Cross-Reference To Related Application

[0001] This application claims the benefit of priority of Singapore Patent Application No. 10201610408S, filed 12 December 2016, the content of it being hereby incorporated by reference in its entirety for all purposes.

Technical Field

[0002] The present disclosure relates to a method for producing an amorphous lithium treated transition metal phosphochalcogenide electrocatalyst. The present disclosure also relates to an amorphous lithium treated transition metal phosphochalcogenide electrocatalyst. The present disclosure further relates to a method for producing a lithium treated transition metal phosphochalcogenide electrocatalyst and such a lithium treated transition metal phosphochalcogenide electrocatalyst. In addition, the present disclosure relates to uses of such electrocatalysts.

Background

[0003] Hydrogen is a clean and sustainable energy carrier for future green energy technologies. To generate hydrogen efficiently, the electrochemically reduction of water may be an attractive approach. Generally, the production efficiency of hydrogen through water reduction depends on the performance of electrocatalytic materials. Therefore, robust electrocatalysts with high activity are in demand. Even though the search for new catalysts has been pursued for a long time, the species of conventionally available catalysts remain limited, particularly electrocatalytically inert materials, which may have been neglected.

[0004] With better understanding of structure-property relationship, designing highly efficient catalysts through structure manipulation of materials have attracted great attention. Nanostructured layered materials, such as transition metal dichalcogenides (TMDs), may be ideal candidates for structural engineering because of their diverse species and highly anisotropy structure. The engineering of layered materials with emphasis on changing their crystal structures, morphologies, and electronic properties, allows improvement to their catalytic performance as well as to develop highly efficient catalyst for hydrogen production. Generally, the structure engineering of inorganic layered materials (e.g. M0S2) tends to be focused on the following aspects: (i) decrease of their thickness, (ii) reduction of their lateral sizes, (iii) creation of sulphur vacancies/defects, (iv) phase engineering, and (v) doping or alloying.

[0005] The aforementioned aspects of structural engineering are capable of altering intrinsic properties of materials, leading to enhanced performance in electrochemical hydrogen evolution reaction (HER).

[0006] What is of particular interest lies in the fact that physicochemical properties of an amorphous material may differ greatly from its crystalline counterpart. Therefore, crystallinity engineering (e.g. transformation from crystalline to amorphous structure) of a material may also be one of the considered approaches to tune the material's properties and functionalities. Referring to conventional TMDs, it may be reasonable to assume that all the aforementioned structure tuning approaches may be achieved in other layered compounds to realize optimized performance for HER. Moreover, the ability to tweak multiple structure engineering (size, shape, defects, doping, crystallinity etc.) in one pot may amplify the structural importance toward HER performance.

[0007] In this connection, ternary palladium phosphosulphide (Pd3P ₂ S ₈ ) is a layered compound with an interlocked layered structure and waved Pd-P-S layers, which was first reported in 1971. However, there is insufficient development and exploration of Pd ₃ P2S ₈ for potential applications, particularly as an electrocatalyst for use in HER and/or oxygen reduction reaction.

[0008] There is thus a need to provide for a method that resolves and/or ameliorates one or more of the issues mentioned above. The method includes producing an electrocatalyst from a transition metal phosphochalcogenide. The method involves structure engineering (e.g. crystallinity engineering) of the transition metal pho sphochalcogenide .

[0009] There is also a need to provide for an electrocatalyst derived from transition metal phosphochalcogenide that addresses one or more of the issues mentioned above, at least for use in HER. Summary

[0010] In one aspect, there is provided for a method for producing an amorphous lithium treated transition metal phosphochalcogenide electrocatalyst, comprising: constructing a cathode comprising transition metal phosphochalcogenide crystalline particles;

configuring the cathode, a lithium anode and a lithium-based electrolyte to form a battery;

subjecting the battery to galvanostatic discharge to transform the transition metal phosphochalcogenide crystalline particles into amorphous lithium treated transition metal phosphochalcogenide particles; and

extracting the amorphous lithium treated transition metal phosphochalcogenide particles from the cathode for use as the amorphous lithium treated transition metal phosphochalcogenide electrocatalyst.

[0011] In another aspect, there is provided for a method for producing a lithium treated transition metal phosphochalcogenide electrocatalyst, comprising:

contacting transition metal phosphochalcogenide crystalline particles with an organolithium solution to form an organolithium suspension, wherein the organolithium solution comprises an organolithium; and

allowing the organolithium suspension to stand for a period of time during which the organolithium dissociates to provide lithium for penetrating into each of the transition metal phosphochalcogenide crystalline particles to form the lithium treated transition metal phosphochalcogenide electrocatalyst.

[0012] In another aspect, there is provided for an amorphous lithium treated transition metal phosphochalcogenide electrocatalyst comprising a transition metal phosphochalcogenide treated with lithium to form an amorphous lithium treated transition metal phosphochalcogenide, wherein the amorphous lithium treated transition metal phosphochalcogenide comprises one or more sulphur atom vacant sites and/or one or more phosphorus atom vacant sites.

[0013] In another aspect, there is provided for a lithium treated transition metal phosphochalcogenide electrocatalyst obtained according to the method as disclosed above. [0014] In another aspect, there is provided for use of an amorphous lithium treated transition metal phosphochalcogenide electrocatalyst as disclosed above for hydrogen evolution reaction and/or oxygen reduction reaction.

[0015] In another aspect, there is provided for use of a lithium treated transition metal phosphochalcogenide electrocatalyst as disclosed above for hydrogen evolution reaction and/or oxygen reduction reaction.

Brief Description of the Drawings

[0016] The drawings are not necessarily to scale, emphasis instead generally being placed upon illustrating the principles of the invention. In the following description, various embodiments of the present disclosure are described with reference to the following drawings, in which:

[0017] FIG. 1A shows an atomic structure of the three layer building blocks of Pd ₃ P ₂ S ₈ along the a axis. Crystalline particles of Pd ₃ P ₂ S ₈ have repeated segments of this atomic structure over a long-range.

[0018] FIG. IB shows a coordination polyhedra of Pd ₃ P ₂ S ₈ viewed along the c axis, illustrating that the waved Pd-P-S layer along the ab plane is composed of the PdS ₄ tetragon and PS ₄ tetrahedron. The S atoms with different sites are labeled as S( _a ) (S atoms at apical position) and S( _S ) (S atoms at side position). The coordination circumstance of Pd with four equivalent S atoms and the coordination circumstance of P with two types of P-S bonds are shown.

[0019] FIG. 1C shows the experimental and simulated X-ray diffraction (XRD) patterns of single crystal Pd ₃ P ₂ S ₈ .

[0020] FIG. 2A shows a photograph of Pd ₃ P ₂ S ₈ single crystals. The scale bar represents 1 cm.

[0021] FIG. 2B shows an optical microscopy image of Pd ₃ P ₂ S ₈ single crystals. The scale bar represents 100 μηι.

[0022] FIG. 2C shows a scanning electron microscopy (SEM) image of Pd ₃ P ₂ S ₈ single crystals. The scale bar represents 1 μηι.

[0023] FIG. 2D shows a SEM image of individual plate-like crystals of Pd ₃ P ₂ S ₈ with cleavage planes. The scale bar represents 100 μπι. [0024] FIG. 2E shows a high resolution transmission electron microscopy (HRTEM) image of individual plate-like crystal of Pd ₃ P2S8. The scale bar represents 2 nm.

[0025] FIG. 3A shows the preparation of Li-PPS (lithium treated Pd ₃ P ₂ S ₈ , also represented herein by simulated LiPd ₃ PS ₇ structure) nanodots (NDs) from layered bulk Pd ₃ P ₂ S8 crystals by using the electrochemical lithiation method.

[0026] FIG. 3B shows two TEM images of Li-PPS NDs. The left image is the TEM image of Li-PPS NDs with scale bar representing 20 nm while the right image is a HRTEM image of Li-PPS NDs with scale bar representing 2 nm. The inset in the left TEM image is a morphology sketch of Li-PPD ND and the inset of the right HRTEM image is a photo of Li-PPS NDs in water.

[0027] FIG. 3C shows an atomic force microscopy (AFM) image of Li-PPS NDs on the left while the right graph illustrates the height profile along the black line indicated in the left AFM image. The scale bar in the left AFM image represents 500 nm.

[0028] FIG. 3D shows two graphs indicating the statistical analysis of the sizes of 100 Li-PPS NDs measured by TEM (left) and the statistical analysis of the heights of 200 Li-PPS NDs measured by AFM (right).

[0029] FIG. 3E shows a HRTEM image of Li-PPS NDs. The scale bar represents 2 nm. The inset is a selected area electron diffraction (SAED) pattern of Li-PPS NDs.

[0030] FIG. 4A shows a plot of the size distribution of Li-PPS NDs in water measured by dynamic light scattering (DLS). The hydrodynamic size of Li-PPS NDs was estimated to be 6.7 nm. Due to the solvent molecules adsorbed on the surface of Li-PPS NDs, the hydrodynamic size measured by DLS (about 6.7 nm) typically may be larger than that measured by TEM, e.g. 2.5 ± 0.9 nm (see FIG. 3D). The narrow and monomodal peak demonstrates the narrow size distribution of Li-PPS NDs.

[0031] FIG. 4B shows a plot of the surface potential of Li-PPS NDs in water (pH=7) measured by DLS. Here, the surface potential of Li-PPS NDs is determined to be about -40 mV.

[0032] FIG. 5A shows the Pd K-edge XANES spectra of Pd ₃ P ₂ S ₈ crystal and Li-PPS NDs.

[0033] FIG. 5B shows the Pd K- dge XAFS Fourier transform (FT) spectra of Pd ₃ P ₂ S ₈ crystal and Li-PPS NDs, and the corresponding fitting curves for the Pd-S peaks. Solid and dashed lines represent the experimental data and the fitting curves, respectively. The data used for FT are in the range of 2-12 A ^"1 , and all FT spectra have not been phase-corrected. The Harming window function was applied and the window parameter was set to 20.

[0034] FIG. 5C shows the Li-PPS nanodots imaged using the nanobeam electron diffraction technique.

[0035] FIG. 5D shows the diffraction pattern acquired from the highlighted region in FIG. 5C (represented by the square with broken outlines).

[0036] FIG. 5E shows the nuclear reaction analysis (NRA) results of Pd ₃ P ₂ S8 crystal and Li-PPS NDs in the top and bottom images, respectively.

[0037] FIG. 5F shows the optimized atomic structure model of LiPd ₃ PS7 along the c axis (left) and b axis (right). The phosphorus vacancy (Vp) and sulphur vacancy (Vs) of LiPd ₃ PS ₇ are located in the same PS ₄ tetrahedron.

[0038] FIG. 5G shows the distribution of the Pd-S bond length (R) in the Pd ₃ P ₂ S ₈ crystal, simulated Pd ₃ P ₂ S ₈ with Li incorporation, simulated Pd ₃ P ₂ S ₈ with sulphur vacancies at S( _a ) sites (Vs(a)) and sulphur vacancies at S( _S ) (Vs(s)), and simulated LiPd ₃ PS ₇ .

[0039] FIG. 6A shows the &¾k) data for Pd-S shell of Li-PPS NDs (solid line) and the fitting curve obtained from the atom distribution function (ADF) method (dotted line). The data were filtered from the FT spectrum of Li-PPS NDs (FIG. 5B) by inverse FT (IFT). The IFT range is 1.5 to 2.2 A, which is mainly contributed by the Pd-S scattering signal.

[0040] FIG. 6B shows the fitting curve in R space using Winxas based on the parameters calculated from the ADF method. Two peaks were clearly observed. Each peak represents one Pd-S subshell. The peak position is the bond length, and the peak area is the coordination number. The bond length difference of the two Pd-S subshells is about 0.13 A. Since the resolution limit of ADF method in the data range of 2 to 12 A ^"1 is 0.05 A, the calculation result is considerably reliable. The crystalline structure of Pd ₃ P ₂ S8 was obtained from the single crystal XRD data (see tables 1 and 2), so the conventional curve fitting based on the least-squares method was carried out using Winxas for Pd ₃ P ₂ S ₈ crystal. Due to the structural distortion of Li-PPS NDs, the ADF method, which has 3 times higher spatial resolution than conventional methods, was used to investigate the structural features of Li-PPS NDs. The fitting curve of Pd- S shell for Li-PPS NDs is shown in FIG. 6A.

[0041] FIG. 6C shows the sulphur iC-edge XANES spectrum of the Pd ₃ P ₂ S ₈ crystal and Li-PPS NDs.

