There and back again - The journey of LiNiO2 as

There and back again - The journey of LiNiO2 as cathode active material

Matteo Bianchinia*, Maria Roca-Ayatsa, Pascal Hartmanna,b, Torsten Brezesinskia, Jürgen Janeka,c* a

Battery and Electrochemistry Laboratory, Institute of Nanotechnology, Hermann-von-

Helmholtz-Platz 1, 76344 Eggenstein-Leopoldshafen, Germany b

BASF SE, 67056 Ludwigshafen, Germany

c

Institute of Physical Chemistry & Center for Materials Science (ZfM/LaMa), Justus-Liebig-

University Giessen, Heinrich-Buff-Ring 17, 35392 Giessen, Germany * Corresponding authors: [email protected], [email protected]

Keywords: Lithium nickel oxide, LNO, lithium-ion batteries, cathode material

1

2

Abstract LiNiO2 (LNO) has been introduced as cathode active material (CAM) for Li-ion batteries in 1990. After years of intensive research, it emerged that several instability issues plague the material, so that it was abandoned in favor of isostructural metal-substituted compounds called NCA (for lithium nickel cobalt aluminum oxide) and NCM (lithium nickel cobalt manganese oxide). These sacrifice a certain amount of specific energy in exchange for stability, durability and safety. With few exceptions, NCA and NCM are nowadays the industrial standard when it comes to automotive applications; however, the continuous push towards electric cars with longer driving range is synonym, for these compounds, with increasing the nickel content (which is already beyond 80%), eventually leading back again to LiNiO2. For this reason we provide here a comprehensive review of the material, almost 30 years after its introduction as CAM. We aim at highlighting its physicochemical peculiarities, which make LNO complex in every aspect. We specifically stress the effect of the Li off-stoichiometry (Li1-zNi1+zO2) on every property of LNO, especially the electrochemical ones. We then focus on the key instability issues that plague the compound and on the strategies implemented so far to overcome them. Finally, in the course of the review we point to open questions that still remain to be addressed by the scientific community, and to which research directions seem more promising to lead LNO to its full exploitation.

3

Table of Contents 1.

Introduction.................................................................................................................................. 6

2.

Structure of LiNiO2 ..................................................................................................................... 8

3.

Phase diagram: the Li2O – NiO – O2 region ....................................................................... 10

4.

5.

6.

7.

8.

9.

3.1

Line I: Over-lithiated Li1-zNi1+zO2 (-1/3 ≤ z ≤ 0) ............................................................... 11

3.2

Line II: Li1-zNi1+zO2 (0 ≤ z ≤ 1) and the Li off-stoichiometry ........................................... 12

3.3

Line III and IV: LixNiO2 (0 ≤ x ≤ 2) .................................................................................... 15

Physical properties .................................................................................................................. 16 4.1

Electronic properties........................................................................................................... 16

4.2

Magnetic properties ............................................................................................................ 17

4.3

Thermodynamics ................................................................................................................ 18

Synthesis methods .................................................................................................................. 18 5.1

Solid-state reaction ............................................................................................................. 19

5.2

Co-precipitation: industrial process .................................................................................. 20

5.3

Sol-gel................................................................................................................................... 21

5.4

Combustion.......................................................................................................................... 21

5.5

Other methods..................................................................................................................... 22

5.6

Washing LNO ...................................................................................................................... 22

Electrochemistry....................................................................................................................... 23 6.1

Phase transitions upon Li (de)intercalation .................................................................... 26

6.2

Nature of the monoclinic domain ...................................................................................... 27

6.3

Nature of the highly oxidized phases ............................................................................... 30

6.4

Critical role of Li off-stoichiometry on the electrochemical performance.................... 32

Challenges.................................................................................................................................. 33 7.1

Mechanical instability ......................................................................................................... 33

7.2

Thermal instability ............................................................................................................... 34

7.3

(Electro)chemical instability ............................................................................................... 36

Elemental substitution / Doping ........................................................................................... 40 8.1

Cobalt ................................................................................................................................... 41

8.2

Iron ........................................................................................................................................ 42

8.3

Manganese .......................................................................................................................... 43

8.4

Titanium................................................................................................................................ 44

8.5

Aluminum ............................................................................................................................. 45

8.6

Magnesium .......................................................................................................................... 46

8.7

Other dopants...................................................................................................................... 46

Future prospects ...................................................................................................................... 49 4

9.1

Coating ................................................................................................................................. 50

9.2

Core-shell and concentration gradient strategies .......................................................... 51

9.3

Related compounds: disordered rock salts .................................................................... 52

10.

Conclusion ............................................................................................................................. 53

Bibliography....................................................................................................................................... 53

5

1. Introduction Lithium-ion batteries (LIBs) are well-established as the leading electrochemical energy storage technology due to their reliability, high energy density and longevity. Since their commercialization in 1991 by Sony Corp., major advances in cell chemistry, design and manufacturing enabled to increase the energy density by a factor of about four to date [1]. While the original graphite anode is still in use today

[2]

the cathode active material is under

continuous development and scrutiny. LiCoO2 (LCO) was first used as introduced by Mizushima et al.

[3]

, and it is still the material of choice in batteries for portable applications.

The theoretical capacity of LiCoO2 is 274 mAh/g for full delithiation, but not all of it can be practically used due to structural instability of LixCoO2 for x < 0.5, although stabilizing coatings allow exchanging up to 0.6 Li/Co [3-4]. Due also to thermal instability issues [5], LCO cannot be used in large cells. Moreover, especially for electric vehicles, LCO is not applicable due to the limited availability and price of cobalt [6]. An obvious alternative to LiCoO2 was already identified in the ´90s in the isostructural compound LiNiO2 (LNO)

[7]

. This layered material is attractive because it has the same high

theoretical energy density of LCO, but with lower cost and higher abundance of nickel with respect to cobalt. However, despite intensive research for more than two decades, LiNiO2 did not become a commercial reality due to the major drawbacks it suffers. A precise control of the material´s synthesis is not straightforward because it is strongly prone to Li off-stoichiometry (z > 0 in Li1-zNi1+zO2) [7a, 8]. Besides, LNO is plagued by stability issues at high potential similar to those of LCO, which are discussed in detail in section 7. For clarity, we divided such instability issues as (electro)chemical, related to the instability of LNO and its surface during handling and when delithiated in a battery cell; mechanical, related to the large volume changes and to the multitude of phase transitions occurring upon cycling; and thermal, related to the decomposition of LNO when heated in the charged state. Research efforts on LNO resulted in the development of several substitution strategies, where Ni is partly replaced with other cations to find a compromise between high energy density, 6

stability (long cycle life and safety) and cost. Based on this approach, new oxides were synthesized, where cobalt and either manganese or aluminum are employed as substituents (or dopants), thereby reducing the total fraction of nickel. The resulting solid solutions are called NCM (also named NMC) and NCA, respectively. Co ions are used to control the 2D character of the layered structure [9], while Al or Mn do not participate in the redox process but stabilize the material structurally and thermally, especially in the charged state. Optimal compositions include LiNi0.8Co0.15Al0.05 and LiNi1/3Co1/3Mn1/3O2 (NCM111) or LiNi0.6Co0.2Mn0.2O2 (NCM622). These materials show sufficiently high performance (q ≈ 200 mAh/g for NCA, q ≈ 180 mAh/g for NCM622 [2.5-4.3 V]

[10]

) to be viable cathodes in the current generation of

electric vehicles [1]. However, the automotive market demands further capacity increase to allow longer driving ranges approaching the 300 miles threshold, and reducing the price of the battery pack

[10-11]

.

For this reason, most efforts go towards the development of Ni-rich NCM and NCA materials, where the Ni content is beyond 80% and the gravimetric capacity is >200 mAh/g. Clearly, a trade-off exists since enriching these materials in Ni is equivalent to approaching again LiNiO2, thus facing stability problems. A huge academic and industrial interest is generated by these two solid solutions and by the capacity/stability (high-Ni/low-Ni) tradeoff, resulting in hundreds of papers and several recent reviews published on the topic [12]. The current battery technology in the Tesla Model 3 is already based on NCA with >80% Ni, clearly indicating that cathode materials will be pushed towards the LNO chemistry within the next years. Therefore, we provide here a comprehensive review on LiNiO2. We discuss the material´s structure, thermodynamics, position in the Li-Ni-O phase diagram, physicochemical and electrochemical properties and stability issues. We then analyze the effect of single dopants on LNO and discuss which directions are promising for future investigation. The critical role of Li off-stoichiometry for the material´s properties, especially electrochemical ones, is highlighted. The review aims at providing a fresh perspective on an apparently “old” material, trying to point out the remaining challenges to be addressed, to clarify discrepancies and the 7

sources of LNO´s drawbacks, in the hope that they could be overcome and LNO be used in the next generation of electric vehicles, with as little modification as possible to its pristine structure.

Figure 1: Crystallographic structure of oxides with ccp oxygen sublattice. From left to right, rock salt structure of composition Li0.3Ni0.7O, layered LiNiO2, and spinel LiNi2O4. Green spheres/octahedra represent Li, grey ones Ni and red ones oxygen. The red double arrows indicate the narrow Ni layers and wider Li interlayers along the [001] direction of the layered structure (c axis), as opposed to the constant layer thickness in the rock salt structure ([111] direction). Note that in the spinel structure Li occupies tetrahedral sites.

2.

Structure of LiNiO2

LiNiO2 is a layered oxide with a cubic close packed (ccp) array of oxygen atoms, first reported in 1954

[13]

. It belongs to the solid solution obtained when concurrently lithiating and oxidizing

NiO. If Li2O is used, the reaction can be written as:

(1 − 𝑧´) NiO +

𝑧´ 2

Li2 O +

𝑧´=(1−𝑧)/2 1 2

𝑧´ 4

O2 → Li𝑧´ Ni1−𝑧´ O →

Li1−𝑧 Ni1+𝑧 O2

(1)

The chemical composition of LNO can be expressed by the general formula Li1-zNi1+zO2 (or as Liz´Ni1-z´O, with z´= (1-z)/2))

[14]

. For low lithium content (z > 0.38) the solid solution maintains

the rock salt structure typical of NiO, with Li and Ni statistically occupying the 4a site. When a 8

higher Li/Ni ratio is reached and a sufficient amount of Ni2+ is oxidized to Ni3+, the structure becomes layered with alternating Li and Ni planes along the [111]cubic direction ([001]hexagonal). Figure 1 shows the crystal structure of the solid solution, before and after the Li/Ni layering is introduced. The cubic-hexagonal transition is driven by steric effects

[15]

, namely by the

increasing size difference between Li+ and Ni during its oxidation from the divalent to the trivalent state (r(Li+) = 0.76 Å, r(Ni2+) = 0.69 Å, r(Ni3+LS) = 0.56 Å according to Shannon

[16]

).

Layering allows to better accommodate Ni3+ in small slabs (thickness S = (2/3 - 2zox) c ≈ 2.11 Å) and Li+ in larger interslabs (thickness I = (c/3 - S) ≈ 2.62 Å)

[17]

. The resulting structure is

rhombohedral, crystallizing in the space group 𝑅3̅𝑚, isostructural with α-NaFeO2. Oxygen atoms occupy the 6c sites (0 0 zox=0.258), and lithium and nickel atoms occupy the octahedral sites (3b (0 0 0.5) and 3a (0 0 0), respectively). It can be denoted as [Li]3b[Ni]3a[O2]6c Diffusion of Li atoms in the interlayer proceeds via intermediate tetrahedral sites

[19]

[7a, 18]

.

