Anode materials for Li / Na ion batteries

Abstract

The present invention relates to active electrode materials and to methods for manufacturing active electrode materials. Such materials are of interest as active electrode materials in lithium-ion or sodium-ion batteries. The present invention provides an active electrode material represented by the general formula [M][Nb] y [O] z wherein the active electrode material is oxygen-deficient; wherein M consists of one of Mg, Cr, W, Mo, Cu, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Sb, Y, La, Hf, Ta, Zn, In or Cd; y satisfies 0.5 ≤ y ≤ 49; and z satisfies 4 ≤ z ≤ 124.

Description

Technical Field

[0001] The present invention relates to electrode active materials and to methods for manufacturing electrode active materials. Such materials are of interest, for example, as electrode active materials in lithium-ion or sodium-ion batteries. Background Art

[0002] Lithium-ion (Li-ion) batteries are a commonly used type of rechargeable battery, with a global market estimated at $40 billion in 2018 and projected to grow to $200 billion by 2030. This large market is divided into various applications, ranging from transportation and utility-scale energy storage to consumer electronics and power tools. Therefore, rechargeable (secondary) Li-ion batteries are currently under intensive research and development to improve their performance to meet industrial technology requirements [Goodenough and Park (2013)]. In particular, Li-ion batteries are the technology of choice for electric vehicles, which have a variety of requirements in terms of technical performance to environmental impact, providing a viable path for the green automotive industry.

[0003] A typical lithium-ion battery consists of multiple cells connected in series or in parallel. Each individual cell generally consists of an anode (negative-polarity electrode) and a cathode (positive-polarity electrode), which are separated by a porous electrical insulating membrane (referred to as a separator) and immersed in a liquid (referred to as an electrolyte) capable of transporting lithium ions.

[0004] In most systems, the electrodes are composed of an electrochemically active material (which means it can chemically react with lithium ions to reversibly store and release them in a controllable manner), and, if necessary, mixed with a conductive additive (such as carbon) and a polymer binder. A slurry of these components is coated as a thin film on a current collector (usually a thin foil of copper or aluminum), thus forming an electrode after drying.

[0005] In known Li-ion battery technologies, the low rate performance of graphite anodes during battery charging severely hinders their application in high-power electronics, automobiles, and industry. Among the various potential alternatives recently proposed, Si, Si alloys, lithium titanate (LTO), and niobium oxide-based materials are the main competitors to replace graphite as the preferred active material for high-power applications.

[0006] The battery charging rate is usually expressed as the "C rate". A 1C charging rate means the charging current at which the battery is fully charged in 1 h, and a 10C charge means the battery is fully charged in 1 / 10 hour (6 minutes).

[0007] Batteries relying on graphite anodes are fundamentally limited in terms of charging rate. Under nominal conditions, lithium ions are inserted into the anode active material during charging. When the charging rate increases, the typical graphite voltage profile results in a high risk that the overpotential causes the potential of the anode to become <0V relative to Li / Li+, which leads to a phenomenon called lithium dendrite plating, whereby lithium ions are deposited as lithium metal on the surface of the graphite electrode. This results in an irreversible loss of active lithium and thus a rapid decay of the battery capacity. In some cases, such dendritic deposits can grow to such large sizes that they pierce the battery separator and cause a battery short circuit. This triggers a catastrophic failure of the battery, leading to fire or explosion. Therefore, the fastest-charging batteries with graphite anodes are limited to a charging rate of 5-7C, but typically lower. Nevertheless, graphite anodes accounted for >90% of the Li-ion battery market in 2018.

[0008] Si and Si alloys offer a large specific capacity, but have short lifetimes when charged and discharged at high rates, and have poor capacity retention when increasing from a low rate (e.g., 0.5C) to a higher rate (e.g., 5C). This is due to the non-uniform lithiation of the active material particles caused by the limited diffusion rate of lithium ions within the particles during charging. The core of the active material particles (usually 1-20μm spheres) may not have time to be lithiated during fast charging because lithium ions need to diffuse from the particle surface to the core, thus resulting in poor capacity retention when increasing the charging rate. In addition, Si and Si alloy active materials physically expand up to 400 volume % during lithiation. Therefore, non-uniform particle lithiation results in internal mechanical stress within the particles, which can cause them to break and electrode pulverization, and thus these electrodes have a short cycle life during fast charging.

[0009] Lithium titanate (LTO) anodes do not undergo dendrite plating at high charging rates due to their high potential (1.6V relative to Li / Li+), and have excellent cycle life due to their lack of volume expansion during lithiation. For these two reasons, LTO batteries are generally considered to be high-safety batteries. However, LTO is a relatively poor electronic and ionic conductor, which results in limited capacity retention at high rates unless the material is nano-sized to increase the specific surface area and is coated with carbon to increase the electronic conductivity. This kind of particle-level material engineering increases the material particle cost and reduces the packing density of the active material LTO powder. This is important because it results in low-density electrodes and a higher fraction of electrochemically inactive materials (such as binders, carbon additives).

[0010] A key measure of anode performance is its volumetric capacity (mAh / cm 3), i.e., the electric charge (i.e., lithium ions) that can be stored per unit anode volume. This is an important factor in determining the total battery energy density (Wh / L) based on volume. The volumetric capacity can be approximated as the product of the electrode density, the specific capacity of the active material, and the fraction of active material in the electrode. LTO anodes typically have a relatively low specific capacity (which is c.170 mAh / g, compared to c.330 mAh / g for graphite), and combined with the low electrode density (usually 1.9 g / cm 3 ) and low fraction of active material (<87%) discussed above results in a very low volumetric capacity (<300 mAh / cm 3 ), and thus results in a low battery energy density and a high $ / kWh cost. Therefore, LTO batteries / cells are typically limited to specific niche applications, despite their long cycle life, fast charging ability, and high safety.

[0011] Mixed niobium oxides (MNOs) were first identified as potential battery materials in the academic literature in the 1980s [Cava et al. (1983); Cava et al. (1984)], but at that time there was limited interest due to the lack of commercially available cathodes to match their rate capabilities.

[0012] Interest in MNO anodes revived in the early 2010s, demonstrating that practical batteries combining TiNb2O7 anodes and commercially available LNMO cathodes showed promising performance in terms of rate capability, cycle life, and energy density [Goodenough and Park (2013)]. Selected MNO anodes such as TiNb2O7 offer characteristics similar to LTOs in terms of a high working potential (1.6 V) relative to Li / Li+ and low volume expansion (<5%), resulting in safe fast charging and long cycle life (>10,000 cycles). A key advantage of MNO anodes is that a significantly higher specific capacity can be achieved compared to LTOs (e.g., c.300 mAh / g for TiNb2O7), thus increasing the battery energy density. However, in MNO materials such as TiNb2O7, the electronic conductivity is usually too low to sustain fast charging rates without particle engineering and carbon coating, which is similar to the limitation of LTOs.

[0013] On the other hand, it has recently been shown that other MNOs for batteries such as Nb2O5, which were also studied in the 1980s, typically have a so-called "Wadsley-Roth" or "bronze" crystal structure that can provide extremely fast lithium ion diffusion rates of 10 -14 -10 -10 cm 2 s -1 (LTO is typically 10 -17 cm 2s -1 ) [Griffith et al. (2016)]. This can potentially improve the electrode density (i.e., > 2.5 g / cm 3 ) and thus improve the volume capacity (> 600 mAh / cm 3 ). However, some challenges limit the commercial application of such materials, such as low electronic conductivity, lifespan issues as "micro-materials" (crystals with a size of about 1 - 10 μm), and "tilted" lithiation voltage curves. The lithiation voltage curve refers to the shape of the anode potential versus the amount of lithium inserted into the anode. LTO and TiNb2O7 have "flat" voltage curves, while materials such as Nb2O5 typically have "tilted" voltage curves. An overly tilted voltage curve results in a large voltage window, which makes full cell balancing challenging in commercial batteries.

[0014] TiNb2O7 has further limitations especially for its application in high-power battery technology. Compared with other MNOs (e.g., Nb 12 MoO 33 = 4.0 x 10 -14 cm 2 s -1 ), it has a limited lithium-ion diffusion rate (8.0 x 10 -16 cm 2 S -1 ) [Zhu 2019], which limits its performance at high power. In particular, this limits the use of the pseudocapacitive charge storage mechanism, which is a key benefit for using MNOs in high power [Yang 2017].

[0015] US9515319B2 discusses TiNb2O7 and considers but does not exemplify the modification of this material. However, the raw materials and processes used in this disclosure are expensive (furnace treatment up to 1500 °C for up to 50 h), and the materials produced show low initial Coulombic efficiency (84.7%, 86.5%). US2015 / 0270543A1 and KR20150131800A disclose the modification of TiNb2O7.

[0016] US2019 / 0288283A1 discloses a lithium niobium composite oxide, in which as a basic feature, some niobium must be replaced by at least one element selected from Fe, Mg, Al, Cu, Mn, Co, Ni, Zn, Sn, Ti, Ta, V, and Mo.

[0017] The present invention is designed in view of the above considerations. Summary of the Invention

[0018] The present inventor has recognized that, despite the obvious challenges in the prior art, it is possible to provide active electrode materials that overcome some or all of the problems presented by the above-described prior art materials.

[0019] Accordingly, in a first aspect, the present invention provides an active electrode material represented by the general formula [M1] x [M2] (1-x) [Nb] y [O] z wherein:

[0020] M1 and M2 are different;

[0021] M1 represents one or more of Ti, Mg, V, Cr, W, Zr, Mo, Cu, Fe, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Te, Se, Si, Sb, Y, La, Hf, Ta, Re, Zn, In or Cd;

[0022] M2 represents one or more of Mg, V, Cr, W, Zr, Mo, Cu, Ga, Ge, Ca, K, Ni, C o , A1, Sn, Mn, Ce, Sb, Y, La, Hf, Ta, Zn, In or Cd; and wherein

[0023] x satisfies 0 < x < 0.5; [[ID=**27]]

[0024] y satisfies 0.5 ≤ y ≤ 49;

[0025] z satisfies 4 ≤ z ≤ 124.

[0026] Materials in which M1 and M2 are different may also be referred to as mixed cation active materials or composite oxide active materials. These terms are used interchangeably in the present disclosure and refer to materials having the general formula set forth above. Compared to non-mixed cation active materials (e.g., materials having the general formula [M] x [Nb] y [O] z where M represents a single ion), such materials may provide improved electrochemical properties.

[0027] In particular, as shown in the examples, the inventors have found that by substituting non-Nb cations to form a mixed cation structure, the entropy in the crystal structure can be increased, thereby reducing the potential energy barrier for lithium ion diffusion by introducing a small amount of defects. The modification carried out by generating a mixed cation structure that maintains the same total oxidation state as the unmodified crystal structure demonstrates the potential improvement by changing the ionic radius, which can cause minor changes in the crystal parameters and Li ion holes, thereby improving the electrochemical properties. For example, by substituting cations with a larger ionic radius, the unit cell can expand relative to the unmodified structure, which can result in a higher lithium ion diffusion rate. The modification carried out by generating a mixed cation structure that results in an increased oxidation state demonstrates potential advantages similar to those of the changed ionic radius, by introducing additional electron holes in the structure to contribute to conductivity and recombination. The modification carried out by generating a mixed cation structure that results in a decreased oxidation state demonstrates potential advantages similar to those of the changed ionic radius, by introducing oxygen vacancies and additional electrons in the structure to contribute to conductivity and recombination. The modification carried out by inducing oxygen deficiency by high-temperature treatment under inert or reducing conditions provides a reduced structure with a greatly increased conductivity. The combination of the mixed cation structure and the induced oxygen deficiency allows for a variety of beneficial effects.

[0028] As described above, M1 represents one or more of Ti, Mg, V, Cr, W, Zr, Mo, Cu, Fe, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Te, Se, Si, Sb, Y, La, Hf, Ta, Re, Zn, In or Cd. M2 represents one or more of Mg, V, Cr, W, Zr, Mo, Cu, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Sb, Y, La, Hf, Ta, Zn, In or Cd. 'Represents one or more' means that M1 or M2 can each represent two or more elements from their respective lists. An example of such a material is Ti 0.05 W 0.25 Mo 0.70 Nb 12 O 33 . Here, M1 represents Ti x’ W x” (where x'+x" = x), M2 represents Mo, x = 0.3, y = 12, z = 33. Another example of such a material is Ti 0.05 Zr 0.05 W 0.25 Mo 0.65 Nb 12 O 33 . Here, M1 represents Ti x’ Zr x” W x”’(where x'+x''+x''' = x), M2 represents Mo, x = 0.35, y = 12, z = 33.

[0029] Optionally, M1 represents one or more of K, Mg, Ca, Y, Ti, Zr, Hf, V, Ta, Cr, Mo, W, Mn, Fe, Co, Ni, Cu, Zn, Al, Ga, Si, Ge, Sn, Sb. M1 can represent one or more of Ti, Mg, V, Cr, W, Zr, Mo, Cu, Ga, Ge, K, Ni, Al, Hf, Ta or Zn. Preferably, M1 represents one or more of Ti, Mg, V, Cr, W, Zr, Mo, Ga, Ge, Al or Zn.

[0030] M2 does not represent Ti. In other words, preferably, Ti is not the main non-Nb cation in the active electrode material. When M1 represents Ti alone, preferably x is 0.05 or less. When M1 represents one or more cations including Ti, preferably, the amount of Ti is 0.05:1 or less relative to the total amount of non-Nb cations.

