PVDF as framework due to its strong binder ability for retraining themorphological instability in the Li-alloy processing, indicating that the function of fluorinated polymer andPEO-LiX as binder is mainly corresponding to sustain morphological stability and supply ionic channels withincomposite electrode respectively. Ionic conductivity predominates in the amorphous rather than crystalline regions

26

Y. Liua, T. Matsumura, A. Hirano, T. Ichikawa, N. Imanishi and Y. Takedao- Z '' / ohm

Voltage / Vof PEO electrolytes. The local relaxation and segmental motion of the polymer chains appears to be a requirement
for lithium ion transport. The effect of operating temperature on mechanical strength and ionic movement ofpolymer electrolytes is contradictory. Therefore, it is reasonable that the SnSb-Li2.6Co0.4N electrode performance
exhibits a strong dependence on the environment temperature, as shown in Fig.24, in which the coulombic
efficiency at the first cycle apparently decreases with increasing the temperature. It is noted that capacity lessen in
the lithium extraction process is mainly attributed to shortness of voltage plateau near 1.1 V, which is related toLi2.6Co0.4N active material. This indicates that the Li2.6Co0.4N may lose some activity at high temperature in
PEO-LiX system and it is the main reason for the reduced charge recovery in the subsequent cycles.
The deterioration of the mechanical performance of polymer PEO electrolyte at high temperature also takes
part consequence. The temperature as a function upon the capacity-cycling behavior was shown in the inset in
Fig.24. Below the melting point of PEO (ca.
465℃), the polymer electrolyte delivers enough
Cycle 2-10
mechanical strength and can hold goodinterfacial morphology stability between
3electrode and electrolyte. In addition, theCycle 1
relative rigidity of PEO-LiX at low temperature
within composite electrode is favorable for2
restraining the volume expands of SnSb hosts
550
500
for Li-insertion. However, low operating
After 1 h
450
After 10 days
400
After 16 h
temperature corresponds to a poor interface
70 C 60 C
350
After 48 h
1
300
80
C
contact and slow Li-transfer in both composite
After 96 h
250
90 C
After 114 helectrode and polymer electrolyte layers,
200
150
resulting in a small capacity and unsteady
100
o
50
60 C
0
cyclability. The capacity increases remarkably
0-50
near PEO melting point due to adhesive-100 0 100 200 300 400 500 600 700 800-100 0 100 200 300 400 500 600 700 800'
interface contact and high ionic conductivity. It
Z / ohm
Z / ohmseems that such a cell system works poorly if
-1
0
100
200
300
400
500
600
the operating temperature continually grows-1
over 75℃. At elevated temperature, sticky or
Capacity / mAh goooeven viscous property of the PEO-LiX in
Fig.25. Charge and discharge profiles of coin type cell using
combination with the drastic volume effect of
LiNi0.8Co0.2O2 cathode and composite anode in solid
active SnSb particles will bring about a server
electrolytes
at 65 . Voltage cut off: 2.1-3.6 V, ic=id=0.1 mA
mechanical instability of the electrode. The
-2
.
The
inset
figure (1) is impedance performances ofcmmorphological change could be aggravated andcomposite
electrodes
storage at 60 at different time. The
hardly be recovered for the PEO-containing
inset
figure(2)
isimpedance
performances of composite
system above 75℃ under strict pressure inside
electrodes
duringstorageatdifferent
operating temperature.
cells. The appropriate operating temperature for
the SnSb-Li2.6Co0.4N composite electrode is
between 65 and 75℃. However such electrodes demonstrate a quite stable interfacial contact with polymer
electrolytes during storage over a wide operating temperature. The interface resistance between composite
electrode and solid electrolytes is acceptable even at 60 ℃, as shown in the inset (1) in Fig.25, because the cell
was preheated at a temperature 10℃ higher than operating one before test. With increasing temperature, theimpedance is greatly reduced, suggesting a significant improvement in the interfacial contact, as shown in theinset (2) in Fig.25. Dispersing BaTiO3 in polymer binders may greatly reinforce the mechanical strength of
PEO-LiX within composite electrode and, as a consequence, improves the morphology and cycling stability ofwhole electrode during charging and discharging. The typical charge and discharge profiles of the solid PEO
Li-ion cell based on SnSb-Li2.6Co0.4N composite anode and the LiNi0.8Co0.2O2 cathode at 65℃ are given in Fig.25.
The LiNi0.8Co0.2O2 cathode can deliver a reversible capacity about 140 mAh g-1 at appropriate working voltage
scales in PEO-LiN(CF3SO2)2-BaTiO3 system, about one forth of that of the composite anodes. For full utilizationof lithium storage capacity of composite anode, the weight of cathode has to be much over that of anode (4-5times). A little increase of voltage plat from the first cycle to the following ones can be observed and it is mainly
linked to the phase transfer of Li2.6Co0.4N from crystal to amorphous state. It is attractive that a high coulombic
efficiency over 94 % and a large reversible anode capacity about 500 mAh g-1 at the first cycle can be obtained inthis system. Moreover, the capacity remains stable of above 400 mAh g-1 after 30 cycles under the limited lithium
insertion. The capacity fade during cycling could be greatly prevented via the suitable charge-discharge control
and the proper cell structure. A further improvement for cycling performance of the cells is still on theNovel negative electrode materials