(also 6c

with ztetra ≈ 0.13, Figure 1). Note that a different choice of space group origin is possible, in which Li/Ni are exchanged between 3a and 3b sites, and zox=0.24. According to the first description of this material, the lattice parameters of the hexagonal structure are a = 2.878 Å and c = 14.19 Å

[13]

and the c/a ratio is a measure of the hexagonal distortion: in the cubic

structure, c/a = 2√6 ≈ 4.9 and this value increases with the material´s layering. A well-layered compound like LiCoO2 has a high value of c/a = 4.99. For LNO instead c/a = 4.93

[7e]

, which

indicates a weaker ordering (indeed, r(Co3+LS) = 0.54 Å, smaller than Ni3+). As discussed in section 3.2, regardless of the synthesis method, LNO tends to be understoichiometric in lithium and contain excess Ni2+ located in the lithium layers (“offstoichiometric” LNO), strongly affecting its physicochemical properties. Hence, in the formula Li1-zNi1+zO2, z represents the fraction of Ni located in the Li layer according to [Li1zNiz]3b[Ni]3a[O2]6c

[8, 20]

. To maintain charge compensation, the average oxidation state of Ni is

3-2z, implying that z Ni2+ are created also in the Ni slab, yielding a formula [Li12+ 2+ 3+ zNi z]3b[Ni zNi 1-z]3a[O2]6c.

The presence of Ni2+ ions shrinks the lithium interslab (r(Ni2+)
r(Ni3+))

[20a, 21]

. Consequently, the Li-O distance

decreases and the Ni-O distance increases, and so do both a and c lattice parameters

[7a, 14b,

9

17-18, 21-22]

. Other possible Li/Ni arrangements in LNO, such as [Li1-zNiz]3b[LizNi1-z]3a[O2]6c (Li/Ni

antisite defects) or [Li1-zNiz]3b[LiδNi1-δ]3a[O2]6c are particularly difficult to probe with standard XRD techniques due to the weak scattering power of Li. Some authors verified these hypotheses using neutron powder diffraction, and confuted them for near-stoichiometric LNO. Yet, compounds far from ideal stoichiometry can display Li in the Ni layer [21a, 22a, 23].

3. Phase diagram: the Li2O – NiO – O2 region Several authors reported the Li-Ni-O phase diagram or subsets of it as obtained from first principles calculations [24]. These yield important information and also include tie lines between the stable phases, but depend on the calculation methodology. For example, stoichiometric LNO is not found to be a stable phase unless peculiar Jahn-Teller (JT) orderings of the ground state are considered (section 4.1). Early on, phase diagrams were also constructed from experimental data [25]. Here, Figure 2 shows a schematic (pseudo) phase diagram of the Li2ONiO-O region. These corners are chosen because they represent typical synthesis precursors (considering that Li and Ni hydroxides, carbonates, etc. eventually decompose to oxides upon annealing, and that typically an oxygen-rich atmosphere is used). The main identified phases are Li2O and Li2O2, rock salt-type NiO, layered LiNiO2 and Li2NiO3, spinel LiNi2O4, and Li2NiO2. The existence of other oxides such as Ni2O3 and Ni3O4 has been postulated, but they were never reproducibly prepared. NiO2, on the other hand, exists as the fully delithiated LiNiO2 and can be prepared chemically or electrochemically

[26]

. Its structure

will be described in detail in section 6.1. It should also be mentioned that recently new oxides of compositions Ni1.25O2, Ni1.5O2, Ni1.75O2 have been predicted to be stable along the NiO-NiO2 line in computed finite-temperature phase diagrams

[24b]

, and their experimental verification

should be sought.

10

Figure 2: Pseudo phase diagram of the Li2O – NiO – O system. The corners represent typical LNO precursors, black points indicate important compounds. The color scale denotes the Ni oxidation state. Arrows represent irreversible reactions, while double arrows represent reversible ones. Lines I to IV are discussed in the text. Along lines I and II, diamonds identify compounds with rock salt structure and triangles layered ones. The grey dot indicates a typical off-stoichiometric Li1-zNi1+zO2 with z = 0.03.

3.1 Line I: Over-lithiated Li1-zNi1+zO2 (-1/3 ≤ z ≤ 0) As previously discussed, LNO is a member of the solid solution Li1-zNi1+zO2 with z = 0. At typical synthesis O2 chemical potentials μ(O2) (patm, ϑ ≈ 700 °C), z = 0 can also be considered as the end member of the solid solution. However, when the oxygen partial pressure is increased, it is possible to stabilize samples in the region -1/3 < z < 0, where the new end member is Li2NiO3 (Ni4+). To synthesize it, either p(O2) up to 15 MPa GPa (at ϑ ≈ 700 °C) were employed

[25b, 27]

or high-pressure synthesis up to 3-4

[28]

. In this compound, still displaying a layered structure,

z lithium ions are in the nickel layer. For z = -1/3, the formula is Li[Li1/3Ni2/3]3aO2, with a Li/Ni ordering in the metal layer, thereby lowering the compound´s symmetry to monoclinic C2/m. 11

As z → -1/3, the unit cell size is reduced (Figure 3) and so are the Li-O and Ni-O bond lengths, as expected for Ni oxidation. The whole series of solid solutions -1/3 < z < 0 has been prepared. Tabuchi et al. [29] recently showed that compounds in the range -0.11 < z < 0 can be stabilized by carefully controlling the annealing temperature. For example, at 625 °C, Li1.09Ni0.91O2 with a small unit cell volume of 100.132(5) Å3 could be prepared. It has hexagonal symmetry with Li occupying the Ni layers, while for z < -0.11, the monoclinic distortion appears.

3.2 Line II: Li1-zNi1+zO2 (0 ≤ z ≤ 1) and the Li off-stoichiometry LiNiO2 is usually obtained via the concurrent lithiation and oxidation of NiO using several methods. This corresponds to moving along the “synthesis” line in the phase diagram of Figure 2. At composition Li0.62Ni1.38O2 (z = 0.38) or Li0.31Ni0.69O (z´ ≈ 0.31), a structural phase transition takes place from rock salt to layered structure (Figure 1) [14, 30]. The insertion of even very small amounts of Li into NiO is sufficient to strongly modify its electrical and magnetic properties, and also its appearance from a green to a black powder. A detailed discussion on the rock salt-type phases can be found elsewhere [31]. Figure 3 shows the decrease of unit cell volume with increasing z, net result of the lithiation and of the size reduction of Ni upon oxidation. A linear behavior (Vegard´s law) exists for 0 ≤ z ≤ 1 and -1/3 ≤ z ≤ 0, with a slope change occurring at z = 0 (LNO) due to the extra Li occupying the Ni site.

12

Figure 3: Evolution of the primitive unit cell volume as a function of z in Li1-zNi1+zO2. Data points are from [14a, 28b, 29a, 30b, 32]. Two linear fits are carried out (for z ≤ 0 and z ≥ 0) and the LNO volume extrapolated is reported. Inset: Selected unit cell volumes (hexagonal setting, Z = 3) and corresponding z value obtained from Rietveld refinement from [17-18, 28a, 33], representing “good” near-stoichiometric LNO samples. Although a clear linear relationship exists for 0 < z ≤ 1 and -1/3 ≤ z < 0, for z → 0, the values reported lack a clear correlation. Ideally, as z → 0, LNO is stabilized. However, the fact that LNO belongs to the Li1-zNi1+zO2 solid solution is a major drawback: the synthesis of real LNO samples results in compounds with z = 0.01-0.02, i.e. at least 1-2% of residual Ni2+ in the Li layer, even under very well controlled conditions, as confirmed by many different groups

[7a, 7d, 8, 17-18, 21b, 34]

. The inset in Figure 3

shows how results on “good” (near-stoichiometric) LNO differ throughout the literature. In this respect, it is likely that a “true” (i.e. stoichiometric) LiNiO2 has never been synthesized. The narrow stability range of LNO

[35]

is due in part to the strong oxidation conditions of the

environment required to stabilize Ni3+, which may result in a mixture of Li1-zNi1+zO2 and residual Li precursor in equilibrium with the environment. On the other hand, controlling the actual stoichiometry is challenging because of Li (or rather Li2O) loss at high temperature, and because the material is not stable at high temperature or for too long annealing times, so that stoichiometric LNO may always face some decomposition to Li1-zNi1+zO2

[32, 36]

. The

decomposition reaction can be written as: LiNiO2 →

1 1+𝑧

Li1−𝑧 Ni1+𝑧 O2 +

𝑧 1+𝑧

Li2 O +

𝑧 2(1+𝑧)

O2

(2)

13

where the product Li1-zNi1+zO2 can be either a layered off-stoichiometric phase (z < 0.38) or a rock salt-type phase (z > 0.38), depending on the extent of decomposition. Although prevented at thermodynamic equilibrium, off-stoichiometric layered and rock salt phases could simultaneously be kinetically stabilized in the products. Quantitatively characterizing the off-stoichiometry in LNO is not trivial. In the early ´90s, Ohzuku et al. evaluated the intensity ratio I003/I104 between the (003) and (104) reflections [7d]. This analysis is based on the fact that the (104) reflection appears in both LiNiO2 with 𝑅3̅𝑚 space group and in the rock salt-type phase with 𝐹𝑚3̅𝑚 space group, while a strong (003) reflection is only present in the former. Thus, the higher the I003/I104 ratio, the more “welllayered” the sample

[7d, 37]

. The I003/I104 ratio was even related to the electrochemical

performance, finding an optimum at ≈1.35. However, this method has several drawbacks, as the peak intensities at significantly different scattering angles can be modified by various experimental setup parameters (peak broadening, peak asymmetry, preferred orientation etc.). Thus, in general this method is not reliable and should be avoided. Another indication of the value of z is the splitting of the (108)-(110) and (006)-(102) reflections because their positions coincide in the rock salt structure, while they drift apart when c/a increases. The unit cell volume is also correlated to the off-stoichiometry (it decreases with z), although for z → 0, this relationship is not trivial (inset in Figure 3). Dahn et al. established the so-called factor R = (I006 + I102) / I101 [7a] based on the fact that the intensity of the (101) reflection is rapidly attenuated as Ni occupies the Li layer. A low R value (0.41) is indicative of stoichiometric samples [7a, 37-38]. R and I003/I104 parameters have been extensively used [39] and they can provide a rule-of-thumb stoichiometry. Nowadays, Rietveld refinement is the method of choice, from XRD or neutron diffraction data [18, 21b, 40]. Rougier et al.[8, 17, 20a, 41] suggested a proper refinement procedure for XRD data by exploiting the atomic isotropic displacement parameter of Li (BLi). If the refinement is performed assuming the ideal LiNiO2 structure, BLi becomes negative, a non-physical effect due to the presence of extra electronic density from Ni in the Li site, unaccounted for by the 14

model. For a proper refinement, BLi can first be set to 1.2 Å2 (typical value for an ideal layered structure at room temperature

[8]

). Then, the Ni atoms present in the Li interslab can be

included. In the last step of Rietveld analysis BLi can also be refined, if data quality allows, and a positive value should be obtained. With this method, z can be evaluated with an accuracy of ±0.01

[20a]

. Unfortunately, X-rays do not allow to properly test other hypotheses, such as the

presence of Li in the Ni layers, because of lithium´s weak scattering. For this purpose, neutrons should be used, which have a much higher contrast due to the negative scattering length of Li (and are also more sensitive to displacement parameters). Indeed, this has been done by some authors, largely confirming the off-stoichiometry of LNO [18, 21a, 23]. The off-stoichiometry z influences almost every property of LNO, as discussed in the following: nature of LNO´s ground state structure (section 4.1), magnetic behavior (section 4.2), electrochemistry (section 6), and finally the overall stability (thermal, mechanical and electrochemical; sections 7.1-7.3).