[0031] Optionally, M2 is selected from one or more of Mo, W, V, Zr, Al, Zn, Ga, Ge, Ta, Cr, Cu, K, Mg, Ni or Hf, and M2 can be selected from one or more of Mo, W, V, Zr, Al, Zn, Ga or Ge. Preferably, M2 is selected from one or more of Mo, W, V or Zr. The inventors have found that when M2 is selected from one of these elements, the active electrode material can have improved electrochemical properties. M2 can consist of a single element.

[0032] When x satisfies 0 < x < 0.5, M2 is the main non-Nb cation in the active electrode material. Preferably, x satisfies 0.01 ≤ x ≤ 0.4, more preferably, x satisfies 0.05 ≤ x ≤ 0.25. For example, x can be about 0.05.

[0033] The exact values of y and z within the defined ranges can be selected to provide a charge-balanced or substantially charge-balanced crystal structure. Additionally or alternatively, the exact values of y and z within the defined ranges can be selected to provide a thermodynamically stable or thermodynamically metastable crystal structure.

[0034] In some cases, z can be defined in the format z = (z’ - z’α), where α is a non-integer value less than 1, for example where α satisfies 0 ≤ α ≤ 0.05. α can be greater than 0, i.e., α can satisfy 0 < α ≤ 0.05. When α is greater than 0, the active electrode material is an oxygen-deficient material, i.e., the material has oxygen vacancies. Such a material will not have an exact charge balance but is considered to be “substantially charge balanced” as described above. Alternatively, α can be equal to 0, in which case the material is not an oxygen-deficient material.

[0035] When α is 0.05, the number of oxygen vacancies is equal to 5% of the total oxygen in the crystal structure. In some embodiments, α can be greater than 0.001 (0.1% oxygen vacancies), greater than 0.002 (0.2% oxygen vacancies), greater than 0.005 (0.5% oxygen vacancies), or greater than 0.01 (1% oxygen vacancies). In some embodiments, α can be less than 0.04 (4% oxygen vacancies), less than 0.03 (3% oxygen vacancies), less than 0.02 (2% oxygen vacancies), or less than 0.1 (1% oxygen vacancies). For example, α can satisfy 0.001 ≤ α ≤ 0.05. When the material is oxygen-deficient, the electrochemical properties of the material can be improved. For example, resistance measurements can show an increased conductivity compared to an equivalent non-oxygen-deficient material. It is understood that the percentage values represented here are atomic percentages.

[0036] The oxygen deficiency in the material (e.g., expressed as a percentage of oxygen vacancies) can be measured, for example, by thermogravimetric analysis (TGA) in an oxygen-rich atmosphere, by measuring the extent to which the mass of the sample changes over time due to oxygen being re-incorporated into the oxygen vacancies. Alternatively or additionally, the oxygen deficiency can be measured qualitatively by evaluating the color of the material relative to a non-oxygen-deficient sample of the same material. For example, non-oxygen-deficient MoNb 12 O 33 has a white, off-white, or yellow color. Oxygen-deficient MoNb 12 O ＜33 has a purple color. When preparing oxygen-deficient crystals of MoNb 12 O ＜33 a color change from white / off-white / yellow to purple can be observed.

[0037] M1 may have an oxidation state equal to or lower than that of M2. Preferably, M1 has an oxidation state lower than that of M2. When more than one element is present as M1 and / or M2, it should be understood that the oxidation state refers to M1 and / or M2 as a whole. For example, if 25 at% of M1 is Ti and 75 at% of M1 is W, the oxidation state of M1 is 0.25×4 (contribution of Ti) + 0.75×6 (contribution of W). Advantageously, when M1 has an oxidation state lower than that of M2, this is compensated by forming oxygen vacancies, i.e., forming an oxygen-deficient active electrode material. The presence of oxygen vacancies is believed to improve the conductivity of the active electrode material and provide further benefits, as demonstrated by the examples. Optionally, M1 comprises at least one cation with a 4+ oxidation state and M2 comprises at least one cation with a 6+ oxidation state. Optionally, M1 has an oxidation state of 4+ and M2 has an oxidation state of 6+. M1 preferably has a different ionic radius from M2, most preferably a larger ionic radius. This results in a change in the unit cell size and local deformation in the crystal structure. It is believed that this changes the Li ion site availability by changing the hole size and reducing the energy barrier for reversible lithiation, thereby improving electrochemical properties such as specific capacity and Coulombic efficiency. The ionic radius can be the Shannon ionic radius for the coordination and valence expected for the ion in the crystal structure of the active electrode material (available in R.D. Shannon, Acta Cryst., A32, 1976, 751-76).

[0038] The active electrode material can be a material selected from the group consisting of:

[0039] (i) M1 x Mo (1-x) Nb 12 O (33-33α)

[0040] M1 x W (1-x) Nb 12 O (33-33α)

[0041] M1 x V (1-x) Nb9O (25-25α)

[0042] M1 x Zr (1-x) Nb 24 O (62-62α)

[0043] M1 x W (1-x) Nb 0.57 O (4.43-4.43α)

[0044] M1 x W(1-x) Nb 0.89 O (5.22-5.22α)

[0045] M1 x Zn (1-x) Nb 17 O (43.5-43.5α)

[0046] M1 x Cu (1-x) Nb 17 O (43.5-43.5α)

[0047] M1 x Al (1-x) Nb 11 O (29-29α)

[0048] M1 x Ga (1-x) Nb 11 O (29-29α)

[0049] M1 x Ge (1-x) Nb 18 O (47-47α)

[0050] M1 x W (1-x) Nb 1.125 O (5.81-5.81α)

[0051] M1 x W (1-x) Nb 3.2 O (11-11α)

[0052] M1 x Al (1-x) Nb 49 O (124-124α)

[0053] M1 x Ga (1-x) Nb 49 O (124-124α) ; or

[0054] (ii)M1 x Mo (1-x) Nb 12 O (33-33α)

[0055] M1 x W (1-x) Nb 12 O (33-33α)

[0056] M1 x V (1-x) Nb9O (25-25α)

[0057] M1 x Zr (1-x) Nb 24 O (62-62α)

[0058] M1 x W (1-x) Nb 0.57 O (4.43-4.43α)

[0059] M1 x W (1-x) Nb 0.89 O (5.22-5.22α)

[0060] Ml x Zn (1-x) Nb 17 O (43.5-43.5α)

[0061] Ml x Al (1-x) Nb 11 O (29-29α)

[0062] Ml x Ge (1-x) Nb 18 O (47-47α) ; or preferably

[0063] (iii)M1 x Mo (1-x) Nb 12 O (33-33α)

[0064] M1 x W (1-x) Nb 12 O (33-33α)

[0065] M1 x V (1-x) Nb9O (25-25α)

[0066] M1 x Zr (1-x) Nb 24 O (62-62α)

[0067] M1 x W (1-x) Nb 0.57 O (4.43-4.43α)

[0068] M1 x W (1-x) Nb 0.89 O (5.22-5.22α)

[0069] Wherein M1 represents one or more of Ti, Mg, V, Cr, W, Zr, Mo, Cu, Fe, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Te, Se, Si, Sb, Y, La, Hf, Ta, Re, Zn, In or Cd; and wherein x satisfies 0 < x < 0.5; and α satisfies 0 ≤ α ≤ 0.05.

[0070] In a particularly preferred aspect, the active electrode material is M1 x Mo (1-x) Nb 12 O (33-33α) . In another particularly preferred aspect, the active electrode material is M1 x W (1-x) Nb 0.57 O (4.43-4.43α) . In another particularly preferred aspect, the active electrode material is M1 x Zn (1-x) Nb 17 O (43.5-43.5α) . In another particularly preferred aspect, the active electrode material is M1 x Al (1-x) Nb 11 O (29-29α) . Examples show that these materials have particularly advantageous properties for use as active electrode materials.

[0071] The materials in groups (i), (ii) and (iii) above and in the particularly preferred aspects represent specific non-mixed cation active materials (i.e., when x = 0), which have been modified into mixed cation active materials by replacing less than half of M2 with different elements M1. Optionally, in these materials, M2 can also be replaced by Nb at non-Nb sites in the crystal structure. That is, M1 can represent one or more of Ti, Mg, V, Cr, W, Zr, Mo, Cu, Fe, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Te, Se, Si, Sb, Y, La, Hf, Ta, Re, Zn, In, Cd or Nb. M1 can also represent the further list of elements described above and in the claims.

[0072] In a second aspect, the present invention provides an active electrode material having the general formula [M][Nb] y [O] zrepresented by; wherein the active electrode material is oxygen-deficient; wherein M consists of one of Mg, Cr, W, Mo, Cu, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Sb, Y, La, Hf, Ta, Zn, In or Cd; y satisfies 0.5 ≤ y ≤ 49; and z satisfies 4 ≤ z ≤ 124.

[0073] The material according to the second aspect of the present invention is an oxygen-deficient analogue of a known "base" material, such as MoNb 12 O 33 , WNb 12 O 33 , W7Nb4O 31 and W9Nb8O 47 . z can be defined as z = (z’ - z’α), where α satisfies 0 < α ≤ 0.05. When z is defined as z = (z’ - z’α), the above comments on the possible range of α stated for the material according to the first aspect also apply here to the material of the second aspect of the present invention. For example, α can satisfy 0.001 ≤ α ≤ 0.05. It has been found that the oxygen-deficient material according to the second aspect has improved properties compared to the stoichiometric 'base' material and is used as an active electrode material. For example, the material according to the second aspect has improved conductivity.

[0074] M can consist of one of Mo, W, Al, Zn, Ga, Ge, Ta, Cr, Cu, K, Mg, Ni or Hf; or M can consist of one of Mo, W, Al, Zn, Ga or Ge; or preferably, M can consist of one of Mo, W, Al or Zn.

[0075] The active electrode material of the second aspect can be represented by the general formula [M] x [Nb] y [O] (z’-z’α) and is selected from the group consisting of:

[0076] MoNb 12 O (33-33α)

[0077] WNb 12 O (33-33α)

[0078] W7Nb4O (31-31α)

[0079] W9Nb8O (47-47α)

[0080] Zn2Nb 34 O (87-87α)

[0081] Cu2Nb 34O (87-87α)

[0082] AlNb 11 O (29-29α)

[0083] GaNb 11 O (29-29α)

[0084] Nb 18 O (47-47α)

[0085] W 16 Nb 18 O (93-93α)

[0086] W5Nb 16 O (55-55α)

[0087] AlNb 49 O (124-124α)

[0088] GaNb 49 O (124-124α)

[0089] Where α satisfies 0<α≤0.05.

[0090] These represent mixed niobium oxides of a specific stoichiometry that have been modified to oxygen-deficient mixed niobium oxides.

[0091] The active electrode material of the second aspect can be represented by the general formula [M] x [Nb] y [O] (z’-z’α) means that it is selected from the group consisting of:

[0092] MONb 12 O (33-33α)

[0093] WNb 12 O (33-33α)

[0094] W7Nb4 (31-31α)

[0095] W9Nb8 (47-47α)

[0096] Zn2Nb 34 O (87-87α)

[0097] AlNb 11 O (29-29α)

[0098] Nb 18 O(47-47α)

[0099] where α satisfies 0 < α ≤ 0.05.

[0100] The active electrode material of the second aspect can be represented by the general formula [M] x [Nb] y [O] (z’-z’α) and is selected from the group consisting of:

[0101] MoNb 12 O (33-33α)

[0102] WNb 12 O (33-33α)

[0103] W7Nb4O (31-31α)

[0104] W9Nb8O (47-47α)

[0105] where α satisfies 0 < α ≤ 0.05.

[0106] Optionally, M is W. That is, the active electrode material can be represented by the general formula [W][Nb] y [O] z For example, the active electrode material can be selected from WNb 12 O (33-33α) , W7Nb4O (31-31α) , W9Nb8O (47-47α) , W 16 Nb 18 O (93-93α) and W5Nb 16 O (55-55α) . Examples show that inducing oxygen vacancies in many different tungsten niobium oxides leads to improved properties, such as increased conductivity, compared to stoichiometric base oxides.

[0107] In a particularly preferred aspect, the active electrode material is MoNb 12 O (33-33α) . In another particularly preferred aspect, the active electrode material is WNb 12 O (33-33α) . In another particularly preferred aspect, the active electrode material is W5Nb 16 O (55-55α) . In another particularly preferred aspect, the active electrode material is W7Nb4O (31-31α) . In another particularly preferred aspect, the active electrode material is Zn2Nb 34 O (87-87α) . In another particularly preferred aspect, the active electrode material is AlNb11 O (29-29α) Examples show that these materials have particularly advantageous properties for use as active electrode materials.

[0108] The present inventors have found that by doping various non-Nb cations to form mixed cation active materials / composite oxide active materials (according to the first aspect of the present invention) and / or by forming oxygen deficiencies (according to the second aspect of the present invention) to modify materials such as MoNb 12 O 33 、WNb 12 O 33 、ZrNb 24 O 62 、VNb9O 25 、W7Nb4O 31 and W9Nb8O 47 active electrode materials with improved electrochemical properties, and in particular improved electrochemical properties when used as anode materials, can be produced.

[0109] The specific capacity / reversible de-lithiation capacity of the active electrode material according to the present invention can be 200 mAh / g or greater, 225 mAh / g or greater, 250 mAh / g or greater, up to about 300 mAh / g or greater. Herein, the specific capacity is defined as the specific capacity measured in the second cycle of a half-cell galvanostatic cycling test at a rate of 0.05C, with a voltage window of 1.1 - 3.0V relative to Li / Li+. Providing a material with a high specific capacity can be advantageous as it can provide improved performance in an electrochemical device containing the active electrode material.