[0042] FIG. 6D shows the measured (the line connecting the dots) and SIMNRA fitted (the black curve) NRA results of Li-PPS NDs in the range of 4.8 to 8.8 MeV. The NRA characterization was used to confirm the existence of lithium in the Li-PPS NDs. As compared with Pd ₃ P ₂ S ₈ crystal, the NRA spectrum of Li-PPS NDs exhibited strong Li signal in the high energy region between 4.8 and 8.8 MeV (FIG. 5E). The Li profile of Li-PPS NDs was fitted through the SIMNRA. The fitting result indicates that the Li content is about 10 at% in the Li-PPS NDs.

[0043] FIG. 6E shows the K-edge XANES spectra of the four-coordinated Pd atoms in the simulated LiPd ₃ PS ₇ and as-prepared Li-PPS NDs.

[0044] FIG. 7 shows the current-voltage (I-V) characteristic curves for the Pd ₃ P ₂ S ₈ crystal and Li-PPS NDs. Compared to Pd ₃ P ₂ Sg crystal, the conductivity of Li-PPS NDs is greatly enhanced in the sweeping voltage of -2 to 2 V.

[0045] FIG. 8A shows the polarization curves (IR compensated) for HER obtained on the bare glassy carbon electrode (GCE), and the GCE modified with Li-PPS NDs, Pd ₃ P ₂ S ₈ crystal, Pd black and conventional Pt/C catalyst. The weights of noble metals in the loaded catalysts were kept the same. Sweep rate was 5 mV s ^"1 .

[0046] FIG. 8B shows Tafel plots for the catalysts derived from the listed for FIG. 8A.

[0047] FIG. 8C shows the atomic models for hydrogen atoms adsorbing at (1) S( _S > site of Pd ₃ P ₂ S ₈ crystal [denoted as Pd ₃ P ₂ S ₈ (H-S)] (left image), and (2) Si and S ₂ sites of LiPd ₃ PS ₇ [denoted as Li-PPS (H-Si) and Li-PPS (H-S ₂ )] (right image). Si and S ₂ represent the three-coordinated and four-coordinated sulphur atoms at the edge of LiPd ₃ PS ₇ . Amorphous LiPd ₃ PS ₇ has an atomic structure as shown in the right image of FIG. 8C, which lacks the long-range order of an arrangement that is present in crystalline Pd ₃ P ₂ S ₈ .

[0048] FIG. 8D shows a calculated free energy diagram for hydrogen evolution at equilibrium (U=0) with adsorbed configurations shown in FIG. 8C. [0049] FIG. 9A shows a comparison of the onset potentials required for the metal sulphide-based and metal phosphide-based catalysts to initiate HER.

[0050] FIG. 9B is a table showing comparison of electrocatalytic parameters of metal sulphide-based and metal phosphide-based catalysts for HER.

[0051] FIG. 10 shows a bar chart comparing the Tafel slopes for some metal sulphide-based and metal phosphide-based catalysts.

[0052] FIG. 11A shows the long-term stability test for Li-PPS NDs in which the polarization curves (IR compensated) were recorded at scan rate of 5 mV s ^"1 after 7500 potential cycles.

[0053] FIG. 1 IB shows a water splitting device powered by two 1.5 V AAA batteries in series, using Li-PPS ND-modified carbon paper as cathode and Ir0 ₂ -modified carbon paper as anode. The scale bar represents 1 cm.

[0054] FIG. 12 depicts a schematic illustration of the growth of Pd ₃ P2S8 single crystals. The chemical state (g/1) of phosphorus (P) and sulphur (S) means that the P and S can exist in the gas or liquid phase.

[0055] FIG. 13 A shows a SEM image of microsized Pd ₃ P ₂ S ₈ crystals. The scale bar represents 10 μηι. The Pd ₃ P ₂ S ₈ flakes were first ground into fragments with size of micrometers, which were then used for the galvanostatic lithiation experiment. The SEM image suggests that the size of ground Pd ₃ P ₂ S8 flakes was less than 20 μηι. The plate-like feature of Pd ₃ P ₂ S ₈ was well preserved after grinding.

[0056] FIG. 13B shows a TEM image of microsized Pd ₃ P ₂ S ₈ crystals. The scale bar represents 20 nm. The plate-like feature of Pd ₃ P ₂ S ₈ was well preserved after grinding.

[0057] FIG. 13C shows the XRD pattern of microsized Pd ₃ P ₂ S ₈ crystals. The XRD pattern in FIG. 13C indicates that the crystal structure of microsized Pd ₃ P ₂ S ₈ flakes matches well with the simulated Pd ₃ P ₂ S ₈ structure.

[0058] FIG. 13D shows a HRTEM image of microsized Pd ₃ P ₂ S8 crystals. The scale bar represents 5 nm. The HRTEM image of ground Pd ₃ P ₂ S ₈ flakes shows a lattice spacing of 0.593 nm, corresponding to the (100) face of Pd ₃ P ₂ S ₈ crystal.

[0059] FIG. 14 shows the galvanostatic discharge curve for the lithiation of Pd ₃ P ₂ S ₈ using microsized Pd ₃ P ₂ S ₈ crystals as the cathode material and Li foil as the anode material. [0060] FIG. 15 shows the ultraviolet- visible (UV-vis) spectrum of Li-PPS NDs dispersed in water.

[0061] FIG. 16 shows the high resolution Pd 3d XPS spectrum of Pd ₃ P2S8 crystal and Li-PPS NDs. To further investigate the effects of lithiation process, the electronic structures of Pd ₃ P ₂ S ₈ crystal and Li-PPS NDs were compared by means of X-ray photoelectron spectroscopy (XPS). The Pd 3d core level XPS spectra of Pd ₃ P ₂ S ₈ crystal exhibited doublet peaks at around 336.4 eV and 341.6 eV (magenta peaks, see FIG. 16), corresponding to the Pd ²⁺ 3d ₅ / ₂ and Pd ²⁺ 3d ₃ / ₂ components of Pd ₃ P ₂ S ₈ , respectively. Once the Pd ₃ P ₂ S ₈ crystal was treated with lithium and transformed to Li- PPS NDs, besides the peaks of Pd ₃ P ₂ S ₈ crystal, two additional peaks appear at the high energy region (337.4 eV and 342.6 eV, cyan peaks) by deconvolution of the Pd 3d spectrum of Li-PPS NDs. Similarly, in the S 2p spectrum of Li-PPS NDs (FIG. 17A), additional peaks at 164.8 eV and 163.0 eV (green peaks) appear besides the known doublet peaks of S 2pm and S 2p ₃ / ₂ peaks at 163.4 eV and 162.0 eV (magenta peaks), respectively. The P 2p core-level XPS spectrum is shown in FIG. 17B, in which the spin-orbit doublet from 2Pi/ ₂ and 2P ₃ / ₂ peaks at 133.6 and 131.8 eV (magenta peaks), respectively, are visible. The XPS survey and energy-dispersive X- ray spectroscopy (EDS) spectra of Li-PPS NDs show intense signals of Pd, P, and S (FIG. 17C, FIG. 17D), indicating that all of the elements in Li-PPS NDs were preserved after lithiation of Pd ₃ P ₂ S ₈ crystal.

[0062] FIG. 17A shows a high resolution XPS spectra of S 2p of Li-PPS NDs.

[0063] FIG. 17B shows a high resolution XPS spectra of P 2p of Li-PPS NDs.

[0064] FIG. 17C shows a XPS survey spectrum of Li-PPS NDs.

[0065] FIG. 17D shows the EDS spectrum of Li-PPS NDs.

[0066] FIG. 18A shows the optimized atomic configurations for Li-incorporated Pd ₃ P ₂ S ₈ along the c axis based on DFT (density functional theory) calculations. The DFT calculations indicate that lithium atoms could intercalate into the interlayer of Pd ₃ P ₂ S ₈ and bind with two S( _a ) atoms and two S( _S ) atoms.

[0067] FIG. 18B shows the optimized atomic configurations for Li-incorporated Pd ₃ P ₂ S ₈ along the b axis based on DFT calculations. The DFT calculations indicate that lithium atoms could intercalate into the interlayer of Pd ₃ P ₂ Ss and bind with two S(a) atoms and two S( _S ) atoms. [0068] FIG. 19 shows the geometries and crystal field splitting diagrams of PdS ₄ . With the framework of crystal field theory, the two -orbitals of Pd ²⁺ close to S atoms in the PdS ₄ tetragon have higher energy, while the other three c/-orbitals have lower energy. The eight 4d-electrons of Pd ²⁺ fill in three low-energy orbitals and one high energy orbital. When sulphur vacancies or Li are introduced into the Pd ₃ P ₂ S ₈ crystal, the number of high energy 4d-electrons of Pd ²⁺ increases, giving rise to the elongated Pd-S bonds. On the contrary, when P vacancies are introduced into Pd ₃ P?S ₈ , the number of high energy electrons decreases and the Pd-S bonds become short. Also as shown in FIG. 5G, when sulphur vacancies or Li is introduced into the Pd ₃ P2S ₈ crystal, the Pd-S bond length is elongated, whereas the P vacancies decrease the Pd-S bonds.

[0069] FIG. 20A relates to the electrochemical double layer capacitance (Cdi) measurements for the Pd ₃ P ₂ S ₈ crystal. CV measurements were in the potential range of 0.048 to 0.248 V (vs. RHE) for the Pd ₃ P ₂ S ₈ crystal. The C _d i of catalyst-modified electrode was measured by CVs, and calculated according to the following equation: Cdi = JIdU/ 2 VvAU

[0070] where V is the geometric area of working electrode, v is the scan rate, and AU is the potential window. The factor of 2 is introduced in the equation, because one CV cycle includes the charge and discharge process, and the integral contains the transferred charges of these two processes.

[0071] FIG. 20B relates to the electrochemical double layer capacitance (Cdi) measurements for the Pd ₃ P ₂ S ₈ crystal. The plot shows the extraction of Cdi for the Pd ₃ P ₂ S ₈ crystal.

[0072] FIG. 20C relates to the electrochemical double layer capacitance (Cdi) measurements for the Li-PPS NDs. CV measurements were in the potential range of 0.048 to 0.248 V (vs. RHE) for Li-PPS NDs. The C _d i of catalyst-modified electrode was measured by CVs, and calculated according to the following equation:

C _di =JIdU/ 2VvAU

[0073] where V is the geometric area of working electrode, v is the scan rate, and AU is the potential window. The factor of 2 is introduced in the equation, because one CV cycle includes the charge and discharge process, and the integral contains the transferred charges of these two processes. [0074] FIG. 20D relates to the electrochemical double layer capacitance (Cdi) measurements for the Li-PPS NDs. The plot shows the extraction of Cdi for Li-PPS NDs.

[0075] FIG. 21 shows the electrochemical impedance spectroscopy (EIS) Nyquist plots for Li-PPS NDs, Pd ₃ P ₂ S ₈ crystal and Pd black. The inset of (A) shows the Nyquist plot at high-frequency range for the three catalysts. Z' is the real impendence and -Z" is the imaginary impedance. All Nyquist plots of the three catalysts consist of a semicircle in the high-frequency region (corresponding to charge transfer resistance, Ret) and a quasi-sloping line in the low-frequency region (corresponding to mass transfer resistance). The smaller i? _c t (diameter of the semicircle) value of the Li-PPS ND-modified GCE suggests its higher charge transport efficiency and faster HER kinetics. The equivalent circuit used in the electrochemical circle fit tool is shown in the inset of (B), where R _s is the value of the uncompensated resistance, CPE is the value of the argument of the constant phase element, and n is the value of the exponent of the constant phase element.

[0076] FIG. 22A shows the CV curves of Li-PPS NDs on rotating disk electrode (RDE) in 0 ₂ -saturated and N ₂ -saturated 0.1 M KOH.

[0077] FIG. 22B shows the linear sweep voltammetry (LSV) curves of Li-PPS NDs at different rotation rates in 0 ₂ -saturated 0.1 M KOH.

[0078] FIG. 22C shows the corresponding Koutecky-Levich plots (J ^-1 versus ω ^~ ° ⁵ ) at different potentials.

[0079] FIG. 22D shows the changes of n values as a function of working potentials.

[0080] FIG. 23 A shows the CV curves of 10% Pt/C modified RDE in 0 ₂ -saturated and N ₂ -saturated 0.1 M KOH.

[0081] FIG. 23B shows the LSV curves of the Li-PPS NDs and 10% Pt/C electrocatalyst modified RDE in 0 ₂ -saturated 0.1 M KOH solution at a rotation speed of l200 rpm.

[0082] FIG. 23C shows the chronoamperometric measurements of Li-PPS NDs and 10% Pt/C in 02-saturated 0.1 M KOH at 0.6 V (vs. RHE).

[0083] FIG. 23D shows the chronoamperometric measurements of Li-PPS NDs and 10%Pt/C at 0.6 V (vs. RHE) with a rotation speed of 1200 rpm.

[0084] FIG. 24A shows a SEM image of FePS ₃ crystals. [0085] FIG. 24B shows a SEM image of NiPS ₃ crystals.

