NiCo2O4 is a p‐type semiconductor consisting of environmentally friendly, abundant elements and higher conductivity compared to NiO. It is shown that the combustion synthesis of spinel NiCo2O4 using tartaric acid as fuel can be used to control the NPs size and provide smooth, compact, and homogeneous functional HTLs processed by blade coating. Study of PVSCs with different NiCo2O4 thickness as HTL reveals a difference on hole extraction efficiency, and for 15 nm, optimized thickness enhanced hole carrier collection is achieved. As a result, p‐i‐n structure of PVSCs with 15 nm NiCo2O4 HTLs shows reliable performance and power conversion efficiency values in the range of 15.5% with negligible hysteresis.The synthesis and characterization of low‐temperature solution‐processable monodispersed nickel cobaltite (NiCo However, in order to achieve the required crystallinity, temperatures above 400 °C are usually required. The need for high temperature is increasing the fabrication cost and limits their potential use for printed electronic applications.Among many processes for the synthesis of nanomaterial compounds, combustion synthesis emerges as an efficient alternative approach. The combustion synthesis, in principles, can be defined as a redox (reduction/oxidation) or electron transfer process, in which the fuel is oxidized (increase of the oxidation state) and the oxidizer is reduced (reduce of the oxidation state) in an exothermic reaction.2O content of the precursor blend.The combustion technique appears to be versatile and effective for the synthesis of high crystallinity solution‐processed metal oxides thin films using low temperature.2O4) is a p‐type transparent conductive oxide semiconductor consisting of abundant and environmentally friendly elements , with a relatively wide optical band gap (≈2.1–2.4 eV), deep‐lying valence band (VB of 5.3 eV) that matches well with the VB of CH3NH3PbI3 perovskite semiconductor and a much better conductivity than NiO and Co3O4 (at least two orders of magnitude higher).2O4 one of the most promising candidates for electronic applications. NiCo2O4 adopts a cubic spinel structure in which all the Ni ions occupy the octahedral sites and the Co ions are distributed between the tetrahedral and octahedral sites.2O4 an appropriate candidate material for introduction as HTL in PVSCs to achieve high‐performance photovoltaic devices. NiCo2O4 derivatives have been used previously in many other applications such as anodic oxygen evolution,Nickel cobaltite  via combustion chemistry proposing for the first‐time tartaric acid as a fuel and nitrate as an oxidizer agent. NiCo2O4 NPs with an average size of ≈4 nm and narrow particle‐size distribution were prepared using a cost‐effective, low‐temperature combustion synthesis method calcinated at 250 °C for 1 h. Those ultrafine NPs enable the formation of compact, very smooth, high electrically conductive, and relatively optically transparent NiCo2O4 films, which were utilized, for the first time, as HTLs in a solution‐processed p‐i‐n PVSC. The effect of NiCo2O4‐NPs HTL thickness on PVSC characteristics is also investigated. A comparative study of devices incorporating different thickness of NiCo2O4‐HTLs reveal a difference in hole extraction efficiency. The photoluminance spectroscopy measurements on perovskite films showed a reduced electron–hole pair recombination for the optimized 15 nm thick NiCo2O4‐HTL. Additional electroimpedance spectroscopy and Mott–Schottky (M‐S) measurements on PVSC confirm the better hole extraction inducing an enhancement in the PVSC characteristics and negligible PCE hysteresis. The corresponding PVSC exhibits a high fill factor (FF) ≈ 80% and a reliable performance with PCE of 15.5%. The results demonstrate the great potentials of applying the low‐temperature combustion synthesis for fabrication of highly reproducible and reliable metal oxide NPs which can be used for the formation of HTLs in variety of solution‐processed printed electronic devices.In this work, we present a one‐step synthesis of low‐temperature solution‐processable nickel cobaltite  come close together.3.In this work, tartaric acid is proposed as a fuel that is critically important to obtain uniform single‐crystalline phase NiCo2.1It is well documented in the literature that thin metal oxide films can be obtained at temperatures lower than bulk‐like powders via combustion synthesis because of the enhanced gas transport mechanisms and the easier out‐diffusion of volatile products.2O4, metal precursors (Ni and Co nitrates) and tartaric acid  were dissolved in 2‐methoxyethanol containing a small amount of 0.24 × 10−3m HNO3  by using the Scherrer's equation and was found to be ≈3.5 nm for the sample annealed at 250 °C and ≈5 nm for the sample annealed at 300 °C. In contrast, the XRD pattern of the material obtained after 200 °C heat treatment showed no diffraction peaks, indicating the formation of an amorphous structure.In order to crystallize the as‐synthesized films to spinel phase, 1 h of heating was applied at different temperatures, i.e., 200, 250, and 300 °C, and the X‐ray diffraction (XRD) results of the obtained materials are shown in 2O4 NPs prepared by the low‐temperature combustion method. Figure 2O4 NPs synthesized at 250 °C. It can be seen that this material is composed of individual NPs with an average diameter of 4 ± 1.3 nm, which is very close to the grain size calculated from XRD patterns. The high‐resolution TEM image shown in Figure 2O4 NPs possess a single‐phase spinel structure with high crystallinity; combined with XRD results, the observed lattice fringes with interplanar distances ≈2.4 and ≈2.8 Å can be assigned to the (331) and (220) crystal planes of spinel NiCo2O4, respectively. The crystal structure of the NiCo2O4 NPs was further studied by selected‐area electron diffraction (SAED). The SAED pattern taken from a small area of the NiCo2O4 NP aggregates  showed an overall Ni:Co atomic ratio close to 1:2, in agreement with the stoichiometry of NiCo2O4 compound .Transmission electron microscope (TEM) verified the high crystallinity and phase purity of spinel NiCo2.22O4 NPs were produced on top of quartz and indium tin oxide (ITO) substrates using the doctor‐blading technique, the processing parameters are described within the Experimental Section.Figure2O4 film fabricated on top of glass/ITO and quartz substrates, as obtained by atomic force microscopy (AFM) scans. On top of ITO substrate 2 vs. photon energy (E)), giving an optical band gap of 2.32 eV. NiCo2O4 thin films of different thicknesses were also fabricated on glass/ITO substrate in order to investigate the transmittance of the front contact at UV–vis spectrum. Figure 2O4 HTLs coated on glass/ITO substrate; it could be seen that NiCo2O4 films thinner than 20 nm reduce only slightly the transparency of the glass/ITO substrate for wavelengths longer than 450 nm, allowing more intense light to reach the absorbing layer.2.32O4 HTLs with three different thicknesses, 30, 20, and 15 nm. On top of each NiCo2O4 HTL, a 230 nm thick perovskite film was solution processed as described in the Experimental Section. The deposited perovskite film exhibits a low roughness of 5.4 nm  and a mean grain size of 0.22 µm with a standard deviation of 0.051 µm , as calculated by AFM topography measurements. To complete the devices, a PC[70]BM film was spin‐coated on the top of the perovskite semiconductor serving as the electron transporting layer followed by a 100 nm thick thermally evaporated Al  of the PVSCs fabricated with 15, 20, and 30 nm thick NiCo2O4 HTLs, and the corresponding solar cell parameters are summarized in TableRs) was extracted from the dark J–V curves  showed the lowest increase (≈8%) for a forward sweep from 18.47 to 19.94 mA cm−2, compared to both Voc and FF. Consequently, the device consisting of a 15 nm thick NiCo2O4 HTL exhibits a PCE as high as 15.5% for the forward sweep. The PCE of devices with thinner NiCo2O4 HTLs were declined (not shown here) exhibiting high leakage currents due to limitations of not fully covered ITO. Figure S4 (Supporting Information) demonstrates the external quantum efficiency (EQE) measurements of the corresponding devices. It is noticed that for 15 nm thick NiCo2O4 film the overall efficiency is increased comparing to thicker layers due to higher transmittance as well as to a better charge collection, as it will be shown below. All the devices show a declined performance at longer wavelength (600–750 nm) which can be attributed to the relatively thin perovskite layer (≈250 nm) and to not optimized electron carrier selectivity of the top electrode (PC[70]BM/Al) that used within this studies.Complete p‐i‐n architectures of PVSCs were fabricated employing NiCoAl Figure5a. Fig2O4 HTL thickness