[0110] Furthermore, the active electrode material according to the present invention can have an appropriate voltage curve, as measured in the second cycle of a half-cell galvanostatic cycling test at a rate of 0.05C, with a voltage window of 1.1 - 3.0V relative to Li / Li+. That is, when the lithiation > 180 mAh / g, the capacity of the material can be between 2.0V and 1.1V, and when the de-lithiation > 180 mAh / g, the capacity of the material can be between 1.1V and 2.0V.

[0111] When formulated or coated as an electrode (optionally with a conductive carbon additive and a binder material), the bulk resistivity of the active electrode material according to the present invention can be 5 kΩ·cm or less, more preferably 2 kΩ·cm or less. The bulk resistivity can be a useful alternative measure of the conductivity of such a material. Providing a material with an appropriately low bulk resistivity can be advantageous as it can provide improved performance in an electrochemical device containing the active electrode material.

[0112] When measured in a Li-ion half button cell having said electrodes, the direct current internal resistance (DCIR) and the resulting area specific impedance (ASI) of the active electrode material can be 90 Ω or less (for DCIR) and 170 Ω·cm 2 or less (for ASI). Providing a material with suitably low DCIR and / or ASI can be advantageous as this can provide improved performance in an electrochemical device incorporating the active electrode material. However, further improvements in the DCIR / ASI values can be seen, for example, in carbon-coated active electrode materials, or where the active electrode material is incorporated in a commercial cell having a cathode, where the electrodes are calendared and prepared in a typically known manner. When measured in a button cell in such an arrangement, the inventors believe the ASI can be as low as, for example, 26 Ω·cm 2 or lower.

[0113] The active electrode material according to the present invention can have a lithium diffusion rate greater than 10 -14 cm 2 s -1 . Providing a material with a suitably high lithium diffusion rate can be advantageous as this can provide improved performance in an electrochemical device incorporating the active electrode material.

[0114] The active electrode material according to the present invention can have an electrode density of 2.5 g / cm 3 or higher after calendering. For example, the electrode density after calendering reaches 3.0 g / cm 3 or higher. Providing a material with such an electrode density can be advantageous as this can provide improved performance in an electrochemical device incorporating the active electrode material. Specifically, when the electrode density is high, a high volume capacity can be achieved since weight capacity × electrode density × active material fraction = volume capacity.

[0115] The initial Coulombic efficiency of the active electrode material according to the present invention can be greater than 88%, more preferably greater than 90%. In some cases, the initial Coulombic efficiency of the active electrode material can be as high as 92% or higher, 93% or higher, or 94% or higher. Providing a material with a suitably high initial Coulombic efficiency can be advantageous as this can provide improved performance in an electrochemical device incorporating the active electrode material. The initial Coulombic efficiency can be measured as the difference in the lithiation and delithiation capacities in the first charge / discharge cycle at C / 10 in a half cell.

[0116] Other optional features of the first and second aspects of the present invention are described below.

[0117] As determined by X-ray diffraction, the crystal structure of the active electrode material of the first aspect can correspond to the crystal structure of the unmodified form of the active electrode material, where the unmodified form is represented by the formula [M2][Nb] y[O] z It is represented that M2 consists of a single element and there is no oxygen vacancy in its unmodified form, and the unmodified form is selected from one or more of the following: M2 I Nb5O 13 、M2 I 6Nb 10.8 O 30 、M2 II Nb2O6、M2 II 2Nb 34 O 87 、M2 III Nb 11 O 29 、M2 III Nb 49 O 124 (M2 III 0.5 Nb 24.5 O 62 )、M2 IV Nb 24 O 62 、M2 IV Nb2O7、M2 IV 2Nb 10 O 29 、M2 IV 2Nb 14 O 39 、M2 IV Nb 14 O 37 、M2 IV Nb6O 17 、M2 IV Nb 18 O 47 、M2 V Nb9O 25 、M2 V 4Nb 18 O 55 、M2 V 3Nb 17 O 50 、M2 VI Nb 12 O 33 、M2 VI 4Nb 26 O 77 、M2 VI 3Nb 14 O 44 、M2 VI 5Nb 16 O 55 、M2 VI 8Nb 18 O 69 、M2VI Nb2O8, M2 VI 16 Nb 18 O 93 , M2 VI 20 Nb 22 O 115 , M2 VI 9Nb8O 47 , M2 VI 82 Nb 54 O 381 , M2 VI 31 Nb 20 O 143 , M2 VI 7Nb4O 31 , M2 VI 15 Nb2O 50 , M2 VI 3Nb2O 14 and M2 VI 11 Nb 12 O 63 , where the numbers I, II, III, IV, V, and VI represent the oxidation states of M2. In this way, it can be confirmed that the unmodified form has been modified without significantly affecting the crystal structure.

[0118] As determined by X-ray diffraction, the crystal structure of the active electrode material of the second aspect can correspond to the crystal structure of the unmodified form of the active electrode material, where the unmodified form is represented by the general formula [M][Nb] y [O] z indicating that the unmodified form has no oxygen vacancies, and the unmodified form is selected from M2 I Nb5O 13 , M2 I 6Nb 10.8 0 30 , M2 II Nb2O6, M2 II 2Nb 34 O 87 , M2 III Nb 11 O 29 , M2 III Nb 49 O 124 , M2 IV Nb 24 O 62 , M2 IV Nb2O7, M2 IV 2Nb 10 O29 , M2 IV 2Nb 14 O 39 , M2 IV Nb 14 O 37 , M2 IV Nb6O 17 , M2 IV Nb 18 O 47 , M2 V Nb9O 25 , M2 V 4Nb 18 O 55 , M2 V 3Nb 17 O 50 , M2 VI Nb 12 O 33 , M2 VI 4Nb 26 O 77 , M2 VI 3Nb 14 O 44 , M2 VI 5Nb 16 O 55 , M2 VI 8Nb 18 O 69 , M2 VI Nb2O8, M2 VI 16 Nb 18 O 93 , M2 VI 20 Nb 22 O 115 , M2 VI 9Nb8O 47 , M2 VI 82 Nb 54 O 381 , M2 VI 31 Nb 20 O 143 , M2 VI 7Nb4O 31 , M2 VI 15 Nb2O 50 , M2 VI 3Nb2O 14 and M2 VI 11 Nb 12 O 63, where the numbers I, II, III, IV, V, and VI represent the oxidation states of M. In this way, it can be confirmed that the unmodified form has been modified without significantly affecting the crystal structure.

[0119] As determined by X-ray diffraction analysis, the crystal structure of the active electrode material can correspond to the crystal structure of one or more of the following:

[0120] (i) MoNb 12 O 33

[0121] WNb 12 O 33

[0122] VNb9O 25

[0123] ZrNb 24 O 62

[0124] W7Nb4O 31

[0125] W9Nb8O 47

[0126] Zn2Nb 34 O 87

[0127] Cu2Nb 34 O 87

[0128] AlNb 11 O 29

[0129] GaNb 11 O 29

[0130] GeNb 18 O 47

[0131] W 16 Nb 18 O 93

[0132] W5Nb 16 O 55

[0133] AlNb 49 O 124

[0134] GaNb 49 O 124 ; or

[0135] (ii) MoNb12 O 33

[0136] WNb 12 O 33

[0137] VNb9O 25

[0138] ZrNb 24 O 62

[0139] W4Nb7O 31

[0140] W9Nb8O 47

[0141] Zn2Nb 34 O 87

[0142] AlNb 11 O 29

[0143] GeNb 18 O 47 ; or preferably

[0144] (iii)MoNb 12 O 33

[0145] WNb 12 O 33

[0146] ZrNb 24 O 62

[0147] VNb9O 25

[0148] W7Nb4O 31

[0149] W9Nb8O 47 。

[0150] Here, the term 'corresponding' is intended to reflect the peaks identified by X-ray diffraction analysis of the active electrode material with one or more reference crystal structures (e.g., MoNb 12 O 33 , WNb 12 O 33 , ZrNb 24 O 62 , VNb9O 25 , W7Nb4O 31 and / or W9Nb8O 47) The corresponding peak shift in the X-ray diffraction analysis does not exceed 0.5 degrees (preferably not exceeding 0.2 degrees, more preferably not exceeding 0.1 degree). Preferably, the crystal structure of the active electrode material does not correspond to the crystal structure of TiNb2O7. For example, preferably, the measured XRD diffraction pattern of the active electrode material does not correspond to the JCPDS crystallographic database entry 00-039-1407 of TiNb2O7. Optionally, the crystal structure of the active electrode material does not correspond to the crystal structure of Ti2Nb 10 O 29 . Optionally, the crystal structure of the active electrode material does not correspond to the crystal structure of M III Nb 11 O 29 , such as FeNb 11 O 29 , GaNb 11 O 29 , CrNb 11 O 29 and A1Nb 11 O 29 .

[0151] At least some of the active electrode material may have a Wadsley-Roth crystal structure and / or a tetragonal tungsten bronze (TTB) crystal structure. Preferably, most of the active electrode material has a Wadsley-Roth crystal structure and / or a tetragonal tungsten bronze (TTB) crystal structure. For example, at least 50%, at least 60%, at least 70%, at least 80% or at least 90% by volume of the active electrode material may have a Wadsley-Roth crystal structure and / or a tetragonal tungsten bronze (TTB) crystal structure. In a preferred embodiment, substantially all of the active electrode material may have a Wadsley-Roth crystal structure and / or a tetragonal tungsten bronze (TTB) crystal structure. When the material has such a crystal structure, it may have improved electrochemical properties.

[0152] The crystal formula of the charge-balanced and thermodynamically stable Wadsley-Roth crystal structure follows the following formula:

[0153] (1)(M1, M2, M3,...) mnp +1O 3mnp-(m+n)p+4

[0154] In this formula, O is oxygen (anion), and M (cation) is any combination of elements selected from Ti, Mg, V, Cr, W, Zr, Nb, Mo, Cu, Fe, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Te, Se, Si, Sb, Y, La, Hf, Ta, Re, Zn, In or Cd. In the material according to the invention, at least one of (M1, M2, M3...) contains Nb.

[0155] Formula (1) is based on the crystal morphology: m and n are the sizes of the edge-sharing superstructure blocks formed, ranging from 3 to 5 (integers). At the corners, the blocks are connected only by edge-sharing into infinite strips (p = ∞), in pairs by partial edge-sharing and partial tetrahedral connection (p = 2), or into isolated blocks only by tetrahedral connection (p = 1). When p is infinite, the formula becomes:

[0156] (2)(M1, M2, M3,...)mnO 3mn-(m+n)

[0157] Formulas (1) and (2) together define a complete composition sample of the Wadsley-Roth crystal structure. Preferably, the total crystal composition should also be charge-neutral and thermodynamically favorable.

[0158] More information can be found in the work of Griffith et al. (2017).

[0159] The tetragonal tungsten bronze (TTB) crystal structure (or simply referred to as the 'bronze' structure) mentioned in this disclosure refers to the tetragonal tungsten bronze (TTB) structure with partially filled tunnels. As described in Montemayor 1998, such phases exist in the framework of NbO6 octahedra, and the octahedra share corners in a way that forms three-, four- and five-sided tunnels. Multiple 5-sided tunnels are filled with W, Nb, O or suitable metal cations to form this structure.

[0160] The active electrode material may also contain Li and / or Na. In other words, the active electrode material can be a lithiated and / or sodiated active electrode material. The active electrode material of the first aspect can be represented by the general formula [Li] λ [M1] x [M2] (1-x) [Nb] y [O] z or [Na] λ [M1] x [M2] (1-x) [Nb] y [O] z represented. The active electrode material of the second aspect can be represented by the general formula [Li] λ [M][Nb]y [O] z or [Na] λ [M][Nb] y [O] z is represented. x, y, and z satisfy the above ranges, and λ is selected to provide a charge-balanced or substantially charge-balanced crystal structure, and / or a thermodynamically stable or metastable crystal structure.

[0161] The BET surface area of the active electrode material can be in the range of 0.1 - 100 m 2 / g, or 0.5 - 50 m 2 / g, or 1 - 20 m 2 / g. Generally, in order to minimize the reaction of the active electrode material with the electrolyte, for example, to minimize the formation of the solid electrolyte interphase (SEI) layer during the first charge-discharge cycle of an electrode containing the material, a low BET surface area is preferred. However, since most of the active electrode material cannot access the metal ions in the surrounding electrolyte, too low a BET surface area results in an unacceptably low charge rate and capacity. The term "BET surface area" refers to the surface area per unit mass calculated from the measurement of the physical adsorption of gas molecules on the solid surface using the Brunauer-Emmett-Teller theory. For example, the BET surface area can be determined according to ISO 9277:2010.

[0162] The active electrode material can include a plurality of primary microcrystals (sometimes referred to as microcrystals or microcrystallites). The average diameter of the primary microcrystals can be 10 nm - 10 μm, preferably 100 nm - 5 μm, although the most ideal diameter of the primary microcrystals can depend on their intended use. For example, when the active electrode material is intended for use in ultra-high power products, a lower primary microcrystal size, such as 50 nm or less or 30 nm or less, can be advantageous. When the active electrode material is intended for the development of "high-energy batteries", a higher microcrystal size, such as 5 μm or more or 7 μm or more, can be advantageous.

[0163] Some or all of these primary microcrystals can agglomerate into secondary particles. Alternatively, the primary microcrystals can be substantially non-agglomerated. When some or all of these primary microcrystals agglomerate into secondary particles, the average diameter of the secondary particles (e.g., D when measured using solid-state powder laser diffraction) 50The (diameter) is from 1 μm to 30 μm, preferably from 2 μm to 15 μm, but the most desired diameter of the secondary particles may depend on their intended use. For example, when the active electrode material is intended for use in ultra-high power products, a lower secondary particle size, such as 4 μm or less, 2 μm or less, or 1.5 μm or less can be advantageous. When the active electrode material is intended for the development of "high energy batteries", a higher secondary particle size, such as 8 μm or greater, 12 μm or greater, or 15 μm or greater can be advantageous. The secondary particles can be porous.

