Heterostructure assemblies

Abstract

The disclosure relates to vertical layered A-[B-A]n type heterostructure assemblies of a 2D ultrathin inorganic solid-state material ('B') and a non-noble metal organic framework (MOF) material ('A'), with bridging MIM-X-MMOF ionic bonds that stabilise heterointerfaces between layers A and B, as well as to methods and processes of synthesising same, and uses of same as electrocatalysts, such as for OER in water splitting.

Description

[0001] HETEROSTRUCTURE ASSEMBLIES

[0002] Cross-Reference

[0003]

[0001] The present application claims priority to Australian provisional patent application AU 2024901113 filed on 19 April 2024, the entire contents of which are incorporated herein by crossreference.

[0004] Technical Field

[0005]

[0002] The invention relates to a heterostructure assemblies suitable as electrocatalysts for water splitting, and which are particularly suitable for OER in alkaline aqueous solutions, including seawater.

[0006] Background

[0007]

[0003] Seawater has emerged as a preferred and unlimited raw material for sustainable I green hydrogen production via electrolysis. The two half-reactions of water splitting are the hydrogen evolution reaction (HER) at the cathode and oxygen evolution reaction (OER) at the anode. Both require highly efficient electrocatalysts to overcome the required overpotentials that make water splitting inefficient. Some excellent electrocatalysts for OER and HER have been developed, but catalysts that are stable and highly active in seawater are still lacking. Overall, H2 production at the cathode is strongly dependent on the efficiency and stability of the OER at the anode. In real seawater, microorganisms and insoluble precipitate impurities inhibit performance by shielding catalytic active sites on the catalyst surface. Chloride products corrode the catalyst and deactivate active sites, thereby interfering with the driving OER and causing rapid decline in activity and stability. OER catalysts for selective seawater oxidation that inhibit chloride ion oxidation reactions are scant due to anodic side reactions in seawater splitting. Seawater spitting is challenging, particularly for OER at the anode, compared to more straightforward HER at the cathode side, and chloride ion oxidation occurs due to the presence of approx. 0.5 M chloride anions in real seawater.

[0008]

[0004] The cathode chemistry is more straightforward than the anode chemistry, as the thermodynamic potentials of competing HER reactions at the cathode are not as close as the thermodynamic potentials for OER and chloride ion oxidation reactions which occur on the anode side. OER, as a four-electron oxidation process involve a high energy barrier reaction pathway with at least two or three intermediates, and thus is much more sluggish reaction than the chlorine evolution reaction (CER) which is a kinetically faster reaction as it involves only a two-electron transfer process with a single catalytic intermediate. CER is pH independent and occurs at relatively high potentials. In contrast, the OER potential is pH dependent and much closer to the CER potential under acidic conditions than under alkaline conditions where the OER potential is much lower. Kinetically, CER is more favourable than OER, so in low pH media, despite its slightly higher redox potential, CER competes with OER under acidic conditions. Indeed, as CER competes with OER in seawater, the anodic reaction efficiencies are poor and suppression of OER at increased current densities becomes the bottleneck in seawater splitting. In alkaline media, as a result of the lower OER potential, there is an overall wider difference in the thermodynamic potentials for OER and chloride ion oxidation at the anode. Indeed, at pH > 7.5, OER is highly selective over CER within an overpotential window of 480 mV. However, at such high pHs in seawater, alongside OER, undesirable hypochlorite anion formation occurs as the dominant competitive anodic reaction to OER. Such competing reactions could be avoided by using an electrocatalyst active for OER below a 480 mV overpotential for alkaline seawater splitting. However, current densities at such low overpotential are typically low. Ideally an industrial current density (1 A. cm-2) below a 480 mV overpotential for alkaline seawater splitting would be desirable. Practical I commercial applications require catalysts to operate at high current densities above 500 mA cm-2which require substantially higher overpotentials, but unfortunately, at such required higher overpotentials, hypochlorite formation is promoted even in alkaline seawater. Ch corrosion means seawater splitting occurs with reduced efficiency and stability and catalyst deactivation occurs rapidly.

[0009]

[0005] Precious metal oxides, e.g. RuC>2 and lrC>2, are the current benchmark catalysts for OER, but their scarcity, high cost and compromised performance in seawater hamper their large-scale application. Engineered non-noble metals-based transition metal compounds are under study to improve OER kinetics to realise direct seawater splitting, but stability is challenging in the presence of the massive concentrations of Ch in seawater.

[0010]

[0006] Therefore, developing improved highly conductive, long term stable anode materials that can operate as water / seawater splitting catalysts at large current densities under alkaline conditions for longer periods at low overpotentials that inhibit chlorine evolution or at higher overpotentials but involve a design that avoids hypochlorite formation is an important goal. Other desirable goals include better performance, stability and great economy compared to commercial RuO2 and lrO2 catalysts for OER. High stability would make electrolysis systems integrable with renewable power sources.

[0011]

[0007] Heterostructuring dissimilar materials can introduce new properties that might overcome the inherent corrosive chemistry of seawater and improve reaction kinetics of OER in seawater. However, known 2D heterostructures are assembled through weak van der Waals (vdW) interactions, and have limitations, such as suppression of electron transfer at the heterointerface during electrochemical water splitting. The weak interactions in typical 2D heterostructures does not alter the chemical states in a way that creates additional active sites. Furthermore, weak interactions at the heterointerface limits long-term stable performance (generally < 100 hours). Moreover, complex synthesis methods make it difficult to engineer the heterointerface on a large scale for hydrogen production. Therefore, another desired goal is to develop a simple and scalable approach to produce improved materials having a well-defined conductive interface that could match or improve existing precious metal oxide catalysts for direct seawater splitting.

[0012]

[0008] Current fabrication techniques for 2D heterostructure materials (2D HSs) involve physical and direct growth synthesis techniques for smaller volume applications and include mechanical transfer, chemical vapour deposition (CVD), epitaxial growth, and atomic layer deposition (ALD). However, much larger quantities of 2D HSs are a prerequisite for energy storage and conversion applications and scale synthesis primarily involves electrostatic self-assembly, layer-by-layer assembly and interface assisted growth. For example, Liu et al. described an interface-assisted direct growth technique to fabricate perpendicularly aligned metal-organic framework (MOF) nanosheets on rGO (a primarily carbon material). A solution assisted heterogeneous nucleation strategy was implemented for the direct growth of the 2D MOF on the flat rGO surface with controlled perpendicular MOF orientation. Pre-grown rGO sheets were used as the solid interface, which provided nucleation sites for perpendicular orientated assembly of the MOF constituents. Liu’s material has vdW interactions at the heterointerface. The weak nature of the vdW interactions limit charge transfer, and lattice mismatching raises the issue of stability when used for electrochemical applications, thus improved materials will require stronger bonding at the interface.

[0013]

[0009] A reference herein to a patent document or other matter which is given as prior art is not to be taken as an admission that the document or matter was known or that the information it contains was part of the common general knowledge as at the priority date of any of the claims.

[0014] Statements of the Invention

[0015]

[0010] In a first aspect the invention provides a vertical layered A-[B-A]ntype heterostructure assembly of a 2D ultrathin inorganic solid-state material (‘B’) and a non-noble metal organic framework (MOF) material (‘A’) that covers surfaces of each layer of the 2D ultrathin inorganic solid-state material (‘B’) in the assembly in a parallel arrangement to the 2D ultrathin inorganic solid-state material (‘B’), wherein n is a positive integer, wherein heterointerfaces of the assembly comprise bridging MIM-X-MMOFionic bonds that stabilise heterointerfaces between the 2D ultrathin inorganic solid-state material (‘B’) and the MOF material (‘A’), where bridging X is -O-, -S-, -Se-, -C-, -N-, -P-, or -B-, MIMis a metal in the inorganic solid-state material, and MMOFis a metal in the non-noble MOF material; wherein the 2D ultrathin inorganic solid-state material comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; and further comprises one or more of: metal-0 (oxide), metal-N (nitride), metal-C (carbide), metal-C / metal N (carbonitride); metal-S (sulfide) and / or M-P (phosphide) functionalities, and optionally wherein the 2D ultrathin inorganic solid-state material is doped with one or more non-metal dopants selected from the group consisting of: B, P, C, N, O, Si, S, As and Se.

[0016]

[0011] As used herein, the term “MOF” will be understood to refer to a non-noble metal organic framework, unless the context clearly indicates otherwise.

[0017]

[0012] For convenience, some elements have been referred to herein as “atoms” for the purposes of identifying them as being present, or doped into, a certain material. However, in such materials, the atoms may have an assigned formal charge (be ions), and so references to atoms in some cases may also be understood as references to the element in its ion form.

[0018]

[0013] In the formula A-[B-A]nherein, n is a positive integer, such as 1 , 2, 3, 4, 5, 6, 7, 8, 9, 10, ... 20, ... , 50, ... , 100, etc. When n = 1 , the assembly has formula A-B-A and the assembly has a tri-layer structure, when n = 2, the assembly has formula A-B-A-B-A, when n = 3, the assembly has formula A- B-A-B-A-B-A, etc. It will be understood that assemblies may form with any number of layers, and within a given composition, there may be a distribution of assemblies having different values of n but still consistent with this vertically layered formula.

[0014] Typically, the MOF is found on the top and bottom surfaces of the vertically layered heterostructure. Typically, the MOF is present on both sides of the basal plane of the inorganic 2D sheets and fully covers the inorganic sheets.

[0019]

[0015] While doping with one or more non-metal dopants is optional, it is preferred as it has been found to induce better ionic connections between the components which enhances catalytic performance. The non-metal dopants also assist in chloride shielding polyanion formation and / or assist in creation of ionic bonds at heterointerface. The polyanions, when present, contribute to a Cl- shielding effect, though most of the Ch shield results from the ionic bonds between the components of the assembly. Desirably, the non-metal dopants result in lattice defects on at least the surfaces of the inorganic solid-state material. The type of defect varies but desirably tunes the electronic properties of the material to provide defects in the material surface that facilitate formation of ionic bonding in the HSs of the invention. B and P are particularly effective non-metal dopants in this regard.

[0020]

[0016] The vertical layered A-[B-A]ntype heterostructure assembly is not a lateral heterostructure. Desirably, the non-noble metal organic framework (MOF) material (‘A’) is also located on the ends of both sides of the vertical layered A-[B-A]ntype heterostructure assembly. Desirably, the inorganic solid-state material (‘B’) is doped with one or more non-metal dopants, preferably selected from the group consisting of: B, C, N, O, P, Si, S, As and Se.

[0021]

[0017] Desirably, the non-metal dopants result in chloride shielding polyanion formation on at least the surfaces of the inorganic solid-state material. One or more of S, B and P are particularly effective in this regard. However, for the electrocatalysts of the invention, shielding against CER is not only because of polyanions, which are supportive, protection also comes from formation of the ionic bond itself which is assisted by the non-metal dopant.

[0022]

[0018] In some examples, the 2D-2D HS materials of the invention are porous, which is helpful in catalysis by improving the mass flow.

[0023]

[0019] In a second aspect the invention provides a process of forming an electrocatalyst comprising the steps of: providing a liquid dispersion of an inorganic solid state material (‘B’) in 2D ultrathin sheet form as an inorganic template for non-noble metal organic framework (MOF) material (‘A’) growth thereon in a parallel arrangement to the 2D ultrathin inorganic solid-state material (‘B’), wherein the inorganic template comprises or is doped with one or more non-noble electrochemically active transition metals preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; and further comprises one or more of: metal- O (oxide), metal-N (nitride), metal-C (carbide), metal-C / metal N (carbonitride); metal-S (sulfide) and / or M-P (phosphide) functionalities, and optionally is doped with one or more non-metal dopants selected from the group consisting of: B, C, N, O, P, Si, S, As and Se; electrostatically associating at least one basic organic coordination ligand species suitable for MOF fabrication with the inorganic template; providing a liquid solution of one or more acidic non-noble transition metal (MMOF) salts to the dispersion of inorganic template having electrostatically associated ligand species; forming a vertical layered A-[B-A]ntype heterostructure assembly of the 2D ultrathin inorganic material (‘B’) with the non-noble metal organic framework MOF material (‘A’) by allowing complexation of acidic MMOFions of the salt with the at least one basic organic coordination ligand species associated with the inorganic template to form the MOF in parallel coverage over surfaces of the inorganic template, whereby during complexation, MMOFforms bridging MIM-X-MMOFionic bonds with metals in the inorganic template, MIM, at heterointerfaces between the inorganic template and MOF, where bridging X is -O-, -S-, -Se-, -C-, -N-, -P-, or -B-, and wherein n is a positive integer. Advantageously no additional steps, such as reduction steps, are required to form the HS.

[0024]

[0020] In the context of the second embodiment and elsewhere herein, the basic organic coordination ligand species refers to the coordination ligand species being a Lewis base.

[0025]

[0021] Desirably, the inorganic solid-state material is doped with one or more non-metal dopants, preferably selected from the group consisting of: B, P, C, N, O, Si, S, As and Se. Desirably, the chloride ion shielding polyanions are one or more of: borate, sulfate, phosphate, chromate, permanganate, carbonate, and nitrate, most preferably borate, sulfate and phosphate polyanions.

[0026]

[0022] Desirably, the inorganic template comprises, or is doped with, one or more electrochemically active transition metal atoms that form the bridged X ionic bonds (MIM-X- MMOF) with the non-noble transition metal of the MOF, wherein bridging X is -O-, -S-, -Se-, -C-, -N-, -P-, or -B-, more preferably bridging X is O, N, C, S, B, and / or P, most preferably bridging X is -O-.

[0027]

[0023] Desirably, MIMmay be selected from the group consisting of: Fe, Ni Co, W, Ti, Cu, V and Mo. Desirably, MMOFmay be selected from the group consisting of: Co, Ni, Fe, Cu, Zn, Zr, and Ti. In one embodiment, MMOFmay be selected from the group consisting of: Co, Ni, Fe, Cu, Zn, Zr, Ti, and combinations thereof. In one embodiment, MMOFis a mixture of one or more metals selected from the group consisting of: Co, Ni, Fe, Cu, Zn, Zr, and Ti. In one embodiment, MMOFis a mixture of one or more metals selected from the group consisting of: Co, Ni, Fe, and Zn. Desirably, the electrochemically active transition metal atoms are Fe, Ni, Co, Cu, Cr, Mo and W, more preferably, Fe, Ni or Co, most preferably Fe. Desirably, the one or more electrochemically active transition metals doped into the inorganic solid-state material are different to the metal(s) in the inorganic solid-state material.

[0028]

[0024] Desirably, the inorganic solid state material / inorganic template is a transition metal oxide (TMO), a transition metal nitride (TMN), a transition metal carbonitride (TMCN), a transition metal sulfide (TMS), a transition metal carbide (TMC), transition metal phosphide (TMP), or a mixed variant thereof, such as a MXene or functionalised MXene and which are inorganic solid state materials that consist of atomically thin layers of TMC, TMN or TMCN. In one embodiment, the inorganic solid state material / inorganic template is a functionalised MXene of formula Mn+iXnTx, where M is an early transition metal (such as a Group 3 to Group 6 transition metal), particularly Ti, Mo or V, X is C and / or N, and T is a functional group (e.g. O, F, OH, Cl). Preferred functionalised MXenes include TisC2O2, which is an O-functionalised titanium carbide comprising a majority, such as at least 50-75%, of functional groups, T, being O.

[0029]

[0025] Desirably, the inorganic solid state material / inorganic template is a transition metal oxide (TMO), a transition metal nitride (TMN), a transition metal carbonitride (TMCN), a transition metal sulfide (TMS), a transition metal carbide (TMC), transition metal phosphide (TMP), or a mixed variant thereof, such as a MXene (Mn+iXnTx), doped with one or more of B, C, N, O, P, Si, S, As and Se. In one embodiment, the inorganic solid state material / inorganic template comprises a transition metal oxide (TMO) or a TMO doped with one or more of B, C, N, O, P, Si, S, As and Se. In one embodiment, the inorganic solid state material / inorganic template comprises a transition metal nitride (TMN) or TMN doped with one or more of B, C, N, O, P, Si, S, As and Se. In one embodiment, the inorganic solid state material / inorganic template is a transition metal carbonitride (TMCN) or a TMCN doped with one or more of B, C, N, O, P, Si, S, As and Se. In one embodiment, the inorganic solid state material / inorganic template is a transition metal carbide (TMC) or a TMC doped with one or more of B, C, N, O, P, Si, S, As and Se. In one embodiment, the inorganic solid state material / inorganic template is a transition metal sulfide (TMS) or a TMS doped with one or more of B, C, N, O, P, Si, S, As and Se. In one embodiment, the inorganic solid state material / inorganic template is a transition metal phosphide (TMP) or a TMP doped with one or more of B, C, N, O, P, Si, S, As and Se. The transition metal in these embodiments may be selected from one or more of Fe, Ni Co, W, Ti, Cu, V and Mo. The transition metal in these embodiments is preferably a non-noble transition metal.

[0030]

[0026] In one embodiment, the inorganic solid state material / inorganic template comprises a MXene or a MXene doped with one or more of B, C, N, P, Si, S, As and Se. In the case of MXenes, the material can be further doped with one or more electrochemically active transition metal atoms that replace some of the early transition metal. A preferred non-metal dopant for MXenes is B. In preferred MXenes, the material is dual doped with one or more electrochemically active transition metal atoms, particularly Fe, and with one or more non-metal dopants, particularly B. Such dual modification / doping desirably breaks the electrochemical neutrality of the MXene through the incorporation of Fe, while B dopant is particularly effective at tuning the electronic properties of the material to provide defects in the material surface which facilitates formation of ionic bonding in the HSs of the invention.

[0031]

[0027] Desirably, the inorganic solid state material / inorganic template comprises a non-metal dopant selected from one or more of B, P, C, N, O, Si, S, As and Se, most preferably one or more of: B, C, S and P; or B and P. Desirably, the polyanion is therefore one or more of borate, sulfate, phosphate, chromate, permanganate, carbonate, and nitrate, most preferably is one or more of: borate, carbonate, sulfate or phosphate functionalities on surfaces of the inorganic solid-state material that repel Ch from the electrochemically active metals of the inorganic solid-state material and / or aid in formation of the ionic bond between the metals in the inorganic material, MIM, bridging X, and MMOFmetals in the non- noble MOF.

[0032]

[0028] Desirably, the inorganic solid state material / inorganic template comprises a non-metal dopant selected from one or more of B, S and P, such that there is B, S and / or P atom substitution or insertion, and / or a borate, sulfate and / or phosphate surface polyanion functionality that repels Ch from the electrochemically active metals of the inorganic solid-state material.

[0033]

[0029] Desirably, the inorganic solid state material / inorganic template is twin doped with B and P, and thus the polyanions are borate and phosphate functionalities that repel Ch from the electrochemically active metals of the inorganic solid-state material, particularly wherein the inorganic solid-state material comprises one or more of Fe, W, and / or Ni.

[0034]

[0030] Desirably, the inorganic solid-state material is transition metal oxide (TMO), a transition metal nitride (TMN), a transition metal carbonitride (TMCN), a transition metal sulfide (TMS), a transition metal carbide (TMC), transition metal phosphide (TMP), or mixed variants thereof, doped with B, or with B and P. MXene examples involving Ti can be doped with one or more electrochemically active metals such as Fe, Ni, Co, Cu, Cr, Mo and W, preferably Fe, W and / or Ni, and most preferably Fe, which may replace some of the surface Ti.

[0035]

[0031] In some examples, the inorganic template in 2D ultrathin form is selected from: non-noble transition metal oxide (TMO), such as iron oxide, such as Fe20s, tungsten oxide such as WO3, vanadium oxide, such as V2O5, molybdenum oxide such as M0O3; non-noble transition metal MXene or functionalised MXene, preferably B-TisC2Tx, where Tx is O; non-noble transition metal sulfide (TMS), such as iron sulfide such as FeS2, cobalt sulfide such as C0S2, or nickel sulfide such as NiS2;

[0036] B-doped non-noble transition metal MXene or B-doped functionalised MXene, preferably B- TisC2Tx or B-Fe-TisC2Tx, preferably where Tx is O, F, OH or Cl, preferably O;

[0037] B doped 2D non-noble transition metal sulfide (TMS), for example, B-doped iron sulfide such as B-FeS2, B-doped cobalt sulfide such as CoS; or B-doped nickel sulfide such as B-NiS2,

[0038] B doped 2D non-noble transition metal oxide (TMO), for example, B-doped iron oxide such as B- Fe20s, B-doped tungsten oxide such as B-WO3, B-doped vanadium oxide such as B-V2O5, B- doped molybdenum oxide such as B-M0O3;

[0039] B doped and / or transition metal doped, such as Ni, Co or Fe doped, non-noble transition metal oxide (TMO), for example, B- and Fe-doped iron oxide such as B-doped iron oxide such as B- Fe20s or B-a-Fe2Os or B- and Fe- doped tungsten oxide such as B-Fe-WOs;

[0040] B doped and Fe doped non-noble transition metal sulfide (TMS), for example, B- and Fe- doped iron sulfide such as B-Fe-FeS2, B- and Fe- doped cobalt sulfide such as B-C0S2, B- and Fe- doped nickel sulfide such as B-Fe-NiS2, B- and Fe- doped tungsten sulfide such as B-Fe-WS2, B- and Fe- doped vanadium sulfide such as B-Fe-V2S5, or B- and Fe doped molybdenum sulfide such as B-Fe-MoS2;

[0041]

[0032] Desirably, the inorganic solid-state material / inorganic template in 2D ultrathin form is selected from: Fe20s; WO3.B2O3; B-Fe20s; B-Fe-Ti3C2Tx, where X is O; B-FeS2; B-NiS2; and P-W8O26.B2O3; MoN, M0S2, B-C0S2.

[0042]

[0033] Desirably, the at least one basic organic coordination ligand species of the 2D porous non- noble transition metal organic framework (MOF) is selected from the group consisting of: dicarboxylic acids, tricarboxylic acids, tetracarboxylic acids, imidazolates, bisimidazolates, triazolates, pyridinates, and combinations thereof. Desirably, the at least one basic organic coordination ligand species of the 2D porous non-noble transition metal organic framework (MOF) is an alkyl imidazole, preferably 2- methylimidazole. Desirably, the acidic non-noble transition metal ion of the MOF is selected from the group consisting of: Co, Ni, Fe, Cu, Zn, Zr, Ti, and combinations thereof. In one embodiment, the acidic non-noble transition metal ion is a combination of two of Co, Ni, Fe, Cu, Zn, Zr, & Ti.

[0043]

[0034] Desirably, the bridging MIM-X-MMOF ionic bonds are selected from Fe-X-Co; W-X-Co; W-X- Fe, particularly where X is O, S, N or P. X is O is preferred. Desirably, the bridging MIM-X-MMOFionic bonds are selected from Fe-O-Co; W-O-Co; W-O-Fe; Fe-N-Co; W-N-Co; W-N-Fe; Fe-S-Co; W-S-Co; W-S-Fe; Fe-P-Co; W-P-Co; and W-P-Fe. Desirably, the heterointerface further comprises one or more additional bonds selected from metal-0 bonds, metal-N bonds, B-0 bonds, and P-0 bonds.

[0044]

[0035] Desirably, the heterostructure assembly is selected from Fe203 / CoMOF; WO3.B2O3 / CoMOF; B-Fe2O3 / NiMOF; B-Fe-Ti3C2Tx / CoMOF; B-FeS2 / NiMOF; B-C0S2 / NiMOF; B-NiS2 / NiMOF; and P-W8O26.B2O31 FeMOF. Desirably, the CoMOF is ZIF-67. ZIF-67 (CoMOF) comprises cobalt (Co2+) ions that are tetrahedrally linked with 'N' atoms from 2-methylimidazole's aromatic rings. ZIF-67 is preferred due to its high porosity, outstanding catalytic activity, superior metal-ligand charge transfer capability, and huge surface area.

[0045]

[0036] Desirably, the process involves using a ratio of the inorganic solid-state material to metal salt that ranges from 1 : 3-10 wt%, preferably 1 : 5.5 wt%. Desirably, the process involves using a ratio of metal salt to organic ligand that ranges from 1 -6 : 1 wt%. For preferred Ni systems that ratio is 2.2 : 1 wt% and for preferred Co systems, that ratio is 3.5 : 1 wt%.

[0046]

[0037] Desirably, the process involves using a solvent for the dispersion which is a slightly polar solvent, such as an alcohol including methanol or ethanol, or an aqueous alcohol solution, preferably methanol in water. Desirably, the basic organic ligand species is provided to the dispersion of inorganic template in solution form, for example, as an aqueous solution of the organic ligand species. Desirably, the acidic non-noble transition metal ions are provided in solution form, for example, an aqueous solution of the acidic non-noble transition metal ions is provided to the dispersion of the inorganic template and organic ligand species. Desirably, the acidic non-noble transition metal ions of the 2D porous non-noble transition metal organic framework (MOF) are preferably derived from a soluble salt, including a chloride or nitrate salt, of the metal. Desirably, the acidic non-noble transition metal ions of the 2D porous non-noble transition metal organic framework (MOF) are Co ions which are preferably derived from a salt including cobalt nitrate (e.g., Co(N03)2.6H20). Desirably, the at least one basic organic coordination ligand species suitable for MOF fabrication is associated with surfaces of the 2D ultrathin sheets of inorganic solid state material template.

[0047]

[0038] In a third aspect, the invention provides a product obtained by or obtainable by the process of the second aspect.

[0048]

[0039] In a fourth aspect the invention provides a vertical layered A-[B-A]ntype heterostructure assembly of a 2D ultrathin inorganic solid-state material (‘B’) which is a non-noble transition metal oxide (TMO) and a non-noble metal organic framework (MOF) material (‘A’) that covers surfaces of each layer of the TMO in the assembly in a parallel arrangement to the 2D ultrathin inorganic solid- state material (‘B’), wherein n is a positive integer, wherein the non-noble transition metal oxide (TMO) comprises or is doped with one or more non- noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo & W; wherein heterointerfaces of the assembly comprise bridging MIM-X-MMOFionic bonds that stabilise heterointerfaces between the TMO and the MOF material, where bridging X is -O-, -S-, -Se-, -C-, -N-, -P-, or -B- and MIMis a metal in the TMO, and MMOFis a metal in the non-noble MOF; and optionally, wherein the TMO is doped with one or more non-metal dopants selected from the group consisting of: B, P, C, N, O, Si, S, As and Se, preferably one or more of: B, S, P, or B and P.

[0040] Desirably, the TMO is an iron oxide such as Fe20s or a tungsten oxide such as WO3; and optionally, the TMO is doped with one or more non-metal dopants selected from B; P; and B in combination with P. Preferably, the Fe20s or WO3 is doped with B and / or P. In some examples, the TMO is B-Fe20s, WO3.B2O3, or P-WsO26-B2O3.

[0049]

[0041] In a fifth aspect the invention provides a vertical layered A-[B-A]ntype heterostructure assembly of 2D ultrathin inorganic solid-state material (‘B’) which is a non-noble transition metal sulfide (TMS) and a non-noble metal organic framework (MOF) material (‘A’) that covers surfaces of each layer of the TMS (‘B’) in the assembly in a parallel arrangement to the 2D ultrathin inorganic solid- state material (‘B’), wherein n is a positive integer, wherein the non-noble transition metal sulfide (TMS) comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; wherein heterointerfaces of the assembly comprise bridging MIM-X-MMOFionic bonds that stabilise heterointerfaces between the TMS and the MOF material, where bridging X is -O-, -S-, -C-, - N-, -P-, or -B-, and MIMis a metal in the TMS, and MMOFis a metal in the non-noble MOF; and optionally, wherein the TMS is doped with one or more non-metal dopants selected from the group consisting of: B, P, C, N, O, Si, S, As and Se, preferably one or more of S, B, and P, or B and P.

[0042] Desirably, the TMS is an iron sulfide such as FeS2 or a nickel sulfide such as NiS2 or a cobalt sulfide such as C0S2; and optionally, wherein the TMS is doped with one or more non-metal dopants selected from the group consisting of: S, B, and P; or B in combination with P. Preferably, the TMS is FeS2, C0S2, or NiS2 is doped with B and / or P. In some examples, the TMS is B-FeS2, B-C0S2, or B- NiS2. Suitably, the range of concentration of B in the TMS, such as B-FeS2, ranges from 0 to 18% w / w%. For example, the B concentration in B-FeS2 may be (6mg_B) 2 w / w%; (12mg_B) 4.28 w / w%; (20mg_B) 7.14 w / w%; and (50mg_B) 17.85 w / w%.

[0050]

[0043] In a sixth aspect the invention provides a vertical layered A-[B-A]ntype heterostructure assembly of 2D ultrathin inorganic solid-state material (‘B’) which is a non-noble transition metal doped MXene or a non-noble transition metal doped functionalised MXene material (‘B’) and a non-noble metal organic framework (MOF) material (‘A’) that covers surfaces of each layer of MXene in the assembly in a parallel arrangement to the 2D ultrathin inorganic solid-state material (‘B’), wherein n is a positive integer, wherein the MXene comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; wherein heterointerfaces of the assembly comprise bridging MMXene-X-MMOFionic bonds that stabilise heterointerfaces between the MXene and the MOF material, where bridging X is -O-, -S-, - Se-, -C-, -N-, -P-, or -B-, and MMXeneis a metal in the MXene, and MMOFis a metal in the non-noble MOF; and optionally, wherein the MXene is doped with one or more non-metal dopants selected from the group consisting of: B, P, C, N, Si, S, As and Se, preferably one or more of: B, S, P; or B and P.

[0051] In a preferred embodiment, the MXene is in the form of freestanding or delaminated 2- dimensional (2D) sheets of at least a dual modified functionalised MXene material, Mn+iXnTx, where n is 1 , 2 or 3, and M is an early transition metal (Group 3 to 6 metal), X is C and / or N, and Txis a hydrophilic surface functional group selected from one or more of O, F, OH, and Cl, wherein

[0052] - the MXene material is functionalised to replace the majority of the surface functional groups with O;

[0053] - the MXene material is modified to replace a portion of M atoms with one or more doped non- metal atoms selected from B, N, C, P, Si, S, As and Se; and

[0054] - the MXene material is further modified to replace at least one surface functional group and / or at least one M with one or more electrochemically active metals selected from the group consisting of: Fe, Ni and Co.

[0055]

[0044] Desirably, the early transition metal M is selected from one or more of: Sc, Ti, V, Cr, Y, Zr, Nb, Mo, Hf, Ta and W. Preferably, the electrochemically active metal is Fe. Suitably the dual modified functionalised MXene is a dual modified functionalised titanium carbide (TMC), TisC2Tx, wherein Txis a surface functional group selected from one or more of O, F, OH, and Cl. Preferably, the dual modified functionalised MXene is titanium carbide (TMC), TisC2O2. Suitably, the dual modifications are (i) non- metal dopant is B and (ii) electrochemically active metal Fe. Preferably, the dual modified functionalised MXene is a dual modified O-functionalised titanium carbide (TMC), TisC2O2, and wherein the non-metal dopant is B and the electrochemically active metal is Fe. In this case the material is designated ‘FBT’. Preferred B-doped MXenes comprise a B ratio range for FBT of from 10 - 30 w / w% B. This is particularly the case for B-doped dual modified titanium carbide (TMC), such as B- doped and Fe- doped TisC2O2.

