Patent Application
20230361288
The present invention relates to active electrode materials and to methods for their manufacture. Such materials are of interest as active electrode materials in lithium-ion or sodium-ion batteries, for example as anode materials for lithium-ion batteries.
Lithium-ion (Li-ion) batteries are a commonly used type of rechargeable battery with a global market predicted to grow to $200 bn by 2030. Li-ion batteries are the technology of choice for electric vehicles that have multiple demands across technical performance to environmental impact, providing a viable pathway for a green automotive industry.
A typical lithium-ion battery is composed of multiple cells connected in series or in parallel. Each individual cell is usually composed of an anode (negative polarity electrode) and a cathode (positive polarity electrode), separated by a porous, electrically insulating membrane (called a separator), immersed into a liquid (called an electrolyte) enabling lithium ions transport.
In most systems, the electrodes are composed of an electrochemically active material—meaning that it is able to chemically react with lithium ions to store and release them reversibly in a controlled manner—mixed if necessary with an electrically conductive additive (such as carbon) and a polymeric binder. A slurry of these components is coated as a thin film on a current collector (typically a thin foil of copper or aluminium), thus forming the electrode upon drying.
In the known Li-ion battery technology, the safety limitations of graphite anodes upon battery charging is a serious impediment to its application in high-power electronics, automotive and industry. Among a wide range of potential alternatives proposed recently, lithium titanate (LTO, particularly spinel-type Li4Ti5O12) and mixed niobium oxide-based materials are the main contenders to replace graphite as the active material of choice for high power applications.
Batteries relying on a graphitic anode are fundamentally limited in terms of charging rate. Under nominal conditions, lithium ions are inserted into the anode active material upon charging. When charging rate increases, typical graphite voltage profiles are such that there is a high risk that overpotentials lead to the potential of sites on the anode to become <0 V vs. Li/Li+, which leads to a phenomenon called lithium dendrite electroplating, whereby lithium ions instead deposit at the surface of the graphite electrode as lithium metal. This leads to irreversible loss of active lithium and hence rapid capacity fade of the cell. In some cases, these dendritic deposits can grow to such large sizes that they pierce the battery separator and lead to a short-circuit of the cell. This can trigger a catastrophic failure of the cell leading to a fire or an explosion. Accordingly, the fastest-charging batteries having graphitic anodes are limited to charging rates of 5-7 C, but often much less.
Lithium titanate (LTO) anodes do not suffer from dendrite electroplating at high charging rate thanks to their high potential (1.55 V vs. Li/Li+), and have excellent cycle life as they do not suffer from significant volume expansion of the active material upon intercalation of Li ions due to their accommodating 3D crystal structure. LTO cells are typically regarded as high safety cells for these two reasons. However, LTO is a relatively poor electronic and ionic conductor, which leads to limited capacity retention at high rate and resultant power performance, unless the material's primary particles are nanosized to increase specific surface area, and carbon-coated to increase electronic conductivity. This level of material engineering increases the porosity and specific surface area of the active material, and results in a significantly lower achievable packing density in an electrode. This is significant because it leads to low density electrodes and a higher fraction of electrochemically inactive material (e.g. binder, carbon additive), resulting in much lower gravimetric and volumetric energy densities. As such, methods that can improve the packing density such as physical mixtures of different active materials and/or particle sizes, are very attractive to improve performance.
A key measure of anode performance is the electrode volumetric capacity (mAh/cm3), that is, the amount of electric charges (that is lithium ions) that can be stored per unit volume of the anode. This is an important factor to determine the overall battery energy density on a volumetric basis (Wh/L) when combined with the cathode and appropriate cell design parameters. Electrode volumetric capacity can be approximated as the product of electrode density (g/cm3), active material specific capacity (mAh/g), and fraction of active material in the electrode. LTO anodes typically have relatively low specific capacities (c. 165 mAh/g, to be compared with c. 330 mAh/g for graphite) which, combined with their low electrode densities (typically <2.0 g/cm3) and low active material fractions (<90%) discussed above, lead to very low volumetric capacities (<300 mAh/cm3) and therefore low battery energy density and high $/kWh cost in various applications. As a result, LTO batteries/cells are generally limited to specific niche applications, despite their long cycle life, fast-charging capability, and high safety.
Mixed niobium oxides (MNO) were first identified as potential battery materials in the academic literature in the 1980's,[2,3] but have only seen a commercial focus since the 2010's with the demonstration of a practical cell combining a TiNb2O7 and a commercially-available LNMO (lithium nickel manganese oxide) cathode showing promising performance in terms of rate capability, cycle life, and energy density.[1] Selected MNO anodes offer characteristics that are similar to LTO in terms of high operating potential vs. Li/Li+ (1.55 V) and low volume expansion (<5%) leading to safe fast-charge and long cycle life (>10,000 cycles). A key advantage of MNO anodes is that practical specific capacities significantly higher than LTO (c. 165 mAh/g) can be achieved (c. 200-300 mAh/g), which improves cell energy density. In contrast to LTO materials (10−17 cm2 s−1), the Li-ion diffusion coefficient is typically much higher for specific MNO compositions that result in so-called “Wadsley-Roth” or “Tetragonal Tungsten Bronze” crystal structures (10−14-10−10 cm2 s−1).[4] This means that Li ions will diffuse across much greater distances through the active material within the same time for MNO materials vs LTO, at a fixed charge/discharge rate. Therefore, MNO materials can be less porous and use larger primary particles/crystals (0.5-10 μm for MNO vs <100 nm for LTO), retaining or improving the high-power charge/discharge performance. This results in higher electrode densities, and volumetric energy densities of cells, leading to a lower $/kWh cost at the application level.
However, the nominal voltage of MNO materials is typically higher than that of LTO (i.e. >1.55 V vs Li/Li+), which acts as a trade-off and decreases achievable energy density in full Li-ion cells. Cost of precursors, in particular Nb-based raw materials, also limits the deployment of MNO materials in commercial products for mass market applications.
US2019/0288283A1 discloses a lithium niobium composite oxide where as an essential feature some of the 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. The document refers to but does not exemplify an electrode comprising the lithium niobium composite oxide and another active material which may be any of lithium titanate having a ramsdellite structure, lithium titanate having a spinel structure, monoclinic titanium dioxide, anatase type titanium dioxide, rutile type titanium dioxide, a hollandite type titanium composite oxide, an orthorhombic titanium-containing composite oxide, and a monoclinic niobium titanium composite oxide.
US2020/0140339A1, US2018/0083283A1, and U.S. Pat. No. 10,096,826B2 disclose titanium niobate materials (based on TiNb2O7 or Ti2Nb10O29). They refer to but do not exemplify mixtures with other active materials such as different forms of titanium dioxide and lithium titanate. Titanium niobate materials exhibit a typically lower lithium ion diffusion coefficient than other MNO materials.
WO2019234248A1 discloses examples of an electrode comprising a mixture of a niobium tungsten oxide (W5Nb16O55) and LTO (Li4Ti5O12). It is believed that the properties of this electrode, in particular the properties of W5Nb16O55, can be improved.
The present invention has been devised in light of the above considerations.
In a first aspect, the invention provides an active electrode material comprising a mixture of (a) at least one lithium titanium oxide and (b) at least one mixed niobium oxide, wherein the mixed niobium oxide is expressed by the formula [M1]x[M2](1-x)[Nb]y[O]z, wherein:
The active electrode material is a mixture of a lithium titanium oxide and a mixed niobium oxide. The mixed niobium oxide has been modified by cation substitution and/or by introducing oxygen deficiency. The inventors have found that mixed niobium oxides that have been modified in this way have improved properties for use as active electrode materials, e.g. in anodes for lithium- and sodium-ion batteries. In particular, the inventors have found that the modified mixed niobium oxides have improved electronic conductivity, improved initial coulombic efficiency, and/or improved capacity retention at high charge/discharge rates, compared to the unmodified ‘base’ mixed niobium oxides. Therefore, the invention represents a way of combining the improved properties of modified mixed niobium oxides with the benefits of lithium titanium oxides for use as active electrode materials. In particular, the combination of mixed niobium oxides and lithium titanium oxides provide advantages versus the individual active materials with regards to improved cost, electrode formulation and ink processing, and various aspects of electrochemical performance.
The active electrode materials of the invention are particularly useful in electrodes, preferably for use in anodes for lithium-ion or sodium-ion batteries. Therefore, a further implementation of the invention is a composition comprising the active electrode material of the first aspect and at least one other component; optionally wherein the at least one other component is selected from a binder, a solvent, a conductive additive, an additional active electrode material, and mixtures thereof. Such a composition is useful for fabricating an electrode. A further implementation of the invention is an electrode comprising the active electrode material of the first aspect in electrical contact with a current collector. A further implementation of the invention is an electrochemical device comprising an anode, a cathode, and an electrolyte disposed between the anode and the cathode, wherein the anode comprises an active electrode material according to the first aspect; optionally wherein the electrochemical device is a lithium-ion battery or a sodium-ion battery.
In a second aspect, the invention provides method for making an active electrode material, wherein the active electrode material is as defined in the first aspect, the method comprising mixing at least one lithium titanium oxide with at least one mixed niobium oxide.
The invention includes the combination of the aspects and features described herein except where such a combination is clearly impermissible or expressly avoided.
The principles of the invention will now be discussed with reference to the accompanying figures in which:
Aspects and embodiments of the present invention will now be discussed with reference to the accompanying figures. Further aspects and embodiments will be apparent to those skilled in the art. All documents mentioned in this text are incorporated herein by reference.
The ratio by mass of (a):(b) may be in the range of 0.5:99.5 to 99.5:0.5, preferably in the range of 2 98 to 98:2. In one implementation the active electrode material comprises a higher proportion of the lithium titanium oxide than the mixed niobium oxide, e.g. the ratio by mass of (a):(b) is at least 2:1, at least 5:1, or at least 8:1. Advantageously, this allows the mixed niobium oxide to be incrementally introduced into existing electrodes based on lithium titanium oxides without requiring a large change in manufacturing techniques, providing an efficient way of improving the properties of existing electrodes. In another implementation the active electrode material comprises a higher proportion of the mixed niobium oxide than the lithium titanium oxide, e.g. such that the ratio by mass of (b):(a) is at least 2:1, at least 5:1, or at least 8:1. Advantageously, this allows for the cost of the material to be reduced by replacing some of the mixed niobium oxide with lithium titanium oxide.
Optionally, the active electrode material may consist of a mixture of (a) at least one lithium titanium oxide and (b) at least one mixed niobium oxide. Additionally, the active electrode material may consist of a mixture of (a) one lithium titanium oxide and (b) one mixed niobium oxide.