[0164] The average diameter of the primary microcrystals and / or secondary particles can be measured using any conventionally known technique, such as using SEM imaging to examine a sample of the material, selecting the number (n) of primary microcrystals and / or secondary particles, and calculating the average diameter as the average diameter of the n primary microcrystals / secondary particles measured, such as where n equals 30.

[0165] An alternative method for measuring the secondary particle size is to use solid-state powder laser diffraction, such as using a Horiba laser diffraction particle analyzer for dry powder, with the air pressure maintained at 0.3 MPa.

[0166] When measured using solid-state powder laser diffraction, the active electrode material can have a D 10 secondary particle size of at least 0.05 μm, or at least 0.1 μm, or at least 0.5 μm, or at least 1 μm. By keeping the D 10 particle size within these ranges, the likelihood of parasitic reactions in the Li-ion battery is reduced due to the reduced surface area, and it is easier to process with less binder in the electrode slurry. The term "D n " refers to the diameter at which a particle population less than n volume % is found.

[0167] When measured using solid-state powder laser diffraction, the active electrode material can have a D 90 secondary particle size of < 50 μm, < 20 μm, < 10 μm or < 5 μm. By keeping the D 90 particle size within these ranges, the proportion of the particle size distribution with large particle sizes is minimized, making it easier to fabricate the material into a homogeneous electrode.

[0168] The active electrode material can include a carbon coating formed on the surface of the primary microcrystals and / or secondary particles. Some suitable methods for forming a carbon coating on the surface of the primary microcrystals and / or secondary particles are described in the literature such as Zhou (2012). Other suitable methods are described below. Based on the total weight of the active electrode material, the carbon coating can be present in an amount of up to 5 wt%. The carbon coating can include graphite carbon.

[0169] In the case where the active electrode has a morphology of multiple primary microcrystals, some or all of which agglomerate into porous secondary particles, the secondary particles may include a carbon coating formed at least on the surface of the pores in the secondary particles.

[0170] In a third aspect, the present invention provides an electrochemical device including an anode, a cathode, and an electrolyte disposed between the anode and the cathode, wherein the anode includes an electrode active material according to the first aspect or the second aspect of the present invention.

[0171] The electrolyte may be a liquid electrolyte. Alternatively or additionally, the electrolyte may be a solid-state electrolyte.

[0172] The anode may further include a conductive additive and / or a binder. For example, the anode may have a composition of about 80 wt% active material, about 10 wt% conductive additive, and about 10 wt% binder. Alternatively, the anode may have a composition of about 91 wt% active material, about 5 wt% conductive additive, and about 4 wt% binder. The amount of the active electrode material in the anode may be in the range of 70 wt% to 99 wt%, more preferably in the range of 75 wt% to 98 wt%, and even more preferably in the range of 85 wt% to 96 wt%.

[0173] In a fourth aspect, the present invention provides a use of an electrode active material according to the first aspect or the second aspect of the present invention, the electrode active material being used as an anode active material or a component of the anode active material in an anode in combination with a cathode and an electrolyte in (i) a lithium-ion battery or (ii) a sodium-ion battery for charging and discharging of the lithium-ion battery or for charging and discharging of the sodium-ion battery.

[0174] In a fifth aspect, the present invention provides a method for processing an electrode active material according to the first aspect or the second aspect of the present invention into an anode active material for the following batteries or processing in the anode active material: (i) a lithium-ion battery, wherein the method includes diffusing lithium ions into the anode active material; or (ii) a sodium-ion battery, wherein the method includes diffusing sodium ions into the anode active material.

[0175] In a sixth aspect, the present invention provides a method for preparing an active electrode material according to the first aspect or the second aspect of the present invention, the method including the following steps:

[0176] Providing one or more precursor materials;

[0177] Mixing the precursor materials to form a precursor material mixture; and

[0178] Heat-treating the precursor material mixture in a temperature range of 400°C - 1350°C to form the active electrode material.

[0179] When preparing the material according to the first aspect of the present invention, preferably, one or more precursor materials include an M1 ion source, an M2 ion source, and a Nb source.

[0180] When preparing the material according to the second aspect of the present invention, preferably, one or more precursor materials include an M source and a Nb source.

[0181] The phrase 'M1 ion source' is used herein to describe a material containing M1 ions / atoms. The phrase 'M2 ion source' is used herein to describe a material containing M2 ions / atoms. For example, the phrase 'Mo / W / Zr / V / Nb source' is used herein to describe a material containing appropriate Mo / W / Zr / V / Nb ions / atoms.

[0182] The precursor material may include one or more metal oxides, metal hydroxides, metal salts, or oxalates. For example, the precursor material may include one or more metal oxides with different oxidation states and / or different crystal structures. Examples of suitable metal oxide precursor materials include, but are not limited to: Nb2O5, NbO2, WO3, TiO2, MoO3, V2O5, ZrO2, and MgO. However, the precursor material may not contain metal oxides, or may contain an ion source other than oxides. For example, the precursor material may contain a metal salt (e.g., NO3 - 、SO3 - ) or other compounds (e.g., oxalates). Preferably, one or more precursor materials include one or more of a Nb source, a Mo source, a W source, a Zr source, and / or a V source.

[0183] Some or all of the precursor materials may be particulate materials. When they are particulate materials, preferably they have an average particle size < 20 μm (e.g., D 50 diameter when measured by solid-state powder laser diffraction). The average particle size may be in the range of, for example, 10 nm to 20 μm. Providing particulate materials with such an average particle size can help promote closer mixing of the precursor materials, thus resulting in a more efficient solid-state reaction during the heat treatment step. However, it is not necessary for the precursor materials to have an initial particle size < 20 μm, as the particle size of one or more precursor materials can be mechanically reduced during the step of mixing the precursor materials to form a precursor material mixture.

[0184] The step of mixing / grinding the precursor materials to form a precursor material mixture can be carried out by a method selected from (but not limited to) the following: dry or wet planetary ball milling, rolling ball milling, high shear grinding, air jet grinding, and / or impact grinding. The force used for mixing / grinding can depend on the morphology of the precursor materials. For example, when some or all of the precursor materials have a large particle size (e.g., an average particle size greater than 20 μm in diameter), the grinding force can be selected to reduce the average particle size of the precursor materials so that the average particle size of the precursor material mixture is reduced to 20 μm or less in diameter. When the average particle size of the particles in the precursor material mixture is 20 μm or less, this can promote a more efficient solid-state reaction of the precursor materials in the precursor material mixture during the heat treatment step.

[0185] The step of heat-treating the precursor material mixture can be carried out for 1 hour to 24 hours, more preferably 3 hours to 14 hours. For example, the heat treatment step can be carried out for 1 hour or longer, 2 hours or longer, 3 hours or longer, 6 hours or longer, or 12 hours or longer. The heat treatment step can be carried out for 24 hours or shorter, 18 hours or shorter, 14 hours or shorter, or 12 hours or shorter.

[0186] In some methods, it can be advantageous to carry out a two-step heat treatment. For example, the precursor material mixture can be heated at a first temperature for a first time length and then heated at a second temperature for a second time length. Preferably, the second temperature is higher than the first temperature. Carrying out such a two-step heat treatment can help the solid-state reaction to form the desired crystal structure.

[0187] The step of heat-treating the precursor material mixture can be carried out in a gas atmosphere. The gas atmosphere can be an inert atmosphere or can be a reducing atmosphere. When it is necessary to prepare an oxygen-deficient material, it is preferred to carry out the step of heat-treating the precursor material mixture in an inert atmosphere or a reducing atmosphere. Suitable gas atmospheres include: air, N2, Ar, He, CO2, CO, O2, H2, and mixtures thereof.

[0188] The method can include one or more post-treatment steps after forming the active electrode material.

[0189] In some cases, the method may include a post-treatment step of heat-treating the active electrode material, sometimes referred to as "annealing". This post-treatment heat-treatment step can be carried out in a gas atmosphere different from the step of heat-treating the precursor material mixture to form the active electrode material. The post-treatment heat-treatment step can be carried out in an inert or reducing gas atmosphere. Such a post-treatment heat-treatment step can be carried out at a temperature higher than 500 °C, for example, at about 900 °C. Including the post-treatment heat-treatment step can be beneficial for forming, for example, vacancies or defects in the active electrode material, such as forming oxygen vacancies. Advantageously, the post-treatment heat-treatment step carried out in an inert or reducing gas atmosphere can increase the conductivity of the active electrode material. In this way, the active electrode material according to the second aspect can be synthesized.

[0190] In some cases, the method may include a post-treatment step of mixing the active electrode material with a carbon source and thereby forming a carbon coating on the active electrode material. Optionally, the mixture of the active electrode material and the carbon source can be heated to form a carbon coating on the active electrode material. Suitable carbon sources include, but are not limited to: carbohydrate materials (such as sugars, polymers); conductive carbon (such as carbon black); and / or aromatic carbon materials (such as pitch carbon).

[0191] A preferred method for forming a carbon coating includes the step of grinding the active electrode material with the carbon source, followed by pyrolyzing the active electrode material and the carbon source (such as in a furnace) in an inert atmosphere or a reducing atmosphere.

[0192] Another preferred method for forming a carbon coating includes mixing the active electrode material with the carbon source, dispersing the active electrode material and the carbon source in an aqueous slurry, and then spray-drying. The resulting powder can be optionally pyrolyzed. When the carbon source is, for example, conductive carbon black, it is not necessary to pyrolyze the material after spray-drying.

[0193] In some cases, the method may include a post-treatment step of grinding the active electrode material to change the particle size of the active electrode material. For example, the active electrode material can be treated by one or more methods including air jet milling, impact milling, high shear milling, sieving or ball milling. This can provide a more suitable particle size for the desired application of the active electrode material.

[0194] In another aspect, the present invention provides the use of a dopant M1 for improving the properties of a base material of an active material used as an anode of a metal ion battery, wherein the base material has the structure M2Nb y O z , and wherein the dopant is used to provide a modified material [M1] x [M2] (1-x) [Nb] y [O] z, wherein the modified material has improved properties compared to the base material. The improved properties can be an improved initial Coulomb efficiency (e.g., the difference in lithiation and delithiation capacities in the first charge / discharge cycle at C / 10 in a half-cell). When comparing high-rate charge / discharge with lower-rate charge / discharge (e.g., 5C or 10C relative to 0.5C), the improved property can be an increased capacity retention rate. The improved property can be an increased specific capacity at low charge / discharge rates (e.g., 0.1C in this text). M1, M2, x, y, and z are as defined herein.

[0195] The present invention includes combinations of the above aspects and features and preferred features, unless such combinations are clearly not permitted or are explicitly avoided. BRIEF DESCRIPTION OF THE DRAWINGS

[0196] Embodiments and experiments illustrating the principles of the present invention will now be discussed with reference to the accompanying drawings, wherein:

[0197] Figure 1 XRD diffraction patterns of Samples 1, 4, 14, 2, 5, 15, 16, 18, and 22 are shown;

[0198] Figure 2 XRD diffraction patterns of Samples 8 and 9 are shown;

[0199] Figure 3 XRD diffraction patterns of Samples 6, 7, 17, 19, and 20 are shown;

[0200] Figure 4 XRD diffraction patterns of Samples 10, 11, and 21 are shown;

[0201] Figure 5 XRD diffraction patterns of Samples 12 and 13 are shown;

[0202] Figure 6 TGA characterization of Sample 3 in air is shown;

[0203] Figure 7 Particle size distributions of Samples 1, 2, 15, and 16 are shown;

[0204] Figure 8 Particle size distribution of Sample 3 is shown;

[0205] Figure 9 SEM image of Sample 3 before pyrolysis and coated with conductive Au for imaging;

[0206] Figure 10 SEM image of Sample 3 after pyrolysis (without conductive coating);

[0207] Figure 11 SEM images of Samples 1 and 2;

[0208] Figure 12 Shows the representative lithiation and delithiation voltage curves obtained by galvanostatic cycling for Samples 1 and 16 in the first 2 cycles in a half-cell configuration, a voltage window of 1.1 - 3.0 V, and a rate of 0.05C;

[0209] Figure 13 Shows the representative lithiation and delithiation voltage curves obtained by galvanostatic cycling for Samples 6 and 7 in the first 2 cycles in a half-cell configuration, a voltage window of 1.1 - 3.0 V, and a rate of 0.05C;

[0210] Figure 14 Shows the lithiation and delithiation capacities obtained by galvanostatic cycling for Samples 1, 4, and 16 in a half-cell configuration, a voltage window of 1.1 - 3.0 V, and current densities of 0.5C, 1C, 2C, 5C (as step changes in the data);

[0211] Figure 15 Shows the lithiation capacities obtained by galvanostatic cycling for Samples 6, 7, and 17 in a half-cell configuration, a voltage window of 1.1 - 3.0 V, and current densities of 0.5C, 1C, 2C, 5C, 0.5C (as step changes in the data);

[0212] Figures 16(a) and (b) show the EIS measurements for Samples 1, 7, and 16 with different axis scales.

[0213] Figure 17 Shows the particle size distribution of Sample 16 before and after post-treatment;

[0214] Figure 18 Is an SEM image of the surface of the electrode made from Sample 22, focusing on the surface of the active material particles;

[0215] Figure 19 Shows the representative lithiation and delithiation voltage curves obtained by galvanostatic cycling for Samples 12 and 13 in the first 2 cycles in a half-cell configuration, a voltage window of 1.1 - 3.0 V, and a rate of 0.05C;

[0216] Figure E1 Shows the XRD diffraction patterns of Samples E1 and E2.