[0056]

[0045] Desirably, the MXene is V2CTX, Ti2CTx, Mo2CTxor the functionalised MXene is TisC2O2, and wherein the MXene is doped with one or more non-metal dopants selected from the group consisting of: B, P, C, N, O, Si, S, As and Se, preferably B, P, B and P.

[0057]

[0046] In some examples, the non-noble transition metal dopant of the non-noble transition metal doped MXene or non-noble transition metal doped functionalised MXene is Fe, Co, Ni, Cr or Cu and wherein the MXene is doped with one or more non-metal dopants for chloride shielding polyanion formation, wherein preferably the dopants are selected from B and P.

[0058]

[0047] Preferably, the 2D ultrathin inorganic solid-state material (‘B’) is Fe-B-Ti3C2Tx, where X is O, such as Fe-B-Ti3C2C>2. In many preferred examples of various aspects above, the MOF is a CoMOF such as ZIF-67, NiMOF, ZnMOF or FeMOF. In some embodiments, the MOF is a bimetallic MOF, such as an FeCo-MOF (“FeCo-OF”), NiCo-MOF, ZnFe-MOF, ZnNi-MOF, NiFe-MOF, or ZnCo-MOF.

[0059]

[0048] In a seventh aspect, invention provides a use of a vertical 2D heterostructure (2D HS) layered inorganic-organic assembly according to the first aspect, or of the third to sixth aspects as an electrocatalyst, preferably for oxygen evolution reaction (OER), especially in alkaline water or alkaline seawater (SW), such as 6 M KOH or 6 M SW. Desirably, the OER is selective OER by inhibiting CER.

[0049] In the aspects described above, the growth of the non-noble metal organic framework (MOF) material (‘A’) over surfaces of each layer of the inorganic solid-state material (‘B’) is preferably parallel growth, that is, growth right across surfaces of the inorganic solid-state material (‘B’) where (‘A’) coverage is in a parallel arrangement to the 2D ultrathin inorganic solid-state material (‘B’). See schematic in Figure 1 (a). This is in contrast to perpendicular growth of the MOF on the inorganic material, where much of the surface of the inorganic material would be exposed. The parallel growth means that substantially all of the surfaces of the inorganic are covered by the MOF component. Therefore, there are minimal, if not no, areas of exposed inorganic material (‘B’) in the assembly.

[0060]

[0050] In an eighth aspect, the invention provides a process involving inclusion of at least one non- metal atom doped atom and at least one electrochemically active metal doped atom in an O- functionalised MXene in a heterostructure to increase the activity of a MXene based OER electrocatalyst in the heterostructure, wherein the one or more doped non-metal atoms are selected from B, N, C, P, Si, S, As and Se; and wherein the one or more electrochemically active metals are selected from the group consisting of: Fe, Ni, Co, Cu, Cr, Mo and W, preferably Fe, Ni and Co.

[0061]

[0051] In a ninth aspect, the invention provides a method of forming a 2D / 2D heterostructured OER electrocatalyst that is selective for OER over CER, comprising forming a vertical layered A-[B-A]ntype heterostructure of a 2D ultrathin inorganic solid-state material (‘B’) and a non-noble metal organic framework (MOF) material (‘A’) that covers surfaces of each layer of the inorganic solid-state material (‘B’) in the assembly in a parallel arrangement to the 2D ultrathin inorganic solid-state material (‘B’), wherein n is a positive integer, wherein heterointerfaces of the assembly comprise bridging MIM-X- MMOFionic bonds that stabilise heterointerfaces between the inorganic solid-state material (‘B’) and the MOF material, where bridging X is -O-, -S-, -Se-, -C-, -N-, -P-, or -B-, MIMis a metal in the inorganic solid-state material, and MMOFis a metal in the non-noble MOF; wherein the inorganic solid-state material comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; and further comprises one or more of: metal-0 (oxide), metal-N (nitride), metal-C (carbide), metal-C / metal N (carbonitride); metal-S (sulfide) and / or M-P (phosphide) functionalities, and optionally wherein the inorganic solid-state material is doped with one or more non-metal dopants selected from the group consisting of: B, P, C, N, O, Si, S, As and Se.

[0062]

[0052] In a tenth aspect, the invention provides a method of improving the OER catalytic selectivity of a 2D / 2D electrocatalyst over CER, wherein the electrocatalyst comprises a vertical layered A-[B- A]ntype heterostructure assembly of a 2D ultrathin inorganic solid-state material (‘B’) and a non-noble metal organic framework (MOF) material (‘A’) that covers surfaces of each layer of the inorganic solid- state material (‘B’) in the assembly in a parallel arrangement to the 2D ultrathin inorganic solid-state material (‘B’), by forming a bridging MIM-X-MMOFionic bonds at the heterointerface wherein the heterointerface stabilising the heterointerfaces between the inorganic solid-state material (‘B’) and the MOF material, where bridging X is -O-, -S-, -Se-, -C-, -N-, -P-, or -B- and MIMis a metal in the inorganic solid-state material, and MMOFis a metal in the non-noble MOF; wherein the inorganic solid-state material comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; and further comprises one or more of: metal-0 (oxide), metal-N (nitride), metal-C (carbide), metal-C / metal N (carbonitride); metal-S (sulfide) and / or M-P (phosphide) functionalities, wherein n is a positive integer, and optionally wherein the inorganic solid-state material is doped with one or more non-metal dopants selected from the group consisting of: B, P, C, N, O, Si, S, As and Se.

[0063]

[0053] Further aspects of the invention appear below in the detailed description of the invention.

[0064] Brief Description of the Figures

[0065]

[0054] Embodiments of the invention will herein be illustrated by way of example only with reference to the accompanying drawings in which:

[0066]

[0055] Figure 1 illustrates formation of ionically anchored heterostructures, a) Schematic representation of the synthesis of Fe2C>3 / CoMOF by solid-liquid interfacial growth technique at room temperature b) TEM image (with inset AFM image) of Fe2C>3 / CoMOF, c) Elemental mapping using TEM of C, N, O, Co and Fe in Fe2C>3 / CoMOF, d) HR-TEM image of Fe2C>3 / CoMOF, e) XRD patterns of Fe2C>3, CoMOF and Fe2C>3 / CoMOF, Top and side views of the optimised CoMOF fragment on the (f, g) Fe-terminated and h) O-terminated Fe20s(001 ) surfaces;

[0067]

[0056] Figure 2 illustrates deconvoluted XPS spectra of a) Co 2p, b) N 1 s, and c) O 1 s present in the Fe2C>3, CoMOF and Fe203 / CoMOF, NEXAFS spectra of d) Co L-edge, e) N K-edge and f) O K-edge;

[0057] Figure 3 illustrates electrocatalytic OER activity of Fe20s, CoMOF and CoMOF / Fe203, lrO2 and Ni-Foam: a) OER polarisation curves recorded by the back-sweep scan method in alkaline seawater electrolyte, b) Plot of current densities vs overpotentials of CoMOF / Fe203, c) Comparison of overpotential of Fe20s, CoMOF and CoMOF / Fe203, lrO2 and Ni-Foam, d) corresponding Tafel plots, e) plots showing double-layer capacitance (Cdi) in alkaline seawater electrolyte of Fe20s, CoMOF and CoMOF / Fe203, f) Linear sweep voltammetry curves before and after the stability test, g) current-time (l-t) curve of heterostructure in seawater electrolyte at 1 .366 V (E vs. RHE) (Tower off and system started again on same setup), and top and side views of the optimised h) CoMOF / Fe2O3(001 )-CI and i) CoMOF / Fe2O3(001 )-OH (the black dashed circle identifies Cl and OH adsorbates).

[0068]

[0058] Figure 4 illustrates a) TEM image, b) HR-TEM image and elemental mapping using TEM of Fe, Co, O, and N, of Fe203 / CoMOF, NEXAFS spectra of c) Co L-edge, and d) O K-edge, of Fe203 / CoMOF, and deconvoluted XPS spectra of e) Co 2p of Fe203 / CoMOF after electrochemical testing;

[0069]

[0059] Figure 5 illustrates construction of ionically bonded heterostructure and its growth mechanism, (a) Schematic illustration of the synthesis of MOF@WB. (b) TEM (inset: SAED) and (c) HRTEM of MOF@WB, (d) XRD of WB, Co-MOF, and MOF@WB. Top and side views of the optimized Co-MOF fragment on the (e) W0s(001 ), (f) B-doped W0s(001 ), and (g) B20s(001 ) surfaces using DFT and AIMD simulations. The atoms within the shaded region were fixed in all calculations. The green circle indicates the location of the B atom;

[0070]

[0060] Figure 6 illustrates bonding and electronic structure of the heterostructure. WB, Co-MOF, and MOF@WB deconvoluted XPS spectra for (a) Co 2p, (b) W 4f, and (c) N 1 s. NEXAFS spectra of (d) Co L-edge (e) B K-edge, and (f) O K-edge;

[0071]

[0061] Figure 7 illustrates OER performance and durability testing of MOF@WB-O.21 in seawater electrolyte in a 3-electrode configuration; OER LSV curves (a) WB of different concentrations in 6M KOH, (b) WB-0.21 , Co-MOF, MOF@WB-O.21 , and lrO2 in 6M alkaline seawater electrolyte, (c) MOF@WB-O.21 in 6M seawater electrolyte using booster (Inset: contact angle measurements). WB- 0.21 , MOF@WB-O.21 , and lrO2 (d) Mass activities in 6M seawater electrolyte, (e) Tafel plots in 6M KOH and seawater electrolyte, (f) Double layer capacitance (Cdi) curves for different WB, Co-MOF, and heterostructures concentrations, (g) Stability tests for MOF@WB-O.21 for 1100 h in 6M seawater electrolyte. (Inset: stability testing after leaving the electrode idle for 5 days and post-stability TEM). Side view of the optimized Co-MOF / WO3(001 ) with adsorbed (h) OH- and (i) Ch and Co- MOF / B2O3(001 ) with adsorbed (j) OH- and (k) Ch. The atoms within the shaded region were fixed atoms in all calculations, (h) and (i) were rotated slightly around the z-axis to improve the visibility of the adsorption. Note: Mass activities were calculated, taking into account other ink components. All the results are reported without iR correction;

[0072]

[0062] Figure 8 illustrates structural and morphology retention of MOF@WB-O.21 post stability. Poststability characterizations of MOF@WB-O.21 (a) TEM (inset SAED) and (b) HRTEM. Deconvoluted XPS spectra for (c) W 4f, (d) O 1s, (e) Co 2p, and (f) N 1s. NEXAFS spectra of (g) O K-edge, (h) Co L-edge, and (i) N K-edge;

[0073]

[0063] Figure 9 illustrates a) Schematic illustration for the synthesis of ZIF-67 / FBT with ionic bonds at the heterointerface, b) XRD patterns of TisC2O2, FBT and ZIF-67 / FBT. c) TEM image of FBT. d-e) SAED of FBT. f) HRTEM image of FBT with lattice fringes (insert), g) TEM image of ZIF-67 / FBT. h) HRTEM image of ZIF-67 / FBT;

[0074]

[0064] Figure 10 illustrates Fitted high-resolution XPS spectra of a) O 1 s of ZIF-67, FBT and ZIF- 67 / FBT, b) N 1 s of ZIF-67 / FBT, c) Co 2p of ZIF-67 / FBT, and d) Fe 2p of FBT and ZIF-67 / FBT. NEXAFS spectra for e) N K-edge of ZIF-67 and ZIF-67 / FBT, f) Co L-edge of ZIF-67 and ZIF-67 / FBT, and g) Fe L-edge of FBT and ZIF-67 / FBT;

[0075]

[0065] Figure 11 illustrates a) FBT, ZIF-67 / FBT, ZIF-67, RuC>2 and ZIF-67+FBT Polarization curves in 6 M alkaline seawater electrolyte, b) Comparison of the overpotentials achieved by FBT, ZIF- 67 / FBT, ZIF-67, RuC>2 and ZIF-67+FBT in 6 M alkaline seawater electrolyte at different current densities, c) Tafel slopes of the electrocatalysts from the polarization curves in a), d) Nyquist plots of OER on TisC2O2, BT, FT, FBT and ZIF-67 / FBT. e) Double layer capacitance (Cdi) curves of TisC2O2, BT, FT, FBT and ZIF-67 / FBT calculated from CV curves in S12-16. f) Current density comparison at 480 mV (equilibrium potential of CER) in an alkaline medium, g) Photo showing the colour of solutions obtained for different samples after the CER test, h) Gas chromatographs obtained after the CER test;

[0066] Figure 12 illustrates a) Long-term Stability test of FBT and ZIF-67 / FBT in 6 M alkaline seawater, b-e) Post stability analysis b) TEM of FBT, c) TEM of ZIF-67 / FBT, d) NEXAFS of FBT and ZIF-67 / FBT, and e) XPS of FBT and ZIF-67 / FBT;

[0076]

[0067] Figure 13 illustrates (a) Schematic presentation of formation of B-FeS2 / MOF HS. The TEM images of (b) B-Fe2Os, (c, the inset shows SAED pattern) B-FeS2 and (d, the inset shows SAED pattern) B-FeS2 / MOF. (e) SEM image of B-FeS2 / MOF. HRTEM images of (f) B-FeS2 and (g) B- FeS2 / MOF. (h) EDS elemental maps of Fe, Ni, S, B and O for B-FeS2 / MOF;

[0077]

[0068] Figure 14 illustrates XRD pattern of FeS2, B-FeS2, B-FeS2 / MOF (a), NEXAFS spectra of (b) Fe L-edge, (c) S K-edge, (d) Ni L-edge. Deconvoluted XPS spectra of (e) Fe 2p, (f) S 2p, (g) B 1s and (h) Ni 2p of B-FeS2and B-FeS2 / MOF;

[0078]

[0069] Figure 15 illustrates (a) B concentration of samples based on FeS2 at different current densities, (b) The OER polarization curves of various electrodes (B-Fe2Os, FeS2, B-FeS2, lrC>2, Ni- MOF, NF, FeS2 / MOF and B-FeS2 / MOF) in alkaline seawater (c) OER Tafel slopes are obtained from the OER polarization curves, (d) the current-time (l-t) curves of B-FeS2 / MOF, FeS2 / MOF and lrO2. The inset (d-l) shows OER polarization curves of the FeS2 / MOF heterostructure before and after stability in the same alkaline electrolyte, (e) The OER polarization curves of the heterostructure before and after stability in the same alkaline electrolyte, (f) EIS spectra for FeS2, B-FeS2 and B-FeS2 / MOF electrodes in alkaline seawater, (g) ECSA is calculated from double-layer capacitance for different catalysts at various scan rates.

[0079]

[0070] Figure 16 illustrates MOF@WPB, Fe-MOF, and WPB (a) XRD (inset: XRD of WB), (b) Raman Spectra. Deconvoluted XPS spectra for (a) W 4f, (b) Fe 2p, (c) B 1 s and (d) N 1s. (e) Fe L-edge, (f) B K-edge, (g) O K-edge, and (h) N K-edge NEXAFS spectra;

[0080]

[0071] Figure 17 illustrates WPB (a) TEM (inset: SAED), (b) HRTEM. MOF@WPB (c) TEM (inset: SAED), (d) HRTEM;

[0081]

[0072] Figure 18 illustrates (a) OER overpotentials for different B concentrations in 6 M KOH at current densities of 0.01 , 0.1 , 0.5, and 1 A cm2. MOF@WPB-O21 , WPB-0.21 , WB-0.21 , and lrO2(b) OER LSV plots and (c) overpotentials at current densities of 0.1 , 0.5, 1 , and 1 .5 A cm-2in 6M seawater (d) Tafel plots, (e) double layer capacitance and (f) electrochemical impedance spectroscopy in seawater, (g) MOF@WPB-O.21 stability tests in seawater;

[0082]

[0073] Figure 19 illustrates MOF@WPB-O.21 post-stability (a) XRD, (b) EDX, (c) TEM, and (d) HRTEM. Deconvoluted (e) W 4f, (f) Fe 2p, (g) O 1s, and (h) P 2p spectra for MOF@WPB-O.21 ;

[0083]

[0074] Figure 20 illustrates (a) LSV of different catalysts in 6 M SW, (b) Comparison of overpotentials at different current densities for the catalysts (c) Stability in 6 M SW at 1 .611 V;

[0084]

[0075] Figure 21 shows (a) XRD patterns of the synthesised materials, (b) TEM image of FeCo- OF / FBT. (c) HRTEM image of FeCo-OF / FBT. (d-f) Fitted high-resolution spectra of d) Fe 2p, e) O 1s, and f) Co 2p of FBT, FeCo-OF and FeCo-OF / FBT; and

[0085]

[0076] Figure 22 shows (a) Back sweep polarisation curves of FBT, FeCo-OF, heterostructures, Ni Foam, and RuC>2. (b) Overpotential comparison of FBT, FeCoOF, FeCo-OF / FBT, RuO2. (c) Tafel plots, (d) Amperometry i-t Curves of FeCo-OF / FBT. (e) Full alkaline seawater splitting polarisation curve and (f) Full alkaline seawater splitting stability test for PtC / / FeCo-OF / FBT-15.

[0086] Detailed Description of the Invention

[0087]

[0077] Strong ionic connections are difficult to create between dissimilar species (e.g., inorganic and metal-organic materials, such as a MOF). The inventors’ discovery that they can be generated via a room temperature (RT) interface process is a significant advance in the art. Furthermore, preferred resultant catalytically active materials with ionic interconnections between dissimilar materials, advantageously favour OH- adsorption, repel Ch, exhibit improved OER performance, even at low overpotentials, and / or provide operation stability even at high current densities and high overpotentials under corrosive seawater splitting conditions, while avoiding catalyst corrosion by inhibiting anodic chlorine oxidation reactions.

[0088]

[0078] In initial work, a first heterostructure was formed having ionic bonds between Fe20s sheets and CoMOF using the process of the invention. The new 2D / 2D HS material Fe2C>3 / CoMOF (MOF@FEO) displays a current density of 0.5 A cm-2at an overpotential of 310 mV for OER in seawater. This result is ~2 times higher performance than commercially used lrO2 in OER. Furthermore, the MOF@FEO catalyst performs stably for 900 h OER in 6 M KOH real seawater. Stably means ±10% of the achievable current density. The ionic connections are important for providing the high OER activities and stable performance described herein.

[0089]

[0079] In subsequent work, to further impact activity, the surface chemistry of other transition metal inorganic template materials were pre-modified by doping with one or more electrochemically active metals such as Fe, and / or doping with one or more non-metals like B and / or P to tune the inorganic material’s properties through the introduction of one or more of atom defects, oxygen vacancies, polyanion functionalities or enhanced active sites I directed synthesis guide creation on the 2D inorganic sheets, prior to parallel MOF growth on the surfaces of the inorganic template in the assembly. In some examples, introduction of electrochemically active metals such Fe into inactive or less active inorganic materials such as Ti MXenes, allowed creation of a direct ionic bond connection at the heterointerfaces of the dissimilar materials through formation of an ionic bond between the doped metal and the metal of the MOF. The ionic bond includes the electrochemical active metal at the stabilised heterointerface. Generally speaking, the polyanion functionalities formed by non-metal doping tends to shield the HS from Cl- present in vast amounts in real seawater and thus allow operation without competing chloride reactions even at higher overpotentials associated with higher OER current densities.

[0090]

[0080] Ionic bond creation is believed to provide one or more of at the following benefits: i) activating the electrocatalyst potential of the inorganic material, even where more electrochemically metals are introduced into a material, and ii) manipulating the inorganic material surface chemistry to assist the formation of the strong ionic connections at heterointerfaces required to develop higher-performance catalysts; (iii) forming Cl- shielding surface functionalities. For the materials described herein, the approach has been very successful. For example, the MOF / WO3.B2O3 of the invention delivers 10 times better performance than commercial lrC>2 for OER and is stable for over 1000 hours without chlorine oxidation in seawater.

[0091]

[0081] A particular problem addressed by the process of the invention is to enable challenging parallel growth of MOF over surfaces of the 2D inorganic template material in a way that allows ionic bond formation and exceptional surface coverage of the inorganic material with the MOF, while resulting in a vertically stacked A-[B-A]narrangement of the dissimilar materials. This is direct contrast to the perpendicular oriented MOF growth on rGO demonstrated by Liu, which fabrication involves an interface assisted synthetic process but with different starting materials and different conditions. The parallel overlapping of dissimilar sheets enabled by the present process of the invention advantageously and unexpectedly provides substantially complete coverage of the inorganic material by the MOF which is stably held in position through formation of ionic binds at heterointerfaces of the dissimilar materials. The arrangement provides access to extra and / or enhanced active sites and a maximum number of connections between dissimilar sheets. Overall, the process reported herein produces materials exhibiting exceptional performance for selective anodic reactions, particularly OER, even in natural I fresh seawater, a most challenging substrate. As will become clear herein, the materials of the invention provide many industrially useful electrocatalysts operable at high current densities of >500 mAcmr2for hundreds of hours and in some cases thousands of hours. Advantageously, when used as an electrocatalyst in seawater, preferred materials avoid chlorine gas evolution or hypochlorite formation by supporting the desirable excellent performance for OER not only at lower overpotentials that inhibit chlorine oxidation reactions at the anode, but at higher overpotentials resulting from OER at higher current densities due to the shielding effect of many of preferred materials of the invention. The materials are particularly stable during operation for a much longer time than conventionally used electrocatalysts as a result of the ionic bonds and in preferred materials, resistance to Cl- based corrosion and deactivation of active sites.

[0092]

[0082] Introduction of preferred dopant atoms into the inorganic material result in tuning I generation of defects (e.g. vacancies, dislocations, misalignments, holes, and / or new surface functionalisations) in the uppermost surface layers of the inorganic material in a way that provides sites that facilitate adsorption of the MMOFmetal around which the MOF is later assembled. For example, the doped atom may relax away from the uppermost surface layers to form a type of vacancy or gap via the doped atom relaxing away from the surfaces of the inorganic material, leaving surface areas which can act as a guided surface which directs or otherwise facilitates MMOFadsorption and indeed ionic bond formation between MIMand MMOF, e.g., through oxygen bridging, as well as subsequent MOF growth. Furthermore, the resultant ionic bonding can prevent leaching of catalytically active sites, hence ensuring long-term structural stability of the material, and providing for enhanced OER activity in application described herein. For example, in the ZIF-67 / FBT material of the invention, the electrochemically active Fe doped atoms which replace some of the Ti metal facilitates ionic bond formation and the electrochemically active Fe is prevented from leaching out of the material as a result of the ionic bond formation. The benefits of both dissimilar materials work in tandem to enhance OER activity, inter alia, due to rapid electrical conductivity for fast redox process during OER in both alkaline seawater and deionised (DI) water.

[0093]

[0083] For many materials, doping, electrochemically active metal doping, as well as P and B doping, and B doping in particular, facilitates the formation of strong ionic bods at the heterointerface between the dissimilar materials, by tuning the electronic I surface properties of the inorganic material in a way that enables strong ionic bond formation which supports charge redistribution between the dissimilar materials which enhances OER activity. The unique heterointerface generated between the 2D materials support rapid charge transfer during electrochemical process. Further, doping improves the electrical conductivity of the material and assists in producing covalent / ionic bonds between dissimilar sheets. Preferred dopants create surface passivation by generating negatively charged polyanions on the inorganic material surface which form a Ch repulsion layer on the inorganic material surfaces to prevent or inhibit chloride oxidation reactions. Preferred polyanions include borate, sulfate, phosphate, carbonate, and combinations thereof, etc. It will be understood that these species may be formed by doping with one or more of B, S, P, N, O, C atoms, or combinations thereof into the inorganic material. These polyanions may provide anti-corrosive properties to the material as the polyanions modulate hydroxyl anions at the interface, e.g., via hydrolysis which repels the chloride anions present in seawater, thus preventing corrosive chlorine chemistry at the anode. Preferred combinations include P and B, or Fe and B.

[0094]

[0084] B doping produces a surface in which B relaxes away from the uppermost layers or sits underneath the top surface layer exposing the top surface of the inorganic template and providing much needed vacancies that guide or direct adsorption of MMOFatoms for growth of the metal organic component and formation of strong ionic bonds between the materials. Such defects / doping allowing ease of construction of the 2D morphology and interlinked heterostructure. The borate species which result from B doping are a weak Lewis acid that acts as a local OH- modulator which preferentially binds OH- to Cl', as the B-OH sites formed result in a layer of Lewis acid protection on the inorganic material.

[0095]

[0085] Furthermore, still the hydrophilic MOF sheets enable quick wetting and favour OH' adsorption, while shielding the active sites / heterointerface from Cl' ions. The B doping assists in better adsorption of MOF to initiate stronger connection and a unique interface between the dissimilar materials. The B doping modulates the electronic properties of the inorganic material and also assists in creating borate species at the interface that results in preferential OH' adsorption over Cl' adsorption. Fe doping into certain inorganic materials, particularly MXenes, enables for the first time, the ability to form ionic bonds in a heterostructure material, e.g., Fe-O-Co bonds in the case of the ZIF-67 / FBT material described herein. Such electrochemically active metal doping breaks the otherwise neutrality of the MXene material and allows its use as an electrocatalyst.

[0096]

[0086] In summary, advantages of doping include one or more of: higher electrical conductivity, assistance in producing covalent / ionic bonds among dissimilar sheets, and yielding Lewis acid species such as borate, a weak Lewis acid, at the interface to act as a local OH' modulator / to regulate hydroxyl ion adsorption in preference to Cl' adsorption; induced changes in surface area, facilitating charge transfer between adjacent carbon (C) atoms of the MOF which enhances catalytic activity; and / or better adsorption of the MOF to initiate stronger connection and a unique interface between two dissimilar materials.

[0097]

[0087] Generally, the new materials are useful electrocatalysts operable in DI water and / or fresh seawater at high current densities of over 500 mA / cm2e.g., as determined by linear scanning voltammetry. Advantageously, when used as an electrocatalyst in fresh seawater, the materials readily avoid chloride ion oxidation (such as chlorine gas or hypochlorite formation) by supporting the necessary excellent current densities for OER at not only lower overpotentials of <480 mV, but also at much higher current density of >1 A cm-2which are typically associated with higher overpotentials of >480 mV where completing chloride ion oxidation reactions are problematic. Due to the excellent Ch repulsion ability of preferred catalysts of the invention, chlorine oxidation reactions and associated by chloride product formation is inhibited even at high current densities and overpotentials >480 mV. Together with the stability of the material, avoidance of catalyst corrosion and supported operation for a much longer time than conventionally used electrocatalysts, such as lrC>2, is achieved. In some embodiments, selective OER at very high current densities of up to 1 .75 A cm-2are achievable. This is a significant advance in the art.

[0098]

[0088] Advantageously, preferred materials of the invention operate extremely well at low and high overpotentials while at same time providing stability due to the strongly connected heterointerface which stabilises performance, while the ionic connections between dissimilar sheets facilitate fast change transfer which improves catalyst OER performance. The 2D ultrathin nature of the inorganic component provides flexible support to the strongly bound MOF layers which prevents them from agglomeration, and allows good parallel orientation surface coverage by the MOF to provide vertically stacked I parallel layers of the dissimilar materials, that allow fast access of the liquid electrolyte to the active sites, resulting in higher current density as compared to individual catalyst component materials. The excellent activity is thought to result from synergistic coupling effects of the ionically dissimilar bonded components, manipulating the electronic structures by facilitating the charge delivery and fast charge and mass transport between dissimilar interfaces. Overall, the durable performance of the heterostructures at higher current densities is attributed to the higher surface areas, ease of access to active sites, avoidance of chloride adsorption and chloride byproduct formation which avoids catalyst corrosion / degradation, as well as synergistic effects and robust ionic connections among the heterointerface. The vertical growth of MOF over surfaces of the suitable transition metal inorganic material templates surfaces provides Cl- corrosion resistance to the system along with more catalytic sites overall, and more active catalytic sites. The MOF coverage make it difficult for Cl- to adsorb to the heterostructure, while OH- adsorption is promoted which contributes to enhanced OER performance.

[0099]

[0089] Advantageously, preferred materials are selective towards OER and inhibition of chloride oxidation reaction such as hypochlorite formation in high pH seawater or chlorine gas evolution in acidic seawater.

[0090] B doping into the materials facilitates the formation of particularly strong bonds, including at least ionic bonds at the heterointerface between the dissimilar materials, by tuning the electronic I surface properties of the inorganic material in a way that enables particularly strong bond formation which supports charge redistribution between the dissimilar materials to enhance OER activity. The unique heterointerface generated between the 2D materials supports rapid charge transfer during electrochemical processes. Furthermore, B doping improves the electrical conductivity of the material and assists in producing ionic bonds between the dissimilar sheets. The surface borate and B-OH species which result from B doping act as a weak Lewis acid at the interface, thereby serving as a local OH- modulator. Furthermore, the resultant more hydrophilic MOF sheets enable quick wetting and favours OH' adsorption, while shielding the active sites I heterointerface from Cl- ions. B doping assists in formation of assembly sites that allow better adsorption of the MOF to initiate stronger connection and a unique interface between the dissimilar materials.

[0100]

[0091] Overall, the materials described herein overcome problems experienced in direct electrolytic seawater splitting which arise from impurities and unwanted side reactions that hinder long term performance of anodes in OER reactions.

[0101]

[0092] Thus, the invention provides a process for forming new materials, such as in the form of a vertical layered 2D-2D metal inorganic-X-metal organic heterostructure assembly having strongly bonded heterointerfaces between the dissimilar materials, where the strong bonds include at least ionic bonds between dissimilar atoms / elements of the component materials forming the heterostructure, and wherein X is -O-, -S-, -Se-, -C-, -N-, -P-, or -B-, preferably O. A vertical heterostructure is a material composed of two heterogeneous (dissimilar) 2D (ultrathin / nanosheet) precursor materials which assemble into a composite material involving direct stacking of individual monolayers of each different materials, for example in a A-[B-A]nvertically stacked orientation. A vertical layered 2D-2D metal inorganic-X-metal organic heterostructure assembly is one in which a first 2D (ultrathin) precursor material is a non-noble transition metal (containing ‘MIM’) inorganic material and a second 2D (ultrathin) precursor material is a metal organic framework (MOF), which are dissimilar materials, wherein X is -O-, -S-, -Se-, -C-, -N-, -P-, or -B-, preferably O. Anchoring strong chemical bonds are located at heterointerfaces between the dissimilar layers, and include, for example, ionically bonded MIM-oxygen-MMOFbridges. Examples of such bridges include Co-O-Fe, Co- O-W, Fe-O-Co and Fe-O-W ionic bonds. In other embodiments, there may be additional covalent bonds between metal of one material and other atoms of the MOF, for example, metal to N covalent bonds in the MOF electron donating organic ligand component.

[0102]

[0093] 2D ultrathin sheet means a 2D nanomaterial which is atomically thick that consist of a single to few layers of atoms. Examples include 2D nanosheets or other types of nanomaterials with high aspect ratios. The lateral size of the 2D nanomaterials can reach up to micrometres, and even longer. Suitably both inorganic material and MOF in the heterostructure are in 2D ultrathin sheet form. Desirably, the 2D ultrathin sheet of inorganic solid state material (‘B’) has a thickness of from 2 nm to 20 nm.

