Copper-Substituted NiTiO3 Ilmenite-Type Materials for Oxygen

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Functional Inorganic Materials and Devices

Copper substituted NiTiO3 Ilmenite type Materials for Oxygen Evolution Reaction amandine guiet, Tran Ngoc Huan, Christophe Payen, Florence Porcher, Victor Mougel, Marc Fontecave, and Gwenaël CORBEL ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b08535 • Publication Date (Web): 05 Aug 2019 Downloaded from pubs.acs.org on August 8, 2019

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ACS Applied Materials & Interfaces

Copper substituted NiTiO3 Ilmenite type Materials for Oxygen Evolution Reaction Amandine Guiet, ‡1Tran Ngoc Huan, ‡2 Christophe Payen,3 Florence Porcher,4 Victor Mougel,2,5Marc Fontecave, *2 Gwenaël Corbel *1

1

Institut des Molécules et Matériaux du Mans, UMR 6283 CNRS, Le Mans Université,

Avenue Olivier Messiaen, 72085 Le Mans Cedex 9, France

2

Laboratoire de Chimie des Processus Biologiques, Collège de France, UMR CNRS

8229, Sorbonne Université, PSL Research University, 11 place Marcelin Berthelot, 75005 Paris, France

3Institut

des Matériaux Jean Rouxel (IMN), Université de Nantes, CNRS, 2 rue de la

Houssinière, BP 32229, 44322 Nantes cedex 3, France

4 Laboratoire

Léon Brillouin, CEA-CNRS, 91191 Gif-sur-Yvette Cedex, France

ACS Paragon Plus Environment

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ACS Applied Materials & Interfaces 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

5Present

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address: Department of Chemistry and Applied Biosciences, Eidgenössische

Technische Hochschule Zürich, 8093 Zurich, Switzerland.

‡both authors contributed equally to this study

* corresponding authors

KEYWORDS. Ni1-xCuxTiO3 solid solution, Ilmenite, nanomaterials, neutron powder diffraction, crystal structure, OER electrocatalysts

ABSTRACT. Single Ni1-xCuxTiO3 (0.05≤x≤0.2) Ilmenite-type phases were successfully prepared through a solid state reaction route by using divalent metal nitrates as precursors and characterized. Their electrocatalytic performance for Oxygen Evolution Reaction (OER) in alkaline media is presented. The Cu content was determined (0.05≤x≤0.2) by X-ray diffraction. A thorough powder neutron diffraction study was carried out to identify the subtle changes caused by copper substitution in the structure of NiTiO3.


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The evolution of the optical and magnetic properties with Cu content was also investigated on the raw micrometer-sized particles. A reduction in particle size down to ≈15 nm was achieved by ball-milling the raw powder prepared by solid state reaction. The best catalytic activity for OER was obtained for nanometer-sized particles of Ni0.8Cu0.2TiO3 drop-casted on Cu plate. For this electrode, a current density of 10 mA.cm-2 for oxygen production was generated at 345 mV and 470 mV applied overpotential with 1M and 0.1M NaOH solution as electrolyte, respectively. The catalyst retained this OER activity at 10 mA.cm-2 for long-term electrolysis with a Faradaic efficiency of 90 % for O2 production in 0.1M NaOH electrolyte.

1

INTRODUCTION Water splitting into hydrogen and oxygen is considered as an ideal way to store

electrical energy obtained from the renewable sources in the form of chemical energy.1,2 This reaction occurs within an electrolyzer: the applied electrical energy (for example provided by a solar panel) allows proton reduction at the cathode (Hydrogen Evolution Reaction or HER) and water oxidation at the anode (Oxygen Evolution Reaction or OER) simultaneously. The two electrode reactions involve multiple electrons and multiple


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protons responsible for kinetic limitations and strong overpotentials. Cheap, stable and efficient catalysts are thus absolutely required to overcome these limitations. The largest contribution to overpotentials in a water splitting electrolyzer comes from the anode. Thus tremendous research efforts have been invested to develop efficient OER electrocatalysts, i.e. with high catalytic activity and stability, in particular based on earth abundant metals as alternatives to common catalysts based on Ir and Ru oxides.2–4 In this context Ti-containing materials/electrocatalysts have attracted great interest. While TiO2 has been extensively studied as an anodic catalytic material, including in very recent work,5,6 most efficient alternative approaches have consisted in combining TiO2 with another metal oxide, resulting in efficient and stable OER electrocatalysts in both acidic and basic solution. In the specific case of OER in acidic conditions, the activity of RuO2 and IrO2 is significantly improved in terms of activity and stability when combined with TiO2.7,8 Similarly, the use of Ti in multimetallic metal oxides (or sulfides) has been identified as an efficient strategy for improving catalytic performances for OER in alkaline media. Of particular relevance for OER in basic media are the recent reports of cobaltdoped black TiO2 nanotube arrays,9 of ultrafine NiO nanosheets stabilized by TiO210 and


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of Ti-Fe mixed sulfide nanoboxes,11 all showing excellent performances for OER in 1 M KOH electrolyte, necessitating applied overpotentials of only 352 mV, 320 and 350 mV respectively to achieve 10 mA.cm-2 current density. Ilmenite oxides appear attractive in this context but the reported NiTiO3 OER catalysts presented high overpotentials. Partial substitution of Ni with Co was recently proposed as an effective strategy to decrease overpotentials in these materials, but these remained significantly higher than state-ofthe-art OER catalysts.12 Following this strategy and motivated by the high performances of Cu-based materials as OER catalysts 13–17 from our and other groups, we report in this paper new Ni1-xCuxTiO3 materials resulting from partial Ni substitution by Cu in NiTiO3. Even if Tursun et al. recently reported the sol-gel synthesis of Ni1-xCuxTiO3 powders with nominal compositions (x = 0.05,0.1),18 neither the extent of the solid solution above x = 0.1 nor the effects of copper substitution on the crystal structure of NiTiO3 have been thoroughly determined by these authors. In the present study, micrometer-sized Ni1-xCuxTiO3 particles (0.05≤x≤0.2) of high purity were successfully prepared through solid state reaction. Because NiTiO3 adopts both the


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cation disordered Corundum and the cation ordered Ilmenite structures, neutron powder diffraction (NPD) was used to probe the cation ordering in Ni1-xCuxTiO3 samples and to address the subtle changes associated with copper substitution. The evolution of the optical and magnetic properties with the copper content were investigated on those raw micrometer-sized particles. Secondly, because the nano-structuration increases the surface area and thereby the density of catalytic active sites, a reduction in particle size was performed by ball-milling the raw powder prepared by solid state reaction. The composition and morphology of those Ni1-xCuxTiO3 nanometer-sized particles were determined by electron microscopies, EDS coupled Scanning Electron Microscopy (SEM) and Transmission Electron Microscopy (TEM). The electro-catalytic activity of both micro- and nanometer-sized particles of Ni1xCuxTiO3

Ilmenite-type phases for the Oxygen Evolution Reaction (OER) under alkaline

and neutral conditions was evaluated. The nanostructured Ni1-xCuxTiO3 oxides are highly active and stable during electrolysis.