[0217] Figure E2 Shows the XRD diffraction patterns of Samples E3, E4, and E5.

[0218] Figure E3 Shows the XRD diffraction patterns of Samples E6, E7, and E8.

[0219] Figure E4 Shows the XRD diffraction patterns of Samples E9 and E10.

[0220] Figure E5 The particle size distributions of samples E2, E4, E7, and E10 are shown.

[0221] Figure E6 Representative lithiation and delithiation voltage curves obtained by galvanostatic cycling during the first 2 cycles of samples E1 and E2 in a half-cell configuration, a voltage window of 1.1 - 3.0 V, and a rate of 0.1C are shown.

[0222] Figure E7 Representative lithiation and delithiation voltage curves obtained by galvanostatic cycling during the first 2 cycles of samples E3 and E5 in a half-cell configuration, a voltage window of 1.1 - 3.0 V, and a rate of 0.1C are shown.

[0223] Figure E8 Representative lithiation and delithiation voltage curves obtained by galvanostatic cycling during the first 2 cycles of samples E6 and E7 in a half-cell configuration, a voltage window of 1.1 - 3.0 V, and a rate of 0.1C are shown.

[0224] Figure E9 Representative lithiation and delithiation voltage curves obtained by galvanostatic cycling during the first 2 cycles of samples E9 and E10 in a half-cell configuration, a voltage window of 1.1 - 3.0 V, and a rate of 0.1C are shown. The x-axis represents the state of charge (SOC) in order to be able to normalize the curves to their maximum capacity and evaluate the curve shape.

[0225] Figure E10 XRD diffraction patterns of E11 - E14 are shown. Detailed Description

[0226] Aspects and embodiments of the present invention will now be discussed with reference to the accompanying drawings. Other aspects and embodiments will be apparent to those skilled in the art. All documents mentioned herein are incorporated herein by reference.

[0227] A number of different materials were prepared and characterized, as summarized in Table 1 below. Broadly speaking, these samples can be divided into several groups:

[0228] Samples 1, 2, 3, 4, 5, 14, 15, 16, 18, and 22 belong to the same family of Wadsley - Roth phases based on MoNb 12 O 33 (M 6+ Nb 12 O 33 ,3×4 octahedral blocks, each with a tetrahedron at a corner). These blocks are connected by the edges shared between NbO6 octahedra and M 6+They are connected to each other by sharing corners between O4 tetrahedra and NbO6 octahedra. Sample 1 is the basic crystal structure, which is modified into a mixed metal cation structure by exchanging one or more cations in Samples 2 to 4 and / or the mixed crystal configuration in Samples 14, 15, 16, 18, and 22 (isomorphic with WNb 12 O 33 blending). Oxygen vacancies are generated in the basic crystal and the mixed metal cation structure 18 of Sample 5. Sample 3 is the spray-dried and carbon-coated form of the crystal prepared in Sample 2, and Sample 22 is the spray-dried and carbon-coated form of the crystal prepared in Sample 16.

[0229] Samples 6, 7, 17, 19, 20 belong to the same family of Wadsley-Roth phases based on ZrNb 24 O 62 (M 4+ NNb 24 O 62 , 3×4 octahedral blocks, each with a half tetrahedron at the corner).

[0230] Samples 8, 9, and E11 belong to the same family of Wadsley-Roth phases based on WNb 12 O 33 (M 6+ Nb 12 O 33 , 3×4 NbO6 octahedral blocks, each with a tetrahedron at the corner).

[0231] Samples 10, 11, and 21 belong to the same family of Wadsley-Roth phases based on VNb9O 25 (M 5+ Nb9O 25 , 3x3 NbO6 octahedral blocks, each with a tetrahedron at the corner).

[0232] Samples 12, 13, and E14 belong to the same family of tungsten tetragonal bronze (TTB) based on W7Nb4O 31 (M 6+ 7Nb4O 31 ). This is a tetragonal tungsten bronze structure, where MO6 (M = 0.4Nb + 0.6W) octahedra share corners completely, with 3-sided, 4-sided, and 5-sided tunnels. Some of these tunnels are filled with -O-M-O- chains, while other tunnels are open for lithium ion transport and storage.

[0233] Samples E1, E2, E13 belong to the same family based on Zn2Nb 34 O 87 (M 2+ 2Nb 34 O 87) of the same family of Wadsley-Roth phases. This orthorhombic phase consists of 3×4 MO6 octahedral blocks (M = Zn +2 / Nb +5 ), where the blocks are connected only by edge-sharing and there are no tetrahedra.

[0234] Samples E3, E4, E5, E12 belong to the same family of Wadsley-Roth phases based on AlNb 11 O 29 (M 3+ Nb 11 O 29 ). This structure belongs to a monoclinic shear structure, where the 3×4 octahedral blocks are connected only by edge-sharing and there are no tetrahedra.

[0235] Samples E6, E7, E8 belong to the same family of Wadsley-Roth phases based on GeNb 18 O 47 (M 4+ Nb 18 O 47 ). The structure is similar to Sample 10, having 3x3 NbO6 octahedral blocks and a tetrahedron connecting the blocks at the corners. However, due to Ge +4 instead of V 5+ , the structure contains intrinsic defects.

[0236] Samples E9, E10 belong to the same family of Wadsley-Roth phases based on W5Nb 16 O 55 (M 6+ 5Nb 16 O 55 ). The structure is made up of 4x5 blocks connected by edge-sharing (W, Nb)O6 on the sides and by WO4 tetrahedra at the corners. The structure is similar to Samples 8 and 9, but has larger block sizes.

[0237]

[0238]

[0239] Table 1: Summary of the different compositions synthesized. Samples marked with * are comparative samples.

[0240] Material synthesis

[0241] The samples listed in Table 1 were synthesized using a solid-state route. In the first step, metal oxide precursor commercial powders (Nb2O5, NbO2, MoO3, ZrO2, TiO2, WO3, V2O5, ZrO2, K2O, CoO, Fe2O3, GeO2, Ga2O3, Al2O3, ZnO, and / or MgO) were mixed in stoichiometric proportions and ball-milled in a zirconia jar and grinding media at a ball-to-powder ratio of 10:1 at 550 rpm for 3 hours. The resulting powder was then heated in air in a static muffle furnace to form the desired crystalline phase. Samples 1 to 5 and 12 to 16, 18, and 22 were heat-treated at 900 °C for 12 h; samples 6 to 9, 17, 19, and 20 were heat-treated at 1200 °C for 12 h, where samples 6, 7, 17, 19, and 20 underwent an additional heat-treatment step at 1350 °C for an additional 4 h; samples 10, 11, and 21 were heat-treated at 1000 °C for 12 h. Samples 3 and 22 were further mixed with a carbohydrate precursor (such as sucrose, maltodextrin, or other water-soluble carbohydrates), dispersed in an aqueous slurry at a concentration of 5, 10, 15, or 20 w / w% with an ionic surfactant, and spray-dried in a laboratory-scale spray dryer (inlet temperature 220 °C, outlet temperature 95 °C, sample introduction rate 500 mL / h). The resulting powder was pyrolyzed in nitrogen at 600 °C for 5 h. Samples 5 and 18 were further annealed in nitrogen at 900 °C for 4 hours.

[0242] Samples E1, E2, E6, E7, E8, E9, E10 were prepared by ball milling as described above and milled to a particle size distribution of D90 < 20 μm by impact milling at 20,000 rpm as required, and then heat-treated in air in a muffle furnace at 1200 °C for 12 h; samples E8, E10, E11, E12, E13 were further annealed in nitrogen at 1000 °C for 4 h; E14 was annealed in nitrogen at 900 °C for 5 h. Samples E3, E4, E5 were heat-treated at 1300 °C for 12 h. After synthesis, samples E11E10 were depolymerized to the desired particle size range by impact milling or jet milling.

[0243] Elemental analysis of the sample

[0244] Elemental analysis was performed by inductively coupled plasma - optical emission spectrometry (ICP - MS / OES). Measurements were carried out on a Thermo Scientific ICP - OES Duo iCAP 7000 series. Samples were digested using 5 ml of nitric acid and 1 ml of HF acid, and internal standards were used to account for any instrumental variations. During this process, the plasma is used to vaporize the material into its elemental atomic / ionic state. The atoms are in an excited state due to the high temperature and decay to the normal state through energy transitions. The characteristic radiation emitted by each excited ion is measured for analysis. The results are listed in Table 2 below.

[0245]

[0246] Table 2: Summary of ICP - OES elemental analysis results for samples 1, 2, 4, 14, 3, 16, 9, 11, and 17

[0247] This elemental analysis table indicates that for each composition tested, the expected cation ratios were essentially achieved.

[0248] XRD characterization of the sample

[0249] The phase purity of some samples was analyzed using a Rigaku Miniflex powder X - ray diffractometer in the 2θ range (10 - 70°) at a scan rate of 1° / min.

[0250] Figure 1 XRD diffraction patterns of the measurements for samples 1, 4, 14, 2, 5, 15, 16, 18, 22 related to comparative study A are shown. All diffraction patterns have peaks at the same positions (within the instrumental error, i.e., 0.1°) and match the JCPDS crystallographic database entry JCPDS 73 - 1322, which corresponds to MoNb 12 O 33 . There is no amorphous background noise, and the peaks are sharp and intense. This means that all samples are of pure phase and crystalline, and according to the Scherrer equation and the crystal structure of MoNb 12 O 33 , the crystallite size is approximately 200 nm.

[0251] Figure 2 XRD diffraction patterns of the measurements for samples 8 and 9 are shown. Figure E10 The XRD pattern of sample E11 is shown. All diffraction patterns have peaks at the same positions (within the instrumental error, i.e., 0.1°) and match the JCPDS crystallographic database entry JCPDS 73 - 1322, which corresponds to WNb 12 O 33。There is no amorphous background noise, and the peaks are sharp and intense. This means that all samples are pure-phase and crystalline, and according to the Scherrer equation and the crystal structure matching WNb 12 O 33 , the microcrystalline size is about 200 nm.

[0252] Figure 3 Shows the measured XRD diffraction patterns of samples 6, 7, 17, 19, 20 related to comparative study B. All diffraction patterns have peaks at the same position (within the instrument error range, i.e., 0.1°), and match the JCPDS crystallographic database entry JCPDS 01-072-1655, which corresponds to ZrNb 24 O 62 。There is no amorphous background noise, and the peaks are sharp and intense. This means that all samples are pure-phase and crystalline, and according to the Scherrer equation and the crystal structure matching ZrNb 24 O 62 , the microcrystalline size is about 200 nm.

[0253] Figure 4 Shows the measured XRD diffraction patterns of samples 10, 11, 21. All diffraction patterns have peaks at the same position (within the instrument error range, i.e., 0.1°), and match the JCPDS crystallographic database entry JCPDS 00-049-0289, which corresponds to VNb9O 25 。There is no amorphous background noise, and the peaks are sharp and intense. This means that all samples are pure-phase and crystalline, and according to the Scherrer equation and the crystal structure matching VNb9O 25 , the microcrystalline size is about 200 nm.

[0254] Figure 5 Shows the measured XRD diffraction patterns of samples 12 and 13. Figure E10 Shows the XRD pattern of sample E14. All diffraction patterns have peaks at the same position (within the instrument error range, i.e., 0.1°), and match the JCPDS crystallographic database entry JCPDS 00-020-1320, which corresponds to W7Nb4O 31 。There is no amorphous background noise, and the peaks are sharp and intense. This means that all samples are pure-phase and crystalline, and according to the Scherrer equation and the crystal structure matching W7Nb4O 31 , the microcrystalline size is about 200 nm.

[0255] Figure E1 Shows the measured XRD diffraction patterns of samples E1 and E2. Figure E10The XRD pattern of sample E13 is shown. All diffraction patterns have peaks at the same position (in the range of 0.1 - 0.2°) and match the JCPDS crystallographic database entry JCPDS 22 - 353. There is no amorphous background noise, and the peaks are sharp and intense. This means that all samples are of pure phase and crystalline, and according to the Scherrer equation and the matching crystal structure of Zn2Nb 34 O 87 , the microcrystalline size is 52 ± 12 nm.

[0256] Figure E2 The measured XRD diffraction patterns of samples E3, E4, and E5 are shown. Figure E10 The XRD pattern of sample E12 is shown. All diffraction patterns have peaks at the same position (in the range of 0.1 - 0.2°) and match the JCPDS crystallographic database entry JCPDS72 - 159 (isostructural Ti2Nb 10 O 29 ). There is no amorphous background noise, and the peaks are sharp and intense. This means that all samples are of pure phase and crystalline, and according to the Scherrer equation and the matching crystal structure of AlNb 11 O 29 , the microcrystalline size is 53 ± 16 nm.

[0257] Figure E3 The measured XRD diffraction patterns of samples E6, E7, and E8 are shown. All diffraction patterns have peaks at the same position (in the range of 0.1 - 0.2°) and match the ICSD crystallographic database entry 72683 (isostructural PNb9O 25 ). There is no amorphous background noise, and the peaks are sharp and intense. This means that all samples are of pure phase and crystalline, and according to the Scherrer equation and the matching crystal structure of GeNb 18 O 47 , the microcrystalline size is 53 ± 3 nm.

