Appendix A: Future Directions
A.7.3 Example for the Interaction of Structural, Aerodynamic,
8.4 Lithium Ion Batteries
Since the use of nonaqueous electrolytes allows the realization of high cell voltages of around 4 V in practical cells, lithium ion cells off er much higher volumetric (W h/L) and gravimetric (W h/kg) energy densities compared to other rechargeable sys- tems as seen in Figure 8.2 [6]. As a result, they have become attractive power sources for portable electronic devices such as laptop computers and cell phones, revolutionizing the electron- ics industry. Although the concept of rechargeable lithium bat- Anode
A
Cathode
− X +
Separator
Electrolyte Electrolyte
e− e−
FIGURE 8.1 Operating principles of a battery, which is an electro- chemical cell.
TABLE 8.1 Major Primary Battery Systems and Th eir Characteristics
Battery Anode Cathode Cell Reaction Cell Voltage (V) Capacity (mA h/g)a
Leclanche Zn MnO2 Zn + 2MnO2→ ZnO . Mn2O3 1.6 224
Magnesium Mg MnO2 Mg + 2MnO2 + H2O → Mn2O3 + Mg(OH)2 2.8 271
Alkaline MnO2 Zn MnO2 Zn + 2MnO2→ ZnO + Mn2O3 1.5 224
Mercury Zn HgO Zn + HgO → ZnO + Hg 1.34 190
Zinc–air Zn O2 Zn + 0.5O2→ ZnO 1.65 658
Li–SO2 Li SO2 2Li + 2SO2→ Li2S2O4 3.1 379
Li–MnO2 Li MnO2 Li + MnO2→ LiMnO2 3.1 286
a Based on active electrode materials (cathode and anode) only.
TABLE 8.2 Major Secondary Battery Systems and Th eir Characteristics
Battery Anode Cathode Cell Reaction Cell Voltage (V) Capacity (mA h/g)a
Lead–acid Pb PbO2 Pb + PbO2 + 2H2SO4→ 2PbSO4 + 2H2O 2.1 120
Nickel–cadmium Cd NiOOH Cd + 2NiOOH + 2H2O → 2Ni(OH)2 + Cd(OH)2 1.35 181
Nickel–hydrogen H2 NiOOH H2 + 2NiOOH → 2Ni(OH)2 1.5 289
Nickel-metal hydride MH NiOOH MH + NiOOH → M + Ni(OH)2 1.35 206
Lithium ion Li Li0.5CoO2 0.5Li + Li0.5CoO2→ LiCoO2 3.7 137
a Based on active electrode materials (cathode and anode) only.
teries was initially demonstrated with a sulfi de cathode such as TiS2 and a metallic lithium anode about three decades ago [7–9], the dendritic short-circuiting and the safety concerns associated with the metallic lithium anode remained as an impediment for several years to commercialize the rechargeable lithium battery technology. It is because of the design and development of smart materials as cathodes and anodes [7–17], the present-day lithium ion battery technology has become a commercial reality, follow- ing the fi rst announcement by Sony in 1991.
Figure 8.3 shows the operating principles involved in a lith- ium ion cell. It involves the reversible extraction and insertion of lithium ions from or into two lithium insertion hosts during the charge–discharge process. For example, in commercial lithium ion cells, the lithium ions are extracted from the lay- ered LiCoO2 cathode and inserted into the layered graphite during charge, while the electrons fl ow from the cathode to the anode through the external circuit. During discharge, exactly the opposite happens. Since the process involves a shuttling or rocking of lithium ions between the cathode and anode during the charge–discharge process, the cell is also some times referred to as rocking chair cell. Although the concept looks simple, the anode and cathode materials should satisfy several criteria outlined below in order for them to be successful:
D
• iff erence in the lithium chemical potential between the cathode and anode should be large to maximize the cell voltage. Th e cell voltage is determined by the energies involved in both the electron transfer and the Li+ trans- fer. While the energy involved in electron transfer is related to the work functions of the cathode and anode, the energy involved in Li+ transfer is determined by the crystal structure and the coordination geometry of the site into/from which Li+ ions are inserted or extracted
[18]. If we consider only the electron transfer, then the Mn+ ion in the insertion compound LixMyOz should have a high oxidation state in order for it to be employed as a cathode and a low oxidation state in order for it to be employed as an anode.
I
• nsertion compound LixMyOz should allow an insertion or extraction of a large amount of lithium x to maximize the cell capacity. Th is depends on the number of available lithium ion sites and the accessibility of multiple valences for M in the insertion host.
I
• nsertion compound LixMyOz should support both high electronic and lithium ion conductivities to off er high power density.
L
• ithium insertion and extraction process should be revers- ible with no or minimal structural changes in the host structure over the entire range x of lithium insertion or extraction in order to provide good cycle life.