on the ITO/NiCo2O4‐NPs/perovskite device performance was evaluated by photoluminescence (PL) spectroscopy , the PL signal of the devices with NiCo2O4 HTL show a quenching of more than 90%, indicating a great reduction in the band‐to‐band charge recombination and, thus, a better hole carrier selectivity of the ITO electrode covered by NiCo2O4. Further, the PL intensity is lower in 15 nm NiCo2O4 film than that of the thicker films, pointing to an efficient suppression of the electron–hole recombination  measurements under solar light and zero bias. Figure2O4 films. The results showed a shape of two frequency responses for PVSCs, where the second semicircle (feature at low frequencies) is been attributed to the recombination resistance (Rrec).2O4 HTL thickness is reduced the radius of the semicircle increases, this parameter implies a higher resistance in the charge recombination, in agreement with the findings from PL measurements. Figure C2 = 0 is attributed to the flat band potential of the device, while the lower slope of the linear region is attributed to the charge accumulation at the interfaces, which impedes an efficient extraction of the charge carriers.2O4 HTL seems to extract faster the charge carriers, increasing the flat band potential, and thus the Voc, as well as the FF of the corresponding device. The direct correlation of faster charge carriers extraction at thinner HTL layers with increased FF has been previously studied by Stolterfoht et al.2O4 HTL is depleted faster than the other devices due to enhanced charge carrier collection, which can also confirm the increase at the FF of the corresponding PVSC. The increased Jsc for the 15 nm NiCo2O4 layer can be ascribed to the higher transparency of the thin film compared to thicker HTLs (shown above) resulting to enhanced photogeneration of electron–hole pairs as well as to the lower series resistance as shown in Table To further understand the charge recombination processes during the hole collection process from perovskite to NiCo2O4‐HTL p‐i‐n PVSCs, compare to our previously reported solvothermal synthesized CuO‐HTL based PVSCs2O4 HTLs. Specifically, as shown in Figure S5 (Supporting Information) the combustion synthesized NiCo2O4‐HTL delivers a ≈14.5% average PCE (16 devices) with 15.5% best performing device, while CuO‐HTL based solar cells give a ≈12.5% average PCE (16 devices) with 15.3% best performing device.Importantly, improved reproducibility and performance reliability of the NiCo32O4 NPs. The size of the NPs was fully controlled due to the usage of tartaric acid leading to the formation of monodispersed NiCo2O4 NPs with a diameter of ≈4 nm. The combustion proceeds under low temperature (250 °C) and within a short reaction time (1 h), produce high‐quality, homogeneous NiCo2O4 NPs films with high electrical conductivity (≈4 S cm−1) and very low roughness (0.56 nm) functional layers were fabricated. The detailed physicochemical characterization of the NiCo2O4 NPs using XRD, EDS, and electron microscopy measurements confirm the high purity, crystallinity, and small grain composition of the NiCo2O4. Furthermore, the proposed synthetic approach allowed the production of compact films using blade coating, which is a large‐scale compatible technique appropriate for the development of printed electronic devices. The impact of NiCo2O4 HTL thicknesses on PVSCs characteristics was also investigated. The optimum thickness is found to be 15 nm showing enhanced charge carrier collection and negligible J–V hysterics, compared to thicker films, delivering reliable p‐i‐n PVSCs with a PCE of 15.5%. We believe that the proposed combustion synthesis method using a tartaric acid as a fuel can provide a route to produce highly reproducible metal oxides suitable for use in a range of advanced materials applications.In conclusion, a low‐temperature combustion synthesis method, using for the first time a tartaric acid as a fuel, was successfully developed and applied for the fabrication of compact films of p‐type NiCo4Materials: Prepatterned glass‐ITO substrates (sheet resistance 4Ω sq−1) were purchased from Psiotec Ltd., Pb(CH3CO2)2.3H2O from Alfa Aesar, methylammonium iodide (MAI) and methylamonium bromide (MABr) from Dyenamo Ltd., PC[70]BM from Solenne BV. All the other chemicals used in this study were purchased from Sigma‐Aldrich.2O4 NPs FilmsSynthesis of NiCo: For the combustion synthesis of NiCo2O4 NPs, 0.5 mmol Ni(NO3)2.6H2O, 1 mmol Co(NO3)2.6H2O, and tartaric acid were mixed in the 15 mL 2‐methoxy ethanol solution. After 150 uL HNO3 (69 wt% HNO3) were added slowly into the mixture, and the solution stirred up to almost complete homogeneity. The whole solution was allowed under stirring for 30 min at 60 °C. The ratio of the total metal nitrates and tartaric acid was 1. Thereafter, the violet colored solution was used for the combustion synthesis of the NiCo2O4 NPs on the various substrates. Doctor blade technique was applied for the fabrication of the precursor films on the various substrates. The resulting light violet colored films were dried at 100 °C for 30 min, and used as a precursor for the combustion synthesis of NiCo2O4 NPs. Subsequently, the obtained films were heated at different temperatures  in ambient atmosphere for 1 h in a preheated oven to complete the combustion process and then left to cool down at room temperature. For UV–vis absorption measurements, the films were fabricated on quartz substrates, while for the transmittance measurements 30, 20, and 15 nm thick films were fabricated on glass/ITO substrates applying 250 °C heating temperature, respectively.Device Fabrication: The inverted solar cells under study was ITO/NiCo2O4‐NPs/CH3NH3PbI3/PC[70]BM/Al. ITO substrates were sonicated in acetone and subsequently in isopropanol for 10 min and heated at 100 °C on a hot plate 10 min before use. The perovskite solution was prepared 30 min prior spin coating by mixing Pb(CH3CO2)2.3H2O:methylamonium iodide (1:3) at 36 wt% in dimethylformamide (DMF) with the addition of 1.5% mole of MABr.−1 in chlorobenzene, was dynamically spin coated on the perovskite layer at 1000 rpm for 30 s. Finally, 100 nm Al layers were thermally evaporated through a shadow mask to finalize the devices giving an active area of 0.9 mm2. Encapsulation was applied directly after evaporation in the glove box using a glass coverslip and an Ossila E131 encapsulation epoxy resin activated by 365 nm UV irradiation.Characterization: TGA and differential thermal analysis (DTA) were performed on a Shimadzu Simultaneous DTA‐TG system (DTG‐60H). Thermal analysis was conducted from 40 to 600 °C in air atmosphere (200 mL min−1 flow rate) with a heating rate of 10 °C min−1. XRD patterns were collected on a PANanalytical X'pert Pro MPD powder diffractometer  using Cu Kα radiation (λ = 1.5418 Å). TEM images and electron diffraction patterns were recorded on a JEOL JEM‐2100 microscope with an acceleration voltage of 200 kV. The samples were first gently ground, suspended in ethanol and then picked up on a carbon‐coated Cu grid. Quantitative microprobe analyses were performed on a JEOL JSM‐6390LV scanning electron microscope equipped with an Oxford INCA PentaFET‐x3 EDS detector. Data acquisition was performed with an accelerating voltage of 20 kV and 60 s accumulation time. Transmittance and absorption measurements were performed with a Schimadzu UV‐2700 UV–vis spectrophotometer. The thickness of the films was measured with a Veeco Dektak 150 profilometer. The current density–voltage (J/V) characteristics were characterized with a Botest LIV Functionality Test System. Both forward and reverse scans were measured with 10 mV voltage steps and 40 ms of delay time. For illumination, a calibrated Newport Solar simulator equipped with a Xe lamp was used, providing an AM1.5G spectrum at 100 mW cm−2 as measured by a certified oriel 91150 V calibration cell. A shadow mask was attached to each device prior to measurements to accurately define the corresponding device area. EQE measurements were performed by Newport System, Model 70356_70316NS. AFM images were obtained using a Nanosurf easy scan 2 controller under the tapping mode. Electrical conductivity measurements were performed using a four‐point microposition probe, Jandel MODEL RM3000. EIS and M‐S measurements were performed using a Metrohm Autolab PGSTAT 302N, where for the EIS a red light‐emitting diode (at 625 nm) was used as the light source calibrated to 100 mW cm−2. For EIS a small AC perturbation of 20 mV was applied to the devices, and the different current output was measured throughout a frequency range of 1 MHz to 1 Hz. The steady state DC bias was kept at 0 V throughout the EIS experiments.The authors declare no conflict of interest.SupplementaryClick here for additional data file.