[0258] Figure E4 The measured XRD diffraction patterns of samples E9 and E10 are shown. All diffraction patterns have peaks at the same position (in the range of 0.1 - 0.2°) and match the JCPDS crystallographic database entry JCPDS 44 - 0467. There is no amorphous background noise, and the peaks are sharp and intense. This means that all samples are of pure phase and crystalline, and according to the Scherrer equation and the matching crystal structure of W5Nb 16 O 55 , the microcrystalline size is 37 ± 11 nm.

[0259] TGA characterization of the sample

[0260] Thermogravimetric analysis (TGA) was performed on some samples using a Perkin Elmer Pyris 1 system in a synthetic air atmosphere. First, the samples were held at 30 °C for 15 min, then heated from 30 °C to 950 °C at 5 °C / min, and finally held at 950 °C for 30 min. TGA was performed on Sample 3 to quantify the carbon content, and TGA was performed on Sample 5 to show the mass increase when the oxygen vacancies were filled.

[0261] Figure 6 The TGA characterization of Sample 3 in air is shown. The sharp mass drop between approximately 400 °C and 500 °C is attributed to the decomposition of the carbon coating. The decomposition temperature corresponds to a mixture of amorphous carbon and graphite carbon. The amount of mass loss indicates that Sample 3 contains a 1.1 wt% carbon coating, which is consistent with the amount expected from the stoichiometry of the precursor.

[0262] Qualitative assessment of oxygen vacancies

[0263] As discussed above, Samples 5 and 18 were heat-treated at 900 °C for 12 h to form the active electrode material, and then further annealed at 900 °C in nitrogen (reducing atmosphere) in a post-treatment heat treatment step. After the post-treatment heat treatment in nitrogen, a color change from white to dark purple was observed, indicating a change in the oxidation state and band structure of the material due to the oxygen deficiency in the samples.

[0264] Samples E8, E10, E11, E12, E13 were further annealed in nitrogen at 1000 °C for 4 h, and Sample E14 was annealed in nitrogen at 900 °C for 5 h. When induced oxygen vacancies were introduced in Sample E8, Sample E7 changed from white to dark yellow; when induced oxygen vacancies were introduced in Sample E10, Sample E9 changed from off-white to blue-gray; in Sample E11, Sample 8 changed from off-white to light blue; in Sample E12, Sample E3 changed from white to gray / black; in Sample E3, Sample E1 changed from white to gray / black; in Sample E14, Sample 12 changed from light yellow to dark blue.

[0265] Particle size distribution analysis of the sample

[0266] The particle size distribution of the dry powder was obtained using a Horiba laser diffraction particle analyzer. The air pressure was maintained at 0.3 MPa. The results are listed in Table 3 below.

[0267]

[0268]

[0269] Table 3: Statistical summary of the particle size distribution of Samples 1, 2, 15, 16, 18, 3 before pyrolysis, 3 after pyrolysis, 16 and 18 after post-treatment, and Samples E1 - E14.

[0270] Figure 7 Shows the particle size distributions of Samples 1, 2, 15, and 16 (the measured particle size is the secondary particle size, not the crystal or microcrystal size), as representative examples of the particle size distributions obtained by the solid-state route in this study, without further treatment or size optimization. The particle size distributions are generally bimodal, with a first mode of approximately 10 μm and a second mode of approximately 90 μm. As Figure 8 shown, due to the spray drying and pyrolysis post-treatment steps, Sample 3 has significant differences in the particle size distribution.

[0271] All particle size distributions can also be refined with further processing steps, such as spray drying, ball milling, high-shear milling, jet milling, or impact milling to reduce the particle size distribution to the desired range (e.g., d90 < 20 μm, < 10 μm, or < 5 μm), as Figure 17 and shown in Table 3. Generally, the particle size distribution is adjusted by optimizing the phase formation process (i.e., the solid-state synthesis route) and the post-treatment steps for the target application. For example, for lithium-ion electrodes with high power, among other considerations, a lower average particle size is typically targeted.

[0272] Figure E5 Shows the particle size distributions of Samples E2, E4, E7, and E10 in their final form, which are then processed into electrode slurries and inks.

[0273] SEM characterization of the sample

[0274] The morphology of some samples was analyzed by scanning electron microscopy (SEM).

[0275] Figure 9 and 10 Shows the SEM images of Sample 3 before and after pyrolysis. A porous microsphere morphology with a carbon coating was observed, where the primary microcrystals are organized into secondary particles. It can be seen that the material has uniform porous particles, which can be effectively packed to form a high-density electrode. Since no conductive coating is required for SEM imaging, this means that the surface conductivity of the material is increased by an order of magnitude, thus greatly improving the conductivity in terms of quality. Figure 18 Is the SEM image of the particle surface in the electrode of Sample 22, where conductive carbon black particles contained in the electrode can also be seen on the right side of the image. This clearly shows evidence of a conformal carbon coating around the MNO material.

[0276] Figure 11 Shows the SEM images of Samples 1 and 2 and confirms the XRD and PSD data, showing micron-sized particles of dense secondary particles composed of primary microcrystals of approximately 200 nm.

[0277] Electrochemical testing of the sample

[0278] Electrochemical tests were conducted in coin half-cells (CR2032 size) for initial analysis. In the coin half-cell tests, the materials were tested in an electrode relative to a lithium metal electrode to evaluate their basic properties. In the following examples, the active material composition to be tested was combined with N-methylpyrrolidone (NMP), carbon black acting as a conductive additive, and poly(vinylidene fluoride) (PVDF) binder, and mixed using a laboratory-scale centrifugal planetary mixer to form a slurry (although an aqueous slurry can also be formed by using water instead of NMP). The non-NMP composition of the slurry was 80 wt% active material, 10 wt% conductive additive, and 10 wt% binder. The slurry was then coated onto an aluminum foil current collector by doctor blading to the desired loading of 1 mg / cm 2 and dried in a vacuum oven for 12 hours. The electrode was punched into the desired size and combined with a separator (Celgard porous PP / PE), lithium metal, and electrolyte (1 M LiPF6 in EC / DEC) in a steel coin cell housing and sealed under pressure. Then, a formation cycle of 2 full charge and discharge cycles was carried out at a low current rate (C / 20). After formation, further cycling can be carried out at a fixed or varying current density as needed. These tests are referred to as "half-cell constant current cycling" for future reference. For samples E1-E10, the electrolyte was changed to 1.3 M LiPF6 in 3:7 EC / DEC, and a formation cycle of 2 charge / discharge cycles was carried out at C / 10 within a voltage limit of 1.1 - 3.0 V. The values shown for these samples are the average of 3 measurements, and the error is the standard deviation.

[0279] A uniform and smooth coating on the current collector foil was also prepared as described above using a centrifugal planetary mixer with a composition of 94 wt% active material, 4 wt% conductive additive, and 2 wt% binder without visible defects. The coating was calendered at 80 °C to a density of up to 3.0 g / cm 2 at a loading of 1.3 - 1.7 mAh / cm 3 to demonstrate a possible volumetric capacity of > 700 mAh / cm 3 in the voltage range of 0.7 - 3.0 V at C / 20 and > 640 mAh / cm 3 in the voltage range of 1.1 - 3.0 V at C / 5. This is an important demonstration of the viability of these materials in electrode power battery formulations of commercial interest, where maintaining performance after calendering to high electrode density allows for high volumetric capacity. Loadings up to and including 1.0, 1.5, 2.0, 2.5, or 3.0 mAh / cm 2 can be used for Li-ion batteries focused on power performance; loadings greater than 3.0, 4.0, or 5.0 mAh / cm 2The loading is useful for the energy concentration performance in Li-ion batteries. These materials are calendared to an electrode porosity value of 35%, and typically in the range of 35 - 40%; defined as the measured electrode density adjusted to its w / w% divided by the average of the true density of each electrode assembly.

[0280] The conductivity of the electrodes made from the samples listed in Table 1 was measured using a 4-point probe thin film resistance measurement device. The slurry was prepared according to the above procedure and coated on a dielectric polyester film at a loading of 1 mg / cm 2 . Disks of electrode size were then punched out and the resistance of the coated film was measured using a 4-point probe. The bulk resistivity can be calculated from the measured resistance using the following equation:

[0281] (3) Bulk resistivity (ρ) = 2πs(V / I); R = V / I; s = 0.1 cm

[0282] = 2π x 0.1 x R (Ω)

[0283] The results of this test are shown in Table 4 below:

[0284] Sample Resistance [kΩ] Volume resistivity [kΩ.cm] 1* 8.5 5.3 2 1.7 1.1 4 3.2 2.0 5 0.52 0.33 6* 0.37 0.23 7 0.52 0.33 13 0.45 0.28 14 2.7 1.7 15 1.2 0.75 16 1.3 0.82 17 0.34 0.21 18 0.89 0.56 19 0.18 0.11 20 0.20 0.13 22 0.33 0.21

[0285] Table 4 - Summary of 4-point probe resistivity measurements for Samples 1, 2, 4, 5, 6, 7, 13 to 20, and 22.

[0286] Samples E1 - E14 were also measured for their 4-point probe resistance to quantify their resistivity. This measurement was made on the coatings on polyester films at a loading of 1.0 mg / cm at 23 °C using different Ossila instruments (T2001A3-UK). The results of the sheet resistance (Ω / square) are listed in Table 4a, with the error based on the standard deviation of 3 measurements. 2

[0287] Sample Sheet resistivity [Ω / square] E1* 1242±156 E2 1041±103 E3* 1396±74 E4 1215±52 E5 1057±35 E6* 1092±52 E7 1009±89 E8 965±83 E9* 1135±92 E10 1113±99 E12 891±61 E13 1027±13 12* 853±51 E14 846±57 6* 880±29

[0288] Table 4a - Summary of 4-point probe resistivity measurements for Samples E1 to E14.

[0289] The direct current internal resistance (DCIR) and the resulting area specific impedance (ASI) are key metrics of the internal resistance of the electrodes in a Li-ion battery. In a typical measurement, a battery that has undergone formation will be cycled 3 times at C / 2. When the electrode is in the delithiated state, a C / 2 discharge current is applied for 1 h to achieve approximately 50% lithiation. The battery is allowed to rest for 30 min to equilibrate at its OCV (open circuit voltage), and then a 5C current pulse is applied for 10 s, followed by a 30-min rest to reach the OCV. During the 10-s pulse, the voltage response is sampled at a higher frequency to accurately determine the average internal resistance. The resistance is then calculated from V = IR using the difference between the OCV (the linear average between the initial OCV before the pulse and the OCV after) and the measured voltage. The resistance is then multiplied by the electrode area to obtain the ASI.

[0290] The results of this test are shown in Table 5 below:

[0291] Sample <![CDATA[ASI / Ω.cm 2 > 1* 141 2 125 4 120 6＊ 126 7 162 13 67 14 99 16 74 17 162 18 75 19 164 22 121

[0292] Table 5 - Summary of DCIR / ASI measurements for Samples 1, 2, 4, 7, 14, 16, and 17.

[0293] The reversible specific capacity C / 20, initial Coulombic efficiency, nominal lithiation voltage relative to Li / Li + at C / 20, 5C / 0.5C capacity retention, and 10C / 0.5C capacity retention were also tested for many samples, and the results are shown in Table 6 below. The nominal lithiation voltage relative to Li / Li+ was calculated by integrating the V / Q curve and dividing by the total capacity of the 2nd cycle C / 20 lithiation. The capacity retention at 10C and 5C was calculated by dividing the specific capacity at 10C or 5C by the specific capacity at 0.5C. It should be noted that the capacity retention was tested by a symmetric cycling test, where the lithiation and delithiation had equal C-rates. When tested with an asymmetric cycling program, a 10C / 0.5C capacity retention greater than 89% was routinely observed.

[0294] Samples E1 - E10 were tested with the minor differences in Table 6a. The reversible specific capacity shown is the 2nd cycle delithiation capacity at C / 10, and the nominal lithiation voltage relative to Li / Li + was at C / 10 in the 2nd cycle. The rate tests were performed with an asymmetric cycling program without a constant voltage step (i.e., constant current), where lithiation was at C / 5 and delithiation was at an increasing C-rate.

[0295]

[0296] Table 6 - Summary of the electrochemical test results of Li - ion half - button cells using multiple samples. Generally (but not exclusively), it is beneficial to have higher capacity, higher ICE, lower nominal voltage, and higher capacity retention rate.

[0297]

[0298]

[0299] Table 6a - Summary of the electrochemical test results of Li - ion half - button cells using multiple samples.

[0300] The modification of the Wadsley - Roth and bronze structures based on mixed niobium oxides as described in the claims demonstrates the applicability of the present invention in improving the performance of active materials in Li - ion batteries. By substituting non - Nb cations to form the mixed cationic structure, the entropy in the crystal structure can be increased (cf disorder), and the potential energy barrier for Li - ion diffusion can be reduced by introducing minute defects (e.g., samples E7, 16). The modification carried out by generating a mixed cationic structure that maintains the same overall oxidation state demonstrates the potential improvement through changing the ionic radius, such as substituting W 6+ for Mo 6+ in sample 14 or substituting Fe 3+ or Ga 3+ for Al 3+ in samples E4 and E5, which can cause minute changes in the crystal parameters and Li - ion holes (e.g., adjusting the reversibility of type VI holes in the Wadsley - Roth structure), which can improve the specific capacity, Li - ion diffusion, and increase the Coulombic efficiency of cycling by reducing Li - ion trapping. The modification carried out by generating a mixed cationic structure that results in an increase in the oxidation state (e.g., substituting Ge 4+ for Zn 2+ in sample E2, or substituting Mo 6+ for Zr 4+ ) demonstrates potential advantages similar to those of the changed ionic radius related to the same capacity and efficiency, which are complexed by introducing additional electron holes in the structure to contribute to conductivity. The modification carried out by generating a mixed cationic structure that results in a decrease in the oxidation state (e.g., substituting K + and Co 3+ for Ge 4+ in sample E7, or substituting Ti 4+ for Mo 6+) demonstrated the potential advantages associated with similar modified ionic radii related to the same capacity and efficiency, which were compounded by introducing oxygen vacancies and extra electrons in the structure to contribute to conductivity. Modification carried out by inducing oxygen deficiency through high-temperature treatment under inert or reducing conditions demonstrated the loss of a small portion of oxygen in the structure, providing a reduced structure with greatly improved conductivity (such as samples 5, E10, and E12 - 14) and improved electrochemical properties such as capacity retention at high C-rates (such as samples 5, E13). The hybrid cationic structure and the induced combination allowed for the compounding of various beneficial effects of oxygen deficiency (such as increased specific capacity, reduced resistance) (such as samples 18, E8).