27

progression.
3.2 Electrochemical behavior of SiO1.1-Li2.6Co0.4N anode with the PEO electrolytes
The charge and discharge profiles of the SiO1.1-Li2.6Co0.4N composite electrode with the PEO electrolytes atdifferent cycles are shown in the inset in Fig.26. The
700
insertion capacity was increased to about 550 mAh g-1
90 C,
85 C
600
75 C,
70 C
at the second cycle, vs. 420 mAh g-1 at the first cycle. It
65 C
is probably attributed to a reduced polarization caused
500
by the improved electrochemical kinetics at the400
electrode interface compatibility. The voltage trends of300
the composite reflect mixing potential characteristics of
1.6
1.4
SiO1.1 and Li2.6Co0.4N active materials. In fact, the200
Li Co N, 1st
1.2
capacity above 0.8 V from Li2.6Co0.4N plays an
1.0100
important role for the compensation for the irreversible
0.8
Cycle 30.6
0Cycle 2
capacity. As a result, a high first cycle efficiency of
Cycle 10.4
ca.100% was obtained. The electrode compositions-100
0.2
have an apparent influence upon the electrochemical
0.0
silicon, 1st
-200
0
100
200
300
400
500
600
behavior of the electrode, similar to that of the-1
Capacity / mAh g-300
SnSb-Li2.6Co0.4N electrode. Fig.26 shows the capacity
0
6
12
18
24
30
change as a function upon the cycle number of the
Cycle number
SiO1.1-Li2.6Co0.4N composite electrodes at various
Fig.26. Cycling performance of the composite
environment temperatures. The reversible capacity inelectrodes at different temperature. The inset figure
the cycling beginning was improved with increasing
is charge and discharge curves of thethe temperature due to the adhesive interface contact
SiO1.1-Li2.6Co0.4N composite electrode at different
and high ionic conductivity. However, highcycles.
temperature fast the capacity fading. It indicates that an
aggravated decay in the mechanical strength of the
PEO electrolytes can hardly endure the morphology instability arising from silicon. By contrast, PEO electrolyte
at low temperature possesses a high mechanical stability with a low ionic conductivity that leads to a small
capacity with good capacity retention. It was further found that dispersing sub-micro ceramic BaTiO3 within theelectrode is in favor for the cycling behavior of the composite electrode. Ceramic fillers, such as BaTiO3 andAl2O3, can improve the mechanical performance of the flexible PEO electrolyte and stabilize the interfacial
properties at elevated temperature. Furthermore, it can reduce the crystallizing tendency of the polymer
electrolyte; thereby it enhances the ionic conductivity. With temperature increasing, the solid PEO-LiN(CF3SO2)2
composite electrolyte turns into be serious sticky and fluid. The function of BaTiO3 is to enhance the mechanicalstrength of the electrolyte components within the composite electrode that could maintain the electrode integrity.
Research reveals that the suitable current density for this composite electrode in the solid PEO electrolytes is
between 0.1 and 0.5 mA cm-2 after tradeoff capacity and capacity retention.
o