I
• nsertion compound LixMyOz should be chemically stable without undergoing any reaction with the elec trolyte over the entire range x of lithium insertion or extraction.
I
• nsertion compound LixMyOz should be low cost and envi- ronmentally benign to be employed in practical cells and commercially feasible.
Considering some of these criteria, transition metal oxides crystallizing in layered LiMO2, spinel LiM2O4, and olivine LiMPO4 structures (M = transition metal) have become appeal- ing as cathodes, while carbon-based materials such as graphite have become attractive anodes for lithium ion batteries. Th is is because while these oxides off er high voltages of 3.5–4.3 V versus metallic lithium, graphite off ers a low voltage of <0.5 V versus metallic lithium, resulting in a cell voltage of 3–4 V on coupling these oxide cathodes and graphite anode. Th e high voltage off ered by these oxide cathodes is due to a large 400
300 200
100 0
0 50 100 150 200 250
Ni/MH Ni/Cd
Lead-acid
Lighter
Smaller
Lithium-ion
Gravimetric energy density (Wh/kg)
Volumetric energy density (W h/L)
FIGURE 8.2 Comparison of the energy densities (gravimetric and volumetric) of various rechargeable battery systems.
Anode Cathode
Charge Discharge
Discharge Charge
Li+
Li+
Li+ Li+
LixC6 Li1−xCoO2 Electrolyte
e− e−
FIGURE 8.3 Operating principles of a lithium ion cell, illustrating the charge–discharge process.
Madelung energy in oxides and the consequent ability to stabilize high oxidation states such as Co3+/4+. However, these oxide cathodes diff er signifi cantly in some of their performance factors such as energy and power densities as well as cost and environmental impact. Also, although the lithium ion battery technology has revolutionized the portable electronics market, they are yet to enter the HEV and PHEV arena. Moreover, the rapid developments in the miniaturization and features of por- table electronics devices demand an increase in the energy density of the lithium ion batteries while lowering the cost of electrode materials. Th e sections below present the recent developments and current status of the smart cathode and anode materials mentioned above.
8.4.1 Layered Oxide Cathodes
Oxides with a general formula LiMO2 (M = Co and Ni) adopt a layered structure in which the Li+ and M3+ ions occupy the alternate (111) planes of a rock salt structure as shown in Figure 8.4 to give a layer sequence of –O–Li–O–M–O– along the c-axis. With an oxygen stacking sequence of ABCABC along the c-axis, this structure is designated as O3 layer structure as the Li+ ions occupy the octahedral sites and there are three
MO2 sheets per unit cell. While the Li+ ions present between the strongly bonded MO2 layers provide facile, reversible Li+- ion diff usion, the edge-shared MO6 octahedral arrangement with a direct M–M interaction provides good electronic con- ductivity. As a result, for example, layered LiCoO2 off ers accept- able power capability (charge–discharge rate) for portable applications such as laptop computers and cell phones. On the other hand, a large work function associated with the highly oxidized Co3+/4+ couple provi des a high cell voltage of around 4 V and the discharge voltage does not change signifi cantly with the degree of lithium extraction/insertion x in Li1−xCoO2 (Figure 8.5). Th e high energy density with acceptable power capability and cycle life has made LiCoO2 an attractive cathode, and com- mercial lithium ion cells presently use predominantly layered LiCoO2 as the cathode.
However, only 50% of the theoretical capacity of LiCoO2, which corresponds to a reversible extraction or insertion of 0.5 lithium per Co and a practical capacity of 140 mA h/g, can be utilized in practical cells as seen in Figure 8.5. Although this limitation was originally attributed to structural distortions around x = 0.5 in Li1−xCoO2, resulting from an ordering of Li+ ions [19], recent characterization of chemically delithiated sam- ples suggest that chemical instability arising from a signifi cant overlap of Co3+/4+:3d energy band with the top of the O2−:2p band may play a critical role in controlling the reversible capacity values [20]. Th is conclusion is consistent with the introduction of holes (removal of electrons) into the O2−:2p band rather than into the Co3+/4+:3d band during electrochemical charge as revealed by x-ray absorption spectroscopic [21] and electron energy loss spectroscopic [22] studies. Additionally, Co is rela- tively expensive and toxic. Moreover, the chemical instability of LiCoO2 leads to safety concerns under the conditions of over- charge, which becomes a serious issue particularly in the case of large batteries such as those necessary for HEV and PHEV applica tions. Also, the power capability of LiCoO2 is not attrac- tive for vehicle applications.
Th e above drawbacks associated with LiCoO2 have created interest in the development of alternative cathode materials.