[0094] It will be understood that ionic bonds are a particularly strong type of chemical bond. Advantageously, the presence of the strong bonds in the materials obtained by or obtainable by the process used in water splitting applications result in excellent performance and stability of the materials of the invention when used as electrocatalysts, for example, in water splitting applications involving water or natural seawater. Preferred materials are sufficiently active and sufficiently stable during electrocatalysis to enable commercial levels of direct electrolytic seawater splitting (typically >500 mA.cmr2current density for OER). This means the materials typically provide acceptable OER performance for >500 hours, >700 hours, >900 hours, and even >1000 hours for particularly preferred materials. Acceptable performance is production of a steady current density of ±10% of the starting current density for the duration of operation.

[0103]

[0095] The durability in OER performance is a result of the strong bonds which stabilise the assembly, particularly during harsh electrocatalytic applications in electrolytes such as seawater splitting. Furthermore, the organic component is arranged over surfaces of the inorganic component in a way (e.g., a parallel fashion) that covers substantially all active surface areas of the inorganic component, shielding the active surfaces from chloride ions when the material is used as an electrocatalyst for seawater splitting. Most preferred materials are selective towards OER in alkaline natural seawater, while avoiding chloride oxidation reactions, such as CER in acidic pH or hypochlorite formation in high pH media (pH>7). To mitigate by product formation during seawater splitting, preferred electrocatalyst must be active enough to generate industrially high current density (typically >500 mA.cmr2), most preferably at overpotential <480 mV in alkaline medium. The invention addresses the difficulty in producing ionic connections between dissimilar sheets which does not occur readily with conventional HS fabrication methods. Other advantages over prior art methods for forming HSs include easy of scale up, and design of specific surface properties / enhanced active sites.

[0104]

[0096] The features imparting stability and excellent electrocatalytic functionality directly result from application of the process of the invention to appropriate starting material inorganic template precursor which, if desired, may be modified before commencement of the process to include suitable electronic and / or surface feature tuning thereof prior to application of those precursor materials to the synthetic process of the invention. The assembly process constructs the precursor components into a vertical layered 2D-2D metal inorganic-O-metal organic heterostructure assembly of the invention that has strongly bonded hetero in terfaces between the dissimilar materials, where the strong bonds include at least ionic bonds between dissimilar atoms / elements of the component materials forming the heterostructure.

[0105]

[0097] Generally speaking, preferred materials of the invention favour hydroxyl ion adsorption over chloride adsorption thereby promoting OER while inhibiting chlorine reactions to prevent corrosion and / or fouling of the electrocatalyst when used in seawater splitting. Besides their stability, the favourable electroactivity is a direct result of the unique electronic features of the materials and / or heterointerface between the dissimilar precursor materials. Preferred electrocatalysts obtained by or obtainable by the process of the invention result in OER in seawater with a remarkably high current density of at least 0.5 A cm-2at low overpotentials of <480mV at ambient conditions. Particularly preferred catalysts of the invention result in OER at higher current densities of over 0.5 A cm-2at overpotentials of >480V while retaining OER selectivity due to the Ch shielding ability. Such performance is typically significantly higher than that of commercially used lrO2 in seawater splitting. The stability of the materials is much higher than commercially used lrO2 in seawater splitting applications.

[0106]

[0098] Advantageously, the process of the invention provides an economically modest / favourable and yet scalable strategy to regulate the chemical states at a 2D / 2D heterointerface and thereby develop robust catalysts for the practical implication of large-scale direct seawater splitting. Desirably, process is carried out at room temperature and / or atmospheric pressures. Higher temperatures and / or pressures can be applied but are not typically needed. Advantageously, the conditions support provision of the electrocatalytic materials of the invention in scaled up quantities, which is attractive for industrial preparation and use.

[0107]

[0099] The solid-liquid interfacial process of the invention forms unique strongly bonded heterointerface which includes at least ionic bonds between the dissimilar atoms I elements of the component materials forming the heterostructure. More specifically, the process forms a stable 2D / 2D transition metal inorganic material / metal-organic framework heterostructured material with unique MIM- X-MMOF(MIM=metal of material B; MMOF=metal of material A) ionic bonds at the interface, where X is - O-, -S-, -Se-, -C-, -N-, -P-, or -B-, which harvest the individual material’s properties and create additional active catalytic sites as well as ease of charge flow. The development of a heterointerface as the HS material is fabricated changes the oxidation state of the metal nodes of the organic / MOF component to create strong ionic interactions between the dissimilar sheets as the organic / MOF component is assembled over the inorganic template.

[0108]

[0100] In general terms, the process of the invention comprises the following main steps. A choice of a 2D ultrathin / nanosheet non-noble transition metal (containing inorganic material and a choice of a metal organic framework (MOF) is made, that is, a choice of MOF metal in salt form and one or more basic electron donating organic ligands, of which the MOF is composed. Ideally, the non-noble transition metal (containing inorganic material is already in 2D ultrathin / nanosheet sheet form. However, if not, it can be modified into the 2D ultrathin / nanosheet formed which is required to serve as inorganic template in a subsequent solid / liquid assisted synthetic process which builds the vertical layered 2D-2D metal inorganic-X-metal organic heterostructure assembly of interest. For example, a suitable precursor material can be converted to 2D ultrathin / nanosheet morphology by suitable processing to form the required 2D ultrathin / nanosheet non-noble transition metal (containing ‘MIM’) inorganic material. For example, a precursor such as FeOOH can be converted to a-Fe2Os, which corresponds to a polymorph which has the required 2D ultrathin sheet morphology. A precursor 2D Fe2C>3 nanosheet can be prepared by annealing of precursor salt template, e.g., an iron chloride salt template, to form 2D a-Fe2Os nanosheets through a salt template method. Likewise, rod-like WO3 materials may be converted to sheets by heat annealing. Methods to convert to ultrathin / nanosheet form include annealing, acid treatment, base treatment, thermal treatment, combinations thereof, and the like.

[0101] If desired, defects, e.g., unoccupied vacancies, particularly oxygen vacancies, dislocations, misalignments, substitutional impurity / doped atoms, holes between atoms, etc., can be introduced into the precursor inorganic material’s lattice structure. Dopant elements / compounds can be introduced into the inorganic material, for example, electrochemically active non-noble transition metals, and / or non-metals such as one or more of B, S, C, P, etc., to generate one or more such defects, and / or charged points, tuned surfaces and variable electronic properties. Such defects and features may serve as convenient nucleation sites on which the MOF component may be grown, as well as chloride ion repulsion I shielding regions and / or positively attenuated or otherwise enhanced catalytic active sites. Preferred non-noble transition metal (containing ‘MIM’) inorganic material is doped with one or more additional metals to generate new defects or surface functionality. Dopants include but are not limited to B, Fe, and / or P. In some materials, it is preferred that B2O3 or B-OH sites are formed in the inorganic template material. Preferred inorganic materials include 2D ultrathin sheet I nanosheet form of a-Fe2 B modified MXene),

[0109]

[0102] In some embodiments, the inorganic material is a doped MXene. MXenes are 2D materials comprising transition metal carbides, carbonitrides or nitrides, having a layered structure of general formula Mn+iXnTx, wherein M is an early transition metal, X is nitrogen or carbon and Tx is a surface functional group such as -OH, -F, -Cl, or -O. Desirably, the MXene is a transition metal carbide (TMC); a transition metal nitride (TNN) or a transition metal carbon nitride (TMN), or functionalised derivative thereof. As an example, a MXene in the form of a titanium carbide (e.g., TisC2Tx) can be prepared by treatment with a strong reducing agent, such as NaOH, to form TisC2O2 (O functionalised titanium carbide MXene) which can then be electronically tuned / engineered to promote surface termination and to break inherent electrochemical neutrality of the MXene by substituting at least some of the surface functionality with electrochemically active metal Fe, and / or by replacing at least some of the Ti with B (designated herein as ‘FBT). The Fe and / or B doping can be readily achieved by thermal treatment of T3C2O2, whereby B doping for some of the Ti atoms and Fe replacement of some of the surface termination, results in Fe3+in surface termination via electrostatic attractions due to its opposite charges with the surface. The Fe interaction is believed to break van der Waals interactions between adjacent MXene nanosheets, to produce freestanding 2D nanosheet, which serve as the doped inorganic template required for solid / liquid assisted synthetic process described below. Desirably, the twin doping strategy of the precursor MXene (with dopant Fe and B produces new material “FBT” provided as a multiphase thin 2D material including a mixture of three phases of (a) hexagonal shaped titanium carbide (TieCs.ys) and (b) iron titanium oxide (Fe2Ti30g; pseudorutile), and (c) tetragonal shaped titanium boron oxide (TB0.024O2; B-doped rutile). In some examples, the freestanding thin 2D nanosheets maybe associated with clusters of smaller hexagonal nanosheets at some of the edges. It is believed that the B substituted some of the Ti in T3C2O2 matrix create defects in the form of unoccupied vacancies that may act as nucleation sites on which MOF may be grown.

[0103] The solid / liquid assisted synthetic process involves providing a solvent dispersion of the precursor non-noble transition metal (containing ‘MIM’) inorganic material in 2D ultrathin sheet (solid) form as an inorganic template for metal organic framework (MOF) growth thereon.

[0110]

[0104] It will be appreciated that the liquid / solid interface synthesis process involves an already grown 2D sheet solid surface that works as an inorganic template or flat interface that is dispersed in liquid media and at the same time supports the parallel growth of the MOF sheets thereon. In short, the inorganic template is used as an in-plane template that directs MOF growth on the template by limiting the main reaction at the interface. Ideally, the non-noble transition metal, MIM, is an acidic metal capable of accepting donated electrons or at a minimum is potentially capable of associating with Lewis base electron donating organic ligands.

[0111]

[0105] The process then involves electrostatically associating at least one basic electron donating organic ligand with the inorganic material template. As referred to above, a basic electron donating organic ligand species is a Lewis base, or a ligand that contains atoms with a filled orbital containing an electron pair which is not involved in bonding but can form a dative or coordinate bond with a Lewis acid (e.g., metal capable of accepting an electron pair) to form a coordination compound or Lewis adduct with a suitably acidic metal centre. Such donating atoms providing Lewis basicity include O, N, P, S. Examples of suitable basic electron donating organic ligand species that are typically used in metal organic frameworks (MOF) include dicarboxylic acids, tricarboxylic acids, tetracarboxylic acids, imidazolium, and combinations thereof. In preferred examples, at least one of the basic electron donating organic ligand species comprises an imidazolium functionality which comprises donating N: atoms. A preferred example is 2-methylimidazole which is a commonly used electron donating organic ligand for many MOFs. The association between the inorganic template and the electron donating organic ligands is thought to involve attractive interactions such as ionic bonding.

[0112]

[0106] The process then continues to the step of forming the vertical layered 2D-2D metal inorganic- X-metal organic heterostructure assembly of interest by allowing coordination of the basic organic coordination ligands with acidic non-noble transition metal ions (MMOF). Desirably, the acidic non- noble transition metal ions (MMOF) are provided in solution form, for example, an aqueous or slightly polar solution of the metal ions, e.g., an aqueous solution of the acidic non-noble transition metal ions (MMOF) in a suitable salt form. The solution is then provided to the dispersion of inorganic template with which the electron donating organic ligand species has associated. The next step involves allowing growth of a 2D ultrathin deposit of a non-noble transition metal (MMOF) organic framework (MOF) over substantially all surfaces of the inorganic template via allowing MOF assembly through coordination of the basic electron donating organic ligands associated with the inorganic template with the acidic non-noble transition metal ions (MMOF) though donor / dative bond formation. The complexing step is allowed for a duration which results in formation of an assembly having a vertically layered arrangement of 2D ultrathin layers of the dissimilar non-noble transition metal (MIM) inorganic material and the non-noble transition metal organic framework (MOF), whereby the arrangement is stabilised via anchoring strong chemical bonds at heterointerfaces between the dissimilar layers that include ionically bonded MIM-X- MMOFbridges.

[0107] Suitably, the process is a solution-based direct synthesis process. Desirably, the process is carried out at room temperature and / or pressures. Higher temperatures and / or pressures can be applied if desired, but are not typically needed. Advantageously, the conditions support provision of the electrocatalytic materials of the invention in scaled up quantities, which is attractive for industrial preparation and use.

[0113]

[0108] Desirably, the solvent for one or more of the dispersion steps is a slightly polar solvent, such as an alcohol such as methanol or ethanol, or an aqueous alcohol solution, preferably methanol in water. Depending on the MMOFsalt used, the solvent for the dispersion may be pure water.

[0114]

[0109] Desirably, the basic electron donating organic ligand species is provided to the dispersion of inorganic template in solution form, for example, as an aqueous or alcoholic solution of dissolved electron donating organic ligand species.

[0115]

[0110] In preferred examples, the inorganic template provided in 2D ultrathin form is a non-noble transition metal (MIM) inorganic material which is doped with one or more additional metals, metalloids, metal compounds or metalloid compounds, preferably those that form Lewis acid species, and especially strong Lewis acid species, within the inorganic material. B doping is preferred. The doped B may be elemental boron which is included in a material as a B defect and / or as surface B2O3 and / or B-OH functionality. In another preferred example, an additional or alternative dopant is Fe and / or P. In other embodiments the dopant may be P. In some particularly preferred embodiments, the dopants are B, Fe, Ni, P, S or combinations thereof.

[0116]

[0111] Generally, the inorganic template in 2D ultrathin / nanosheet form is not particularly limited, once it comprises one or more non-noble transition metals MIMhaving d-orbitals and / or other orbitals available for bonding through ionic, covalent or coordinate bonding. Typically, the inorganic material is crystalline, having crystal lattice with one or more repeating units. Other elements in the material may include one or more elements selected from B, O, N, S, P, Se, Si, As, Sb, etc. Suitably non-noble transition metals MIMinclude one or more metals such as Fe, W, V, Mo, Ti, Co, Ni and Cu. The material may be an oxide, sulfide, MXene or doped MXene of these metals or combinations thereof. Polymorphs of these materials that present as, or facilitate conversion to, 2D ultrathin / nanosheet form are particularly preferred.

[0117]

[0112] In some examples, the non-noble transition metal (MIM) inorganic material may be selected from non-noble transition metal oxides (TMOs) in 2D ultrathin form. Examples of TMOs include iron oxides, tungsten oxides, vanadium oxides, molybdenum oxides, cobalt oxide, nickel oxide, copper oxide manganese oxides.

[0118]

[0113] Other examples include B doped 2D non-noble transition metal oxide (TMOs) in 2D ultrathin / nanosheet form, for example, B- doped iron oxides, B- doped tungsten oxides, B- doped vanadium oxides, B- doped molybdenum oxides, B- doped cobalt oxide, B- doped nickel oxide, B- doped copper oxide, B- doped manganese oxides, such as B-Fe20s, B-a-Fe2Os, B-WO3, B-V2O5, B- M0O3, WO3.B2O3.

[0119]

[0114] Further doped examples include twin or tri doped derivatives of the non-noble transition metal (MIM) inorganic material. In some embodiments, twin doped inorganic materials are preferred as they include more active sites for OER activity, and / or provide different functionality (more active sites, stability, chloride ion shielding effect) that synergistically improve overall performance.

[0120]

[0115] The dopes may be any combination of the above-mentioned doped metal elements, or metalloid metals, metalloids, metal compounds or metalloid compounds, preferably those that form Lewis acid species, and especially strong Lewis acid species, within the inorganic material. Desirably the dopants are a combination of B and Fe; P and B.

[0121]

[0116] Desirably, the B doped and / or Fe doped non-noble transition metal oxide (TMOs), for example, B-Fe2C>3, B-a-Fe2C>3 or B-Fe-WOs, B-Fe-V2Os, B-Fe-MoOs. In some examples, a-Fe2Os and doped version thereof such as B-a-Fe2Os and B2O3-a-Fe2Os, B-B2O3-V2O5, B-B2O3-M0O3 are desirable, particularly when they are in 2D ultrathin sheet form. Desirably, the B doped and P doped non-noble transition metal oxide (TMO) is, for example, 2D ultrathin PW8O26.B2O3.

[0122]

[0117] In some examples, the non-noble transition metal (MIM) inorganic material may be selected from non-noble transition metal sulfides (TMSs) in 2D ultrathin form. Examples of TMSs include FeS2, C0S2, or NiS2 Doped examples include B doped 2D non-noble transition metal sulfide (TMS) in 2D ultrathin form for example, B-FeS2, B-C0S2, or B-NiS2.

[0123]

[0118] In some examples, the non-noble transition metal (MIM) inorganic material may be selected from doped non-noble transition metal MXenes or doped functionalised MXenes, in 2D ultrathin / nanosheet form. Examples of MXenes include TisC2Tx, where Txis O, F, OH, Cl, preferably O. An example of functionalised MXene is TisC2O2 (an O functionalised MXene). An example of a doped functionalised MXene is B-Ti3C2O2. An example of a twin doped functionalised MXene is B-Fe- TisC2O2.

[0124]

[0119] Preferably, the 2D non-noble transition metal inorganic material template is selected from Fe20s, a-Fe2Os, WO3, B doped MXene, B and Fe doped MXene or B doped functionalised MXene, B and Fe doped functionalised MXene, FeS2, B-FeS2, B-Fe20s, B-a-Fe2Os, B-WO3, B-B2O3-Fe2Os, B- B2O3-a-Fe2O3, B-B2O3-WO3, B-Fe-Fe2Os, B-Fe-a-Fe2O3, B-Fe-WOs, P-WsO26.B2O3, B-Fe20s, B-NiS2, B-FeS2, FBT, C0S2, B-CoS2-

[0125]

[0120] The doped materials are preferred, particularly those doped with boron, sulfur and / pr phosphorus, most preferably B, or B and P together. In particularly preferred examples, the 2D non- noble transition metal inorganic material is a boron doped 2D non-noble transition metal inorganic material or a twin doped 2D non-noble transition metal inorganic material with boron doping an at least one other metal, such as Fe. P and B doping is also preferred such as PW8O26.B2O3.

[0126]

[0121] Boron doping is particularly desired as it impacts several desirable attributes to the materials of the invention which include higher electrical conductivity, production of additional and / or further covalent / ionic bond among the dissimilar ultrathin sheets, and formation of borate species which are weak Lewis acid species at the interface which desirably act as a local OH- modulator, or hydrolysed borate functionality B-OH. In general terms, the B-doping modulates the electronic properties of the inorganic material and creates borate species at interface of heterostructure that regulate hydroxyl ion adsorption. The B doping also induces desirable changes in the inorganic material surface which facilitate charge transfer between adjacent carbon (C) atoms of MOF which is believed to enhance catalytic activity. Furthermore, the B-doping assists in better and strong connections of the MOF to the inorganic 2D material which induces better activity and supporting stability of the unique interface between two dissimilar materials. As a result of at least these attributes, introducing the boron, e.g., in a form of B dopant atom or boron oxide into the material has a positive effect on the OER catalysis.

[0127]

[0122] Desirably, the acidic non-noble transition metal ion of the 2D porous non-noble transition metal organic framework (MOF) comprises a transition metal, preferably a first-row transition metal, for example, selected from the group consisting of: Co, Ni, Zn, Fe, Ti, Cr and Cu and combinations thereof. These particular metals are preferred because of the presence of empty d-orbitals, which help with the absorption of OH ions. However other metals, such as Zr and Mo, could also be used. In one embodiment, two or more different acidic non-noble transition metal ions are used to construct the MOF. In one embodiment, two different acidic non-noble transition metal ions are used to construct a bimetallic MOF.

[0128]

[0123] Desirably, the acidic non-noble transition metal ions of the 2D porous non-noble transition metal organic framework (MOF) are derived from a salt of the metal, for example, nitrate salts, chloride salts, sulfate salts, etc. Desirably, cobalt nitrate (e.g., Co(N03)2.6H20) or nickel nitrate (Ni(NO3)3-6H2O) are preferred, where a CoMOF or NiMOF is required. The nitrate salt is particularly preferred.

[0129] Materials per se

[0130]

[0124] The invention extends to a vertical layered 2D / 2D A-[B-A]ntype metal inorganic-X-metal organic heterostructure having strongly bonded heterointerfaces obtained by or obtainable by the process of the invention. The invention provides new and useful electrocatalytic materials perse, with particularly good utility as OER electrocatalysts, especially for water spitting in alkaline seawater. The materials are in the form of a vertical layered 2D-2D A-[B-A]ntype metal inorganic-X-metal organic heterostructure having strongly bonded heterointerfaces, the heterostructure comprising: a 2D heterostructure (HS) layered assembly with vertically stacked ultrathin layers of dissimilar materials with strong chemical bonds, that include ionic bonds, between the dissimilar materials, and wherein dissimilar materials include: (a) a 2D non-noble transition metal inorganic material, and (b) a 2D porous non-noble transition metal organic framework (MOF).

[0131]

[0125] As will be evident from the above description, the heterostructure assembly is composed of two heterogenous (dissimilar) 2D (ultrathin) precursor materials, a 2D non-noble transition metal (MIM) inorganic material and a 2D porous non-noble transition metal (MMOF) organic framework (MOF). Desirably, the assembly involves direct stacking of alternating individual monolayers of each different materials, for example, in a A-[B-A]nvertically stacked orientation. Anchoring strong chemical bonds, e.g., ionic bonds and particularly MIM-O-MMOFionic bonds, are located at heterointerfaces between the stacked layers of dissimilar materials, where strong chemical bonds include ionic bonds, for example, ionically bonded MIM-oxygen-MMOFbridges. In some cases, direct bonds between the N and / or O atoms of the organic linker may be present.

[0126] In some examples, covalent bond and / or coordinate bonds may be present instead of covalent bonds or in addition to covalent bonds. The strong bonds are located at heterointerfaces, for example, at gaps and / or junctions between adjacent layers of the dissimilar materials.

[0132]

[0127] As explained above, the 2D porous non-noble transition MOF includes a metal cluster of a tetrahedrally coordinated acidic non-noble transition metal ion (e.g., M = Co, Zn, Ni, Fe) coordinated to the one or more basic organic coordinating ligands.

[0133]

[0128] Preferably, the 2D-2D HS does not have a foam like morphology. Preferably, the MOF is not associated with the 2D ultrathin inorganic material in a perpendicular fashion (MOF perpendicular to plane of the inorganic sheet), as is the case for the material disclosed in Liu which is discussed herein.

[0134]

[0129] Suitably, complex covalent linkers, such as bridging molecules have two or more different atoms, between the dissimilar materials are not required. As a result, the layers of the HS materials of the invention have a smaller interlayer distance than heterostructure materials involving such known linker facilitated bond between dissimilar materials.

[0135]

[0130] Importantly, the materials are a true heterostructure whereby components remain as distinct / dissimilar materials evidenced by retention of typical individual component crystallographic features remaining discernible in the HS material. All elements of both materials are uniformly dispersed throughout the HS material indicating excellent MOF coverage over the surfaces of the inorganic template. Desirably, the basal planes of the nanosheets are strongly linked and grown on each other. The elements I atoms of both components are uniformly disturbed through the heterostructure which indicates successful coverage of the inorganic template by the MMOF-organic component and confirming that the sheets grow on top of each other. The lattice spacing of the HS indicate sheets of the components are vertically stacked in an alternating layered arrangement. In some examples, the MIMmetal of the inorganic template and / or the MMOFof the metal-organic component exists in more than one oxidation state as a result of formation of the ionic bond. Typical ionic bonds formed include X (-O-, -S-, -Se-, -C-, -N-, -P-, or -B-) bridging between MIMand MMOFmetals. Other bonds include new MIM-N (ligand) interactions created between the dissimilar 2D sheets. In typical examples, the HS material includes tetrahedrally coordinated MOF metal ions as a result of ionic bond formation between the materials. Octahedral coordination may also be present as an indicator of ionic bond formation. In many materials, point defect oxygen atoms are present. The organic component of, e.g., a MOF, also makes the WO3.B2O3 (WB) in the HS highly hydrophilic, increasing its wettability. High wettability and low charge transfer resistance results in better kinetics by the quick adsorption of ions and their conversion to molecular species. Heterostructuring constituent components can enhance their individual stabilities in water splitting applications.

[0136]

[0131] MOF also makes the inorganic component of the HS highly hydrophilic, increasing its wettability. Thus, high wettability and low charge transfer resistance results in better kinetics by the quick adsorption of ions and their conversion to molecular species.

[0137]

[0132] Desirably, the one or more basic organic coordinating ligands of the 2D porous non-noble transition metal organic framework (MOF) are selected from the group consisting of: dicarboxylic acids, tricarboxylic acids, tetracarboxylic acids, imidazolate, and combinations thereof. Desirably, the 2D ultrathin / nanosheet non-noble transition metal inorganic material is selected from the options described above in terms of the process of the invention.

[0138]

[0133] Preferred 2D ultrathin / nanosheet non-noble transition metal inorganic materials are described above. Preferably selected from: a transition metal oxide (TMO), such as Fe2C>3 (see Example 1 involving Fe20s), WO3, V2O5, M0O3, and a B doped transition metal oxide (TMO), such as a B-Fe20s, B-WOs (see Example 2 involving B- WO3), B-V2O5, B-M0O3; a B doped transition metal sulfide (TMS), such B-FeS2 (see Example 4 involving B-FeS2), B-C0S2, or B-NiS2, and a B and P doped transition metal oxide (TMO), such as B and P doped WO3 (see Example 5 involving PW8O26-B2O3) a B doped MXene, B and Fe doped MXene or B doped functionalised MXene, B and Fe doped functionalised MXene wherein the MXene is a transition metal carbide (TMC) (see Example 4 involving FBT); a transition metal nitride (TNN) or a transition metal carbon nitride (TMN).

[0139]

[0134] In one preferred example, the 2D non-noble transition metal inorganic material is a Fe and B doped MXene or functionalised MXene, and the 2D porous non-noble transition metal organic framework (MOF) is a CoMOF. Desirably, the MXene or functionalised MXene is a transition metal carbide (TMC); a transition metal nitride (TNN) or a transition metal carbon nitride (TMN), such as TisC2Tx, particularly twin doped Fe- and B- doped Ti3C2Tx. An example is twin doped TisC2O2, particularly B and Fe doped TisC2O2, designated herein as ‘FBT’. For MXenes, preferred non-metal dopants are one or more of B, S, N, and P. In terms of electrochemically active metal doping, the metal can be one or more of Fe, Co and Ni, particularly Fe. Preferred MXenes are described above.

[0140]

[0135] Desirably, the 2D ultrathin I nanosheet noble transition metal inorganic material template is selected from Fe20s, WO3, MXene or functionalised MXene such as FBT, FeS2, B-FeS2, B-Fe20s, B- WO3, B-FeS2, B-NiS2, V2O5, M0O3, P-WsO26-B2O3, C0S2, B-C0S2

[0141]

[0136] Desirably, the MOF is a CoMOF, a NiMOF, a FeMOF, a ZrMOF, a TiMOF, a CrMOF, a CuMOF, a ZnMOF, or combinations thereof. Any of one of the inorganic material templates can be combined with and one of the MOFs listed. However, the following combinations are particularly preferred. Desirably, the 2D non-noble transition metal inorganic material is a transition metal oxide (TMO, e.g., Fe20s, WO3), and the 2D porous non-noble transition metal organic framework (MOF) is a CoMOF. Desirably, the 2D non-noble transition metal inorganic material is a B doped transition metal oxide (B-TMO, e.g., WO3.B2O3), and the 2D porous non-noble transition metal organic framework (MOF) is CoMOF or FeMOF. Desirably, the 2D non-noble transition metal inorganic material is a B doped transition metal sulfide (B-TMS, such as B-FeS2, B-C0S2 or B-NiS2), and the 2D porous non-noble transition metal organic framework (MOF) is a NiMOF. Desirably, the 2D non-noble transition metal inorganic material is a P and B doped transition metal oxide (P-B-TMO, e.g., PW8O26-B2O3), and the 2D porous non-noble transition metal organic framework (MOF) is a FeMOF. Desirably, the at least one basic electron donating organic ligand species of the 2D porous non-noble transition metal organic framework (MOF) comprises an alkylimidazole species, such as a C1-C5 alkylimidazole, for example 2-methylimidazole. In some examples, the ultrathin / nanosheet 2D porous non-noble transition metal organic framework (MOF) is ZIF-67. ZIF stands for zeolitic imidazolate framework-67 which is an assembly of 2-methylimidazole cobalt salts.

[0142]

[0137] In another preferred embodiment, the invention relates to a vertical layered 2D-2D metal inorganic-X-metal organic heterostructure having strongly bonded heterointerfaces, the heterostructure comprising:

[0143] 2D heterostructure (HS) layered assembly with vertically stacked ultrathin layers of dissimilar materials with strong chemical bonds, that include ionic bonds, between the dissimilar materials at heterointerfaces at gaps and / or junctions between the dissimilar materials, and wherein dissimilar materials include: (a) a 2D non-noble transition metal inorganic material which a transition metal oxide (TMO), a B doped transition metal oxide (B-TMO), or a P doped and B doped transition metal oxide (P-B-TMO) and (b) a 2D porous non-noble transition metal organic framework (MOF) including a metal cluster of a tetrahedrally coordinated non-noble transition metal ion (e.g., M = Co, Zn, Ni, Fe) coordinated to one or more electron donating organic ligands, such as dicarboxylic acids, tricarboxylic acids, tetracarboxylic acids, and imidazolates.

[0144]

[0138] Desirably, the B doped 2D ultrathin / nanosheet non-noble transition metal inorganic material includes B2O3 and or B-OH functional groups. Desirably, the MOF is a CoMOF, a NiMOF, a FeMOF, a ZrMOF, a TiMOF, a CrMOF, a CuMOF, a ZnMOF, or combinations thereof. In one embodiment, the MOF comprises a combination of two or more different metals each independently selected from the group Co, Ni, Fe, Zr, Ti, Cr, Cu and Zn. In one embodiment, the MOF is a bimetallic MOF comprising a combination of two different metals each independently selected from the group Co, Ni, Fe, Zr, Ti, Cr, Cu and Zn. In one embodiment, the MOF is a bimetallic MOF comprising a combination of two different metals each independently selected from the group Co, Ni, Fe, and Zn. In one embodiment, the MOF is a bimetallic MOF, such as an FeCo-MOF (“FeCo-OF”), NiCo-MOF, ZnFe-MOF, ZnNi- MOF, NiFe-MOF, or ZnCo-MOF. Preferably, the MOF grown thereon is a CoMOF or a FeMOF, preferably comprising imidazolate ligands. In one embodiment, the MOF grown thereon is a CoFe- MOF.

[0145]

[0139] Desirably, the 2D ultrathin I nanosheet non-noble transition metal inorganic material is Fe20s, and the 2D porous non-noble transition metal organic framework (MOF) grown thereon is a CoMOF, preferably comprising imidazolate ligands. In one example, the HS material is Fe203 / CoMOF, without any dopant atoms / elements. In another example, the HS material is B-WO3 / C0MOF. In another example, the HS material is P-B-doped WOs / FeMOF, particularly, P-WsC^e-E^Os / FeMOF.