3.3 Line III and IV: LixNiO2 (0 ≤ x ≤ 2) 0 ≤ x ≤ 1 is the domain relevant for electrochemical application. Upon delithiation from LNO, a number of phases characterized by Li-vacancy orderings are stabilized, which will be discussed in section 6.1. Here, we shall mention the Li0.5NiO2 phase obtained halfway during delithiation of LNO. This is a layered structure, but it has the typical spinel stoichiometry LiNi2O4. Indeed, Thomas et al. [42]

showed that by gentle heating above 150 °C the layered phase converts irreversibly into

the spinel one, which is stable up to 300 °C. The layered and spinel structures are closely related as both have ccp oxygen sublattices, and the layered one becomes spinel by cooperative migration of ¼ of the Ni to the neighboring Li planes, with concurrent displacement of all the lithium ions to tetrahedral sites (Figure 1). Interestingly, the spinel structure is calculated to be the ground state at Li0.5NiO2 composition [24b]. Moreover, the spinel phase was shown to reversibly take up Li according to LixNi2O4 (1 < x < 2) around 2 V vs. Li+/Li

[18]

.

15

Electrochemical lithium extraction has not been reported, but spinel phases LixNi2O4 with x < 0.5 have been obtained by heating the corresponding layered compounds [18]. Finally, Li2NiO2 can also be synthesized [43]. Although initially indexed as orthorhombic (Immm), the compound was later shown to be tetragonal (𝑃3̅𝑚) [7a, 44], with cell parameters a = b = 3.095 Å, and c = 5.07 Å, isostructural to Li2MnO2 and Ni(OH)2. It has an hcp oxide ion sublattice, in which Ni occupies octahedral sites Li tetrahedral sites. If Li is intercalated into LNO, Li2NiO2 forms via a two-phase reaction, which is reversible but with large hysteresis; the transformation between the two phases occurs via displacement of the O-Ni-O sheets normal to the c axis.

4. Physical properties 4.1 Electronic properties Electrical measurements demonstrated that LNO is a semiconductor with a small bandgap. While NiO is an insulator, its electrical conductivity increases dramatically with the introduction of small amounts of Li. A value of 10-1 S/cm is reported for nearly-stoichiometric LNO

[14a, 22b,

45]

. In LNO, Ni is commonly considered to be in the low-spin trivalent state, with an electronic

configuration t2g6 eg1

[46]

. This should induce a JT distortion to remove the ground state

degeneracy, analogous to what happens for NaNiO2, which has a long-range cooperative distortion with octahedra aligned along their longer axes, lowering the crystal symmetry to monoclinic C2/m

[13]

. However, for LNO, such long-range distortion has never been observed

experimentally. Two long (2.09 Å) and four short (1.91 Å) Ni-O bonds have instead been observed on the local scale by EXAFS and neutron PDF studies

[23, 46-47]

, but their relative

arrangement is still under debate, despite first principles calculations performed by a number of authors to find the ground state of LNO

[24b, 48]

. All results confirm that a JT distorted ground

state is more stable than a non-distorted one, and a zigzag arrangement of JT octahedra in P21/c space group is more stable than a collinear arrangement. However, other ground states with Ni2+/Ni4+ disproportionation or with Ni trimers have been suggested and cannot be completely ruled out

[23, 49]

. In this regard, the true nature of the LNO ground state is yet to be 16

determined experimentally. The discrepancy between computations and experiments can be ascribed to the off-stoichiometry present in real LNO samples that might prevent the stabilization of the ground state structures predicted by DFT methods. Several authors also pointed out that Ni in LNO is likely not in a pure trivalent state, and some even suggested that a hole on the oxygen atom, rather than on the nickel, may exist

[50]

.

Although a large number of experiments indeed confirms the trivalent nature of Ni in LNO, a significant amount of charge redistribution occurs with oxygen ligands [51]. In this respect, there is a huge difference between LiCoO2, where Co is without doubt purely trivalent in the stable low-spin t2g6 configuration, and LiNiO2.

4.2 Magnetic properties Magnetic properties of LiNiO2 are strongly affected by the stoichiometry and ordering degree, and this has led to several different interpretations of the results obtained

[14b, 38, 46b, 52]

. In

particular, the fraction of Ni present in the Li layer plays a key role in the magnetic behavior [53]

. An ideal stoichiometric LNO shows spin glass-like behavior, with a typical Curie-Weiss

dependence above Tc [38, 52d, 53]. In fact, every magnetic Ni layer is separated from the next one by three non-magnetic layers (O-Li-O), and thus only in-plane magnetic interactions are possible. Ni3+LS has S = ½ and a spin-only magnetic moment of 2√S(S + 1)μB = 1.73μB. Since Ni is placed on a triangular lattice, the antiferromagnetic coupling of spins is frustrated and no long-range order arises, resulting in spin glass-like behavior. However, when Ni2+ (S = 1) is located in the Li layer, a magnetic transition around 200-300 K appears

[46b, 52d]

. Ni2+ in the

interslab interacts with the Ni in the slab, locally stabilizing magnetic clusters (although their ferrimagnetic or ferromagnetic nature is unclear)

[53a, 54]

.

Thus, magnetic measurements are a useful tool to determine the stoichiometry of LNO. Figure 4 shows the reciprocal magnetic susceptibility as a function of T for Li1-zNi1+zO2. Nearstoichiometric LNO shows typical linear Curie-Weiss behavior. When the fraction of Ni in the 17

Li interslab increases, the linearity for χ-1 vs. T is lost, indicating the occurrence of magnetic ordering [17, 55] at a temperature increasing with z.

Figure 4: Variation of reciprocal magnetic susceptibility vs. temperature with the offstoichiometry of different LNO samples. Data reproduced from [17].

4.3 Thermodynamics Thermodynamic data for LNO is scarce, especially that obtained experimentally. Wang and Navrotsky evaluated the enthalpy of formation of LNO, with data obtained from calorimetry measurements, reporting a value of ΔHf° = -593 kJ/mol at 298 K and 1 atm from elements, and ΔHf° = -56.2 kJ/mol from oxides

[56]

. A Gibbs energy change for formation from elements of

ΔGf0 = -514.96 kJ/mol [57] or -600 kJ/mol [58] at 300 K was also reported from calculations. The heat capacity of LNO as a function of temperature was measured and found to be significantly higher than for LiCoO2 [59].

5. Synthesis methods Different methods have been employed over the years in order to obtain high-quality LNO. Although often claimed, samples with a thorough structural characterization, where a low z value was demonstrated, are less common. In most cases, solid-state synthesis yields the best results, especially if precursors are prepared by co-precipitation methods. In this section, the main synthesis techniques are reviewed, with special attention to the critical parameters (temperature, gas flow, precursors etc.), allowing to obtain near-stoichiometric LNO. We note 18

that the vast majority of samples are obtained as powders, while only one report exists on the synthesis of LNO single crystals (actually obtained as Li0.92Ni1.08O2)[60].

5.1 Solid-state reaction The solid-state method is the most commonly employed for the synthesis of LNO. Solid lithium and nickel precursors are intimately mixed and subsequently calcined. Ohzuku et al.

[7d]

and

Rougier et al. [17] performed some of the first systematic studies using this method, evaluating the key factors that must be controlled to obtain near-stoichiometric samples. The major findings can be simply summarized: common nickel precursors contain Ni2+, thus oxidation must occur. Air or O2 flow are used, with the latter consistently yielding better results. The effect of O2 flow rate has been studied, showing that a sustained flow is needed [61]. However, that is mainly because mass transport is required during the synthesis (for example, H2O or CO2 removal from the precursors´ decomposition). The temperature of the calcination process is crucial; high temperatures (800 °C and above) could be used to obtain well-crystallized samples in short times, but also imply a lower μ(O2), incompatible with the need for full oxidation to Ni3+. Furthermore, even if LNO forms, at high calcination temperature it is unstable and decomposes towards off-stoichiometric Li1-zNi1+zO2 (eq. 2)

[62]

. For this reason, lower

temperatures in the vicinity of 700 °C are often reported as optimal. Even lower temperatures are possible, but may lead to products with remaining impurity phases of unreacted precursors, long calcination times and small particles of poor crystallinity. Kanno et al. reported the formation of mixtures of NiO and Li2NiO3-y for calcination temperatures below 500 ºC [18], while Lin et al. detected the presence of unreacted NiO and Li2CO3 below 600 ºC

[39a]

. Therefore,

most LNO samples have been made for 650 °C < ϑ < 750 °C, with 700 °C being the most widespread choice [7d, 20a, 39a]. The nature of precursors is also crucial and complementary to the choice of calcination temperature. To obtain LNO, the decomposition temperature of the lithium precursor must be as low as possible, to promptly react with nickel without the need for high temperatures, where 19

Li (or rather Li2O) may be lost. For this reason, hydroxides, nitrides and oxides/peroxides of lithium are preferred precursors. Li2CO3, often used in the synthesis of NCM, should be avoided, as it decomposes above 725 °C

[7d, 20a]

. The loss of Li during the synthesis is an

important limiting factor, caused by the high vapor pressure of Li2O. It has been shown

[32, 63]

that Li2O reacts with O2 to form Li2O2, which is easily vaporized at high temperature. Hence, an excess of Li precursor is commonly used to minimize off-stoichiometry and also to compensate for the Li loss during the calcination process, in amounts varying from 1% to a 4fold excess. High Li/Ni ratios have the advantage that Li off-stoichiometric Li1-zNi1+zO2 becomes less likely, but require a washing process after calcination to remove unreacted precursors. This may have a non-negligible effect on LNO and especially on its surface, as we discuss in section 5.6. Arai et al. were the first to report a synthesis of LNO with excess Li. They used a large excess (up to n(Li):n(Ni) = 4:1) and the sample was washed after calcination to remove the remaining Li salts, mainly present in the form of surface carbonates [34]. Recently, Xu et al. also studied the effect of Li excess, remarking that a 2% or 10% Li excess does not make a difference in the bulk LNO structure and its stoichiometry, but that it affects the surface properties of the compound [64]. One should also note that when large amounts of Li excess are employed, the resulting material can be placed in the region of the phase diagram that contains both Li2NiO2 and Li2NiO3, depending on the oxidative power of the medium (μ(O2)). Consequently, products other than LNO may be obtained [7a, 65].

5.2 Co-precipitation: industrial process Methods involving a solution-based mixing of precursors are usually employed to improve homogeneity and are especially important when attempting doping of LNO. Among these, coprecipitation was used by several authors

[40b, 66]

and is particularly noteworthy because of its

scalability to produce huge batches of layered materials, and is commonly employed in industry [67]

. 20

Co-precipitation typically involves different steps [68]. First, an aqueous solution is made, where nickel is added in the form of sulfate or nitrate and it is pumped to a continuously stirred tank reactor. Then, NaOH is added to precipitate Ni(OH)2 (most divalent metals or metal mixtures can also be precipitated in the form of hydroxides). A NH4OH solution can also be used as chelating agent. Careful control of solution concentration, temperature, stirring, pH etc. are crucial to yield homogeneous hydroxides with controlled particle shape and size. The coprecipitated precursors are then mixed with the Li precursor, followed by a calcination step analogous to the one described in section 5.1.