The term “mixed niobium oxide” (MNO) refers to an oxide comprising niobium and at least one other cation. MNO materials have a high redox voltage vs. Lithium (Li/Li+) >0.8V, enabling safe and long lifetime operation, crucial for fast charging battery cells. Moreover, niobium cations can have two redox reactions per atom, resulting in higher theoretical capacities than, for example, LTO.
The MNO is expressed by the formula [M1]x[M2](1-x)[Nb]y[O]z, wherein:
By ‘one or more of’, it is intended that either M1 or M2 may each represent two or more elements from their respective lists. An example of such a material is Ti0.05W0.25Mo0.70Nb12O33. Here, M1 represents Tix′Wx″ (where x′+x″=x), M2 represents Mo, x=0.3, y=12, z=33. Another example of such a material is Ti0.05Zr0.05W0.25Mo0.65Nb12O33. Here, M1 represents Tix′Zrx″Wx″ (where x′+x″+x′″=x), M2 represents Mo, x=0.35, y=12, z=33.
M2 may be selected from one or more of Mo, W, V, Zr, P, Al, Zn, Ga, Ge, Ta, Cr, Cu, K, Mg, Ni, and Hf; or one or more of Mo, W, V, Zr, P, Al, Zn, Ga, and Ge; or one or more of Mo, W, V, and Zr. Preferably, M2 consists of a single element. M2 does not represent Ti. In other words, preferably, Ti is not the major non-Nb cation in the mixed niobium oxide. Where M1 represents Ti alone, preferably x is 0.05 or less. Where M1 represents one or more cations including Ti, preferably the amount of Ti relative to the total amount of non-Nb cations is 0.05:1 or less.
M1 is a cation which substitutes for M2 in the crystal structure. M1 may be selected from one or more of P, B, Ti, Mg, V, Cr, W, Zr, Mo, Cu, Fe, Ga, Ge, K, Ni, Co, Al, Hf, Ta, and Zn; or one or more of P, B, Ti, Mg, V, Cr, W, Zr, Mo, Ga, Ge, Al, and Zn; or one or more of Ti, Zr, V, W, and Mo. M1 may have an equal or lower oxidation state than M2. Preferably, M1 has a lower oxidation state than M2. When more than one element is present as M1 and/or M2 it will 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 (the contribution from Ti)+0.75×6 (the contribution from W). Advantageously, when M1 has a lower oxidation state than M2 this is compensated for by the formation of oxygen vacancies, i.e. forming an oxygen deficient mixed niobium oxide. The presence of oxygen vacancies is believed to improve the conductivity of the mixed niobium oxide and to provide further benefits, as evidenced 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 than M2, most preferably a larger ionic radius. This gives rise to changing unit cell size and local distortions in crystal structure. This is believed to improve electrochemical properties such as specific capacity and Coulombic efficiency through altering the Li ion site availability by varying cavity size and reduction of energy barriers to reversible lithiation.
x defines the amount of M1 which replaces M2 in the mixed niobium oxide. Since x is <0.5, M2 is the major non-Nb cation in the mixed niobium oxide. Preferably, x>0. x may satisfy 0<x<0.5, 0.01<x<0.4, or 0.05≤x≤0.25.
y represents the amount of niobium in the mixed niobium oxide. z represents the amount of oxygen in the mixed niobium oxide. The precise values of y and z within the ranges defined may be selected to provide a charge-balanced structure. The precise values of y and z within the ranges defined may be selected to provide a charge balanced, or substantially charge balanced, crystal structure. Additionally or alternatively, the precise values of y and z within the ranges defined may be selected to provide a thermodynamically stable, or thermodynamically metastable, crystal structure, e.g. based on the unmodified crystal structures disclosed herein.
In some cases, z may 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 may be greater than 0, i.e. a may satisfy 0<α≤0.05. When α is greater than 0, the mixed niobium oxide is oxygen deficient, i.e. the material has oxygen vacancies. Such a material would not have precise charge balance, but is considered to be “substantially charge balanced” as indicated above. Alternatively, α may equal 0, in which case the mixed niobium oxide is not oxygen deficient. Preferably, the mixed niobium oxide is oxygen deficient. In particular, when x=0 preferably the material is oxygen deficient.
When α is 0.05, the number of oxygen vacancies is equivalent to 5% of the total oxygen in the crystal structure. In some embodiments, a may 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, a may 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, a may satisfy 0.001≤α≤0.05. When the material is oxygen deficient, the electrochemical properties of the material may be improved, for example, resistance measurements may show improved conductivity in comparison to equivalent non-oxygen deficient materials. As will be understood, the percentage values expressed here are in atomic percent.
Oxygen vacancies may be formed in a mixed niobium oxide by the sub-valent substitution of a base material. For example, oxygen vacancies may be formed by substituting some of the Mo(6+) cations in MoNb12O33 with cations of a lower oxidation state, such as Ti(4+) and/or Zr(4+) cations. A specific example of this is the compound Ti0.05Zr0.05W0.25Mo0.65Nb12O33-δ which is derived from the base material MoNb12O33 and includes oxygen vacancies. Oxygen vacancies may also be formed by heating a mixed niobium oxide under reducing conditions (for instance, heating under nitrogen atmosphere at e.g. 800-1350° C.). A specific example of this is the compound MoNb12O33-δ. The mixed niobium oxide may have induced oxygen deficiency. Induced oxygen deficiency may be understood to mean that the mixed niobium oxide contains additional oxygen vacancies, e.g. in addition to oxygen vacancies already present in the mixed niobium oxide due to sub-valent substitution of M2 with M1.
A number of methods exist for determining whether oxygen vacancies are present in a material. For example, Thermogravimetric Analysis (TGA) may be performed to measure the mass change of a material when heated in air atmosphere. A material comprising oxygen vacancies can increase in mass when heated in air due to the material “re-oxidising” and the oxygen vacancies being filled by oxide anions. The magnitude of the mass increase may be used to quantify the concentration of oxygen vacancies in the material, on the assumption that the mass increase occurs entirely due to the oxygen vacancies being filled. It should be noted that a material comprising oxygen vacancies may show an initial mass increase as the oxygen vacancies are filled, followed by a mass decrease at higher temperatures if the material undergoes thermal decomposition. Moreover, there may be overlapping mass loss and mass gain processes, meaning that some materials comprising oxygen vacancies may not show a mass gain (and sometimes not a mass loss or gain) during TGA analysis.
Other methods of determining whether oxygen vacancies are present include electron paramagnetic resonance (EPR), X-ray photoelectron spectroscopy (XPS, e.g. of oxygen 1s and/or and of cations in a mixed oxide), X-ray absorption near-edge structure (XANES, e.g. of cations in a mixed metal oxide), and TEM (e.g. scanning TEM (STEM) equipped with high-angle annular darkfield (HAADF) and annular bright-field (ABF) detectors). The presence of oxygen vacancies can be qualitatively determined by assessing the colour of a material relative to a non-oxygen-deficient sample of the same material. For example, stoichiometric MoNb12O33 has a white, off-white, or yellow colour whereas oxygen-deficient MoNb12O33-δ has a purple colour. The presence of vacancies can also be inferred from the properties, e.g. electrical conductivity, of a stoichiometric material compared to those of an oxygen-deficient material.
When the mixed niobium oxide is oxygen deficient it may be selected from MoNb12O(33-33α), WNb12O(33-33α), Mo3Nb14O(44-44α), VNb9O(25-25α), ZrNb24O(62-62α), Zn2Nb34O(87-87α), Cu2Nb34O(87-87α), WNb4O(31-31α), W9Nb8O(47-47α), W5Nb16O(55-55α), W16Nb18O(93-93α), AlNb11O(29-29α), GaNb11O(29-29α), FeNb11O(29-29α), AlNb49O(124-124α), GaNb49O(124-124α), FeNb49O(124-124α), and GeNb18O(47-47α) wherein α satisfies 0<α≤0.05. These are examples of materials where x=0 and M2 consists of a single element. Preferably when the mixed niobium oxide is oxygen deficient it is selected from MoNb12O(33-33α), WNb12O(33-33α), VNb9O(25-25α), ZrNb24O(62-62α), W5Nb16O(55-55α), W7Nb4O(31-31α), and W9Nb8O(47-47α) wherein α satisfies 0<α≤0.05.
The mixed niobium oxide may be selected from M1xMo(1-x)Nb12O(33-33α), M1xW(1-x)Nb12O(33-33α), M1xMo(1-x)Nb4.667O(14.567-14.667α) (i.e. Mo3Nb14O44 base structure), M1xV(1-x)Nb9O(25-25α), M1xZr(1-x)Nb24O(52-62α), M1xZn(1-x)Nb17O(43.5-43.5α) (i.e. Zn2Nb34O87 base structure), M1xCu(1-x)Nb17O(43.5-43.5α) (i.e. Cu2Nb34O87 base structure), M1xW(1-x)Nb0.571O(4.429-4.429α) (i.e. W7Nb4O31 base structure), M1xW(1-x)Nb0.889O(5.222-5.222α) (i.e. W9Nb8O47 base structure), M1xW(1-x)Nb3.2O(11-11α) (i.e. WNb16O55 base structure), M1xW(1-x)Nb1.125O(5.813-5.813α) (i.e. W16Nb18O93 base structure), M1xAl(1-x)Nb11O(29-29α), M1xGa(1-x)Nb11O(29-29α), M1xFe(1-x)Nb11O(29-29α), M1xAl(1-x)Nb49O(124-124α), M1xGa(1-x)Nb49O(124-124α), M1xFe(1-x)Nb49O(124-124α), and M1xGe(1-x)Nb18O(47-47α) wherein α satisfies 0≤α≤0.05 and x and/or α is >0. x is as defined above. These represent modified versions of the ‘base’ mixed niobium oxide (i.e. when x=α=0). When x>0 the oxide is modified by cation substation of M1. When α>0 the oxide is modified by oxygen deficiency. Preferably the mixed niobium oxide is selected from M1xMo(1-x)Nb12O(33-33α), M1xW(1-x)Nb12O(33-33α), M1xV(1-x)Nb9O(25-25α), M1xZr(1- x)Nb24O(62-62α), M1xZn(1-x)Nb17O(43.5-43.5α) (i.e. Zn2Nb34O87 base structure), M1xAl(1-x)Nb11O(29-29α), M1xW(1- x)Nb0.571O(4.429-4.429α) (i.e. W7Nb4O31 base structure), and M1xGe(1-x)Nb18O(47-47α) wherein α satisfies 0≤α≤0.05 and x and/or α is >0. Most preferably, the mixed niobium oxide is selected from M1xMo(1-x)Nb12O(33-33α), M1xW(1-x)Nb12O(33-33α), M1xV(1-x)Nb9O(25-25α), M1xZr(1-x)Nb24O(62-62α), and M1xW(1-x)Nb0.571O(4.429-4.429α) (i.e. W7Nb4O31 base structure) wherein α satisfies 0≤α≤0.05 and x and/or α is >0.