[0301] Figure 12 , 13 and 19 show the representative lithiation / delithiation curves of unmodified and modified MoNb 12 O 33 ( Figure 12 - samples 1 and 6), ZrNb 24 O 62 ( Figure 13 - samples 6 and 7) and W7Nb4O 31 ( Figure 19 - samples 12 and 13) at a C / 20 rate during their first two formation cycles. In Figure 12 , approximately 90% of the specific capacity of the shown sample 16 is shown in a narrow voltage range of about 1.2 - 2.0 V, and in Figure 13 , approximately 90% of the capacity of the shown sample 7 is shown in a narrow range of about 1.25 - 1.75 V; these data highlight the attractive voltage curves obtainable with MNO crystals based on the Wadsley - Roth crystal structure. In Figure 19 , approximately 90% of the specific capacity of sample 13 is shown in a narrow range of about 1.2 - 2.2 V; this indicates that an attractive voltage curve is achieved with MNO crystals based on the tetragonal bronze crystal structure. Secondly, compared with the unmodified crystal samples 1, 6, and 12, the composite metal oxide samples 7, 16, and 13 show improved specific capacity. This is because the cations contained in the composite structure increase the number of sites in the crystal that can accommodate lithium ions due to their different ionic radii and oxidation states, thus increasing the capacity. An increase in ICE was observed between samples 1 and 16 and between samples 12 and 13, which further demonstrates that when the lithium ion sites are modified to enable their insertion / extraction, the lithium ions embedded in the modified crystal structure can be more effectively delithiated.

[0302] Figure E5 shows the particle size distributions of samples E2, E4, E8, E11, which mainly contain a single peak with a narrow distribution, i.e., the values of D 10 and D 90 are similar to D50 This is beneficial for processing the materials in the electrode slurry to effectively fill the materials and maintain uniform electrochemical properties (e.g., due to the shorter diffusion distance, smaller particles will be fully lithiated before larger particles).

[0303] Figure E6 Shows the advantages in the modified sample E1, especially regarding the increase in the observed specific capacity by replacing Zn 4+ cations with higher-valent Ge 2+ cations. Figure E7 Demonstrates the improved specific capacity observed on the modified sample E6 by replacing Ge4+ with K and Co cations, i.e., cations with a lower valence. Figure E9 Shows the improvement in ICE and the possible reduction in the nominal lithiation voltage by introducing induced oxygen vacancies, which reduce the polarization effect by increasing the conductivity and by improving the reversibility of the lithiation / delithiation process.

[0304] Among all the materials tested, each material according to the present invention shows an improvement relative to the unmodified 'base' crystal structure. This can be inferred from resistivity / impedance measurements carried out by two different methods and electrochemical tests in Li-ion half coin cells, especially the capacity retention at increased current density (see rate, Table 6, Figure 14 and 15 ). Without wishing to be bound by theory, the inventors propose that this is the result of an increase in the ionic and electronic conductivity of the material upon introduction of defects or by changing the lattice by altering the ionic radius; the DCIR / ASI (Table 5) and EIS (Figure 16) measurements also demonstrate a reduction in resistance or impedance after material modification. The Li-ion diffusion rate in the materials according to the present invention may also increase compared to the unmodified 'base' materials. As shown in Table 6, in some cases, the specific capacity itself may also increase because doping / exchanging metal ions of different sizes can expand or contract the lattice and allow more or greater reversibility of Li-ion insertion or Li-ion insertion than possible in the unmodified structure.

[0305] The data in Table 4 show a significant reduction in resistivity between Sample 1 (comparison) and Samples 2, 4, 5, 14, 15, 16, 18, 22, indicating the effect of the embodiments of the present invention on improving the conductivity of the crystal structure through cation exchange, oxygen deficiency, and carbon coating. Samples 17, 19, and 20 also show low resistivity similar to Sample 6. When 0.05 equivalents of V substance is introduced into the base crystal of Sample 7, the resistivity slightly increases, however, an improvement in specific capacity is observed due to the change in the available Li-ion sites in the lattice, which may be due to the different ionic radius of V relative to Zr (see Table 6).

[0306] The data in Table 5 show that DCIR / ASI decreases significantly from Sample 1 (comparison) to Samples 2, 4, 14, 16, 18, and 22, which reflects the trend shown in Table 4. Samples 7, 17, and 19 are confirmed by DCIR to be higher than these samples, however these involve different underlying crystal structures. Without wishing to be bound by theory, the inventors hypothesize that Samples 7, 17, and 19 show an increase in DCIR / ASI compared to the comparative material of Sample 6 (ZrNb 24 O 62 ), which is due to lattice changes caused by the introduction of cations with different ionic radii. However, in terms of the conductivity of these structures of Samples 17 and 19, it is still beneficial because the resistivity is reduced, as shown in Table 4, thereby minimizing Joule heat and enabling a more uniform current distribution across the material, which in turn improves the safety and lifespan of the lithium-ion system. For Sample 7, although the exchange of V with Zr did not show improvement, the specific capacity increased, as discussed above.

[0307] In Table 6, in most samples, there is a trend for the materials according to the present invention to show improvement in specific capacity, initial Coulomb efficiency (ICE), nominal lithiation voltage relative to Li / Li + , and importantly capacity retention at 5C and 10C relative to 0.5C, compared to the comparative 'base' materials (e.g., Samples 1, 6, 8, 10, 12). For example, compared to Sample 1, Samples 2, 3, 4, 5, 14, 15, 16, 18, 22 all show improvement in one or more of these parameters. The same is true for Samples 7, 17, 19 compared to Sample 6 in multiple parameters; Samples 11 and 21 compared to 10, where improvement in specific capacity or capacity retention was observed; Sample 9 compared to 8, where ICE and capacity retention were improved; and Sample 13 compared to 12 * , where improvement in all parameters was observed.

[0308] Figure 14 and 15 show that the materials according to the present invention (Samples 4, 16, 7, 17) have improved capacity retention at higher cycling rates compared to the comparative materials (Samples 1 and 6).

[0309] Electrochemical impedance spectroscopy (EIS) measurements were also carried out to further understand the impedance present in the electrodes of Li-ion batteries. In a typical measurement, the battery was prepared as for the DCIR measurement at about 50% lithiation, and then the frequency of the alternating charge / discharge current pulses was varied while measuring the impedance. By plotting the real and imaginary parts as axes and varying the AC frequency, a Nyquist plot was generated. From this plot of the Li-ion battery, different types of impedance in the battery can be identified, but it is usually difficult to interpret. For example, the ohmic resistance can be partially separated from the electrochemical double layer effect and also from the diffusion effect.

[0310] Figures 16(a) and (b) show the EIS spectra of (comparative) sample 1 and samples 16 and 7 (samples according to the invention).

[0311] ***

[0312] Features disclosed in the foregoing description, or in the following claims, or in the drawings, in their specific forms, or in terms of a means for performing the disclosed function, or a method or process for obtaining the disclosed result, may, where appropriate, be used alone or in any combination of such features in different forms to implement the invention.

[0313] Although the invention has been described in connection with the above exemplary embodiments, many equivalent modifications and variations will be apparent to those skilled in the art when the present disclosure is given. Accordingly, the above exemplary embodiments of the invention are considered illustrative and not restrictive. Various changes may be made to the described embodiments without departing from the spirit and scope of the invention.

[0314] For the avoidance of any doubt, any theoretical explanations provided herein are provided to enhance the reader's understanding. The inventors do not wish to be bound by any of these theoretical explanations.

[0315] Any section headings used herein are for organizational purposes only and should not be construed as limiting the subject matter described.

[0316] In this specification (including the appended claims), unless the context requires otherwise, the words "comprise" and "include" and variations such as "comprises", "comprising" and "including" will be understood to imply the inclusion of the stated integer or step or group of integers or steps, but not the exclusion of any other integer or step or group of integers or steps.

[0317] It should be noted that, as used in the specification and the appended claims, unless the context clearly indicates otherwise, the singular forms "a / an" and "the" include plural referents. Ranges may be expressed herein as from "about" one particular value and / or to "about" another particular value. When expressing such a range, another embodiment includes from one particular value and / or to another particular value. Similarly, when values are expressed as approximations by use of the antecedent "about", it should be understood that the particular value forms another embodiment. The term "about" associated with a numerical value is optional and means, for example, + / - 10%.

[0318] Numbered embodiments

[0319] The following numbered embodiments form part of this specification.

[0320] 1. An active electrode material represented by the general formula [M1] x [M2] (1-x) [Nb] y [O] z wherein:

[0321] M1 and M2 are different;

[0322] M1 represents one or more of Ti, Mg, V, Cr, W, Zr, Nb, Mo, Cu, Fe, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Te, Se, Si, Sb, Y, La, Hf, Ta, Re, Zn, In or Cd;

[0323] M2 represents one or more of Mg, V, Cr, W, Zr, Nb, Mo, Cu, Fe, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Te, Se, Si, Sb, Y, La, Hf, Ta, Re, Zn, In or Cd; and wherein

[0324] x satisfies 0 < x < 0.5;

[0325] y satisfies 0.5 ≤ y ≤ 49;

[0326] z satisfies 4 ≤ z ≤ 124

[0327] 2. The active electrode material according to embodiment 1, wherein M2 is selected from one or more of Mo, W, V or Zr.

[0328] 3. The active electrode material according to embodiment 2, wherein the [M1] x [M2] (1-x) [Nb] y [O] zSelected from the group consisting of the following materials:

[0329] M1 x Mo (1-x) Nb 12 O (33-33α)

[0330] M1 x W (1-x) Nb 12 O (33-33α)

[0331] M1 x V (1-x) Nb9O (25-25α)

[0332] M1 x Zr (1-x) Nb 24 O (62-62α)

[0333] M1 x W (1-x) Nb 0.57 O (4.43-4.43α)

[0334] M1 x W (1-x) Nb 0.89 O (5.22-5.22α)

[0335] wherein M1 represents one or more of Ti, Mg, V, Cr, W, Zr, Nb, Mo, Cu, Fe, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Te, Se, Si, Sb, Y, La, Hf, Ta, Re, Zn, In or Cd; and wherein

[0336] x satisfies 0 < x < 0.5; and

[0337] α satisfies 0 ≤ α ≤ 0.05.

[0338] 4. The active electrode material according to any one of the foregoing embodiments, wherein the active electrode material is oxygen-deficient.

[0339] 5. An active electrode material represented by the general formula [M] x [Nb] y [O] (z'-z'α) selected from the group consisting of:

[0340] MoNb 12 O (33-33α)

[0341] WNb 12 O (33-33α)

[0342] VNb9O (25-25α)

[0343] ZrNb 24 O (62-62α)

[0344] W7Nb4O (31-31α)

[0345] W9Nb8O (47-47α)

[0346] wherein α satisfies 0 < α ≤ 0.05.

[0347] 6. The active electrode material according to any one of the foregoing embodiments, wherein at least some of the material has a Wadsley-Roth crystal structure and / or a tetragonal tungsten bronze crystal structure.

[0348] 7. The active electrode material according to any one of the foregoing embodiments, wherein the active electrode material comprises a plurality of primary microcrystals, and some or all of the primary microcrystals are optionally agglomerated into secondary particles.

[0349] 8. The active electrode material according to embodiment 7, wherein the average diameter of the primary microcrystals is from 10 nm to 10 μm.

[0350] 9. The active electrode material according to embodiment 7 or embodiment 8, wherein some or all of the primary microcrystals are agglomerated into secondary particles, and the average diameter of the secondary particles is from 1 μm to 30 μm.

[0351] 10. The active electrode material according to any one of the foregoing embodiments, wherein the active electrode material comprises a carbon coating formed on the surface of the primary microcrystals and / or the secondary particles.

[0352] 11. The active electrode material according to embodiment 10, wherein the carbon coating is present in an amount of up to 5 w / w% based on the total weight of the active electrode material.

[0353] 12. The active electrode material according to any one of the foregoing embodiments, wherein the crystal structure of the active electrode material corresponds to the crystal structure of one or more of the following, as determined by X-ray diffraction analysis:

[0354] MoNb 12 O 33

[0355] WNb 12 O 33

[0356] ZrNb24 O 62

[0357] VNb9O 25

[0358] W7Nb4O 31

[0359] W9Nb8O 47 。

[0360] 13. The active electrode material according to any one of the foregoing embodiments further comprises Li and / or Na.

[0361] 14. An electrochemical device comprising an anode, a cathode, and an electrolyte disposed between the anode and the cathode, wherein the anode comprises the electrode active material according to any one of Embodiments 1 to 13.

[0362] 15. Use of an electrode active material according to any one of Embodiments 1 to 13, the electrode active material being used as an anode active material or a component of an anode active material in an anode in combination with a cathode and an electrolyte in (i) a lithium-ion battery or (ii) a sodium-ion battery for charging and discharging of the lithium-ion battery or for charging and discharging of the sodium-ion battery.