[0086] FIG. 24C shows a SEM image of MnPS ₃ crystals.

[0087] FIG. 24D shows a SEM image of C0PS3 crystals.

[0088] FIG. 24E shows a SEM image of ZnPS ₃ crystals.

[0089] FIG. 24F shows a SEM image of Co ₂ PS ₃ crystals.

[0090] FIG. 25 shows the polarization curves of raw and lithium treated Co ₂ PS ₃ crystals modified GCE toward HER. The lithium treated Co ₂ PS ₃ crystals were derived from treatment with hexane solution containing n-butyllithium (2.5 M). The lithium treatment improves the catalytic performance of Co ₂ PS ₃ crystals.

[0091] FIG. 26 shows the polarization curves of FePS ₃ crystal and lithium treated FePS ₃ catalyst toward HER. The lithium treated FePS ₃ catalyst was derived from treatment with hexane solution containing n-butyllithium (2.5 M).

[0092] FIG. 27 shows the polarization curves of N1PS3 crystal and lithium treated NiPS ₃ catalyst toward HER.

[0093] FIG. 28 A shows the atomic models for hydrogen atoms adsorb at (1) S( _S ) site of pristine Pd ₃ P ₂ S ₈ (denoted as Pristine (H-S( _S ))), (2) S( _S ) site of Pd ₃ P ₂ S ₈ with Vs(a) (denoted as V _(a ) (H-S _(s) )), (3) S _(s) site of Pd ₃ P ₂ S ₈ with V _S ( _S ) (denoted as V _(s) (H-S _(s) )), (4) S(a) site of Pd ₃ P ₂ S ₈ with Vs( _S ) (denoted as V( _S ) (H-S( _a ))), (5) S( _a ) site of Li-adsorbed Pd ₃ P ₂ S ₈ (Li (H-S(a))) and (6) S _(s) site of Li-adsorbed Pd ₃ P ₂ S ₈ (denoted as Li (H-S _(s) )).

[0094] FIG. 28B shows the calculated free energy diagram for hydrogen evolution at equilibrium (U=0).

Detailed Description

[0095] The following detailed description refers to the accompanying drawings that show, by way of illustration, specific details and embodiments in which the invention may be practised.

[0096] Features that are described in the context of an embodiment may correspondingly be applicable to the same or similar features in the other embodiments. Features that are described in the context of an embodiment may correspondingly be applicable to the other embodiments, even if not explicitly described in these other embodiments. Furthermore, additions and/or combinations and/or alternatives as described for a feature in the context of an embodiment may correspondingly be applicable to the same or similar feature in the other embodiments.

[0097] In the present disclosure, the material structure engineering of layered transition metal phosphochalcogenides (e.g. Pd ₃ P ₂ S ₈ ) crystal through an electrochemical lithiation process is provided. The lithiation process transforms, for example, Pd ₃ P ₂ S ₈ crystalline particles to amorphous ultra-small lithium incorporated palladium phosphosulphide (Li-PPS) particles, which introduces sulphur vacancies or defects, and/or incorporate lithium into the Pd-P-S framework (where the incorporated lithium has a strong interaction with Pd, and the incorporated lithium atoms may be intercalated into the interlay er of Pd ₃ P ₂ S ₈ crystalline particles, and may be bound to one or more sulphur atoms). In other words, the present method involves treating the transition metal phosphochalcogenides crystalline particles with lithium to obtain amorphous lithium treated transition metal phosphochalcogenides particles. The crystalline particles may be called crystals in the present disclosure. The particles or amorphous particles obtained after lithiation of the crystalline particles may be referred to as nanodots (NDs) in the present disclosure, where the nanodots may have sizes of about 2 run to 10 nm. Advantageously, such structure tuning activates the Pd ₃ P ₂ S ₈ for electrocatalytic HER, and the Li-PPS NDs demonstrate unprecedented HER activity and outstanding long-term stability. The activation of the electrocatalytic activity of Pd ₃ P ₂ S ₈ may be attributed to its morphology and structure changes (e.g. crystallinity, sulphur vacancy and Li incorporation) induced by the lithiation process. As defined herein, the term "amorphous" refers to a material which is assembled into a state that exhibits no readily perceptible organization, regularity, or orientation of its constituent elements. Meanwhile, the term "crystalline" refers to a material which is assembled into a state that exhibits readily perceptible organization, regularity, or orientation of its constituent elements. Amorphous materials also lack the long-range of repeated structures formed by atoms, ions or molecules in a crystalline material, may have vacancies, and different bond lengths, which are not characteristic of crystalline materials.

[0098] Conventionally, the most efficient HER electrocatalyst is Pt. However, the scarcity and high price of Pt limit its commercial application. One alternative way is to replace Pt with other cheaper metals/compounds and reduce its consumption. In the present disclosure, the catalytic activity of Pd ₃ P ₂ S8 for electrochemically HER is activated. The price of Pd is much lower than Pt, which makes the present electrocatalyst economically viable for commercialization.

[0099] As mentioned above, in order to activate and/or enhance the catalytic performance of inorganic layered materials, structure engineering has been mainly focused on the following aspects: (i) decrease of their thickness, (ii) reduction of their lateral sizes, (iii) creation of the sulphur vacancies/defects, (iv) phase engineering, and (v) doping or alloying. However, single structure engineering processes are typically too weak for tuning their properties from the pristine states, and it may be difficult to simultaneously realize most, if not all, of the structural engineering aspects as described in a single material. In the present method and electrocatalyst, most of the above structure engineering aspects are simultaneously realized for Pd-P-S framework. The ability to tweak multiple structure engineering (size, shape, defects, doping and crystallinity) at the same time amplifies the material's properties for electrocatalysis and to boost the material's catalytic performance. In this regard, the physicochemical property of an amorphous material can differ greatly from its crystalline counterpart. However, excluding the preparation of amorphous molybdenum (or tungsten) sulphide by chemical/electrochemical method, only a handful of works have sought to study other layered materials. In the present disclosure, the crystalline engineering of Pd ₃ P ₂ S ₈ , and other transition metal phosphochalcogenides, have been utilized to convert such transition metal phosphochalcogenides from a crystal to amorphous structure.

[00100] Apart from being able to activate transition metal phosphochalcogenides for electrocatalysis and to convert it from a crystalline to an amorphous material, the present method also has a production yield, e.g. approximately 100% of Li-PPS NDs.

[00101] Due to the facile preparation procedures (the lithiation process can be finished in several hours) and the high production yield, the present method is easily more advantages for synthesizing the highly efficient HER catalyst.

[00102] In another consideration, conventional nanomaterials synthesized by chemical methods are usually covered by organic ligands, which significantly limit their electron transfer efficiency and applications. In contrast, the present method and electrocatalyst circumvent this issue as no organic ligands are required and hence no organic ligands are formed on the surface of the resultant amorphous NDs.

[00103] The obtained NDs of the present disclosure are highly dispersable in water with long-term stability (e.g. more than 1 month). The obtained NDs also have a different structure from the conventional nanomaterials, which can be distinguished via their X-ray diffraction (XRD) patterns. The obtained nanodots may be smaller in size relative to the conventional nanomaterials.

[00104] In a further consideration, an efficient electrocatalyst usually requires good electrical conductivity. The incorporation of lithium into the transition metal phosphochalcogenide (e.g. Pd-S-P) framework provides the feasibility to regulate electrical property of the transition metal phosphochalcogenide (e.g. Pd ₃ P ₂ S ₈ ) crystals.

The conductivity of the lithium treated transition metal phosphochalcogenide NDs

(e.g. Li-PPS NDs) is much higher than its crystal form (Pd ₃ P ₂ S ₈ ).

[00105] The resultant electrocatalyst' s performance for hydrogen evolution reaction was tested and unexpectedly found to be quite efficient with an onset potential of 52 mV and a Tafel slope of 29 mV/dec, which are better than conventional HER electrocatalysts.

[00106] The good solubility of the transition metal phosphochalcogenide NDs (e.g. Li-PPS) enables it to be mixed with other organic/inorganic materials for the requirements of various applications. They can be easily hybridized with other materials such as metal organic frameworks (MOFs), organic polymers etc.

[00107] The present disclosure therefore describes a method for producing an amorphous lithium treated transition metal phosphochalcogenide electrocatalyst. The present disclosure also describes an amorphous lithium treated transition metal phosphochalcogenide electrocatalyst. The present disclosure further describes the use of such an amorphous lithium treated transition metal phosphochalcogenide electrocatalyst for hydrogen evolution reaction (HER) and/or oxygen reduction reaction (ORR).

[00108] The present disclosure also describes another method for producing a lithium treated transition metal phosphochalcogenide electrocatalyst and such a lithium treated transition metal phosphochalcogenide electrocatalyst. Such a method and an electrocatalyst differ from the earlier method and electrocatalyst in that such a lithium treated transition metal phosphochalcogenide is not amorphous but may be crystalline.

[00109] Having outlined various advantages of the present electrocatalyst and method, definitions of certain terms are first discussed before going into details of the various embodiments.

[00110] The term "chalcogenide" refers to an element from group 16 of the chemical periodic table. Non-limiting examples include oxygen, sulphur and selenium. The term "phosphochalcogenide" refers to a chalcogenide with phosphorus.

[00111] The word "substantially" does not exclude "completely" e.g. a composition which is "substantially free" from Y may be completely free from Y. Where necessary, the word "substantially" may be omitted from the definition of the invention.

[00112] In the context of various embodiments, the articles "a", "an" and "the" as used with regard to a feature or element include a reference to one or more of the features or elements.

[00113] In the context of various embodiments, the term "about" or "approximately" as applied to a numeric value encompasses the exact value and a reasonable variance.

[00114] As used herein, the term "and/or" includes any and all combinations of one or more of the associated listed items.

[00115] As used herein, the phrase of the form of "at least one of A and B" may include A or B or both A and B. Correspondingly, the phrase of the form of "at least one of A and B and C", or including further listed items, may include any and all combinations of one or more of the associated listed items.

[00116] Unless specified otherwise, the terms "comprising" and "comprise", and grammatical variants thereof, are intended to represent "open" or "inclusive" language such that they include recited elements but also permit inclusion of additional, unrecited elements.

[00117] Having defined the various terms as mentioned above, details of the various embodiments are now described below.

[00118] In the present disclosure, there is provided for a method for producing an amorphous lithium treated transition metal phosphochalcogenide electrocatalyst, comprising: constructing a cathode comprising transition metal phosphochalcogenide crystalline particles; configuring the cathode, a lithium anode and a lithium-based electrolyte to form a battery; subjecting the battery to galvanostatic discharge to transform the transition metal phosphochalcogenide crystalline particles into amorphous lithium treated transition metal phosphochalcogenide particles; and extracting the amorphous lithium treated transition metal phosphochalcogenide particles from the cathode for use as the amorphous lithium treated transition metal phosphochalcogenide electrocatalyst.

[00119] The method as disclosed above specifies two electrodes (i.e. anode and cathode) with lithium ion battery construction.

[00120] The present method may, generally, involve two steps. The first step lies with the synthesis of the transition metal phosphochalcogenide crystalline particles. The second step may involve lithiation, i.e. treating the transition metal phosphochalcogenide crystalline particles with lithium.

[00121] The transition metal phosphochalcogenide crystalline particles may be prepared by a chemical vapour transport (CVT) method as described in the examples. The CVT method allows for formation of pure crystalline solid particles, where their crystal structure can be determined by diffraction methods. Compared to solution based methods, there is more control for the CVT method and hence composition of crystalline particles resulting from CVT can be controlled more easily. In various embodiments, the present method may involve a CVT step for providing the transition metal phosphochalcogenide crystalline particles. Such a step may comprise forming the transition metal phosphochalcogenide crystalline material in a reactor comprising a first reaction zone and a second reaction zone. The temperature of the first reaction zone may be the same as or different from the temperature of the second reaction zone throughout the CVT step.

[00122] In some embodiments, the forming may comprise a first step of heating a transition metal, a chalcogenide, phosphorus and a transporting agent in the first reaction zone while maintaining the second reaction zone at a higher temperature than the first reaction zone. The first step of heating in the first reaction zone may occur at 650°C to 850°C while maintaining the second reaction zone at the higher temperature of 750°C to 900°C for 48 hours or less. The forming may comprise a second step of heating after the first step of heating, wherein the second step of heating may comprise switching the first reaction zone to a higher temperature than the second reaction zone. In some embodiments, the first reaction zone may be switched to the higher temperature of 750°C to 1050°C for 5 days or less. The transporting agent may comprise halogen such as bromine, chlorine, iodine etc. The halogen can exist as a halide in a halogen compound, for example, hydrogen halides such as hydrogen chloride, hydrogen bromide and hydrogen iodide. The transporting agent may comprise a bromide, a chloride or an iodide according to various embodiments. Iodine is commonly used as the transporting agent for sulphides (as for selenides and tellurides) as it exhibits the same transport behaviour for the metal and chalcogen constituents of a transition metal phosphochalcodenide. In some embodiments, the forming may comprise cooling the reactor after the second step of heating to form the transition metal phosphochalcogenide crystalline particles (i.e. crystals).