5.3 Sol-gel In a typical sol-gel synthesis, nickel and lithium precursors (usually nitrates, hydroxides and/or acetates) are dissolved in water or a mixture of water and ethanol/methanol

[69]

. A chelating

agent, such as adipic acid [70], oxalic acid [39b], citric acid [39b, 71], acrylic acid [39b], polyvinyl butyral (PVB)

[72]

or triethanolamine (TEA)

[39b]

is commonly added to stabilize the sol, but syntheses

without any chelating agents have been reported as well [40a, 73]. During the process, and after a proper stirring, the solution is evaporated until the formation of a gel. This is usually preheated at 400-600 °C in air in order to combust the organic compounds, and finally calcined under the conditions described above for the solid-state method. Due to the intimate powder mixing, this synthesis method reduces calcination times, often yielding close-to-stoichiometric materials with low crystallite size[39b, 40a]. The Pechini method, where the precursors are mixed with citric acid and ethylene glycol in order to develop a polyesterification reaction

[74]

, was

employed as well.

5.4 Combustion Several papers are devoted to the synthesis of LNO by the combustion method, consisting in heating a mixture of metal precursors (usually nitrides) and a fuel with high combustion enthalpy. Thus, it is possible to calcine at low temperatures, as the combustion of the fuel 21

provides the necessary energy [75]. In recent years, urea has been the preferred fuel [75b, 76], but glycine is also suitable [75a].

5.5 Other methods Other synthesis routes include hydrothermal methods (often involving ion exchange from NiOOH)

[77]

or the utilization of microwaves as the energy source

[78]

. Mechanochemical

synthesis results in rock salt Li1-zNi1+zO2, which layers only after annealing above 600 °C

[79]

.

The preparation of thin films of LNO by different techniques has also been reported [80].

5.6 Washing LNO The use of excess Li precursors likely leads to some Li remaining on the surface of LNO, usually in the form of unreacted Li2O (swiftly reacting with moisture to form LiOH or Li2CO3). The presence of the latter on the surface of LNO appears to be unavoidable, as even small amounts of CO2 react quickly with the layered oxide surface. XPS analysis of samples synthesized under different conditions and atmospheres always revealed some fraction of Li2CO3 at the surface [64, 81], which is also known to be a cause of capacity decay in NCM and NCA cathodes [82]. Although this is rarely discussed in the literature, washing them (and LNO) to remove such residues is often crucial for the electrochemical performance [34, 83]. The bulk of the LNO primary particles is not affected by washing, as observed by XRD [34], but their surface may react with water, so that a carefully controlled washing process should be devised. Excess water was proposed to lead to severe Li extraction from the layered structure, while ethanol was suggested to be a promising (but more costly) alternative be noted that LNO is structurally related to NiOOH polymorphs

[84]

. It should also

[77c]

, so that washing using

water may result in Li/H exchange, as recently suggested for NCM materials

[85]

. In general,

the above considerations about water apply also to the moisture sensitivity of the compound, thus handling in dry atmosphere is crucial. 22

6. Electrochemistry Dahn et al.[7a] first showed that LiNiO2 can be oxidized via delithiation according to: LiNiIII O2 → NiIV O2 + Li+ + e−

(3)

The complete reaction yields a high theoretical specific capacity of 275 mAh/g. The deintercalation reaction occurs at an average voltage close to 3.8 V (Figure 5a), which is well within the stability window of most electrolytes. LNO is typically cycled between 3.0 and 4.3 V. Although values achieved in the years 1990-2000 were mostly lower (qch = 220-240 mAh/g, qdis = 200-220 mAh/g), recent work shows capacity values achieved for LNO in the 1st cycle exceeding qch = 255 mAh/g and qdis = 245 mAh/g

[66c, 86]

, very close to the theoretical values.

Nevertheless, the commercial use of LNO has been hindered by: a) the difference between initial charge and discharge capacity, typically ≈10-15%, limiting the Coulombic efficiency, as discussed in the following. b) fast capacity fading upon cycling (observed mostly when cycling above 4.1 V)

[86a, 87]

,

discussed in sections 7.1 and 7.3.

c) safety issues due to thermal instability of the deintercalated phases LixNiO2 [5, 88], discussed in section 7.2.

23

Figure 5: a) Typical charge-discharge curve of LixNi1.02O2/Li cell and b) inverse derivative curve –dx/│dV│ from [89]. c) Mechanism causing the capacity loss in LNO during the 1st cycle. Reprinted from [8], Copyright 1997, with permission from Elsevier. Upon lithium deintercalation, the unit cell parameters change as is characteristic of layered compounds: a decreases due to oxidation of Ni3+ to Ni4+ (r(Ni4+) < r(Ni3+LS)), while c expands due to the increasing repulsion between NiO2 slabs (which are screened by a decreasing amount of Li). However, when the fraction of lithium becomes too low, the interlayer collapses and c steeply decreases (Figure 6). The net result is a large volume decrease ΔV/V ≈ -10% [7a, 7d, 8, 40c, 90]

, which is one of the factors adversely affecting the long-term cycling behavior of the

material. A large difference between the initial charge and discharge is typically observed in LNO, and has initially been ascribed to the formation of “inactive domains”

[91]

later understood in relation to the presence of extra Ni2+ in the Li interslab

. This has been

[92]

, remembering

that off-stoichiometric LNO should be written as [Li1-zNi2+z]3a[Ni2+zNi3+1-z]3b[O2]6c. Careful electrochemical experiments have evidenced that, upon charge, a small reversible 24

intercalation range is observed in the vicinity of the starting composition, whose extent increases with z. If a critical fraction of Li is deintercalated, a part of reversibility is lost in the discharge. The reversibility loss increases with z and, importantly, the capacity lost can be recovered by slow discharge at low voltage (a potentiostatic process at 2 V for 1 month)

[92]

.

These observations indicate that the phenomenon is a kinetic effect, and that it depends on the value of z. As shown in Figure 5c, when a sufficient amount of Li is deintercalated, z Ni2+ cations in the interslab are oxidized to smaller Ni3+. Upon discharge, this small cation in the interslab will hinder Li diffusion in its six neighboring sites due to shrinkage of the local environment. A small z has a large effect on the discharge capacity of LNO, since for e.g. z = 0.02, Li access to 7z sites is hindered, resulting in a 14% capacity loss. Hence, the importance of taking extra care in the synthesis to minimize z. On the other hand, it was recently shown that lithium excess (≈10%) rather than deficiency does not negatively affect the electrochemical performance, but instead improves the capacity retention of LNO

[29a]

.

However, reports on the over-lithiation of LNO are limited and no clear conclusions on the effect can be drawn today. We note that, regardless of the value of z, significant kinetic hindrance is reported for LNO upon discharge at a potential close to 3.5 V, where a peak in the dQ/dV plots appears after the first charge, but only at sufficiently low current rate (already at C/5, the peak is suppressed and the capacity lost)

[86c]

. It is well-known that in layered compounds Li diffusion between

neighboring octahedral sites occurs via an intermediate tetrahedral site (Figure 1) [19]. In LNO, at least two factors contribute to increase the activation barrier for Li diffusion at high Li content. Firstly, Ni in the Li layers decreases the interslab space, thus decreasing the tetrahedron height and increasing the activation barrier for Li diffusion [93]. Secondly, as shown by Van der Ven et al.

[94]

, the lowest energy barrier is obtained for diffusion via di-vacancies in the Li lattice, but

they are only present at a Li content lower than ≈0.8. At higher lithiation, the small interlayer spacing and the low probability to encounter di-vacancies (and the fact that an ordered composition, undetected to date, may exist at x = 3/4 or 6/7, for example, related to the dQ/dV peak at 3.5 V), all cooperate to hinder Li kinetics. 25

Figure 6: a) Phase diagrams proposed from experimental and computational studies from [34, 89-90, 95] . White and blue areas denote single- and two-phase regions, respectively. b) and c) Evolution of unit cell volume and unit cell parameters (aH and cH for hexagonal phases, aM, bM, cM and β for the monoclinic phase) from in situ XRD and the corresponding voltage profile. Reproduced with permission from [86c], Copyright 2018, The Electrochemical Society.

6.1 Phase transitions upon Li (de)intercalation Lithium deintercalation from LNO has been studied in detail by a number of authors, especially by means of ex situ [7d, 34] and in situ XRD [7a, 86c, 90, 96]. Several first-order phase transitions were reproducibly observed, consisting of two-phase regions separating single-phase ones, as depicted in Figure 6. These can be well-correlated with the electrochemical characteristics of the materials and especially with the inverse derivative curve -dx/│dV│(Figure 5b). As lithium is extracted from LixNi1+zO2 (0 < x < 1-z), it is widely accepted that the reversible sequence of stable phases is as follows: H1 (pristine LNO) → M → H2 → H3 (the transition to H4 is discussed in section 6.3), where H denotes a hexagonal phase and M a monoclinic one 26

(Figure 6a). In some works, the H letter is replaced by R for rhombohedral. As previously mentioned, there is a significant overall change of unit cell parameters (Figure 6c). For low degrees of deintercalation (x > 0.85), the original rhombohedral structure is maintained, labeled as H1. At x ≈ 0.8, the monoclinic phase M starts to crystallize, until x ≈ 0.75, where it becomes a single phase. The monoclinic phase is observed by a splitting of the (101) and (104) Bragg reflections, forming the new (201) - (110) and (202) - (111), respectively. This M phase is stable up to x ≈ 0.5, where a 2nd hexagonal phase H2 starts to form. H2 (a ≈ 2.8221(1) Å, c ≈ 14.404(1) Å [40c]) is stable until high delithiation states (x ≈ 0.25), where a two-phase domain leads to the 3rd hexagonal phase H3 (highly delithiated). In this case, both a and (especially) c lattice parameters are considerably smaller (a ≈ 2.8154(7) Å and c ≈ 13.363(6) Å [40c]), similar to the LiCoO2 case [97]. The H2-H3 transition involves a shrinkage of the unit cell and a large volume decrease (Figure 6b), inducing significant strain in the material [86a, 98]. Importantly, observations on the width of the M phase are not very consistent between different authors. Yang et al.[96b] did not observe the monoclinic phase at all. This was shown to be due to the off-stoichiometry. Indeed, the monoclinic region is mainly stabilized by in-plane Livacancy orderings (see next section). Since Ni in the Li interslab strongly affects and disrupts this ordering, and the Ni2+ ions act like pillars in the Li layer, making the structure less flexible, it also prevents the phase transitions that occur during Li (de)intercalation [21, 46b]. This induces a “smoother” voltage curve (i.e. plateaus are less well-defined) and a solid solution-like behavior for LNO samples far from the stoichiometric one. Such dependence of structural evolution during (de)intercalation on the stoichiometry has led to some discrepancies in the literature, even recently

[86a, 86b, 96b]

. We believe this issue has now been well-understood and

that well-stoichiometric LNO always shows a monoclinic domain.