[0146]

[0140] For the Fe203 / CoMOF material, the CoMOF sheets cover the 2D Fe20s via strong ionic bonds while keeping the morphological features. The heterointerface of the material is anchored with ionic bonds between metals and oxygen (Co-O-Fe) of dissimilar sheets to get stable structure and electrocatalytic properties. The new ionic connections between dissimilar sheets ensure enhanced stability of the developed heterostructure. Desirably, in addition to (MOF)Co-O-Fe(surface) bridging bonds between the dissimilar materials, one or more of the following additional bonds may be present as indicated by one or more of FTIR, XPS, NEXAFS, DFT calculations: (surface)Fe-O [ionic bond], (surface)Fe-N, Co-0 [ionic bond], Co-N(imidazole), direct Co-Fe bond. The development of heterointerface changes the oxidation state of the metal nodes of the CoMOF to create strong ionic interactions between the sheets. New Co-0 bonds are formed in the heterostructure due to bridging between Fe20s and CoMOF. During electrocatalytic OER in real seawater the heterostructured catalyst achieved 0.5 and 1 A cm-2current density at overpotential of 310 and 410 mV, respectively. Catalyst stability of more than 900 hours was demonstrated. The Co MOF component acts as a protective shield that repels chlorine and prevents corrosion of the catalyst supporting a demonstrated operating time in seawater. In particular, it is believed that OH- adsorption is favoured at heterointerface due to tuned electronic and surface properties as well as a result of the ionic bond between the nanosheets of the HSs. The ionic bond imparts stability via creation of a strongly bonded 2D heterostructure, but also contribute to fast charge transfer.

[0147]

[0141] Suitably, the C0-MOF / WO3.B2O3 (MOF / WB) desirably comprises monoclinic WO3. The material comprises borate species such as B2O3. The material comprises B dopant in the WO3 sheets which is relaxed away from the surface layers and creating a guided surface for the CoMOF growth. The material comprises a uniform distribution of all elements across the sheets, confirming the presence of both the MOF and WB in the heterostructure. The material comprises Co-O-W ionic bonds. W6+and W4+oxidation states are both present in the HS. The material includes tetrahedrally coordinated Co.

[0148]

[0142] In preferred embodiment, the invention relates to a vertical layered 2D-2D metal inorganic-X- metal organic heterostructure having strongly bonded heterointerfaces, the heterostructure comprising:

[0149] 2D heterostructure (HS) layered assembly with vertically stacked ultrathin layers of dissimilar materials with strong chemical bonds, that include ionic bonds and covalent bonds, between the dissimilar materials at heterointerfaces at gaps and / or junctions between the dissimilar materials, and wherein dissimilar materials include: (a) a 2D non-noble transition metal inorganic material which is a doped MXene, a doped functionalised MXene, a B doped MXene, a B doped functionalised MXene, a twin Fe and B doped MXene, or twin Fe and B doped functionalised MXene, and (b) a 2D porous non-noble transition metal organic framework (MOF) including a metal cluster of a tetrahedrally coordinated non-noble transition metal ion (e.g., M = Co, Zn, Ni, Fe) coordinated to one or more electron donating organic ligands, such as dicarboxylic acids, tricarboxylic acids, tetracarboxylic acids, and imidazolates, such as alkylimidazolates. The MXene may be doped with one or more of B, S, N and P. The MXene may also be doped with one or more of electrochemically active Fe, Co and Ni.

[0150]

[0143] Desirably, the 2D non-noble transition metal inorganic material template comprises metals or is doped with metals that form Lewis acid species, such as boron, aluminium, sulfur, and / or phosphorus. In preferred examples, the 2D non-noble transition metal inorganic material template is doped with at least boron or phosphorous or combinations thereof.

[0151]

[0144] Desirably, the 2D non-noble transition metal inorganic material is a Fe and B doped MXene or functionalised MXene, and the 2D porous non-noble transition metal organic framework (MOF) grown thereon is a CoMOF, preferably comprising imidazolate ligands. One preferred example of a suitable MXene is ZIF-671 FBT (FBT@MXene) as described elsewhere herein which comprises a twin doped B and Fe doped Ti3C2O2.

[0152]

[0145] Desirably, wherein the 2D non-noble transition metal inorganic material is a Fe and B doped functionalised MXene, TisC2Tx wherein X is O and the 2D porous non-noble transition metal organic framework (MOF) is ZIF-67. Desirably the functionalised MXene is TisC2O3 and the MOF is a CoMOF, e.g., ZIF-67 (CoMOF). NiMOF or Fe MOF can also be used with the MXene examples, particularly FBT.

[0153]

[0146] In one general example, the invention relates to a vertical layered 2D / 2D metal inorganic-O- metal organic heterostructure having strongly bonded heterointerfaces, the heterostructure comprising:

[0154] 2D heterostructure (HS) layered assembly with vertically stacked ultrathin layers of dissimilar materials with strong chemical bonds, that include ionic bonds and covalent bonds, between the dissimilar materials at heterointerfaces at gaps and / or junctions between the dissimilar materials, and wherein dissimilar materials include: (a) a 2D non-noble transition metal inorganic material which is a B doped functionalised MXene, preferably TigC2O3, or a twin Fe and B doped functionalised MXene, preferably TisC2O3; and (b) a 2D porous non-noble transition metal organic framework (MOF) which is a CoMOF, preferably ZIF-67.

[0155]

[0147] The twin doping strategy of a precursor MXene (TisC2Tx, where Tx is surface termination and can be Fe, B or both) produces new material “FBT” preferably provided as a multiphase thin 2D material. In some examples, the FBT precursor may include a mixture two or more phases of (a) hexagonal shaped titanium carbide (TigCs.ys) and (b) iron titanium oxide (Fe2TigOg; pseudorutile), and (c) tetragonal shaped titanium boron oxide (TiBo.02402; B-doped rutile). Preferably, the inorganic material is in thin 2D sheets. Suitably, clusters of dense particles with different shapes (rectangular & hexagonal) may be present. In some embodiments, the material comprises a mixture of hexagonal shaped titanium carbide (TieCs.ys) and iron titanium oxide (Fe2TigOg), and tetragonal shaped titanium boron oxide (TiBo.02402).

[0156]

[0148] A MOF, such as a NiMOF, FeMOF, or CoMOF, particularly ZIF-67 MOF, is grown on the FBT sheets to form a 2D / 2D HS heterostructure with strong ionic bonds at heterointerfaces. In preferred examples involves CoMOF, such as ZIF-67 MOF being grown on the FBT sheets to form a MOF@FBT HS according to the invention (ZIF-67 / FBT, MOF@FBT). The preferred MOF@FBT material comprises Co3+species as a result of Co-O-Fe ionic interactions between the dissimilar materials which are formed during heterostructure formation, which are particularly aided by the vacancies generated when B replaces some of the Ti in Ti3C2O2-MXene in the precursor material synthesis. Desirably, the material comprises Fe-O-Co bonds. In the preferred MOF@FBT material, C-Ti-Tx, Ti(ll)-Tx (where Tx is surface termination and can be Fe, B or both) is present, e.g., as confirmed by XPS analysis identifying presence of Ti in different oxidation states including 0, +2, and +4. Suitably, the material may comprise one or more of: TiC>2 (metal-O), Ti-O-Tx, C-Ti-(OH)Xand C-0 groups. Suitably, the material may comprise one or more of: Ti, O, C-C, C-N, C-O, and O-C=O groups. Desirably, the material comprises Ti-B-0 bonds. Desirably, the material comprises one or more of Fe2+and Fe3+species. Desirably, the material comprises Co, and C-N bonds. Desirably, the material comprises Co-N(imidazole) bonds. The material may also comprise TiC>2 (lattice oxygen), (Ti or B)-O- Fe (defective oxygen vacancy sites due to low oxygen coordination), C-Ti-(OH)Xand C-0 respectively were identified. The Fe active sites are thought to be mainly responsible for the improved OER activities, while B doping generate the much-needed unoccupied oxygen site vacancies (by substituting some of the Ti) that aid MOF growth without destroying the 2D morphology, thus exposing more active sites for OER activities in the inorganic components. Suitably, the Fe dopant is confined to incorporation of the some of the surface termination of the MXene. The ionic bonds and particular Co-0 bond of Co-O-Fe, exposes more metal sites for OER activity. The material includes Fe in the +2 and +3 oxidation states.

[0157]

[0149] In terms of OER performance, ZIF-67 / FBT mitigates chloride oxidation reactions due to its ability to generate high current densities at overpotential <480 mV in alkaline 6 M KOH DI and 6 M KOH seawater medium, demonstrating exclusive selectivity towards OER while avoiding chlorine oxidation byproduct formation. This exclusive OER selectivity is attributed to the material’s impressively high OER activities in seawater, hence making them an excellent catalyst for industrial seawater splitting.

[0158]

[0150] MXene materials are conventionally considered to have limited activity and stability as an anode in seawater and exhibit poor OER performance. It is believed that the electrochemical neutrality of MXenes, caused by poor redox potential and propensity to undergo surface oxidation especially during the OER process, has traditionally limited their application in OER application in a complex and corrosive system like seawater. As such there is need for a doping strategy that improves MXene performance for OER, e.g., by promoting surface termination / breaking electrochemical neutrality while preserving its 2D morphology while improving the material flexibility without negatively impacting the electrochemical activity and in a way in which prevents the surface termination from leaching.

[0159]

[0151] The twin doped functionalised MXene material of the invention address this problem by exhibiting vastly improved OER performance in seawater splitting. The doping strategy described herein break the otherwise OER inactivity of MXene to drive ampere-level selective OER in seawater. The exemplary material ZIF-67 I FBT (CoMOF I FBT) of the invention demonstrates an exceptional OER performance in both fresh and seawater splitting, and in some examples, achieves OER at a current density of 500, 1000, and 1500 mA cm-2at overpotentials of 338, 430 and 520 mV respectively in 6 M deionised water (DI) water and 387, 508, and 613 mV in 6 M KOH seawater, as determined by LSV. A Tafel slope <66.89 mV dec1shows improved kinetics of the materials. This activity is thought to be due to abundant accessible electroactive sites on the FBT inorganic material that are further enhanced by presence of the 2D morphology. Further, the unique ionically bonded heterointerface connection between the ZIF-67 and FBT components further enhanced performance and stability. Incorporation of Fe and B into the MXene and / or the unique heterointerface is thought to significantly boost electrochemical performance by enabling and enhancing rapid charge redistribution between the dissimilar materials which enhances OER activity.

[0152] In another preferred embodiment, the invention relates to a vertical layered 2D-2D metal inorganic-O-metal organic heterostructure having strongly bonded heterointerfaces, the heterostructure comprising:

[0160] A 2D heterostructure (HS) layered assembly with vertically stacked ultrathin layers of dissimilar materials with strong chemical bonds, that include ionic bonds and covalent bonds, between the dissimilar materials at heterointerfaces at gaps and / or junctions between the dissimilar materials, and wherein dissimilar materials include: (a) a 2D non-noble transition metal inorganic material which is a transition metal sulfide (TMS) or B doped transition metal sulfide (TMS), and (b) a 2D porous non- noble transition metal organic framework (MOF) including a metal cluster of a tetrahedrally coordinated non-noble transition metal ion (e.g., M = Co, Zn, Ni, Fe) coordinated to one or more electron donating organic ligands, such as dicarboxylic acids, tricarboxylic acids, tetracarboxylic acids, and imidazolates.

[0161]

[0153] Desirably, the 2D non-noble transition metal inorganic material is a 2D non-noble transition metal inorganic material which is B doped FeS2, B doped C0S2 or B doped NiS2, and the 2D porous non-noble transition metal organic framework (MOF) grown thereon is a NiMOF or a FeMOF, preferably comprising one or more imidazolate organic ligands, such as alkyl imidazolate organic ligand including 2-methylimidazole.

[0162]

[0154] In a related aspect, the invention provides a vertical 2D heterostructure (2D HS) layered inorganic-organic assembly comprising:

[0163] 2D heterostructure (HS) layered assembly with vertically stacked ultrathin layers of dissimilar materials with strong chemical bonds, that include ionic bonds and covalent bonds, between the dissimilar materials at heterointerfaces at gaps and / or junctions between the dissimilar materials, and wherein dissimilar materials are: (a) a 2D non-noble transition metal inorganic material which is B doped FeS2’ B doped C0S2 or B doped NiS2, and (b) 2D porous non-noble transition metal organic framework (MOF) which is NiMOF.

[0164]

[0155] The invention provides new heterostructured boron-doped iron disulfide (B-FeS2), B doped C0S2 (B-C0S2) or boron-doped nickel disulfide (B-NiS2) HSs with transition metal MOF sheets.

[0165]

[0156] An exemplary material B-doped FeS2 / Ni MOF having ionic and covalent bonds between the dissimilar materials, demonstrates an exceptional OER performance in both 6 M KOH in DI and seawater as electrolytes, achieving a current density of 1 A cm-2for OER at an overpotential of just 1 .67 V in seawater. This corresponds to an 85% improvement compared to lrO2. Further, the system has been demonstrated to be stable over 230 operating hours. Furthermore, the material potentially realises H2 directly from seawater.

[0166]

[0157] The composite material exhibits improved surface reactions that improve catalyst behaviours in water splitting. The metal salt complexes of the MOF cover the surfaces of the catalytically active areas of the material and function as a protective shield. B doping introduces changes in surface area thereby facilitating charge transfer between adjacent carbons of the MOF which enhances catalytic activity. The presence of borate anion in the material assists in OER catalyst performance by enables deeper comprehension of in the catalytic activity in the HS. Pristine FeS2 has limited OER performance due to electronegative of the S atom, that hampers interaction between OH- ions and metal site, leading to repulsion between S and O orbitals which restrict binding of OH groups and subsequent OER kinetics.

[0167]

[0158] Desirably, the B-FeS2 material comprises micron scale ultrathin sheets. Desirably, the B-FeS2 material is highly porous as a result of the thermal conversion step. It is believed that the catalyst performance benefits from the low thickness and layered structure of the B-FeS2 nanosheet, along with interconnected metal ions and ligands in the MOF.

[0168] Applications

[0169]

[0159] The heterostructured materials of the invention have utility in many applications, including hydrogen electrochemistry, water electrolysis, fuel cells, and their H related reactions. In a preferred application, the invention extends to a use of a vertical 2D heterostructure (2D HS) layered inorganic- organic assembly of the invention as an OER catalyst for water splitting, preferably alkaline water, alkaline seawater, deionised water, tap water or wastewater.

[0170] Preferred Embodiments of the Invention

[0171]

[0160] Various 2D / 2D ionically bonded heterostructures (HSs) were prepared using the interface- assisted synthesis technique of the invention at room temperature and pressure. First, a 2D ultrathin non-noble transition metal inorganic material was provided as an inorganic template for directed parallel MOF growth thereon, and then, a sheet of a 2D non-noble transition metal MOF was constructed in over surfaces / in parallel to sheets of the inorganic template to provide vertically stacked or arranged heterostructures of dissimilar materials which were stabilised by ionic bonds at heterointerfaces of the dissimilar materials. The resulting new 2D / 2D heterostructure was then characterised and electrocatalytic OER activity tested. Preferred heterointerfaces were anchored together via the ionic bonds between bridged metal-X-metal atoms of the dissimilar sheets, preferably where X is -O-, -S-, -Se-, -C-, -N-, -P-, or -B-, etc., (e.g., MIM-O-MMOF) which stabilise the structure in a vertically stacked arrangement. Suitably, the inorganic solid-state material comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W. The inorganic material may further comprise one or more of: metal-0 (oxide), metal-N (nitride), metal-C (carbide), metal-C / metal N (carbonitride); metal-S (sulfide) and / or M-P (phosphide) functionalities. While not necessarily essential but preferable, to enhance ionic connections between the components, wherein the inorganic solid-state material is doped with one or more non-metal dopants for chloride shielding polyanion formation, wherein preferably the dopants are selected from the group consisting of: B, P, C, N, O, Si, S, As and Se. The ionic bond and optional additional surface modification introduced into the inorganic template through doping prior to HS formation enhance electrocatalytic activity of the resulting 2D / 2D HS material. Electrocatalytic OER in real seawater can be achieved with the materials at excellent industrial level current densities at low overpotentials, as well as at higher overpotentials that otherwise would not avoid chloride oxidization reactions. The strong ionic connections, quality and number of activity sites, and surface functionality of preferred HSs that repel Cl- but is attractive to OH- support catalyst activity and stability at the excellent current densities for up to and in some cases more than 1000 hours, far surpassing commercial noble catalyst performance. The synthetic method enables easy and low cost scale up to produce quantities suitable for commercial application.

[0172]

[0161] Unless the context indicates otherwise, where the terms “comprise”, “comprises” and “comprising” are used in the specification (including the claims) they are to be interpreted as specifying the stated features, integers, steps or components, but not precluding the presence of one or more other features, integers, steps or components, or group thereof.

[0173] Embodiments

[0174]

[0162] Embodiment 1. A vertical layered A-[B-A]ntype heterostructure assembly of a 2D ultrathin inorganic solid-state material (‘B’) and a non-noble metal organic framework (MOF) material (‘A’) that covers surfaces of each layer of the inorganic solid-state material (‘B’) in the assembly in a parallel arrangement to the 2D ultrathin inorganic solid-state material (‘B’), wherein n is a positive integer, wherein heterointerfaces of the assembly comprise bridging MIM-X-MMOFionic bonds that stabilise heterointerfaces between the inorganic solid-state material (‘A’) and the MOF material, where bridging X is -O-, -S-, -Se-, -C-, -N-, -P-, or -B-, MIMis a metal in the inorganic solid-state material, and MMOFis a metal in the non-noble MOF; wherein the inorganic solid-state material comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; and further comprises one or more of: metal-0 (oxide), metal-N (nitride), metal-C (carbide), metal- C / metal N (carbonitride); metal-S (sulfide) and / or M-P (phosphide) functionalities; and optionally wherein the inorganic solid-state material is doped with one or more non-metal dopants, preferably selected from the group consisting of: B, P, C, N, O, Si, S, As and Se.

[0175]

[0163] Embodiment 2. A process of forming an electrocatalyst comprising the steps of: providing a liquid dispersion of an inorganic solid state material (‘B’) in 2D ultrathin sheet form as an inorganic template for non-noble metal organic framework (MOF) material (‘A’) growth thereon in a parallel arrangement to the 2D ultrathin inorganic solid-state material (‘B’), wherein the inorganic template comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; and further comprises one or more of: at least metal-0 (oxide), metal-N (nitride), metal-C (carbide), metal-C / metal N (carbonitride); metal-S (sulfide) and / or M-P (phosphide) functionalities, and optionally is doped with one or more non-metal dopants selected from the group of non-metals consisting of: B, C, N, O, P, Si, S, As and Se; electrostatically associating at least one basic organic coordination ligand species suitable for MOF fabrication with the inorganic template; providing a liquid solution of one or more acidic non-noble transition metal (MMOF) salts to the dispersion of inorganic material template having electrostatically associated ligand species; forming a vertical layered A-[B-A]ntype heterostructure assembly of the 2D ultrathin inorganic material (‘B’) with MOF material (‘A’) by allowing complexation of acidic MMOFions of the salt with the at least one basic organic coordination ligand species associated with the inorganic template to form the MOF in coverage over surfaces of the inorganic template, whereby during complexation, MMOFforms bridging MIM-X-MMOFionic bonds with metals in the inorganic material, MIM, at heterointerfaces between the inorganic template and MOF, wherein bridging X is -O-, -S-, -Se-, -C-, -N-, -P-, or -B-, and wherein n is a positive integer.

[0176]

[0164] Embodiment 3. The heterostructure assembly of Embodiment 1 or the process of Embodiment 2, wherein non-metal dopants form chloride ion shielding polyanions selected from one or more of: borate, sulfate, phosphate, chromate, permanganate, carbonate, and nitrate.

[0177]

[0165] Embodiment 4. The heterostructure assembly of Embodiment 1 or the process of Embodiment 2 or Embodiment 3, wherein the inorganic material template comprises, or is doped with, one or more electrochemically active transition metal atoms that form the bridged X ionic bonds (MIM- X-MMOF) with the acidic non-noble transition metal of the MOF.

[0178]

[0166] Embodiment 5. The heterostructure assembly or process of any one of the preceding Embodiments, wherein electrochemically active transition metal atoms are Fe, Ni, Co, Cu, Cr, Mo and W, preferably Fe.

[0179]

[0167] Embodiment 6. The heterostructure assembly or process of any one of the preceding Embodiments, wherein the inorganic solid-state material is a transition metal oxide (TMO); a transition metal nitride (TMN); a transition metal sulfide (TMS); a transition metal carbide (TMC), preferably a MXene doped with at least one electrochemically active metal; transition metal phosphide (TMP); or mixed variants thereof.

[0180]

[0168] Embodiment 7. The heterostructure assembly or process of any one of the preceding Embodiments, wherein the inorganic solid-state material is a transition metal oxide (TMO); a transition metal nitride (TMN); a transition metal sulfide (TMS); a transition metal carbide (TMC), preferably a MXene doped with at least one electrochemically active metal; transition metal phosphide (TMP); or mixed variants thereof, doped with one or more of B, C, N, O, P, Si, S, As and Se; and optionally further doped with at least one electrochemically active metal, such as Fe, Ni, Co, Cu, Cr, Mo and W.

[0181]

[0169] Embodiment 8. The heterostructure assembly or process of any one of the preceding Embodiments, wherein the non-metal dopant is one or more of B, C, S and P, and the polyanion is one or more of borate, sulfate, phosphate, chromate, permanganate, carbonate, and nitrate, preferably borate, carbonate, sulfate and / or phosphate functionalities on surfaces of the inorganic solid-state material that repel Ch from the electrochemically active metals of the inorganic solid-state material.

[0182]

[0170] Embodiment 9. The heterostructure assembly or process of any one of the preceding Embodiments, wherein the inorganic solid-state material is a transition metal oxide (TMO), a transition metal nitride (TMN), a transition metal sulfide (TMS), a transition metal carbide (TMC), a transition metal carbonitride (TMCN); transition metal phosphide (TMP), or a mixed variant thereof, doped with B, or with B and P.

[0183]

[0171] Embodiment 10. The heterostructure assembly or process of any one of the preceding Embodiments, wherein MMOFis selected from the group consisting of: Co, Ni, Fe, Cu, Zn, Zr, and Ti, and MIMis selected from the group consisting of: Fe, Ni Co, W, Ti, Cu, V and Mo.

[0184]

[0172] Embodiment 11. The heterostructure assembly or process of any one of the preceding Embodiments, wherein the non-metal dopant is B and P, forming is borate and phosphate polyanion functionalities that repel Cl- from the electrochemically active metals of the inorganic solid-state material.

[0185]

[0173] Embodiment 12. The heterostructure assembly or process of any one of the preceding Embodiments, wherein the non-metal dopant is B and P, and the polyanion is borate, and phosphate functionalities that repel Ch from the electrochemically active metals of the inorganic solid-state material, and wherein the inorganic solid-state material comprises one or more of Fe, Co, Ti, V, Mo, W, and / or Ni.

[0186]

[0174] Embodiment 13. The heterostructure assembly or process of any one of the preceding Embodiments, wherein the inorganic template in 2D ultrathin form is selected from: non-noble transition metal oxide (TMO), iron oxide, such as Fe20s, tungsten oxide such as WO3, vanadium oxide, such as V2O5, molybdenum oxide such as M0O3; non-noble transition metal MXene or functionalised MXene, preferably B-TisC2Tx, where Tx is O; non-noble transition metal sulfide (TMS), such as iron sulfide such as FeSa, cobalt sulfide such as C0S2, nickel sulfide such as NiS2;

[0187] B doped non-noble transition metal MXene or B-doped functionalised MXene, preferably B- TisC2Tx or B-Fe-TisC2Tx, preferably where Tx is O, F, OH or Cl, preferably O;

[0188] B doped 2D non-noble transition metal sulfide (TMS), for example, B-doped iron sulfide such as B-FeS2, or B-doped cobalt sulfide such as B-C0S2, or B-doped nickel sulfide such as B-NiS2, B doped 2D non-noble transition metal oxide (TMO), for example, B-doped iron oxide such as B- Fe20s, B-doped tungsten oxide such as B-WO3, B-doped vanadium oxide such as B-V2O5, B- doped molybdenum oxide such as B-M0O3;

[0189] B doped and / or Fe doped non-noble transition metal oxide (TMO), for example, B- and Fe-doped iron oxide such as B-doped iron oxide such as B-Fe20s or B-a-Fe2Os or B- and Fe- doped tungsten oxide such as B-Fe-WOs;

[0190] B doped and Fe doped non-noble transition metal sulfide (TMS), for example, B- and Fe- doped iron sulfide such as B-Fe-FeS2, B- and Fe- doped nickel sulfide such as B-Fe-NiS2, B- and Fe- doped tungsten sulfide such as B-Fe-WS2, B- and Fe- doped vanadium sulfide such as B-Fe- V2S5, or B- and Fe doped molybdenum sulfide such as B-Fe-MoS2, B- and Fe- doped cobalt sulfide such as B-Fe-CoS2.

[0191]

[0175] Embodiment 14. The heterostructure assembly or process of any one of the preceding Embodiments, wherein the inorganic template in 2D ultrathin form is selected from: Fe20s; WO3.B2O3; B-Fe2Os; B-Fe-Ti3C2Tx, where X is O; B-FeS2; B-C0S2; B-NiS2; and P-WsCke-E^Os; MoN, M0S2, B- C0S2.

[0192]

[0176] Embodiment 15. The heterostructure assembly or process of any one of the preceding Embodiments, wherein the at least one basic organic ligand of the 2D porous non-noble transition metal organic framework (MOF) is selected from the group consisting of: dicarboxylic acids, tricarboxylic acids, tetracarboxylic acids, imidazolates, bisimidazolates, triazolates, pyridinates and combinations thereof.

[0177] Embodiment 16. The heterostructure assembly or process of any one of the preceding Embodiments, wherein the at least one basic organic ligand species of the 2D porous non-noble transition metal organic framework (MOF) is an alkyl imidazole, preferably 2-methylimidazole.

[0193]

[0178] Embodiment 17. The heterostructure assembly or process of any one of the preceding Embodiments, wherein the acidic non-noble transition metal ion of MOF is selected from the group consisting of: Co, Ni, Fe, Zn, Cu, Ti, Zr, and combinations thereof.

[0194]

[0179] Embodiment 17a. The heterostructure assembly or process of any one of the preceding Embodiments, wherein the bridging MIM-X-MMOFionic bonds are selected from Fe-O-Co; W-O-Co; W- O-Fe.

[0195]

[0180] Embodiment 18. The heterostructure assembly or process of any one of the preceding Embodiments, wherein the heterointerface further comprises one or more additional bonds selected from metal-N bonds, B-O, P-0 bonds.

[0196]

[0181] Embodiment 19. The heterostructure assembly or process of any one of the preceding Embodiments, wherein the heterostructure assembly is selected from Fe20s I CoMOF; WO3.B2O3 1 CoMOF; B-Fe2O3 / NiMOF; B-Fe-Ti3C2Tx / CoMOF; B-FeS2 / NiMOF; B-C0S2 / NiMOF B-NiS2 / NiMOF; and P-W8O26.B2O3 / FeMOF.

[0197]

[0182] Embodiment 20. The heterostructure assembly or process of any one of the preceding Embodiments, wherein the metal organic framework (MOF) is a CoMOF, and the CoMOF is ZIF-67.

[0198]

[0183] Embodiment 21. The process of any one of Embodiments 2 to 20, wherein the solvent for the dispersion is a slightly polar solvent, such as an alcohol including methanol or ethanol, or an aqueous alcohol solution, preferably methanol in water.

[0199]

[0184] Embodiment 22. The process of any one of Embodiments 2 to 21 , wherein the basic organic ligand species is provided to the dispersion of inorganic template in solution form, for example, as an aqueous solution of dissolved organic ligand species.

[0200]

[0185] Embodiment 23. The process of any one of Embodiments 2 to 22, wherein the acidic non- noble transition metal ions are provided in solution form, for example, an aqueous solution of the acidic non-noble transition metal ions is provided to the dispersion of the inorganic template and organic ligand species.

[0201]

[0186] Embodiment 24. The process of any one of Embodiments 2 to 23, wherein the acidic non- noble transition metal ions of the 2D porous non-noble transition metal organic framework (MOF) are derived from a nitrate salt, for example, cobalt nitrate (e.g., Co(N03)2.6H20) for a CoMOF.

[0202]

[0187] Embodiment 25. A product obtained by or obtainable by the process of any one of Embodiment 2 to 24.

[0203]

[0188] Embodiment 26. A vertical layered A-[B-A]ntype heterostructure assembly of a 2D ultrathin inorganic solid-state material (‘B’) which is a non-noble transition metal oxide (TMO) and a non-noble metal organic framework (MOF) material (‘A’) that covers surfaces of each layer of the TMO in the assembly in a parallel arrangement to the 2D ultrathin inorganic solid-state material (‘B’), wherein n is a positive integer, wherein the non-noble transition metal oxide (TMO) comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; wherein heterointerfaces of the assembly comprise bridging MIM-X-MMOFionic bonds that stabilise heterointerfaces between the TMO and the MOF material, where bridging X is -O-, -S-, - Se, -C-, -N-, -P-, or -B-, MIMis a metal in the TMO, and MMOFis a metal in the non-noble MOF; and optionally, wherein the TMO is doped with one or more non-metal dopants selected from the group consisting of: B, P, C, N, O, Si, S, As and Se, preferably B, S, P or B and P.

[0204]

[0189] Embodiment 27. The heterostructure assembly of Embodiment 26, wherein the TMO is an iron oxide such as Fe20s or a tungsten oxide such as WO3; and optionally, wherein the TMO is doped with one or more non-metal dopants selected from B, S, and P, or wherein the TMO is doped with B and P.

[0205]

[0190] Embodiment 28. The heterostructure assembly of Embodiment 26 or Embodiment 27, wherein the TMO is doped with B and / or P.

[0206]

[0191] Embodiment 29. The heterostructure assembly of any one of Embodiments 26 to 28, wherein the TMO is B-Fe20s, WO3. B2O3, or P-WsO26-B2O3.

[0207]

[0192] Embodiment 30. A vertical layered A-[B-A]ntype heterostructure assembly of 2D ultrathin inorganic solid-state material (‘B’) which is a non-noble transition metal sulfide (TMS) and a non-noble metal organic framework (MOF) material (‘A’) that covers surfaces of each layer of the TMS (‘B’) in the assembly in a parallel arrangement to the 2D ultrathin inorganic solid-state material (‘B’), wherein n is a positive integer, wherein the non-noble transition metal sulfide (TMS) comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; wherein heterointerfaces of the assembly comprise bridging MIM-X-MMOFionic bonds that stabilise heterointerfaces between the TMS and the MOF material, where bridging X is -O- , -S-, -Se-, -C-, -N-, -P-, or -B-, MIMis a metal in the TMS, and MMOFis a metal in the non-noble MOF; and optionally, wherein the TMS is doped with one or more non-metal dopants, preferably selected from the group consisting of: B, P, C, N, O, Si, S, As and Se, preferably B, S, P, or B and P.

[0208]

[0193] Embodiment 31. The heterostructure assembly of Embodiment 30, wherein the TMS is an iron sulfide such as FeS2 or a cobalt sulfide such as C0S2 or a nickel sulfide such as NiS2; and optionally, wherein the TMS is doped with one or more non-metal dopants selected from the group consisting of: B, P, B and P.

[0209]

[0194] Embodiment 32. The heterostructure assembly of Embodiment 30 or Embodiment 31 , wherein the TMS is FeS2, C0S2 or NiS2 is doped with B and / or P.

[0210]

[0195] Embodiment 33. The heterostructure assembly of any one of Embodiments 30 to 32, wherein the TMS is B-FeS2, B-C0S2 or B-NiS2.