[00123] Once the transition metal phosphochalcogenide crystalline particles are obtained, they may be subjected to a lithiation or lithium treatment process to transform the crystalline particles into amorphous materials according to some embodiments. The amorphous materials may be in the form of an amorphous layer comprising the amorphous nanodots, or may be in the form of amorphous nanodots.

[00124] According to some embodiments, in the lithiation process of the present method, the transition metal phosphochalcogenide crystalline particles may be used to construct an electrode for assembling into a battery, for the lithiation process. The electrode may be a cathode. The present lithiation process, which utilizes a battery for lithiation, is advantageous as it keeps the lithium treatment time short, and the extent of lithiation can be controlled by adjusting a cut-off voltage. Accordingly, constructing the cathode may comprise mixing the transition metal phosphochalcogenide crystalline particles with a conductive filler and a binder before contacting the transition metal phosphochalcogenide crystalline particles with an organic solvent to form a mixture. The mixture may then be deposited onto a conductive substrate to form the electrode, e.g. the cathode, which may be dried before use as a component for forming the battery.

[00125] The conductive filler may comprise or consist of acetylene black, carbon black or carbon nanotube. The use of a conductive filler enhances the transport efficiency of conductive ions. For example, during the charge-discharge process, the electrode materials may aggregate. If there is no conductive filler, the charge- discharge process may be stopped or disrupted by aggregation of the electrode materials. Other suitable conductive filler may be used. The binder may comprise or consist of polyvinylidene fluoride (PVDF), polyvinyl alcohol (PVA), polytetrafluoroethylene (PTFE) or carboxymethyl cellulose (CMC). Other suitable binder may be used. The organic solvent may comprise or consist of N- methylpyrrolidone (NMP) or dimethylformamide (DMF). Other suitable organic solvent may be used.

[00126] In some embodiments, the constructing of the cathode may comprise depositing the mixture onto a conductive substrate. The constructing of the cathode may comprise drying the mixture on the conductive substrate in vacuum for at least 24 hours at 80°C to 100°C. The conductive substrate may comprise a carbon paper, copper foil disc, nickel foil disc or zinc foil disc. The conductive substrate, according to various embodiments, refers to any suitable substrate for use in constructing an electrode, e.g. the cathode, and conducts electricity.

[00127] According to the present method, the cathode formed using the transition metal phosphorus crystalline particles may be used as one of the components in forming a battery. Accordingly, in some embodiments, the configuring step (of the cathode) of the present method may comprise disposing or assembling the cathode and the anode in the lithium-based electrolyte in an inert environment and at a temperature of 20°C to 40°C. Non-limiting example of an inert environment may be to assemble the battery in a glove box filled with argon. Other inert gases, apart from argon, such as nitrogen gas may be used. The inert environment prevents explosion of lithium based materials during assembly of the battery. The lithium based electrolyte may comprise LiPF ₆ , LiBF ₄ or LiC10 ₄ . The lithium based electrolyte may be formed by dissolving LiPF ₆ , LiBF ₄ or LiC10 ₄ in a mixture of organic solvents. For example, 1 M of LiPF ₆ may be dissolved in a mixture of ethyl carbonate (EC) and dimethyl carbonate (DMC), where the EC to DMC ratio may be 1 : 1. Other lithium based compounds and organic solvents suitable for forming the lithium based electrolyte may be used.

[00128] According to the present method, subjecting the battery to galvanostatic discharge may comprise discharging a current of 0.02 to 0.03 mA, 0.025 mA etc. This drives the lithium in the lithium-based electrolyte into the crystalline particles of the cathode. The treating of the cathode with lithium then converts the crystallline particles into its amorphous form. Subjecting the battery to galvanostatic discharge may also comprise discharging at a cut-off voltage of 0.9 V. Once this is carried out, according to the present method, the amorphous lithium treated transition metal phosphochalcogenide particles (i.e. NDs) may be extracted. The extracting of the amorphous lithium treated transition metal phosphochalcogenide particles may comprise washing the amorphous lithium treated transition metal phosphochalcogenide particles with water or an alcohol under sonication after removing the cathode from the battery. The alcohol may be any suitable alcohol such as methanol, ethanol etc. The extracting may include or comprise centrifuging the amorphous lithium treated transition metal phosphochalcogenide nanoparticles at 2500 rpm to 4000 rpm (e.g. 3000 rpm) for at least 10 minutes after the washing of the amorphous lithium treated transition metal phosphochalcogenide particles.

[00129] The present method, in some embodiments, provides for a method for producing a lithium treated transition metal phosphochalcogenide electrocatalyst, comprising: contacting transition metal phosphochalcogenide crystalline particles with an organolithium solution to form an organolithium suspension, wherein the organolithium solution comprises an organolithium; and allowing the organolithium suspension to stand for a period of time during which the organolithium dissociates to provide lithium for penetrating into each of the transition metal phosphochalcogenide crystalline particles to form the lithium treated transition metal phosphochalcogenide electrocatalyst. The lithium treated transition metal phosphochalcogenide produced according to this method is not amorphous but may be crystalline. Such a lithium treated transition metal phosphochalcogenide may comprise or consist of a mixture of nanodots and nanosheets. The nanosheets may be about 10 wt% of the mixture or less. According to such embodiments, a lithium treated transition metal phosphochalcogenide electrocatalyst obtained according to such a method is provided.

[00130] Such embodiments differ from the above in that no construction of a battery is required, the lithium becomes incorporated into the transition metal phosphochalcogenide crystalline particles by contacting the latter with an organolithium. [00131] In such embodiments, the organolithium suspension may be prepared by immersing transition metal phosphochalcogenide crystalline particles (e.g. in their powder form) into a solution containing the organolithium (i.e. organolithium solution). The organolithium solution may comprise an organolithium. The organolithium solution may consist of an organolithium and an organic solution. Non- limiting examples of the organic solution include pentane, hexane etc.

[00132] In such embodiments, the organolithium suspension may be allowed to stand for a period of time of at least 1 day, for example, 3 days. This provides time for the lithium to dissociate from the organolithium, and to penetrate and become incorporated into the transition metal phosphochalcogenide crystalline particles, thereby transforming the latter into the lithium treated transition metal phosphochalcogenide.

[00133] Such embodiments may further comprise mixing the organolithium suspension with water after forming the lithium treated transition metal phosphochalcogenide electrocatalyst. Such embodiments may further comprise sonicating the organolithium suspension at a temperature of 0°C to 5°C. The sonication process helps to reduce the size and thickness of the resultant lithium treated transition metal phosphochalcogenide, which is beneficial for catalysis.

[00134] The organolithium in such embodiments may be an alkyllithium. Non- limiting examples of alkyllithium may include isopropyllithium, n-butyllithium and tert-butyllithium.

[00135] The present disclosure also provides for an amorphous lithium treated transition metal phosphochalcogenide electrocatalyst comprising a transition metal phosphochalcogenide treated with lithium to form an amorphous lithium treated transition metal phosphochalcogenide, wherein the amorphous lithium treated transition metal phosphochalcogenide comprises one or more sulphur atom vacant sites and/or one or more phosphorus atom vacant sites.

[00136] Various embodiments of the present method, and advantages associated with various embodiments of the present method, as described above may be applicable to the present electrocatalyst, and vice versa.

[00137] In various embodiments, the amorphous lithium treated transition metal phosphochalcogenide may be formed as nanodots. Nanodots, in the context of the present disclosure, may refer to small sized nanoparticles of 2 nm to 10 nm. Due to such significant small sizes, the nanodots may be treated as zero dimensional materials unlike conventional three dimensional particles (e.g. microparticles or nanoparticles) or two dimensional nanosheets. This is in contrast to conventional lithiation of layered crystals that leads to crystallized two dimensional nanosheets. The nanodots also form an amorphous arrangement. Accordingly, lithiation of Pd ₃ P ₂ S8 crystals leads to small sized amorphous "zero dimensional" Li-PPS nanodots.

[00138] As nanodots of the present disclosure are significantly smaller than conventional lithium treated nanoparticles, that is, lithiated nanosheets, the nanodots have properties that differ from lithiated nanosheets. For example, nanodots may have better solubility, more tunability in physiochemical properties, and better amenability to structural/compositional transformation and functionalization with other nanomaterials. Particularly, the rich edge sites of nanodots can function as chemically active sites for many applications, such as catalysis and bio-applications.

[00139] The use of the term "nanodot", in this disclosure, may be used to refer to the morphology of a material while the phrase "layered structure" may be used to describe the structure of a material.

[00140] In various embodiments, the transition metal phosphochalcogenide of the resultant electrocatalyst may comprise at least one chalcogenide. The transition metal phosphochalcogenide of the resultant electrocatalyst may comprise a transition metal, phosphorus, and at least one chalcogenide. The transition metal may be selected from the group consisting of cobalt, iron, manganese, nickel, palladium and zinc. The at least one chalcogenide may comprise or consist of sulphur and/or selenium.

[00141] In the resultant electrocatalyst, the lithium may be directly attached to one or more atoms of the at least one chalcogenide. The lithium may also be intercalated between and/or in the layers of the amorphous lithium treated transition metal phosphochalcogenide. Further, the lithium may have both of these characteristics (i.e. the direct attachment and the intercalation).

[00142] The present disclosure further provides for the use of an amorphous lithium treated transition metal phosphochalcogenide electrocatalyst, and the use of a lithium treated transition metal phosphochalcogenide electrocatalyst, as disclosed above, for hydrogen evolution reaction and/or oxygen reduction reaction. [00143] Various embodiments of the present method and present electrocatalyst, and advantages associated with various embodiments of the present method and present electrocatalyst, as described above may be applicable to the use of the present electrocatalyst, and vice versa.

[00144] While the methods described above are illustrated and described as a series of steps or events, it will be appreciated that any ordering of such steps or events are not to be interpreted in a limiting sense. For example, some steps may occur in different orders and/or concurrently with other steps or events apart from those illustrated and/or described herein. In addition, not all illustrated steps may be required to implement one or more aspects or embodiments described herein. Also, one or more of the steps depicted herein may be carried out in one or more separate acts and/or phases.

Examples

[00145] The present disclosure relates to a method of producing an amorphous lithium treated transition metal phosphochalcogenide electrocatalyst. The present disclosure also relates to such an amorphous lithium treated transition metal phosphochalcogenide electrocatalyst. Such an electrocatalyst may be composed of an amorphous material containing lithium treated transition metal phosphochalcogenide nanoparticles. The present electrocatalyst may, in such instances, be composed of amorphous lithium treated transition metal phosphochalcogenide nanoparticles. Such nanoparticles may be called nanodots (NDs) in the present disclosure. Such nanoparticles, which may be amorphous in nature, are derivable from crystalline particles in the present disclosure. The crystalline particles may be called crystals in the present disclosure. In some instances, the present disclosure further relates to a method of producing a lithium treated transition metal phosphochalcogenide electrocatalyst. In such instances, the present electrocatalyst may be composed of a lithium treated transition metal phosphochalcogenide that is not amorphous in nature. Such lithium treated transition metal phosphochalcogenide may be in the form of crystalline nanodots.

[00146] Referring to bulk Pd ₃ P ₂ S ₈ crystal, which was first discovered in 1971, specific experiments on its applications are rarely, if not at all, reported till date, and

[00150] With activation, the morphology of bulk Pd ₃ P2S ₈ crystals is changed from crystalline to amorphous Li-PPS NDs through the present electrochemical tuning process (see FIG. 3A to FIG. 3E). The obtained Li-PPS NDs can be highly dispersed in water with negative surface charges (see FIG. 4A and FIG. 4B).

[00151] The present electrochemical turning process introduces sulphur vacancies in Pd ₃ P ₂ S ₈ , and lithium into the Pd-P-S framework after lithium treatment (FIG. 5A to FIG. 5G and FIG. 6A to FIG. 6B). After lithium treatment, the electrical conductivity of Li-PPS NDs is dramatically enhanced compared to bulk Pd ₃ P ₂ S ₈ crystals (FIG. 7).

[00152] The Li-PPS NDs exhibit excellent electrocatalytic activity toward HER in acidic electrolyte with onset potential as low as 52 mV and Tafel slope of 29 mV/dec (FIG. 8A and FIG. 8B). The onset potential of Li-PPS NDs for HER is better than most metal sulphide and metal phosphide HER electrocatalysts (FIG. 9A and 9B). In the Tafel slope approach, the current Li-PPS NDs have better Tafel slope compared to conventional platinum-carbon (Pt/C) catalyst and other non-Pt catalysts (FIG. 10 and FIG. 9B).