6.2 Nature of the monoclinic domain The nature of the monoclinic phase M and of the end member phase H3 (and H4) obtained at the end of charge have been object of extensive investigation. The former is a distortion of the 27

hexagonal cell, shown in Figure 7, which can be described with a transformation matrix (H → M): −1 1 1/3 [ 1 1 −1/3] 0 0 −1/3 This leads to unit cell parameters a = 4.99 Å, b = 2.83 Å, c = 5.07 Å and β = 109.7°, where the distortion in the basal plane means that a ≠ √3b and the interlayer distance is not simply c/3, but rather (c/3)/sin(β). The values above are approximate, as the monoclinic region extends over a large Li content and the cell parameters vary accordingly (see Figure 6). Although early work attributed the formation of a monoclinic region to the JT distortion of Ni3+ [34]

, the explanation is not convincing, since the pristine LNO, which contains the largest fraction

of Ni3+, does not display such distortion (unlike NaNiO2). This hypothesis was also discredited by X-ray absorption experiments

[99]

. It was later clarified experimentally that the monoclinic

region is constituted by various domains where Li-vacancy orderings are stabilized. Specifically, Delmas and coworkers used EXAFS, XRD and electron diffraction (ED)

[8, 100]

to

investigate a sample of composition Li0.63Ni1.02O2, seen as the intermediate of the monoclinic region, extending roughly between 0.5 < x < 0.75. Superstructure reflections were observed by ED and could be indexed either with an 8-times bigger supercell of the monoclinic one (2a x 2b x 2c, F-centering) or with a 4-times bigger one, with a standard C-centered monoclinic cell. The latter supercell, described in space group C2/m (a = 10.138 Å, b = 5.656 Å, c = 8.223 Å and β = 145.178°), contains three independent Li sites and allows to explain the cause of the monoclinic distortion. Indeed, in this superstructure, the limit compositions x = 0.5 and x = 0.75 can be described as Li0.5□0.25□´0.25NiO2 and Li0.5Li0.25□´0.25NiO2. So, the monoclinic distortion is driven by the strong Li-vacancy ordering present between these three sites

[100]

(Figure 7c). Later on, other Li-vacancy orderings were found by the same group using ED, occurring at compositions Li0.33NiO2 and Li0.25NiO2

[89, 101]

. Furthermore, twinned crystals have

also been observed. The authors propose that these compositions stabilize hexagonal supercells, characterized by a´ = √3a and a´ = 2a, related to 2:1 and 3:1 Li-vacancy orderings, 28

respectively. However, we note that to date no other study has experimentally confirmed the models of these two structures, nor shown that the supercells observed locally by ED are extended on the long-range, for example, by neutron diffraction.

Figure 7: a) Structural relationship between hexagonal and monoclinic setting in LNO. Reprinted from [58], Copyright 2012, with permission from Elsevier. b) Electron diffraction pattern of Li0.63Ni1.02O2 along the [001] zone axis, showing characteristic superstructure reflections. c) Electrostatic distribution of Li ions and vacancies among the Li layers in LixNiO2 (0.5 < x < 0.75). The small unit cell is deduced from X-ray refinement (normal monoclinic) and the larger one characterizes the superstructure (in the 8-times larger Ftype Bravais lattice). Reprinted from [100], Copyright 1999, with permission from Elsevier.

Theoretical calculations were later performed by Arroyo and Ceder

[95, 102]

. They obtained the

phase diagram of the material from first principles and specifically evaluated the stability of the monoclinic region, which is found to be more pronounced around compositions Li0.75NiO2 and Li0.4NiO2. Their results confirm that Li-vacancy orderings drive the formation of the monoclinic phase, but their ground state structures also indicate a contribution of the Ni3+ JT-active cations to their stabilization. In fact, the relevant interactions are the in-plane Li-Li repulsion, but also 29

a long-range Li-Li attractive interaction across planes, via JT-active Ni3+. Ordered phases are found at room temperature for x = 1/4, 1/3, 2/5, 1/2 and 3/4. The first two are supported by the above-mentioned ED studies

[89, 101]

, 2/5 was reported for the first time and to date has not

been experimentally confirmed, while 1/2 and 3/4 are indeed the limits of the superstructure region found by Peres et al. [100]

6.3 Nature of the highly oxidized phases Most early literature identifies the phase H3 found at the end of charge with the end member NiO2 (Figure 6a). Actually, the highly delithiated states of LNO have been shown to have a rather complex behavior, involving at least two phases called H3 and H4

[26, 40c, 103]

. Upon

charge, the H3 phase forms first, having an O3-type oxygen packing and the same overall structure as H1 and H2, but with a significantly smaller interlayer distance. Its composition can be seen as LiεNiO2, with ε ≈ 0.1. This is also in agreement with the earlier phase diagrams shown in Figure 6. The interlayer collapse between H2 and H3 is dictated by a change in the prevailing interaction. For x > 0.3 in LixNiO2, the electrostatic effects prevail and the O-O repulsion causes the interlayer to expand (since Li that screens such repulsion is gradually removed). For smaller x, the system behaves more similarly to transition-metal sulfides and selenides, where the anions are less electronegative. The swift decrease in the c lattice parameter upon Li deintercalation indicates that steric effects start to prevail over the electrostatic ones, thus suggesting that the structure becomes strongly covalent

[26, 40c]

.

Upon further charge, the H3 phase transforms to the fully delithiated H4 (LiεNiO2, with ε ≈ 0), having an O1-type (ABAB) oxygen packing. H4 was found to have a distorted CdI2 structure, isostructural to CoO2 [104], with hexagonal close packed oxygen atoms, C2/m space group and cell parameters a = 4.8754(3) Å, b = 2.8141(2) Å, c = 5.5820(3) Å and β = 125.836(4)°

[26]

. In

general, because of their close compositions, it is particularly difficult to isolate H3 and H4 as single phases. Tarascon et al.

[26]

devised a particularly careful electrochemical oxidation

protocol in order to isolate the H4 phase. Croguennec et al. [103b], although without providing a 30

full structural solution, reported for the same phase the smallest unit cell parameters (ahex = 2.8145(6) Å and chex = 13.039(6) Å), suggesting the lowest reported Li content. The H3-H4 phase transition occurs via layer gliding and is thus prone to creating layer displacement faults. In fact, O1 regions in the O3 phase and O3 regions in the O1 phase (Figure 8a) were commonly observed in XRD as a broadening of certain Bragg reflections

[40c]

and more recently also

confirmed by high-resolution TEM [86a].

Figure 8: a) Structures of the H3 (R3) and H4 phases seen through a section along the (110) plane. Parallelograms represent nickel (in black), lithium (hatched) and vacancies (in white). b) and c) Difference in interactions between the p oxygen orbitals through the van der Waals gap in the AB and ABCABC packings. Reproduced from Ref.[40c] with permission from The Royal Society of Chemistry.

The nature of these faults is related to the presence of Ni in the Li layer, which impedes a proper gliding, and thus for small z creates the defects. However, as z increases above 0.07, the H3-H4 transition is totally eliminated

[40c, 103b]

. Moreover, z is suggested as the cause for

the observation of a distorted CdI2-type structure for the H4 phase, while a perfect LNO would supposedly have the undistorted CdI2 structure

[26]

. The reason for the occurrence of the H3-

H4 transition, namely for a layer gliding that modifies a fcc packing to a hcp one in NiO2, was 31

identified by Croguennec et al.[40c] (Figure 8b-c), who showed how the oxygen p orbitals overlap is minimized by having an O1-type stacking in the H4 structure. In the O3 one, instead, the orbitals would be directly facing each other via an intermediate empty octahedral site, resulting in higher electrostatic repulsion. Clearly, there is a strong energy gain to be obtained by this mechanism only when no Li remains in the interslab. Finally, it should be noted that the reversibility of the H3-H4 phase transition has been studied [103a, 103b]

. Different behaviors were observed in the material, suggesting fluctuations in

composition of LNO not only at the local scale (those responsible for the creation of stacking faults), but also at the crystallite scale. While the H3-H4 transition is mostly reversible, the H4 phase is observed by XRD to gradually decompose irreversibly (e.g. during a 1000 h CV step at 4.45 V) to a phase with O3-type stacking called H3´. This decomposition is dictated by Ni3+ migration into the interslab space, as indicated by the interslab distance of H3´, which is clearly smaller than that of H3 (4.45 Å for H3 and 4.41 Å for H3´). It was suggested that crystallites closest to the ideal stoichiometry form an H4 phase stable at high potential, while the others lead to an H4 phase transforming into the H3´ phase via Ni migration from the slab to the interslab space, thereby affecting the overall reversibility of the material.

6.4 Critical role of Li off-stoichiometry on the electrochemical performance The effect of off-stoichiometry z on the electrochemical performance of LNO is briefly summarized here. For LNO samples with z > 0, the theoretical specific capacity decreases due to the lower Li content. The presence of Ni2+ in the interlayer induces a large irreversibility in the 1st cycle, causing capacity loss due to kinetic reasons. This further worsens the poor Li mobility in LNO during discharge at high lithiation. While LNO undergoes sharp two-phase transitions, indicated by flat regions in the charge-discharge curve, increasing z results in a smoother curve and less well-defined features. The monoclinic domain, mainly stabilized by Li-vacancy orderings, is destabilized by Ni in the Li layer, and can be entirely suppressed. Likewise, the large volume change across the H2-H3 transition is reduced by z. Highly oxidized 32

LNO transforms to an hcp phase called H4, while this transition first creates faults, and then it is entirely prevented, as z increases.

7. Challenges As mentioned above, the synthesis of LNO with a perfect layered structure and stoichiometry already presents significant challenges. In general, the formation energy for different kinds of defects (Li/Ni site exchange, interstitials, O vacancies) is found to be very low in LNO [35-36, 105]. Moreover, the material suffers from several other drawbacks that have so far hindered its commercialization, and these are all related to different forms of instability. In this section, such issues will be reviewed.

7.1

Mechanical instability

LNO is delithiated via a sequence of first-order phase transitions (section 6.1). The most critical ones in terms of stability are those occurring at low Li content, namely the formation of the H3 and H4 phases. Firstly, the H2 → H3 transition (4.15 V vs. Li+/Li) induces a volume change of ≈7% over a narrow composition range [86a, 90]. Such a large ΔV between the two phases means that significant strain is created at the interface between them, mainly at the primary particle level, but also affecting the secondary ones. The strain is then released via formation of cracks during cycling. The cracks along the secondary particles lead to a loss of electrical contact between the primary particles, which has a strong repercussion on the electrochemical performance of the material

[106]

. In addition, some electrolyte may enter through the cracks,

further reacting with the freshly-formed LNO surfaces. The crack generation on LNO seems to start already in the first charge/discharge cycle and to increase with cycling, leading to pulverization of the secondary particles

[86a, 98]

. Mechanical instability, although to a lower

degree, also negatively affects NCM cathodes of any nickel content [106a, 107].

33

Figure 9: Dependence of capacity retention on the upper cut-off voltage and correlation with the secondary particles cracking from SEM imaging. Reprinted with permission from [86a] , Copyright 2017 American Chemical Society.

This mechanical degradation (and the subsequent chemical degradation of the freshly created surfaces) is directly correlated with the upper cut-off voltage (Figure 9). If the region where the H2 → H3 transition takes place is avoided (cycling below 4.1 V), the cracking is minimized [86a, . This corresponds to ≈75% delithiation and would reduce the specific capacity well below

98]

200 mAh/g even for the best performing LNO samples. Under this condition, long-term cycling of LNO has already been reported in the ´90s. Broussely and coworkers

[87]

demonstrated

LNO/graphite cells retaining 60% of their initial capacity after 1200 cycles at 1C rate, by limiting the voltage range to 2.5 - 4.1 V at room temperature. Finally, we already mentioned the instability of the H4 phase in section 6.3. This phase can irreversibly transform to a related hexagonal one called H3´ via Ni3+ migration into the interslab, thus making part of the material inactive. It was also shown that this process likely occurs only in crystallites whose stoichiometry is more far off from ideal LNO [103a].