Patiria pectinifera (phylum Echinodermata). Furthermore, analysis of transcriptome sequence data from the starfish Asterias rubens revealed two PP/OK‐type precursors: an SMP‐type precursor  and a second precursor (ArPPLNP2). We reported previously a detailed analysis of ArPPLNP1 expression in A. rubens and here we report the first functional characterization ArPPLNP2‐derived neuropeptides. Sequencing of a cDNA encoding ArPPLNP2 revealed that it comprises eleven related neuropeptides (ArPPLN2a‐k), the structures of several of which were confirmed using mass spectrometry. Analysis of the expression of ArPPLNP2 and neuropeptides derived from this precursor using mRNA in situ hybridization and immunohistochemistry revealed a widespread distribution, including expression in radial nerve cords, circumoral nerve ring, digestive system, tube feet and innervation of interossicular muscles. In vitro pharmacology revealed that the ArPPLNP2‐derived neuropeptide ArPPLN2h has no effect on the contractility of tube feet or the body wall‐associated apical muscle, contrasting with the relaxing effect of ArPPLN1b (ArSMP) on these preparations. ArPPLN2h does, however, cause dose‐dependent relaxation of cardiac stomach preparations, with greater potency/efficacy than ArPPLN1b and with similar potency/efficacy to the SALMFamide neuropeptide S2. In conclusion, there are similarities in the expression patterns of ArPPLNP1 and ArPPLNP2 but our data also indicate specialization in the roles of neuropeptides derived from these two PP/OK‐type precursors in starfish.Molluscan pedal peptides (PPs) and arthropod orcokinins (OKs) are prototypes of a family of neuropeptides that have been identified in several phyla. Recently, starfish myorelaxant peptide (SMP) was identified as a PP/OK‐type neuropeptide in the starfish  Thus, the evolutionary origin of PP/OK‐type neuropeptides can be traced to the bilaterian common ancestor of protostomes and deuterostomes Jekely, . HoweverAplysia californica on account of its expression by neurons in the pedal ganglia  that encode precursors of PP‐like neuropeptides  provided the first evidence that PP/OK‐type peptides also occur in deuterostomes  was identified as a PP/OK‐type neuropeptide  and we have named the other A. rubens PP/OK‐type precursor A. rubens PP‐like neuropeptide precursor 2 (ArPPLNP2).A homolog of the A. rubens using mRNA in situ hybridization and immunohistochemistry  with a diameter >4 cm were collected at low tide from the Thanet coast  or were obtained from a fisherman based at Whitstable . These animals were maintained in a circulating seawater aquarium at ∼12°C in the School of Biological and Chemical Sciences at Queen Mary University of London and were fed on mussels (Mytilus edulis). Smaller juvenile specimens of A. rubens (diameter 0.5–1.5 cm) were collected at the University of Gothenberg Sven Lovén Centre for Marine Infrastructure  and fixed in Bouin's solution.Starfish  of PP‐like neuropeptide precursor 2 (ArPPLNP2), which comprises a 31‐residue N‐terminal signal peptide and seven putative neuropeptides bounded by dibasic cleavage sites  and a contig (GAUS01023767.1) comprising the 3′ region of PPLNP2 was identified in Asterias forbesi. Combining the partial PPLNP2 transcript sequences from A. rubens and A. forbesi, forward and reverse primers  were designed to enable PCR amplification of a full‐length ArPPLNP2 cDNA from A. rubens.Analysis of The methods used for RNA extraction, cloning and sequencing were as described previously  visualization of ArPPLNP2 transcripts in sections of A. rubens were the same as reported previously in a study investigating expression of the relaxin‐like gonad‐stimulating peptide precursor in A. rubens (AruRGPP)  production of anti‐sense and sense DIG‐labeled RNA probes, (b) preparation of sections of fixed specimens (diameter 4–6 cm) of 2.5The peptide ArPPLN2h (GRTSLSGSSGLTHLSSGFH) was selected as a target antigen for production of rabbit antibodies because it is the most abundant of the ArPPLN2‐derived neuropeptides, occurring in triplicate whilst other peptides occur singly in the precursor. Furthermore, ArPPLN2h shares between 58 and 95% sequence identity with the other ArPPLNP2‐derived neuropeptides. A peptide comprising the C‐terminal octapeptide of ArPPLN2h but with the addition of N‐terminal lysine and tyrosine residues (KYTHLSSGFH) was designed as an antigen (ag) peptide (ArPPLN2h‐ag) and synthesized by Peptide Protein Research Ltd. . Rabbit immunization with the ArPPLN2h‐ag–thyroglobulin conjugate and serum collection was performed by Charles River Labs , as described previously for ArPPLN1b , employing the same methods as described previously for ArPPLN1b  visualization of ArPPLN2h expression in sections of A. rubens using immunohistochemistry (including pre‐absorption tests) and (c) photography of immunostained sections and preparation of montages were the same as those reported previously for ArPPLN1b  preparation of sections of fixed specimens (diameter 4–6 cm) of 2.7A. rubens, using the same methods reported previously for ArPPLNP1b , which is a known relaxant of the cardiac stomach, apical muscle, and tube feet in A. rubens  with 100 µL of 1 mM ArPPLN2h.Consistent with the selection of ArPPLN2h as a representative neuropeptide for generation of antibodies (see above), ArPPLN2h was also selected as a representative ArPPLNP2‐derived peptide for in vitro and in vivo pharmacological experiments. ArPPLN2h was synthesized by Peptide Protein Research Ltd.  and was tested on in vitro preparations of the cardiac stomach, apical muscle and tube foot from 33.15 located in the N‐terminal region and Leu14 and Phe18 located in the C‐terminal region . The ArPPLNP2 ORF encodes a 563 amino acid residue protein comprising a 31‐residue signal peptide and 13 copies of predicted PP/OK‐type peptides, which we have named ArPPLN2a‐k according to their position in the precursor protein. There are three copies of ArPPLN2h and one copy of each of the other peptides  on three in vitro preparations from A. rubens: apical muscle, tube feet and cardiac stomach. No effects on the contractile state of apical muscle and tube foot preparations were observed, whereas ArPPLN2h caused relaxation of cardiac stomach preparations. ArPPLN2h caused dose‐dependent relaxation of cardiac stomach preparations and comparison of the effects of ArPPLN2h with the SALMFamide neuropeptide S2 revealed that the potency and efficacy of ArPPLN2h was very similar to that of S2  observed at 10−6 M was 40.05 ± 5.07% and 48.03 ± 8.17% reversal of KCl‐induced contraction for ArPPLN2h and S2, respectively.ArPPLN2h was tested at a high concentration . ArPPLNP2 comprises 13 copies of 11 PP/OK‐type peptides (ArPPLN2a‐k), with three copies of ArPPLN2h and one copy each of the other ten peptides. Alignment of ArPPLNP2‐derived peptides with other PP/OK‐type neuropeptides reveals several conserved structural characteristics , although it is thought to comprise sensory neurons, interneurons and motor neurons   and in A. planci  suggests that both of these PP/OK‐type precursors would have been present in the common ancestor of all extant starfish. Previous studies have shown that PPLN1‐type (SMP‐type) peptides act as muscle relaxants in P. pectinifera  and therefore it seems likely, based on asteroid phylogeny . It will be interesting, therefore, to investigate if PPLN2‐type peptides also act as muscle relaxants in starfish that belong to other starfish orders—for example, in P. pectinifera and/or A. planci, both of which are species belonging to the order Valvatida. Looking beyond starfish, it would be of interest to determine if PP/OK‐type neuropeptides act as muscle relaxants in other echinoderms Addressing this issue is feasible because precursors of PP/OK‐type neuropeptides have been identified in other echinoderms, including two precursors in the sea urchin S. purpuratus (SpPPLNP1 and SpPPLNP2; Rowe & Elphick, A. japonicus (Rowe, Achhala, & Elphick, This study of PPLN2‐type peptides in As discussed in detail previously (Lin, Egertová et al., The authors declare that they have no conflict of interest.


In the present study, we employed flat, thin CoPx electrodes prepared through the thermal phosphidation (PH3) of Co3O4 films made by plasma‐enhanced atomic layer deposition to evaluate their stability in acidic water electrolysis by using a multi‐technique approach. The films were found to be composed of two phases: CoP in the bulk and a P‐rich surface CoPx (P/Co>1). Their performance was evaluated in the HER and the exchange current density was determined to be j0=−8.9 ⋅ 10−5 A/cm2. The apparent activation energy of HER on CoPx (Ea=81±15 kJ/mol) was determined for the first time. Dissolution of the material in 0.5 M H2SO4 was observed, regardless of the constantly applied cathodic potential, pointing towards a chemical instead of an electrochemical origin of the observed cathodic instability. The current density and HER faradaic efficiency (FE) were found to be stable during chronoamperometric treatment, as the chemical composition of the HER‐active phase remained unchanged. On the contrary, a dynamic potential change performed in a repeated way facilitated dissolution of the film, yielding its complete degradation within 5 h. There, the FE was also found to be changing. An oxidative route of CoPx dissolution has also been proposed.Cobalt phosphides are an emerging earth‐abundant alternative to platinum‐group‐metal‐based electrocatalysts for the hydrogen evolution reaction (HER). Yet, their stability is inferior to platinum and compromises the large‐scale applicability of CoP It is also a promising clean energy carrier to replace fossil‐derived fuels in the near future.+ and atomic H, while still being able to efficiently desorb H2.