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2

EXPERIMENTAL SECTION

2.1

Synthesis of Ni1-xCuxTiO3 microsized powders

Copper oxide (CuO, Alfa Aesar), nickel oxide (NiO Alfa Aesar) and titanium oxide (TiO2, Sigma Aldrich) and nitric acid (HNO3, VWR) were employed as the starting agents. One gram of Ni1-xCuxTiO3 (x = 0.05, 0.1, 0.15 and 0.2) polycrystalline sample was successfully prepared as single phase through a modified solid state reaction route using copper and nickel nitrates as precursors. A stoichiometric amount of CuO, NiO and TiO2 was first dissolved in concentrated nitric acid (68%). After homogenization for 15 min, the asobtained solution was mildly heated at 120°C until complete evaporation. The resulting bright-green nitrate mixture was then manually grinded for 5 min and the powder was shaped into pellet (Ø = 13mm). The pellet placed in a platinum crucible was annealed for 12h at 500°C in air to decompose nitrates into oxides (heating rate of 1°C min−1). After annealing, the pellet was manually grinded for 5 min and the powder was shaped again into pellet for the final annealing at high temperature in a platinum crucible. For coppercontaining Ilmenite, a single phase was obtained upon annealing the oxide mixture for 18h (heating rate of 5°C/min) at 950°C for x=0.05 and x=0.10, 900°C for x=0.15 and


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850°C for x=0.2. For each studied Ni1-xCuxTiO3 (x = 0.05, 0.1, 0.15 and 0.2) composition, the above process was performed ten times in order to get the 10 g of sample required for the neutron diffraction study.

2.2

Synthesis Ni1-xCuxTiO3 nanosized powders

Typically, 300 mg of micrometer-sized particles of Ni1-xCuxTiO3 were placed into 45cm3 agate bowl with ten agate balls of 10mm in diameter. Particles were then ball-milled for 2 hrs with a FRITSCH planetary micromillpulverisette 7 apparatus. To prevent overheating of the bowl and of the powder, ball-milling was carried out for 8 alternations of 15 min milling sequences at 400 rpm with 15 min pause in between.

2.3

Preparation of the CuxNi1-xTiO3/Cu plates electrodes

2 mg of Ni1-xCuxTiO3 materials were firstly added into 200 µL ethanol. 10 µL nafion (5%) solution was subsequently added and the resultant mixture was sonicated for 30 minutes. Then, the solution was deposited onto a Cu plate electrode by drop-casting and dried in air before measurements of the electro-catalytic activity for the Oxygen Evolution Reaction (OER) in alkaline and neutral media.


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2.4

Characterization methods

X-Ray Powder Diffraction (XRPD). The phase purity was checked by recording XRPD patterns at room temperature on the PANalytical / Bragg-Brentano Empyrean diffractometer (CuK1+2 radiations) equipped with the PIXcel1D detector. Data were collected in the [5-135°] 2θ scattering angle range for a total acquisition time of 5h30 with a 0.0131° step. The integral-breadth apparent size εβ of nanometer-sized particles was determined from diffraction data by the Integral Breadth method19 (the integral breadth β is defined as the integral (= area) of a peak divided by the peak height). In the present microstructural analysis, the peak broadening is assumed to only originate from the nanometer size of particles ("size broadening" effect) rather than from the lattice distortion ("microstrain broadening" effect). In this approach, the "size broadening" effect is considered being isotropic in nature because the variation of line breath due to the reduction in size of the grains is similar whatever the lattice direction. The size of coherently diffracting domains thereby corresponds to the volume-weighted mean diameter of spherical particles. The instrumental resolution function of the Empyrean diffractometer was first determined


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on the NIST Standard Reference Material LaB6 (NIST SRM 660b). The XRPD pattern of LaB6 was recorded with the same data collection conditions used for the analysis of Ilmenite-type samples. The pattern was then refined by the Le Bail method

20

with the

Fullprof program21 using a modified Thompson-Cox-Hastings pseudo-Voigt profile function (named TCH-Z function)22. In a second step, the XRPD patterns of Ilmenite-type samples were then modelled by the Le Bail method20 using this instrumental resolution TCH-Z function. The two parameters (Y and F) of the Lorentzian component of this function were refined to calculate the average integral-breadth apparent size εβ for all (hkl) reflections. The diameter of the spherical particles is related to εβ by the following formula = 4/3×εβ.19 Neutron Powder Diffraction (NPD). 10 g of Ni1-xCuxTiO3 (x = 0.05, 0.1, 0.15 and 0.2) were loaded in a cylindrical vanadium can (10mm) for data collection at room temperature. High-resolution neutron powder diffraction patterns were recorded on the high resolution two axis 3T2 diffractometer at the Orphée reactor of the Laboratoire Léon Brillouin (CEA Saclay, France) with an incident monochromatic wavelength of 1.229891 Å (Ge curved monochromator). The intensities were measured by a bank of 50 3He detectors in the


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[4.50°-120.35°] 2θ scattering angle range for approximately 22 h with a 0.05° step. The program FullProf was used for Rietveld refinements.21 The line-shape was modelled by a Thompson-Cox-Hastings pseudo-Voigt function (refined parameters U, V, W, X).22 The background intensity was estimated from linear interpolation between up to 45 points manually selected in regions free from Bragg reflections of space group R-3 (n°148). UV-Visible-NIR spectroscopy. The UV−vis diffuse reflectance spectra were recorded at room temperature from 400 to 1000 nm with a 1 nm step using a Shimadzu UV-2700 spectrophotometer. The 100% reflectance was obtained with a Halon powder reference. The reflectivity spectra (R) were transformed to absorption α/S spectra by using a Kubelka−Munk (KM) transformation. The optical band gap value was determined from the intersection of the energy axis thanks to the equation [𝐹(𝑅)ℎ𝜈]2 = 𝐶1(ℎ𝜈 ― 𝐸𝑔). Infrared (IR) spectroscopy. IR spectra were collected in air at room temperature with a Bruker ALPHA FT-IR spectrometer equipped with the Platinum QuickSnap ATR sampling module. The spectral resolution is 4 cm-1 in the 400-4000 cm-1 range. 25 consecutive scans were averaged to obtain a single spectrum. A reference IR spectrum was collected


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in the same conditions with an empty cell and subtracted to specimen spectra to remove H2O(g) and CO2(g) contributions. Scanning Electronic Microscopy (SEM). SEM images of the powders were obtained using a JEOL microscope (JSM 6510 LV). Acceleration voltages varied between 20 and 30 kV as a function of the analyzed samples. Elementary quantitative microanalyses were performed using an Energy Dispersive X-ray (EDX) OXFORD detector (AZtec software). Transmission Electron Microscopy (TEM).The TEM study (SAED and HRTEM) was performed on a JEOL JEM 2100 HR electron microscope operating at 200 kV and equipped with a side entry ±35° double-tilt specimen holder. The samples for transmission electron microscopy investigation were prepared by ultrasonically dispersing the raw powder in ethanol, depositing a drop of the resulting suspension onto a holey carbon-coated copper grid and finally drying the grid in air. N2 sorption. N2 sorption analysis were conducted at 77 K using a TriStar II 3020 (Micrometrics). The powders were degassed under vacuum at 100 °C for 12 h prior measurement. The surface areas were calculated using the Brunauer-Emmett-Teller (BET) method.