7.2 Thermal instability The poor thermal stability of LNO is a major drawback, as it adversely affects the safety of LIBs. Indeed, ideal cathode materials should undergo little or no exothermic reaction when heated well above 200 °C because of the presence of flammable electrolytes. The instability 34

of cathodes may also result in the release of O2 into an electrolyte heated above its flash point, leading to a violent reaction. Stability in the pristine state is necessary, but not sufficient; electrode materials should be thermally stable even during oxidation in the delithiated state. In this respect, LNO was found to be thermally stable in the pristine state [5], but very unstable when delithiated [108]. The limited thermodynamic stability of LixNi1+zO2, especially for x < 0.5, does not come as a surprise. Computational results

[24b]

demonstrated that no compound on the LiNiO2 – NiO2 line in the

phase diagram is stable as a layered structure at 0 K. Even at composition Li0.5NiO2 the computed stable phase has a spinel structure (see section 3.3). So all delithiated LixNi1+zO2 layered phases are at most metastable, and the situation only worsens with increasing temperature, where the oxygen chemical potential contribution becomes more dominant (entropy term), and thus destabilizes highly oxidized Ni4+ in favor of Ni2+. The first rational study on the thermal stability of LNO was performed by Dahn et al.

[5]

They

compared the stability of several materials under an inert gas atmosphere using thermogravimetric analysis-mass spectrometry (TGA-MS) and observed that Li0.3NiO2 is less stable than Li0.4CoO2 and λ-MnO2, as it releases O2 at lower temperatures. They also reported that lower lithiation levels correlate with lower temperatures at which the maximum amount of O2 release occurs. It was proposed [5, 42, 108] that for highly delithiated materials (LixNi1+zO2; x ≤ 0.5) the decomposition follows a two-step process: around 200 ºC, the layered structure transforms exothermically to the spinel phase (releasing O2 for x < 0.5, but not for x ≥ 0.5) and at temperatures of ≈300 ºC, this spinel transforms to the rock salt structure, accompanied by more O2 release [5, 108-109]. Later on, other authors confirmed the thermal instability of delithiated LNO by TGA, differential scanning calorimetry (DSC) and in situ/ex situ XRD [109-110]. As schematically depicted in Figure 10, regardless of the lithiation degree (Li content 0.5 or 0.3), first Ni migrates to the interslab, followed by the displacement of Li into tetrahedral sites. This arrangement stabilizes a defective spinel structure (or pseudo-spinel). Then, concurrent to O2 evolution and via a two35

phase region a disordered layered structure is observed, with Li/Ni mixed in the slab and interslab and finally a fully disordered rock salt structure is formed.

Figure 10: Sequence of steps observed for the thermal decomposition of delithiated LNO. The overall reaction follows eq. (4). Adapted from [110b].

The effect of electrolyte on the heat release of delithiated cathodes was also studied by MacNeil et al.

[88]

. The authors used 1M LiPF6 in EC:DEC (1:1 ethylene carbonate/dimethyl

carbonate) and classified various electrode materials based on the amount of heat released and its onset temperature. The results confirmed that LNO is the least safe material due to strong O2 and heat release already at 200 °C, starting from 4 V vs. Li+/Li. Oxygen can exothermally react with the electrolyte, causing a violent vent

[111]

. Zhang et al. observed by

DSC that if the electrolyte is removed from the partially delithiated samples, the exothermic decomposition peaks almost disappear, while the process becomes exothermic again when fresh electrolyte is reintroduced

[112]

. This supports the fact that thermal decomposition of

delithiated LNO does not release a significant amount of heat itself, but that it is the electrolyte taking part in the reaction releasing most of the heat. It was also shown that large primary particles are more thermally stable than smaller ones, while secondary particles size plays no role [113].

7.3

(Electro)chemical instability

36

As mentioned above, delithiated LNO is thermodynamically only metastable. Ni4+ formed in highly delithiated states may react when in contact with the electrolyte. The mechanism of Ni4+ reduction is related to the O2 release and both are a result of the position of LixNi1+zO2 on the phase diagram. Each delithiated phase will likely decompose to a stable LiδNi1-δO rock salt phase (commonly called NiO-like), O2 and possibly Li2O (eq. 2). Depending on the delithiation level, spinel LiNi2O4 can also be one of the products. Several authors have indeed shown that the highly reactive Ni4+ is reduced to Ni2+ (mainly forming NiO-like rock salt phases) in LNO

[64, 86a, 86b]

(Figure 11a) and related compounds (in

particular, LiNi0.8Co0.2O2 cathodes were subject to intensive investigation in the US followed by NCM and NCA of different stoichiometry (Figure 11b-d)

[10, 12b, 106b, 115]

[114]

,

). The rock

salt-type phase formation mainly occurs at the particle´s surface, where O2 can freely evolve. While O2 is released from the layered LNO structure, a densified zone close to the surface grows. This can be as thick as 20 nm and mainly formed of rock salt structure (Figure 11) [86a].

37

Figure 11: a) HRTEM image of a primary particle from the LNO cathode cycled to 4.1 V and the corresponding FFT images from the selected regions. From III to I, there is a progressive transformation from the original 𝑅3̅𝑚 structure to the NiO-like 𝐹𝑚3̅𝑚 structure at the surface. Reprinted with permission from [86a]. Copyright 2017, American Chemical Society. b) Atomic resolution ADF-STEM images of NCM particles after 1 full cycle (2.0-4.7 V). The blue arrow indicates the surface reconstruction layer. (c,d) Corresponding FFT results showing the surface reconstruction layer (𝐹𝑚3̅𝑚 [110] zone axis) and the bulk layered structure (𝑅3̅𝑚 [100] zone axis). Reproduced with permission [116] , Copyright 2014, Macmillan Publishers Limited.

The formation of a rock salt phase is accompanied by release of O2 from the lattice, according to: Li𝑥 NiO2 → (1 + 𝑥) Li

𝑥 1+𝑥

Ni

1 1+𝑥

O+

1−𝑥 O2 2

(4) 38

or, if the possible loss of lithium in the form of Li2O is considered (see also eq. (2)): 𝑥

δ

1

Li𝑥 NiO2 → [2 − 2(1−δ)] Li2 O + 1−δ Liδ Ni1−δ O +

2−𝑥−δ(3−𝑥) O2 4(1−δ)

(5)

Note that Li2O would be hardly detected, as it rapidly reacts with water or CO2 to form LiOH and Li2CO3, respectively. Although there is no doubt about the decomposition of the LNO surface to yield a rock salttype phase, the characterization of its structure is challenging

[117]

and the exact Li/Ni ratio

mostly unknown. This is significant because the properties of rock salt LiδNi1-δO strongly depend on the Li concentration. Urban and coworkers

[118]

showed that a Li excess of 10% is

sufficient to enable macroscopic Li diffusion in a fully disordered fcc lattice, so in principle a percolating network of Li could also be present in LiδNi1-δO and the rock salt phase could passivate the electrode while retaining reasonable ionic conductivity. Moreover, new nickel oxides have recently been predicted and could result from the decomposition of LNO, some possessing good Li diffusivity [24b]. However, so far experimental observations suggest that the formation of rock salt LiδNi1-δO at the surface is deleterious, as it increases the cathode´s impedance and ultimately limits the capacity. It is difficult to separate the impedance increase due to the rock salt-type phase, and the one due to a different issue related to the reactivity of the LNO surface towards the electrolyte. As shown for different NCMs and NCAs, Ni4+ ions can induce electrolyte oxidation, forming several species, most of them solid, that coat the surface of the cathode material

[10]

. In LNO, the Ni

content is higher and so is the reactivity. Indeed, recent computational results revealed that polar facets with exposed oxygen lower the Fermi energy of the material, facilitating electron transfer from the electrolyte

[119]

. The same authors also showed that in general Ni-exposed

surfaces are the least stable ones, while Li-exposed ones are the most stable. The formation of solid species such as LiF, LixPFy and LixPFyOz has been reported as well as a significant release of gas (CO and CO2)

[114c, 114e, 120]

. In addition, considerable amounts of

Li2CO3 and polycarbonates were detected as a consequence of the interaction of reactive 39

oxygenated species, CO2 and Li+ ions

[114c, 121]

. All these solid phases form a cathode solid

electrolyte interphase (cSEI) [106b]. However, the main drawback of such cSEI formation is that instead of acting like a protective coating, as in the case of graphite, it grows with cycling, contributing to the progressive impedance increase and capacity decay of LNO [115b]. The limited thermodynamic stability of LNO and its surface also has important implications regarding handling and storage. The exposure of fresh LNO to air leads to modification of lattice parameters, capacity decrease and to growth of Li2CO3 on the surface [84, 122]. Although different mechanisms were suggested that assume the oxidation of LNO upon carbonation [82b]

, the process is instead more likely to induce Ni reduction [122] according to: LiNiO2 +

𝑧 1+𝑧

CO2 →

1 1+𝑧

Li1−𝑧 Ni1+𝑧 O2 +

𝑧 1+𝑧

Li2 CO3 +

𝑧 2(1+𝑧)

O2

(6)

Liu et al. reported that the initial capacity of LNO was reduced from 215 mAh/g to 165 mAh/g after 1 month storage in air and to almost complete electrochemical inactivity after 1 year [122]. In the same work, an increase in Ni2+ fraction at the surface of the air-exposed samples was detected by XPS as well as an important modification of the O 1s signal, attributed by the authors to the presence of active oxygenated species. They proposed that such species and the spontaneous reduction of Ni3+ to Ni2+, accompanied by an increase of the disorder degree in the layered structure, make the material highly reactive towards H2O and CO2 present in air, leading to the formation of Li2CO3 and LiOH. Furthermore, they reported a deterioration even for the samples stored under an Ar atmosphere, although considerably slower

[122]

. As

previously mentioned, surface Li salts can be eliminated by careful washing procedures (section 5.6).

8. Elemental substitution / Doping Quickly after realizing that LNO presents the issues we summarized in the previous sections, several research groups implemented strategies to mitigate them. The most developed one is metal substitution in LNO with foreign cations, either replacing Ni or Li. Substantial fractions of 40

Ni were initially substituted by other metals, while nowadays rather small fractions are substituted (≈1 mol%), which can be referred to as “doping”. For simplicity, in the following we will use “substituents” and “dopants” interchangeably, and denote them with y. The first ones to be investigated were Mn

[123]

and Co

[7e, 124]

, followed by Fe

[125]

, Al [110a] and several others.

Al, Co, Mg and Ti have been the most studied to date and reviews on the topic are available [20a, 126]

. Here, we will briefly discuss the strategies that in our opinion are significant. Special

emphasis is placed on the use of transition metals. Table 1 summarizes key properties of doped LNO. One should note that, to the best of our knowledge and although included in various patent applications, Sc and V have not been reported in the scientific literature as dopants, while Cr was attempted, but without success [127]. Table 1: Summary of widely investigated dopants (D) and their effect on LiNi1-yDyO2. Columns contain the maximum y reported (preserving the layered structure), whether the end member LiDO2 has a layered structure (M = exists but is metastable), the variation of a and c unit cell parameters, c/a ratio and unit cell volume (arrows indicate increase/decrease), the cations occupying the Li site (3b) and the Ni site (3a). D

Solubility limit: ymax

LiDO2 layered?

a

c

c/a

Vol.

Cations on Li site

Cations on Ni site

Refs.

Co

1

Yes

↓

↓

↑

↓

Li, Ni (y < 0.2), Li (y ≥ 0.2)

Ni, Co

[7e, 15, 26,

Li, Ni, Fe

Ni, Fe (y < 0.3) Ni, Fe, Li (y > 0.3)

[20a, 39d, 125,

Mn, Ni (small y and no Li excess). Mn, Ni, Li (large y) Ni, Ti

[123, 130]

66a, 88, 114e, 114f, 128]

Fe

0.3

No (M)

↑

↑

slight↑

↑

Mn

0.5

No (M)

↓

↑

↑

≈ const.