In this regard, the nickel-rich LiNi0.85Co0.15O2 that has the lay- ered structure shown in Figure 8.4 became appealing a few years ago [23] as it off ers a higher capacity of 180 mA h/g as seen in Figure 8.5 and Ni is less expensive than Co. However, more in-depth investigation has shown that LiNi0.85Co0.15O2
suff ers from impedance growth during cycling at elevated temperatures, which may partly be related to the migration of Ni2+/3+ ions from the transition metal plane to the lithium plane at elevated temperatures [24]. Recently, layered compo- sitions such as LiNi1/3Mn1/3Co1/3O2 [25–27] and LiNi1/2Mn1/2O2
[28,29] have become attractive as they off er capacities close to 200 mA h/g at a lower cost compared to LiCoO2. However, their power capability is lower than that of LiCoO2 due to some cation disordering between the transition metal and lithium planes, and the power capability may not be adequate for vehicle applications.
FIGURE 8.4 Crystal structure of layered LiMO2 (M = Co or Ni).
More recently, complex layered oxide solutions between Li[M]
O2 (M = Ni, Mn, and Co) and Li[Li1/3Mn2/3]O2 that have the same O3 layered structure shown in Figure 8.4 have become appealing as they exhibit high capacities of up to 280 mA h/g [29–32], which is two times higher than that of LiCoO2. However, these oxides encounter an irreversible loss of oxygen gas from the lattice during fi rst charge and a huge irreversible capacity loss of 40–100 mA h/g in the fi rst cycle. Also, they need a high cutoff charge voltage of 4.8 V and slow charge–discharge rate (low power capability). Th e irreversible capacity loss could be sup- pressed to some extent by a surface modifi cation of the oxides with inert oxides like Al2O3 [32]. Th e amount of oxygen loss from the lattice and the reversible capacity values have been found to increase with increasing lithium content in the transition metal layer in such Li[M1−hLih]O2 solid solutions [33,34]. Also, the irreversible oxygen loss from the lattice has been found to be sensitive to the substitution of other elements like Al3+ for Li+ and F− for O2− [35]. Optimization of such complex oxide solid solutions as well as development of robust electrolytes that are stable up to 4.8 V could make these high-capacity cathodes attractive for portable electronic devices although their low power capability due to cation disorder may not be adequate for vehicle applications.
8.4.2 Spinel Oxide Cathodes
Although layered oxides exhibit high capacity, some of their per- formance parameters such as power capability are oft en limited by cation disorder between the transition metal and lithium planes as well as structural and chemical instabilities occurring during cycling. Th e compositions with low cation disorder tend to transform from the initial O3 structure to, for example, O1 structure at deep charge [20,27,36]. Th e limited power capability as well as the safety and cost concerns particularly in composi- tions with high Co contents makes them less appealing for HEV and PHEV applications. In this regard, LiMn2O4 crystallizing in the spinel structure has appealing for vehicle applications as Mn is inexpensive and environmentally benign. Also, the three- dimensional spinel structure (Figure 8.6) with LiO6 tetrahedra and edge-shared MnO6 octahedra provides good structural
stability during the cycling while the lying of the Mn3+/4+:3d band well above the O2−:2p band off ers good chemical stability unlike, for example, layered LiCoO2.
However, LiMn2O4 that has been investigated extensively over the years has been plagued by severe capacity fade particularly at elevated temperatures. Several mechanisms such as Mn dissolu- tion from the lattice [37], structural (Jahn–Teller) distortion during over discharge [38], formation of two cubic phases during the charge–discharge process [39], and loss of crystallinity have been proposed to account for the capacity fade [40]. Several approaches such as cationic and anionic substitutions as well as surface modifi cation have been pursued to improve the capacity retention, but they could not completely overcome the problem.
More recently, from a systematic investigation of a number of singly and doubly substituted spinel compositions, initial Mn valence and the lattice parameter diff erence Δa between the two cubic phases formed during the charge–discharge process have been found to play a critical role in controlling the capacity retention [41–43].
2.5 3 3.5 4 4.5
0 50 100 150 200 Capacity (mA h/g)
Cell voltage (V)
LiCoO2 LiMn2O4
LiNi0.85Co0.15O2 LiFePO4
FIGURE 8.5 Comparison of the discharge curves of layered LiCoO2, layered LiNi0.85Co0.15O2, spinel LiMn2O4, and olivine LiFePO4.
FIGURE 8.6 Crystal structure of spinel LiMn2O4, illustrating the three-dimensional framework with LiO4 tetrahedra and MnO6 octahedra.