[00153] In addition to the high electrocatalytic activity, the Li-PPS NDs exhibit remarkable durability throughout the electrochemical operation. As shown in FIG. 11 A, the polarization curve of Li-PPS NDs almost overlays with the initial one with negligible loss of cathodic current after 7500 potential cycles, suggesting superior durability and stability under testing condition.

[00154] In the present method, activating the intrinsic electrocatalytically inert layered materials that possess no or very poor intrinsic catalytic activity is a promising way to develop new catalysts. However, as mentioned above, the intrinsic electrocatalytically inert materials have been neglected, thereby omitting potential electrocatalysts. In the present disclosure, activation of electrocatalytically inert Pd ₃ P2S ₈ crystals renders it highly efficient for electrocatalysis.

[00155] The present method and the present electrocatalyst are described in the examples below.

[00156] Example la: Materials (Chemicals)

[00157] Palladium powder (99.9%), phosphorus powder (99.9%), sulphur powder (99.99%), iodide (99.99%), acetylene black, polyvinylidene fluoride (PVDF), and N- methylpyrrolidone (NMP) were purchased from Sigma- Aldrich (Germany). Lithium and copper foils were purchased from ACME Research Support Pte. Ltd. (Singapore). Lithium-ion battery electrolyte, i.e. 1 M of LiPF ₆ dissolved in a mixture of ethyl carbonate (EC) and dimethyl carbonate (DMC) (volume ratio of 1 : 1), was purchased from Charslton Technologies Pte. Ltd. (Singapore). Ethanol (99.9%) and acetone (Tech Grade) were purchased from Merck (Germany). All chemicals and materials were used as received without further purification. The Milli-Q water used in experiments was obtained from a Milli-Q System (USA).

[00158] High purity elemental powders of other transition metals, phosphorus and sulphur were used as starting materials throughout the synthesis of every ternary transition metal phosphosulphide crystals. Iron powder (Aldrich, >99%), nickel powder (Alfa Aesar, -325 mesh, 99.8%), manganese powder (Alfa Aesar, -325 mesh, 99.3%), cobalt powder (Alfa Aesar, 99.9%), zinc powder (Alfa Aesar, 99.8%), phosphorus lump, red (Alfa Aesar, 99.999%, Puratronic ^® ) and sulphur pieces (Alfa Aesar, 99.999%, Puratronic ^® ) were used as purchased without further purification for the synthesis of respective compounds.

[00159] Example lb: Synthesis of Transition Metal Phosphochalcogenide Crystals

[00160] The transition metal phosphochalcogenide crystalline material or crystals were obtained via vapour phase growth using chemical vapour transport (CVT) technique. Stoichiometric amounts of transition metal (Ni, Fe, Mn, Co, Zn etc.) along with phosphorus and sulphur were sealed in quartz ampoules with an internal pressure in the range of 10 ^"5 to lO ^"6 Torr (0.0133 to 0.00133 Pa). In addition, iodine (2 mg/cm ³ ) was generally incorporated inside the quartz ampoule as a transport agent unless otherwise specified. Iodine was sealed inside a small capillary to avoid contamination of the vacuum pump. The sealed tubes were then placed in a two-zone horizontal tube furnace (see FIG. 12).

[00161] Based on the above, in one example, single crystals of Pd ₃ P ₂ S8 were prepared by CVT method using iodide as transporting agent (see FIG. 12). The stoichiometric amounts of Pd:P:S used was 1 :1 :3 based on a total weight of 400 mg (i.e. 80 mg Pd powder, 80 mg P powder and 240 mg S powder), mixed with additional 35 mg iodide (or 2 mg/cm ³ ). This mixture was sealed in an evacuated 20 cm long quartz tube under vacuum of 10 ^"6 Torr (0.00133 Pa). The sealed tube was placed in a two-zone furnace and pretreated at 850°C for 30 hours while the growth zone (i.e. second reaction zone) was kept at 900°C to prevent transportation of samples. This also means the first reaction zone was heated at 850°C. The temperature of the first reaction zone was then programmed to 1030°C while the growth zone was kept at 900°C, for five days, in order to generate the temperature gradient at the region where growth of the single crystals occur. Finally, the furnace was naturally cooled down to room temperature and red single crystals of Pd ₃ P ₂ S ₈ were collected in the growth zone. The first reaction zone may be called a source zone.

[00162] In another example, the source zone (first reaction zone) was set at 650°C and the growth zone (second reaction zone) was set at 750 ^° C for 48 hours. This allows the constituents (e.g. transition metal, phosphorus, chalcogenide) to react completely and prevent the back transport or formation of undesired additional phases. After this, the source zone was gradually increased to 750°C and the growth zone was lowered to 700°C. This continued for another 120 hours, after which the temperatures of both zones were lowered to room temperature, and the ampoules were taken out to retrieve the transition metal phosphochalcogenide crystals for characterization and subsequent processes (e.g. lithiation).

[00163] In another example, Co ₃ P ₂ S3 crystals were synthesized from cobalt, phosphorus and sulphur. Stoichiometric amounts of these starting materials were sealed in an evacuated quartz tube (ampoule). Initially, temperature of the furnace was raised to 350°C for 12 hours and later raised to 800 ^° C. This continued for another 336 hours. Polycrystalline compounds of Co ₃ P ₂ S ₃ were then retrieved and obtained after lowering temperature of the furnace to room temperature. The longer duration is critical for achieving pure crystalline phase and complete reaction between the starting materials.

[00164] Example 2; Lithiation of the Transition Metal Phosphochalcogenide Crystals

[00165] In one example, the lithium treated or lithium incorporated palladium phosphosulphide nanodots (Li-PPS NDs), prepared from bulk Pd ₃ P ₂ S ₈ crystals, was based on the present electrochemical lithiation method.

[00166] Briefly, the Pd ₃ P ₂ S ₈ crystals were ground to micrometers and mixed with acetylene black and PVDF binder, in a mass ratio of 8: 1 :1, dispersed in N-Methyl-2- pyrrolidone (NMP). After magnetic stirring for 24 hours, the homogeneously mixed slurry was then uniformly coated on Cu foil discs and dried in vacuum for 24 hours at 90°C. The as-prepared electrodes were then used as cathode for assembling battery in an argon filled glove box at room temperature using Li foil as anode and 1 M LiPF ₆ as electrolyte. The galvanostatic discharge process was performed in a Neware battery test system at a discharge current of 0.025 mA. An example of this setup is shown in FIG. 3 A. After finishing the discharge process, the electrodes were taken out from the battery cell and washed with acetone to remove any residual LiPF ₆ electrolyte. Then the material was dispersed in 10 ml distilled water (or ethanol) by ultrasonication in ice water. The suspension was then centrifuged at 3000 rpm for 15 minutes, the supernatant (nanodots) was collected for further characterization and application.

[00167] In another example, the lithiation of FePS ₃ and N1PS3 were realized by chemical intercalation and exfoliation process using n-butyllithium as the intercalate. Briefly, 50 mg sample powders (i.e. crystals) were immersed in 2.5 M n-butyllithium (in hexane) solution operated in a glove box. After 3 days, the dispersion was washed with 100 mL hexane. The dispersion was then taken out from the glove box and dispersed in 20 mL water and sonicated for 30 minutes in ice water. After centrifugation (1000 rpm, 15 minutes), the supernatant was collected for further characterization, and for device fabrication.

[00168] Example 3a: Characterization

[00169] The single crystal X-ray diffraction (XRD) data were collected at room temperature from Rigaku SCXmini with Mercury CCD automatic diffractometer equipped with a graphite-monochromated Mo Ka radiation (λ = 0.71073 A). The data were corrected for Lorentz factors, polarization, air absorption, and absorption due to variations in the path length through the detector faceplate. The SADABS program was used for the absorption correction based on a multi-scan technique. The space group P3mi (No. 164) was determined according to the systematic absences, E-value statistics, and subsequent successful refinements of the crystal structure. All structures were solved by the direct methods and refined by full-matrix least-squares fitting on F ² by SHELX-97.

[00170] The Li-PPS NDs suspension was dropped onto an ultrathin carbon coated holey copper grid for transmission electron microscopy (TEM) characterization. The same suspension was dropped onto clean Si substrates for scanning electron microscopy (SEM), X-ray photoelectron spectroscopy (XPS) and nuclear reaction analysis (NRA) characterizations. The same suspension was dropped onto a clean Si0 ₂ /Si substrate for atomic force microscopy (AFM) characterization and onto a carbon paper for X-ray absorption spectrum (XAS) measurement.

[00171] TEM images and energy dispersive X-ray spectroscopy (EDS) were recorded with JEOL-2100F at an acceleration voltage of 200 kV. SEM images were obtained using a field emission scanning electron microscope (SEM, JEOL JSM-7600F). Powder XRD was conducted using a Siemens D-500 X-ray diffractometer (Bruker AXS, Inc., Madison, USA). Optical microscopy images were taken on a Nikon Eclipse LV100D microscope. AFM (AFM, Cypher, Asylum Research, USA) was used to characterize the Li-PPS NDs in tapping mode in air. Ultraviolet-visible (UV- vis) absorption spectrum was recorded on a UV-2700 (Shimadzu) with QS-grade quartz cuvettes (1 1 1 -QS, Hellma Analytics) at room temperature. Dynamic light scattering (DLS) was performed at room temperature using Zetasizer Nano ZS (Malvern, England). XPS measurements were performed using Kratos Axis-ULTRA X-ray photoelectron spectroscopy instrument equipped with a monochromatic Al Ka (1486.7 eV) X-ray source with emission of 10 mA and anode HT (high tension) of 15 KV. The sample analysis chamber pressure was about 10 ^"9 Torr (1.33 x 10 ^"6 Pa) during the spectrum acquisition. In the current-voltage (I-V) characterization, Au electrodes with the same separation distances were fabricated on the Pd ₃ P ₂ S ₈ crystal and Li-PPS NDs film on Si0 ₂ by thermal evaporation through shadow masks. The I-V characteristics were measured under ambient conditions by using a Keithley 4200 semiconductor characterization system with a shielded probe station.

[00172] XAS of Pd K-edge was recorded using the BL14W1 beamline at the Shanghai Synchrotron Radiation Facility (SSRF) in China, operated at 3.5 GeV with beam current of 300 mA. A Si (311) double crystal was used to monochromatize the X rays from the storage ring. Pd foil was used as reference for the energy calibration, and all samples were measured in transmission mode. The X-ray absorption near edge structure (XANES) measurements of S K-edge were carried out at the XAFCA beamline of Singapore Synchrotron Light Source. The fitting of the Pd-S shell for the Pd ₃ P ₂ S ₈ crystal was carried out using conventional curve fitting by Winxas, and the atom distribution function (ADF) method was applied to investigate the structural features of Li-PPS NDs.

[00173] NRA was carried out at the 3.5 MeV Singletron accelerator at CIBA. The ⁷ Li(p,a) reaction was used because it provided an essentially background-free Li- signal at energies around 8 MeV. After checking the system calibration using a thick LiNbCb standard sample, the NRA data were collected with a 2 MeV proton beam using perpendicular incidence at scattering angle of 144°. The resulting spectra were then fitted using the simulation code SIMNRA.

[00174] The nanobeam electron diffraction (NBED) patterns were collected at an operation voltage of 300 kV using a FEI Titan ³ cubed 60-300 electron microscope and fitted with a Schottky field emission source and double correctors. Different from the incoherent electron beam used in conventional convergent beam electron diffraction, a nearly parallel coherent electron beam produced by a small condenser aperture with a diameter of 10 μηι was used as a nanoprobe. By using this aperture, a 3 nm-diameter electron beam was produced. Therefore, an individual nanodot could be located under conventional scanning transmission electron microscopy (STEM) mode. The convergence angle was approximate 0.7 mrad. The NBED patterns were then acquired around the nanodots by using a scanning function of the STEM system. The simulation of NBED pattern was carried out using the multislice method.