[0363] 16. A method for processing an electrode active material according to any one of Embodiments 1 to 13 into an anode active material for the following batteries or processing in the anode active material: (i) a lithium-ion battery, wherein the method comprises diffusing lithium ions into the anode active material; or (ii) a sodium-ion battery, wherein the method comprises diffusing sodium ions into the anode active material.

[0364] 17. A method for preparing an active electrode material according to any one of Embodiments 1 to 13, the method comprising the following steps:

[0365] Providing one or more precursor materials;

[0366] Mixing the precursor materials to form a precursor material mixture; and

[0367] Heat-treating the precursor material mixture in a temperature range of 400°C - 1350°C to form the active electrode material.

[0368] 18. The method for preparing an active electrode material according to Embodiment 17, wherein the one or more precursor materials comprise a Mo source, a W source, a Zr source, or a V source and a Nb source.

[0369] 19. The method for preparing an active electrode material according to embodiment 17 or embodiment 18, wherein the one or more precursor materials include an M1 ion source, an M2 ion source, and a Nb source, and wherein the resulting active electrode material is a material as defined in any one of embodiments 1 to 4 or embodiments 6 to 13 subordinate to embodiment 1.

[0370] 20. The method for preparing an active electrode material according to embodiment 17, wherein the precursor material includes one or more metal oxides, metal hydroxides, metal salts, or oxalates.

[0371] 21. The method according to any one of embodiments 17 to 20, wherein the one or more precursor materials are particulate materials, optionally having an average particle size < 20 μm.

[0372] 22. The method according to any one of embodiments 17 to 21, wherein the step of mixing the precursor materials to form a precursor material mixture is carried out by a method selected from dry or wet planetary ball milling, rolling ball milling, high shear grinding, air jet grinding, and / or impact grinding.

[0373] 23. The method according to any one of embodiments 17 to 22, wherein the step of heat-treating the precursor material mixture is carried out for a time of 1 to 14 h.

[0374] 24. The method according to any one of embodiments 17 to 23, wherein the step of heat-treating the precursor material mixture is carried out in a gas atmosphere selected from air, N2, Ar, He, CO2, CO, O2, H2, and mixtures thereof.

[0375] 25. The method according to any one of embodiments 17 to 24, wherein the method includes one or more post-treatment steps selected from the following:

[0376] (i) Heat-treating the active electrode material;

[0377] (ii) Mixing the active electrode material with a carbon source and optionally further heating the mixture to form a carbon coating on the active electrode material;

[0378] (iii) Spray-drying the active electrode material; and / or

[0379] (iv) Grinding the active electrode material to change the particle size of the active electrode material.

[0380] References

[0381] Numerous publications are cited above in order to more fully describe and disclose the present invention and the prior art to which the present invention pertains. The complete citations of these references are provided below. The entire contents of each of these references are incorporated herein by reference.

[0382] Goodenough and Park, “The Li-on Rechargeable Battery: A Perspective”, Journal of the American Chemical Society 2013 135(4), 1167-1176, DOI: 10.1021 / ja3091438

[0383] Griffith et a., “High-Rate Intercalation without Nanostructuring inMetastable Nb2O5 Bronze Phases. Journal of the American Chemical Society 2016138(28), 8888-8899, DOI: 10.1021 / jacs.6b04345

[0384] Griffith et a., “Structural Stability from Crystallographic Shear inTiO2Nb2O5 Phases: Cation Ordering and Lithiation Behavior of TiNb 24 O 62 ”InorganicChemistry(2017), 56, 7, 4002-4010

[0385] Montemayor et al., “Lithium insertion in two tetragonel tungstenbronze type phases, M8W9O47(M = Nb and Ta)”, Journal of Material Chemistry(1998), 8, 2777-2781

[0386] Zhou et al., “Facile Spray Drying Route for the Three-Dimensional Graphene Encapsulated Fe2O3 Nanopartices for Lithium lon Battery Anodes”, Ind. Eng. Chem. Res. (2013), 52, 1197-1204

[0387] Zhu et al., “MoNb 12 O 33 as a new anode material for high-capacity, safe, rapid and durable Li + storage: structural characteristics, electrochemieal properties and working mechanisms”, J. Mater. Chem. A. (2019), 7, 6522-6532

[0388] Yang et al., “Porous ZrNb 24 O 62 Nanowires with Pseudocapacitive Behavior Achieve High-Performance Lithium-lon Storage”. J. Mater. Chem. A. (2017) 5.10.1039 / C7TA07347J.

Claims

1. An active electrode material, which is represented by the general formula [M][Nb] y [O] z ; wherein the active electrode material has oxygen vacancies; wherein M consists of one of Mg, Cr, W, Mo, Cu, Ga, Ge, Ca, K, Ni, Co, Al, Sn, Mn, Ce, Sb, Y, La, Hf, Ta, Zn, In, or Cd; y satisfies 0.5 ≤ y ≤ 49; and z satisfies 4 ≤ z ≤ 124; wherein at least some of the material has a Wadsley-Roth crystal structure by volume.

2. The active electrode material according to claim 1, wherein z is defined as z = (z’ - z’α), where α satisfies 0 < α ≤ 0.

05.

3. The active electrode material according to claim 1, wherein (i) M consists of one of Mo, W, Al, Zn, Ga, Ge, Ta, Cr, Cu, K, Mg, Ni, or Hf; or (ii) M consists of one of Mo, W, Al, Zn, Ga, or Ge; or (iii) M consists of one of Mo, W, Al, or Zn.

4. The active electrode material according to claim 1, which is represented by the general formula [M] x [Nb] y [O] (z’-z’α) and is selected from the group consisting of: MoNb 12 O (33-33 α) , WNb 12 O (33-33α) , Zn2Nb 34 O (87-87 α) , Cu2Nb 34 O (87-87 α) , AlNb 11 O (29-29 α) , GaNb 11 O (29-29 α) , GeNb 18 O (47-47 α) , W5Nb 16 O (55-55 α) , AlNb 49 O (124-124 α) ,or GaNb 49 O (124-124 α) ; where α satisfies 0 < α ≤ 0.

05.

5. The active electrode material according to claim 1, which is represented by the general formula [M] x [Nb] y [O] (z’-z’α) and is selected from the group consisting of: MoNb 12 O (33-33 α) , WNb 12 O (33-33α) , Zn2Nb 34 O (87-87 α) , AlNb 11 O (29-29 α) , GeNb 18 O (47-47 α) ; where α satisfies 0 < α ≤ 0.

05.

6. The active electrode material according to claim 1, which is represented by the general formula [M] x [Nb] y [O] (z’-z’α) and is selected from the group consisting of: MoNb 12 O (33-33 α) , WNb 12 O (33-33α) ; where α satisfies 0 < α ≤ 0.

05.

7. The active electrode material according to claim 1, wherein the active electrode material has the formula MoNb 12 O (33-33 α) , where α satisfies 0 < α ≤ 0.

05.

8. The active electrode material according to claim 1, wherein the active electrode material has the formula WNb 12 O (33-33 α) , where α satisfies 0 < α ≤ 0.

05.

9. The active electrode material according to claim 1, wherein the active electrode material has the formula W5Nb 16 O (55-55 α) , where α satisfies 0 < α ≤ 0.

05.

10. The active electrode material according to claim 1, wherein the active electrode material has the formula Zn2Nb 34 O (87-87 α) , where α satisfies 0 < α ≤ 0.

05.

11. The active electrode material according to claim 1, wherein the active electrode material has the formula AlNb 11 O (29-29 α) , where α satisfies 0 < α ≤ 0.

05.

12. The active electrode material according to claim 1, wherein the active electrode material is represented by the general formula [W][Nb] y [O] z ; optionally, wherein the active electrode material is selected from WNb 12 O (33-33α) and W5Nb 16 O (55-55 α) , where α satisfies 0 < α ≤ 0.

05.

13. The active electrode material according to claim 1, wherein, as determined by X-ray diffraction, the crystal structure of the active electrode material corresponds to the crystal structure of the unmodified form of the active electrode material, wherein the unmodified form is represented by the general formula [M][Nb] y [O] z wherein the unmodified form is not oxygen-deficient, and wherein the unmodified form is selected from M2 I Nb5O 13 、M2 I 6Nb 10.8 O 30 、M2 II Nb2O6、M2 II 2Nb 34 O 87 、M2 III Nb 11 O 29 、M2 III Nb 49 O 124 、M2 IV Nb 24 O 62 、M2 IV Nb2O7、M2 IV 2Nb 10 O 29 、M2 IV 2Nb 14 O 39 、M2 IV Nb 14 O 37 、M2 IV Nb6O 17 、M2 IV Nb 18 O 47 、M2 V Nb9O 25 、M2 V 4Nb 18 O 55 、M2 V 3Nb 17 O 50 、M2 VI Nb 12 O 33 、M2 VI 4Nb 26 O 77 、M2 VI 3Nb 14 O 44 、M2 VI 5Nb 16 O 55 、M2 VI 8Nb 18 O 69 、M2 VI Nb2O8、M2 VI 20 Nb 22 O 115 、M2 VI 82 Nb 54 O 381 、M2 VI 31 Nb 20 O 143 、M2 VI 15 Nb2O 50 、M2 VI 3Nb2O 14 and M2 VI 11 Nb 12 O 63 ; wherein the numbers I, II, III, IV, V and VI represent the oxidation state of M.

14. The active electrode material according to claim 1, wherein all of the active electrode material has a Wadsley-Roth crystal structure.

15. The active electrode material according to claim 1, wherein the active electrode material comprises a plurality of primary microcrystals, and some or all of the primary microcrystals are optionally agglomerated into secondary particles.

16. The active electrode material according to claim 15, wherein the average diameter of the primary microcrystals is from 10 nm to 10 μm.

17. The active electrode material according to claim 15, wherein some or all of the primary microcrystals are agglomerated into secondary particles, and the average diameter of the secondary particles is from 1 µm to 30 µm.

18. The active electrode material according to claim 15, wherein the active electrode material comprises a carbon coating formed on the surface of the primary microcrystals and / or secondary particles.

19. The active electrode material according to claim 18, wherein the carbon coating is present in an amount of up to 5 wt% based on the total weight of the active electrode material.

20. The active electrode material according to claim 1, wherein the BET surface area of the active electrode material is in the range of 0.1-100 m 2 / g, or 0.5-50 m 2 / g, or 1-20 m 2 / g.

21. The active electrode material according to claim 1, wherein as determined by X-ray diffraction analysis, the crystal structure of the active electrode material corresponds to one or more of the following crystal structures: (i) MoNb 12 O 33 , WNb 12 O 33 , Zn2Nb 34 O 87 , Cu2Nb 34 O 87 , AlNb 11 O 29 , GaNb 11 O 29 , GeNb 18 O 47 , W5Nb 16 O 55 , AlNb 49 O 124 , GaNb 49 O 124 ; or (ii) MoNb 12 O 33 , WNb 12 O 33 , Zn2Nb 34 O 87 , AlNb 11 O 29 , GeNb 18 O 47 ; or (iii) MoNb 12 O 33 , WNb 12 O 33 。 22. The active electrode material according to claim 1, which further comprises Li and / or Na.

23. An electrochemical device, comprising an anode, a cathode, and an electrolyte disposed between the anode and the cathode, wherein the anode comprises the active electrode material according to any one of claims 1 to 22.

24. Use of the active electrode material according to claim 1, wherein the active electrode material is used as an anode active material or a component of the anode active material in an anode combined with a cathode and an electrolyte in (i) a lithium-ion battery or (ii) a sodium-ion battery for charging and discharging of the lithium-ion battery or for charging and discharging of the sodium-ion battery.

25. A method for processing the active electrode material according to claim 1 into an anode active material for the following batteries or processing it in the anode active material: (i) a lithium-ion battery, wherein the method includes diffusing lithium ions into the anode active material; or (ii) a sodium-ion battery, wherein the method includes diffusing sodium ions into the anode active material.

26. A method for preparing the active electrode material according to claim 1, the method comprising the following steps: providing one or more precursor materials; mixing the one or more precursor materials to form a precursor material mixture; and heat-treating the precursor material mixture in a temperature range of 400 °C - 1350 °C to form the active electrode material; wherein the method includes a post-treatment step of heat-treating the active electrode material in an inert or reducing gas atmosphere at a temperature higher than 500 °C to form oxygen vacancies in the active electrode material.

27. The method for preparing an active electrode material according to claim 26, wherein the one or more precursor materials include an M source and an Nb source.

28. The method for preparing an active electrode material according to claim 26, wherein the one or more precursor materials include one or more metal oxides, metal hydroxides, metal salts or oxalates.

29. The method according to claim 26, wherein the one or more precursor materials are particulate materials, optionally having an average particle diameter < 20 µm.

30. The method according to claim 26, wherein mixing the one or more precursor materials to form a precursor material mixture is carried out by a method selected from dry or wet planetary ball milling, rolling ball milling, high-shear grinding, air jet grinding and / or impact grinding.

31. The method according to claim 26, wherein the heat-treatment of the precursor material mixture is carried out for 1 to 14 h.

32. The method according to claim 26, wherein the heat-treatment of the precursor material mixture is carried out in a gas atmosphere selected from air, N2, Ar, He, CO2, CO, O2, H2 and mixtures thereof.

33. The method according to claim 26, wherein the method includes one or more of the following post-treatment steps: (i) mixing the active electrode material with a carbon source and optionally further heating the mixture to form a carbon coating on the active electrode material; (ii) spray-drying the active electrode material; and / or (iii) grinding the active electrode material to change the particle size of the active electrode material.

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