o

o

o

2.6

0.4

Voltage / V

Dis.capacity/ mAh g-1o

3.3 Electrochemical behavior and thermal performance of the Co3O4-Li2.6Co0.2Cu0.2N
The Co3O4-Li2.6Co0.2Cu0.2N composite demonstrate reasonable interfacial compatibility with PEO
electrolytes and thus good electrochemical behavior. The capacity changes, as a function of cycling, for different
anodes with the solid PEO-based polymer electrolytes are shown in Fig. 27. The reversible capacity of graphite
was below 200 mAh g-1 and the charge and discharge efficiency in the first several cycles was extremely low.
Moreover, its capacity retention suffered from an obvious fading during cycling. The SnSb-Li2.6Co0.4N anode hada large capacity of over 600 mAh g-1 and high coulombic efficiency of ca. 100 % in the first cycle. However, agradual deterioration in the cycling performance was always observed due to the volume effects of alloys. The
capacity fade was significantly suppressed for the Li2.6-xCo0.2Cu0.2N-Co3O4 anode and estimated to be only about
0.37%/cycle during cycling. The charge and discharge profiles of the Li2.6-xCo0.2Cu0.2N-Co3O4 anode at 65 ℃ as
shown in the inset in Fig.27 indicate that the charge-discharge efficiency was almost 100% and lithium can beintercalated and extracted with a high reversibility. In general, the safety of small size lithium-ion cells undernormal use is well established. In contrast, the safety of the large size lithium-ion batteries is still questionable,
especially in cases of abusive use. Safety of lithium ion batteries is mainly related to the thermal reactivity of theelectrode components. We found that specific reaction heats, which indicates the thermal stability and safety of thebattery system, for the Li2.6Co0.2Cu0.2N-Co3O4 anode charged up to 0.01 V vs. Li/Li+ was estimated to be 1.1 J
mAh-1. Those for Li2.6Co0.2N-SnSb anode was 2.75 J mAh-1. The specific reaction heat for lithium metal and PEO
electrolyte (1:1 weight ratio) is 1.24 J mAh-1. However, in practice, for the case of a lithium metal anode, an
excess amount of lithium metal, at least four times compared to the cathode capacity, should be used because of

Y. Liua, T. Matsumura, A. Hirano, T. Ichikawa, N. Imanishi and Y. Takeda

3.4 Conclusions

2.5

1400

2.0

1000

Filled: Li insertion
Open: Li extraction

Voltage / V

1600

1200
-1

the dendrite formation on the anode.
Therefore, the reaction heat of lithium
metal with PEO electrolytes was estimatedto reach 4.96 J mAh-1 or higher. Therefore,
the Li2.6Co0.2Cu0.2N-Co3O4 anode shows
the lowest reaction heat, indicating that thenovel anode has potential applications in
the large size rechargeable lithium-ion
batteries for EV & HEV.