It will be understood that the discussion of the variables of the active electrode material is intended to be read in combination. For example, M2 may be selected from one or more of Mo, W, V, Zr, P, Al, Zn, Ga, Ge, Ta, Cr, Cu, K, Mg, Ni, and Hf and M1 may be selected from P, B, Ti, Mg, V, Cr, W, Zr, Mo, Cu, Fe, Ga, Ge, K, Ni, Co, Al, Hf, Ta, and Zn. M2 may be selected from one or more of Mo, W, V, Zr, P, Al, Zn, Ga, and Ge and M1 may be selected from one or more of P, B, Ti, Mg, V, Cr, W, Zr, Mo, Ga, Ge, Al, and Zn. M2 may be selected from one or more of Mo, W, V, and Zr and M1 may be selected from one or more of Ti, Zr, V, W, and Mo. Optionally M1 and M2 are not Fe.
In one particular example M2 is selected from one or more of Mo, W, V, Zr, P, Al, Zn, Ga, Ge, Ta, Cr, Cu, K, Mg, Ni, and Hf; M1 is selected from P, B, Ti, Mg, V, Cr, W, Zr, Mo, Cu, Fe, Ga, Ge, K, Ni, Co, Al, Hf, Ta, and Zn; x satisfies 0.01<x<0.4. In a further example M2 is selected from one or more of Mo, W, V, Zr, P, Al, Zn, Ga, and Ge; M1 is selected from one or more of P, B, Ti, Mg, V, Cr, W, Zr, Mo, Ga, Ge, Al, and Zn; and 0.05≤x≤0.25. In a further example the mixed niobium oxide is selected from M1xMo(1-x)Nb12O(33-33α) and M1XW(1-x)Nb0.571O(4.429-4.429α) (i.e. W7Nb4O31 base structure) wherein α satisfies 0≤α≤0.05, x satisfies 0.01<x<0.4, wherein M1 is selected from Ti, Zr, V, W, and Mo.
The mixed niobium oxide may have a ReO3-derived MO3-x crystal structure. Preferably, the mixed niobium oxide has a Wadsley-Roth or Tetragonal Tungsten Bronze (“TTB” or “bronze”) crystal structure. Both Wadsley-Roth and bronze crystal structures are considered to be a crystallographic off-stoichiometry of the MO3 (ReO3) crystal structure, with simplified formula of MO3-x. As a result, these structures typically contain [MOB] octahedral subunits in their crystal structure alongside others. Mixed niobium oxides with these structures are believed to have advantageous properties for use as active electrode materials, e.g. in lithium-ion batteries.
The open tunnel-like MO3 crystal structure of MNOs also makes them ideal candidates for high capacity and high rate intercalation. The crystallographic off-stoichiometry that is introduced in MO3-x structures causes crystallographic superstructures such as the Wadsley-Roth shear and the Bronze structures. These superstructures, compounded by other qualities such as the Jahn-Teller effect and crystallographic disorder by making use of multiple mixed cations, stabilise the crystal and keep the tunnels open and stable during intercalation, enabling extremely high rate performance.
The crystal formula of a charge balanced and thermodynamically stable Wadsley-Roth crystal structure obeys the following formula:
(M1,M2,M3, . . . )mnp+1O3mnp−(m+n)p+4 (1)
In this formula, O is oxygen (the anion) and M (the cation) can be any alkali metal, alkali earth metal, transition element, semi-metal, or non-metal if the correct proportions are used to provide a stable structure. In the MNO, at least one of (M1, M2, M3 . . . ) comprises Nb.
Formula (1) is based on crystal topography: m and n are the dimensions of the formed edge sharing superstructure blocks, ranging from 3-5 (integers). At the corner, blocks are connected into infinite ribbons (p=∞) only by edge-sharing, into pairs (p=2) by partly edge-sharing and partly tetrahedra or into isolated blocks only by tetrahedra (p=1). When p is infinity the formula becomes:
(M1,M2,M3, . . . )mnO3mn−(m+n) (2)
More information can be found in work by Griffith et al.[5]
Together, formula (1) and (2) define the full composition range for Wadsley-Roth crystal structures. The total crystal composition should also be charge neutral and thermodynamically favourable to follow the above description. Structures partially deficient in their oxygen content through introduction of oxygen vacancy defects are preferable when reducing the material's electrical resistance such that MxOy becomes MxOγ-δ where 0%<δ≤5%, i.e. the oxygen content is reduced by up to 5 atomic % relative to the amount of oxygen present.
Tetragonal tungsten bronze crystal structures are phases formed of a framework of [MO6] octahedra sharing corners linked in such a way that three, four and five sided tunnels are formed (Montemayor et al.,[6] e.g. M8W9O47). A bronze structure does not have to include tungsten[7]. A number of 5-sided tunnels are filled with (M1, M2, M3 . . . ), 0, or a suitable cation to form the pentagonal columns. In the structure the pentagonal bipyramid MO7 shares edge with five MO6 octahedra. In the MNO, at least one of (M1, M2, M3 . . . ) comprises Nb. Structures partially deficient in their oxygen content through introduction of oxygen vacancy defects are preferable when reducing the materials electrical resistance such that MxOy becomes MxOy-δ where 0%≤δ≤5%, i.e. the oxygen content is reduced by up to 5 atomic % relative to the amount of oxygen present.
The crystal structure of a material may be determined by analysis of X-ray diffraction (XRD) patterns, as is widely known. For instance, XRD patterns obtained from a given material can be compared to known XRD patterns to confirm the crystal structure, e.g. via public databases such as the ICDD (JCPDS) crystallography database. Rietveld analysis can also be used to determine the crystal structure. Therefore, the mixed niobium oxide may have a Wadsley-Roth or Tetragonal Tungsten Bronze crystal structure, as determined by X-ray diffraction.
Optionally, the crystal structure of the mixed niobium oxide, as determined by X-ray diffraction, corresponds to the crystal structure of the unmodified form of the mixed niobium oxide, wherein the unmodified form is expressed by the formula [M2][Nb]y[O]z wherein M2 consists of a single element and wherein the unmodified form is not oxygen deficient, wherein the unmodified form is selected from one or more of: M2INb5O13, M2IBNb10.8O30, M2IINb2O6, M2II2Nb34O37, M2IIINb11O29, M2IIINb49O124 (M2III0.5Nb24.5O62), M2IVNb24O62, M2IVNb2O7, M2IV2Nb10O29, M2IV2Nb14O39, M2IVNb14O37, M2IVNb6O17, M2IVNb18O47, M2VNb9O25, M2V4Nb18O55, M2V3Nb17O50, M2VINb12O33, M2VI4Nb23O77, M2VI3Nb14O44, M2VI5Nb16O55, M2VI8Nb18O59, M2VINb2O3, M2VINb18O93, M2VI20Nb22O115, M2VI9Nb8O47, M2VI82Nb54O381, M2VI31Nb20O143, M2VI7Nb4O31, M2VI15Nb2O50, M2VI3Nb2O14, and M2VINb12O63, wherein the numerals I, II, III, IV, V, and VI represent the oxidation state of M2. In this way, it can be confirmed that the unmodified mixed niobium oxide has been modified without significantly affecting the crystal structure. Preferably the crystal structure of the mixed niobium oxide, as determined by X-ray diffraction, corresponds to the crystal structure of one or more of M2II2Nb34O87, M2IIINb11O29, M2IIINb49O124, M2IVNb24O62, M2IVNb18O47, M2VNb9O25, M2VINb12O33, M2VI7Nb4O31, M2VI9Nb8O47, M2VI5Nb16O55, and M2VI16Nb18O93.
The crystal structure of the mixed niobium oxide, as determined by X-ray diffraction, may correspond to the crystal structure of one or more of: MoNb12O33, WNb12O33, Mo3Nb14O44, VNb9O25, ZrNb24O62, Zn2Nb34O7, Cu2Nb34O87, W7Nb4O31, W9Nb8O47, W5Nb16O55, W16Nb18093, AlNb11O29, GaNb11O29, FeNb11O29, AlNb4O124, GaNb49O124, FeNb49O124, and GeNb13O47. Preferably, the crystal structure of the mixed niobium oxide, as determined by X-ray diffraction, corresponds to the crystal structure of one or more of MoNb12O33, WNb12O33, ZrNb24O62, Zn2Nb34O87, VNb9O25, W5Nb16O5, AlNb11O29, GeNb18O47, W7Nb4O31, and W9Nb8O47. Most preferably the crystal structure of the mixed niobium oxide, as determined by X-ray diffraction, corresponds to the crystal structure of one or more of MoNb12O33, WNb12O33, ZrNb24O62, VNb9O25, and W7Nb4O31.
Here the term ‘corresponds’ is intended to reflect that peaks in an X-ray diffraction pattern may be shifted by no more than 0.5 degrees (preferably shifted by no more than 0.25 degrees, more preferably shifted by no more than 0.1 degrees) from corresponding peaks in an X-ray diffraction pattern of the material listed above (e.g. MVINb12O33 where MVI=Mo etc.). This comparison may be performed with respect to the strongest peaks in the pattern, for example the three strongest peaks. Optionally, the crystal structure of the mixed niobium oxide does not correspond to the crystal structure of TiNb2O7, for example, optionally the measured XRD diffraction pattern of the mixed niobium oxide does not correspond to the JCPDS crystallography database entry database 00-039-1407, for TiNb2O7 Optionally, the crystal structure of the mixed niobium oxide does not correspond to the crystal structure of Ti2Nb10O29. Optionally, the crystal structure of the mixed niobium oxides does not correspond to the crystal structure of MIIINb11O29 for example FeNb11O29, GaNb11O29, CrNb11O29, and AlNb11O29.
The mixed niobium oxide and/or the lithium titanium oxide may further comprise Li and/or Na. For example, Li and/or Na may enter the crystal structures when the active electrode material is used in a metal-ion battery electrode.
The mixed niobium oxide may have a lithium diffusion rate of greater than 10−14 cm2 s−1. It may be advantageous to provide materials having a suitably high lithium diffusion rate, as this can provide improved performance in an electrochemical device comprising the active electrode material. For example, the lithium diffusion rate may be determined by cyclic voltammetry.