+ acceptor) “dilutes” metal atoms (hydride acceptors) resulting in a weakening of the M−H bond (i. e. ensemble effect).Electrocatalysts play a critical role in the performance of electrochemical water splitting devices.x in acid electrolyte. We deposited 20 nm thick cobalt oxide films on flat Au‐coated Si substrates by plasma enhanced atomic layer deposition (PE‐ALD). We chose PE‐ALD to deposit cobalt oxide films because of its low processing temperatures, sub‐nm thickness control and 3D conformity. Then, we converted the cobalt oxide films to cobalt phosphide (CoPx) by gas‐phase phosphidation. With such a low mass loading on a flat surface, we were able to study the intrinsic stability and activity of CoPx. The apparent stability of CoPx measured by repeated LSV is much lower than that measured by CA. Combining these results with compositional analysis of both electrodes and electrolytes after stability testing, we propose a possible degradation mechanism of cobalt phosphide electrodes in acidic electrolytes. We also studied the temperature‐dependent kinetics from which the apparent activation energy for HER over CoPx in acidic electrolyte has been derived. In particular, On‐Line Electrochemical Mass Spectroscopy (OLEMS) was employed to follow gaseous products evolution simultaneously to the applied electrochemical treatment. Combined EC‐OLEMS measurements are commonly used in the field of electrochemistry for an accurate evaluation of electrode performance, stability, selectivity and reaction kinetics.In the present work, we studied the intrinsic HER activity and stability of CoP22.1x films were thermally prepared by phosphidation of smooth PE‐ALD fabricated Co3O4 films with an initial thickness of 20 nm. The Co 2p XP spectrum of as‐prepared cobalt oxide  confirms its chemical composition by showing the characteristic Co3O4 profile and the matching binding energy (BE) positions of spectral lines (BE(Co 2p3/2)=779.6 eV) and shake‐up satellites =789.9 eV).3O4 films with a uniform coverage. Additionally, atomic force microscopy (AFM) shows that the average surface roughness of the as‐deposited Co3O4 film is σ=0.7 nm – the same as for the bare Au/Ti/Si substrate .CoP3O4 undergoes both chemical and morphological changes. The morphology analysis of the CoPx film (400 °C) shows that the film is composed of spherical particles with an average lateral size of 84±10 nm and an average height of 18±4 nm  shows that phosphidation of Co3O4 involves agglomeration of the particles composing the initial film.Upon thermal phosphidation treatment conducted at temperatures ranging from 350–450 °C, Co3O4 (111) and (511), while other reflections are either not observable or overlapping with substrate peaks . Phosphidation leads to disappearance of oxide reflections, although no visible reflections attributed to any other new phase  were found to be formed. This can be explained by both low thickness and crystallinity of the phosphide film formed during the relatively mild thermal phosphidation. Thus, additional characterization was performed to determine the structural and chemical composition of the CoPx film.The XRD pattern of the untreated oxide film reveals the presence of two weak reflections attributed to Cox phase (BE(Co 2p3/2)=778.2±0.2 eV) in which Co is assumed to be mostly covalently bound to P with corresponding charges δ+ (∼1+) and δ− (∼1−), respectively  >779.6 eV) is most likely attributed to a native oxide formed on the surface of CoPx upon air exposure. The P 2p XP spectral region reveals that the major (76±4 %) fraction of the surface phosphorous is present in Pδ− state (BE(P 2p3/2)=128.6±0.1 eV) with some part of phosphide being oxidized to form a mixture of P0 (BE(P 2p3/2)=130.2±0.2 eV) and PO43− (BE(P 2p3/2)=133.3±0.1 eV) . Thus, all further measurements were conducted on the sample prepared at 400 °C.XPS analysis of the Co 2p spectral region reveals the formation of a dominant P‐rich CoPrges δ+ ∼+ and δ− x has been further studied by XPS depth profiling (DP‐XPS). Depth profiling of the fresh sample points to a difference between surface and bulk composition of CoPx . Although, cobalt phosphides normally crystalize in the form of CoP and Co2P – the formation of P‐rich polyphosphides is a known phenomenon in phosphorous chemistry.3 and lower phosphides in the top layer of the film. Thus, we assume that formation of polyphosphides as well as other P‐containing species (e. g. P0) is the reason for the observed phosphorus excess. Unfortunately, an exact thickness analysis of DP‐XPS requires information on the sputtering rates of the studied materials, which is unknown for CoPx. Thus, we could only estimate the percentage of the P‐rich fraction  in the overall CoPx film . Qualitative analysis of DP‐XPS reveals the presence of oxidized Co species in the bulk (tS=4 min) which can be attributed to unreacted Co3O4 . Together with that, the presence of P0 and residuals of CoPOx were also observed.The bulk structure of CoPx Figure a. Partic3O4 structure  and the appearance of a new set of peaks corresponding to CoPx  and HeNe laser impurities . The Raman peaks, which we attributed to CoPx, match with the Raman spectrum of electrodeposited CoP/CoPxOy mixed electrodes reported by Saadi et al.3O4 powder  allowing us to measure the Raman spectrum of CoP . Based on this, we assign CoP to be the major component in the bulk of the CoPx film.The Raman spectrum of the as‐prepared PE‐ALD film also confirms its Co8P8 cluster with the structure depicted in Figure S6a matching literature data.−1. We notice a consistent but small shift of about 10–20 cm−1 to higher wavenumbers for every peak in the computed spectrum as compared to the measured spectrum. This can be explained by methodological differences between VASP and Gaussian software utilized for structural optimization and Raman spectrum simulation, respectively. Moreover, the computations were performed on an isolated cluster leading to different electronic properties as compared to the actually measured bulk CoP phase. However, we aim at a qualitative description of the Raman spectrum and the correspondence between the measured and computed spectrum is considered to be good.Additionally, we utilized DFT calculations to obtain a simulated Raman spectrum of CoP. Calculations were done on a Co2.2x films, faradaic efficiency (FE), the exchange current density and the apparent activation energy were measured by CA and OLEMS. For this purpose, a series of potential pulses with a duration of 10 s was applied and both current density and m/z=2 (H2+) ion current were recorded simultaneously. Pulsed chronoamperometric measurements were chosen over continuous voltammograms since it enables longer accumulation of reaction products which is especially crucial for OLEMS measurements of FE at lower overpotentials.In order to evaluate the HER performance of the as‐synthesized CoP2 which are proportional to m/z=2 peak area were plotted versus values of negative charge generated during the pulses to make sure that faradaic efficiency is not changing during the experiment . In the case of the T‐dependent pulse CA treatment of CoPx cathodes accompanied by in‐between protective potentials , a single Faraday slope was observed. Thus, the FE of HER on CoPx was found to be constant within ±10 % under the chosen experimental conditions, while in a reference experiment, where no protective currents were applied between the temperature steps, a clear alteration of the Faraday slopes could be observed .The amounts of evolved HT‐dependent I–V curves were converted to Tafel plots in order to assess the kinetics of the HER on CoPx  can be achieved on nanostructured CoPx cathodes with high surface area, while flat electrodes had similar values to what is reported in our study.The obtained x Figure a. The Taj0) of the T‐dependent measurements were determined by extrapolation of the linear Tafel region to η=0 V. j0 at 25 °C (8.9 ⋅ 10−5 A/cm2) was found to be comparable with reported values for CoP (0.4–1.2 ⋅ 10−4 A/cm2).−5 (28 °C), 1.6 ⋅ 10−4 (31 °C) and 2.5 ⋅ 10−4 (35 °C) A/cm2). The j0 values were plotted in the form of an Arrhenius plot  of the HER, based on Equation (4), where Bk and h are the Boltzmann and Planck constants, z and e are the number of electrons and the elementary charge and Nas is the number of surface active sites.In addition to the Tafel slope values, the exchange current densities ,x electrocatalyst, different strategies could be applied, including an increase of the P content in the material (e. g. CoP3), an introduction of unsaturated P sites and nanostructuring of the surface.Although, the HER performance of CoP2.3x films under HER conditions, a series of experiments with different types of electrochemical treatments was performed. At first, the stability of CoPx under continuous operation was evaluated by subjecting the electrode to a constant potential (E=−0.12 V) for 5 h .In order to assess the stability of CoP2 production  but did not lead to an overproduction of H2 bubbles which would disturb the measured current. In the second experiment, a series of repeated linear cathodic sweeps  was applied to a fresh CoPx sample (rLSV). The starting potential was set equal to the open circuit potential (OCP) of fresh CoPx in order to avoid anodic currents, while the end potential lies within the Tafel region of the voltammogram. rLSV treatment allows evaluation of the electrode stability under fast potential scanning (20 mV/s).This potential value was chosen since it was found to be negative enough to assure a decent rate of Hx under constant cathodic polarization. Only a slight current decrease (17 %) was observed in the first 2 hours of electrolysis arriving at a plateau current density of −0.35 mA/cm2 . In about one hour, the cathodic current increased and reached its maximum value of −0.74 mA/cm2. After passing this point, the trend changed to the opposite and the cathodic current began to decrease, eventually arriving at a constant value of −0.37 mA/cm2  at 4.3 h. Together with that, unlike in the CA case, the appearance of the Au layer (i. e. substrate) was observed visually and confirmed by XPS after 5 h of rLSV treatment . In the case of CoPx electrodes, a cathodic current increase can potentially be attributed to a change of HER faradaic efficiency, electrode dissolution or capacitive charging of the surface. Structural and compositional changes of the spent electrodes caused by the performed electrochemical treatments were analyzed by a comprehensive set of characterization techniques.On the other hand, the currents derived from the rLSV experiment Figure b point tx electrode reveals only a slight change in intensity of the Coδ+ peak (BE(Co 2p3/2)=778.1 eV), while the position and line shape were found to be preserved  appearing at higher BE (780.6 eV). A similar trend was found in the P 2p spectral region =132.5 eV (oxidation state of P is below 5+) while the fresh sample had mainly PO43− species with a higher BE value (133.3 eV). Other components of P 2p spectra are matching with the fresh sample (Pδ−: 