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Magnetic measurements. A commercial SQUID magnetometer (MPMS-XL7, Quantum Design) was used to investigate the magnetic properties of the samples. Temperature dependences of the Zero-Field-Cooled (ZFC) susceptibility were measured down to 2 K. The susceptibility, χ, was defined as the ratio of the DC magnetization M to the applied field H (i.e., χ = M/H). Magnetization data were acquired under a magnetic field of 5 kOe. Electrochemical measurements. Electrocatalytic measurements and electrolysis experiments were performed in a three-electrode two-compartment cell allowing for separation of gas phase products at the anodic and cathodic compartments using a Biologic SP300 potentiostat. Ag/AgCl was used as the reference electrode and placed in the same compartment as the working electrode. A platinum counter electrode was placed in a separate compartment connected by a P5 glass-frit and filled with the electrolytic solution. The geometrical surface of the working electrode was 1 cm2 for all this study. The potentials were referenced versus RHE by using the equation below: ERHE = Evs Ag/AgCl + 0.195 + 0.059*pH The results of linear sweep voltammetry (LSV) were compensated with Ohmic drop. Faradaic efficiency was obtained by comparing the theoretical amount of oxygen that


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should be produced based on the charge consumption to the amount of oxygen determined using gas chromatography. Bulk electrolysis was performed under an applied fixed current density of 10 mA/cm-2 in a solution of 0.1M KOH (pH 13). The solution of the anodic compartment (10 mL total volume) was constantly renewed using a peristaltic pump at a flow of 0.5 mL/min, during the 10 h experiment time.

3

RESULT AND DISCUSSION

3.1

Synthesis of Ni1-xCuxTiO3 micro- and nanomaterials

The micrometer-sized particles of the parent compound NiTiO3 are usually synthesized by a conventional solid state reaction route. This synthetic route requires annealing of a stoichiometric mixture of metallic Ni or NiO and TiO2 at temperatures higher than 1000°C.23,24 However, single NiTiO3 phase is difficult to obtain and the presence of secondary phases such as Bunsenite (NiO) and/or Rutile (TiO2) is often observed. This phase purity problem is not restricted to NiTiO3 and concerns most of the Titaniumcontaining Ilmenite-type materials, likely deriving from the poor reactivity of divalent metal oxides with respect to TiO2.25 In this context, copper and nickel nitrates were used as


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precursors to successfully prepare, as a single phase, the copper-substituted NiTiO3 (Ni1xCuxTiO3

with 0.05≤x≤0.20) polycrystalline materials. Indeed, the thermal decomposition

of nickel (or copper) nitrate can serve to in-situ produce nickel (or copper) oxide particles of few hundred nanometers in diameter (Ø ~ 200nm) thus having a higher reactivity than commercial micrometer-sized NiO (or CuO) particles (Ø ~1-2 μm) (Figure S1). Furthermore, metal nitrates often are hydrated compounds and the hydration degree of commercial nitrates is not always constant. Thereby, a mixture of copper and nickel nitrates and TiO2 was directly prepared by dissolving the corresponding oxides weighted in stoichiometric proportion in concentrated nitric acid (68%) to perfectly control the copper content of the targeted composition. The solution was mildly heated at 120°C until complete evaporation giving a mixture of copper and nickel nitrates together with TiO2 as shown by X-Ray Power Diffraction (Figure S2a). After the decomposition of those nitrates at 500 °C in air, an intimate mixture of elementary NiO, CuO and TiO2 oxides was obtained (Figure S2b). Particle diameter of around 200 nm is preserved after decomposing the nitrates as shown by SEM in Figure S1. To promote a solid state reaction between elementary oxides, calcination in air at temperature higher than 500°C


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is required. A single phase of Ni1-xCuxTiO3 was obtained upon annealing the oxide mixture for 18h at 950°C for x = 0.05 and x = 0.10, 900°C for x = 0.15 and 850°C for x = 0.2. For the parent composition NiTiO3 (x = 0), some traces of NiO (~3.5 mol%) always remain present whatever the annealing temperature used. SEM images reveal micrometer-sized grains for all Ni1-xCuxTiO3 compositions, as expected for oxides synthesized by solid state reaction at high temperature (Figure S1). Contrary to Co1-xNixTiO3,12 the solid solution studied here is not continuous in the whole composition range but limited to 20 mol% in copper (x = 0.2). Indeed, for a copper content higher than x = 0.25, simple oxides (TiO2 and NiO-type) are systematically detected in addition to a major NiTiO3-type phase whatever the annealing temperature used. Nanometer-sized particles of Ni1-xCuxTiO3 oxides were then directly obtained by ball-milling those micrometer-sized particles.

3.2

Structural analysis of micrometer-sized Ni1-xCuxTiO3 powders.

Electron, X-ray and Neutron diffraction techniques were performed to demonstrate the successful substitution of nickel by copper and to probe the cation ordering in Ni1-xCuxTiO3 samples as well as to investigate subtle changes in the structure resulting from the copper substitution.


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2.2.1 Crystal structures of the pristine NiTiO3: a crystallographic reminder Pristine NiTiO3, as several other A2+Ti+4O3 titanates (A= Mg, Mn, Fe, Co and Zn), adopts the Ilmenite-type structure at room temperature. The Ilmenite structure (space group R-3 n°148) can be described as a pseudo-compact ABAB stacking of oxygen layers along the

c axis (in hexagonal setting) with cations occupying only two thirds of the octahedral interstices (Figure 1a). A2+ and Ti4+cations occupy two distinct crystallographic 6c sites in the structure named C1 (0,0,≈1/3) and C2 (0,0,≈1/6) sites. The C1 sites, as well as the C2 ones, form honeycomb layers of edge-sharing octahedra parallel to the (a,b) plane (Figure 1b). In the structure, honeycomb C1-type and C2-type layers successively alternate along the c hexagonal axis, each C1-type layer being translated from the 



adjacent C2-type one by the vector 1 2 a  1 2 b . In terms of connection between honeycomb layers, each occupied octahedron of one C1-type layer shares a triangular face with one occupied octahedron of a first C2-type layer and the opposite triangular face with a "vacant" octahedron corresponding to the "comb" of a second C2-type layer (Figure 1a). This face-sharing connexion of octahedra induces second-nearest neighbour