The specific capacity of the active electrode material may be 162 mAh/g or more. Here, specific capacity is defined as that measured in the 2nd cycle of a half cell galvanostatic cycling test at a rate of 0.1 C with a voltage window of 1.1-3.0V vs Li/Li+. It may be advantageous to provide materials having a high specific capacity, as this can provide improved performance in an electrochemical device comprising the active electrode material. The specific capacity may be targeted to a certain value by varying the proportion of the mixed niobium oxide and the lithium titanium oxide. Values of above 200 mAh/g can be achieved by using a high proportion of mixed niobium oxide, as shown by the present examples.
The mixed niobium oxide is preferably in particulate form. The mixed niobium oxide may have a Doo particle diameter in the range of 0.1-100 μm, or 0.5-50 μm, or 1-25 μm. These particle sizes are advantageous because they are easy to process and fabricate into electrodes. Moreover, these particle sizes avoid the need to use complex and/or expensive methods for providing nanosized particles. Nanosized particles (e.g. particles having a D50 particle diameter of 100 nm or less) are typically more complex to synthesise and require additional safety considerations.
The mixed niobium oxide may have a D10 particle diameter of at least 0.05 μm, or at least 0.1 μm, or at least 0.5 μm, or at least 1 μm. By maintaining a D10 particle diameter within these ranges, the potential for parasitic reactions in a Li ion cell is reduced from having reduced surface area, and it is easier to process with less binder in the electrode slurry.
The mixed niobium oxide may have a D50 particle diameter of no more than 200 μm, no more than 100 μm, no more than 50 μm, or no more than 30 μm. By maintaining a D50 particle diameter within these ranges, the proportion of the particle size distribution with large particle sizes is minimised, making the material easier to manufacture into a homogenous electrode.
The term “particle diameter” refers to the equivalent spherical diameter (esd), i.e. the diameter of a sphere having the same volume as a given particle, where the particle volume is understood to include the volume of any intra-particle pores. The terms “Dn” and “Dn particle diameter” refer to the diameter below which n % by volume of the particle population is found, i.e. the terms “D50” and “D50 particle diameter” refer to the volume-based median particle diameter below which 50% by volume of the particle population is found. Where a material comprises primary crystallites agglomerated into secondary particles, it will be understood that the particle diameter refers to the diameter of the secondary particles. Particle diameters can be determined by laser diffraction. For example, particle diameters can be determined in accordance with ISO 13320:2009.
The lithium titanium oxide is in preferably in particulate form. The lithium titanium oxide may have a D50 particle diameter in the range of 0.1-50 μm, or 0.25-20 μm, or 0.5-15 μm. The lithium titanium oxide may have a D10 particle diameter of at least 0.01 μm, or at least 0.1 μm, or at least 0.5 μm. The lithium titanium oxide may have a D90 particle diameter of no more than 100 μm, no more than 50 μm, or no more than 25 μm. By maintaining a D90 particle diameter in this range the packing of lithium titanium oxide particles in the mixture with mixed niobium oxide particles is improved.
Lithium titanium oxides are typically used in battery anodes at small particle sizes due to the low electronic conductivity of the material. In contrast, the mixed niobium oxide as defined herein may be used at larger particle sizes since it typically has a higher lithium ion diffusion coefficient than lithium titanium oxide. Advantageously, in the active electrode material the lithium titanium oxide may have a smaller particle size than the mixed niobium oxide, for example such that the ratio of the D50 particle diameter of the lithium titanium oxide to the D50 particle diameter of the mixed niobium oxide is in the range of 0.01:1 to 0.9:1, or 0.1:1 to 0.7:1. In this way, the smaller lithium titanium oxide particles may be accommodated in the voids between the larger mixed niobium oxide particles, increasing the packing efficiency of the active electrode material.
The mixed niobium oxide may have a BET surface area in the range of 0.1-100 m2/g, or 0.5-50 m2/g, or 1-20 m2/g. The lithium titanium oxide may have a BET surface area in the range of 0.1-100 m2/g, or 1-50 m2/g, or 3-30 m2/g. In general, a low BET surface area is preferred in order to minimise the reaction of the active electrode material with the electrolyte, e.g. minimising the formation of solid electrolyte interphase (SEI) layers during the first charge-discharge cycle of an electrode comprising the material. However, a BET surface area which is too low results in unacceptably low charging rate and capacity due to the inaccessibility of the bulk of the active electrode material to metal ions in the surrounding electrolyte. The the ratio of the BET surface area of the lithium titanium oxide to the BET surface area of the mixed niobium oxide is in the range of 1.1:1 to 20:1, or 1.5:1 to 10:1.
The term “BET surface area” refers to the surface area per unit mass calculated from a measurement of the physical adsorption of gas molecules on a solid surface, using the Brunauer-Emmett-Teller theory. For example, BET surface areas can be determined in accordance with ISO 9277:2010.
The mixed niobium oxide may comprise a carbon coating. The coating may be present in an amount of up to 10 wt %, or 0.05-5 wt %, or 0.1-3 wt %, based on the total weight of the mixed niobium oxide and the coating. It has been found that a carbon precursor comprising polyaromatic sp2 carbon provides a particularly beneficial carbon coating on mixed niobium oxides. Preferably the carbon coating comprises polyaromatic sp2 carbon. Such a coating is formed by pyrolysing a carbon precursor comprising polyaromatic sp2 carbon since the sp2 hybridisation is largely retained during pyrolysis. Typically, pyrolysis of a polyaromatic sp2 carbon precursor under reducing conditions results in the domains of sp2 aromatic carbon increasing in size. Accordingly, the presence of a carbon coating comprising polyaromatic sp2 may be established via knowledge of the precursor used to make the coating. The carbon coating may be defined as a carbon coating formed from pyrolysis of a carbon precursor comprising polyaromatic sp2 carbon. Preferably, the carbon coating is derived from pitch carbons.
The presence of a carbon coating comprising polyaromatic sp2 carbon may also be established by routine spectroscopic techniques. For instance, Raman spectroscopy provides characteristic peaks (most observed in the region 1,000-3,500 cm1) which can be used to identify the presence of different forms of carbon. A highly crystalline sample of sp3 carbon (e.g. diamond) provides a narrow characteristic peak at ˜1332 cm−1. Polyaromatic sp2 carbon typically provides characteristic D, G, and 2D peaks. The relative intensity of D and G peaks (ID/IG) can provide information on the relative proportion of sp2 to sp3 carbon. The mixed niobium oxide may have an ID/IG ratio as observed by Raman spectroscopy within the range of 0.85-1.15, or 0.90-1.10, or 0.95-1.05.
X-ray diffraction may also be used to provide information on the type of carbon coating. For example, an XRD pattern of a mixed niobium oxide with a carbon coating may be compared to an XRD pattern of the uncoated mixed niobium oxide and/or to an XRD pattern of a pyrolysed sample of the carbon precursor used to make the carbon coating.
The carbon coating may be semi-crystalline. For example, the carbon coating may provide a peak in an XRD pattern of the mixed niobium oxide centred at 2θ of about 26° with a width (full width at half maximum) of at least 0.20°, or at least 0.25°, or at least 0.30°.
The lithium titanium oxide preferably has a spinel or ramsdellite crystal structure, e.g. as determined by X-ray diffraction. An example of a lithium titanium oxide having a spinel crystal structure is Li4Ti5O12. An example of a lithium titanium oxide having a ramsdellite crystal structure is Li2Ti3O7. These materials have been shown to have good properties for use as active electrode materials. Therefore, the lithium titanium oxide may have a crystal structure as determined by X-ray diffraction corresponding to Li4Ti5O12 and/or Li2Ti3O7. The lithium titanium oxide may be selected from Li4Ti5O12, Li2Ti3O7, and mixtures thereof.
The lithium titanium oxide may be doped with additional cations or anions. The lithium titanium oxide may be oxygen deficient. The lithium titanium oxide may comprise a coating, optionally wherein the coating is selected from carbon, polymers, metals, metal oxides, metalloids, phosphates, and fluorides.
The lithium titanium oxide may be synthesises by conventional ceramic techniques, for example solid-state synthesis or sol-gel synthesis. Alternatively, the lithium titanium oxide may be obtained from a commercial supplier.
A method of making a mixed niobium oxide for use in the invention comprises the steps of: providing one or more precursor materials; mixing said precursor materials to form a precursor material mixture; and heat treating the precursor material mixture in a temperature range from 400° C.-1350° C. to form the mixed niobium oxide.
The one or more precursor materials may include an M1 source, an M2 source, and a source of Nb. It will be understood that the sources may be contaminated by impurities. For example, Ta is a typical impurity present in sources of Nb which may thus be present in a mixed niobium oxide.
The phrase ‘M1 source’ is used herein to describe a material comprising M1 ions/atoms. The phrase ‘M2 source’ is used herein to describe a material comprising M2 ions/atoms. The phrase ‘a source of Nb’ is used herein to describe a material comprising Nb ions/atoms, as appropriate.
The precursor materials may include one or more metal oxides, metal hydroxides, metal salts or oxalates. For example, the precursor materials may include one or more metal oxides of different oxidation states and/or of different crystal structure. 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 materials may not comprise a metal oxide, or may comprise ion sources other than oxides. For example, the precursor materials may comprise metal salts (e.g. NO3−, SO3−) or other compounds (e.g. oxalates).
Some or all of the precursor materials may be particulate materials. Where they are particulate materials, preferably they have D50 particle diameter of <20 μm in diameter. The D50 particle diameter may be in a range from e.g. 10 nm to 20 μm. Providing particulate materials with such a particle size can help to promote more intimate mixing of precursor materials, thereby resulting in more efficient solid-state reaction during the heat treatment step. However, it is not essential that the precursor materials have an initial D50 particle diameter of <20 μm, as the particle size of the one or more precursor materials may be mechanically reduced during the step of mixing said precursor materials to form a precursor material mixture.
The step of mixing/milling the precursor materials to form a precursor material mixture may be performed by a process selected from (but not limited to): dry or wet planetary ball milling, rolling ball milling, high shear milling, airjet milling, and/or impact milling. The force used for mixing/milling may depend on the morphology of the precursor materials. For example, where some or all of the precursor materials have larger particle sizes (e.g. a D50 particle diameter of greater than 20 μm), the milling force may be selected to reduce the particle size of the precursor materials such that the such that the D50 particle diameter of the precursor material mixture is reduced to 20 μm or lower. When the D50 particle diameter of 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.