128.7 eV and P0: 130.8 eV). In contrast, the rLSV‐treated sample shows the absence of phosphide‐related spectral components with only a minor amount of residual PO43− (133.6 eV), supporting the statement that most of the CoPx film was dissolved during rLSV treatment.The sample subjected to the CA treatment shows a similar amount of overall P species compared to the fresh sample, with a slight change in the composition of the native oxide. Particularly, the P 2p spectrum of the CA‐treated CoPx dissolution rates, ICP‐OES measurements of spent electrolytes were performed. These measurements reveal the presence of Co ions in all tested solutions, meaning that electrodes were corroded to a certain degree regardless the type of applied EC treatment. However, the P content could not be evaluated accurately due to the much lower detection limit of phosphorus in ICP‐OES. In the case of the rLSV‐treated electrode, the content of Con+ ions was found to be 181±6 nmol/cm2 which can roughly be attributed to the overall Co content of the initial CoPx film. This was also confirmed by XPS depth profiling  which only shows negligible amounts of Co and P present on the surface of the rLSV‐treated electrode.In order to estimate CoP2), which is only 7.2±0.8 % of the cobalt released in the rLSV experiment. XPS depth profiling of this sample also shows a slight decrease in the amount of the P‐rich fraction  compared to the fresh sample (tS=2.4 min) . The Raman spectrum of the CA‐treated electrode suggests preservation of the CoPx structure with only a change in intensity of the peak at 320 cm−1 . The dissolution rate at the initial stage of electrolysis (t<2.5 h) was not accessible in this way due to the low concentration of Co ions in the electrolyte. We found that Co leaching occurred throughout the entire experiment with a slightly higher rate (∼17 %) in the first sampled period (2.5 hours) of electrolysis. High initial dissolution rates typically originate from dissolution of native oxides and can be quantified by e. g. online ICP‐OES or ‐MS analysis.2P, while following potentiostatic treatment led to almost zero Co leaching.To assess the temporal evolution of Co leaching during CA treatment at x under resting conditions , a reference experiment was conducted in which a fresh CoPx substrate was immersed in 0.5 M H2SO4 solution for 5 h without any additional EC treatment. A constant dissolution rate of CoPx  was found which is comparable with that for the CA‐treated sample . From this, we can conclude that CoPx dissolution, observed in the case of the CA treatment, proceeds mostly via a chemical route with additional involvement of anodic oxidation during the HER in accordance with the mixed potential theory.2, but no PH3 was detected by OLEMS  during an applied potential sweep means that CoPx dissolution most likely does not proceed via a Coδ+ reduction route as observed for the electrochemical dissolution of other metal phosphides (e. g. InP+3H++3e−→In0+PH3↑).et al. who proposed for NiP2 that an anodic dissolution mechanism should be more likely than a cathodic degradation (NiP2+2H++3O2+2 H2O→Ni2++2H3PO4).x case. Adding to the preceding compositional analysis, an SEM and AFM analysis of the CA‐treated electrode reveals moderate pitting of the surface . On the other hand, the rLSV‐treated sample was subjected to a repeated sequence of potential sweeps in which the potential was gradually changed from OCP (E=+0.24 V) to HER potential (E=−0.26 V) with a sweep rate of 20 mV/s (25 s per scan). Dissolution during the rLSV treatment can be partially explained by minor anodic currents occurring at higher potentials . However the major part of Co (105 nmol/cm2) was dissolved in cathodic region of LSV and with that is still higher than that of CA treatment (13 nmol/cm2). This might be explained by the fact that such a frequent change of load  creates additional stress to the electrode by forming reverse currents upon repeated sweeping. A similar phenomenon is widely observed in commercial electrolyzers and fuel cells which suffer from reverse (or shunt) currents created by the shut‐down/start‐up of electrolyzer.2 produced during the preceding HER and accompanied by anodic degradation of the electrode. Indeed, in case of a constant cathodic treatment, no additional electrode corrosion was observed since the electrode was subjected to a constant negative current.Additionally, in order to evaluate the stability of CoPx is its faradaic efficiency. We conducted OLEMS measurements in order to evaluate the change in CoPx faradaic efficiency upon constant and interrupted electrolysis, similar to the performed CA and rLSV experiments. Faraday plots represent Faraday's law of electrolysis and can be used to detect possible electrode (de)activation processes related to surface changes. In the ideal case of a stable electrode, Faraday plots of the same electrode surface must have the same Faraday slope independent of the applied treatment (i. e. faradaic and gas collection efficiencies are constant). However, in cases where the surface of the electrode overcomes transformations or other side‐reactions begin to contribute to the measured faradaic charge, faradaic efficiency of the main reaction will also change resulting in a deviation of the Faraday slope.Another important parameter which characterizes HER performance of CoPx was subjected to a series of consecutive potential pulses  repeated 4 times. Short repeated pulses were preferred over continuously applied potential since OLEMS cannot assure an adequate linear detection of large quantities of evolved H2. The plot shows a reproducible dependence between charge and evolved H2, which can be interpreted as constant faradaic efficiency for HER  starts to be exposed to the electrolyte as well.In the first experiment, conforming CA, a fresh CoPR Figure a. Despitx was first subjected to a series of consecutive potential pulses  repeated twice to validate the previously observed constant FE. Then, a series of cyclic voltammograms  was applied to initiate changes similar to what was observed in the rLSV experiment and subsequently, another Faraday plot was recorded  CoPx phase at the surface. Formation of the P‐rich phase in the top‐most 23 % of the film was observed by XPS depth profiling and was attributed to a diffusion limitation of P during the thermal phosphidation resulting in a P gradient. Formation of a bulk CoP phase was also confirmed by Raman and validated by additional DFT calculations and XRD experiments. Morphology analysis confirmed roughening and agglomeration of the initial Co3O4 film upon thermal phosphidation. The exchange current density of the electrode was found to be comparable with reported values (j0=8.9 ⋅ 10−5 A/cm2) and the apparent activation energy for HER (Eapp=81±15 kJ/mol) of this system was measured for the first time. The large value of the Tafel slope (∼100 mV/dec) together with the fact that P is the HER‐active H+‐adsorbing site suggested that the reaction is limited by the number of P sites.In this work, we synthesized model CoP2SO4 with a constant dissolution rate (Co) of 2.4 nmol h−1 cm−2 even when no additional electrochemical treatment was applied. The constant cathodic treatment (CA) was found to have no influence on the dissolution rate, meaning that dissolution does not involve cathodic corrosion of CoPx. Moreover, OLEMS could not confirm the formation of PH3 upon cathodic potential cycling. The faradaic efficiency towards HER was found to be constant within ±10 % when a protective cathodic current was applied in between the temperature steps. However, an interruption under resting conditions (OCP) resulted in a significant decrease of the faradaic efficiency. These findings confirm that an anodic corrosion of CoPx is prevailing in accordance with literature. Dynamic potential changes performed in a repeated way (repeated cathodic linear sweep voltammetry) were found to increase dramatically the dissolution rate of the film leading to its complete degradation in 5 hours. This observation was explained by the fact that frequent change between HER and open circuit potentials increases the electrochemical stress applied to the electrode yielding a flow of transient reverse (i. e. anodic) currents which facilitate oxidation of CoPx. A similar issue is commonly detected in commercial electrolyzers upon electrolysis interruption, but was not yet observed for TMP systems.A detailed stability study revealed that the material is unstable in 0.5 M HIn this study, we show the importance of employing a multi‐technique characterization approach for an unambiguous evaluation of electrode stability. We stress out and showcase that current density constancy is not necessarily indicative for electrode stability and changes of faradaic efficiency cannot be accurately reflected in chronoamperometric measurements solely. Moreover, we show that the assessment of the chemical stability of electrodes under standby conditions is as important as stability under applied potentials.3O4 films were grown on commercial Au(100 nm)/Ti/Si(100) (Sigma Aldrich) substrates by plasma‐enhanced atomic layer deposition (PE‐ALD). CoCp2  was used as Co precursor and an O2 plasma  was used as oxidant. The deposition was carried out at 100 °C. The growth rate of Co3O4 was set to 0.05 nm/cycle. A single deposition cycle was composed of the following steps: precursor dosing (4 s), purging (6 s), plasma treatment (8 s), and second purging (6 s). No additional post‐treatment was applied to as‐deposited Co3O4 films. More details on the technical parameters of the utilized PE‐ALD setup and Co3O4 deposition process are given elsewhere.20 nm thick Co3O4 films was carried out in a horizontal calcination tube furnace. The substrates with an average area of 1 cm2 were positioned in the middle of the tube together with a quartz boat containing 0.5 g of sodium hypophosphite monohydrate NaH2PO2⋅H2O  used as P‐precursor. The boat was located upstream relative to the gas flow consisting of 4 : 1 Ar/H2, 150 mL/min STP. The reactor was heated to 260 °C (20 °C/min) to dehydrate the sodium hypophosphite followed by a further temperature increase to the target T (350–450 °C) with a slower heating rate (2 °C/min). The reactor was then kept at the target T for 30 min. Upon thermal decomposition, sodium hypophosphite releases PH3 which further reacts with Co3O4 to form cobalt phosphide. Unreacted phosphine was