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electrostatic repulsions between C1 and C2 cations. To minimize repulsions, C1 and C2 cations move from the octahedron barycentre (off-centering) towards the vacancies or the "combs" of the neighbouring C2-type and C1-type layers, respectively (Figure 1c). As crystallographic 6c sites (0,0,z) are located on the threefold axes running along the hexagonal c axis, the displacements of C1 and C2 cations residing on these sites are prone to take place along these axes and in opposite directions from the shared triangular face (the [O4O5O6] face in Figure 1c). Within each honeycomb C1-type and C2-type layers (Figure 1b), positive and negative deviations  from the octahedron barycentre are thus noticed: z=1/3±1 and z = 1/6±2 for C1 and C2 sites, respectively. The deviation magnitude |δ| from the octahedron barycentre is indicative of the electrostatic repulsion between the cations. Among all titanates, only NiTiO3 undergoes a second order phase transition from Ilmenite-type to Corundum-type structures at around 1570 K.26 Both structures are built up from the same stacking of honeycomb octahedral layers. However, all octahedral layers in the high temperature Corundum structure (space group R-3c n°167) are randomly populated by both Ni2+ and Ti4+cations. The Ilmenite structure can be thus


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described as an ordered derivative of the Corundum one. When comparing X-ray powder diffraction patterns of Ilmenite and Corundum forms, the cation ordering induces the appearance of small superstructure (003) and (101) diffraction peaks at low 2θ scattering angles (Figure S3).

Figure 1: a) Projection onto the (1-10) plane of portion of the Ilmenite structure illustrating the cation ordering. b) Projection along the hexagonal c axis of the honeycomb C1-type and C2-type layers. c) Off-centering of C1 and C2 cations within their coordination octahedra sharing the triangular [O4O5O6] face. The solid black line is the trace of the unit cell on both projections, in the background of octahedra for a) and in the front of octahedra for b).

2.2.2 Structural study of the micrometer-sized Ni1-xCuxTiO3 materials


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Electronic diffraction was first performed on several crystallites of Cu0.15Ni0.85TiO3 (Figure S4). The reciprocal space reconstitution by Selected Area Electron Diffraction (SAED) technique confirms the A2+B4+O3 structure with a hexagonal cell, a ≈ 5.1 Å and c ≈ 13.8 Å, corresponding to the classical parameters of the NiTiO3-type phase unit cell. The reflection conditions are compatible with the R− − extinction symbol. However, this study does not allow to make the difference between the R-3 (n°148) and R-3c (n°167) space groups of Ilmenite and Corundum structure respectively. Indeed, the main difference between those two space groups relies on the 00l reflection conditions (00l = 3n for R-3 and 00l = 6n for R-3c) which cannot be differentiated due to the double diffraction phenomenon, typical in electron diffraction. Therefore, X-ray and neutron powder diffraction were subsequently performed. The X-ray diffraction patterns of Ni1-xCuxTiO3 (x = 0.05, 0.1, 0.15 and 0.2) powders show that a single NiTiO3-type phase is obtained for all compositions (Figure S3). The two characteristic superstructure (003) and (101) peaks of the Ilmenite-type phase are very weak in intensity and undetectable, respectively. Thereby, a doubt subsists on the effective and complete stabilization of an Ilmenite-type structure at room temperature for


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Ni1-xCuxTiO3. Indeed, one can interpret this situation as resulting from the coexistence of a major Corundum-type form with a minor Ilmenite-type one. For oxides exhibiting a firstorder structural phase transition from an ordered low-temperature form to a disordered high-temperature one upon heating, cationic substitutions can stabilize, above a certain level, the disordered form at room temperature, like when lanthanum (or molybdenum) is partially substituted by europium23 (or tungsten24) in La2Mo2O9. A stabilization of the Corundum form of NiTiO3 at room temperature could therefore occur above a certain copper content x in the present Ni1-xCuxTiO3 series due the cationic disorder caused by the random substitution of Ni2+ ions. Differential Thermal Analysis (DTA) cannot be used to determine the substitute content leading to the complete stabilization of the Corundum form at room temperature because NiTiO3 exhibits a second-order transition (no endothermic peak). In this context, neutron diffraction is a suitable technique to study the phase transition in NiTiO3 as diffraction patterns of Ilmenite (cation ordered) and Corundum (cation disordered) phases significantly differ from each other (Figure 2a).26 Indeed, the intensity of the two superstructure (003) and (101) peaks at low 2θ scattering angles is large, thus


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making easy the distinction between the two forms of NiTiO3. This strong difference between X-ray and neutron diffraction patterns arises from the very different neutron scattering lengths of Ni (10.3 fm) and Ti (-3.438 fm) atoms27 (while they are almost indistinguishable by X-rays). The neutron scattering length of Cu atom being also positive (7.718 fm), a neutron diffraction study can be therefore carried out on the Ni1-xCuxTiO3 series to probe the cation ordering and consequently identify the crystallographic form of NiTiO3 at room temperature (Ilmenite versus Corundum). Neutron diffraction can also allow to detect the subtle changes induced by copper substitution. Indeed, divalent copper is known to exhibit Jahn-Teller distortion of its octahedral coordination. As neutrons are much more sensitive to oxygen than X-rays, neutron diffraction study allows the accurate positioning of oxygen and provides a deep insight on the way copper ion distorts the geometry of C1-type octahedra initially occupied by Ni2+ alone and changes the off-centering of Ti4+ in the neighbouring C2-type octahedra of the structure of NiTiO3. In all neutron diffraction patterns in Figure 2a, the two superstructure (003) and (101) peaks are now clearly detected at 15.4° and 17.0° in 2θ, thus implying that all Ni1-xCuxTiO3


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samples adopt the Ilmenite-type structure at room temperature. To confirm the complete stabilization of an Ilmenite-type structure and therefore the absence of a fraction of the Corundum form, the crystal structure has been refined from the NPD pattern by the Rietveld method using solely the Ilmenite structure of NiTiO326 as a starting structural model for each composition (ordered model hereafter named model 1). In this full ordered model (model 1), C1 (0,0, 1/3±δ1) and C2 (0,0,1/6±δ2) sites are only occupied by divalent Ni/Cu and Ti cations, respectively. The C1 site occupation factors for Ni and Cu atoms were kept fixed to (1-x) and x, respectively, according to the nominal Cu content x of the analysed sample. The possibility of a cationic exchange between the honeycomb C1-type and C2-type layers was also tested: exchange of divalent Ni (model 2) or Cu (model 3) cations with Ti4+ ions and vice versa to preserve the stoichiometry. When the anisotropic temperature factors βij of C1 and C2 sites were separately refined, Rietveld refinements give non-positive definite matrix of atomic displacements βij for the C2 site whatever the structural models used. When the anisotropic temperature factors βij of C2 site were constrained to that of C1 one, the fraction of divalent cations on the C2-site and reciprocally of Ti4+ ions on the C1 site was around 1% and sometimes below 0.5% with