The step of heat treating the precursor material mixture may be performed for a time of from 1 hour to 24 hours, more preferably from 3 hours to 14 hours. For example, the heat treatment step may be performed for 1 hour or more, 2 hours or more, 3 hours or more, 6 hours or more, or 12 hours or more. The heat treatment step may be performed for 24 hours or less, 18 hours or less, 14 hours or less, or 12 hours or less.
In some methods it may be beneficial to perform a two-step heat treatment. For example, the precursor material mixture may be heated at a first temperature for a first length of time, follow by heating at a second temperature for a second length of time. The second temperature may be higher than the first temperature. Performing such a two-step heat treatment may assist the solid state reaction to form the desired crystal structure.
The step of heat treating the precursor material mixture may be performed in a gaseous atmosphere. The gaseous atmosphere may be an inert atmosphere, or may be a reducing atmosphere. Where it is desired to make an oxygen-deficient material, preferably the step of heat treating the precursor material mixture is performed in an inert or reducing atmosphere. Suitable gaseous atmospheres comprise: air, N2, Ar, He, CO2, CO, O2, H2, and mixtures thereof.
The method may include one or more post-processing steps after formation of the mixed niobium oxide.
In some cases, the method may include a post-processing step of heat treating the mixed niobium oxide, sometimes referred to as ‘annealing’. This post-processing heat treatment step may be performed in a different gaseous atmosphere to the step of heat treating the precursor material mixture to form the mixed niobium oxide. The post-processing heat treatment step may be performed in an inert or reducing gaseous atmosphere. Such a post-processing heat treatment step may be performed at temperatures of above 500° C., for example at about 900° C. Inclusion of a post-processing heat treatment step may be beneficial to e.g. form deficiencies or defects in the mixed niobium oxide, for example to form oxygen deficiencies.
In some cases, the method may include a post-processing step of mixing the mixed niobium oxide with a carbon source, and thereby forming a carbon coating on the mixed niobium oxide. Optionally, the mixture of the mixed niobium oxide and the carbon source may be heated to thereby form the carbon coating on the mixed niobium oxide. Suitable carbon sources include but are not limited to: carbohydrate materials (e.g. sugars, polymers); conductive carbons (e.g. carbon black); and/or aromatic carbon materials (e.g. pitch carbon).
One method of forming a carbon coating includes a step of milling the mixed niobium oxide with a carbon source, followed by pyrolysis of the mixed niobium oxide and carbon source (e.g. in a furnace) under an inert or reducing atmosphere.
Another preferred method of forming a carbon coating includes mixing of the mixed niobium oxide with a carbon source, dispersion of the mixed niobium oxide and carbon source in an aqueous slurry, followed by spray drying. The resulting powder may optionally be pyrolysed. Where the carbon source is e.g. conductive carbon black, it is not necessary to pyrolyse the material post spray-drying.
In some cases, the method may include a post-processing step of milling the mixed niobium oxide to modify the mixed niobium oxide particle size. For example, the mixed niobium oxide may be treated by one or more processes including air jet milling, impact milling, high shear milling, sieving, or ball milling. This may provide a more suitable particle size for use in desired applications of the mixed niobium oxide.
It has been found that a carbon precursor comprising polyaromatic sp2 carbon provides a particularly beneficial carbon coating on mixed niobium oxides for use in the invention. Therefore, a method of making a coated mixed niobium oxide may comprise the steps of: combining a mixed niobium oxide with a carbon precursor comprising polyaromatic sp2 carbon to form an intermediate material; and heating the intermediate material under reducing conditions to pyrolyse the carbon precursor forming a carbon coating on the mixed niobium oxide and introducing oxygen vacancies into the mixed niobium oxide.
The intermediate material may comprise the carbon precursor in an amount of up to 25 wt %, or 0.1-15 wt %, or 0.2-8 wt %, based on the total weight of the mixed niobium oxide and the carbon precursor. The carbon coating on the mixed niobium oxide may be present in an amount of up to 10 wt %, or 0.05-5 wt %, or 0.1-3 wt %, based on the total weight of the mixed niobium oxide and coating. These amounts of the carbon precursor and/or carbon coating provide a good balance between improving the electronic conductivity by the carbon coating without overly reducing the capacity of the mixed niobium oxide by overly reducing the proportion of the mixed niobium oxide. The mass of carbon precursor lost during pyrolysis may be in the range of 30-70 wt %.
The step of heating the intermediate material under reducing conditions may be performed at a temperature in the range of 400-1,200° C., or 500-1,100° C., or 600-900° C. The step of heating the intermediate material under reducing conditions may be performed for a duration within the range of 30 minutes to 12 hours, 1-9 hours, or 2-6 hours.
The step of heating the intermediate material under reducing conditions may be performed under an inert gas such as nitrogen, helium, argon; or may be performed under a mixture of an inert gas and hydrogen; or may be performed under vacuum.
The carbon precursor comprising polyaromatic sp2 carbon may be selected from pitch carbons, graphene oxide, graphene, and mixtures thereof. Preferably, the carbon precursor comprising polyaromatic sp2 carbon is selected from pitch carbons, graphene oxide, and mixtures thereof. Most preferably, the carbon precursor comprising polyaromatic sp2 carbon is selected from pitch carbons. The pitch carbons may be selected from coal tar pitch, petroleum pitch, mesophase pitch, wood tar pitch, isotropic pitch, bitumen, and mixtures thereof.
Pitch carbon is a mixture of aromatic hydrocarbons of different molecular weights. Pitch carbon is a low cost by-product from petroleum refineries and is widely available. The use of pitch carbon is advantageous because pitch has a low content of oxygen. Therefore, in combination with heating the intermediate material under reducing conditions, the use of pitch favours the formation of oxygen vacancies in the mixed niobium oxide.
Other carbon precursors typically contain substantial amounts of oxygen. For example, carbohydrates such as glucose and sucrose are often used as carbon precursors. These have the empirical formula Cm(H2O)n and thus contain a significant amount of covalently-bonded oxygen (e.g. sucrose has the formula C12H22O11 and is about 42 wt % oxygen). In some instances the pyrolysis of carbon precursors which contain substantial amounts of oxygen may prevent or inhibit reduction of a mixed niobium oxide, or even lead to oxidation, meaning that oxygen vacancies may not be introduced into the mixed niobium oxide. Accordingly, the carbon precursor may have an oxygen content of less than 10 wt %, preferably less than 5 wt %.
The carbon precursor may be substantially free of sp3 carbon. For example, the carbon precursor may comprise less than 10 wt % sources of sp3 carbon, preferably less than 5 wt % sources of sp3 carbon. Carbohydrates are sources of sp3 carbon. The carbon precursor may be free of carbohydrates. It will be understood that some carbon precursors used may contain impurities of sp3 carbon, for example up to 3 wt %.
The invention also provides a composition comprising the active electrode material of the first aspect of the invention and at least one other component, optionally wherein the at least one other component is selected from a binder, a solvent, a conductive additive, an additional active electrode material, and mixtures thereof. Such a composition is useful for preparing an electrode, e.g. an anode for a lithium-ion battery.
The invention also provides an electrode comprising the active electrode material of the first aspect of the invention in electrical contact with a current collector. The electrode may form part of a cell. The electrode may form an anode as part of a lithium-ion battery. Preferably, the active electrode material is in the form of an active layer on the current collector, wherein the active layer has a density of 2.00-3.75 g cm−3.. It may be advantageous to provide materials having such an electrode density, as this can provide improved performance in an electrochemical device comprising the active electrode material. Specifically, when the electrode density is high, high volumetric capacities can be achieved, as gravimetric capacity x electrode density x active electrode material fraction=volumetric capacity.
The invention also provides the use of the active electrode material of the first aspect of the invention in an anode for a metal-ion battery, optionally wherein the metal-ion battery is a lithium-ion battery.
A further implementation of the invention is an electrochemical device comprising an anode, a cathode, and an electrolyte disposed between the anode and the cathode, wherein the anode comprises an active electrode material according to the first aspect of the invention; optionally wherein the electrochemical device is a lithium-ion battery or a sodium-ion battery. Preferably, the electrochemical device is a lithium-ion battery having a reversible anode active material specific capacity of greater than 165 mAh/g at 20 mA/g, wherein the battery can be charged and discharged at current densities relative to the anode active material of 200 mA/g or more, or 1000 mA/g or more, or 2000 mA/g or more, or 4000 mA/g or more whilst retaining greater than 70% of the initial cell capacity at 20 mA/g. It has been found that use of the active electrode materials of the first aspect of the invention can enable the production of a lithium-ion battery with this combination of properties, representing a lithium-ion battery that is particularly suitable for use in applications where high charge and discharge current densities are desired. Notably, the examples have shown that active electrode materials according to the first aspect of the invention have excellent capacity retention at high C-rates.
Preferably, the electrochemical device is a lithium-ion battery cell. The anode active material mixture in the cell preferably having an initial coulombic efficiency greater than 88% or greater than 90%. Initial coulombic efficiency has been measured as the difference in the lithiation and de-lithiation capacity on the 1st charge/discharge cycle at C/10 in a half-cell. It may be advantageous to provide materials having a suitably high initial coulombic efficiency, as this can provide improved performance in an electrochemical device comprising the active electrode material.
In the second aspect, the invention provides a method for making an active electrode material, wherein the active electrode material is as defined in the first aspect, the method comprising mixing at least one lithium titanium oxide with at least one mixed niobium oxide. The lithium titanium oxide and the mixed niobium oxide are as defined above.
The step of mixing at least one lithium titanium oxide with at least one mixed niobium oxide may comprise low to high energy powder mixing/blending techniques, such as rotational mixing in multiple directions, rotational V-type blending over a single axis, planetary mixing, centrifugal planetary mixing, and high shear mixing.
The step of mixing at least one lithium titanium oxide with at least one mixed niobium oxide may comprise mixing in a carrier solvent.
Prior to mixing, the method may include the step of milling and/or classifying the lithium titanium oxide, e.g. to provide any of the particle size parameters given above. Prior to mixing, the method may include the step of milling and/or classifying the mixed niobium oxide, e.g. to provide any of the particle size parameters given above. The method may include a step of milling and/or classifying the mixture of the lithium titanium oxide and the mixed niobium oxide. The milling and/or classifying may be performed by impact milling or jet milling.
Optionally, the method for making an active electrode material, wherein the active electrode material is as defined in the first aspect, comprises the steps of: combining a mixed niobium oxide with a carbon precursor comprising polyaromatic sp2 carbon to form an intermediate material; heating the intermediate material under reducing conditions to pyrolyse the carbon precursor forming a carbon coating on the mixed niobium oxide and introducing oxygen vacancies into the mixed niobium oxide; and mixing at least one lithium titanium oxide with the coated mixed niobium oxide.