collected at the outlet by a liquid trap containing a concentrated aqueous CuSO4 solution. The sample was let to cool to room temperature in the same gas composition. Thermally treated substrates were rapidly transferred through air and stored in a glovebox. Such transfer results in the formation of a thin native passivation layer composed of PO43− and P0 species.Phosphidation of PE‐ALD fabricated Cox/Au/Ti/Si substrates with an average size of 0.1 cm2 were fixed on FTO glass slides (1.5×0.5 cm2) by conductive Ag paint (RS components) with additional top‐contact. To avoid the electric response of substrate and Ag, all conductive surfaces except CoPx were covered with a layer of insulating epoxy . The working electrode (WE) slides were mounted into a PEEK holder and contacted from the top by a gold wire enclosed in a quartz tube. The holder was additionally covered by several layers of Teflon tape (Swagelok) to prevent electrolyte interaction with the contact wire. The OLEMS tip was made of a PEEK tube with a porous hydrophobic Teflon plug inside and fixed in another quartz tube. The OLEMS tip was approached 20 μm above the surface of the working electrode by the micrometer screw with help of a video microscope . A Pt plate (10 cm2) and a Red Rod electrode  were utilized as counter (CE) and reference (RE) electrodes, respectively, and connected to an Ivium Compactstat potentiostat (Ivium Technologies). Ar‐purged 0.5 M H2SO4  was used as electrolyte. Before the experiments, the EC cell was cleaned with chromic acid (85 mmol K2Cr2O7+9.2 mol H2SO4). Milli‐Q water (18.2 MΩ ⋅ cm) was used in all cleaning and dilution steps. The temperature of the cell (25–35 °C) was maintained by a thermostat.In our experiments, we used a standard water jacketed electrochemical (EC) glass cell equipped with an OLEMS tip. CoPiR‐corrected reversible hydrogen electrode (RHE) scale, if not mentioned otherwise and current densities are normalized on the geometric surface area of the electrodes. More detailed information on the cell construction and OLEMS operation principles are given elsewhere.All potentials reported in this work are plotted on x electrode was continuously polarized at E=−0.12 V for 5 h and the electric current was recorded . In the second experiment, a series of linear sweep voltammograms was repeatedly applied to a fresh CoPx electrode (repeated LSV=rLSV). The potential range of voltammograms was selected starting from +0.24 V  till −0.26 V with a scan rate of 20 mV/s. The total duration of the rLSV experiment was also set to 5 h. In the third (reference) experiment, a fresh CoPx substrate was immersed into 0.5 M H2SO4 for 5 h without additional EC treatment. The (electro)chemical stability of the electrodes was then evaluated by ICP‐OES, XPS and Raman spectroscopy.HER conditions in acidic electrolyte, a series of experiments was performed. In the first experiment, a fresh CoPT‐dependent kinetic measurements were performed on a fresh CoPx electrode. The electrode was subjected to a series of cathodic potential pulses ranging from the resting potential (−0.20 V vs. Red Rod) till E=−0.40 V vs. Red Rod in 20 mV steps . Each pulse was applied for 10 s and the time between separate pulses was kept at 60 s to remove excess of produced H2. During a temperature change, a protective potential (−0.36 V vs. Red Rod) was applied for 40 min. The same potential sequence was repeated for each temperature . The temperature of the electrolyte was measured by a thermometer positioned near the WE. The temperature range was limited to 35 °C to reduce possible electrode degradation by temperature increase. Potentials were later normalized on the thermal drift of the RE and are reported on the RHE scale. Along with electrochemical treatment, OLEMS ion currents of the following m/z values were continuously recorded: 2 (H2+), 28 (N2+), 31 (P+), 32 (O2+ and PH+), 33 (PH2+) and 34 (PH3+). For comparison, the same experiment was repeated without applying a protective potential.In order to evaluate the change of faradaic efficiency (FE) under facilitated dissolution conditions (rLSV), a series of cyclic voltammograms was applied prior to the kinetic measurements. Constant FE of yet untreated sample was first validated by recording two consecutive Faraday curves which showed the same values of the slope. Voltammetric treatment led to a decrease of Faraday slope.Inductively Coupled Plasma Optical Emission Spectrometry (ICP‐OES) measurements were performed on a SPECTROBLUE EOP spectrometer equipped with an axial plasma source (Ar). Sample uptake rate was set to 2 ml/min. The emission intensity of Co ions was measured at 228.6 and 238.9 nm. The content of P was found to be below the detection limit of the spectrometer.Kα=1486.6 eV). The background pressure inside the analysis chamber was kept below 8 ⋅ 10−8 mbar reaching maximum of 3 ⋅ 10−7 mbar during the measurements due to the flow of low energy Ar+ ions involved in the charge neutralization process. High‐resolution spectra of core levels  and wide‐range survey spectra were recorded with a pass energy of 50 eV and 200 eV, respectively. Due to significant overlap between Au 4f peaks and the Co 3p loss feature, the Au 4d5/2 level was used for quantification instead. XP spectra were taken at the different positions of the surface to average the obtained quantitative results. XPS depth profiles were recorded on the same spectrometer. Ar+ sputtering (p=1 ⋅ 10−7 mbar) was done at 1 kV with time steps of 10 s (for the surface region) and 30 s (for bulk). XP spectra were taken after each sputtering step. XP spectra of spent electrodes were taken shortly after performed treatments. Electrodes were dismounted from the EC cell, rinsed with copious amounts of Milli‐Q water, dried with compressed air and immediately introduced to the vacuum chamber of an XP spectrometer. Total duration of air exposure was estimated to be approximately 15 min (before and after each experiment) .X‐ray photoelectron spectroscopy (XPS) measurements were carried out on a K‐Alpha XP spectrometer (Thermo Scientific), equipped with a monochromatic small‐spot (400 μm) X‐ray source, a 180° double focusing hemispherical analyzer with a 128‐channel detector and an Al anode (Al Raman spectra were recorded on a Horiba Raman microscope equipped with an Olympus LMPlanFl 100x 0.80 lens and HeNe laser (632.8 nm) with a diffraction grating of 600 gr/mm. Each spectrum was collected for 20 s (2×10 s).Tapping mode atomic force microscopy (AFM) measurements were carried out on NT‐MDT Next microscope. Both height and phase images were recorded simultaneously with gold‐coated Si probes with a curvature radius of 10 nm . Scanning electron micrographs (SEM) were recorded on a FEI Quanta 3D FEG microscope at an accelerating voltage of 5 kV without additional coating of the surface.The authors declare no conflict of interest.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


The Mg insertion into spinel lattices occurs via “intercalation and push‐out” process to form a rocksalt phase in the spinel mother phase. For example, by utilizing the valence change from Co(III) to Co(II) in MgCo2O4, Mg insertion occurs at a considerably high potential of about 2.9 V vs. Mg2+/Mg, and similarly it occurs around 2.3 V vs. Mg2+/Mg with the valence change from Mn(III) to Mn(II) in MgMn2O4, being comparable to the ab initio calculation. The feasibility of Mg insertion would depend on the phase stability of the counterpart rocksalt XO of MgO in Mg2X2O4 or MgX3O4 . In addition, the normal spinel MgMn2O4 and MgCr2O4 can be demagnesiated to some extent owing to the robust host structure of Mg1−xX2O4, where the Mg extraction/insertion potentials for MgMn2O4 and MgCr2O4 are both about 3.4 V vs. Mg2+/Mg. Especially, the former “intercalation and push‐out” process would provide a safe and stable design of cathode materials for polyvalent cations.On the basis of the similarity between spinel and rocksalt structures, it is shown that some spinel oxides (e.g., MgCo Currently, lithium ion batteries (LIBs) are widely used for a lot of practical applications, and their energy density has been enlarged year by year, but its growing rate tends to be saturated recently. If lithium metal itself could be used as an anode material instead of carbonaceous materials currently used, LIBs would have shown significantly high energy densities, but this cannot be done due to the well‐known fatal problem, “dendritic growth” of Li metal on charge that leads to dangerous short circuits.−1) largely exceeds that for the current carbonaceous anode materials (ca. 370 mAh g−1). Especially, it has been reported that Mg electrodeposition occurs with non‐dendritic formation;−1 (currently 370 mWh g−1 in electrode energy density of LiCoO2 vs graphite). Therefore, unless more talented cathode materials that can accommodate polyvalent cations are sought out, PSBs would not be comparable to Li ion batteries in terms of the energy density. Thus, in order to change the energy storage paradigm, we have to seek cathode materials for polyvalent cations.As an alternative to Li metal‐anode battery, polyvalent‐metal  storage batteries (PSBs) have recently attracted increased attention owing to their large capacities; for example, in the case of Mg, its capacity , while those in the rocksalt structure are denoted as 16c, 16d (cation sites) and 32e (oxygen sites) when it is assigned to the same space group. Thus, a spinel structure can be regarded as a rocksalt whose 16c sites are vacant and instead the 8a sites are usually occupied by cations. Therefore, it is expected that Mg cations can be inserted onto 16c vacant sites in the spinel lattice, as well as the Li insertion mechanism in spinel oxide materials.Here we focus Mg spinel oxides as candidates for cathode materials of MRBs. As shown in 2O4, MgMn2O4, MgFe2O4, MgCr2O4, and Co3O4, we demonstrate that some of spinel oxides can allow the insertion of Mg cations at high potentials (about 3 V vs. Mg2+/Mg) via “intercalation and push‐out” mechanism as shown in Figure In this work, with several spinel oxides MgCo22.1Figure2/CsTFSA binary ionic liquid, but the electrolytic dissociation of Mg(TFSA)2 in the ionic liquid takes place, which is judged from the fact that the Mg cations can be inserted into the Chevrel compounds.2+/Mg was tentatively determined to be “0.5 V vs. Li+/Li in RE ≈ 0 V vs. Mg2+/Mg”, as shown in Figure +/Li in RE; see Figure