[0211]

[0196] Embodiment 34. A vertical layered A-[B-A]ntype heterostructure assembly of 2D ultrathin inorganic solid-state material (‘B’) which is a non-noble transition metal doped MXene or a non-noble transition metal doped functionalised MXene material (‘A’) and a non-noble metal organic framework (MOF) material (‘A’) that covers surfaces of each layer of MXene in the assembly in a parallel arrangement to the 2D ultrathin inorganic solid-state material (‘B’), wherein n is a positive integer, wherein the MXene comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; wherein heterointerfaces of the assembly comprise bridging MMXene-X-MMOFionic bonds that stabilise heterointerfaces between the MXene and the MOF material, where bridging X is -O-, -S-, -Se-, -C-, - N-, -P-, or -B-, MMXeneis a metal in the MXene, and MMOFis a metal in the non-noble MOF; and optionally, wherein the MXene is doped with one or more non-metal dopants, preferably selected from the group consisting of: B, P, C, N, Si, S, As and Se, more preferably B, S, P, or B and P.

[0212]

[0197] Embodiment 35. The heterostructure assembly of Embodiment 34, wherein the MXene is V2CTX, Ti2CTx, M02CTX or the functionalised MXene is TisC2O2, and wherein the MXene is doped with one or more non-metal dopants selected from the group consisting of: B, S, and P, or is doped with B and P.

[0213]

[0198] Embodiment 36. The heterostructure assembly of Embodiment 34 or Embodiment 35, wherein non-noble transition metal dopant of the non-noble transition metal doped MXene or non- noble transition metal doped functionalised MXene is Fe, Co, Ni, Cr or Cu and wherein the MXene is doped with one or more non-metal dopants for chloride shielding polyanion formation, wherein preferably the dopants are B and P.

[0214]

[0199] Embodiment 37. The heterostructure assembly of any one of Embodiments 34 to 36, wherein the 2D ultrathin inorganic solid-state material (‘B’) is Fe-B-Ti3C2Tx, wherein X is O.

[0215]

[0200] Embodiment 38. The heterostructure assembly of any one of Embodiments 34 to 37, wherein the MOF is: a CoMOF such as ZIF-67, an NiMOF, a ZnMOF or an FeMOF.

[0216]

[0201] Embodiment 39. Use of a heterostructure assembly of any one of Embodiments 1 or 25 to 38 as an electrocatalyst.

[0217]

[0202] Embodiment 40. Use of a heterostructure assembly of any one of Embodiments 1 or 25 to 38 as an OER electrocatalyst for water splitting.

[0218]

[0203] Embodiment 41. Use of a heterostructure assembly of any one of Embodiments 1 or 25 to 38 as an OER electrocatalyst for water splitting of seawater, preferably alkaline seawater, DI water, tap water or waste-water.

[0219]

[0204] Embodiment 42. The use of Embodiment 41 , wherein the OER is selective OER by inhibiting CER.

[0220]

[0205] Embodiment 43. The process of Embodiment 2, wherein the inorganic template for non-noble metal organic framework (MOF) material (‘A’) is an MXene material in freestanding or delaminated 2- dimensional (2D) sheet form, which is prepared by a method of forming a MXene-based OER electrocatalyst, comprising the steps of forming freestanding or delaminated 2-dimensional (2D) sheets of the MXene, Mn+iXnTx, where n is 1 , 2 or 3, and M is an early transition metal (Group 3 to 6 metal), X is C and / or N, and Txis a hydrophilic surface functional group selected from one or more of O, F, OH, and Cl, by dual modifying a functionalised MXene material, wherein the modifying steps involve: functionalising the MXene such that a majority of the surface functional groups in the MXene material are oxygen; replacing at least a portion of the M atoms in the MXene with one or more doped non-metal atoms selected from B, N, C, P, Si, S, As and Se; and replacing at least one surface functional group and / or at least one M with one or more electrochemically active metals selected from the group consisting of: Fe, Ni and Co.

[0206] Embodiment 44. A process involving inclusion of at least one non-metal dopant atom and at least one electrochemically active metal dopant atom in an O-functionalised MXene in a heterostructure to increase the activity of a MXene based OER electrocatalyst in the heterostructure, preferably wherein the one or more dopant non-metal atoms are selected from B, N, C, P, Si, S, As and Se; and preferably wherein the one or more electrochemically active metal dopants are selected from the group consisting of: Fe, Ni and Co.

[0221]

[0207] Embodiment 45. A method of forming a 2D / 2D heterostructured OER electrocatalyst that is selective for OER over CER, comprising forming a vertical layered A-[B-A]ntype heterostructure of a 2D ultrathin inorganic solid-state material (‘B’) and a non-noble metal organic framework (MOF) material (‘A’) that covers surfaces of each layer of the inorganic solid-state material (‘B’) in the assembly in a parallel arrangement to the 2D ultrathin inorganic solid-state material (‘B’), wherein heterointerfaces of the assembly comprise bridging MIM-X-MMOFionic bonds that stabilise heterointerfaces between the inorganic solid-state material (‘A’) and the MOF material, where bridging X is -O-, -S-, -Se-, -C-, -N-, -P-, or -B-, MIMis a metal in the inorganic solid-state material, and MMOFis a metal in the non-noble MOF; wherein the inorganic solid-state material comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; and further comprises one or more of: metal-0 (oxide), metal-N (nitride), metal-C (carbide), metal-C / metal N (carbonitride); metal-S (sulfide) and / or M-P (phosphide) functionalities, wherein n is a positive integer, and optionally wherein the inorganic solid-state material is doped with one or more non-metal dopants, preferably selected from the group consisting of: B, P, C, N, O, Si, S, As and Se.

[0222]

[0208] Embodiment 46. A method of improving the OER catalytic selectivity of a 2D / 2D electrocatalyst over CER, wherein the electrocatalyst comprises a vertical layered A-[B-A]ntype heterostructure of a 2D ultrathin inorganic solid-state material (‘B’) and a non-noble metal organic framework (MOF) material (‘A’) that covers surfaces of each layer of the inorganic solid-state material (‘B’) in the assembly in a parallel arrangement to the 2D ultrathin inorganic solid-state material (‘B’), by forming a bridging MIM-X-MMOFionic bonds at the heterointerface wherein the heterointerface stabilising the heterointerfaces between the inorganic solid-state material (‘A’) and the MOF material, where bridging X is -O-, -S-, -Se-, -C-, -N-, -P-, or -B-, MIMis a metal in the inorganic solid-state material, and MMOFis a metal in the non-noble MOF; wherein the inorganic solid-state material comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; and further comprises one or more of: metal- O (oxide), metal-N (nitride), metal-C (carbide), metal-C / metal N (carbonitride); metal-S (sulfide) and / or M-P (phosphide) functionalities, wherein n is a positive integer, and optionally wherein the inorganic solid-state material is doped with one or more non-metal dopants, preferably selected from the group consisting of: B, P, C, N, O, Si, S, As and Se.

[0223] Examples

[0224]

[0209] The present invention is described with reference to the following examples. It is to be understood that the examples are illustrative of and not limiting to the invention described herein. Example 1 : vertically layered Fe2O3 / CoMOF HS

[0225]

[0210] 2D / 2D ionic bonded Fe2C>3 / CoMOF was prepared using the facile interface-assisted synthesis technique of the invention at room temperature (RT). Initially, non-layered a-Fe20s was synthesised through a salt-template scalable approach and then, 2D CoMOF was constructed over and in parallel orientation to the surfaces of the 2D a-Fe2Os to produce a HS which is stabilised by strong ionic Fe- O-Co bonds at hetero in terfaces of the dissimilar component materials. The heterointerface was anchored with ionic bonds between metals and oxygen of dissimilar sheets (Fe-O-Co) to stabilise the structure and electrocatalytic properties. During electrocatalytic OER in real seawater, the heterostructured catalyst achieved 0.5 and 1 A cm-2current density at overpotential of 310 and 410 mV, respectively. Catalyst stability of more than 900 hours was demonstrated. For electrocatalytic performance testing, the synthesised catalysts were drop cast on the Ni foam to make the working electrodes. The inorganic template in this example was not doped with any other atoms, including metal atoms and / or non-metal atoms.

[0226] Preparation of 2D Fe2Os nanosheets through salt template method

[0227]

[0211] In a typical synthesis process, 1 g of iron(lll) nitrate nonahydrate was dispersed in 20 mL of ethanol followed by 30 min magnetic stirring. The obtained precursor solution was added dropwise to 310 g of NaCI salt and mixed for ~20 min at ~80 °C on a hotplate. The material was separated and annealed in a tube furnace under N2 at 400 °C for 2 hours to obtain 2D a-Fe2Os after filtration.

[0228] Growth of 2D / 2D Fe2O3 / CoMOF (MOF@FEO) heterostructure

[0229]

[0212] The growth of 2D / 2D Fe2C>3 / CoMOF HS is as follows. Typically, 48 mg of 2D a-Fe2Os was dispersed in a 192 mL of methanol / water mixed solution (2:1 in volume) through 1 .5 hours sonication to get solution A. Then, 96 mL methyl imidazole (25 mM) solution and 96 mL Co(NC>3)2.6H20 solution (25 mM) solution (1 :1 ) were added to the solution A one by one, respectively. The obtained solution B has 2D / 2D HSs of CoMOF supported on a-Fe2Os template at room temperature (RT). The product was centrifuged, followed by methanol and water washes to remove the unreacted reagents and finally vacuum dried. For electrocatalytic performance testing, the synthesised catalysts were drop cast on the Ni foam to make the working electrodes.

[0230] Electrochemical testing and Ink preparation for the electrodes

[0231]

[0213] Electrochemical testing was done in 6 M KOH DI and 6 M KOH real seawater electrolyte at ambient temperature and pressure using a CHI 760D workstation. Saturated Hg / HgO and graphite rod was used as the RE and CE, respectively. The WE was prepared by drop-casting the samples on nickel foam which was washed with dilute (1 M) HCI, acetone and water before material drop-cast. The polarisation curves were recorded after iR correction except testing performed on Metrohm Autolab workstation. 20 mL carbon solution was prepared by adding 20 mg of carbon black to a 20 mL mixture of isopropyl alcohol (IPA) and water in 4:1 , respectively. For making ink, 1.25 mL of prepared carbon solution was sonicated for 90 min, and then 5 mg of catalyst powder and 5 pL of PTFE was added and sonicated for 20 min to obtain homogenous solution. 100 pL homogenous ink solution was drop-cast on Ni foam. The catalyst loading was 1 .6 mg cm-2on working electrode.

[0232] Results & Discussion for Example 1

[0214] Figure 1(a) outlines the construction strategy for 2D / 2D Fe2C>3 / CoMOF HS with a functional heterointerface. TEM images of Fe2C>3 / CoMOF HS, a-Fe2Os and CoMOF are presented in (Figure 1(b)) which indicate the development of the sheets for the prepared samples. Figure 1(b) HS reveals that the CoMOF is successfully grown as a parallel sheet over the Fe20s to develop HS through the solid-liquid interfacial synthesis. The basal planes of the two dissimilar nanosheets are strongly linked and grown on each other. AFM demonstrates nanosheets of the 2D Fe2Os, 2D CoMOF and the HS with a thickness of -3, -4 and -7 nm, respectively. The increase in thickness for 2D / 2D HS indicates the successful growth of CoMOF on the Fe2Os nanosheets.

[0233]

[0215] EDS unveils uniform distribution of C, Co, N, O and Fe elements throughout the HS, confirming the successful substantial parallel coverage of the 2D Fe2Os solid template by the CoMOF (Figure 1(c)). The Fe and Co are homogeneously distributed in the HS, indicating the presence of Fe2Os and CoMOF in Fe203 / CoMOF respectively and confirming the sheets are grown on top of each other. HR- TEM image in Figure 1(d) presents a lattice spacing of 0.27 nm corresponding to the (104) of Fe2Os in the HS, confirming the development of the heterointerface between two sheets.

[0234]

[0216] XRD (Figure 1(e)) pattern indicates the formation of pure hematite (a-Fe2Os) and characteristics peaks of Fe2Os are well-matched in the case of Fe203 / CoMOF indicating the existence of dissimilar sheets which retain crystallographic features of the components in case of the HS. The strong characteristic diffraction peak of CoMOF nanosheets at 10.44° (002) is shifted to 9.95° in the case of Fe203 / CoMOF HS. The characteristics peaks 33.4° and 35.84° attributed to the (104) and (110) planes of Fe2Os are present at the same place in the case of Fe203 / CoMOF. It confirms the presence of both CoMOF and Fe2Os sheets while retaining the individual crystallographic features of both distinct constituents.

[0235]

[0217] FTIR analysis is also performed to see the structure, coordination of organic linker with the metal centres and connection between the dissimilar sheets. Organic components in CoMOF and Fe203 / CoMOF HS are observed through peaks between 1000-1650 cm-1. The peak between 1350- 1500 cm-1is attributed to the stretching ring of 2-methylimidazole, also seen in Fe203 / CoMOF. For Fe2Os, peaks at 717.7 and 538.3 cm-1are ascribed to the Fe-0 stretching. For CoMOF, the peak at 507.8 cm-1shows the presence of the Co-N vibrations. FTIR indicates the presence of Fe-O, Co-O, and Co-N bonds through peaks at 717.7, 536.6 and 432.5 cm-1and shifts in the metal-nitrogen or metal-oxygen peaks in the case of the HS.

[0236]

[0218] To understand the nature of bonding at heterointerface, DFT calculations were performed. The relaxed structures of the CoMOF fragment on the Fe-terminated and O-terminated surfaces of the Fe2Os substrate are shown in Figure 1(f)-(h). For the Fe-terminated CoMOF / Fe203 system, the fragment can adsorb in two different locations and orientations on the substrate, as shown in Figure 1(f) and (g). In Figure 1(f), the Co atom forms bonds with three surface O atoms, with bond lengths calculated to be -2.10 A, which is consistent with the Co-0 bond length in bulk cobalt (II) oxide (2.13 A), indicating the formation of ionic bonds. The N atoms in each methylimidazole group form bonds to a surface Fe atom. The methylimidazole ring that is oriented perpendicular to the surface has an N- Fe bond length of 2.05 A, which is close to the N-Fe bond length in the nitrogen-doped Fe2Os surface. The other methylimidazole ring lies parallel to the surface, resulting in a slightly longer N-Fe bond length of 2.23 A.

[0237]

[0219] For the optimised Fe-terminated CoMOF / Fe2C>3 system (Figure 1(f)), the Co atom is bonded to one surface Fe atom, having a bond length calculated to be 2.35 A. The only other direct bond to the surface is between one of the N atoms and a surface Fe atom having an N-Fe bond length of 2.06 A. The CoMOF fragment is more strongly bound in the first orientation due to the greater number of bonds formed with the surface. The presence of the Co-0 bonds agrees with the experimental characterisation, where the main interactions observed between the CoMOF fragment and the substrate were between the Co atom and a surface O atom.

[0238]

[0220] To experimentally observe the bonding as predicted by DFT, XPS analysis is performed on prepared structures. The Co 2p deconvoluted spectra peaks (Figure 2(a)) at 781 .6 and 797.4 eV for CoMOF and 780.9 and 796.8 eV for Fe203 / CoMOF ascribed to the Co-N bonds showing the formation of the CoMOF coordination between the inorganic part and imidazole ligands. For Fe203 / CoMOF, the Co valence state changes from Co2+to Co3+’ attributing to the new Co-0 bond formation between the different interfaces to develop the HS. Peaks located at 781 .39 and 797.0 eV belong to the Co3+state and appeared by the Co-0 bond. In addition, deconvoluted N 1 s (Figure 2(b)) peaks 399.4 eV (Fe203 / CoMOF) and 399.3 eV (CoMOF) are attributed to the coordinated bonding of nitrogen in the electron donating organic ligand with the metal centres. The change in the peak intensity appears in the case of the HSs due to the formation of new N-Fe bonds. The peak at 400.8 eV (Fe203 / CoMOF) and 400.7 eV (CoMOF) is attributed to the N in the imidazole linker. The O 1 s XPS spectra (Figure 2(c)) at 529.8 and 529.9 eV for Fe203 / CoMOF and Fe2Os, respectively, display the interaction of oxygen with the metal centre. The peak positioned at 531 .6, 532.5 and 531 .5 eV in Fe203 / CoMOF, CoMOF and Fe2Os ascribed to the adsorbed moisture and oxygen. XPS data confirms the chemical composition and supports the presence of Fe2Os and CoMOF in the HS while maintaining the characteristics of the individual 2D sheets. The development of heterointerface changes the oxidation state of the metal nodes of the CoMOF to create strong ionic interactions between the sheets.

[0239]

[0221] XPS C1s deconvoluted spectra of CoMOF and Fe203 / CoMOF indicate that it is divided into three peaks. In Fe203 / CoMOF, the peaks located at 284.6, 285.2, and 288.2 eV are assigned to the presence of C-C / C=C, C-N and C=N groups, respectively. Similar three peaks for these groups are present in CoMOF at 284.7, 286.0, and 288.3 eV, respectively. This shows the presence of the characteristic group belonging to imidazole linker in MOFs and the negative peak shift in the Fe203 / CoMOF sample due to the local environment change. A doublet peaks of Fe 2p at 710.9 and 724.5 eV for Fe203 / CoMOF and 710.8 and 724.4 eV for Fe2Os indicates the presence of the Fe-O / Fe- N and Fe-O, respectively. The surface site interaction with nitrogen at basal planes slightly shifts the binding energy. The satellite peak at 719.0 eV between Fe 2ps / 2 and Fe 2pi / 2 is assigned to Fe3+in hematite Fe2Os.

[0240]

[0222] NEXAFS was used to extend the understanding of the local atomic and electronic structures of metals and non-metals in CoMOF, Fe2Os and Fe203 / CoMOF. The Co L-edge NEXAFS spectra (Figure 2d) of Co L3 and L2 split edges displayed the prominent peaks 779.9 eV and 794.6 eV, respectively for Fe2C>3 / CoMOF and 779.9 eV and 794.8 eV, respectively for CoMOF due to the transition of 2ps / 2, 2pi / 2 into vacant 3d orbitals dipole transitions (Co2+) by coupling of 2p core-electron spin-orbital suggesting change in the electronic state. The difference in split Ls-edges characteristics of 3d-electrons indicated the development of new bond formation (Co3+) of Co atoms by disappearing of some peaks and shifting of the peak positions towards the high photon energy in HS, i.e., 777.1 eV and 779.0 eV as compared to pristine CoMOF, i.e., 776.8 eV and 778.8 eV respectively. These changes are attributed to ligating new Co-0 bonds in HS due to bridging between Fe2Os and CoMOF. Moreover, in the case of HS, the intensity of the peaks changes due to new interactions of Co metal with O of Fe2Os along with Co-N. These new ionic connections between dissimilar sheets ensuring enhanced stability of the developed HS. The N K-edge NEXAF spectra (Figure 2(e)) depicted that the 388.5 eV, 389.3 eV, 390.0 eV, 391 .1 eV, 396.8 eV, and 397.4 eV peaks in both HS and CoMOF are well matched with each other except the reduction of the intensity due to bonding between heterointerface. The peaks at 389.3 eV and 397.4 eV result from overlapping N K-edge features with Co to exhibit second harmonic Co Ls and L2 edge signals, respectively. These peaks below photon energy of 400 eV ascribed to the IT* resonance peak attributed to the presence of pyridinic nitrogen in the ligand. The peaks after 400 eV exhibited the o* peaks due to the pyrrolic N in ligand structure and the interaction of N with Co to form the coordination between organic and inorganic parts to develop the MOF structure. The peaks in the HS were located at 401 .5 eV and 405.2 eV, respectively, whereas in CoMOF, they were located at 401 .4 eV and 405.0 eV, respectively. The positive shift in the peak positions is attributed to new Fe-N interactions created as well between dissimilar 2D sheets.

[0241]

[0223] Furthermore, the O K-edge show the two peaks at 530.6 and 531 .8 eV, which, after the MOF growth, leads to broadening and change in intensity might be due to the competition between the Co and Fe for the oxygen and formation of new ionic connections of Co with the oxygen (Fe-O-Co). Thus, supporting the DFT findings of ionic connections development for HS. At the same time, the O K-edge in the case of MOF shows a high-intensity peak compared to Fe2Os due to loosely adsorbed environmental oxygen on the surface of MOF. At the Fe L-edge spectra (Figure 2(f)), there are mainly two regions, i.e., L3- and l_2-edge, with prominent peaks at 709.6 eV and 722.9 eV owing to the electronic transition of Fe 2ps / 2 and Fe 2pi / 2 core electrons correspond to the interaction of -13.3 eV spin-orbit splitting. The overall resemblance in the L3- and l_2-edge in Fe2Os and Fe203 / CoMOF HS are due to the similarity in the 3d ground-state configuration, p-d Coulomb magnitude and interaction of the crystal-field. The lowering of spectral intensity and splitting of the peak at 721.3 eV in Fe2Os into two peaks, 721.2 eV and 721.4 eV in Fe2C>3 / CoMOF HS, indicates the formation of new interactions as Fe-N.

[0242]

[0224] For the accurate comparison of the OER catalytic performance in the 6M KOH in DI water and real seawater, the electrodes comprising various prepared catalysts are designed at the same conditions. The electrolytic OER performance of CoMOF, Fe2Os, Fe203 / CoMOF and lrO2 were measured by LSV. The Fe203 / CoMOF HS displayed excellent performance compared to the individual nanosheets and lrO2, as presented in (Figure 3) as a result of synergistic coupling effects from the two ionically bonded components, manipulating the electronic structures by facilitating the charge delivery and fast charge and mass transport between dissimilar interfaces. Interestingly, Fe2C>3 / CoMOF HS catalyst demonstrates exceptional OER activity by achieving the current density of 0.5, 0.8, 1 and 1.2 A cm-2at an overpotential of 310, 370, 410 and 450 mV, respectively, in direct seawater electrolytes (Figure 3(a) and 3(b)) and 161 , 269, 380 and 485 mV, respectively in alkaline DI electrolyte, as determined by LSV. The HS can achieve ampere level current densities not only in DI water but also in seawater electrolytes. Moreover, Figure 3(c) compares the CoMOF, Fe20s, Fe2C>3 / CoMOF and lrC>2 catalysts overpotentials at 500, 800 and 1000 mA cm-2current densities showing that the HS offers the best OER activity among these catalysts with low overpotentials of 310 and 410 mV for 500 and 1000 mAcmr2current densities. Moreover, HS requires low energy of 310 and 370 mV for 500 and 800 mA cm-2higher current densities, making it suitable for commercial scale seawater splitting.

[0243]

[0225] Besides, the Tafel slope of HS (61 .6 mV dec-1) is much smaller than Fe2Os (73.7 mV dec-1), CoMOF (101.5 mV dec-1) and lrO2 (118.4 mV dec-1) as shown in Figure 3(d) revealing fast reaction kinetics. The electrochemical double-layer capacitance (Cdi) of the catalysts obtained from CV measurement are presented in Figure 3(e). 2D MOF with (701 mF cm-2) has a Cdi higher than 2D Fe20s (558 mF cm-2), and the additional presence of Co-based active species contribute to enhance the OER capabilities of HS in seawater. The hybridisation of CoMOF with Fe2Os sheets to form the HS increased the Cdi further to 1221 mF cm-2because the Fe2Os nanosheets might provide flexible supports to the CoMOF layers to prevent them from agglomeration and thus maximise the exposed active surface resulting in fast access of the liquid electrolyte to the active sites and achieving higher current density as compared to individual catalysts.

[0244]

[0226] Long-term stability is a vital index to assess the capability of the catalyst, and especially in the complex seawater environment, it is more challenging to avoid corrosion and maintain performance for practical applications. The developed heterostructure shows durable performance even after 900 plus hours (Fig. 3 (f) & (g)).). The catalytic performance also remains stable as displayed by the LSV testing before and after the stability testing (Figure 3(f)). The ionic-connected HS displays exceptional stability compared to individual counterparts, even at more than 300 mA cm-2current density. Interestingly, Fe2Os decays quicker than the MOF during seawater oxidation. That means ionic bonds along with surface chemistry of MOF synergistically provides stability to the HS. The excellent durability of the ionic bonded HS suggests that catalyst does not degrade by avoiding the CER. To confirm this, continuous gas is generated at the anode electrode and that gas is passed through the acetone solution at room temperature. If chlorine is present acetone turns instantly into chloroacetone and HCL The acetone + gas mixture is analysed by GC-MS. Analysis shows only acetone peaks and no chloroacetone peak in the acetone + gas mixture. Hence, confirming the selectivity of anode towards OER and inhibition of the CER, even at higher potentials. Overall, the durable performance of the HS at higher current densities I overpotentials is attributed to the higher surface area to allow easy access to active sites, while avoiding CER which avoid catalyst degradation, synergistic effects and robust ionic connections among the heterointerface.

[0227] To comprehensively reveal that the MOF in the OER prevents CER from competing with the OER, DFT was performed by adsorbing Cl- and OH- on the surface in different locations. To balance the charge, a Na+ion was added to the supercell whose position was fixed far enough from the surface to not interact with it or the anion. Figures 3(h) & 3(i) show the relaxed structures of the CoMOF / Fe2C>3 with adsorbed OH- or Ch. For the CoMOF / Fe2C>3-OH system (Figure 3(h)), the OH- bonds, via the O atom, to a C atom on the methylimidazole group that is oriented perpendicular to the surface. The O- C bond length is 1 .40 A. For the CoMOF / Fe2C>3-CI system (Figure 3(i)), the Ch did not adsorb on the MOF fragment but instead preferred to adsorb directly on the surface to a Fe atom, forming a Fe-CI bond that is 2.20 A. The Cl atom is 3.31 A away from the CoMOF fragment. This indicates that it is unlikely the Ch will bind to the MOF, reducing the likelihood of the CER occurring. These results further supported the hypothesis that the growth of the CoMOF over the Fe20s surfaces provides Ch corrosion resistance to the system along with more active sites, which is displayed in the experimental results. The calculated binding energies of CoMOF / Fe203-OH system is -2.66 eV, while that of the CoMOF / Fe203-CI system is -2.00 eV. While the Ch is more strongly bound, it is bonded to the surface, not the MOF. However, the experiments indicate that the MOF covers the surface; hence, the Ch is not likely to adsorb to the HS. Therefore, the CER cannot occur in these ionically designed HSs, but OH- is expected to adsorb and contribute to the enhancement of OER performance.

[0245]

[0228] To examine the impact and influence of the electrochemical testing on the morphology and structure of the catalyst, TEM, EDS, NEAXFS and XPS were performed. TEM image (Figure 4(a)) indicates that the morphological structure is well maintained even after the electrochemical testing. MOF sheets cover the 2D Fe20s via strong ionic bonds while keeping the morphological features. Moreover, HR-TEM in Figure 4(b) shows the lattice spacing of 0.25 nm corresponding to the (110) plane of Fe2Os in the HS having a heterointerface (highlighted by white dotted lines) with the CoMOF. EDS analysis (Figure 4(b) inset) also confirms the presence of Fe20s and CoMOF-related elements. The elemental maps show the uniform distribution of Fe, Co, O, and N in the sample, indicating that the HS features are retained, and sheets are well intact even after catalytic testing. The ex-situ NEXAFS data indicates that the Co L-edge NEXAFS spectra (Figure 4(c)) of Co Lsand L2 split edges displayed the almost similarly located prominent peaks 780.1 eV and 794.6 eV, respectively, for HS. The split l_3-edges show the presence of the Co3+state due to Co-O-Fe in HSs, supporting the stability results that the HS retained their features even after the testing. The O-K edge spectra (Figure 4(d)) demonstrate the presence of a peak at 542.3 eV at a similar location but shift and change in the peak intensity below 535 eV due to the surface oxidation and adsorption of environmental oxygen. The XPS spectra of Co (Figure 4(e)) also highlight the chemical stability of the HS even after the electrochemical testing. The presence of the Co3+and Co2+oxidation states in the catalyst indicates that the CoMOF structure and ionic connections are present. These ex-situ characterisations support the maintaining of electrochemical performance in the complex seawater system and proving HS as a stable catalyst which can operate at ampere level current densities without the CER.

[0246] Summary - Example 1

[0229] 2D HS formed via the liquid-solid interfacial growth at room temperature (RT) using the process of the invention. CoMOF is grown at the basal plane of the a-Fe2Os to form a 2D-D Fe2C>3 / CoMOF HS. A change in the oxidation state of the metal present in nodes develops the ionic bonded interface. The MOF grown over the metal oxide provides catalyst protection from Cl- attack and encourages OH adsorption over the Cl ’ resulting in better catalytic activity and stability. Heterointerface offers better electron transfer, new active sites and firmly bonded, more stable structures. Heterointerface provides the pathway for the rapid charge transfer between CoMOF and Fe20s to display better reaction kinetics as compared to pristine CoMOF and Fe20s. The OER performance (LSV, electrochemical surface area and stability) is enhanced as compared to the individual 2D pristine materials in alkaline solution in DI water and seawater electrolytes attributed to the interaction of the Co and O bonding between CoMOF and Fe20s providing fast charge and mass transport between dissimilar interfaces and creation of active sites at the heterointerface.

[0247] Example 2: vertically layered B-doped WOa / CoMOF HS

[0248]

[0230] C0-MOF / WO3.B2O3 (MOF / WB) HS was constructed via the solid-liquid interfacial growth process of the invention at room temperature (RT). As predicted by theoretical calculations, B relaxes away from the uppermost surface layers, exposing the top W-0 surface as the best possible site for the adsorption of Co atoms within the MOF to construct the 2D MOF morphology and ionically (W-O- Co) interlinked HS. The DFT studies and experimental observation confirms the materials interact via formation of strong ionic bonds, which unite the MOF and substrate as single new material rather than simple composite or hybrid structures. This unique construction is key to creating new material properties that provide an exceptional surface and interface to inhibt chlorine chemistry and activate the catalyst for extraordinary performance. The catalyst’s preference to bond hydroxyl ions over chloride ions not only prevails at the surface but also at the interface where introduced B2O3 polyanion functionality acts as a B-OH modulator. To obtain such an interface, it was observed that the doping concentration of B can also play a role, with the doping position adjacent to W just beneath the top layer suggested to be the best position. The HS was able to operate stably for more than 1100 h in harsh seawater and maintained stable operation even after sitting idle in the electrochemical setup for 5 days.

[0249] Synthesis of WO3.B2O3

[0250]

[0231] WO3.B2O3 was synthesized by a facile salt template method using KCI salt. Boric acid (30, 65,130, and 195 mg) was added to 12 mL DI water and mixed via sonication for 2 mins to achieve the different B concentrations. Then, a mixture of 300 mg ammonium metatungstate in 6 mL ethanol was added to the boric acid solution and again sonicated for 2 mins to obtain proper dispersion. Simultaneously, 30 g of KCI salt was heated on a hot plate for 5 mins at a temperature of 80 °C, after which the precursor mixture was added dropwise and stirred continuously until the template had dried completely. The precursor salt was then left for 2 days for the growth of the sheets and then annealed at a temperature of 650 °C for 3 h at a ramping time of 3 h. The annealed material was filtered and washed with ethanol and left for overnight drying in a vacuum oven at a temperature of 50 °C. The as- synthesized samples were named WB-0.1 , 0.21 , 0.43, and 0.65 based on the boric acid and ammonium metatungstate ratio.