Li, Ni, Mn

Ti

0.2

No

↑

↑

↓

↑

Li, Ni

129]

[33a, 33b, 130e, 131]

Al

0.5

No (M)

↓

↑

↑

↓

Li, Ni

Ni, Al

[66b, 110, 132]

Mg

0.2

No

↑

↑

slight↓

↑

Li, Mg (y < 0.05); Li, Mg, Ni (y ≥ 0.05) -

Ni (y < 0.05) Ni, Mg (y ≥ 0.05)

[55, 88, 132e,

-

[39c, 39e, 134]

Ga

0.05

No (M)

↑

↑

≈ const.

↑

Other dopants and maximum y: W (0.01), Zr (0.014), Cu (0.1), B (0.05), Ce (0.2), Y (0.05), Nb (0.01), Sb (0.25), Zn (0.01), In (0.05), Tl (0.05), Ca (0.03).

133]

[39c, 66c, 132e, 135]

8.1 Cobalt

41

Due to the structural similarity of LCO and LNO and the similar ionic sizes of trivalent Co (0.545 Å) and Ni (0.56 Å), the full solid solution LiNi1-yCoyO2 can be synthesized [7e, 15, 66a]. The products are seen as homogeneous solid solutions by XRD, although a tendency to Co clustering has been observed by NMR

[128a]

. Co3+ is a 3d6LS cation particularly stable in octahedral

environment and slightly smaller than Ni3+, so with a larger size difference with Li+. Hence, LCO is a well-layered oxide that does not display a tendency to off-stoichiometry towards the parent rock salt phase, as opposed to LNO. For this reason, Co is known to facilitate the synthesis of well-layered LNO and Li off-stoichiometry decreases with increasing Co content . However, a relatively high fraction of Co (≈ 20%) is needed in order to fully suppress the

[66a]

presence of Ni on the Li site [15, 128b, 128c]. Doping LNO with Co decreases both the a and c lattice parameters, and thus the unit cell volume, but the c/a ratio is increased. Co also (slightly) improves the thermal and mechanical stability of delithiated LNO

[26, 88]

. In terms of

electrochemical performance, LiNi1-yCoyO2 cathodes are reported to perform well and mitigate the issues of both end members. Since the Co3+/Co4+ redox couple is active, in principle, no capacity loss results from the substitution. The best compositions are found close to LiNi0.8Co0.2O2 [66a, 128c]. First discharge capacities up to 210 mAh/g have been demonstrated (y = 0.26, 4.3 V cut-off) [128d], while prototype full cells showed 157 mAh/g at C/2 (y = 0.18, 4.1 V cut-off) [128e]. However, despite the improved cycling stability as compared to LNO, a large pilot study

[114e, 114f]

clearly demonstrated that the material is still not stable enough for commercial

applications and suffers from the same electrochemical stability issues as LNO.

8.2 Iron The doping of LNO with Fe was initially demonstrated by Reimers et al.

[125]

. By solid-state

synthesis, a layered compound was only obtained for y < 0.3, while for y > 0.5 the material disorders with a rock salt structure. For 0.3 < y < 0.5, the layered and rock salt phases coexist. This can be understood by the fact that LiFeO2 is most commonly found as a disordered rock salt structure (called α). It can be obtained as a metastable layered phase only by soft 42

chemistry routes

[129a]

. In fact, by ion exchange the full layered solid solution could be made

[129b]

. For any y, a significant fraction of transition metal occupies the Li site, in a quantity

increasing with y

[129c]

. By anomalous X-Ray scattering and

57

Fe Mössbauer spectroscopy, it

was shown that it is mostly Fe occupying the Li sites, so that effectively the substitution worsens LNO´s off-stoichiometry [125, 129c]. Indeed, Fe3+ is in high-spin configuration and it is a large cation, almost matching the size of Ni2+, thus being well accommodated in the Li site. As a consequence, Fe doping significantly worsens the electrochemical performance of LNO [129d]. Still, about 190 mAh/g could be achieved by some authors in the first discharge for y = 0.15 [39d]

. Interestingly, upon Li deintercalation, Mössbauer spectroscopy revealed that Fe3+/Fe4+

oxidation, often hard to achieve, occurs concurrently to Ni oxidation. On the other hand, in Codoped LNO, cobalt is only oxidized after nickel, as observed by X-ray absorption measurements

[136]

. This was understood showing how the cation-oxygen environment

imposed by the structure can modify the crystal fields and the stability of a given oxidation state [20a]

.

8.3 Manganese Mn substitution has also been thoroughly investigated. A recent study of the Li-Ni-Mn-O phase diagram clearly evidenced its complexity

[130a]

. Concerning layered phases, a solid solution

LiNi1-yMnyO2 has been isolated for 0 ≤ y < 0.5 [123]. Similar to the Fe case, the limited solubility is due to the fact that layered LiMnO2 is only metastable [130b], while the thermodynamic ground state structure is orthorhombic

[130c]

. Moreover, one should remember that Li2MnO3

(Li[Li1/3Mn2/3]O2) exists as a stable layered phase

[137]

, so that Li excess may result in the

formation of this latter compound. In fact, Mn and Ni share the Ni site, but Rossen et al.

[123]

showed that increasing y results in a cation mixing of Li in the Mn/Ni site, and also Mn/Ni in the Li site. Moreover, increasing y resulted in a higher average Li/metal ratio, evidence of the formation of an increasing amount of Mn4+. In general, it was shown that the amount of Li that can be reversibly cycled is reduced by Mn substitution, so that LNO´s electrochemistry is 43

negatively affected by it

[123, 130d]

. However, by limiting the Mn content to 10%, Arai et al.

[130e]

could achieve good specific capacity of 190 mAh/g in the 1st discharge, if using significant Li excess; otherwise only 159 mAh/g were obtained. In any case, the samples also exhibited quite high polarization and capacity loss after 10 cycles (worse than the Co-doped counterpart). The improved thermal stability of the Mn-doped compounds was also demonstrated [130e, 130f].

8.4 Titanium Substitution of Ni with Ti has been studied for 0 < y < 0.5, but only y ≤ 0.2 leads to samples with a rhombohedral symmetry

[130e, 131a]

. For higher Ti content, disordering of the cations

towards the rock salt structure occurs. In the hexagonal phases, titanium is incorporated as Ti4+ [131b], thus the average oxidation state of nickel is reduced. As a consequence, higher annealing temperatures can be used. On the other hand, a higher fraction of Ni2+ is left in the sample for charge compensation, likely occupying the Li site. Indeed, despite opposite claims[131b], Croguennec et al. showed that the Li off-stoichiometry increases with y and that it is Ni which occupies the Li site, while Ti is only placed in the slab

[33b]

. This is also confirmed

by the fact that Ti4+ ions, larger than Ni3+, increase both the a and c unit cell parameters, but the c/a ratio decreases. As a result, the amount of titanium should be minimized. Kim et al. [131b]

demonstrated promising electrochemical performance for samples with y = 0.025: initial

discharge capacity of 235 mAh/g and a rather low capacity fading over 100 cycles. With increasing y, the specific capacity decreases, while its retention improves

[131b-d]

. Analogously

to off-stoichiometric LNO samples, the voltage profile becomes smoother as the amount of Ti increases because of the presence of increasing Ni2+ in the interslab space, destabilizing Livacancy orderings. For the same reason, the cell polarization increases with y [33b]. The thermal stability of Ti-doped LNO is also enhanced [33a].

44

8.5 Aluminum Aluminum can substitute Ni in LNO, as first shown by Ohzuku et al. prepared by-precipitation

[110a]

Typical samples are

[66b]

. Being redox-inactive, clearly the substitution results in a lower

theoretical specific capacity, somewhat mitigated by the low molar mass of Al. Only y < 0.5 can be obtained, as for y = 0.5 impurities of other Li-Al-O phases appear. Due to the smaller size of Al3+ (0.535 Å) with respect to Ni3+, the unit cell volume contracts upon doping. In fact, the size difference is sufficient to make a homogeneous solid solution hard to achieve, as commonly observed via a broadening of (11l) Bragg reflections

[110a]

. Aluminum segregation

within the slabs was confirmed by a combination of synchrotron XRD, ED, EDX and EELS within single particles at the nm scale, but also among different particles

[132a]

. In this respect,

the claims of homogenous Al doping and that the presence of Al accelerates the formation of the LNO layered structure are surprising [132b, 132c]. Structural studies revealed that Al does not significantly improve the layered character of LNO, and a significant fraction of Ni2+ (≈ 5%), but no Al, is always found in the interlayer

[110b]

. In fact, while the a lattice parameter decreases,

the interlayer distance increases with the Al content. The 1st cycle specific capacity and efficiency achieved for Al-doped LNO are decreased with respect to non-doped samples, while the capacity retention seems to be only slightly improved [66b, 132c-f]

. On the other hand, the average voltage is clearly increased [132g] due to the increased

covalence of the M-O bonds

[132h]

, as is the cell polarization because Al hinders the charge

redistribution within the slab. The incorporation of 10% Al has been reported to suppress the phase transitions that LNO suffers during Li (de)intercalation

[66b, 110a]

. Moreover, the strain

induced by the Al3+/Ni3+ size difference is relieved as Ni is oxidized. Overall, a decrease in particle cracking and improved cycling performance can be achieved. The main advantage of Al doping is the improved thermal stability, as several authors reported by DSC a smaller amount of heat evolution, with higher onset temperature

[110]

. It can be rationalized by the fact

that the decomposition occurs via an intermediate pseudo-spinel structure, and cationic

45

migration occurs via tetrahedral sites. The high stability of Al in these sites delays the decomposition, thus stabilizing the layered structure [110b].

8.6 Magnesium The doping of LNO with Mg2+ is peculiar. Due to steric effects, Mg is preferentially placed in the larger interslab space

[55, 133a]

. For low doping amounts y < 0.1, only Mg is found on the Li

site, so no Ni off-stoichiometry is present (but the presence of Mg2+ still results in charge compensating Ni2+ in the slab, according to [Li1-yMgy][NiIII1-yNiIIy]O2). On the other hand, for y > 0.1 some Mg is also occupying the Ni layer, with creation of compensating Ni4+. The inactive electrochemical character of Mg2+ as substituting cation and its size similarity to Li+ contribute to significantly improve the cycling stability of Mg-doped LNO. Mg2+ ions behave as pillars in the Li layer, preventing the structure from collapsing during deintercalation. 5% doping is sufficient to yield a complete solid solution behavior [55, 133b]. It was observed that upon charge all Mg2+ ions migrate to the interslab, which is responsible for a small capacity loss, but does not decrease the stabilization effect on the structure. As for discharge capacity, it is reduced with increasing doping, so an optimal Mg content is 5% or below

[55, 132e, 133b-d]

. Finally, Mg2+-

Ti4+ co-doping was also reported to improve both the reversible specific capacity and thermal stability [88, 133e].

8.7 Other dopants Other elements proposed as possible cation dopants are Cu 134]

,Y

[135c]

, Sb

[135d, 135e]

, Zn

[135f]

, In and Tl

[39c]

, Ca and Nb

[135a]

[135g]

,B

[132e]

, Ce

[135b]

, Ga

[39c, 39e,

. Special emphasis should be

placed on a very recent work devoted to doping with low amounts of W

[66c]

and Zr [135h, 135i].

46

Although tungsten may seem a surprising dopant, Kim et al.