S1, Supporting Information, for details.A typical construction of beaker cells used here is illustrated in 2O4. Usually, one would consider the conventional reaction, i.e., the reaction of Mg extraction from the host material, MgCo2O4 ⇔ Mg1−xCo2O4 +x(Mg2+ + 2e−). However, such a reaction may occur above 4.4 V vs. Li+/Li in RE, but this cation extraction from MgCo2O4 appears to be difficult in terms of the structural stability and anodic limit of the electrochemical window of the binary ionic liquid. Apart from this usual deintercalation, it is clearly seen that direct Mg insertion can occur into the host MgCo2O4 without a pre‐charge process, and then the cation extraction is observed during a charge process; the equilibrium redox potential is about 2.9 V vs. Mg2+/Mg (3.4 V vs. Li+/Li in RE), which is in agreement with the ab initio calculation (3.0 V vs. Mg2+/Mg); see the later section on the ab initio calculation. Thus, the insertion/extraction of Mg cations are found to be drastically facilitated by elevating temperature. After the electrochemical tests, we confirmed Mg insertion semi‐quantitatively by the energy‐dispersive X‐ray (EDX) spectroscopy analysis (not presented here), and also we have performed the inductively coupled plasma (ICP) analysis after the electrochemical performance test for the Mg‐Li rocking‐chair type dual‐salt battery.Figure 3O4, MgMn2O4, MgCr2O4, and MgFe2O4, were also investigated. Figures 3O4 is very similar to that of MgCo2O4, where the extraction of Co(II) cations is not observed during the anodic scan from the open circuit potential (OCP). The elemental analysis by EDX suggested that Mg cations were inserted into the spinel Co3O4 (note presented here). In contrast to MgCo2O4, the extraction of Mg2+ ions from MgMn2O4 is observed in Figure 2+ ions are observed at around 3.4 V and 2.3 V vs. Mg2+/Mg; the former reaction would correspond to the valence change of Mn(IV) to Mn(III), whereas the latter would correspond to that of Mn(III) to Mn(II). According to the present ab initio calculations of the energy differences between MgMn2O4 and demagnesiated‐spinel Mn2O4 and between MgMn2O4 and magnesiated‐spinel  Mg2Mn2O4, the average redox potentials were estimated to be 2.9 V vs. Mg2+/Mg for the former and 1.8 V vs. Mg2+/Mg for the latter, being in fairly agreement with the experimental results.As candidates of cathode active materials for MRB systems, other spinel oxides, Co2O4, that is, the valance change from Cr(IV) to Cr(III) was observed, but in contrast, the valence change of Cr(III) to Cr(II) was hardly observed in this case. In the case of MgFe2O4, as well as the above materials, we confirmed that MgFe2O4 can be used as a cathode material for MRBs, but marked redox peaks were not observed in Figure 2O4; the insertion/extraction potential is about 2.7 V vs. Li+/Li in RE, being lower than the redox potential (about 3.4 V vs. Li+/Li) of Fe cations in the olivine LiFePO4. This is further supported by the XRD and XANES measurements in the next section.As to the former reaction, the similar trend was observed in Figure 2O4 at relatively lower temperatures, by using the Mg(TFSA)2/triglyme electrolyte, which was recently developed for MRBs.2O4 in a triglyme electrolyte containing 0.5 m Mg(TFSA)2 at room temperature (RT) and 100 °C. The insertion of Mg2+ ions into the spinel MgCo2O4 was observed in each case (60 mAh g−1 for RT and 105 mAh g−1 for 100 °C) below 3.4 V vs. Li+/Li in RE. After the cathodic sweep to 1.5 V vs. Li+/Li in RE, the anodic current corresponding to the extraction of Mg2+ ions was markedly observed above 3.4 V vs. Li+/Li in RE, but it is found that the extraction of Mg2+ below 100 °C is considerably laborious within the electrochemical window of the triglyme. Thus, the triglyme electrolyte can be used for the evaluation of active materials at RT and/or higher temperatures.In the above cyclic voltammetry tests, the temperature was set at 150 °C. Here, we show the CV profiles of MgCo2.22O4 before/after Mg insertion/extraction tests to comprehend the cation‐insertion mechanism. The analyses have been done for MgCo2O4, Co3O4, MgMn2O4, and MgFe2O4. As shown in XRD profiles in Figure−1, the active material contains two phases, i.e., spinel and rocksalt phases. Further insertion of Mg cations up to about 210 mAh g−1 substantially forms a rocksalt single phase. The fact that the spinel phase disappears even at such an incomplete discharge amount less than the theoretical value (260 mAh g−1) means that the rocksalt phase includes a certain amount of cation vacancies, that is, a solid‐solution phase of off‐stoichiometry exists. When charging after the discharge of 120 mAh g−1, the structure completely reverts to the spinel structure. By measuring the corresponding XANES spectra around the Co K‐edge in Figure −1 and again oxidized to Co(III) after charge; compare to the XANES profiles of Co3O4 containing Co(II) and Co(III), CoO with only Co(II), and LiCoO2 with only Co(III).We conducted structural analyses for mainly MgCo2O4, we have also conducted XRD and XANES measurements for Co3O4 in Figure 2O4 in Figure 2O4 in Figure 3O4 in Figure 2O4. Incidentally, a certain amount of Si was mixed to gain a sufficient sample volume in the capillary for the measurements. Namely, by the insertion of Mg cations (200 mAh g−1), the rocksalt phase is formed in the spinel mother phase. Comparing to MgCo2O4, the spinel phase tends to remain in the case of such an incomplete discharge amount. On the other hand, we observe not only the Mg insertion into a pristine MgMn2O4 but also the Mg extraction from such a pristine MgMn2O4, as seen in Figure 2O4 shifts to a lower energy after discharge and shifts to a higher energy after charge. These behaviors are consistent with those of the reference samples corresponding to the various valence states of Mn. Finally, although MgFe2O4 did not show an excellent cathode property in Figure 2O4, FeO and MgO are very close, so that we cannot clearly see the peak shifts in the broad XRD profile of pristine MgFe2O4. However, after the Mg insertion, the XRD peaks tend to move to lower angles (probably due to the influence of FeO) and the white line (Fe K‐edge) tends to shift lower. As seen in the XANES profiles of Fe2O3 and Fe3O4, there is only a little change in the respective white lines, and the present XANES profile change after Mg insertion is rather similar to this trend, unlike the Fe K‐edge profile in the olivine LiFePO4 structure.As well as MgCo2O4 sample and one after partial discharge (120 mAh g−1) in the Mg battery system that corresponds to Figure −1) was not taken into account for the Rietveld analysis. The fitting results and detailed parameters for crystal structures are shown in Figure S3 and Table S1, Supporting Information. From the Rietveld analysis, as‐synthesized MgCo2O4 takes a disordered spinel structure with a degree of disorder of about 0.43, being consistent with our previous works.−1/260 mAh g−1), we need to consider the presence of vacancies in the rocksalt crystal. Thus, the insertion of one Mg atom induces the spinel to rocksalt transition in a larger region than one unit cell of the rocksalt structure.As a representative of spinel oxide cathode materials for MRBs, the crystal structure parameters were determined by Rietveld refinement with the program RIETAN‐FP2O4 spinel lattice is expressed as2Co2O4 takes a rocksalt structure, and the rocksalt structure is formed via “intercalation and push‐out” process in Figure From the present structure analysis, the Mg cations are inserted into 16c sites in the spinel structure, and the original cations located at the 8a sites of the spinel structure are pushed out to the 16c sites, eventually to form a rocksalt structure. The Mg insertion mechanism, “intercalation and push‐out” process, is close to the Li‐insertion mechanism in spinel oxide materials.2.3+/Li in RE,2O4 and MgMn2O4 for the cathode performance tests. Figure2O4, and 1/20 C and 1/50 C for b) MgCo2O4 and for c) MgMn2O4.We have successfully obtained various evidences of the Mg insertion/extraction into/from the spinel oxides. Then, let us demonstrate the high potentials of these cathode materials by constant‐current battery performance tests. In this Mg battery system using the (Mg10/Cs90)‐TFSA ionic liquid, the Mg anode is readily passivated, so that its anodic dissolution comes to occur above 1.5 V vs. Li2O4, the battery test can also start from the discharge process. As shown in Figure 2+/Mg) of the electrochemical window of the electrolyte. Consequently, after the charge process, the discharge amount became smaller than the first one, but this is not the essential problem of the cathode material. Moreover, the cyclability is significantly affected by the thermal stability of the PVDF binder  that cannot be endurable around 150 °C. To return to the subject, in the 1st discharge process, the converted working potential is shown to be 2–2.5 V vs. Mg2+/Mg, which is much higher than that of the Chevrel compounds,−1 above 2 V vs. Mg2+/Mg and also amounts to 170 mAh g−1 above 1 V vs. Mg2+/Mg at a rate of 1/10 C.As to the cathode material MgCo2O4 to MgMn2O4 at slower rates, 1/20 C and 1/50 C. As to MgMn2O4, the battery test was started from a charge process (over 200 mAh g−1 but including the oxidation decomposition of the electrolyte). In the case of 1/20 C, the MgCo2O4 cathode material shows a longer plateau region around 2.0–2.5 V vs. Mg2+/Mg in the potential versus capacity curve, and the discharge amount reaches about 200 mAh g−1. Even for a slower rate of 1/50 C, the capacity was substantially unchanged, indicating that a certain repulsive interaction associated with the push‐out process is influenced on the Mg insertion process when the cation density is increased, which is a certain kind of freezing phenomenon like glass/jamming transition.At a slower C rate, we can obtain more excellent discharge behavior. Figure 2O4 cathode material displays two stages  in the discharge process, which are in accordance with the redox potentials observed in the CV profile for MgMn2O4 in Figure −1) can be attained for a slower rate in MgMn2O4  of this cathode material would be a fascinating characteristics.In contrast, the MgMn2.4U in order to supplement the experimental findings. Although MgCo2O4 and MgFe2O4 are disordered spinels, we considered only normal spinel configurations as it is found in our recent study2O4 does not significantly