Capacity / mAh g

28

Cycle 1
Cycle 2
Cycle 6

1.5

η1= 99.6%

1.0
0.5

800

SnSb-Li2.6Co0.4N

600

0.0
0

100

200

300

400

500

600

-1

Capacity / mAh g

400

Anode material, which possesses
Graphite
Co3O4-Li2.6Co0.2Cu0.2N
200
large capacity density, high first cycle
efficiency and low reaction heating with the0
PEO electrolytes, is the key for developing
0
10
20
30
40
50
lithium polymer rechargeable batteries that
Cycle number
proposed for extremely demanding
applications such as electric vehicles (EV)Fig.27.
Cycling
performanceof
SnSb-Li2.6Co0.4N,
and hybrid electric vehicles (HEV). By this
O
-Li
Co
Cu
N
and
graphitewithPEOelectrolytes
vs. Li
Co
3
4
2.6
0.2
0.2
motivation, we have developed series novel
at
65℃.
The
inset
figureischargeanddischarge
profilesofthecomposite anode which show good
O
-Li
Co
Cu
N
electrode.
Co
3
4
2
6
0
2
0
2
electrochemical behavior with the PEO
electrolytes. We are at present investigating the performance of the all solid-state PEO cells with these proposednovel anodes and some typical high potential cathodes
Acknowledgements
Part of the research work was carried out under a collaboration program of Mie University and Genesis
Research Institute (Nagoya), and part was sponsored by new Energy and Industrial Technology Development
Organization, Japan.
References[1] J.-M. Tarascon and M. Armand, Nature 414,359-367 (2001)[2] S. Eckfeld, J.E. Manders, M.W. Stevenson, J. Power Sources 5105 (2003) 1-8[3] S. Levaesseur,M.Menetrier,C. Delmas, J. Electrochem. Soc. 149 (2002) A1533.
[4] A.K. Padhi, K.S. Nanjundaswamy, J.B. Goodenough, J. Electrochem.Soc. 144 (1997) 1188
[5] K. Sato, M. Noguchi, A. Demachi, N. Oki, M. Endo, Science 264 (1994) 556[6] M. Winter, J.O. Besenhard, Electrochimica Acta 45, (2003) 31[7] M. Nishijima, T. Kagohashi, N. Imanishi, Y. Takeda, O. Yamamoto and S. Kondo, Solid State Ionics, 83(1996) 107[8] P. Poizot, S. Laruelle, S. Grugeon, L. Dupont, J-M. Tarascon, Nature 407 (2000) 496[9] Y. Idota, T. Kubota, A. Matsufuji, Y. Maekawa, T. Miyasaka, Science 276 (1997) 1395
[10] A. N. Dey, J. Electrochem. Soc., 118, 1547 (1971).
[11] J. O. Besenhard, M. Hess and P. Komenda, Solid State Ionics, 40/41 (1990) 525[12] J. Yang, M. Winter and J. O. Besenhard, Solid State Ionics, 90 (1996) 281[13] W.B. Xing, A.M. Wilson, K. Eguchi, G. Zank, J.R. Dahn, J. Electrochem. Soc. 144, 2410 (1997)[14] A.M. Wilson, G. Zank, K. Eguchi, W. Xing, J.R. Dahn, J. Power Sources 68, 195 (1997)[15] H. Kim, J. Choi, H.J. Sohn, T. Kang, J. Electrochemical Society 146, 4401 (1999)[16] H. Li, X.J. Huang, L.Q. Chen, G.W. Zhou, Z. Zhang, D.P. Yu, Y.J. Mo, N. Pei, Solid State Ionics 135, 181(2000)[17] A. Netz, R.A. Huggins, W. Weppner, J. Power Sources 5271, 1 (2003)[18] O. Mao, R.L. Turner, I.A. Courtney, B.D. Fredericksen, M.I. Buckett, L.J. Krause, J.R. Dahn, Electrochem.
Solid-State Lett., 2 (1999) 3[19] II-S. Kim, P.N. Kumta, G.E. Blomgren, Electrochem. Solid-State Lett., 3 (2000) 493[20] II-S. Kim, G.E. Blomgren, P.N. Kumta, Electrochem. Solid-State Lett., 6 (2003) A157
[21] II-S. Kim, G.E. Blomgren, P.N. Kumta, J. Power Sources, 130 (2004) 275