However, cationic substitutions oft en result in an increase in the oxidation state of Mn and a consequent decrease in capacity values to unattractive levels of <100 mA h/g. Th is problem has been overcome by a partial substitution of F− for O2− in stabilized cat- ion-substituted spinel oxide compositions. For example, stabilized spinel oxyfl uoride compositions such as LiMn1.8Li0.1Ni0.1O3.8F0.2
exhibit excellent capacity retention at elevated temperatures with a capacity of >100 mA h/g, high power capability, superior storage properties, and low irreversible capacity loss in the fi rst cycle [44]. Th e excellent electrochemical performance is found to be due to a smaller lattice mismatch (low Δa) between the two cubic phases formed during the charge–discharge process, sup- pressed manganese dissolution, and an initial Mn valence of
>3.6+. From a systematic investigation of several spinel composi- tions, the capacity fade is found to bear a clear relationship to initial Mn valence, Δa, and Mn dissolution [45]. Th e Mn dissolu- tion could also be suppressed further by mixing the stabilized spinel oxyfl uoride compositions with a layered oxide like LiCoO2
and charging to high enough voltages (4.7 V) in the fi rst cycle [46]. Th is is because the overcharged layered Li1−xCoO2 traps the trace amounts of protons present in the electrolyte and thereby suppresses the disproportionation of Mn3+ and the Mn dissolu- tion from the spinel lattice. Th is composite strategy involving a spinel + layered oxide mixture could become a viable strategy to adopt the spinel cathodes for automotive applications.
In addition to the 4 V spinel cathodes described above, 5 V spinel cathodes based on LiMn1.5Ni0.5O4 [47] are appealing par- ticularly for high-power applications. As in the case of 4 V spinel cathodes, the electrochemical performance of 5 V spinel also bears a clear relationship to the lattice mismatch among the three cubic phases formed during the charge–discharge process [48]. Stabilized spinel compositions such as LiMn1.42Ni0.42Co0.16O4
that involve a much smaller lattice mismatch among the cubic phases exhibit better cyclability. However, development of robust electrolyte compositions that off er long-term stability at 5 V is critical to realize the full potential of these 5 V cathodes.
8.4.3 Olivine Oxide Cathodes
Another candidate that exhibits excellent structural and chemi- cal stabilities along with superior safety characteristics is LiFePO4
crystallizing in the olivine structure as shown in Figure 8.7 [17].
Th e excellent safety is due to the chemically more stable Fe2+/3+
redox couple and the covalently bonded PO4 groups. Despite the Fe2+/3+ couple, LiFePO4 exhibits a high discharge voltage of 3.6 V as seen in Figure 8.5 due to the inductive eff ect caused by the countercation P5+ and the lowering of the Fe2+/3+ redox energy [49,50]. However, the major drawback with olivine LiFePO4 is the low electronic and lithium ion conductivities and the consequent decrease in the power capability. Th e low electronic conductivity is due to the highly localized Fe2+/3+
redox couple, the presence of FeO6 octahedra, and PO4 tetrahe- dra, and the coexistence of LiFePO4 and FePO4 as line phases during the charge–discharge process without any detectable mixed valent Fe2+/3+.
However, this problem has been overcome signifi cantly by making the LiFePO4 powder at the nanoscale to reduce the lith- ium diff usion length as well as by an intimate mixing or coating with conductive carbon to improve the electrical conductivity [51–53]. Although conventional LiFePO4 powder with a larger particle size undergoes a two-phase reaction involving LiFePO4
and FePO4 without any solid solubility range during the charge–
discharge process, reducing the particle size to the nanoscale is believed to shift the miscibility gap to lower temperatures and provide some solid solubility range and mixed valency. With nanoscale powder and conductive carbon along with associated engineering, signifi cant progress has been made recently by A123 company to employ LiFePO4 for high-power applications.
LiFePO4 is currently employed in lithium ion batteries used in power tools, and it is intensively being pursued for automotive applications.
8.4.4 Carbon Anodes
With a lightweight and low electrochemical potential lying close to that of metallic lithium, graphite has become an attractive anode, and commercial lithium ion cells currently use mostly the carbon anodes [54,55]. It off ers a theoretical capacity of 372 mA h/g, which corresponds to an insertion of one lithium per six carbon atoms (x = 1 in LixC6). One of the drawbacks with the carbon anodes is the occurrence of signifi cant amount of irreversible capacity loss during the fi rst discharge–charge cycle due to unwanted, irreversible side reactions with the electrolyte.
Although natural graphite cannot be charged with electrolytes consisting of propylene carbonate (PC) as it leads to an evolution of gas at around 1 V, this problem could be overcome with elec- trolytes consisting of other solvents such as ethylene carbonate FIGURE 8.7 Crystal structure of olivine LiFePO4, consisting of FeO6 octahedra and PO4 tetrahedra.
(EC) and diethyl carbonate (DEC). Additionally, hard carbons (glassy carbon) obtained by a thermal decomposition of pheno- lic and epoxy resins and products from petroleum pitch show higher capacities than graphite as they could accommodate extra lithium at the edges and surfaces of the graphene sheets and into nanometer size cavities. Also, hard carbons can be used with PC-based electrolytes unlike graphite. However, the hard car- bons show a sloping discharge profi le between 0 and 1 V unlike graphite, which shows a nearly f lat discharge profile between 0 and 0.3 V.