[00175] Example 3b: Electrochemical Measurements

[00176] Hydrogen evolution reaction (HER) measurements were performed on the Autolab workstation (PGSTAT12) using a conventional three-electrode system with a graphite rod as the counter electrode, Ag/AgCl (3 M KC1) and catalyst-modified glassy carbon electrode (GCE) were used as the reference electrode and work electrode, respectively. Ag/AgCl electrode was calibrated with respect to a reversible hydrogen electrode (RHE). The working electrode was prepared by casting 5 μΐ, Li- PPS ND suspension (4 mg mL ^'1 ) prepared in a mixed solvent (ethanol and water, volume ratio of 1 : 1) onto a prepolished GCE (3 mm in diameter). The catalyst- modified GCE was dried at room temperature and then 5 xL nation ethanolic solution (0.5 wt%) was dropped on its surface to protect the catalyst. By using the same procedure, the Pd ₃ P ₂ S8 crystal, Pt/C and Pd black were used to modify GCEs. The loading amount of noble metals of all catalysts on GCE were kept same. The HER test was carried out in H ₂ saturated 0.5 M H ₂ S0 ₄ aqueous solution. Linear sweep voltammetry (LSV) was conducted at a scan rate of 5 mV s ^_1 . The stability test was carried out by taking continuous potential cycling in the potential window of -0.152 to 0.248 V (vs. RHE) at a scan rate of 100 mV s ^"1 . After potential cycling, the LSV curve was recorded at a scan rate of 5 mV s ^_I .The LSV curves were corrected for ohmic potential drop (iR) losses. To evaluate the electrochemical active surface area (ECSA), the cyclic voltammetry (CV) was conducted from 0.048 to 0.248 V (vs. RHE) in N ₂ saturated 0.5 M H ₂ S0 ₄ aqueous solution at different scan rates, i.e. 10, 20, 40, 80, 120 mV s ^"1 for Li-PPS NDs, and 40, 80, 120, 160, 200, 240 mV s ^"1 for the Pd ₃ P ₂ S ₈ crystal. Electrochemical impedance spectroscopy (EIS) was performed at - 0.102 V (vs. RHE) with frequency from 0.01 to 100,000 Hz and amplitude of 10 mV.

[00177] Example 3c: Computational Details

[00178] The density functional theory (DFT) method was performed to predict the atomic models of the NDs. In one example, the DFT method was applied to Li-PPS NDs. The generalized gradient approximation (GGA) of Perdew-Burke-Ernzerhof (PBE) functional was performed using the projector augmented wave method as implemented in the Vienna Ab-initio Simulation Package (VASP). Van der Waals corrections to the total energy were included using the DFT-D2 method. Plane wave basis set with an energy cutoff of 400 eV was employed, and the Brillouin zone was sampled with a 4 4x3 Monkhorst-Pack k-point mesh for the Pd ₃ P ₂ S ₈ unit cell. The structures were fully relaxed by conjugate gradient algorithms until the maximum force on each atom was less than 0.01 eV/A, and the convergence criterion of self- consistent calculations was 10 ^"6 eV. All the simulated structures were screened by calculation of the K-edge XANES spectra of the Pd atoms using FEFF program. The structure exhibiting identical XANES spectra with the experimental one was chosen to represent the structure of Li-PPS NDs.

[00179] The DFT calculations were performed to investigate the HER activity by using Vienna ab initio simulation package. Intermediate states for HER of adsorbed hydrogen atoms on Pd ₃ P ₂ S ₈ crystal were modelled by creating 2x2x 1 supercells of Pd ₃ P ₂ S ₈ with an optimized lattice constant of 6.879 A with relaxing all the atomic positions, together with a vacuum region of thickness larger than 15 A. A kinetic energy cutoff of 400 eV and a 4 4x 1 Monkhorst-Pack grid were adopted. The maximum force exerted on each atom was relaxed to less than 0.005 eV/A.

[00180] The free energy for atomic hydrogen adsorption (AGH) was calculated according to equation (1):

AGH = AEH + AEZPE - TASR (1) where AEk is the hydrogen chemisorption energy, AEzPE is the correction of zero point energy, ASn is the entropy difference between the adsorbed state and the gas state, and T is temperature being assumed to 300 K. The AEu was calculated according to equation (2):

AEu = (EPd3P2S8+«H - EPd3P2S8 - «E(H ₂ )/2) x \ln (2) where Epd3P2S8+«H is the total energy of the Pd3P2Ss system with n hydrogen atoms adsorbed on the surface, Epd3P2S8 is the total energy of Pd ₃ P2Ss, and £(¾) represents the energy for a ¾ molecule in the gas phase.

[00181] Here, a single hydrogen atom adsorbed in a 2 2 1 Pd ₃ P ₂ S ₈ supercell was considered. AEzPE = £ΖΡΕ(Η*) - £ΖΡΕ(Η2)/2, where £ZPE(H*) can be obtained by calculating the vibrational frequency associated with the adsorbed hydrogen atom (H*), and £ΖΡΕ(Η2) is the zero point energy of hydrogen molecule (0.27 eV). Calculations showed that the vibrational frequency of H* is nearly site independent, and the EzPE was calculated to be 0.02 eV. For the ASn which is defined as the difference of the entropy between adsorbed H* and H ₂ gas molecule, it is well accepted that the configuration entropy of H* is small and can be neglected. Therefore, AS = S(H ₂ )/2, where S(H2) is the entropy of hydrogen in the gas phase at standard condition. Based on the aforementioned calculations, AGH can be expressed as: AGH = AEu + 0.22 eV.

[00182] The AGH on LiPd ₃ PS ₇ was calculated using a similar process with that of Pd ₃ P2S ₈ . Based on the creation of slab model for simulating the (001) surface of LiPd3PS ₇ , various H adsorption sites were examined and two different sulphur sites with different coordination were identified to be highly HER active.

[00183] Example 4: Results and Discussion - Crystal Growth and its Characterization

[00184] In this discussion, the Pd3P2S ₈ crystal prepared by the present CVT technique (see FIG. 12 and example lb) is referred to as a non-limiting example. [00185] The single crystal XRD analysis revealed that the Pd ₃ P ₂ S ₈ crystalline material crystallized in a trigonal symmetry with the space group of P3 mi (No. 164; a = b = 6.8420 A, c = 7.2466 λ, α = β = 90°, y = 120° and V= 259 A3 A ³ ). Its structure parameters are listed in table 1 above. The powder XRD pattern was consistent with the simulated pattern of single-crystal Pd ₃ P ₂ S ₈ (FIG. 1C), indicating the pure phase of the obtained Pd3P ₂ S ₈ crystals.

[00186] As shown in FIG. 1A, Pd ₃ P ₂ S8 possesses an interlocked layered structure with weak interlayer Van der Waals interactions. Unlike graphene and ternary metal dichalcogenides (TMDs), Pd ₃ P ₂ S ₈ has waved Pd-P-S layers, which are composed of PdS ₄ tetragon and PS ₄ tetrahedron (FIG. IB). Importantly, there are two non- equivalent sulphur sites, i.e. the singly-coordinated sulphur at the apical position (S( _a )) and the three fold-coordinated sulphur at the ring side position (S( _S >) (FIG. IB). Each Pd atom is surrounded by four S( _S ) atoms in a square-planar configuration to form a PdS ₄ tetragon while each P atom occupies the center of a tetrahedron formed by one S(a) and three S( _S ) atoms (FIG. IB). The S( _S ) atoms in PdS ₄ tetragon are shared and connected together with the PS ₄ tetrahedron, forming the wavy layer (FIG. 1A). The PdS ₄ square is regular with Pd-S=2.342 A (FIG. IB, table 2), which is consistent with that of PdS ₂ , while the PS ₄ unit features two types of non-equivalent P-S bonds with bond lengths of P-S _(a )= 1.907 and P-S _(s )=2.11 1 A. The PS tetrahedron is slightly distorted (FIG. IB, table 2), in which the S-P-S angle changes from 97.57(9)° to 1 19.71(7)° as compared to the ideal one. The atomic coordinates and the equivalent isotropic displacement parameters for the Pd ₃ P ₂ S ₈ crystal are listed in table 3 below.

[00187] Table 3 - Atomic Coordinates and Equivalent Isotropic Displacement Parameters of Single Crystal Pd ₃ P ₂ S ₈ .

aocc is defined as the atom site occupancy ^h U _eq is defined as one-third of the trace of the orthogonalized t/jj tensor.

[00188] FIG. 2A shows the synthesized Pd ₃ P ₂ S ₈ crystals with size up to 1 cm. The well defined orientation of Pd ₃ P ₂ S ₈ crystal is reflected in its plate-like shape (FIG. 2B and FIG. 2C). In addition, the SEM image of FIG. 2D clearly shows the layered structure of Pd ₃ P ₂ S ₈ . As shown in the high resolution transmission electron microscopy (HRTEM) image (FIG. 2E), the as-grown Pd ₃ P ₂ S ₈ single crystals showed a high degree of crystallinity with no obvious defects. The lattice spacing of 0.593 nm corresponds to the (100) face of Pd ₃ P ₂ S ₈ .

[00189] Example 5: Results and Discussion - The Lithiation Process of the Present Method

[00190] Referring to the Pd ₃ P ₂ S ₈ crystal as an example, the engineering of Pd ₃ P ₂ S ₈ crystal was carried out by a controllable galvanostatic lithiation process (FIG. 3A). In a typical experiment, the crystalline Pd ₃ P ₂ S ₈ flake was first ground into fragments of micrometer size (FIG. 13 A to FIG. 13D), which were then assembled into a columnar lithium-ion battery cell as the cathode, where the lithium foil served as the anode (FIG. 3A, example 2). The galvanostatic discharge curve was recorded to control the lithiation process. As shown in FIG. 14, the voltage dropped monotonically from 2.95 to 1.87 V before reaching a well defined plateau at 1.82 V, during which the Li ⁺ ions intercalated into the Pd ₃ P ₂ S ₈ crystal and induced the structure change of Pd ₃ P ₂ S ₈ . A cut-off voltage of 1.70 V was selected to stop the reaction (FIG. 14). After completion of the lithiation process, the intermediate product was washed in water (or ethanol) by sonication.

[00191] Example 6: Effects of Lithiation on Crystal Structure and Electronic State

[00192] Following through with the same non-limiting example described in example 5, once the delithiation process was completed, the Pd ₃ P ₂ S ₈ crystal was disintegrated and transformed to ultra-small Li-PPS NDs with high production yield (nearly 100%). The TEM image indicated that the NDs with lateral size of 2.5±0.9 nm were obtained (left image of FIG. 3B and left image of FIG. 3D). The DLS result revealed a narrow size distribution of Li-PPS NDs (FIG. 4A). Surprisingly, different from exfoliated TMD nanosheets and NDs with organized crystal structures, the obtained ultra-small Li-PPS NDs of the present disclosure were amorphous with disordered crystal lattice orientations (right image of FIG. 3B, and FIG. 3E). The selected area electron diffraction (SAED) pattern gives a weak signal (inset in FIG. 3E), further indicating the loss of crystallinity of Li-PPS NDs. The AFM result gives the topographic morphology of Li-PPS NDs (left image of FIG. 3C). The statistical analysis of AFM data provides the height distribution profile with average thickness of 1.2 ± 0.8 nm (right image of FIG. 3C and right image of FIG. 3D). The morphology sketch of a Li- PPS ND is presented in the inset of left image of FIG. 3B. The obtained Li-PPS NDs were negatively charged with surface potential of about -40 mV in water at pH 7.0 (FIG. 4B), so they can be well dispersed in water and stable for more than one month. The aqueous solution of Li-PPS NDs showed a dark brown color (inset in right image of FIG. 3B) and exhibited a broad absorption band in the ultraviolet- visible region (FIG. 15).

[00193] The X-ray absorption spectrum (XAS) was conducted to investigate the lithiation effect toward to the atomic structure change of Pd ₃ P ₂ S ₈ . The obtained Pd K- edge XANES spectra for the Pd ₃ P ₂ S ₈ crystal and Li-PPS NDs showed similar results (FIG. 5A), indicating their similar Pd valence state and crystal environment. The Fourier Transform (FT) analysis for the Pd K-edgo, extended X-ray absorption fine structure (EXAFS) was then used to verify their distinct local structures. As shown in FIG. 5B, the FT spectrum of Pd ₃ P ₂ S ₈ crystal centered at 1.86 A corresponds to the nearest Pd-S coordination. A similar peak associated with the Pd-S coordination in the Li-PPS NDs was also observed. However, compared to the Pd ₃ P ₂ S ₈ crystal, the Pd-S peak intensity obtained from Li-PPS NDs was weak and this peak position shifts to a small bond length (R) of about 0.03 A. The aforementioned difference suggested that the Pd-S shell of Li-PPS NDs was distorted and its structural homogeneity was reduced due to the lithiation process. Moreover, the peak centered at 3.27 A in Pd ₃ P ₂ S ₈ crystal was almost invisible for the Li-PPS NDs, which is likely due to the small size of Li-PPS NDs. To obtain the quantitative structural parameters around the Pd atoms, the curve fitting of Pd-S shell was performed (FIG. 6A). The obtained structural parameters around Pd atoms for the Pd ₃ P ₂ S ₈ crystal and Li-PPS NDs are listed in table 4 below.

[00194] Table 4 - The Structure Parameters Around Pd Atoms in the Pd ₃ P ₂ S ₈ Crystal and Li-PPS NDs.