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an erratic evolution along the series of samples. The cationic exchange was then considered to be negligible and the models 2 and 3 were therefore ruled out. Finally, the model 1 with a constraint on the anisotropic temperature factors βij of C1 and C2 sites was selected. All neutron diffraction patterns were satisfactorily modelled by least-squares fitting this structural model consisting in 5 atomic coordinates, 12 anisotropic thermal factors and the 2 cell parameters a and c (in hexagonal setting). The values of atomic parameters (positions coordinates, equivalent isotropic temperature factors together with conventional reliability factors are reported in Table S1 for the four samples x = 0.05, 0.1, 0.15 and 0.2. Figure 2b shows the final observed, calculated, and difference diffraction patterns of Ni0.8Cu0.2TiO3. The evolution of the unit cell parameters a and c as a function the copper content x is displayed in Figure 2c. Along the series, the linear expansion of both a and c parameters with increasing the copper content x (usual Vegard’s law) reflects that cupric cation (ionic radius = 0.73 Å (CN = 6))28 is larger in size than Ni2+ ion (ionic radius = 0.69 Å (CN = 6).28 The existence of these two Vegard's laws attests the formation of a substitutional solid solution in the whole composition range investigated 0 ≤ x ≤ 0.2. This expansion of the unit cell caused by copper substitution is


24

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not in agreement with the study of Tursun et al. carried out on NiTiO3(x = 0), Ni0.95Cu0.05TiO3 (x = 0.05) and Ni0.9Cu0.1TiO3 (x = 0.1) samples.18 Without any trend in the series of Tursun et al., their values of the a and c cell parameters are, respectively, lower and higher than those determined on the same compositions in the present study.


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Figure 2: a) NPD patterns of Ni1-xCuxTiO3 raw powder samples collected at room temperature. Comparison of the NPD patterns of the Ilmenite form of NiTiO3at 300 K with that of the its Corundum form at 1575K, both calculated from crystallographic data published in reference 26 with the neutron radiation wavelength used in the present study. b) Variation of the unit cell parameters of Ni1-xCuxTiO3 raw powder samples with the copper content x. c) Comparison of the observed NPD pattern of Ni0.8Cu0.2TiO3 sample with the pattern calculated using the structural model reported in Table S1. Vertical markers give Bragg peak positions of the space group R-3 (n°148). c) The cell parameters determined for NiTiO3 (x=0) 26 are added for reference (open triangles).

To detect the effect of copper substitution on the crystal structure of NiTiO3, a thorough analysis of the evolution of inter-atomic distances with the copper content x is required. These distances have been calculated from the experimental data reported in Table S1. Within the coordination octahedra of C1 sites randomly occupied by divalent Cu/Ni cations, a slight elongation of the O7-O8 edge clearly occurs as the copper content x increases (≈+0.005 Å over the range 0.05≤x≤0.2) whereas no change in length of the O4O5 edge is observed (Figure S5). O8-O9 and O7-O9 edges are both equivalent in length to that of the O7-O8 one due to the threefold inversion axis (Figure 1). The same situation exists for the O4-O5, O5-O6 and O4-O6 edges (Figure 1). The copper substitution thus induces an isotropic expansion in size of the triangular [O7-O8-O9] face of the [C1O6] octahedron while leaving unchanged the size of the opposite triangular [O4-O5-O6] face shared with the [C2O6] octahedron. An asymmetric deformation of the [C1O6] octahedron then occurs in the (a,b) plane as the copper content x increases. In addition to this distortion, a slight increase of the distance h1 between the opposite triangular [O4-O5-


26

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O6] and [O7-O8-O9] faces of the [C1O6] octahedron is shown in Figure 3a (Δh1 ≈ +0.004 Å over the range 0.05 ≤ x ≤ 0.2). Thus, the [C1O6] octahedron slightly expands along the c axis as it deforms upon copper substitution. These asymmetric deformation and trigonal expansion are both responsible for the elongation of both O4-O7 and O4-O8 edges with x, observed in Figure S5. The octahedral crystal field splits the five d-orbitals of metal ion into two sets, one set (t2g) being triply degenerate (corresponding to dxy, dyz and dxz orbitals) and the other (eg) being doubly degenerate (corresponding to 𝑑𝑥2 ― 𝑦2 and 𝑑𝑧2 orbitals directing electron density toward the oxide ligand). The t2g orbitals are stabilized while the eg orbitals are destabilized relative to their energies in a spherical field. For the d9 electronic configuration of cupric ion ((t2g)6(eg)3), the t2g orbitals are fully occupied whereas unequal occupancies of the two eg orbitals are noted. When a Jahn-Teller distortion occurs, the singly occupied eg orbital is destabilized by the same energy as the doubly occupied eg orbital is stabilized, resulting in a net stabilization of energy. The single occupancy of the 𝑑𝑧2 or 𝑑𝑥2 ― 𝑦2 orbitals of Cu2+ results in compressed [2+4] or elongated [4+2] coordination octahedron along its 4-fold symmetry axis, respectively. Here, these asymmetric deformation and trigonal expansion of the coordination octahedron are clearly


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not related to a Jahn-Teller effect of cupric ions. In the compact structure of Ilmenite, the only way to deform the [C1O6] octahedron is to expand in size the triangular [O7-O8-O9] face which is not shared with a neighbouring [C2O6] octahedron. The impossibility for cupric ions to impose a Jahn-Teller distortion of the [C1O6] octahedron for geometrical reasons could explain why the solubility limit of copper in NiTiO3 does not exceed 20 mol%. However, despite these subtle changes in the geometry of the [C1O6] octahedron upon copper substitution, no evolution of the C1-O7 bond length with x is detected in Figure 3b (C1-O8 and C1-O9 bond are both identical in length to the C1-O7 one by symmetry). In order to leave unchanged those C1-O7,8,9 bond lengths, the C1 cation must necessarily move away from the octahedron barycentre in the direction of the triangular [O7-O8-O9] face as it expands with x. Figure 3c shows that the distance between C1 site and this octahedron

barycentre

effectively

increases

upon

copper

substitution

(Δ|δ1|

≈ + 0.007 Å over the range 0.05 ≤ x ≤ 0.2). The effect of the [O7-O8-O9] face expansion on the C1-O7,8,9 bond lengths is here fully balanced, in magnitude, with the effect of the displacement of C1 cations towards this face. In addition, a lengthening of