[0251] Synthesis of Co-MOF@WO3.B2O3

[0252]

[0232] C0-MOF / WO3.B2O3 (MOF@WB-x) was synthesized using the obtained WO3.B2O3 as a template for the growth of Co-MOF sheets at room temperature. 30 mg of as-synthesized WO3.B2O3 was added to a 44 mL mixture of methanol and water (2:1 ) and sonicated for 1 h. Simultaneously, a mixture of 45.46 mg of 2-methyl imidazole linker and 159.86 mg of Co(NC>3)2.6H20 was prepared separately in 22 mL DI water. The linker was added to the sonicated WO3.B2O3 solvent, followed by the addition of cobalt salt, and left undisturbed for growth. After 3 hours, the obtained mixture was washed first with methanol, then with water, and finally with methanol. The C0-MOF / WO3.B2O3 obtained was dried for 4 hours in a vacuum oven at 50 °C.

[0253] Synthesis of Co-MOF

[0254]

[0233] A mixture of 159.86 mg of Co(NC>3)2.6H20 in 22 mL DI water and 45.46 mg of 2- methylimidazole in 22 mL of DI water was added to a 44 mL mixture of methanol and water (2:1 ) and allowed to react at room temperature without stirring for a minute. The product was then centrifuged and washed with methanol, water, and methanol again and then dried for 4 h in a vacuum oven at a temperature of 50 °C.

[0255] Electrochemical measurements

[0256]

[0234] The electrochemical testing was done in 6.0 M KOH and real seawater electrolyte using a CHI 760D workstation, Autolab PGSTAT204 with BOOSTER10A (Metrohm) in a three-electrode system at room temperature (RT) where a graphite rod was used as the CE, and a saturated calomel electrode (SCE) was used as the RE. The WE was prepared by drop-casting the samples on nickel foam, which was washed with dilute HCI, water, and ethanol before electrode preparation. The polarization curves were recorded with a scan rate of 5 mV s-1without iR correction. The powdered sample was added to the carbon black solution (4:1 ), followed by the addition of PTFE solution (60% in H2O, sigma) and sonicated to form a homogeneous ink. The carbon black solution was prepared by adding 20 mg carbon to a 20 mL mixture of isopropyl alcohol (IPA) and water (4:1 ). The solution was then sonicated for 1 h to acquire proper dispersion. Then, 200 pL of the ink was drop-cast onto an area of 0.25 cm-2on nickel foam. For determining the electrochemical double-layer capacitance (Cdi), electrochemical CV measurements were performed over a range of scan rates (10, 20, 30, 40, and 50 mV s-1). The stability tests were also performed at a fixed potential. All the potentials are reported vs RHE.

[0257] Results & Discussion for Example 2

[0258]

[0235] To facilitate creating of the ionic bond, the composition and surface of WO3 were first tuned by introducing B and B2O3 functionality. Then, the inherent rod-like structure of WO3 was converted to sheets through a series of temperature-dependent experiments where a sheet-like morphology was obtained at 650°C. TEM images reveal that with increasing temperature, the thin rods merged together to produce a 2D platform with a monoclinic structure which created an ideal template for parallel growth of 2D Co-MOF sheets hereon to realize well-defined MOF@WB HSs (Figure 5(a)). The B dopant assists in better adsorption of the MOF constituents on the solid WO3 sheets by relaxing away from the uppermost surface layers, creating a guided surface at the liquid interface for growth of the 2D MOF and ionic bond formation. TEM images confirm complete integration of the WB and ultrathin MOF sheets (Figure 5(b)). Moreover, microstructure analysis using selected area electron diffraction (SAED, inset of Figure 5(a)) and a lattice spacing of 0.191 nm corresponding to the (004) of WO3 observed by high-resolution TEM (HRTEM) (Figure 5(c)) confirmed that the monoclinic WB structure is well preserved even after chemical bonding with Co-MOF via ionic bonds.

[0259]

[0236] The bulk structural analysis of WB, Co-MOF, and MOF@WB done using powder (P-XRD, Figure 5d) also confirmed the monoclinic structure of WB and its HS. The peaks around 15, 27, and 36° verify the existence of B2O3 and borate polyanion functionality in both samples. The reduced peak intensity for MOF@WB, except at 11 ° (MOF characteristic peak), confirmed the integration of MOF sheets with WB to construct the HS. Raman spectroscopy analysis showed O-W-O bending (251 , 300, and 320 cm-1) and 0-W6+-0 stretching (717 and 780 cm-1) peaks in both samples, peaks which are characteristic of monoclinic WO3. Similarly, the peak at 878 cm-1is attributed to borate species in both samples. The new peaks at 470 and 670 cm-1in the HS correspond to the surface of the Co-MOF. EDS maps showed a uniform distribution of all elements across the sheets, confirming the presence of both the MOF and WB in the HS.

[0260]

[0237] DFT and ab initio molecular dynamics (AIMD) calculations were conducted to understand the impact of B on the structure of WO3 and its role in creating a unique heterointerface via ionic bonding. The optimized structures of the Co-MOF unit on the three possible surfaces of WO3 are shown in Figure 5(e)-(g). For Co-MOF / WO3(001 ), the Co-MOF fragment adsorbs on the surface, forming a bond between a surface O atom and the MOF Co atom. One of the methylimidazole rings lies parallel to the surface, while the other is orientated perpendicular to the surface (Figure 5(e)). The Co-0 bond length was calculated to be 1 .83 A, which is close to the Co-0 bond length in bulk cobalt (II) oxide (2.13 A), indicating formation of an ionic bond. On the B-doped W0s(001 ) surface, the B atom relaxes away from the uppermost surface layers so that it is situated below the outermost layer of O atoms (Figure 5(f)). After adsorption of the Co-MOF unit, the binding is similar to that on the pristine WO3 surface.

[0261]

[0238] On the B20s(001 ) surface, the Co-MOF fragment bonds to the surface through formation of a Co-B bond (1 .99 A) and two C-B bonds (1 .65 A each), as shown in Figure 5(g). At the same time a proton from the methyl group of one of the methylimidazole rings is removed and bonds to a surface B atom (1 .2 A). The Co-MOF fragment was calculated to bind more strongly on the B20s(001 ) surface, followed by the B-doped W0s(001 ) surface and then the WO3 surface. This indicates that the presence of B in the WO3 surface enhances the bonding of the Co-MOF to WO3. The presence of strong Co-0 bonds shows that the main interaction is between the Co atom from Co-MOF and an O atom from WO3. Further, the presence of B facilitates the establishment of a surface with the capability to provide an ideal platform for forming a unique ionically bonded heterointerface with enhanced interactions between Co-MOF and the surface, leading to improved structure and properties of the resulting HS.

[0262] Ionic bonding and electronic structure

[0239] XPS was carried out to understand the surface chemistry of the MOF@WB and individual components. Figure 6(a) shows deconvoluted Co 2p XPS spectra of MOF@WB with peaks at 780.6 and 796.5 eV attributing to Co3+and peaks at 782 and 797.9 eV corresponding to the Co2+state which can be assigned to Co-0 and Co-N bonds, respectively. While Co 2p spectra of Co-MOF indicate only Co2+oxidation state and satellite peaks due to Co-N bonds between Co atoms and N from the imidazole groups. Interestingly, the presence of the Co-0 bond only in the HS confirms that the interaction between the MOF and WB occurs between the Co in MOF and O in WB. The shift in binding energy of the Co 2p spectra also supports formation of ionic bonds between the Co-MOF and WB, which was also concluded by DFT calculations. The deconvoluted W 4f spectra show a well-resolved doublet at 35.2 and 37.3 eV, attributing to the W6+state and an extra peak at 33.12 eV in the WB, which is attributed to the W4+state. The increase in W6+states in the HS and a shift in binding energy indicate charge transfer between WB and Co-MOF as a result of the formation of ionic Co-O-W bonds (Figure 6(b)). The O 1s spectra show a peak at 530.4 eV corresponding to M-0 bond in WB which is absent in the MOF verifying no Co-0 bond existed. However, a shift is observed in the O 1s spectrum of the HS is thought to be due to formation of Co-O-W bond.

[0263]

[0240] The B 1s deconvoluted spectra of WB show peaks around 191 and 186 eV corresponding to the B-0 species and elemental B. However, no such peaks were observed in the HS, which is thought to be due to concealment by Co-MOF. Therefore, to confirm the presence of B in the HS, ICP-MS was done, demonstrating similar B concentration in both WB (~69 ppb) and MOF@WB (~66 ppb). The N 1 s spectra of Co-MOF (Figure 6(c)) showing Co-N bond and pyrrolic N which confirms the coupling between Co, N, and C as supplemented by the C 1s spectra of Co-MOF which shows C=N bond, while the N 1 s spectra of the HS show pyridinic N, Co-N, and -NO2 species. These modulations of the N bonds arise due to the creation of the W-O-Co ionic bond in the HS, which alters the electronic situation of N.

[0264]

[0241] NEXAFS analysis was carried out to probe the redistribution of charge states and regulation of electron density of different elements in response to B concentration and ionic bond formation. The Co L-edge NEXAFS spectra (Figure 6(d)) show Co L3 and L2 edges around 780 and 794 eV for the Co2+state resultant of the 2ps / 2 and 2pi / 2 to 3d transition. The peak at 776.9 eV for both the Co-MOF and the HS is attributed to the octahedrally coordinated Co, while the peak around 778 eV in the HS corresponds to the tetrahedrally coordinated Co. The existence of octahedral and tetrahedral coordination in the HS indicates the presence of an additional bond, i.e., the W-O-Co bond. The feature at 778.7 eV is assigned to metal-to-ligand charge-transfer transitions, while the wide shoulder around 781 .4 eV is due to the electric dipole-allowed transition. The shift in the Co peak positions in the HS compared to the Co-MOF demonstrates the transfer of electrons between both constituents via Co-O- W bonds. N K-edge spectra show peaks around 389 and 397.2 eV corresponding to second-order L3- edge and l_2-edge for Co, where L3 is more prominent than L2, suggesting the bonding between Co and N. The peaks around 401 .4 and 405 eV are attributed to the pi* and sigma* peaks for C-N from the imidazole ring, confirming the coupling between Co, N, and C. The intensities of these features are significantly reduced after the growth of Co-MOF over WB as the ionic connection induces a change in the C-N bond, supporting the observation made in N 1s XPS spectra.

[0265]

[0242] The B K-edge spectra of WB show a sharp peak at 191 .85 eV corresponding to the B1 s state to pi* resonance absorption (Figure 6(e)). Two weak peaks at 192.5 and 193.15 eV represent point defect oxygen atoms and the peak at 193.85 eV corresponds to the pi* resonance peak of B2Oswith the most significant intensity in WB-0.21 (0.21 represents B concentration). Small intensity peaks at 193.85 eV in the HSs also confirmed the presence of B2O3 close to the interface between WB and Co- MOF. The O K-edge (Figure 6(f)) for WO3 has two peaks, where the first main peak at 530.6 eV demonstrates the O 2p-state in the d(t2g) band arising from the 5d W and 2p O orbitals. The larger intensity of the first peak in comparison to its shoulder peak at 531 .9 eV is an indication of a more covalent W, i.e., W6+state. The shoulder peak arises due to the first coordination shell O atoms, which are non-equivalent, and the intensity of these peaks depends on the empty 2p O states. So, after B introduction (which can cause a partial positive charge on W / O), the second peak increased in intensity and shifted 0.2 eV towards lower energy, which might be due to the competition between W and B in the acquisition of O. The O K-edge of the Co-MOF shows peak in the region of 528-534 eV associated with the excitation of O 1s electrons to 0-2p / Co-3d orbitals while the broader peak around 540 eV corresponds to the excitation to 0-2p / Co-4sp orbitals demonstrating the presence of Co2+oxidation state, as observed in the Co L-edge and XPS Co-MOF spectra as well. The O K-edge results for the HSs demonstrate peaks present in both WB and Co-MOF with reduced intensity and a 0.1 eV shift towards higher energy than WO3, implying interfacial interactions between Co-MOF and WB resultant of stronger electronic interactions due to the ionic Co-O-W bond. The NEXAFS findings, thus, corroborate the DFT results that MOF and WB are bonded via ionic Co-O-W bonds.

[0266] Electrochemical reaction at the anode

[0267]

[0243] To probe the efficiency and role of various components in the HS, initially, WB samples with different boron concentrations were tested, as shown in Figure 7(a). Prominent oxidation peaks (~1 .29 V) were observed in the linear sweep voltammetry (LSV) curves for WB-0.1 and -0.65 while no such peaks were observed for WB-0.21 and -0.43 suggesting that regulated B concentration prevents unwanted surface oxidation. Further, heterostructuring WB with MOF surpasses not only individual WB but also lrO2 in both deionized water and real seawater (Figure 7(b)). MOF@WB-O.21 achieved a current density of 500 mA cm-2at 462 mV, whereas the benchmark lrO2 required 776 mV in seawater (Figure 7(b)). The HS outperforms commercial lrO2 in terms of efficiency by ~10x and can deliver a current density of 2 A cm-2at only 1 .97 V in seawater (Figure 7(c)), meeting the industry needs. The mass activities were also calculated as 198.1 , 154.3, and 20.8 Ag-1for MOF@WB-O.21 , WB-0.21 , and lrO2, respectively (Figure 7(d)). A comparison clearly shows that at similar mass loading, HS delivers 177.3 amperes more than lrO2 at a certain potential, highlighting its exceptional performance. It is also worth mentioning that the HS is not only better in terms of performance than lrC>2 but also cost- effective, as an economic analysis showed using MOF@WB-O.21 as a catalyst is -65% cheaper.

[0268]

[0244] The lowest Tafel value of 41.68 mVdec-1for MOF@WB-O.21 among all the tested samples showed the best kinetics and highest energy-to-fuel conversion efficiency (Figure 7(e)). This is further evident from a lower charge transfer resistance of HS compared to the WB as recorded by EIS. Furthermore, the MOF also makes the WB in the HS highly hydrophilic, increasing its wettability (inset of Figure 7(c)). Thus, high wettability and low charge transfer resistance results in better kinetics by the quick adsorption of ions and their conversion to molecular species. To assess the access to active sites, the electrochemical surface area (ECSA) of the Co-MOF, WBs, and HSs was measured by recording double-layer capacitance (Cdi) as shown in Figure 7(f). It is found that the initial increase in B concentration (up to -0.43) brought higher active surface (Cdi= 23.7 mFc r2), which decreases (Cdi= 20.9 mFcmr2) with further increase in the B concentration (WB-0.65). Upon building the HSs, the Cdi was quickly increased to 52.7 mFcm-2for MOF@WB-O.21 , which is more than twice that of WB, confirming that incorporation of the MOF provides faster access to active sites by exposing them to electrolytes, taking advantage of the MOF’s porous structure and extended ionic connection among the two dissimilar sheets.

[0269]

[0245] Stable long-term operation is a key hurdle in realizing seawater, especially considering the nature of renewable power supply systems to ensure green energy. The MOF@WB-O.21 HS showed exceptional stability over 1000 h continuous operation without any current losses (Figure 7(g)). In contrast, WB-0.21 , Co-MOF, MOF@WB-0.43, and lrO2 lost more than 50% of their initial current density after only 50, 100, 100, and 10 h respectively, showing the importance of optimized interfacial connection. It is worth noting that WB and Co-MOF are not individually stable; therefore, it is unique ionic bonds that stabilize the HS. Therefore, even though the Cdi of the HS is less than that of the Co- MOF, the ionic bonds formed during the creation of the HS are vital to its stable operation in seawater. Furthermore, a higher Cdi and poor stability of MOF@WB-0.43 demonstrate that controlled B concentration impacts the ionic bonds between the Co-MOF and WB and obtaining optimized B2O3 at the interface.

[0270]

[0246] To further ascertain the outstanding stability of the catalyst, the electrode was left idle in the electrochemical setup for 5 days after testing it for 1100 h, and it was then retested. It demonstrated phenomenal stability by retaining more than 90% of its initial current density after 3 days of operation (Figure 7(g) inset). MOF@WB-O.21 was further tested at a high current density of >250 mAcm-2and remained stable for 100 h, demonstrating extraordinary stability and performance. No doubt, the HS has displayed exceptional stability; however, a gas chromatograph mass spectroscopic (GC-MS) analysis was conducted to ensure it prevents CER. No associated peaks were identified in the GC- MS spectrum, confirming that chlorine was not generated during the reaction.

[0271]

[0247] DFT calculations were employed to understand the role of the ionically bonded heterointerface in safeguarding the catalyst against chlorine chemistry and to ascertain its selectivity for hydroxyl groups over chloride anions. To determine the potential of the unique heterointerface to mitigate competition between the CER and OER, a Cl- and an OH- ions were positioned in different locations on the CO-MGF / W03(001 ) and Co-MOF / B2O3(001 ) surfaces. To balance the charge, a cation (Na+) was introduced into the supercell and placed at a sufficient distance from both the surfaces and the anions to ensure no interaction occurred. Figure 7(h)-(k) display the relaxed structures of the Co- MOF / WO3(001 ) and Co-MOF / B2O3(001 ) with the adsorbed OH- or Cl- ions. The calculated binding energies and bond lengths show for both systems, the OH- and Ch bonds to the Co atom, with the Co- OH bonds formed being 1.99 and 1.87 A, while the Co-CI bonds are 2.31 and 2.30 A for the Co- MOF / WC>3(001 ) and Co-MOF / B2C>3(001 ) surfaces, respectively. Importantly, the binding of OH- was stronger than Ch on both HSs. Hence, this confirms that the creation of the ionically bonded HS can not only safeguard the catalyst against chlorine chemistry but also improve the adsorption of OH- to boost the catalytic activity.

[0272] Structural, morphological, and electronic stability

[0273]

[0248] XRD demonstrated similar peaks were present as for the original HS before and after testing, assuring structural stability of the HS in the harsh seawater. Moreover, EDS maps confirmed no deposition of any foreign material from the seawater during the electrochemical process, indicating that the surface has high selectivity towards OER. The TEM image in Figure 7(a) clearly shows the WB sheets sandwiched by Co-MOF, and the monoclinic microstructure is well preserved, as observed by HRTEM (Figure 7(b)) and SAED (Figure 7(a) inset). The unchanged and preserved electronic states of W prove the operational stability of the HS for long operations (Figure 7(c)), deconvoluted XPS spectrum of W 4f). The peaks at 35.13 and 37.20 eV attributed to the doublet W 4f of the WO3 (W6+), indicate an unaffected oxidation state of W. The retention of a higher oxidation state favors the establishment of a more active oxyhydroxide phase during OER, supporting unchanged activity for an extended duration. Moreover, the deconvoluted 01s spectra (Figure 7(d)) peak at 531 .12 eV confirms retention of the metal-to-oxygen bond, assuring the heterointerface prevails in the harsh environment. Similarly, the unaffected Co oxidation states (Figure 7(e)) and slight variation in N 1s states (Figure 7(f)) ensure not only is WB stable, but the ionic bond also maintains the MOF structure over the continuous selective reaction.

[0274]

[0249] The post-stability O K-edge NEXAFS spectra of MOF / WB-O.21 (Figure 7(g)) show a shift in the peak position towards higher energy with the increased intensity in the first peak arising from the surface oxidation of the catalyst and formation of oxy-hydroxide phase. Co L-edge (Figure 7(h)) shows a similar behaviour with a positive shift in the photon energy, while the N K-edge (Figure 7(i)) shows a negligible shift in energy with small intensity peaks, apart from the ones at 402 and 405 eV, corresponding to the pi* and sigma* state for C-N from the imidazole ring, which might be due to the affected C-N and Co-N bond during formation of the oxy-hydroxide phase. No notable peaks for B K- edge NEXAFS spectra were obtained, similar to that of the original MOF@WB-O.21 HS confirming Co-MOF is well maintained over the WB sheets, making it challenging to detect B. ICP-MS verified that B (60.029 ppb) is present in the catalyst. Furthermore, no B2O3 peaks were observed, which might be due to hydrolysis of the B-O-B bond to generate and modulate B-OH species at the interface, contributing to improved catalytic activity.

[0275] Discussion - Example 2

[0276]

[0250] As predicted by theoretical calculations, B relaxes away from the uppermost surface layers, exposing the top W-0 surface as the best possible site for the adsorption of Co atoms within the MOF to construct the 2D morphology and interlinked HS. The DFT studies and experimental observation confirm that both materials interact through the formation of strong ionic bonds, which unite the MOF and substrate as single new material rather than simple composite or hybrid structure. This unique construction creates new material properties that provide an exceptional surface and interface to tackle problematic chlorine chemistry while activating the catalyst for extraordinary performance. The catalyst’s preference to bond with hydroxyl ions over chloride not only prevails at the surface but also at the interface where B2O3 acts as a B-OH modulator. To obtain such an interface, it was observed that the doping concentration of B may also play a role, with the doping position adjacent to W just beneath the top layer suggested to be preferred, as with either a low concentration (not enough B to modulate W-0 bond) or high concentration (self-competition to locate right place) such properties are not achieved. However, further probing of such doping to modulate and realize ionically bonded HSs with optimized properties needs to be studied so that this unique chemistry can be explored for other applications than OER. Ionic bonds are the most active and stable bonds where high precision conversion reactions are required, e.g., catalysts for water splitting for stable long operation. This understanding is confirmed as the HS operates stably for more than 1100 h in harsh seawater and maintained stable operation even after sitting idle in the electrochemical setup for 5 days.

[0277] Example 3: vertically layered B-Fe-doped MXene CFBT')@CoMOF HS

[0278]

[0251] A 2D-2D ionic bonded HS material with an ultrathin twin (Fe and B) doped engineered TisC2Tx- MXene (‘FBT’) component and an ultrathin CoMOF (ZIF-67) component at room temperature (RT) was prepared using the facile interface-assisted synthesis technique of the invention to form the unique 2D-2D HS material FBT / CoMOF (MOF@FBT). The B- and Fe- modified Ti3C2Tx-MXene was engineered by substituting regular Ti3C2Tx-MXene surface functionality with electrochemically active Fe and partially replacing Ti with B. The FBT material was then used in the HS fabrication process of the invention described herein, resulting in interfacial growth of ultrathin ZIF-67 MOF on FBT, where the process generated ionic bonding at the heterointerface between the dissimilar materials through Fe-O-Co and N-Fe formation. The presence of vacancies and surface termination resulting from the doping in the modified MXene, and direct growth of the MOF on the modified MXene using the process of the invention creates strong interactions including ionic bonding at heterointerface of the 2D / 2D HS, which advantageously prevents leaching of catalytically active sites, hence ensuring long-term structural stability of the material, and providing for enhanced OER activity in application described herein. The benefits of both dissimilar materials work in tandem to enhance OER activity. In an electrochemical process, the presence of ionic bond at the heterointerface helps modulate the electronic charge redistribution and facilitates rapid charge transfer for improved activity. The presence of ionic bonding at the heterointerface, which was confirmed by XPS and NEXAF analysis, also ensure rapid electrical conductivity for fast redox process during OER in both alkaline seawater and deionised (DI) water. OER evaluation shows that both FBT and ZIF-67 / FBT HS required low overpotentials of 521 and 508 mV respectively to achieve 1 A cm-2in seawater. It is believed that both FBT and ZIF-67 contributed OER activities to the overall performance of the catalyst. However, only ZIF-67 / FBT was able demonstrate stability for 500 h, and this was attributed to material stability induced by the presence of ionic bonding at the heterointerface.

[0252] Providing B underneath the top layer generated much-needed vacancies that helped in the better adsorption of the ZIF-67 MOF to create strong ionic bonds with surface Fe through generation of Fe-O-Co bonds. The ionically bonded heterointerface in ZIF-67 / FBT tuned the active sites for enhanced OER, prevented the leaching of active iron species, and increased the selectivity of OER over CER. The presence of B-0 bond improves reaction kinetics by modulating the adsorption of OH- ions at the interface. As a result, both FBT and ZIF-67 / FBT required very low overpotentials of 521 .2 and 508 mV respectively to deliver 1 A cm-2, however, only ZIF-67 / FBT was stable (>500 h) in seawater, while FBT loss their activity in few hours. Here the invention harnesses the hidden potential of MXenes, changing them from simple conductive substrates to a high performance OER catalysts for alkaline seawater splitting.

[0279] MXene (Ti3C2Tx) synthesis

[0280]

[0253] 2 g of TisAIC2 (Sigma) powder was gradually added to 40 mL of 10% HF solution (Sigma) and the mixture was stirred at 500 rpm and 35 °C for 24 h. The resultant mixture divided into three 50 mL centrifuge tubes that already contains about 25 mL of milli-Q water. The tubes were then sealed, placed inside a centrifuge and centrifuged at 5000 rpm for 5 min. The supernatant was discarded and milli-Q water was added to the sediment for more washing and the process was repeated until a pH of ~7 was obtained. To ensure all the F- was removed, the neutral sediment (wet TisC2Tx) was dispersed in 1 L of milli-Q water and the solid recovered via vacuum filtration and using 0.2-micron PES filter membrane. After filtration, the product was vacuum dried at 50 °C overnight.

[0281] Ti3C2Txsurface modification using 0.1 M NaOH

[0282]

[0254] 2g of dried Ti3C2Txwas treated with 20 mL of 0.1 M NaOH (Sigma) solution with stirring at room temperature for 24 hours. The solid product was recovered by centrifuging and followed with multiple washing with DI water until a pH of ~7 was obtained. The neutral solid was subsequently redispersed in DI water and sonicated for 15 min. The solid product was recovered by vacuum filtration, dried overnight under vacuum at 50 °C and the final TisC2O2 was labelled T.

[0283] Twin Doping of TisC2O2 (T) with Iron and Boron

[0284]

[0255] Calculated amount of T, Fe(NO3)3-9H2O (Sigma) and B(OH)3 (Sigma) (mass ratio: T: Fe(NO3)3-9H2O = 1 :1 .5 and Fe(NO3)3-9H2O: B(OH)s= 3:10) were dispersed in DI water (for T = 50 mg, DI water = 20 mL) and stirred for 90 min at room temperature (RT). After stirring, the water was evaporated until dryness, and the resultant power collected and annealed. The annealing was done at 600 °C for 5 hours, at a heating rate of 5 °C / min in a tubular furnace under Ar flow. Upon cooling to room temperature, the product was collected, washed with warm water (60 - 70 °C) followed by absolute ethanol and dried at overnight under vacuum at 50 °C. The resultant ‘FBT’ product was collected.

[0285] ZIF-67 / FBT synthesis

[0286]

[0256] The ZIF-67 / FBT by directly growing ZIF-67 on FBT. Briefly, 30 mg of FBT was dispersed in 45 mL solvent (methanol (30 mL) + DI water (15 mL)), sonicated for 60 min and titled Sample A. Sample B was formed by dissolving 45.46 mg of 2-methylimidazole (Sigma) in 22 mL DI water while Solution C was formed by dissolving 159.86 mg of Co(N03)2.6H2Q (Sigma) in 22 mL of DI water. Upon completing the sonication of Solution A, Solution B was gradually added to Solution A gradually, followed by the gradual addition of Solution C, and was allowed to remain in a steady state for 2 hours without any disturbance. Afterwards, solid product was recovered by centrifugation and followed by the washing with methanol, DI water and methanol. The resultant wet residue was dried at 50 °C under vacuum for 1 hour and the product labelled ‘ZIF-67 / FBT’.

[0287] Electrochemical Testing

[0288]

[0257] The electrochemical testing was conducted on an electrochemical workstation (CHI760D) in a standard 3-electrode system at room temperature (RT) using 6 M KOH in seawater as electrolyte. Graphite rod and Hg / HgO were employed as the counter and reference electrodes respectively. The working electrodes were prepared by drop-casting of prepared electrode ink on washed nickel foam (washed with 1 M HCI, DI water and ethanol). To prepare the electrodes ink, firstly, carbon black (CB) solution was prepared by dispersing 20 mg of carbon powder in 20 mL solvent (Isopropyl alcohol: water = 1 :4) and sonicated for 1 hour. Next, 4 mg of FBT or ZIF-67 / FBT, 1 mL of CB solution and 8 pL of PTFE (Sigma) binder was mixed and sonicated for about 30 - 40 min for homogeneous ink. The WEs were then prepared by drop-casting 200 pL of the prepared ink on an area of 0.25 cm2on the washed and dried nickel form, then dried in a vacuum over at 50 °C for 24 hours. All the LSV were performed at a scan rate of 5 mV s-1without iR compensation using back-sweep method. Also, the potentials were converted to RHE using E(V + RHE) = E(Hg / HgO) + (0.059*pH) + 0.098.

[0289] Results & Discussion for Example 3

[0290]

[0258] Tuning I engineering Ti3C2C>2-MXene such that B replaces some of the Ti, while Fe replaces some of the surface oxygen can effectively activates dormant OER potential. Ti3C2C>2-MXene was first synthesised and then treated with NaOH solution to generate O-terminations which aided simultaneously Fe and B introduction into the matrix of Ti3C2C>2-MXene to result in the new FBT material.

[0291]

[0259] FBT was then used in the process of the invention which involves interfacial growth of the CoMOF (ZIF-67) on the FBT to form a 2D / 2D ZIF-67 / FBT HS having strong bonding, including ionic bonds, at the heterointerface of the dissimilar material as illustrated in Figure 9(a). TEM image revealed that the precursor Ti3C2C>2-MXene possesses the characteristic multilayered nanosheets morphology associated with MXene, while XRD confirmed the successful synthesis of MXene through presence of characteristic MXene XRD peaks (Figure 9(b) - bottom is TiC2C>2, middle is FBT and top is 2D / 2D HS).

[0292]

[0260] During formation of FBT, the O-terminations at the surface of TisC2O2 (after NaOH treatment) aided B and Fe adsorption to TisC2O2 during the initial deposition, with Fe3+interacting with the surface termination via electrostatic attractions due to its opposite charges with the surface. Thereafter, Fe and B are incorporated into the TisC2O2 upon thermal treatment, with B doping replacing some of Ti, while Fe replaced some of the surface termination. The doping breaks van der Waals interactions between adjacent MXene nanosheets, generating freestanding sheets of the new FBT material. The freestanding modified MXene nanosheets was confirmed by the disappearance of (002) peak in the XRD of FBT (Figure 9(b)). The retention of (006), (200) and (110) peaks located at 2© = 36.2°, 41 .7° and 61 ° respectively for FBT confirms the preservation of most of original TisC2O2 crystal structure upon Fe and B incorporation (Figure 9(b)). However, the disappearance of (002) peak indicates the disruption of ordered stacking of multilayered MXene induced by the presence of Fe at the surface. Further, whilst XRD could not independently confirm the presence of Fe and B, the evolution of new peaks at 2© = -32°, -54.6° and -56.8° could be attributed to their influence upon incorporation into the Ti3C2C>2-MXene structure. For 2D / 2D ZIF-67 / FBT HS, all its XRD peaks matches well with that of FBT, indicating low crystallinity of ZIF-67 and that the ZIF-67 / FBT HS retains the FBT structure.