Sample Bond R (A) CN σ ² (10- ³ Α ² )

Pd-S 2.34 (±0.02) 4* 3.8 (±0.1) crystal

beam for NBED is as small as 0.7 mrad to form a nearly parallel electron probe, which allowed location and characterization of an individual ND (FIG. 5C). The representative patterns are shown in FIG. 5D, in which the variable distances matched well with those from the XAFS characterization (table 4).

[00196] To understand the unusual Pd-S distortion in Li-PPS NDs, the S K-edge XANES spectra of the Pd ₃ P ₂ S ₈ crystal and Li-PPS NDs were recorded (FIG. 6C). The absorption edge at about 2470 eV, attributed to the characteristic of S ^2" , appears in both Pd ₃ P ₂ S ₈ crystal and Li-PPS NDs. However, another absorption edge at about 2482 eV, corresponding to S ⁶⁺ , showed notable enhancement only in the Li-PPS NDs. The electronic structures of Li-PPS NDs were also characterized by means of XPS (FIG. 16 and FIG. 17A to FIG. 17D). In the S 2p core level XPS spectrum of Li-PPS NDs (FIG. 17A), a broad peak at around 168.3 eV appeared, indicating the formation of S ⁶⁺ after the lithium treatment, which was consistent with the XANES result (FIG. 6C). In addition, in the P 2p core-level XPS spectrum, additional peaks appear at high energy range (136.6 eV and 134.1 eV, FIG. 17B), corresponding to the phosphate. The S ⁶⁺ and phosphate were likely attributed to the sulphates and phosphate absorbed on the surface of Li-PPS NDs, respectively. It means some sulphur and phosphor segments were dissociated from the bulk material, resulting in the formation of sulfur vacancies (Vs) and phosphor vacancies (Vp) in the Li-PPS NDs. The density functional theory (DFT) calculations showed that the lithium atoms could form chemical bonds with the sulphur atoms (FIG. 18A and FIG. 18B). Thus the subsequent delithiation process in water could induce the loss of sulphide and phosphor segments and create a large number of vacancies (V, including Vs and Vp). In addition, the calculation results showed that the introduction of vacancies (V) led to strong distortion of Pd-S bond length (R) in the Pd ₃ P ₂ S ₈ crystal (FIG. 19).

[00197] Moreover, the lithium signal in Li-PPS NDs has been detected by NRA, as shown in FIG. 5E. As compared with the Pd ₃ P ₂ S ₈ crystal, the NRA spectrum of Li- PPS NDs exhibited stronger Li signal in the high energy region between 4.8 and 8.8 MeV. The lithium content was determined to be 10 atomic percent (at%) by fitting the Li profile of NRA (FIG. 5E and FIG. 6D), suggesting that there was still lithium trapped in the Li-PPS NDs even after the product was delithiated in water. Since there were Li and V in Li-PPS NDs, the structural distortion of Li-PPS NDs can be attributed to the combined effects of the V formation and Li incorporation. In order to understand the structure of Li-PPS NDs, the structure model of lithium-incorporated palladium phosphosulphide was built based on the DFT calculations. All the simulated structures were screened by calculation of the K-edge XANES spectra of Pd atoms. It was found that the XANES pattern of the four-coordinated Pd atoms in the structure model with formula of LiPd ₃ PS (as shown in FIG. 5F) reproduced well with that of Li-PPS NDs (FIG. 6E). Based on the DFT simulation results, the crystal data and Pd-S bond length (R) distribution of the Pd ₃ P ₂ Sg crystal, Pd ₃ P ₂ Sg with Li incorporation, Pd ₃ P ₂ S ₈ with vacancies at S( _a ) sites (Vs(a)) and vacancies at S( _S ) sites (Vs(s)), and LiPd ₃ PS ₇ are summarized in table 5 below, and in FIG. 5G, respectively.

[00198] Table 5 - Structure Data of Pd ₃ P ₂ S ₈ crystal, Simulated Pd ₃ P ₂ S ₈ with

P (deg.) 90.000 91.819 90.000 94.006 92.143 γ (deg.) 120.000 120.470 120.000 1 18.244 127.992

[00199] As shown in FIG. 5G, when Vs (including Vs(a) and Vs( _S )) or Li was introduced into the Pd3P ₂ S ₈ crystal, the Pd-S bond length elongated, whereas the P vacancies decreased the Pd-S bonds (FIG. 19). The LiPd ₃ PS ₇ has variable Pd-S bond lengths (R) with elongated or shorten Pd-S bond length as compared to Pd ₃ P ₂ S8, consistent with the EXAFS fitting results (FIG. 5B, table 4). Based on the aforementioned results, the LiPd ₃ PS ₇ could be used as the representative structure of Li-PPS NDs for the subsequent theoretical calculations.

[00200] As known, an efficient electrocatalyst usually requires good electrical conductivity. The Li incorporation into the Pd-S-P framework provides the feasibility to regulate the electrical property of Pd ₃ P ₂ S ₈ crystal. Therefore, the effect of lithiation process on the electrical conductivity was investigated. As shown in FIG. 7, the electrical conductivity of Li-PPS NDs was much higher compared to the Pd3P ₂ S8 crystal, indicating that the electrical conductivity of Pd ₃ P ₂ S ₈ crystal have been dramatically improved after the lithiation treatment.

[00201] Example 7: Effect of Lithiation on HER Performance

[00202] As a proof-of-concept application, the HER catalytic property of Li-PPS NDs was evaluated and compared with the Pd ₃ P2S ₈ crystal, commercial Pd black, and Pt/C catalyst (10% Pt on activated charcoal). As shown in FIG. 8A, the reductive sweep of Li-PPS ND catalyst began at exceptionally low potential with an onset potential value of -52 mV, smaller than that of the Pd ₃ P ₂ S ₈ crystal (-175 mV) and Pd black (-92 mV) (FIG. 8A). Impressively, this exceptionally low onset potential was among the best results obtained in the metal sulphide-based and metal phosphide-based HER catalysts (FIG. 9A and FIG. 9B), suggesting the outstanding electrocatalytic activity of the synthesized Li-PPS NDs toward HER. As shown in FIG. 8B, the Li-PPS ND catalyst exhibited substantially higher activity with a Tafel slope of 29 mV dec ^"1 compared to Pd ₃ P ₂ S ₈ crystal (123 mV dec ¹ ) and Pd black (107 mV dec ^"1 ), which was close to that of the conventional Pt/C (28 mV dec ^"1 ). Conventionally, except for the Pt-based metal catalysts, Tafel slope with less than 30 mV dec ^"1 are rarely reported. The Tafel slope of Li-PPS NDs (29 mV dec ^"1 ) obtained here was the lowest among all the metal sulphide-based or metal phosphide-based HER catalysts (FIG. 9B and FIG. 10). Such a small Tafel slope suggests a two-electron transfer process through the Volmer-Tafel mechanism toward HER, in which recombination of chemisorbed hydrogen atoms on the Li-PPS ND surface is the rate-limiting step in hydrogen evolution. The fast HER kinetics enable the Li-PPS ND catalyst to reach a high current density at a low overpotential, i.e. 10 mA cm ^"2 at -126 mV, 20 mA cm ^"2 at -139 mV, and 50 mA cm ^"2 at -169 mV.

[00203] As discussed above, the Li-PPS NDs exhibit excellent HER performance whilst the Pd ₃ P ₂ S ₈ crystal shows negligible HER catalytic activity (FIG. 8 A). To better understand the lithiation effect toward the electrocatalytic activity, the structure and property changes of Li-PPS NDs were systematically evaluated, including the reduced size, V formation, Li incorporation and enhanced conductivity, toward the HER activity.

[00204] First, the electrochemical active surface areas (ECSA) of the Li-PPS NDs and Pd ₃ P ₂ S ₈ crystal were compared by calculating the electrochemical double-layer capacitance (Cdi), which was linearly proportional to the effective active surface area. As shown in FIG. 20A to FIG. 20D, the Li-PPS NDs possess C _d i of 2.5 mF cm ^"2 , which was much higher than that of Pd ₃ P ₂ S ₈ crystal with Cdi of 12.2 μΡ cm ^"2 . The large Cdi of Li-PPS NDs indicated that the lithiation of Pd ₃ P ₂ S ₈ crystal gave rise to a high ECSA of Li-PPS NDs, increasing the active sites for producing hydrogen via electrocatalysis.

[00205] To uncover the effect of V and Li species on the HER activity, the first- principle calculations were performed to investigate the energetics of hydrogen adsorption. The free energy for atomic hydrogen adsorption (AGR), which is a reasonable descriptor for the ability of hydrogen evolution, was evaluated. It is known that an active catalysis for HER must comply with the thermoneutral requirement, that is Δ(¾~0, where atomic hydrogen is bound neither too strongly nor too weakly to allow the concomitant efficient hydrogen adsorption and release. As shown in FIG. 8C, FIG. 8D, the lowest energetic configuration for Pd ₃ P ₂ S ₈ crystal, i.e. hydrogen atom adsorbing at the S _s site, gave AGu of 1.15 eV, indicating the Pd ₃ P ₂ S ₈ crystal was essentially inert with a large thermodynamically uphill step of the hydrogen adsorption. However, the AGu could be greatly decreased once the Pd ₃ P ₂ S ₈ crystal transformed to Li-PPS NDs. As mentioned above, the LiPd ₃ PS ₇ structural model could be used to present the Li-PPS NDs (FIG. 5G). It was found that the surface of distorted LiPd ₃ PS ₇ could significantly promote the adsorption of hydrogen atoms. The A H was reduced from 1.15 eV for the Pd ₃ P ₂ S ₈ crystal [FIG. 8C-(1): left image of FIG. 8C, FIG. 8D-(1): solid line at top] to -0.09 eV for LiPd ₃ PS ₇ when H atom binds at the S ₂ atom (four coordination) at the surface [FIG. 8C-(2): right image of FIG. 8C, FIG. 8D-(2)]. Furthermore, the free energy nearly reached thermoneutral [AGH-0.01 eV, FIG. 8C-(2), FIG. 8D-(2)], which was comparable to Pt, when H atom bound at the Si atom (three coordination). The nearly thermoneutral free energy indicated that the H atoms were prone to bind with both of the sulphur sites, giving rise to high activity of HER. Since the HER on the Li-PPS ND surface proceeded via the Volmer- Tafel mechanism, the large amount of active S atoms could chemisorb the hydrogen atoms and promote recombination of hydrogen atoms on its surface. The easy recombination of hydrogen atoms efficiently reduced the activation energy for hydrogen evolution and thus decreased the overpotential for HER. A more in depth investigation of AGH on pristine Pd ₃ P ₂ S ₈ (FIG. 28A-(1)) and Pd ₃ P ₂ S ₈ monolayer with Vs (FIG. 28A-(2), FIG. 28A-(3), FIG. 28A-(4)) (based on first-principle calculation), and evaluation of Vs effect toward the HER activity are as follow.

[00206] Due to the weak interlayer interaction of Pd ₃ P ₂ S ₈ , the calculations were based on the monolayer Pd ₃ P ₂ S ₈ . Here, after various binding sites for hydrogen atom on the pristine Pd ₃ P ₂ S ₈ were calculated, it was observed that the hydrogen atom did not bind to the sheet and the adsorption of hydrogen was endothermic. The AGH on the pristine Pd ₃ P ₂ S ₈ with the lowest energetic configuration, i.e. hydrogen atom adsorbed at the S _s site, referred to as pristine Pd ₃ P ₂ S ₈ (H-S _s ), was 1.15 eV (FIG. 28A- (1), FIG. 28B-(1)), suggesting that the pristine Pd ₃ P ₂ S ₈ was essentially inert with a large thermodynamically uphill step of the hydrogen adsorption. However, its catalytic activity could be activated by introducing Vs. There are two different kinds of Vs in Li-PPS NDs, i.e. Vs(a) and Vs( _S ). Based on the calculation results in FIG. 28B, the AGH of Pd ₃ P ₂ S ₈ withVs(a) (FIG. 28A-(2) and FIG. 28B-(2)) was higher than those of Pd ₃ P ₂ S ₈ with Vs(s) (FIG. 28A-(3), FIG. 28A-(4)), indicating that the V _S (a) was less effective in improving the catalytic activity of Pd ₃ P ₂ S ₈ . In particular, when the hydrogen atom anchors at the nearest S( _S ) atom of Vs(s) (FIG. 28A-(3)), the free energy decreased to nearly thermo-neutral (AGH=0.05 eV, FIG. 28B-(3)), which was comparable to platinum, indicating the comparable catalytic activity. Based on the aforementioned discussion, Vs( _S ) specifically leads to a significantly promoted catalytic activity even though the formation of the less-coordinated Vs(a) is energetically favoured with about 0.9 eV lower formation energy compared to that of Vs(s). In addition, when hydrogen atom adsorbs on the near S( _a > atom (FIG. 28A-(4)), the created Vs (H-S( _a )) is also effective for decreasing AGu, producing a moderate adsorption energy (AGH=0.22 eV), as shown in FIG. 28B-(4). Regarding the effect of Li incorporation, the AGu was investigated by examining different anchoring sites of Li atoms in the pristine Pd ₃ P ₂ S ₈ sheet. It was found that once the Li atoms formed a coplanar structure with three S atoms (FIG. 28A-(5), FIG. 28A-(6)), improved HER activities with AGu of -0.14 eV and -0.21 eV occured when the hydrogen atoms bonded to the S _(a ) (FIG. 28A-(5), FIG. 28B-(5)) and S _(s) atoms (FIG. 28A-(6), FIG. 28B-(6)), respectively. The low AGu indicated that Li incorporation could also directly modulate the HER activity of Pd ₃ P ₂ S8. Therefore, the calculation results revealed that both of the vacancy formation and Li incorporation in pristine Pd ₃ P ₂ S ₈ provide for positive effects in activating the HER performance of Pd ₃ P ₂ S ₈ .