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the C1-O4 bond with x is observed in Figure 3b (C1-O5 and C1-O6 bond are both identical in length to the C1-O4 one by symmetry). C1 cations move away from the opposite [O4O5-O6] face shared with the [C2O6] octahedron. Despite the trigonal expansion of the [C1O6] octahedra along the c axis, the off-centering of C1 cations in these octahedra therefore increases, in average, upon copper substitution. When considered two neighbouring [C1O6] octahedra sharing a [O-O] edge within an honeycomb C1-type layer (Figure 1b), the deviation of C1 cations from the octahedra barycentres is of the same magnitude |δ1|and is alternately positive (+δ1) in the former and negative (-δ1) in the latter. Because this deviation |δ1| increases in magnitude with x, two nearest C1 cations move away from each other, thus increasing the C1-C1 distance upon substitution as observed in Figure 3d (ΔdC1-C1 ≈ +0.004 Å over the range 0.05 ≤ x ≤ 0.2). Within the coordination octahedra of C2 sites only populated by Ti4+ cations, no change in length of the O1-O2, O1-O3 and O2-O3 edges (both equivalents in length due to the threefold inversion axis) is noted. Contrary to what was observed for the free [O7-O8-O9] face of the [C1O6] octahedron, the free [O1-O2-O3] face of the [C2O6] one does not expand upon copper substitution. As stated above, no change in size of the opposite


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triangular [O4-O5-O6] face, shared with the [C1O6] octahedron, was also noticed as increasing the copper content x. A very slight decrease of the distance h2 between these opposite triangular [O1-O2-O3] and [O4-O5-O6] faces of the [C2O6] octahedron is shown in Figure 3a(Δh2 ≈ -0.002 Å over the range 0.05 ≤ x ≤ 0.2). Thereby, the copper substitution only induces a slight trigonal compression of the [C2O6] octahedron along the c axis (contraction in length of O1-O5 edge with x in Figure S5). When the octahedron does compress upon substitution, the C2-O1 bond length (C2-O2 and C2-O3 bond are both identical in length to the C2-O1 one by symmetry) remains however unchanged as increasing x (Figure 3b). The C2-O1,2,3 bond lengths are much shorter than the sum of the ionic radii of Ti4+ (0.605 Å28) and O2- (1.38Å28) in sixfold and fourfold coordination, respectively. Because the C2-O1,2,3 bond lengths cannot physically be shorter than they already are, C2 cation must move closer to the octahedron barycentre in the direction of the [O4-O5-O6] face when the [C2O6] octahedron does compress. In Figure 3c, the distance between C2 site and the octahedron barycentre effectively “decreases” with x (Δ |δ2| ≈ ― 0.003 Å over the range 0.05 ≤ x ≤ 0.2), thus “shortening” the C2-O4 bond length (C2-O5 and C2-O6 bond are both identical in length to the C2-O4 one by


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symmetry). Within an honeycomb C2-type layer of edge-sharing [C2O6] octahedra (Figure 1b), the re-centering of C2 cation in octahedra remains too small in magnitude to significantly decrease the C2-C2 distance upon the substitution (ΔdC2-C2 ≈ -0.0007 Å over the range 0.05 ≤ x ≤ 0.2 in Figure 3d). The distances between the C2 site and neighbouring O4 atom (Figure 3b) or between C2 site and the octahedron barycenter (δ2 in Figure 3c) exhibit standard deviations larger than the magnitude of their evolutions over the compositional range. One could speak of a trend rather than a shortening/decrease of those distances in the above discussion. As earlier mentioned, (see Rietveld refinement strategy), unphysical temperature factors were found for C2 site when their atomic displacements βij were refined independently from those of C1 site. The unphysical temperature factors could originate from a spread distribution of nuclear density around the C2 site, thus making its location within the octahedron difficult to determine accurately without any constraints. In Figure 3d, the standard deviation on the C1-C2 distance is much larger than that on the C1-C1 distance. This difference in standard deviations reflects that the C2 site location remains less accurate than the C1 site one, despite the constraints on atomic displacements βij used


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in the final Rietveld refinements, thus both increasing the standard deviations on C2-O4 and δ2 distances. The copper substitution only has a second-neighbor effect on those C2O4 andδ2 distances. Because the magnitude of this effect is directly dependent on the copper content x which does not exceed x = 0.2 (20 mol%), the standard deviations on those distances are higher in magnitude than the distance evolutions with x. Both remarks point out that the limit of the diffraction technique is probably attained here for the accurate location of C2 site or/and the detection of the second-neighbor Cu2+ effect. Because Cu2+ is larger in size than Ni2+, the [CuO6] octahedra in the structure of Ni1xCuxTiO3

oxides are larger in volume and more distorted than the [NiO6] ones. The

increase of the hexagonal cell a parameter clearly originates from the average expansion in size of the triangular [O7-O8-O9] face of [C1O6] octahedra which are parallel to the (a,b) plane. Because of the limited extent of the solid solution (x ≤ 0.2), the proportion of distorted [CuO6] octahedra within the honeycomb C1-type layer is weak, thus making very small in magnitude the increase of the a parameter upon substitution (Δa ≈ +0.003 Å over the range 0.05 ≤ x ≤ 0.2). In comparison, the expansion of the hexagonal cell along the c axis is, in magnitude, almost twice as large as along the a axis (Δc ≈ +0.005 Å over the


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range 0.05≤x≤0.2). However, the introduction of the large cupric ions in the honeycomb C1-type layer does not expand the hexagonal c parameter as much as it should do. Indeed, along the c axis, the trigonal compression of the [C2O6] octahedron (Δh2 ≈ -0.002 Å over the range 0.05 ≤ x ≤ 0.2) half reduces the effect of the trigonal expansion of [C1O6] octahedron (Δh1 ≈ +0.004 Å over the range 0.05 ≤ x ≤ 0.2).

Figure 3: Evolution, as a function of the copper content x in Ni1-xCuxTiO3 samples, of the octahedra heights (a), of the cationoxygen distances (b) of the deviations || from the octahedron barycentre (c) and of the cation-cation distances (d). All these values have been calculated from the crystallographic data reported in Table S1. Black lines = linear regressions. These lines must be


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considered as guides for eyes in the evolutions of C2-O4 and C1-C2 distances and deviation |𝟐|with the copper content x. The values calculated from the crystallographic data of NiTiO3 (x=0)26 are added for reference (open symbols).

3.3

Magnetic susceptibility and optical properties of the micrometer-sized Ni1-xCuxTiO3 samples

Besides the structural changes caused by copper substitution, continuous evolutions in the magnetic properties and in the optical properties were also observed. From the magnetic susceptibility curves, a linear decrease of the molar Curie constant with increasing Cu content is observed which is fully consistent with chemical substitution of S=1/2 Cu2+ magnetic moments for S=1 Ni2+ ions in NiTiO3 (Figure S6, Table S2). The FT-IR spectra of Ni1-xCuxTiO3 powders show three main absorption bands (marked as I, II, III in the Figure S7). Absorption band I at 424 cm-1 is assigned to the stretching vibration ν of Ni-O bonds, band II at 504 cm-1 to the stretching vibration ν of Ti–O bonds, and band III at 617 cm-1 to the bending vibration δ of O–Ti–O bonds; these band assignments are in agreement with those previously reported for NiTiO3.18 Furthermore, no absorption bands ascribed to the bending vibration δ of water molecules at around 1600 cm−1 and to the stretching vibration ν of hydroxyl group at around 3700 cm−1 are detected in all the IR spectra collected. No water molecules are therefore absorbed at the surface of grains.