The following reference examples demonstrate the improvement in properties of a modified mixed niobium oxide (i.e. a cation substituted and/or oxygen deficient oxide) compared to the unmodified ‘base’ mixed niobium oxide. The reference examples test the oxides as the sole active electrode material. It would be expected that the same improvements would be seen when the oxides are tested in combination with a lithium titanium oxide in accordance with the invention, i.e. that a mixture of a modified mixed niobium oxide and a lithium titanium oxide will have improved properties for use as an active electrode material compared to a mixture of the unmodified ‘base’ mixed niobium oxide and the lithium titanium oxide.
A number of different materials were prepared and characterised, as summarised in Table 1, below. Broadly, these samples can be split into a number of groups:
Samples 1, 2, 3, 4, 5, 14, 15, 16, 18, and 22 belong to the same family of Wadsley-Roth phases based on MoNb12O33. Sample 1 is the base crystal structure, which is modified to a mixed metal cation structure by exchanging one or multiple cations in samples 2 to 4, and/or in a mixed crystal configuration (blending with isostructural WNb12O33) in samples 14, 15, 16, 18, and 22. Oxygen deficiencies are created in the base crystal in sample 5 and in the mixed metal cation structure 18. Sample 3 is a spray-dried and carbon-coated version of the crystal made in sample 2, and sample 22 is a spray-dried and carbon-coated version of the crystal made in sample 16.
Samples 6, 7, 17, 19, 20 belong to the same family of Wadsley-Roth phases based on ZrNb24O62 (M4+Nb24O62, 3×4 block of octahedra with half a tetrahedron at each block corner).
Samples 8, 9 and R11 belong to the same family of Wadsley-Roth phases based on WNb12O33 (M6+Nb12O33, a 3×4 NbO6 octahedra block with a tetrahedron at each block corner).
Samples 10, 11 and 21 belong to the same family of Wadsley-Roth phases based on VNb9025 (M5+Nb9O25, a 3×3 NbOB octahedra block with a tetrahedron at each block corner).
Samples 12, 13 and R14 belong to the same family of tungsten tetragonal bronzes (TTB) based on W7Nb4O31 (M6*7Nb4O31). This is a tetragonal tungsten bronze structure, where MO6 (M=0.4 Nb+0.6 W) octahedra are exclusively corner-sharing, with 3, 4, and 5-sided tunnels. Some of these tunnels are filled with —O-M-O— chains whereas others are open for lithium ion transport and storage.
Samples R1, R2, R13 belong to the same family of Wadsley-Roth phases based on Zn2Nb34O87 (M2+2Nb34O87). This orthorhombic phase consists out of 3×4 blocks of MOB octahedra (M=Zn+2/Nb+5), where the blocks are connected exclusively by edge-sharing and have no tetrahedra.
Samples R3, R4, R5, R12 belong to the same family of Wadsley-Roth phases based on AlNb11O29 (M3+Nb11O29). The structure belongs to monoclinic shear structure with 3×4 octahedra blocks connected through exclusively edge-sharing and have no tetrahedra.
Samples R6, R7, R8 belong to the same family of Wadsley-Roth phases based on GeNb18O47 (M4+Nb13O47). The structure is similar to sample 10 with 3×3 NbO6 octahedra blocks and one tetrahedron connecting blocks at corners. However, the structure contains intrinsic defects due to Ge+4 instead of V5+.
Samples R9, R10 belong to the same family of Wadsley-Roth phases based on W5Nb16O55 (M6+5Nb16O55). The structure is made of 4×5 blocks connected at the sides by edge-sharing (W,Nb)O6 and connected at the corners by WO4 tetrahedra. This structure is similar to Sample 8 and 9 but with a larger block size.
Samples listed in Table 1 were synthesised using a solid-state route. In a 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 stochiometric proportions and planetary ball-milled at 550 rpm for 3h in a zirconia jar and milling media with a ball to powder ratio of 10:1. The resulting powders were then heated in a static muffle furnace in air in order to form the desired crystal phase. Samples 1 to 5 and 12 to 16, 18 and 22 were heat-treated at 90000 for 12h; samples 6 to 9, 17, 19, and 20 were heat-treated at 1200° C. for 12h, with samples 6, 7, 17, 19 and 20 undergoing a further heat treatment step at 1350° C. for an additional 4h; samples 10, 11 and 21 were heat-treated at 1000° C. for 12h. Sample 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 concentrations of 5, 10, 15, or 20 w/w % with ionic surfactant, and spray-dried in a lab-scale spray-drier (inlet temperature 220° C., outlet temperature 95° C., 500 mL/h sample introduction rate). The resulting powder was pyrolyzed at 600° C. for 5h in nitrogen. Sample 5 and 18 were further annealed in nitrogen at 900° C. for 4 hours.
Samples R1, R2, R6, R7, R8, R9, R10 were prepared by ball milling as above, and impact milling at 20,000 rpm as needed to a particle size distribution with D90<20 μm, then heat-treated as in a muffle furnace in air at 1200° C. for 12 h; samples R8, R10, R11, R12, R13 were further annealed in nitrogen at 1000° C. for 4 h; R14 was annealed in nitrogen at 900° C. for 5 h. Samples R3, R4, R5 were heat-treated at 1300° C. for 12 h. Samples R1-R10 were de-agglomerated after synthesis by impact milling or jet milling to the desired particle size ranges.
The phase purity of some samples was analysed using Rigaku Miniflex powder X-ray diffractometer in 2θ range (10-70°) at 1°/min scan rate.
Thermogravimetric Analysis (TGA) was performed on some samples using a Perkin Elmer Pyris 1 system in a synthetic air atmosphere. Samples were first held for 15 min at 30° C., then heated from 30° C. to 950° C. at 5° C./min, and finally held for 30 min at 950° C. TGA was performed on sample 3 to quantify carbon content.
As discussed above, sample 5 and 18 were heat-treated at 900° C. for 12h to form the mixed niobium oxide, and was then further annealed in nitrogen (a reducing atmosphere) at 900° C., in a post-processing heat treatment step. A colour change from white to dark purple was observed after the post-processing heat treatment in nitrogen, indicating change in oxidation states and band structure of the material, as a result of oxygen deficiency of the sample.
Samples R8, R10, R11, R12, R13 were further annealed in nitrogen at 1000° C. for 4 h, sample R14 was annealed in nitrogen at 900° C. for 5 h. Sample R7 transitions from a white colour to a deep yellow colour upon introduction of induced oxygen deficiencies in sample R8; sample R9 transitions from an off-white colour to a blue-grey colour upon introduction of induced oxygen deficiencies in sample R10; sample 8 transitions from off-white to light blue in R11; sample R3 transitions from white to grey/black in R12; sample R1 transitions from white to grey/black in R3; sample 12 transitions from light yellow to dark blue in R14.
Particle Size Distributions were obtained with a Horiba laser diffraction particle analyser for dry powder. Air pressure was kept at 0.3 MPa. The results are set out in Table 2, below.
All particle size distributions can also be refined with further processing steps, for example 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 shown in
The morphology of some samples was analysed by Scanning Electron Microscopy (SEM).
It can be seen that the material has with homogeneous porous particles that can pack efficiently to form a high-density electrode. Qualitatively the conductivity is vastly improved as a conductive coating does not need to be applied for SEM imaging to be carried out, implying an order of magnitude improvement in material surface conductivity.
Electrochemical tests were carried out in half-coin cells (CR2032 size) for initial analysis. In half-coin tests, the material is tested in an electrode versus a Li metal electrode to assess its fundamental performance. In the below examples, the active material composition to be tested was combined with N-Methyl Pyrrolidone (NMP), carbon black acting as a conductive additive, and poly(vinyldifluoride) (PVDF) binder and mixed to form a slurry using a lab-scale centrifugal planetary mixer (although it is also possible to form aqueous slurries by using water rather than NMP, with binders such as CMC:SBR or alginate). The non-NMP composition of the slurries was 80 w. % active material, 10 w. % conductive additive, 10 w. % binder. The slurry was then coated on an Al foil current collector to the desired loading of 1 mg/cm2 by doctor blade coating and dried in a vacuum oven for 12 hours. In this way, extremely thin coatings were achieved that enabled assessment of material fundamental properties, rather than those driven by electrode quality such as excess impedance, or poor packing of materials. Electrodes were punched out at the desired size and combined with a separator (Celgard porous PP/PE), Li metal, and electrolyte (1 M LiPF6 in EC/DEC) inside a steel coin cell casing and sealed under pressure. Formation cycling was then carried out at low current rates (C/20) for 2 full charge and discharge cycles. After formation, further cycling can be carried out at a fixed or varied current density as required. These tests have been termed “half-cell galvanostatic cycling” for future reference. For samples R1-R10, the electrolyte was altered to 1.3 M LiPF6 in 3:7 EC/DEC, and the formation cycling was carried out at C/10 for 2 charge/discharge cycles in the limits 1.1-3.0 V. The values shown for these samples is an average of 3 measurements, with the error being the standard deviation.
Homogeneous, smooth coatings on current collector foil, the coatings being free of visible defects were also prepared as above with a centrifugal planetary mixer to a composition of 94 w. % active material, 4 w. % conductive additive, 2 w. % binder. The coatings were calendared at 80° C. to a density of up to 3.0 g/cm3 at loadings of 1.3-1.7 mAh/cm2 in order to demonstrate possible volumetric capacities >700 mAh/cm3 in the voltage range 0.7-3.0 V at C/20, and >640 mAh/cm3 in the voltage range 1.1-3.0 V at C/5. This is an important demonstration of these materials being viable in a commercially focussed electrode power cell formulation, where retaining performance after calendaring to a high electrode density allows for high volumetric capacities. Loadings of up to and including 1.0, 1.5, 2.0, 2.5, or 3.0 mAh/cm2 may be useful for Li-ion cells focussed on power performance; loadings greater than 3.0 mAh/cm2 are useful for energy-focussed performance in Li ion cells.
Electrical conductivity of electrodes made with the samples listed in Table 1 was measured using a 4-point probe thin film resistance measurement apparatus. Slurries were formulated according to the procedure described above and coated on a dielectric mylar film at a loading of 1 mg/cm2. Electrode-sized discs where then punched out and resistance of the coated-film was measured using a 4-point probe. Bulk resistivity can be calculated from measured resistance using the following equation:
The results of this test are shown in Table 3, below:
Samples R1-R14 also had their 4-point probe resistance measured to quantify their electrical resistivity. This was carried out with a different Ossila instrument (T2001A3-UK) at 2300 for coatings on mylar films at loadings of 1.0 mg/cm2. The results for sheet resistance (Ω/square) are outlined in Table 3a, with error based on the standard deviation of 3 measurements.