[0293]

[0261] EDX mapping confirms the uniform distribution of Fe, B, Ti, C and O demonstrating successful incorporation of Fe and B, into the MXene matrix. TEM analysis revealed that FBT possess freestanding thin 2D nanosheets with centre of smaller hexagonal nanosheets at some of the edges (Figure 9(c)), and these shapes were confirmed through microstructural analysis of selected area electron diffraction (SAED) imaging (Figure 9(d) & (e)). HR-TEM imaging of FBT (Figure 9(f)) reveals the 2D nanosheets are arranged in definite order. The magnified image (insert) reveals a section of the atomic structure of FBT nanosheet consisting of Ti and C atoms and sectional distance of -0.92 nm. Furthermore, a defect is apparent in the area marked with golden colour (most right square in image) which was attributed to unoccupied vacancies left behind when B substituted some of the Ti in TisC2O2 matrix, and which may act as nucleation sites for growth of ZIF-67. SEM and TEM (Figure 9(g)) images of 2D / 2D ZIF-67 / FBT HS shows the presence of ultrathin 2D nanosheets of ZIF-67 on the surface of FBT, indicating the successful synthesis which was further confirmed by HRTEM imaging (Figure 9(h)). Lattice fringes were not detected for the ZIF-67, due to its low crystallinity, which agrees with the XRD result.

[0294]

[0262] A XPS survey scan also confirmed successful synthesis of Ti3C2©2-MXene via identification of Ti, C, and O in the X-ray photoelectron spectroscopy (XPS) survey scan and the presence of Ti-C bond in the high-resolution C 1 s spectrum. Upon doping with only B, there was decrease in the intensity of Ti 2p peak confirming the successful replacement of some of the Ti in Ti3C2©2-MXene, which is supported by the presence B-0 bond in the high-resolution B 1 s spectrum. Upon incorporation of only Fe, there was no significant observable decrease in the intensity of Ti 2p, confirming that Fe only interacted with the surface terminations of the Ti3C2©2.

[0295]

[0263] The presence of Ti, C, O, B and Fe in the survey scan of FBT confirm the successful incorporation of Fe and B into the Ti3C2©2-MXene matrix. High-resolution O 1 s spectrum of FBT (Figure 10(a)), reveals four binding energy peaks located at 529.75, 531.30, 532.15 and 533.42 eV corresponding to TiC>2 (lattice oxygen), (Ti or B)-O-Fe (defective oxygen vacancy site due to low oxygen coordination), C-Ti-(OH)Xand C-0 respectively were identified. Compared to Ti3C2©2, there was a decrease in the intensity of Ti©2 peak in FBT, indicating its conversion to (Ti or B)-O-Fe bonds. According to the principle of metal oxide defect chemistry, the presence of dopants often resulted into different lattice defects, causing oxygen vacancies and this depends on oxidation state. For example, if a cation of lesser oxidation state replaces Ti4+in its oxide, oxygen vacancies are created. Therefore, when B with a maximum oxidation state of +3 replaces some of the Ti4+(from TiC>2 peak), oxygen vacancies are created, explaining the decrease in Ti©2 peak intensity and the increase of (Ti or B)-O- Fe peak intensity. Furthermore, the peak area ratio of defective oxygen to lattice oxygen for TisC2O2, and FBT was calculated to be 0.33 and 1 .78 respectively, thus, confirming more oxygen vacancies in FBT which could subsequently act as nucleation sites for the growth of ZIF-67.

[0296]

[0264] In the high-resolution O 1s spectrum (Figure 10(a)), the O-M peak located at 531 .05 eV for ZIF-67 / FBT is negatively shifted by 0.25 eV compared to FBT, confirming electron transfer during Co- O bond formation in the HS. Further, there was an increase in the intensity of O-metal peak, caused by the formation of Co-0 bond and indicating the exposure of more metal sites for OER activity. Moreover, the peak at 399.88 eV in the high-resolution N 1 s spectrum (Figure 10(b)) was due to coordination between nitrogen from 2-methylimidazole and metal centres (N-Co for ZIF-67 and N-Fe for ionic bond formation in HS). Co2+-metal centres are understood to often coordinated with four N- atoms in ZIF-67, but both Co2+(at 781.70 eV) and Co3+(780.58 eV) peaks are found in the high- resolution Co 2p spectrum of ZIF-67 / FBT (Figure 10(c)). The Co3+peak is due to the formation of new Co-O-Fe bond caused by ionic interactions between ZIF-67 and FBT during HS formation, aided by the vacancies generated when B replaces some of the Ti in Ti3C2C>2-MXene. Moreover, from the high- resolution Fe 2p XPS spectra (Figure 10(d)), there was a significant decrease in the intensity of the Fe2+peak and a positive shift of 0.67 eV in the binding energy of Fe3+in ZIF-67 / FBT as compared to that of FBT. The intensity decrease in the Fe2+peak is due to Co-O-Fe bond formation, while the positive shift of binding energy position of Fe3+confirms electronic transfer between Fe, O and Co during the ionic bond formations.

[0297]

[0265] NEXAFS provided more insight into the successful growth of ZIF-67 on the FBT and the nature of interactions between the dissimilar materials. In the N K-edge spectrum (Figure 10(e)), both ZIF- 67 and ZIF-67 / FBT displayed similarities in peak positions, shape and number of peaks except for change in intensity caused by N-Fe formation on the HS. Peaks at <400 eV are attributed to the TT*- bonds that originated from the graphitic and pyridinic nitrogen from the 2-methylimidazole. Peaks located at -402.5 and -406 eV are attributed to the o*-bonds that originated from bonding between pyrrolic nitrogen and metal centres (N-Co / N-Fe). Comparing the o* and ir*-bonding region, there was a significant increase in the intensity of the o*-bond located at -406 eV for the ZIF-67 / FBT and this was due to the strong ionic bonding at the heterointerface of the HS. Similarly, the Co L-edge spectrum exhibited identical peaks features, though with some observable changes, for both ZIF-67 and ZIF- 67 / FBT, with electron transitions from 2p to 3d states resulting into peaks at L2 and L3 regions (Figure 10(f)). The disappearance of the shoulder edge indicated with arrow around 780 eV and change in the L2 shape is due to new bond formation (Co-O-Fe) upon hybridization of ZIF-67 with FBT which occurs as a result of the process of the invention. Moreover, the observed decrease in intensity of ZIF- 67 / FBT indicates higher Co 3d electron occupation caused by electron transfer from FBT to ZIF-67, and confirming strong electronic interactions. In addition, Fe L-edge spectrum (Figure 10(g)) of FBT and ZIF-67 / FBT displayed similar features, except little noticeable differences. The L3 and L2-edge peak regions are due to electronic transitions from Fe 2p core level to empty 3d level that is highly hybridized with oxygen 2p. Noticeably, the Fe3+concentration increases at the expense of Fe2+, which was triggered by the strong electronic transfer between FBT and ZIF-67 during Fe-O-Co formation in the HS and in agreement with the XPS results.

[0298]

[0266] The 2D / 2D HS material OER activity in alkaline seawater was compared to a commercial RuC>2 catalyst. Both FBT and ZIF-67 / FBT delivered ampere level current density performances at lower overpotentials, and outperformed the commercial RuC>2 catalyst (Figure 11(a) & (b)). Remarkably, FBT achieved the current density of 1 A cm-2at low overpotentials of 521 .2 mV in alkaline seawater. This performance was attributed to the redox potential of FBT that was triggered by the presence of Fe and B, and the 2D morphology that could enhance electrochemical active surface area (ECSA). Further, the ZIF-67 / FBT equally displayed exceptionally OER activities by achieving the current density of 1 A cm-2at low overpotentials of 508 mV in alkaline seawater. Moreover, the ZIF-67 / FBT exhibited the excellent OER kinetics by displaying the least Tafel slop values of 27.9 mV dec-1is alkaline seawater (Figure 11(c)). The superior activity of the ZIF-67 / FBT HS could be attributed to the synergistic effect of the combined starting materials and strong bonds including ionic bonding between the ZIF-67 and FBT that accelerates conductivity for rapid redox process for enhance OER.

[0299]

[0267] The influence of Fe and B doping was also investigated by measuring the OER activities of Ti3C2O2-MXene, BT (no Fe doping) and FT (no B doping), and comparing the results to that of FBT and ZIF-67 / FBT in alkaline DI water electrolyte. The comparison makes it is apparent that Fe active sites are mainly responsible for improved OER activities. Conversely, B generated the much-needed vacancies that aided ZIF-67 growth without destroying the 2D morphology, thus exposing more active sites for OER activities in FBT. Also, the improved performance of FBT over FT could be attributed to its impressive large 2D morphology (Figure 9(b)) as compared to that of FT. Further, from the Nyquist plots (Figure 11(d)), FBT and ZIF-67 / FBT exhibited the smallest semicircle diameter than TisC2O2, BT and FT, indicating the least charge transfer resistance at the electrocatalyst-electrolyte interface, thus improved charge transfer capability and OER activities. Moreover, it was observed from Figure 11(e) that both FBT and ZIF-67 / FBT possess higher ECSA than FT, BT and Ti2CsO2, indicating the presence of more active sites for OER activity, hence improved performances.

[0300]

[0268] In seawater electrolysis, exclusive selectivity towards OER while avoiding CER is highly desirable for electrocatalysts for industrial application. To mitigate CER during seawater splitting, the electrocatalyst must be active enough to generate high current density at overpotential <480 mV (equilibrium potential of CER) in alkaline medium. As shown in Figure 11(f), both FBT and ZIF-67 / FBT could potentially mitigate CER due to their ability to generate high current densities at 480 mV. Moreover, we investigated the selectivity of FBT and ZIF-67 / FBT towards OER in seawater, by passing the gas generated during electrochemical testing through acetone to probe the evolution of chlorine gas though analysis for chloroacetone generation. Gas generated during 1 h of continuous electrochemical testing was passed through acetone, and only acetone could be detected on GC analysis (Figure 11(g) & (h))), indicating the absence of chlorine gas. This exclusive OER selectivity was attributed to the materials’ impressively high OER activities in seawater, hence making them a catalyst of choice for seawater splitting in the industry.

[0269] Results for chronoamperometry for FBT and ZIF-67 / FBT in seawater in probing stability (Figure 12(a)) show after a period of 500 h, ZIF-67 / FBT exhibited only a current drop of 14%, which was far better than FBT that already exhibited 50% drop in current after just 100 h, and is likely due to the high stability of 2D / 2D ZIF-67 / FBT HS. TEM analysis of the spent electrodes after stability test shows that both FBT and ZIF-67 / FBT retained its 2D morphology (Figure 12(b) & (c)). XPS and NEXAF analysis (Figure 12(d) & (e)) reveals depletion of Fe from FBT via dissolution during stability test, while ZIF-67 / FBT still retain its Fe component. ZIF-67 / FBT’s ability to retain its Fe active component is based on the presence of strong ionic bonding at the heterointerface of the HS. This observation shows that, although the OER activity of Ti3C2C>2-MXene has been significantly enhanced to the point of outperforming commercial RuC>2 catalysts, however, there is still a need for improved active site protection, which is facilitate by the HS material structure. Conversely, the ionic bonds between FBT and ZIF-67 in the 2D / 2D HS prevent Fe dissolution pointing to good stability.

[0301] Conclusion - Example 3

[0302]

[0270] OER activity of bare Ti3C2O2-MXene both in alkaline seawater and DI water can be enhanced using the doping method of the invention. The post synthesis analysis of the dopant modified MXene, FBT, reveals that B substituted some of the Ti in TisC2O2 forming B-0 bond underneath the outer layer, while Fe species substituted the surface termination. B underneath the top layer generated the much-needed oxygen vacancies that helped in better adsorption of ZIF-67 MOF to create strong ionic bonds with surface Fe through Fe-O-Co. The ionically bonded heterointerface in ZIF-67 / FBT tuned the active sites for enhanced OER, prevented the leaching of active iron species, and increased the selectivity of OER over CER. Remarkably, both FBT and ZIF-67 / FBT required very low overpotentials (521.2 and 508 mV respectively) to deliver 1 A cm-2, however, only ZIF-67 / FBT was stable (>500 h) in alkaline seawater. The strong ionic bonding at the heterointerface of ZIF-67 / FBT HS protected the dissolution of Fe, which was identified as the main active site, hence ensuring long-term stability. The major observation here is that, unlike other MXene-based HS OER catalysts where the MXene only acted as support that supplies electrical conductivity, the new FBT material contribute to both OER activity and conductivity. Hence, this work demonstrated a new approach through which MXene could be modified for enhanced applications, such as OER activities both in alkaline seawater and DI water.

[0303] Example 4: vertically layered B-FeS2 / NiMOF HS

[0304]

[0271] The inventors designed a unique interface by heterostructuring boron-doped iron disulfide (B- FeS2) with NiMOF sheets (MOF@BFS). The boron-doping leads to high electrical conductivity, assist in producing ionic bond among dissimilar sheets, and yield borate species, a weak Lewis acid, at the interface which act as local OH- modulator. Further, the hydrophilic nature of MOF sheets enables quick wetting and favours OH- adsorption and shielding from Cl- ions. As a result, this HS only requires an overpotential of 450 mV to derive a current density of 1 A cm-2as an OER catalyst in seawater far superior to commercially available lrO2. Moreover, the system is stable over 200 h, presenting an ideal pathway to realize H2 via materials engineering directly from seawater.

[0305]

[0272] Interlayer interactions between B-FeS2 and Ni-MOF were created to develop HS. This is achieved by introducing boron doping on the 2D FeS2, which enhances the adsorption of Ni-MOF onto the surface of B-FeS2 and help to create strong ionic connections between the interfaces. The appearance of B-OH sites at the interface introduced a Lewis acid layer of protection on FeS2. Also, the presence of the borate anion in the electrocatalyst, the unsatisfied S atoms in B-FeS2, reinforce the system’s defence mechanisms by optimizing the adsorption of MOF in HS design. Moreover, the integration of Ni-MOF efficiently shielded the B-FeS2 catalyst layer, preventing potential interactions with chloride anions on the surface to avoid chloride oxidation chemistry.

[0306] Synthesis of B-FeS2 precursor template

[0307]

[0273] Initially, the precursor solution was prepared by stirring a mixture of 300 mg of Fe(NC>3)3-9H2O and 20 mg of boric acid in 10 mL of ethanol absolute for 20 min to ensure homogeneity. Subsequently, the prepared solutions were slowly added and mixed with 31 g of NaCI in 500 mL glass beakers. This mixture was heated on a hotplate at 80°C for 30 min to evaporate the solvent and enhance thermodynamic mixing. NaCI was employed as the salt-based template to facilitate the growth process. After the mixing process, the material, were annealed at 400 °C for 2 hours at ramp rate of 2 °C min-1under nitrogen to obtain B-Fe2Os after filtration and later obtained powder of B-Fe2Os converted into B-FeS2.

[0308] Synthesis of B-FeS2 / MOF (MOF@BFS)

[0309]

[0274] 15 mg of B-FeS2 material was dispersed in a mixture of methanol (29.5 mL) and DI water (14.5 mL) (2:1 ) in a 300 mL beaker, followed by sonication for one hour to obtain solution A. Ni(NC>3)3-6H2O (100 mg) and 2-methylimidazole (45.5 mg) electron donating organic ligand were dissolved in two separate beakers in 22 mL of deionized water (DI). Then, linker solution was added to solution A followed by nickel salt solution. Then, the final obtained solution was centrifuged at 4400 r.p.m for 10 min and washed a few times with DI water and absolute ethanol. Finally, the obtained powder of HS was dried in an oven at 50 °C for 24 h.

[0310] Results and Discussion

[0311]

[0275] A room temperature (RT) salt template method was opted to develop the sacrificial B-doped sheets of iron oxyhydroxide, which was annealed prior to sulfurization to obtain the final B-FeS2 sheets.

[0312]

[0276] To develop the strong ionic connections, the surface energy of the Fe2Os is manipulated by doping with Bto obtain B-Fe2Os which is then converted into the active B-FeS2thin sheets. The sheets of B-FeS2 act as a platform to grow the ionically linked NiMOF via the interfacial growth strategy of the invention at room temperature (RT) as illustrated in Figure 13(a). B-doping develop strong connections between the two dissimilar materials, B-FeS2 and MOF by modulating the electronic properties of the inorganic template. These modulations generate borate species that preferentially adsorb hydroxyl ions at the interface making it a highly active electrocatalyst. TEM imaging shows the method produces micron scale ultrathin sheets of B-Fe2Os (Figure 13(b)) which retains its features even after sulfurization. However, thermal conversion reaction yielded porous features in the sheets which are clearly visualised by TEM image in Figure 13(c). The HS growth as shown in Figure 13(d) depicts the successful growth of MOFs sheet covering the B-FeS2 sheets to form B-FeS2 / MOF HS. This unique ionically bonded HS benefits from the strong heterointerface interaction along with thin and layered structure of the B-FeS2 nanosheets and interconnections between metal ions and ligands within the MOF. The stacking of numerous uniform layers upon each other creates a HS of B- FeS2 / MOF as further confirmed by SEM imaging (Figure 13(e)).

[0313]

[0277] The lattice spacing of 0.27 and 0.24 nm observed via the high-resolution TEM (HRTEM) images of both B-FeS2 (Figure 13(f)) and B-FeS2 / MOF (Figure 13(g)) corresponding to (200) and (111 ) plane of pyrite FeS2 according to JCPDS No.42-1340, confirming its crystalline nature. However, the SAED pattern (inset of Figure 13(c), (d)) shows polycrystalline growth obvious from highly porous sheet structure. In addition, EDS elemental mappings of B-FeS2 / MOF shows the homogeneous distribution of all the elements throughout the HS (Figure 13(h)).

[0314]

[0278] XRD was used to analyse the crystal structures of FeS2, B-FeS2, B-FeS2 / MOF and Ni-MOF (Figure 14(a)). The FeS2 and B-FeS2 XRD peaks are well matched with the JCPDS 42-1340 showing the successful formation of FeS2. In case of the HS, in addition to the main FeS2 XRD peaks, additional peaks were observed, indicating the presence of Ni-MOF (JCPDS No:00-014-0117). These catalysts have been characterized by utilizing NEXAFS analysis (Figure 14(b)-(d)). This method enabled the precise identification of the chemical states of Fe, Ni and S and reveal that the strong ionic connections are developed between the component at the heterointerface. The co-existence of elements from B- FeS2 and MOF in the HS (Figure 14(e)-(h)) was revealed by XPS, wherein the chemical states of different elements in samples were also investigated. The XPS data confirms the presence of the constituent’s elements in the HS and the pristine materials. The B is successfully doped in to the FeS2 sheets and further strongly connected heterointerface is developed between the B-FeS2 and MOF. The shifted peak of B in B-FeS2 / MOF HS demonstrates the interactions which effectively improved the OER activity.

[0315]

[0279] To confirm the electrocatalytic performance of the strongly bonded HS catalyst, electrocatalytic LSV performance of the prepared catalysts and commercial lrO2 / NF was considered using a three- electrode configuration in a 6 M KOH DI water and seawater (Figure 15(a), (b)). The OER polarization curves of the catalysts were measured with varying dopant concentrations of boron on FeS2 base (Figure 15(a)). Under optimized conditions, the FeS2 with a 1 .2% of B dopant ratio insertion exhibited the lowest overpotential of only 534 mV to achieve current density of 1 A cm-2in B-FeS2 / MOF (Figure 15(a)). Meanwhile, the B values in B-FeS2 (0%, 0.2%, 0.6% and 2.4%) showed the higher overpotentials as shown in Figure 15(a). Additionally, Figure 15(b) shows that the B-FeS2 / MOF HS reached a remarkable 1 A cm-2current density at overpotential of 489 mV in seawater, surpassing the lrC>2 benchmark of 707 mV in the same condition. Although, OER performance was slightly reduced in alkaline seawater electrolyte compared to DI water, it remained very impressive.

[0316]

[0280] Low concentration of B atoms within the B-FeS2 / MOF interface can facilitate reduction of Fe2+by acting as weak Lewis acid site, which interacts with Lewis basic anions (SO42created in the electrolyte. Also, the improvement can be attributed to the borate species with B-0 bonds in the B- FeS2 / MOF catalyst (inset Figure 15(a)) which act as promoters for catalytic activity by facilitating of electron transfer kinetics on the B-FeS2 / MOF surface, without changing the pH. Therefore, the B- FeS2 / MOF HS led to several synergistic effects, including increased surface area, enhanced surface adsorption capability, and decreased charge transfer resistance, contributing to its improved electrochemical catalytic performance. In Figure 15(c), the OER reaction kinetics were evaluated using Tafel plots. The smallest Tafel slope 36 mVdec-1of the B-FeS2 / MOF shows efficient electrocatalytic kinetics and a rapid charge transfer in alkaline seawater. This results in a swift increase in current density with lower overpotential. Moreover, the generated gases were passed through acetone for 1 h and resulted product was analysed GC-MS which confirmed there was no peak observed for chloroacetone (a product of acetone + chlorine) confirming inhibition of CER and selectivity of catalyst to drive only the OER at anode.

[0317]

[0281] Long-term stability of B-FeS2 / MOF catalyst in seawater for a duration of 500 h is demonstrated, Figure 15(d) &(e). The B-FeS2 / MOF HS showed exceptional performance, sustaining a remarkable and stable current density of 1 .5 A cm-2with a slight decrease of 5% during the initial stage due to bubble adsorption, which hindering certain active sites. Therefore, no significant degradation in current density of B-FeS2 / MOF HS was observed. The B-FeS2 / MOF HS sample exhibiting the similar LSV curves even after 500 hours (Figure 15(e)). In addition, the surface of B-FeS2 / MOF was passivated by negatively charged polyanions. This was facilitated by the co-existence of SO42" and carbonate ions from alkaline solution at the anode interface of B-FeS2 / MOF. This interaction effectively neutralized the impact of Cl- ions in the seawater electrolyte, leading to an enhanced passivation effect, ultimately mitigating the corrosion process. However, the current density of FeS2 / MOF, B-FeS2, and pure Ni-MOF dropped by more than 30%, 45% and 41 % after 50 h, 70 h, and 15 h, respectively.

[0318]

[0282] EIS evaluated the charge transfer kinetics of pyrite FeS2, B-FeS2, and B-FeS2 / MOF catalysts. Figure 15(f), the Nyquist plots reveal diameter of the semicircle and an inclined line in low frequency regions. The B-FeS2 / MOF HS exhibits the smallest semicircle compared to FeS2 and B-FeS2, corresponding to lower contact and charge transfer impedance at the electrode / electrolyte interface. This confirms the higher charge transfer kinetics in the HS than in FeS2 and B-FeS2.

[0319] Conclusion - Example 4

[0320]

[0283] B-FeS2 / MOF electrocatalysts exhibits optimised OER performance in seawater. The synthesis method leveraged the interplay of Fe, S, and B, resulting in a network of B-FeS2-doped 2D sheet that effectively hindered agglomeration and re-stacking. Moreover, the introduction of a MOF enhanced durability and stability of B-FeS2 in alkaline electrolyte by producing strongly connected interfaces. The HS components are firmly attached and provide higher electrocatalytic activities with stable performance in the corrosive seawater electrolyte. Specifically, Ni-MOF along with B-FeS2 created an ionic bond which enhanced electrical conductivity and stability of system. The remarkable electrochemical performance of B-FeS2 / MOF catalysts can significantly improve charge transfer efficiency within the B-FeS2 / MOF structure. Furthermore, the MOF architecture proved effective in protecting active sites against chloride corrosion during electrochemical processes in alkaline media. Additionally, the incorporation of HS nanosheets optimized the surface electronic structure and adsorption energy, thereby enhancing the electrocatalytic activity. Overall, this study demonstrates a promising pathway for developing efficient and stable catalysts for seawater electrolysis.

[0321] Example 5 - vertically layered / FeMOF HS

[0284] FeMOF / PW8O26-B2C>3 (MOF@WPB) HS with ionic bonds between Fe-MOF and WPB have been synthesized via a facile solid-liquid interfacial growth process of the invention. The ionic bonds provide strong bonding and interactions that result in enhanced stability such that MOF@WPB operates stably for over 1000 h without any notable current losses while the individual materials were not stable in harsh seawater. The ionic bond formation is assisted by boron while B2O3 helps in modulating the hydroxyl ions at the interface via the hydrolysis of B-0 bonds. Additionally, phosphate ions present at the interface also provides anti-corrosive properties by repelling the chloride ions.

[0322]

[0285] Fe-MOF / PW8O26-B2C>3 (MOF@WPB) with ionic bonds at the heterointerface caters for increased transfer of electrons, better charge transfer, and enhanced stability of the material. B2O3 engineered HS helps in modulating the B-OH bonds at the interface via the B-O-B bond hydrolysis leading to the generation of OH- which along with the phosphate anions helps in repelling Cl’, thus preventing chlorine evolution reaction (CER). As a result, MOF@WPB achieves an industrial scale current density of 1.75 A cm-2at 2 V without CER. The material was also stable for over 1000 h of continuous operation demonstrating that ionic bonds at the interface can be critical in improving the performance as well as stability of the electrocatalysts in harsh seawater environment.

[0323] Synthesis of tungsten oxide phosphate borate (WPB)

[0324]

[0286] PW8O26-B2O3 was synthesized by a facile salt template method using KCI salt. Different concentrations of boric acid (30, 65, 130, and 195 mg) were added to 12 mL DI water and sonicated for 2 mins to acquire samples with different B concentrations. Simultaneously, 300 mg ammonium metatungstate was added in 6 mL ethanol which was then mixed with the boric acid solution followed by sonication for 2 mins to obtain proper dispersion. Separately, 30 g of KCI salt was heated on a hot plate for 5 mins at a temperature of 80 °C and then add the mixture of ammonium metatungstate and boric acid was added dropwise. The template was stirred continuously until it had dried completely. The salt with precursor was then left for 2 days for the sheets to grow. The template was then annealed at a temperature of 650 °C for 3 h at a ramping time of 3 h. The annealed material was filtered and washed with ethanol and left for overnight drying in a vacuum oven at a temperature of 50 °C. The as- synthesized samples were then named as WB-0.1 - 2.3 w / w%; WB-0.21 - 5.09 w / w%; WB-0.43- 10.18 w / w%; WB-0.65- 15.28 w / w% (range 0 -16 w / w%) and were then phosphorized in a tube furnace at a temperature of 350 °C for 3 h at a ramping rate of 3 °C and then named as WPB.

[0325] Synthesis of Fe-MOF@tungsten oxide phosphate borate (WPB)

[0326]

[0287] WPB was used as a template for the growth of Fe-MOF sheets at room temperature (RT). 30 mg of WPB was mixed with a 44 mL mixture of methanol and water (2:1 ) and sonicated for 1 h. Simultaneously, a mixture of 159.86 mg of Fe(NO3)3-9H2O and 45.46 mg of 2-methyl imidazole organic ligand was prepared in 22 mL DI water. The organic ligand was added to the sonicated WPB solvent followed by the addition of iron nitrate solution and left undisturbed for growth. After 3 hours, the obtained mixture was washed first with methanol, then with water, and finally with methanol. The Fe- MOF I WPB obtained was dried for 4 hours in a vacuum oven at 50 °C. Pure Fe-MOF was grown by following the same procedure without the addition of WPB.

[0327] Results and Discussion - Example 5

[0288] To create ionic bonds, the structure of the base material i.e., PWsC^e-E^Os (WPB) is first analyzed by using powder XRD showing unique arrangements of different components such as PWsO26 and B2O3 as shown in Figure 16(a) matching well with JCPDS No. 50-0660 demonstrating a monoclinic structure while the peaks around 23.4, 27.8, 36, 48.4, 57.6, and 66.2 confirms the presence of B2O3 (JCPDS No. 13-0570). The small shift in the peak position for WPB was observed due to the presence of borate species. The as-synthesized WPB was then used as a support for the construction of Fe-MOF, resulting in the formation of MOF@WPB HS, also matching with the XRD of WPB but with reduced peak intensity probably due to the amorphous Fe-MOF. The chemical structure of the MOF@WPB was then analysed by using Raman spectroscopy (Figure 16(b)) showing peaks around 250 and 297 cm-1corresponding to the O-W-O bending modes and the peak at 756 cm-1corresponds to the 0-W6+-0 stretching bond of the monoclinic WO3. The borate species are confirmed by the peak at 883 cm-1which attributes to the B-0 stretching modes while the broad band around 950 cm-1was observed attributing to the symmetric P-0 stretching modes of POzf tetrahedra. The presence of negatively charged POzf improve the activity and provide resistance against chloride anions in seawater by preventing them from reaching the catalyst surface. No notable peaks were observed in the Raman spectra of amorphous Fe-MOF which was supported by the fact that MOF@WPB had the same peaks as WPB but with reduced intensity which also corroborates the XRD results.

[0328]

[0289] The presence of ionic bonds between WPB and Fe-MOF is then verified by using XPS. The deconvoluted W 4f spectra (Figure 16(c)) have peaks around 35.2 and 37.4 eV attributing to the W6+state for MOF@WPB, while an extra peak was observed for WPB at 34.1 eV corresponding to W4+state. The deconvoluted Fe 2p spectra for MOF@WPB (Figure 16(d)) show peaks at 710.8 and 724.4 eV corresponding to Fe2+state and peaks at 713.6 and 725.7 eV corresponding to Fe3+state with a satellite peak at 718.5 eV. A shift of 0.4 and 0.2 eV was observed in the binding energy of W 4f and Fe 2p spectra for MOF@WPB, respectively while the shift in oxidation state from W4+to higher W6+state is observed resulting from the transfer of electrons between the two constituent materials due to the formation of the ionic bonded heterointerface and ionic Fe-O-W bond.

[0329]

[0290] Figure 16(e) shows the deconvoluted B 1 s spectra for MOF@WPB having peaks at 186.3 and 190.7 eV corresponding to the elemental B and B-0 peaks indicating the presence of B2O3. The peaks were shifted by 0.3 eV as compared to WPB. The deconvoluted P 2p spectra of MOF@WPB show peaks at 133.5 and 130.09 eV attributing to the phosphate group formation and M-P bond. The peaks are shifted by 0.1 eV in WPB. The deconvoluted N 1 s spectra of MOF@WPB and Fe-MOF (Figure 16(f)) show peak at 397.5 eV attributing to M-N bond while the additional peak at 400.1 eV in HS corresponds to the oxidized-N species. The C 1s spectra shows peaks around 284.3, 285.8, and 288.1 eV corresponding to the C-C, C-O / C-N, and O-C=O with a shift in binding energy in comparison to Fe- MOF demonstrating interactions between C, N, and O. The O 1 s spectra of MOF@WPB show peaks at 530, 531.1 , and 532.76 eV corresponding to the M-0 bond, P-0 bond due to phosphate, and atmospheric adsorbed oxygen, respectively while the peak at 531.5 eV in Fe-MOF attributes to the adsorbed OH group. It is hypothesized that the phosphate anions on the surface of WB can prevent its corrosion by repelling the negatively charged chloride anions. The M-0 and P-0 peaks in HS were shifted by 0.5 eV as compared to WPB, due to the charge transfer between Fe-MOF and WPB, further confirming the ionic bond formation between the two constituents. The XPS results, thus, confirm the interactions between Fe-MOF and WPB via ionic -Fe-O-W- bond.

[0330]

[0291] NEXAFS analysis examines how the formation of ionic bond affects the charge states and electron density of the constituent elements. The Fe L-edge spectra (Figure 16(g)) shows typical structure of electronic transition 2p to 3d-t2g and 2p to 3d-egin Fe oxides with a mixture of Fe2+and Fe3+species at 708.3 (a) and 709.9 (b) eV, respectively which is in agreement with the XPS results. The spectra demonstrated that the peak positions and intensities are in correlation with the characteristic spectra of Fe2Os and FesC attributing to the presence of trivalent state and both di- / trivalent states, respectively and presence of both oxidation states indicate the capability of Fe to make additional bonds resulting into the formation of ionic Fe-O-W bond. Furthermore, the reduced intensity of b indicates increased bonding via the Fe3+along with a shift in the photon energy, justifying the HS formation.