affect the redox potential caused by the Mg insertion; see Supporting Information for details. We calculated the redox potentials of MgX2O4 by Mg insertion asE(A) denotes the total energy of phase A, X = Cr, Mn, Fe, or Co, and e is the elementary charge. The redox potential of Co3O4 by Mg insertion was also calculated asMg2X2O4 and MgCo3O4 models were created by displacing the Mg or Co cations located at the 8a sites to neighboring 16c sites and inserting Mg cations to the remaining 16c sites. On the other hand, the redox potential of Mg extraction was calculated asX2O4 phases were made by simply removing the 8a‐site Mg cations. The calculated potential vs. Mg2+/Mg for the insertion/extraction of Mg cations into/from the host spinel structure are summarized in TableAb initio calculations were performed using GGA+2.52Co2O4 and MgCo3O4 of a random‐solution type rocksalt phase can be formed for MgCo2O4 and Co3O4, since both of MgO and CoO have similar rocksalt structures and both phases are thermally stable in the ambient condition. Similarly, since the rocksalt MnO is stable, Mg cations can easily be inserted into the MgMn2O4 host material. However, since the lattice mismatch of MnO and MgO is fairly large  is the d3 ion and t2g orbitals in the octahedral crystal field are fully filled for the majority spin component, the valence change from Cr(III) to Cr(II) is quite unfavorable energetically. In such a case, the Mg insertion into the spinel host would be laborious.Here we discuss the feasibility of Mg insertion and cation extraction into/from the spinel oxides. Table 2O4, but the potentials for MgFe2O4, MgMn2O4, and MgCr2O4 are within 3–4.5 V vs. Mg2+/Mg. Thus, the latter spinel oxides are expected to be demagnesiated relatively easily, and actually, among these oxides the Mg extraction tends to occur in MgMn2O4 and MgCr2O4. This is because these host materials take normal spinel structures so that the demagnesiated Mg1−xMn2O4 or Mg1−xCr2O4 structures would be rather robust. Since the valence change from 4 to 3 of transition metals can be utilized in this case, we can obtain a relatively higher redox potential than that by utilizing the valence change from 3 to 2. If the stability of the demagnesiated structure is ensured, one had better utilize the former valence (4 to 3) change. Furthermore, the Mg insertion with the spinel to rocksalt transition is a fascinating characteristic redox reaction. Thus, MgMn2O4 can be a high potential cathode material for MRBs in that the both redox reactions  can be utilized.On the other hand, the feasibility of Mg extraction from the original spinel oxides can basically be judged from the magnitude of the extraction potential. According to the ab initio calculation, the extraction potential is relatively high for MgCo32O4, MgMn2O4, MgFe2O4, MgCr2O4, and Co3O4, toward Mg rechargeable batteries (MRBs). In some of the spinel oxides , the Mg insertion and extraction can be clearly observed, which is facilitated by elevating temperature (about 150 °C). From various viewpoints, for example, electrochemistry, structural analysis, and ab initio calculation, we have substantiated the eccentric mechanism on Mg insertion into spinel‐oxide lattices, termed “intercalation & push‐out” process. This Mg insertion into a spinel structure occurs with an atomic‐level coherent phase transition, where dual‐phase reaction of the spinel and rocksalt phases proceeds. For example, in the case of spinel MgCo2O4, Mg insertion occurs at a significantly high potential of about 2.9 V vs. Mg2+/Mg, being consistent with ab initio calculation, and its capacity approximately amounts to 200 mAh g−1 . The feasibility of Mg insertion into spinel oxides would depend on the stability of both MgO and XO rocksalt phases in MgX2O4 or X3O4. In contrast, the normal spinel oxides, MgMn2O4 and MgCr2O4, can be demagnesiated to some extent, where the Mg insertion/extraction potentials of MgMn2O4 and MgCr2O4 are both about 3.4 V vs. Mg2+/Mg.In conclusion, based on the structural similarity in spinel and rocksalt structures, we have investigated cathode properties of spinel oxides, MgCo2+/Mg (for MgCo2O4), the electrode energy density experimentally amounts to about 400 mWh g−1 . Thus, we are sure that the “intercalation and push‐out” mechanism provides a new strategy for designing future cathode materials for polyvalent cations such as Mg cations. In addition, the spinel oxide cathode materials  can allow not only Mg‐cation insertion but also Li‐cation insertion, which enables us to design a new type of rechargeable battery, “rocking‐chair type Mg‐Li dual‐salt battery”, which is discussed in our another paper.Since the valence change from 4 to 3 can be utilized in the latter redox reaction, a relatively higher potential is delivered, while a relatively lower potential due to the valence change from 3 to 2 is utilized in the former redox reaction. Nevertheless, even though the former redox reaction is used, since the Mg‐insertion potentials are as high as about 2–3 V vs. Mg4Sample Preparation: All spinel oxides were synthesized by the inverse co‐precipitation method.m Mg(II), 0.160 m X(II), X = Co, Mn, Cr, Fe) were prepared by dissolving Mg(NO3)2·6H2O and X(NO3)2·nH2O, etc, in deionized water. A sodium carbonate solution  for pH control and precipitation was also prepared. These solutions were heated to 70–80 °C under vigorous stirring (500 rpm). The metallic nitrate salt solutions were added dropwise into the sodium carbonate precipitation solution. The resulting suspensions were stirred at 70–80 °C for 30 min and then filtered. The filtered precipitates (precursors) were rinsed with deionized water (300 cm3) at 80 °C to remove completely Na‐containing by‐products, and air‐dried for 24 h at 80 °C. The precursors were followed by calcination in air at 350–750 °C for 2–24 h.Electrochemical Tests: Each composite cathode was prepared by coating an Al plate collector with a mixture of the active material, carbon black (as conductive agents), and PVDF (binder) in a weight percent of 80:10:10. Mainly we used CsTFSA‐based ionic liquids containing Mg(TFSA)2 (and/or LiTFSA) salt established by Hagiwara et al.,3SO2)2−, and sometimes used an 0.5 m Mg(TFSA)2/triglyme electrolyte established recently.N,N‐diethyl‐N‐methyl‐N‐(2‐methoxyethyl)ammonium bis(trifluoromethanesulfonyl)amide (DEMETFSA) for CsTFSA‐based mixed ionic liquid and the same triglyme solvent for the 0.5 m Mg(TFSA)2/triglyme electrolyte, and the solution for the reference electrode was LiTFSA in both cases. As shown in Figure 2+/Mg couple in the (Mg10/Cs90)‐TFSA ionic liquid was estimated to be about 0.5 V vs. Li+/Li in RE, where RE means the Li reference electrode in a glass tube separated by a ceramic filter. Moreover, the redox potential of Mg2+/Mg couple in the 0.5 m Mg(TFSA)2 in triglyme electrolyte was estimated to be about 0.8 V vs. Li+/Li in RE (not presented here). All the electrochemical tests and beaker‐cell construction were done with galvanostatic/potentiostatic apparatuses  in the glove box whose dew point was below −72 °C.Structural Analysis: The structure and valence state of the active material were investigated by X‐ray diffraction (XRD) and X‐ray absorption near edge structure (XANES), respectively, at a synchrotron radiation facility, SPring‐8. XRD patterns were acquired by the Debye‐Scherrer method with a Lindeman glass capillary and a rotating stage, where the wavelength of λ =1.672 Å and 0.5 Å were used. The discharged samples for the X‐ray measurements were rinsed by tetrahydrofuran (THF), dried in the glove box, followed by encapsulating into a capillary in the air atmosphere, since we confirmed that the open circuit potentials of these samples were within the electrochemical window of water. On the other hand, the charged samples were wholly prepared in the glove box by using triglyme and dimethyl carbonate instead of THF for washing to circumvent the self‐discharge of the electrode samples. The obtained capillaries were sealed by a vacuum grease and kept in an Ar atmosphere. Si powder was added to some of the samples so as to assist crushing the composite electrode before encapsulating to the capillary. The samples for the XANES measurement were prepared in the same procedure by using the same electrode samples for XRD measurements; the electrode‐shape sample with an Al collector was used for the XANES measurement. As for the charged sample, the electrode was packed by the polyimide tape to circumvent the exposure to the air atmosphere, while the bare electrode was used for the XANES measurements of discharged samples. The Rietveld analysis of XRD was performed with RIETAN‐FP.Ab Initio Calculation Procedure: To complement the experimental findings, ab initio calculations were performed using the GGA+U approach; the detailed procedure was described in our recent paper.vasp.U scheme proposed by Lichtenstein et al.U values on X‐3d  orbitals in various oxides using a selfconsistent scheme: 4.64–5.09, 4.90, 4.91–6.34 eV for Mn(III), Fe(III), and Co(III), respectively, which can well reproduce the voltage of Li ion batteries within a few tenth eV.U = 5 eV for Mn‐ and Fe‐3d orbitals and U = 6 eV for Co‐3d orbitals. For Cr‐3d, we used a typical value of U = 5 eV. J was set at a value of 0.88 eV for all cases. Wave functions were expanded using a plane‐wave basis set with a cutoff energy of 550 eV. Spin polarization was considered for all cases. The lattice constants and internal atomic positions were fully optimized in all calculations until the residual stresses and forces converged to less than 0.24 GPa and 0.02 eV/Å, respectively.The authors would like to thank Ms. Chiharu Hirao for her experimental help. This work was supported by the Advanced Low Carbon Technology Research and Development Program (ALCA), Grant‐in‐Aid from the Special Coordination Funds for Promoting Science and Technology commissioned by JST, MEXT of Japan, and the MEXT Elements Strategy Initiative to Form Core Research Center. Computing resources of ACCMS at Kyoto University were used in this work. This achievement is based on the significant work on the ionic liquids by Professor Hagiwara et al. (Kyoto University) at the Research and Development Initiative for Scientific Innovation of New Generation Batteries (RISING) project from New Energy and Industrial Technology Development Organization (NEDO) of Japan.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.