The direct current internal resistance (DCIR) and the resultant area specific impedance (ASI) is a key measurement of internal resistance in the electrode in a Li-ion cell. In a typical measurement, a cell that has already undergone formation will be cycled at C/2 for 3 cycles. With the electrode in its delithiated state a C/2 discharge current is applied for 1 h to achieve ˜50% lithiation. The cell is rested for 30 mins to equilibrate at its OCV (open circuit voltage), and then a 5 C current pulse is applied for 10 s, followed by a 30 mins rest to reach the OCV. During the 10 s pulse the voltage response is sampled at a higher frequency to determine the average internal resistance accurately. 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 afterwards) and the measured voltage. The resistance is then multiplied by the area of the electrode to result in the ASI.
The results of this test are shown in Table 4, below:
The reversible specific capacity C/20, initial coulombic efficiency, nominal lithiation voltage vs Li/Li+ at C/20, 5C/0.5 C capacity retention, and 10 C/0.5 C capacity retention for a number of samples were also tested, the results being set out in Table 5, below. Nominal lithiation voltage vs Li/Li+ has been calculated from the integral of the V/Q curve divided by the total capacity on the 2nd cycle C/20 lithiation. Capacity retention at 10 C and 50 has been calculated by taking the specific capacity at 100 or 5 C, and dividing it by the specific capacity at 0.5 C. It should be noted that the capacity retention was tested with symmetric cycling tests, with equivalent C-rate on lithiation and de-lithiation. Upon testing with an asymmetric cycling program, 10 C/0.5 C capacity retention greater than 89% is routinely observed.
Samples R1-R10 were tested with minor differences in Table 5a, the reversible specific capacity shown is the 2nd cycle delithiation capacity at C/10, the nominal lithiation voltage vs Li/Li+ is at C/10 in the 2nd cycle, the rate tests were carried out with an asymmetric cycling program with no constant voltage steps (i.e. constant current), with lithiation at C/5 and delithiation at increasing C-rates.
The modification of mixed niobium oxide-based Wadsley-Roth and Bronze structures as shown in the reference examples demonstrates the applicability of the modification to improve active material performance in Li-ion cells. By substituting the non-Nb cation to form a mixed cation structure as described, the entropy (cf disorder) can increase in the crystal structure, reducing potential energy barriers to Li ion diffusion through minor defect introduction (e.g. samples R7, 16). Modification by creating mixed cation structures that retain the same overall oxidation state demonstrate the potential improvements by altering ionic radii, for example replacement of an Mo6+ cation with W6+ in sample 14 or Fe3+ or Ga3+ for Al3+ in samples R4 and R5, which can cause minor changes in crystal parameters and Li-ion cavities (e.g. tuning the reversibility of Type VI cavities in Wadsley-Roth structures) that can improve specific capacity, Li-ion diffusion, and increase Coulombic efficiencies of cycling by reducing Li ion trapping. Modification by creating mixed cation structures that result in increased oxidation state (e.g. Ge4+ to replace Zn2+ in sample R2, or Mo6+ for Zr4+ in sample 19) demonstrate similar potential advantages with altered ionic radii relating to capacity and efficiency, compounded by introduction of additional electron holes in the structure to aid in electrical conductivity. Modification by creating mixed cation structures that result in decreased oxidation state (e.g. K+ and Co3+ to replace Ge4+ in sample R7, or Ti4+ to replace Mo6+ in sample 2) demonstrate similar potential advantages with altered ionic radii relating to capacity and efficiency, compounded by introduction of oxygen vacancies and additional electrons in the structure to aid in electrical conductivity. Modification by inducing oxygen deficiency from high temperature treatment in inert or reducing conditions demonstrate the loss of a small proportion of oxygen from the structure, providing a reduced structure of much improved electrical conductivity (e.g. sample 5, R10 and R12-14) and improved electrochemical properties such as capacity retention at high C-rates (e.g. sample 5, R13). Combination of mixed cation structures and induced oxygen deficiency allows multiple beneficial effects (e.g. increased specific capacity, reduced electrical resistance) to be compounded (e.g. samples 18, R8).
Across all materials tested, each modified (cation substituted and/or oxygen deficient) material demonstrates an improvement versus the unmodified ‘base’ crystal structure. This is inferred from measurements of resistivity/impedance by two different methods, and also electrochemical tests carried out in Li-ion half coin cells, particularly the capacity retention at increased current densities (cf. rates, Table 5). Without wishing to be bound by theory, the inventors suggest that this is a result of increased ionic and electronic conductivity of the materials as defects are introduced, or by alterations to the crystal lattice by varying ionic radii; also evidenced by DCIR/ASI (Table 4) and EIS (
The data in Table 3 show a large reduction in the resistivity between sample 1 (comparative) and samples 2, 4, 5, 14, 15, 16, 18, 22, demonstrating the effect of improving electrical conductivity of the crystal structures through both cation exchange, oxygen deficiencies, and carbon coating. Samples 17, 19, and 20 also show a similarly low resistivity versus sample 6. The resistivity slightly increased upon incorporation of 0.05 equivalents of V species in the base crystal in sample 7, however an improvement in specific capacity was observed due to the changes in available Li-ion sites in the crystal lattice likely as a result of the differing ionic radius of V over Zr (see Table 5).
The data in Table 4 shows a large reduction in the DCIR/ASI from sample 1 (comparative) to samples 2, 4, 14, 16, 18 and 22, reflecting the trends shown in Table 3. Samples 7, 17, and 19 demonstrate a higher than these by DCIR, however these relate to a different base crystal structure. Without wishing to be bound by theory, the inventors hypothesise that samples 7, 17, and 19 demonstrate an increase in DCIR/ASI as compared with the comparative material of sample 6 (ZrNb24O62) due to the changes in the crystal lattice with the introduced cations of different ionic radii. However, it remains beneficial in terms of conductivity for these structures for samples 17 and 19 as the electrical resistivity is decreased as shown in Table 3, thereby minimising joule heating and enabling a more uniform current distribution across the material, which in turn can enable improved safety and lifetime of a Li ion system. For sample 7, whilst there is no demonstrated improvement utilising V to exchange with Zr, there is an increase in specific capacity, as discussed above.
In Table 5, across most samples there is a trend for improved specific capacities, initial Coulombic efficiencies (ICE), nominal lithiation voltage vs Li/Li+, and importantly capacity retention at 5 C and 10 C vs 0.5 C for modified materials versus the comparative ‘base’ materials (e.g. samples 1, 6, 8, 10, 12). For example samples 2, 3, 4, 5, 14, 15, 16, 18, 22 all demonstrate improvements in one or more of these parameters vs sample 1. This is also the case for samples 7, 17, 19 versus sample 6 across multiple parameters; sample 11 and 21 versus 10 where an improvement in specific capacity or capacity retention is observed; sample 9 versus 8 where ICE and capacity retention are improved; and sample 13 versus 12 where ICE and capacity retention are improved.
Electrochemical impedance spectroscopy (EIS) measurements were also carried out to gain a further understanding on the impedance present in the electrode in a Li-ion cell. In a typical measurement, the cell is prepared as for DCIR measurements to ˜50% lithiation and then the frequency of alternating charge/discharge current pulses is varied whilst measuring the impedance. By plotting the real and imaginary components as the axes, and varying the AC frequency, a Nyquist plot is generated. From this plot for a Li-ion cell different types of impedance in the cell can be identified, however it is typically complex to interpret. For example, Ohmic resistance can be partially separated from electrochemical double layer effects and also separated from diffusion effects.
Commercial-grade LTO (Li4Ti5O12) was purchased from Targray Technology International Inc with properties outlined in Table E1 (Sample E1). The modified and carbon-coated Wadsley-Roth and Bronze materials were synthesised in-house by a solid-state route. In a first step precursor materials (e.g. Nb2O5, WO3, ZrO2, TiO2, MoO3, Cr2O3, ZnO, Al2O3 etc.) were mixed in stoichiometric proportions (200 g total) and ball-milled at 550 rpm with a ball to powder ratio of 10:1 for 3 h. The resulting powders were heat treated in an alumina crucible in a muffle furnace in air at T1=800-1350° C. for 24 h, providing the desired Wadsley-Roth or Bronze phase. An additional heat treatment step was also applied under a N2 atmosphere at T2=800-1350° C. for 5 h to result in minor oxygen deficiencies in the base crystal structure for samples E2, E3, E4. For Sample E2 the above synthesis was carried out with T1=9000, T2=9000. For Sample E4 the above synthesis was carried out with T1=11000, T2=11000. For sample E5 the above synthesis was carried out with T1=11000, repeated twice with an intermediary grinding step by impact milling at 20,000 rpm. For sample E6, the above synthesis was carried out with T1=11000 for 24 h.
Sample E2 (98 g) was then combined with petroleum pitch (2 g) (ZL 118M available from Rain Carbon) by high energy impact mixing/milling. The mixture was heat treated in a furnace under reducing conditions at T=9000 for 5 h to provide Sample E3, which was a free-flowing black powder. A final de-agglomeration step was utilised for each sample by impact milling or jet milling to adjust to the desired particle size distribution. Specifically, the material was de-agglomerated by impact milling at 20,000 RPM for 10 seconds.
Active electrode material mixtures of MNO and LTO were obtained by low to high energy powder mixing/blending techniques, such as by rotational mixing in multiple directions, rotational V-type blending over a single axis, planetary mixing, centrifugal planetary mixing, high shear mixing, and other typical mixing/blending techniques. In this case, mixing was achieved with a centrifugal planetary mixer on 5 g batches of materials, mixed at 2000 rpm for 3 mins, 10 times.
The phase purity of samples was analysed using a Rigaku Miniflex powder X-ray diffractometer in 20 range (20-70°) at 1°/min scan rate.
Diffraction patterns in Sample E2 has peaks at the same locations (within instrument error, that is 0.10) and match JCPDS crystallography database entry JCPDS 73-1322, which corresponds to MoNb12O33. Sample E3 has some changes to its peaks due to the introduced oxygen-deficiency beginning to induce minor crystallographic distortions due to significant quantities of vacancy defects, and additional peaks relating to the carbon at ˜26° and ˜40°. There is no amorphous background noise and the peaks are sharp and intense. This means that all samples are crystalline, with crystallite size 38±4 nm for Sample E2 and 32±12 nm for Sample E3 according to the Scherrer equation and crystal structure matching MoNb12O33. This confirms the presence of a Wadsley-Roth crystal structure.