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As shown in Figure 4, the optical properties characterization was conducted by diffuse reflectance spectroscopy and the experimental spectra were analyzed using the Kubelka– Munk (K–M) model.29 For x=0, the absorption spectral feature is in accordance with the yellow color associated to Ni2+ ions (3d8 ion) in octahedral C1 sites of NiTiO3 Ilmenite.30,31 Three wide bands corresponding to the spin-allowed (∆S = 0) transitions from the 3A2g (3F) ground state are observed: a sharp band in the blue region around 450 nm (ν3: 3A2g (3F) → 3T1g (3P)), a broader band in the red region centered around 750-900 nm ν2: 3A2g (3F) → 3T1g (3F)) and a last broad band in the near IR centered between 1100 and 1400 nm (ν1: 3A2g (3F) → 3T2g (3F)). In addition, the absorption shoulders observed around 500 nm and 740 nm may be assigned to spin-forbidden transitions to 3A1g (1G) + 1T2g (1D) and 1E

g

(1G) terms. Moreover, the spectra present also a sharp band in the near UV

associated to the typical charge transfer (CT) transitions O2−–Ni2+ (320–390nm), which partially overlaps with ν3 band in the blue region. As the positions of those three ν bands do not change with the Cu content x, it can be first assumed that there is no change neither in the coordination nor in the oxidation state of Ni. However, subtle changes in the UV-vis absorption spectrum are induced by copper


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substitution (region highlighted by a circle and a square in Figure 4a). Since the octahedral ligand field stabilization energy of Cu2+ (0.6 in units of ΔO)32 is twice lower than that of Ni2+ (1.2 in units of ΔO)32, the LMCT between oxide ions and divalent cations as well as the d-d electronic transitions take place at wavelength higher for cupric ions than for nickel ones. The consequence is twofold. First, the overlapping of bands associated to LMCT transitions O2−–M2+ for Ni2+ and Cu2+ ions transforms the sharp band for NiTiO3 oxide into a broad massif for Ni1-xCuxTiO3 compounds. As the copper content x increases, a change in the relative proportion of the LMCT bands occurs, which shifts in position the massif to higher wavelength, thus inducing a strong overlapping with the ν3 band. Secondly, the depression in the absorption spectrum at around 600 nm (and the absorption edge at around 650 nm), red-shifts with increasing the Cu content x which is certainly at the origin of the color darkening of powders from yellow to light brown with x (Figure S8). This red-shift of the depression at around 600 nm is the sign of a reduction of the direct optical band gap. For each sample, the optical band gap was determined from the Tauc plot (K/S*h* ν)2 vs.hν).33 From the analysis of those plots (Figure S8), a linear decrease of the optical band gap from 2.5 to 2.05 eV was thus noted with increasing


36

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the Cu content from x = 0 to 0.2. Finally, a new band centered at 1050 nm grows in intensity as the copper content x increases. A peak deconvolution in the wavenumber range 7000-16000 cm-1 was carried out in order to extract the band area values of ν2 band and of this new band, hereafter named A2 and A respectively. The band area ratio A2/A increases linearly with the copper content x (Figure 4b) and c)). To conclude, both the linear decrease of the molar Curie constant (magnetic properties) and the evolution of the band area ratio A/A2 together with the color darkening (optical properties) give further experimental evidences that copper ions really substitute for nickel ones in the [C1O6] octahedra of the Ilmenite structure. This conclusion is consistent with the neutron diffraction study. The formation of a substitutional solid solution in the composition range 0 ≤ x ≤ 0.2 is now definitively established.


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Figure 4: a) UV-vis absorption spectra of the micrometer-sized Ni1-xCuxTiO3 particles. b) Peak deconvolution of the absorption spectrum of Ni0.8Cu0.2TiO3 sample in the wavenumber range 7000-16000 cm-1 as representative example and c) Plot of the band area ratio A/A2 (A2= area of the ν2 band, A= area of the new band) as a function of x.

2.3. Nanometer-sized Ni1-xCuxTiO3 particles Nanometer-sized Ni1-xCuxTiO3 particles were obtained by ball-milling the corresponding micrometer-sized particles previously prepared by solid state reaction. The corresponding XRD patterns show that the Ilmenite-type structure is preserved upon ball-milling. However, such harsh conditions also induce a minor sample pollution with Quartz SiO2 coming from the Agate grinding jar used (2 ≈26.6° in Figure 5d, Quartz (PDF No. 461045)). High-resolution TEM (HRTEM, Figure 5a-b as well as selected area electron diffraction (SAED) confirmed the good crystallinity of the Ni1-xCuxTiO3 particles. The lattice fringe spacing measured in Figure 5b) is in good agreement with the inter-reticular distance between consecutive (104) planes of the Ilmenite structure. For each composition, the unit cell parameters and the diameter of the spherical particles were determined by the Le Bail method from the XRPD pattern (see experimental part). A satisfactory example is given in Table S3. The diameter of ≈15 nm extracted from the analysis of XRPD data is in good agreement with the particle size measured by TEM. N2 sorption measurements were also carried out on both micro- and nanometer-sized Ni1-


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xCuxTiO3

particles to determine the specific surface area (SABET) of samples according to

the Brunauer–Emmett–Teller (BET) theory. As expected, the measured surface areas for micrometer-sized particles were less than 2 m2.g-1. However, the SABET was drastically increased above 20 m2.g-1 after reducing the particle size down to around 15 nm by ballmilling. (See table in Table S4).

Figure 5: TEM (a), HRTEM (b) images and the corresponding SAED pattern (c) for Ni0.8Cu0.2TiO3-nano. (d) Le Bail refinement of the XRD pattern of the ball-milled Ni0.8Cu0.2TiO3 powder.

Electrocatalysis for Oxygen Evolution Reaction (OER) The Ni1-xCuxTiO3 materials were evaluated for electrocatalytic OER in various pH conditions. The electrodes were prepared by drop-casting the sample powder onto a Cu plate (see experimental part for details) and the Ni1-xCuxTiO3-nano/Cu plate electrodes


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were characterized by SEM and EDX measurement (Figure S9). The presence and homogeneous distribution of all elements contained in the deposited material (O, Cu, Ni, Ti) were confirmed by EDX spectroscopy. Figure 6a shows the Linear Sweep Voltammetry (LSV) curves (after Ohmic drop correction) of the different Ni1-xCuxTiO3 /Cu electrodes in 0.1 M NaOH solution. The bare Cu plate electrode exhibited a very low catalytic activity with a high onset potential of 620 mV vs RHE. For x=0.2, the electrocatalytic activity of the nanometer-sized particles significantly increased in comparison with the corresponding micrometer-sized particles before ball-milling. Based on capacitance measurements (Figure S10), we have shown that the nanostructuration led to a significant increase of the electrochemical surface area (ECSA): the capacitance values were determined as 0.87 F.cm-2 and 0.42 F.cm-2 for Ni0.8Cu0.2TiO3-nano and Ni0.8Cu0.2TiO3-micro, respectively. Since the specific capacitance is expected to be the same for both materials, it means that the ECSA value for Ni0.8Cu0.2TiO3-nano is 2.1 times larger than that for Ni0.8Cu0.2TiO3-micro. This could partly explain the greater activity of the former. In addition, the LSV data indicated an increase of the catalytic activity for OER following Ni substitution by Cu for both micrometer- and