[0331]

[0292] Figure 16(h) shows the B K-edge spectra demonstrating a peak at 193.3 eV attributing to the triagonal B species bonding with O confirming the presence of B2O3 which play a crucial role in modulating the B-OH sites at the interface. The differences in the peak intensities and width of WPB and HS with respect to oxide arise due to the additional coordination with P and Fe-MOF, respectively. The O K edge spectra in Figure 16(i) demonstrate spectral features mentioned as A, B, C, D, and E. The first peak at 530.1 eV consists of main peak A and shoulder peak B reflecting the oxygen 2p- states in the d(t2g) band formed by 5d tungsten and 2p oxygen orbitals and non-equivalent oxygen atoms of the first coordination shell and their peak intensities depend on the number of empty 2p oxygen states. The higher intensity peak A corresponds to the higher oxidation state (W6+) of WO3. The intensity of second peak increased in WB and WPB due to the new B-0 and P-0 bond formation, which can affect the interactions between W and O. For Fe-MOF, the O K-edge spectra shows the characteristic features of Fe2Os and FesO4, as observed in the Fe L-edge spectra while the broad peak around 540.5 eV arise due to the multiple scattering paths of Fe atoms away from the O atoms at the centre. The O K-edge of HS show peaks which resemble both Fe-MOF and WPB with a positive shift in binding energy, implying interactions between Fe-MOF and WPB via the ionic Fe-O-W bond.

[0332]

[0293] The N K-edge spectra (Figure 16(j)) mainly consisted of two features representing 1s to pi* transition around 400 eV and 1 s to sigma* transition for C-N above 406 eV which can be attributed to the electronic transition from N1 s to the unoccupied hybridized states of Fe 3d and N2p orbitals. After the interactions between WPB and Fe-MOF, the C-N bond orientation is affected along with the considerable drop in the peak intensity, which occurs due to the bonding between Fe and W via O. The P K-edge spectra of WPB show a strong band at 2151 .6 eV attributing to the lower symmetry around the phosphorus atoms with P-0 and P-W bond, as observed in XPS as well. A positive shift in the photon energy was observed after interactions with Fe-MOF which might be due to modulations in the bonding of P with O and W. The variations in the electronic states and features of the elements thus confirm the interactions between Fe-MOF and WPB via the ionic Fe-O-W bond.

[0333] Morphological features and development of ionic bonded HS

[0294] For the construction of the 2D MOF@WPB HS sheets, the inherent rod-like structure of WO3 was first converted into sheets via heat treatment at different temperatures. The sheet-like morphology was obtained at a temperature of 650 °C, which was confirmed by using TEM demonstrating the conversion of WO3 rods into sheets. The morphology was retained after the addition of boric acid to acquire WB as well as after its phosphorization to obtain WPB (Figure 17(a)) which also demonstrated sheet-like morphology. The 2D WPB then served as an ideal substrate for the development of the 2D Fe-MOF which then resulted in the formation of 2D MOF@WPB HS as shown by the TEM image in Figure 17(c). The HS construction was further confirmed by the high-resolution TEM image of MOF@WPB showing the WPB sheets covered by the Fe-MOF. Furthermore, the lattice spacing of 0.1712 nm for MOF@WPB (Figure 17(d)) corresponds to (024) plane and 0.1861 nm for WPB (Figure 17(b)) corresponds to the (014) plane of monoclinic PWsChe and the selected area electron diffraction (SAED, inset of Figure 17(a) & (c)) also confirms the existence of the monoclinic structure, verifying the XRD results. EDS was conducted to further demonstrate the uniform distribution of elements across the sheets of MOF@WPB.

[0334] Anodic reaction and durability testing of the catalyst

[0335]

[0295] After establishing the structure, oxidation states, and morphology of the catalyst, samples with different B concentration were initially tested for OER in 6.0 M KOH. Figure 18(a) shows the OER overpotentials and LSV curves for different concentrations achieving a current density of 10 mA cm-2(q ) at 1 .23, 1 .422, 1 .434, and 1 .231 V in 6.0 M KOH. However, large oxidation peaks were observed in WB-0.1 and 0.65 which is not a correct indication of the performance. As a results WB-0.21 was phosphorized followed by the construction of the HS with Fe-MOF. The as-synthesized MOF@WPB- 0.21 was then tested in real seawater, achieving a current density of 1 A cm-2at 1.8 V while WPB- 0.21 required a potential of 1.88 V to achieve the same current density (Figure 18(b)-(c)) and Fe- MOF achieved only 0.9 A cm-2at 1 .8 V. In contrast, benchmark lrO2 achieved only 0.2 A cm-2at 1 .8 V demonstrating the huge difference between their electrocatalytic performance. Tafel plots (Figure 18(d)) were then used to understand the reaction kinetics showing Tafel values of 78.71 mV dec-1for MOF@WPB-O.21 which are lower than WPB-0.21 (83.17 mV dec1), WB-0.21 (90.01 mV dec1), Fe- MOF (96.16 mV dec1), and lrO2 (110 mV dec1) in seawater showing fast reaction kinetics of MOF@WPB-O.21.

[0336]

[0296] The electrochemical surface area (ECSA) was also compared to understand the access to active sites by measuring double-layer capacitance (Cdi) of MOF@WPB-O.21 , WPB-0.21 , Fe-MOF and WB-0.21 as shown in Figure 18(e). The Cdi for MOF@WPB-O.21 (536.3 mF cm-2) is higher than the individual WPB-0.21 (225 mF cm2), WB-0.21 (148.9 mF cm2), Fe-MOF (114.5 mF cm2), and lrO2(99 mF cm-2) demonstrating that interfacial interactions between the individual constituents leads to an enhanced access to active sites. These results were further supplemented by EIS where smallest semi-circle was observed for MOF@WPB-O.21 demonstrating lower charge transfer in seawater than constituent individual materials (Figure 18(f)).

[0337]

[0297] MOF@WPB-O.21 demonstrated stable operation for more than 1000 h in 6M seawater without any significant current losses as shown in Figure 18(g). In contrast WPB-0.21 lost almost 30% of its initial current density within 24 h while the commercial lrC>2 lost more than 50% of its initial current density in under 10 h. Furthermore, Fe-MOF was also unstable in seawater as it lost 50% in 10 h. These results indicate that the ionic bond connecting WPB and Fe-MOF is critical in providing strong interactions which makes the HS stable in seawater. GC-MS analysis was carried out and showed the absence of any peaks for chloroacetone, further confirming that no chlorine was generated during the reaction.

[0338]

[0298] The morphological and structural integrity of the electrocatalyst after long-term stability was analyzed using characterizations by XRD (Figure 19(a)) showing similar peaks as that of the HS before operation in seawater confirming no deposition of any unwanted / foreign materials on the electrocatalyst. These results were further supported by EDS results (Figure 19(b)) showing no foreign deposition. The morphological integrity was confirmed by TEM images (Figure 19(c)) showing Fe-MOF over the surface of WPB and HRTEM further confirmed the findings as shown in Figure 19(d). Figure 19(e) demonstrate deconvoluted W 4f spectra showing the doublet peaks at 34.3 and 36.7 V corresponding to the doublet W 4f for W6+states. Similarly, deconvoluted Fe 2p spectra (Figure 19(f)) shows peaks at 713.86 eV and 721 .5 eV corresponding to the Fe3+state and its satellite showing that the oxidation states of W and Fe remained unaffected even after continuous operation in harsh seawater conditions. The presence of higher oxidation states during and after the reaction helps in establishing a more active oxyhydroxide for OER which caters for prolonged activity and stability. The deconvoluted O 1 s spectra (Figure 19(g)) shows peaks at 530.7 and 531 .8 eV corresponding to the M-0 and P-0 bonds while the deconvoluted P 2p spectra in Figure 19(h) shows the phosphate peak at 132.7 eV confirming the presence of P-0 bond. Deconvoluted B 1 s spectra show peak at 187.6 eV corresponding to the elemental B while the absence of B-0 bond is probably due to its hydrolysis to modulate B-OH sites at the interface. The shift in the binding energies were observed for each element due to the conversion of oxyhydroxide phase during the reaction. Therefore, unchanged oxidation states and structural and morphological retention of the catalyst shows its durability, making it suitable for tough industrial conditions.

[0339] Conclusion - Example 5

[0340]

[0299] A 2D MOF@WPB HS was synthesized where Fe-MOF and PW8O26-B2O3 are connected via ionic bonds at the interface leading to improved transfer of electrons, charge transport, and improved stability provided by the strong bonding between the two constituent materials. Addition of B helps in providing anti-corrosive properties to the material as B2O3 modulates the hydroxyl anions at the interface via B-O-B hydrolysis which along with phosphate anions repels the chloride anions present in seawater, thus preventing corrosive chlorine chemistry at the anode. As a result, the electrocatalyst was able to achieve a current density of 1 .75 A cm-2at 2 V while also being stable over 40 days of continuous operations, thus opening a pathway for highly efficient and stable HSs in seawater by creating ionic bonds at the interface.

[0341] Example 6 - Comparative study of 2D-2D HSs and Liu’s rGO / MOF material

[0342]

[0300] Figure 20(a) & 20(b) shows the LSV of the different catalysts and corresponding overpotentials at different current densities, using ZIF-67 / FBT that contains ionic bond at the heterointerface as the reference. Liu’s work shows the least OER activities as compared to the other catalysts of the invention. FBT + ZIF-67 MIX, WO3 + ZIF-67 MIX and Fe2Os + ZIF-67 MIX were synthesised by directly mixing the constituent 2D materials together, and are expected to be connected together by either electrostatic attractions, covalent bonding or both. All these materials outperformed Liu’s work, with FBT + ZIF-67 MIX being the best performing amongst them. However, comparing the OER activity of FBT + ZIF-67 MIX with that of ZIF-67 / FBT, the ZIF-67 / FBT exhibited superior activity despite having same 2D materials as building block. The superior activity of ZIF- 67 / FBT can further be seen from Figure 20(b) as it exhibited the least overpotentials at different current densities as compared to other catalysts. The superior activity of ZIF-67 / FBT is due to the presence of strong ionic bonds at the heterointerface, which gives the material new properties by taking advantages of the constituent 2D materials properties, while mitigating their shortcomings. Further, the stability test of FBT + ZIF-67 MIX shown in Figure 20(c) indicates that it loses over 57% of its activity after just 24 hours as against ZIF-67 / FBT that retains its activity for over 500 hours, thus, showing the importance and influence of ionic bonding at the heterointerface of the 2D / 2D HS.

[0343]

[0301] As acknowledged above, Liu’s 2D MOF on the flat rGO surface with controlled perpendicular orientation provides a solid / liquid assisted synthesis of a material presenting perpendicular growth of a metal-organic frameworks (MOFs) on a rGO surface. Notably, however, the template for growth is a purely carbon-based material in contrast to the inorganic materials of the invention and particularly the tuned inorganic materials. Liu does not recognise or suggest that superior materials in terms of electrocatalytic activity could be formed by selection of certain 2D inorganic materials templates, nor that these materials would result in strongly bonded interfaces and accompanying advantages described herein.

[0344]

[0302] Liu does not consider or attempt to explain the nature of the bonding created via the direct synthesis strategy between the heterointerface, and in particular, does not consider whether the bonds are strong or weak. There is nothing in Liu indicating role of the bonding (ionic / covalent) at an heterointerface in enhancing electrochemical performance. Further, given the nature of Liu’s starting materials, the interactions between Liu’s HS cannot be ionic in nature or strong enough to sustain operation under the harsh electrolytic conditions of strong alkalinity I seawater. There is nothing in Liu that would suggest any utility as a water splitting electrocatalyst suitable for industrial use.

[0345]

[0303] Furthermore, an evident significant challenge associate with Liu’s HS material for seawater catalysis is the exposure of base materials due to the perpendicular growth of MOF sheets over the carbon structure. The perpendicular arrangement of the dissimilar materials means the MOF cannot provide the required complete shield of the rGO from the corrosion or parasitic attack of chlorine present in seawater. This is one particularly desirable benefit of the method of the invention, which is based on the discovery that a solution-based direct synthesis process involving a 2D ultrathin transition metal inorganic material with a 2D ultrathin transition metal MOF material can produce vertically stacked HS composite material having strongly bonded heterointerface between the materials and ensure that the required complete shielding of the active TM surface is present. The resulting strong connections (ionic and covalent) have been found to promote the electrochemical performance of the materials as catalysts even, in harsh conditions like seawater electrolytes.

[0346] Example 7 - Bimetallic MOF / FBT

[0347] Synthesis of FeCo-OF / FBT

[0348]

[0304] FeCh, CO(N03)2and BDC linker (at a mole ratio of 1 :1 :1 ) were dissolved in a solution containing FBT (dual modified O-functionalised titanium carbide (TMC), Ti3C2O2, wherein the non- metal dopant is B and the electrochemically active metal is Fe) + DMF (dimethylformamide) (amount of FBT = 15 wt.% of total metal composition in the salts) followed by the addition of ethanol, water and TEA (triethylamine) in a HDPE bottle. After that, the solution was stirred overnight for ~15 h, allowed to age at room temperature and a steady state for another 3 h, and the solid product was recovered by centrifugation. The solid product was washed with DI water (4 times) and ethanol once. The resultant wet residue was dried at 50 °C under vacuum, and the product was labelled FeCo-OF / FBT- 15. A similar procedure was repeated for 0%, 10% and 20% FBT and the product labelled FeCo-OF, FeCo-OF / FBT-10 and FeCoOF / FBT-20, respectively.

[0349] Discussion - Example 7

[0350]

[0305] The XRD pattern of pristine FeCo-OF shows that it is an amorphous material, as there is no diffraction (Figure 21a). Conversely, FBT is a crystalline material that preserves much of the original structure of the parent Ti3C2TxMXene, as evident by the XRD patterns, where reflections from the (006), (200), and (110) planes (Figure 21a) are observed. On the other hand, the XRD pattern of FeCo-OF / FBT (Figure 21a) show the presence of peaks corresponding to FBT, indicating that both FBT and FeCo-OF / FBT have a similar structure. Further, the TEM of FeCo-OF / FBT (Figure 21b) reveals the presence of sheet-on-sheet morphology, indicating that FBT also serves as a platform for FeCo-OF growth, leading to 2D / 2D HS formation. The crystalline nature of FeCo-OF / FBT was also confirmed by the HRTEM analysis (Figure 21c), where the presence of lattice fringes (Figure 21c insert) is observed.

[0351]

[0306] The chemical states of the elements in the materials were investigated by X-ray photoelectron spectroscopy (XPS). The Fe 2p spectrum of FeCo-OF / FBT (Figure 21d) shows peaks at 711 .05 and 713.3 eV, corresponding to Fe2+and Fe3+, respectively. The Fe2+shows a positive 0.31 eV peak shift, while the Fe3+shows a positive 0.22 eV peak shift compared to the FBT. These positive peak shifts signify the donation of electrons from the Fe-atom in FBT to the neighbouring oxygen atom at the interface to FeCo-OF / FBT. Conversely, the Fe3+of FeCo-OF / FBT experiences a negative peak shift of - 178 0.23 eV compared to FeCo-OF, suggesting the collection of electrons from the FeCo-OF / FBT heterostructure interface. These observations suggest the transfer of electrons from FBT to FeCo-OF through the heterointerface, indicating the presence of ionic bonding.

[0352]

[0307] In the O 1s spectrum (Figure 21e), FeCo-OF / FBT exhibits three distinct peaks at 529.86, 531.56, and 533.21 eV, corresponding to metal-oxygen (M-O), metal-hydroxide (M-OH)Zdefective oxygen, and surface water adsorption bonds, respectively. The M-OH peak of FeCo-OF / FBT shows a peak shift of -0.4 eV relative to FBT, confirming the collection of electrons by surface O-atom from Fe-atom of FBT for onward donation to FeCo-OF through O-bridge at the interface. Further, FeCo- OF / FBT shows a peak shift of +0.12 eV relative to FeCo-OF, indicating the transfer of electrons from the O-bridge at the interface to FeCo-OF, confirming the presence of Fe— >0— >FeCo-OF electrons transfer channel.

[0353]

[0308] Moreover, the Co 2p spectrum of FeCo-OF / FBT (Figure 21f) displays the presence of Co2+and Co3+peaks at 784.79 and 781 .49 eV, respectively. The Co2+and Co3+peaks of FeCo-OF / FBT exhibit negative peak shifts of -0.83 and -0.55 eV, respectively, implying the collection of electrons from the interface through O— >Co transfer channels [18, 21], These observations confirm the presence of Fe— >0— >Co electrons transfer channels in the FeCo-OF / FBT heterostructure, confirming the presence of ionic bonds at the heterointerface.

[0354]

[0309] The synthesised catalysts' OER performance was evaluated in an alkaline seawater electrolyte, with results compared to commercial RuC>2. OER data were collected at a scan rate of 5 mV s-1using the back-sweep technique to avoid oxidation peaks that are often present during forwardsweep scans with nickel foam. As shown in (Figure 22a), all heterostructure catalysts outperformed FBT, FeCo-OF, and RuO2, highlighting the benefits of heterostructure design. This enhanced performance is attributed to abundant catalytic active sites and rapid charge transfer triggered by the ionic bond at the heterostructure interface. The FeCo-OF / FBT-15 heterostructure, in particular, required a low overpotential of 330 and 450 mV to reach a current density of 500 and 1000 mA cm-2, respectively, outperforming FBT (420 and 540 mV) and FeCo-OF (390 and 530 mV), and RuC>2 (630 and 770 mV) (Figure 22b).

[0355]

[0310] Interestingly, above 1000 mA cm-2, FBT requires lesser overpotentials to attain higher current density than FeCo-OF (Figure 22a and b). This observation suggests that besides rapid charge transfer, FBT aided the FeCo-OF / FBT-15 in sustaining its OER activity at a higher current density beyond 1000 mA cm-2. Also, FeCo-OF / FBT-15 demonstrated the best kinetics, achieving the lowest Tafel slope of 28.93 mV dec-1(Figure 22c), compared to FBT (54.36 mV dec-1), FeCo-OF (53.06 mV dec-1), and RuO2 (109.97 mV dec-1).

[0356]

[0311] For industrial applications, the long-term durability of a catalyst is a key factor in determining its effectiveness for electrochemical processes. To evaluate this, we assessed the stability of FeCoOF / FBT-15 using the amperometry i-t curve method in alkaline seawater at 1 .9 V vs RHE. As shown in Figure 22d, the FeCo-OF / FBT-15 catalyst is stable for 2000 h at 1 .5 A cm-2, which is 3 times the industrial requirement. To further investigate the robustness of the FeCo-OF / FBT-15 catalyst, full alkaline seawater splitting was carried out using a commercial Pt / C as the counter electrode. The FeCo-OF / FBT-15 catalyst delivers a high current density of 1 A cm-2at a low cell voltage of 1 .57 V (Figure 22e). Further, the catalyst operates continuously for over 250 h and above the industrial required current density of 0.5 A cm-2(Figure 22f). These results demonstrate the efficiency and robustness of the developed FeCo-OF / FBT-15 catalyst.

[0357]

[0312] It will be apparent to the person skilled in the art that while the invention has been described in some detail for the purposes of clarity and understanding, various modifications and alterations to the embodiments and methods described herein may be made without departing from the scope of the inventive concept disclosed in this specification.

Claims

Claims1. A vertical layered A-[B-A]ntype heterostructure assembly of a 2D ultrathin inorganic solid-state material (‘B’) and a non-noble metal organic framework (MOF) material (‘A’) that covers surfaces of each layer of the 2D ultrathin inorganic solid-state material in the assembly in a parallel arrangement to the 2D ultrathin inorganic solid-state material, wherein n is a positive integer, wherein heterointerfaces of the assembly comprise bridging MIM-X-MMOFionic bonds that stabilise heterointerfaces between the 2D ultrathin inorganic solid-state material and the non- noble MOF material, where bridging X is -O-, -S-, -Se-, -C-, -N-, -P-, or -B-, MIMis a metal in the 2D ultrathin inorganic solid-state material, and MMOFis a metal in the non-noble MOF material; wherein the 2D ultrathin inorganic solid-state material comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; and further comprises one or more of: metal-0 (oxide), metal-N (nitride), metal- C (carbide), metal-C / metal N (carbonitride); metal-S (sulfide) and / or M-P (phosphide) functionalities; and optionally wherein the 2D ultrathin inorganic solid-state material is doped with one or more non-metal dopants, preferably selected from the group consisting of: B, P, C, N, O, Si, S, As and Se.

2. The heterostructure assembly of claim 1 , wherein one or more non-metal dopants form chloride ion shielding polyanions selected from one or more of: borate, sulfate, phosphate, chromate, permanganate, carbonate, and nitrate.

3. The heterostructure assembly of claim 1 or claim 2, wherein the non-noble electrochemically active transition metals are selected from Fe, Ni, Co, Cu, Cr, Mo and W, preferably Fe.

4. The heterostructure assembly of any one of the preceding claims, wherein the 2D ultrathin inorganic solid-state material is a transition metal oxide (TMO); a transition metal nitride (TMN); a transition metal sulfide (TMS); a transition metal carbide (TMC), preferably a MXene doped with at least one electrochemically active metal; transition metal phosphide (TMP); or mixed variants thereof, optionally doped with one or more non-metal dopants selected from B, C, N, O, P, Si, S, As and Se, preferably doped with B, or with B and P; and optionally further doped with at least one non-noble electrochemically active metal, optionally selected from Fe, Ni, Co, Cu, Cr, Mo and W.

5. The heterostructure assembly of any one of the preceding claims, wherein the non-metal dopant is one or more of B, C, S and P, forming polyanions selected from one or more of borate, sulfate, phosphate, chromate, permanganate, carbonate, and nitrate, preferably borate, carbonate,sulfate and / or phosphate, on surfaces of the 2D ultrathin inorganic solid-state material that repel Ch from the one or more non-noble electrochemically active metals of the 2D ultrathin inorganic solid- state material.

6. The heterostructure assembly of any one of the preceding claims, wherein MMOFis selected from the group consisting of: Co, Ni, Fe, Cu, Zn, Zr, Ti, and combinations thereof, and MIMis selected from the group consisting of: Fe, Ni Co, W, Ti, Cu, V and Mo.

7. The heterostructure assembly of any one of the preceding claims, wherein the non-metal dopant is B and P, forming borate and phosphate polyanion functionalities that repel Ch from the non- noble electrochemically active metals of the inorganic solid-state material, optionally wherein the inorganic solid-state material comprises one or more of Fe, Co, Ti, V, Mo, W, and / or Ni.

8. The heterostructure assembly of any one of the preceding claims, wherein the 2D ultrathin inorganic solid-state material is selected from:- non-noble transition metal oxide (TMO), iron oxide, such as Fe20s, tungsten oxide such as WO3, vanadium oxide, such as V2O5, molybdenum oxide such as M0O3;- non-noble transition metal MXene or functionalised MXene, preferably B-TisC2Tx, where Tx is O;- non-noble transition metal sulfide (TMS), such as iron sulfide such as FeS?, cobalt sulfide such as C0S2, nickel sulfide such as NiS2;- B doped non-noble transition metal MXene or B-doped functionalised MXene, preferably B- TisC2Tx or B-Fe-TisC2Tx, preferably where Tx is O, F, OH or Cl, preferably O;- B doped 2D non-noble transition metal sulfide (TMS), for example, B-doped iron sulfide such as B-FeS2, or B-doped cobalt sulfide such as B-C0S2, or B-doped nickel sulfide such as B- NiS2,- B doped 2D non-noble transition metal oxide (TMO), for example, B-doped iron oxide such as B-Fe2Os, B-doped tungsten oxide such as B-WO3, B-doped vanadium oxide such as B-V2O5, B-doped molybdenum oxide such as B-M0O3;- B doped and / or Fe doped non-noble transition metal oxide (TMO), for example, B- and Fe- doped iron oxide such as B-doped iron oxide such as B-Fe2Os or B-a-Fe2Os or B- and Fe- doped tungsten oxide such as B-Fe-WOs;- B doped and Fe doped non-noble transition metal sulfide (TMS), for example, B- and Fe- doped iron sulfide such as B-Fe-FeS2, B- and Fe- doped nickel sulfide such as B-Fe-NiS2, B- and Fe- doped tungsten sulfide such as B-Fe-WS2, B- and Fe- doped vanadium sulfide such as B-Fe-V2Ss, or B- and Fe doped molybdenum sulfide such as B-Fe-MoS2, B- and Fe- doped cobalt sulfide such as B-Fe-CoS2.

9. The heterostructure assembly of any one of the preceding claims, wherein the 2D ultrathin inorganic solid-state material is selected from: Fe2Os; WO3.B2O3; B-Fe2Os; B-Fe-Ti3C2Tx, where X is O; B-FeS2; B-C0S2; B-NiS2; P-WsO26-B2O3; MoN, M0S2, and B-C0S2.

10. The heterostructure assembly of any one of the preceding claims, wherein the bridging MIM- X-MMOFionic bonds are selected from Fe-O-Co; W-O-Co; and W-O-Fe, optionally wherein the heterointerface further comprises one or more additional bonds selected from metal-N bonds, B-0 bonds, and P-0 bonds.

11. The heterostructure assembly of any one of the preceding claims, wherein the heterostructure assembly is selected from Fe20s I CoMOF; WO3.B2O3 1 CoMOF; B-Fe20s I NiMOF; B-Fe-Ti3C2TxI CoMOF; B-FeS2 / NiMOF; B-C0S2 / NiMOF B-NiS2 / NiMOF; and P-W8O26.B2O3 / FeMOF.

12. The heterostructure assembly of any one of the preceding claims, wherein the 2D ultrathin inorganic solid-state material (‘B’) is a non-noble transition metal oxide (TMO) that comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; optionally wherein the TMO is Fe20s or WO3; wherein MIMis a metal in the TMO; and optionally, wherein the TMO is doped with one or more non-metal dopants selected from the group consisting of: B, P, C, N, O, Si, S, As and Se, preferably B, S, P or B and P.

13. The heterostructure assembly of any one of claims 1 to 11 , wherein the 2D ultrathin inorganic solid-state material (‘B’) is a non-noble transition metal sulfide (TMS) that comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; optionally wherein the TMS is FeS2, C0S2 or NiS2; wherein MIMis a metal in the TMS; and optionally, wherein the TMS is doped with one or more non-metal dopants, preferably selected from the group consisting of: B, P, C, N, O, Si, S, As and Se, preferably B, S, P, or B and P.

14. The heterostructure assembly of any one of claims 1 to 11 , wherein the 2D ultrathin inorganic solid-state material (‘B’) is a non-noble transition metal doped MXene or a non-noble transition metal doped functionalised MXene material (‘B’) that comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; optionally wherein the MXene is V2CTX, Ti2CTx, Mo2CTxor the functionalised MXene is TisC2O2, wherein heterointerfaces of the assembly comprise bridging MMXene-X-MMOFionic bonds that stabilise heterointerfaces between the MXene and the MOF material, where bridging X is -O-, -S-, - Se-, -C-, -N-, -P-, or -B-, MMXeneis a metal in the MXene; andoptionally, wherein the MXene is doped with one or more non-metal dopants, preferably selected from the group consisting of: B, P, C, N, Si, S, As and Se, more preferably B, S, P, or B and P.

15. Use of the heterostructure assembly of any one of claims 1 to 14 as an electrocatalyst, optionally as an OER electrocatalyst for water splitting, optionally as an OER electrocatalyst for water splitting of seawater.

16. A process of forming an electrocatalyst comprising the steps of: providing a liquid dispersion of an inorganic solid state material (‘B’) in 2D ultrathin sheet form as an inorganic template for non-noble metal organic framework (MOF) material (‘A’) growth thereon in a parallel arrangement to the 2D ultrathin inorganic solid-state material, wherein the inorganic template comprises or is doped with one or more non-noble electrochemically active transition metals, preferably selected from Fe, Ni, Co, Cu, Cr, Mo and W; and further comprises one or more of: metal-0 (oxide), metal-N (nitride), metal-C (carbide), metal-C / metal N (carbonitride); metal-S (sulfide) and / or M-P (phosphide) functionalities, and optionally is doped with one or more non-metal dopants selected from the group of non-metals consisting of: B, C, N, O, P, Si, S, As and Se; electrostatically associating at least one basic organic coordination ligand species suitable for MOF fabrication with the inorganic template; providing a liquid solution of one or more acidic non-noble transition metal (MMOF) salts to the dispersion of inorganic template having electrostatically associated ligand species; forming a vertical layered A-[B-A]ntype heterostructure assembly of the 2D ultrathin inorganic material (‘B’) with the non-noble metal organic framework MOF material (‘A’) by allowing complexation of acidic MMOFions of the salt with the at least one basic organic coordination ligand species associated with the inorganic template to form the MOF in coverage over surfaces of the inorganic template, whereby during complexation, MMOFforms bridging MIM-X-MMOFionic bonds with metals in the inorganic material, MIM, at heterointerfaces between the inorganic template and MOF, wherein bridging X is -O-, -S-, -Se-, -C-, -N-, -P-, or -B-, and wherein n is a positive integer.

17. The process of claim 16, wherein the liquid dispersion comprises a slightly polar solvent, such as an alcohol including methanol or ethanol, or an aqueous alcohol solution, preferably methanol in water, wherein the basic organic coordination ligand species is provided to the liquid dispersion of inorganic template in solution form as an aqueous solution of dissolved basic organic coordination ligand species, and wherein the acidic non-noble transition metal salts are provided to the dispersion of the inorganic template and basic organic coordination ligand species in an aqueous solution.

18. The process of claim 16 or claim 17, wherein the acidic non-noble transition metal is selected from the group consisting of: Co, Ni, Fe, Zn, Cu, Ti, Zr, and combinations thereof.

19. The process of any one of claims 16 to 18, wherein the at least one basic organic coordination ligand species is selected from the group consisting of: a dicarboxylic acid, a tricarboxylic acid, a tetracarboxylic acid, an imidazolate, a bisimidazolate, a triazolate, a pyridinate and combinations thereof, or is preferably an alkyl imidazole, or is more preferably 2-methylimidazole.

20. The process of any one of claims 16 to 19, wherein the inorganic template is a dual modified functionalised MXene, Mn+iXnTx, where n is 1 , 2 or 3, M is an early transition metal (Group 3 to 6 metal), X is C and / or N, and Txis a hydrophilic surface functional group selected from one or more of O, F, OH, and Cl, wherein the functionalised MXene is in freestanding or delaminated 2D sheet form, and wherein the functionalised MXene is prepared by: optionally, functionalising an MXene such that a majority of surface functional groups in the MXene material are oxygen; modifying the functionalised MXene by replacing at least a portion of the M atoms in the functionalised MXene comprising a majority of surface functional groups that are oxygen with one or more doped non-metal atoms selected from B, N, C, P, Si, S, As and Se, preferably B; and modifying the functionalised MXene by replacing at least one surface functional group and / or at least one M in the MXene with one or more electrochemically active metals selected from the group consisting of: Fe, Ni and Co, thereby forming the dual modified functionalised MXene.

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