Diffraction patterns in Sample E4 has peaks at the same locations (within instrument error, that is 0.1°) and match JCPDS crystallography database entry database JCPDS 00-020-1320, which corresponds to W7Nb4O31. There is no amorphous background noise and the peaks are sharp and intense. This means that the sample is phase-pure and crystalline, with crystallite size 43±10 nm according to the Scherrer equation and crystal structure matching W7Nb4O31. This confirms the presence of a Tetragonal Tungsten Bronze crystal structure.
Sample E5 presented a phase mixture between the orthorhombic and monoclinic forms of the 3×4×∞ Wadsley-Roth structure, corresponding to crystallography database entry JCPDS 28-1478 and PDF card: 04-021-7859. There is no amorphous background noise and the peaks are sharp and intense. This means that the samples are phase-pure and crystalline, with crystallite size 49±6 nm according to the Scherrer equation and crystal structure matching Zn2Nb34O87. This confirms the presence of a Wadsley-Roth crystal structure.
Diffraction patterns in Sample E6 have peaks at the same locations (with some shift due to crystal modification, up to around 0.2°), and match crystallography database entry JCPDS 22-009. There is no amorphous background noise and the peaks are sharp and intense. This means that the sample is crystalline, with crystallite size 45 nm according to the Scherrer equation and crystal structure matching AlNb11O29. This confirms the presence of a Wadsley-Roth crystal structure.
Thermogravimetric Analysis (TGA) was performed on Sample E3 using a Perkin Elmer Pyris 1 system in an air atmosphere. Samples were heated from 30° C. to 900° C. at 5° C./min, with an air flow of 20 mL/min.
Particle Size Distributions were obtained with a Horiba laser diffraction particle analyser for dry powder. Air pressure was kept at 0.3 MPa. The results are set out in Table E1. BET surface area analysis was carried out with N2 on a BELSORP-miniX instrument at 77.35 K and are set out in Table E1.
The electrochemical characterisation of the Examples was performed under different conditions to the Reference Examples. Therefore, in some instances the electrochemical characterisation of the Examples may not be directly comparable to the electrochemical characterisation of the Reference Examples.
Li-ion cell charge rate is usually expressed as a “C-rate”. A 1 C charge rate means a charge current such that the cell is fully charged in 1 h, 10 C charge means that the battery is fully charged in 1/10th of an hour (6 minutes). C-rate hereon is defined from the reversible capacity of the anode within the voltage limits applied, i.e. for an anode that exhibits 1.0 mAh cm−2 capacity within the voltage limits of 1.1-3.0 V, a 1 C rate corresponds to a current density applied of 1.0 mA cm-2.
Electrochemical tests were carried out in half-coin cells (CR2032 size) for analysis. In half-coin tests, the active material is tested in an electrode versus a Li metal electrode to assess its fundamental performance. In the below examples, the active material composition to be tested was combined with N-Methyl Pyrrolidone (NMP), carbon black acting as a conductive additive, and poly(vinyldifluoride) (PVDF) binder and mixed to form a slurry using a lab-scale centrifugal planetary mixer. The non-NMP composition of the slurries was 90 wt % active material, 6 wt % conductive additive, 4 wt % binder. The slurry was coated on an Al foil current collector to the desired loading of 5.7-6.6 mg cm−2 by doctor blade coating and dried. The electrodes were then calendared to a density of 2.00-3.75 g cm−3 (dependent on material density) at 80° C. to achieve targeted porosity of 35-42%. Porosity was calculated as the measured electrode density divided by the weighted average density of each component of the composite electrode coating film. Electrodes were punched out at the desired size and combined with a separator (Celgard porous PP/PE), Li metal, and electrolyte (1.3 M LiPF6 in EC/DEC) inside a steel coin cell casing and sealed under pressure. Cycling was then carried out at low current rates (C/10) for 2 full cycles of lithiation and de-lithiation between 1.1-3.0 V. Afterwards, the cells were tested for their performance at increasing current densities. During rate tests, the cells were cycled asymmetric, with a slow lithiation (C/5, with a CV step at 1.1V to C/20 current density) followed by increasing de-lithiation rates for de-lithation rate tests. All electrochemical tests were carried out in a thermally controlled environment at 23° C.
The first cycle efficiency was calculated as the fraction of de-lithiation capacity/lithiation capacity in the first cycle at C/10. The nominal voltage at each C-rate was determined by integrating the voltage-capacity curves and then by dividing it by the total capacity.
To quantify the significance of the differences in data observed, an error calculation was carried out and applied to the values for specific capacity. The error for these was approximated as the largest error possible with the microbalance used (±0.1 mg), and the lowest loading electrode (5.7 mg cm−2) on a 14 mm electrode disc. This provides an error of t 1.1%, which has been applied to each capacity measurement. Error in Coulombic efficiency, capacity retention, and voltage were assumed to be negligible as the instrument accuracy far exceeds the stated significant figures, and the values are independent of the balance errors.
Sample E2 has a Wadsley-Roth 3×4 block shear crystal structure based on a MVINb12O33 crystal structure where all blocks are connected by tetrahedra, that has been made oxygen-deficient through heat treatment in an inert atmosphere and through cation exchange. The combination of induced oxygen-deficiency and cation exchange leads to improved electrochemical performance versus a material such as WNb12O33 or MoNb12O33.
As shown in tests A* and B*, Sample E2 has a higher specific capacity, lower ICE, lower capacity retention at higher C-rates, and higher nominal voltage at each C-rate than Sample E1. Therefore, by providing a physical mixture of the two materials in suitable proportions, the disadvantages of each can be alleviated. Due to the material design having suitable mixing characteristics, such as particle size distribution and surface chemistry, homogeneous powdered mixtures and subsequently homogeneous coated electrodes can be produced having an intimate mixture of the 2 components.
Tests E and G demonstrate mixtures with a high proportion of Sample E1 at 90 and 95 wt % respectively, and Tests F and H demonstrate mixtures with a high proportion of Sample E2 in a similar fashion. Tests E and G show increased specific capacity vs test A*, and improved initial Coulombic efficiency (ICE) vs test B; exemplifying the advantages of providing a physical mixture of LTO with Wadsley-Roth MNO materials. Tests F and H show increased specific capacity vs test A*, and improved ICE vs test B* in a similar fashion.
The de-lithiation capacity retention further demonstrates advantages to the mixture of active materials, with increased retention for tests E and G vs test B*. It is expected in a full cell arrangement with a cathode active material, that similar benefits will be observed for lithiation of the anode active material. The nominal de-lithiation voltage can be improved for the mixture of materials vs the individual active materials, with tests E, F, G, and H all showing a reduced nominal voltage compared to test B*.
Sample E5 is a Wadsley-Roth 3×4 block shear crystal structure based on a MII2Nb34O37 crystal structure composed of octahedra and no tetrahedra. Similar advantages can be observed in test M versus tests A* and K*.
Sample E6 is a Wadsley-Roth crystal structure based on MIINb11O29. Similar advantages can be observed in test N versus tests A* and L*. Notably, test N, a 50:50 mixture of Samples E1 and E6, was found to provide improved capacity retention at 1 C, 2 C, and 5 C compared to both tests A* and L*, the respective individual materials E1 and E6.
It is expected that similar benefits will be observed with all Wadsley-Roth crystal structures containing Nb as described in the claims mixed with lithium titanate as described above for use in Li-ion cells.
Sample E3 is a modified form of Sample E2, which has been coated with pitch-carbon by high energy milling, and then pyrolysed in an inert atmosphere to provide increased oxygen deficiency, and a polyaromatic sp2-based carbon coating based on a pitch precursor. This provides advantages in reducing impedance, reducing nominal voltage, and improving performance at high rate vs Sample E2. It carries further advantages such as improved surface electrical conductivity of the active material crystal, and improved mixing with other components of the electrode such as the carbon additive (typically carbon black, graphite, etc).
Test J shows an improved specific capacity vs test A*, and an improved ICE vs test C. The de-lithiation capacity retention is greater for test J than for both test A* and C*, implying that providing the mixture of these 2 different active materials is advantageous for high rate performance over either individual materials of Sample E1 and E3. This could be due to a combination of favourable surface chemistry of Sample E1 and E3 leading to enhanced electrode quality (adhesion, cohesion, conductivity), or a favourable combination of the electrochemical properties that can prevent impedance more effectively than the individual materials. Furthermore, the nominal de-lithiation voltage is decreased for test J vs test C*.
It is expected that similar benefits will be observed with all MNO materials that are coated with carbon and are oxygen-deficient as described mixed with lithium titanate for use in Li-ion cells.
Sample E4 has a Bronze crystal structure that has been made partially oxygen-deficient by a heat treatment in an inert atmosphere and by cation exchange. Specifically, the MVI7Nb4O31 base crystal structure has 3, 4, and 5 sided tunnels with a low degree of filled tunnels, resulting in a high availability of Li-ion intercalation sites. In this case the combination of cation exchange and oxygen deficiency provides improved electrochemical performance versus materials such as W7Nb4O31.
Test I demonstrates increased specific capacity and increased ICE versus test A*. Test I further demonstrates great improvement in de-lithiation capacity retention vs test D, more so than expected with inclusion of Sample E1 as the minor component (5 w/w %).
It is expected that similar benefits will be observed with all Bronze (Tetragonal Tungsten Bronze) crystal structures containing Nb as described in the claims mixed with lithium titanate for use in Li-ion cells.
While the invention has been described in conjunction with the exemplary embodiments described above, many equivalent modifications and variations will be apparent to those skilled in the art when given this disclosure. Accordingly, the exemplary embodiments of the invention set forth above are considered to be illustrative and not limiting. Various changes to the described embodiments may be made without departing from the spirit and scope of the invention.
For the avoidance of any doubt, any theoretical explanations provided herein are provided for the purposes of improving the understanding of a reader. The inventors do not wish to be bound by any of these theoretical explanations.
Any section headings used herein are for organizational purposes only and are not to be construed as limiting the subject matter described.
A number of publications are cited above in order to more fully describe and disclose the invention and the state of the art to which the invention pertains. Full citations for these references are provided below.
The entirety of each of these references is incorporated herein.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2013576.0 | Aug 2020 | GB | national |
| 2104508.3 | Mar 2021 | GB | national |
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/GB2021/052231 | 8/27/2021 | WO |