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nanometer-sized materials. Indeed, a maximum activity is measured for the Ni0.8Cu0.2TiO3-nano/Cu electrode as demonstrated by the observed low applied overpotential values (340 mV vs RHE at 1.0 mA.cm−2 and 470 mV at 10 mA.cm−2). These values are among the best ones for OER in 0.1 M NaOH when compared to other polymetallic electrocatalytic materials containing Ti (Table S5). The higher electrocatalytic activity of Ni0.8Cu0.2TiO3/Cu plate is consistent with a lower value of the Tafel slope, as shown in Figure 6b. The Tafel slope in the case of the Ni0.8Cu0.2TiO3-nano/Cu electrode in 0.1M NaOH is 125 mV dec−1, a lower value than those obtained for the Ni0.8Cu0.2TiO3-micro/Cu and NiTiO3-nano/Cu electrodes, 165 and 153 mV dec−1, respectively. This result indicates a more favorable kinetic and a superior OER catalytic power for Ni0.8 Cu0.2TiO3-nano/Cu electrode. Electrolysis at a controlled potential of 1.78 V vs RHE using Ni0.8Cu0.2TiO3-nano/Cu electrode was carried out over 90 minutes (Figure 6d). A stable current density of 20 mA.cm−2 was obtained corresponding to a large passed charge (Figure S11) and oxygen was formed with a high Faradaic efficiency of 90 % (Figure 6 and Figure S12). The stability of the electrode was probed under long-term electrolysis at a constant current of


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10 mA.cm−2for OER in 0.1 M NaOH solution (Figure S13). The electrode material proved well stable since the current density at 460 mV applied overpotential remained stable during 10h of electrolysis. The OER catalytic activity of Ni0.8Cu0.2TiO3-nano/Cu plate electrode was also investigated in electrolytes at different pHs (Figure 6c). In 1.0 M NaOH, Ni0.8Cu0.2TiO3nano/Cu plate electrode required only 345 mV applied over-potential in order to reach a current density of 10.0 mA.cm-2. At pH=11 and 9, an applied overpotential of 450 and 510 mV, respectively, is required to reach 1.0 mA.cm-2 current density.


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Figure 6: a) Catalytic activity of different material electrode for O2 evolution in 0.1M NaOH, b) the Tafel plots of Ni0.8Cu0.2TiO3micro/Cu plate (red), NiTiO3-nano/Cu plate (blue) and Ni0.8Cu0.2TiO3-nano/Cu plate (black) c) catalytic activity of Ni0.8Cu0.2TiO3nano/Cu plate in different pH solution, d) electrolysis for O2 evolution of Ni0.8Cu0.2TiO3-nano/Cu plate.

The Ni0.8Cu0.2TiO3-nano material displayed an excellent performance in electrochemical oxidation of water likely due to a high proportion of incomplete coordination spheres at the metal active sites on the surface, which synergistically promote OH- adsorption and facilitate charge transfer.5,14,34 As established for other Ni-based OER catalyst, the active sites in NiTiO3 are Ni ions in the C1 site, allowing activation of adjacent adsorbed OH


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groups (at a distance of 2.92 Å) to form O-O bond (Figure S14). While further studies are required to understand how Cu ions, partly replacing Ni ions, stimulate OER activity, we propose that a partial charge transfer between Ni and Cu provides Ni active sites with greater OH adsorption potential, resulting in enhanced OER activity.

4

Conclusion In the present study, micrometer-sized particles of single Ni1-xCuxTiO3 Ilmenite-type

phase were prepared by solid state reaction from nitrate precursors. The extent of the solid solution (0.05 ≤x ≤ 0.2) was determined for the first time. A thorough structural analysis carried out by neutron powder diffraction on those micrometer-sized particles allows to highlight the subtle changes caused by the copper substitution in the structure of NiTiO3. Copper substitution in NiTiO3 induces an asymmetric deformation of the [(Ni,Cu)O6] octahedron in the (a,b) plane together with the expansion in height of the octahedron along the c axis. Along this c axis, [TiO6] octahedron endures, at reverse, a compression in height which is in magnitude lower than the expansion of the [(Ni,Cu)O6]


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octahedron. These subtle structural changes are at the origin of linear increases of both hexagonal cell a and c parameters observed with the copper content x. A reduction in particle size down to ≈15 nm was achieved by ball-milling the raw powder prepared by solid state reaction. Those unique nano-materials were tested in electrocatalysis for OER. In particular, Ni0.8Cu0.2TiO3-nano oxide is a highly active and stable material for O2 evolution with overpotentials of 345 mV and 470 mV at 10 mA cm-2 observed in 1.0 M and 0.1 M KOH respectively. This high performance in electrochemical oxidation of water is due both to the high electrochemically active surface area of the material and to the cupric ions.

ASSOCIATED CONTENT

Supporting information CCDC 1915348 (Ni0.95Cu0.05TiO3), CCDC 1915335 (Ni0.90Cu0.10TiO3), CCDC 1915349 (Ni0.85Cu0.15TiO3),

CCDC

1915350

(Ni0.8Cu0.2TiO3)

contain

the

supplementary

crystallographic data for this paper. These data can be obtained free of charge via


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www.ccdc.cam.ac.uk/data_request/cif, by emailing [email protected], or by contacting The Cambridge Crystallographic Data Centre, 12 Union Road, Cambridge CB2 1EZ, UK; fax: +44 1223 336033.

AUTHOR INFORMATION

Corresponding Author *E-mail:[email protected]

*E-mail:[email protected]

ORCID Amandine Guiet: 0000-0001-7590-1119 Christophe Payen: 0000-0003-4189-6838 Florence Porcher: 0000-0002-8873-6498 Victor Mougel:0000-0003-4136-7442 Marc Fontecave: 0000-0002-8016-4747 Gwenaël Corbel: 0000-0003-2605-7702


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Author Contributions ‡AG and TNH contributed equally to this work.

ACKNOWLEDGMENT AG and GC greatly acknowledge the "X-ray Diffusion and Diffraction" and the "Electron Microcopy" technical platforms of IMMM (Le Mans University). AG and GC acknowledge Dr Brigitte BOULARD for her help in peak deconvolution of UV-vis absorption spectra and Marie-Pierre Crosnier-Lopez for electronic diffraction. AG and GC are grateful to the M Sc. trainee Valere Audric WOUAMBA for preliminary syntheses.

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Copper-Substituted NiTiO3 Ilmenite-Type Materials for Oxygen

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