[66c]

clearly demonstrated its

beneficial effect when used in small fractions (0.5-1 mol%). W fosters the rock salt phase formation during synthesis, but more importantly it promotes the segregation of this phase at the particle surface, thus effectively stabilizing it. By first principles calculations, it was shown that, although W doping barely affects the bulk energy of LNO, it decreases the energy of the (110) rock salt surface. W-LNO has only slightly lower initial specific discharge capacity than LNO (245 mAh/g in half cells) and does not significantly change the shape of the dQ/dV curve, yet remarkably, it shows 210 mAh/g after 100 cycles (90% capacity retention as opposed to 74% for LNO) and 145 mAh/g after 1000 cycles (in full cells, Figure 12) [66c]. The stable surface also induces a higher thermal stability, and no particles degradation is observed after cycling. Importantly, the same strategy also improves NCM materials. One should note that the same

Figure 12: (a) HRTEM image of 1 mol% W-LNO with surface rock salt 𝐹𝑚3̅𝑚 and bulk 𝑅3̅𝑚 phases. The insets show Fourier filtered images of each phase at an atomic resolution. (b) ASTAR TEM phase mapping and the corresponding bright field images for LNO, 0.5 mol% and 1 mol% W-LNO, showing the distribution of the rock salt phase. (c) Long-term cycling performance of 1 mol% W-doped Li[NixCoyMn1-x-y]O2 with x = 0.8, 0.89, 0.9 and 1.0; y = 0.15, 0.11, 0.05, and 0 and Li[Ni0.82Co0.14Al0.04]O2 (NCA82, commercial Ni-rich benchmark cathode) for comparison, tested using graphite-based pouch cells at 1 C rate [3.0-4.2 V]. Reproduced from Ref.[66c] with permission from The Royal Society of Chemistry. 47

authors analyzed the effect of Ti, Zr and Mo doping, and found W was the most beneficial dopant for LNO. Zr substitution modifies LNO in a surprising way: 0.4 mol% Zr was shown to promote Ni/Li exchange, normally not observed in LNO (about 2% of anti-site defects, with Li-Ni ordering seen by TEM) [135h]. The Ni/Li exchange does not have a negative effect on the electrochemical performance, since Zr-doped LNO has the same discharge capacity and rate capability of LNO, and even delayed capacity loss upon cycling (81% retention after 100 cycles from the initial value of 246.5 mAh/g). Hence, Li/Ni exchange (LNO closer to a rock salt structure, but without off-stoichiometry) is not necessarily detrimental to LNO´s performance. Moreover, a higher Zr fraction (1.4 mol%) results in simultaneous Zr doping and Zr segregation at the surface of the particles, forming a protective Li2ZrO3-like coating that stabilizes LNO (86% retention after 100 cycles from the initial 233 mAh/g)

[135i]

. In particular, the peak in the dQ/dV curve associated

with the H2-H3 transition is lost upon long-term cycling of LNO, while this is not the case in Zrdoped samples, where the deleterious phase transition is delayed, the charge transfer resistance remains low and no secondary particle fracture is observed after cycling.

Figure 13: Common dopant cations (D) in LNO as a function of their ionic radius and valence. Ni and Li are represented by circles, dopants by diamonds. Increasing trends are denoted by arrows.

Figure 13 summarizes the effect of different cations doping on LNO. Clearly, not all the complexity of each doping can be captured in a single scheme. However, some trends appear 48

and can be recognized: (i) LNO has little tolerance for small cations that can only be placed on the Ni site, and a tendency to inhomogeneous doping (clustering) appears even at small doping fractions. Large cations can be easily accommodated, because of the availability of both Li and Ni sites. (ii) Cations of increasing size and valence ≤ 3 results in a tendency for the dopant to occupy the Li site. (iii) Among tested substituents, only Mn combines a smaller ion size and a larger valence, making it unique. Mn tends to cluster and to cause Li migration to the Ni site, hence forming stable Li2MnO3. (iv) Cations larger than Ni3+ and possessing higher valence create charge-compensating Ni2+ and combine the tendency to place Li in the Ni site and dopant/Ni2+ on the Li site, which ultimately results in formation of the rock salt structure. This is also confirmed by cation doping with low solubility limit (W, Zr), forming rock salt phases at the particle surface. Anion substitution has also been reported in the literature. In this case, oxygen is partially replaced by anions such as halides (F, Cl) or sulfur. F was investigated by a few authors, although in early work it is unclear whether the substitution was actually successful, also because fluorination was attempted in parallel to Li excess [138]. F substitution was later shown to be possible up to LiNiO1.8F0.2. However, upon doping, more Ni2+ is available to occupy the Li interslab, thus increasing the samples off-stoichiometry

[139]

. The c lattice parameter

decreases with F content, while a increases. The (003) reflection also weakens, confirming the worse layering of the material. As expected, the above results in lower reversible specific capacity, suppressed phase transitions with cycling and capacity retention is not significantly improved, if at all [138a, 140]. Sulfur-doped samples have also been synthesized up to LiNiO1.7S0.3. S was shown to ameliorate the capacity retention, but affecting the initial capacity [141]. In good agreement with these results, a systematic DFT study on anion doping shows that the most relevant anion dopants play a conflicting role on the properties of LNO, improving some properties while making others worse [142].

9. Future prospects 49

Until recently, substitution has been the sole strategy thoroughly developed. A proper “stabilization” has not been achieved by a single dopant without sacrificing a significant amount of energy density (with possibly the exception of the recent reports on W and Zr doping). A different strategy, currently under development, consists in modifying the surface of LNO instead of the bulk, since most of the chemical, thermal and mechanical degradation is especially affecting the material´s surface. Three main strategies can be distinguished, namely coating LNO, synthesizing core-shell/concentration gradient materials and controlling the disordered rock salt structure at the surface. The key challenge in each of these methods is to minimize the amount of extra processing required, its cost and the loss of specific capacity, while maximizing the beneficial stabilizing effect.

9.1 Coating A coating is a thin layer of a stable compound that passivates the surface of LNO, thereby reducing the amount of reactions occurring such as O2 release and the formation of a cSEI. Different metal oxides can be used for this purpose (MgO, Al2O3, SiO2, TiO2, ZnO, SnO2, ZrO2, Li2O·2B2O3-glass etc.), and several reports exist on coatings of various NCM materials [143]. As for coating of LNO, remarkable results were obtained by Cho et al. using ZrO2. The characterization results suggest that a thin layer of LiNi1-yZryO2 is formed at the surface of the cathode, in addition to the zirconia layer. The typical phase transitions of LNO were suppressed, resulting in 98% retention of the initial specific capacity (190 mAh/g) after 70 cycles [144]. Silica and lanthanum oxides have also been tested as possible coatings [145]. While the former did not show promising results, a 2 wt% coating with La2O3 yielded a 94% retention of the initial specific capacity (198 mAh/g) after 60 cycles, again due to suppression of the phase transitions of LNO. More recently, a first principles study proposed a monolayer of alumina as a possible candidate coating to suppress the O2 release and improve the thermal stability of LNO [146].

50

Figure 14: (a) Initial (dis)charge curves of LNO and core-shell LNO (with high-Ni NCM shell, LiNi0.8Co0.1Mn0.1O2) cathodes at 0.1 C and (b) their cycling performance at 0.5 C. (c) Dark and (d) bright field STEM images of a core-shell LNO particle. Adapted with permission from [147] Copyright 2017, American Chemical Society.

9.2 Core-shell and concentration gradient strategies By core-shell (CS) particle in this context we mean a LNO core encapsulated in a doped-LNO shell, such as a NCM material. In this case, a LNO (or high-Ni NCM) core provides the high capacity, and it is then coated by a Ni-poor and especially Mn-rich NCM shell, providing thermal and mechanical stabilization. Moreover, if the transition between core and shell is not abrupt but continuous, the result is a particle with a Ni concentration gradient. The key point is that the shell is created during the same synthesis procedure of the core, by tuning the coprecipitation parameters and the precursors´ solution delivery to the reactor. Thus, a thorough control of the whole co-precipitation process is paramount to the success of this method. Recent examples include the synthesis of a LiNi0.8Co0.1Mn0.1O2 core encapsulated in a LiNi0.5Mn0.5O2 shell, or a LiNi0.86Co0.10Mn0.04O2 core - LiNi0.70Co0.10Mn0.20O2 shell. The authors demonstrated in both cases initial capacities in excess of 200 mAh/g and good cycling and thermal stability, entirely dominated by the stability of the shell material [148]. A pure LiNiO2 core 51

has been encapsulated in protective shells of LiCoO2, LiCo1-yMnyO2 [149] or LiNi0.8Co0.1Mn0.1O2 (Figure 14) [147]. The resulting compound is for 75% of the volume LNO, resulting in an average composition of LiNi0.95Co0.025Mn0.025O2. With this strategy, retention of 90% of the initial capacity (236 mAh/g) after 100 cycles was reported. Although this likely also arises from the favorable morphology of the shell, composed of radially aligned rod-shaped primary particles, it is clearly shown that less structural degradation occurs in CS LNO as opposed to the pristine material.

9.3 Related compounds: disordered rock salts In conclusion of this prospective section, we think it is important to mention the role of disordered rock salt compounds, as they are structurally related to LNO and LNO´s surface degrades towards a rock salt structure. Although disordered rock salt compounds had been disregarded as possible cathodes, recent literature intensively reevaluated them [150]. They are relevant for LNO because certain materials possessing similar stoichiometry of doped LNO samples (but different structure), such as LiNi0.5Ti0.5O2 have been reported

[151]

. Recent

achievements have clearly shown that (i) a disordered rock salt material can effectively conduct lithium if a threshold Li concentration is exceeded (about 10% excess with respect to the metal(s)), creating a percolating Li network increased by cation disordering

[152]

[118]

; (ii) the voltage of compounds based on Ni is

and (iii) the fact that a given lithium transition-metal oxide

may be stable as a cation-disordered phase, even in the presence of cations with large size difference, is strongly related to the electronic structure of the involved elements. All disordered rock salt cathodes reported so far indeed contain elements such as Mo6+, Cr6+, V5+, Nb5+, Ti4+, or

Zr4+ not only as charge compensators, but also because a cation with d0 electronic

configuration is needed to stabilize the structure

[153]

. Transition metals with d0 configuration

are least sensitive with respect to local site distortions and can tolerate them at very low energy cost. In this context, we believe that the recent achievements on W 6+-doped LNO [66c] and Zr4+doped LNO [135h] are not coincidental, but may be due to close relationship between the layered

52

and disordered rock salt structures, and thus the two fields of research converge in the search of strategies to stabilize the surface of LNO and Ni-rich layered oxide materials.

10. Conclusion Herein, we reviewed almost three decades of work on LiNiO2. This compound revealed several challenging aspects, born from the electronic structure of trivalent nickel and its particular position in the Li-Ni-O phase diagram. We clarified the well-established properties of LNO and put special emphasis on those that are to date not fully understood. Among these, we highlight the experimental verification of the ground state structure of stoichiometric LNO and of the missing Li-vacancy orderings in the monoclinic region. Furthermore, the exact structure and stoichiometry of the rock salt NiO-like phases forming at the surface of LNO during electrochemical cycling are of utmost importance to understand its structural instability and should be object of further investigation. Likewise, the recent substitution/doping strategies that affect the formation of a disordered rock salt phase at the surface of LNO are particularly promising, as they exploit different mechanisms than “standard” bulk substitution, which has been already thoroughly investigated and resulted in the well-established NCA and NCM materials. We believe that these strategies may finally lead to

the commercialization of

the ultimate Ni-richest layered oxide cathode material.

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