Novel negative electrode materials

29[22] H.Y. Lee, S.M. Lee, J. Power Sources, 112 (2002) 649[23] N. Dimov, K. Fukuda, T. Umeno, S. Kugino, M. Yoshio, J. Power Sources 5030, 1 (2002)[24] X.Q. Yang, J. McBreen, W.S. Yoon, M. Yoshio, H.Y. Wang, K.J. Fukuda, T. Umeno, Electrochemistry
Communications 4, 893 (2002)[25] N. Dimov, S. Kugino, M. Yoshio, Electrochimica Acta 48, 1579 (2003)[26] A.M. Wilson, J.R. Dahn, J. Electrochem. Soc. 142, 326 (1995)[27] D. Larcher, C. Mudalige, A.E. George, V. Porter, M. Gharghouri, J.R. Dahn, Solid State Ionics 122, 71 (1999)[28] Y. Liu, K. Hanai, J. Yang, N. Imanishi, A. Hirano, Y. Takeda, Solid State Ionics, 168 (2004) 61.
[29] Y. Liu, K. Hanai, J. Yang, N. Imanishi, A. Hirano, Y. Takeda, Electrochem. Solid-State Lett., 7 (2004) A369.
[30] Y. Liu, K. Hanai, T. Matsumura, N. Imanishi, A. Hirano and Y. Takeda, Electrochem. Solid-State Lett., 7(2004) 12[31] K. Hanai, Y. Liu, N. Imanishi, T. Ichikawa, A. Hirano and Y. Takeda, J. Power Sources, in press[32] Y. Liu, K. Hanai, N. Imanishi, A. Hirano, T. Matsumura and Y. Takeda, Solid State Ionics, in press.
[33] G.X. Wang , L. Sun, D.H. Bradhurst, S. Zhong, S.X. Dou, H.K. Liu, Journal of Alloys and Compounds 306(2000) 249[34] J. Yang, Y.Takeda, N.Imanishi, C.Capiglia, J.Y.Xie, O.Yamamoto, Solid State Ionics 152-153 (2002) 125[35] M.N. Richard, J.R. Dahn, J. Electrochem. Soc. 146 (1999) 2068
[36] J. Yamaki, H. Takatsuji, T. Kawamura, M. Egashira, Solid State Ionics 148 (2002)241[37] M. Nishijima, T. Kagohashi, N. Imanishi, Y. Takeda, O. Yamamoto, J. Power Sources, 68 (1996) 510[38] S. Suzuki, T. Shodi and J. Yamaki, J. Phys. Chem. Solids, 59 (1998) 331[39] T. Shodi, S. Okada, S. Tobishima and J. Yamaki, J. Power Sources 68 (1997) 515[40] T. Shodi, Y. Sakurai and T. Suzuki, Solid State Ionics 122 (1999) 85[41] T. Y. Kim, M. G. Kim, J. M. Lee, T. Kang and H. J. Sohn, Electrochemical and Solid-State Letters, 5 (2002)
A103[42] Y. Takeda, M. Nishijima, M. Yamahata, K. Takeda, N. Imanishi and O. Yamamoto, Solid State Ionics, 130(2000) 61[43] J. Yang, Y. Takeda, N. Imanishi and O. Yamamoto, J. Electrochem. Soc., 147 (2000) 1671
[44] J. Yang, Y. Takeda, N. Imanishi and O. Yamamoto, Electrochimica Acta, 46 (2001) 2659
[45] Y. M. Kang, S. C. Park, Y. S. Kang, P. S. Lee and J. Y. Lee, Solid State Ionics 156 (2003) 263[46] Y. Liu, K. Horikawa, M.Fujiyosi, N. Imanishi, A. Hirano and Y. Takeda, J. Electrochem. Soc., 151 (2004)
A1450[47] Y. Liu, T. Matsumura, N. Imanishi, T. Ichikawa, A. Hirano, Y. Takeda, Electrochemistry Communications 6(2004) 632[48] Y. Liu, K. Horikawa, M. Fujiyosi, N. Imanishi, A. Hirano and Y. Takeda, Electrochimica Acta 49 (2004) 3487
[49] Y. Liu, K. Horikawa, M. Fujiyoshi, T. Matsumura, N. Imanishi and Y. Takeda, Solid State Ionics, 172 (2004)
69[50] Y. Liu, N. Imanishi, A. Hirano, T. Matsumara and Y. Takeda, Electrochem. Solid-State Lett., in submission[51] Y. Liu, K. Hanai, K. Horikawa, N. Imanishi, A. Hirano and Y. Takeda, Materials Chemistry & Physics, inpress[52] Y.Aihara, M.Kodama, K.Nakahara, H.Okise, K. Murata, J.Power Sources. Vol.65 pp.143-148 (1997).
[53] F. Croce, B. Scrosati, J. Power Sources.Vol.43-45, pp.9-12 (1993).
[54] I.M. Ismail, U. Kadiroglu, N. D. Gray, J. R. Owen, in: Electrochemical Society Meeting Abtracts, San
Antonio, TX, Oct. 6-12 Vol. 96-2,1999, Abstract 63.
[55] D. Fauteux and R. Koksbang, J. Appl. Electrochem. Vol.23, pp.1-6 (1993).
[56] J. O. Besenhard, J. Yang and M. Winter, J. Power Sources 68 (1997) 87[57] A. Belanger, M. Robitaille, US patent 4652506, 1989.
[58] M. Hiratani, K. Miyauchi, T. Kudo, in: Int.Conf Solid State Inoics, Garmish 28-30 (1988) 1431
[59] J. Yang, Y. Takeda, Q. Li, N. Imanishi, O. Yamamoto, J. Power Sources 90 (2000) 64[60] J. Yang, Y. Takeda, Q. Li, N. Imanishi, O. Yamamoto, J. Power Sources 97-98 (2001) 779[61] J. Yang, M. Wachtler, M. Winter, J. O. Besenhard, Electrochem. Solid State Lett. 2 (1999) 161.
[62] Y. Liu, J. Yang, N. Imanishi, A. Hirano, Y. Takeda, and O. Yamamoto, J. Ceramic Society of Japan 112 (2004)s638.
[63] Y. Liu, J. Yang, N. Imanishi, A. Hirano, Y. Takeda and O. Yamamoto, J. Power Sources, in press