[00207] The EIS for Li-PPS NDs and Pd ₃ P ₂ S ₈ crystal were recorded and compared under the HER process. As shown in FIG. 21 , the Nyquist plot of Li-PPS NDs has apparently smaller charge transfer resistance (R _c t) than that of Pd ₃ P ₂ S ₈ crystal, and even smaller than that of Pd black, indicating the ultrafast Faradaic process and thus a superior HER kinetics of Li-PPS NDs. This small R _c t could be attributed to the excellent conductivity and unique structure of Li-PPS NDs. As shown in FIG. 7, the excellent conductivity of Li-PPS NDs could accelerate the electron transfer during the electrocatalytic reaction and thus resulted in the high response current toward HER. Moreover, the ultra-small Li-PPS NDs with amorphous structure and large amount of edge sites could accelerate the electron transport from the Li-PPS NDs to the GCE surface, significantly decreasing the resistance during the HER process.

[00208] In addition to the high electrocatalytic activity, the Li-PPS NDs also exhibited remarkable durability during HER. The long-term cycles were performed by taking continuous cyclic voltammetry between -0.152 and 0.248 V (vs. RHE) in H ₂ - saturated 0.5 M H ₂ S0 ₄ aqueous solution. As shown in FIG. 11 A, the polarization curve of Li-PPS NDs for HER after 7,500 potential cycles was similar to the initial one with negligible loss of cathodic current, suggesting the super durability and stability under the testing condition.

[00209] The outstanding HER performance of Li-PPS NDs inspired exploring its feasibility for water splitting. A prototype electrolyzer cell was fabricated at room temperature with 0.5 M H ₂ S0 ₄ aqueous solution as the electrolyte. The Li-PPS NDs and commercial Ir0 ₂ were loaded onto carbon paper (1 cm ^χ 1 cm) as the cathode and anode, respectively. Because of the remarkably small overpotential of Li-PPS NDs (-52 mV) for hydrogen evolution, this water electrolyzer cell can be powered simply using two 1.5 V AAA batteries arranged in series (nominal voltage of about 3.0 V; FIG. 1 IB). All the aforementioned results demonstrated that the present Li-PPS NDs derived from the Pd ₃ P ₂ S ₈ crystal can be used as an excellent electrocatalyst for practical water splitting under low overpotential.

[00210] Example 8: Application of Lithium Incorporated Transition Metal Phosphochalcogenide Nanodots for Oxygen Reduction Reaction

[00211] Apart from using the present NDs as an electrocatalyst for HER, it can also be used for oxygen reduction reaction (ORR). Hence, the present lithium incorporated transition metal phosphochalcogenide nanodots may be a multifunctional electrocatalyst.

[00212] ORR measurements were performed on the Autolab electrochemical workstation (PGSTAT12) using a conventional three-electrode system with Pt wire and Ag/AgCl (3.0 M KCl) as the counter and reference electrodes, respectively. A glassy carbon rotating disk electrode (RDE) (3 mm in diameter) was used as the working electrodes. The catalyst suspension was prepared by dispersing 5 of the catalyst ink (4 mg mL ^'1 ) onto the prepolished RDE. Then, 5 of 0.5% Nation ethanol solution was dropped on the surface of catalyst modified RDE to protect the catalyst. The cyclic voltammogram (CV) and Linear sweep voltammograms (LSV) for the ORR were obtained using a RDE, corrected by iR-compensation in 0 ₂ - saturated 0.1 M KOH solution at a scan rate of 5 mV s ^"1 . A Pt/C commercial catalyst having 10 wt% Pt was used as a reference for evaluation of the electrocatalytic performance of Li-PPS in ORR. [00213] As shown in FIG. 22A, the CV curve of Li-PPS shows an onset potential of about -0.10 V (vs. Ag/AgCl) and peak potential of -0.28 V (vs. Ag/AgCl), which were comparable to 0 V and -0.14 V (vs. RHE) of conventional Pt/C catalyst (FIG. 23 A, FIG. 23 B), respectively, indicating the excellent electrocatalytic activity of Li-PPS. To gain further insight into the electron transfer kinetics during ORR, the behaviours in the rotating disk voltammetry at different rotating rates were tested. The rotating disk electrode (RDE) measurement of Li-PPS showed the increase of current density with the rotation rate (FIG. 22B), which was due to the enhanced diffusion of electrolyte at higher rotation rate. The electron transfer number of Li-PPS in the ORR process was calculated to be about 3.5 at the potential range from -0.8 to -0.4 V (vs. Ag/AgCl) based on the Koutecky-Levich plots derived from the RDE measurement (FIG. 22C, FIG. 22D).

[00214] Moreover, the electrocatalytic stability of Li-PPS toward ORR was examined by the chronoamperometric measurement. As shown in FIG. 23C, the current density of Li-PPS-modified GCE decays more slowly than that of conventional Pt/C-modified GCE, demonstrating the excellent electrocatalytic stability of Li-PPS for ORR in the alkaline aqueous solution. Moreover, the resistance to crossover effect is another important factor to determine the catalytic performance in fuel cells. The resistance to crossover effect of CH ₃ OH for the Li-PPS was also tested and compared with commercial Pt/C catalysts by using the chronoamperometric measurement. As shown in FIG. 23D, after CH ₃ OH was injected into the 0 ₂ -saturated 0.1 M KOH aqueous solution, the current response of commercial Pt/C-modified GCE decreased obviously, while there was no obvious current decrease for the Li-PPS-modified GCE, indicating that Li-PPS has a remarkable tolerance to the crossover effect of methanol.

[00215] Example 9: SEM Characterizations of Various Transition Metal Phosphochalcogenide Crystals

[00216] The SEM images of the as-obtained ternary metal phosphorsulphide crystals of FePS ₃ , NiPS ₃ , MnPS ₃ , C0PS3, ZnPS ₃ and Co ₂ PS ₃ are shown in FIG. 24A to FIG. 24F, respectively.

[00217] Example 10: Various Transition Metal Phosphochalcogenide and Their Lithium Treated Samples for Use As HER Electrocatalyst [00218] Raw and lithium treated C02PS3 crystalline particles were also demonstrated as an electrocatalyst for HER. As shown in FIG. 25, the raw Co ₂ PS3 crystals showed a low onset potential value of -50 mV and low j\ 0 of -127 mV, both of which were smaller than most of the reported state-of-art metal sulphide or phosphide HER catalysts. Moreover, after lithium treatment, the catalytic activity toward HER is further improved. As C02PS3 crystal is attractive for commercialization due to its low price compared to catalysts containing noble metal elements and the lithium treatment improves the catalytic performance of C02PS3 crystals, lithiated C02PS3 crystalline particles are therefore suitable for HER electrocatalyst.

[00219] FePS3 and N1PS3, and their corresponding lithium-treated materials were tested for their suitability for HER. After lithium treatment, both of the lithium treated FePS3 and N1PS3 showed enhanced performance toward HER (FIG. 26 and FIG. 27).

[00220] Example 11: Commercial and Potential Applications

[00221] Via the structure engineering strategy of the present disclosure based on the current lithiation process, the Pd ₃ P ₂ S8 crystal was transformed into amorphous Li-PPS NDs. The atomic stmcture features, including the vacancy formation and lithium incorporation, were investigated in detail as described above. The unique structure and superior electrical conductivity of the amorphous Li-PPS NDs resulted in an unexpectedly high electrocatalytic HER performance in acidic solution, which is superior to conventional metal sulphide-based and metal phosphide-based catalysts. The current structure engineering strategy therefore provides a new avenue to tune the crystal structures and prepare materials with unique properties and enhance performance, which can be used for various promising applications.

[00222] The method and Li-PPS NDs of the present disclosure have also been demonstrated for their applications in hydrogen evolution. The facile fabrication process, high production yield and the method being environmentally sustainable, make it appealing for industrial usage. The overpotential of Li-PPS NDs is as low as 52 mV and the Tafel slope is 29 mV/dec. This makes Li-PPS a promising and highly efficient HER catalyst. The above results showed that the Li-PPS NDs exhibit significantly improved HER performance compared to Pd3P ₂ S8 crystals. The outstanding HER performance of Li-PPS NDs makes it feasible for water splitting. Based on this, a prototype electrolyzer cell was fabricated at room temperature. As shown in FIG. 1 IB, the water electrolysis cell can be powered by two AAA batteries in series with a nominal voltage of about 3.0 V. The results demonstrated that Li-PPS NDs derived from Pd ₃ P ₂ Ss crystals can be an excellent catalyst for water splitting.

[00223] The method of the present disclosure offers opportunities for structure engineering of nanomaterials, with the introduction of active sites, engineering crystallinity of materials, and commercialization of highly efficient HER catalysts.

[00224] The Li-PPS NDs can also be used as building blocks for integration with other two dimensional nanomaterials (graphene, MoS ₂ , WS ₂ etc.), which may provide further advantages for high performance solar cells and fuel cells.

[00225] The present method is also demonstrated for other transition metal phosphochalcogenide crystals. Such transition metal phosphochalcogenide crystals were lithiated and demonstrated for ORR, apart from HER.

[00226] Two non-limiting examples include Li-PPS nanodots (Lithium treated Pd ₃ P ₂ S ₈ ) and lithium treated Co ₂ PS ₃ crystals. Both of these materials showed excellent HER performance, which were comparable to conventional Pt/C catalyst. Other crystals include FePS ₃ , NiPS ₃ , MnPS ₃ , CoPS ₃ , ZnPS ₃ etc. The lithiation process of the present method is easy and completable in several hours. The obtained nanodots, e.g. Li-PPS NDs, are highly dispersed in water with long-term stability (more than 1 month). Moreover, the production of amorphous Li-PPS nanodots from Pd ₃ P ₂ S ₈ crystals has a high yield and quantity, which is attractive for commercialization into mass production.

[00227] The onset potential of Li-PPS nanodots for HER is much lower than most metal sulphide and metal phosphide HER electrocatalysts. The Tafel slope approach was also much lower than that of the commercial Pt/C catalyst, and is the best among non-Pt catalyst. The catalytic performance of Li-PPS nanodots is superior or at least comparable to conventional Pt/C catalyst, which makes it an alternative to Pt/C catalyst.

[00228] Co ₂ PS ₃ also showed excellent catalytic performance that was comparable with that of Pt/C catalyst. Most importantly, there is no noble metal in the present electrocatalyst and hence its price remains low.

[00229] Currently, the most efficient HER electrocatalyst seems to be Pt. However, the scarcity and high price of Pt limit its large scale production and commercial application. One alternative way is to replace Pt with other cheaper metals/compounds and reduce its consumption. In the present disclosure, the catalytic activity of Pd ₃ P ₂ S ₈ for electrochemically HER was activated by the present method. The price of Pd is lower than Pt, which makes it potentially attractive over the latter. The price comparison of Pt, Pd, and Pd ₃ P ₂ S ₈ is listed in table 6 below.

[00230] Table 6 - Price Comparison of Pt, Pd, and Pd ₃ P ₂ S ₈

[00231] For materials like Co ₂ PS ₃ , FePS ₃ , NiPS ₃ , MnPS ₃ , CoPS ₃ , ZnPS ₃ , FePSe ₃ , FePS _{0 5} Se _{2 5} , FePSSe ₂ , FePS ₂ Se ₂ , the central metal atoms are all non-noble metals. Thus, these materials have much lower price than Pt/C and Pd ₃ P ₂ S8.

[00232] While the invention has been particularly shown and described with reference to specific embodiments, it should be understood by those skilled in the art that various changes in form and detail may be made therein without departing from the spirit and scope of the invention as defined by the appended claims. The scope of the invention is thus indicated by the appended claims and all changes which come within the meaning and range of equivalency of the claims are therefore intended to be embraced.