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Tuning Surface Chemistry of TiC Electrodes for Lithium−Air Batteries Anna Ya. Kozmenkova,†,■ Elmar Yu. Kataev,†,■ Alina I. Belova,† Matteo Amati,‡ Luca Gregoratti,‡ Juan Velasco-Vélez,§ Axel Knop-Gericke,§ Boris Senkovsky,∥ Denis V. Vyalikh,⊥,# Daniil M. Itkis,*,† Yang Shao-Horn,∇ and Lada V. Yashina† †

Lomonosov Moscow State University, Leninskie gory, 119991 Moscow, Russia Elettra - Sincrotrone Trieste S.C.p.A., Area Science Park, I-34012 Basovizza, Trieste, Italy § Fritz-Haber-Institut der Max-Planck-Gesellschaft, Faradayweg 4−6, 1495 Berlin, Germany ∥ Universität zu Köln, Zülpicher Strasse 77, 50937 Köln, Germany ⊥ IKERBASQUE, Basque Foundation for Science, 48011 Bilbao, Spain # Donostia International Physics Center (DIPC), Departamento de Fisica de Materiales and CFM-MPC UPV/EHU, 20080 San Sebastian, Spain ∇ Materials Science and Engineering Department, Massachusetts Institute of Technology, Cambridge, Massachusetts 02139, United States ‡

S Supporting Information *

ABSTRACT: One of the key problems hindering practical implementation of lithium−air batteries is caused by carbon cathode chemical instability leading to low energy efficiency and short cycle life. Titanium carbide (TiC) nanopowders are considered as an alternative cathode material; however, they are intrinsically reactive toward oxygen, and its stability is controlled totally by a surface overlayers. Using photoemission spectroscopy, we show that lithium−air battery discharge product, lithium peroxide (Li2O2), easily oxidizes clean TiC surface. At the same time, TiC surface, which was treated by molecular oxygen under ambient conditions, shows much better stability in contact with Li2O2 that can be explained by the presence of a surface layer containing a significant amount of elemental carbon in addition to oxides and oxycarbides. Nevertheless, such protective coatings produced by room temperature oxidation are not practically useful as one of its components, elemental carbon, is oxidized in the presence of lithium−air battery discharge intermediates. These results are of critical importance in understanding of TiC surface chemistry and in design of stable lithium−air battery electrodes. We postulate that dense, uniform, carbon-free titanium dioxide surface layers of 2−3 nm thickness on TiC will be a promising solution, and thus further efforts should be taken for developing synthetic protocols enabling preparation of TiO2/TiC core−shell structures.

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INTRODUCTION Lithium−oxygen chemistry can potentially enable development of rechargeable batteries demonstrating few-fold increased specific energy in comparison to lithium-ion ones.1 The practical implementation of this idea, however, faces a number of significant challenges, including reactivity of the electrolyte components,2 polymeric binders,3 and positive electrode materials toward species forming inside the cell. During lithium−air battery (LAB) operation, oxygen reduction reaction (ORR) occurs at the positive electrode, and oxygen dissolved in the electrolyte is reduced to superoxide (O2−)4 that in the presence of Li+ forms ionic pairs (LiO2), the stability of which depends on electrolyte solvent properties,5−8 additives in electrolyte,8 and nature of anions.9,10 Superoxide species are then converted to a final discharge product lithium peroxide (Li2O2) after transfer of a second electron from the electrode or by chemical disproportionation reaction.4,11 The abovementioned chemical reactivity of positive electrode materials, © 2016 American Chemical Society

electrolyte, and other components toward reduced oxygen species results in a number of side reactions leading to undesired byproducts (e.g., Li2CO3, LiOH, etc.) passivating electrode surface and contaminating electrolyte. This study focuses on resolving the cathode chemical stability problem. Among others, carbon materials received major attention as a positive electrode in Li−O2 cells as they are freely available, porous, highly conductive, and lightweight. Unfortunately, carbon electrodes were found to be insufficiently stable in cell operation conditions.12−14 Although thermodynamically permitted,13 carbon reactivity with Li2O2 was found to be negligible.15 Superoxide intermediate species, nevertheless, promote carbon degradation,14,15 thus limiting sustainable cycling of the positive electrode. Received: August 2, 2016 Revised: October 10, 2016 Published: October 11, 2016 8248

DOI: 10.1021/acs.chemmater.6b03195 Chem. Mater. 2016, 28, 8248−8255

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Chemistry of Materials

electrolyte was performed inside an Ar-filled glovebag directly attached to a spectrometer load lock. The glovebag was evacuated and purged with high-purity Ar (Linde group, 99.9999%) five times prior to each cell assembly/disassembly. For bubbling of the electrolyte with highpurity oxygen (Logika, 99.9999%), the glovebag was equipped with a gas inlet/outlet with hydroseal disabling back diffusion of environmental air inside the cell or the glovebag. Cyclic voltammetry measurements were carried out using Bio-logic SAS SP-200 and SP-300 potentiostats at a sweep rate of 100 mV/s. Cells for in situ XPS studies were assembled similarly to those reported previously.14 The positive electrode was prepared by rubbing a small amount of TiC nanopowder (30−50 nm, 99%, Alfa Aesar) onto a solid glass-ceramic NASICON-type electrolyte (Li1.5Al0.5Ge1.5(PO4)3) piece with nickel strips that were preliminarily deposited by RF sputtering and served as electric contacts. A twoelectrode cell was assembled in an Ar-filled glovebox by stacking the cathode film on the solid electrolyte, a porous polypropylene separator soaked with 1 M lithium bis(trifluoromethane)sulfonylimide (LiTFSI, Sigma-Aldrich) in 1-ethyl-3-methylimidazolium bis(trifluoromethylsulfonyl)imide (EMI TFSI, ≥98%, Sigma-Aldrich) electrolyte, and lithium foil, which was employed as anode. The stack was placed between two stainless steel plates (one with a hole on cathode side). Galvanostatic discharge of the cell was carried out by applying a constant current of 1 μA for a certain periods of time. A SP200 potentiostat/galvanostat (BioLogic SAS Instruments) was used for that purpose. Oxygen pressure of about 0.1 mbar was kept up in the chamber during cell operation. After each discharge step, the photoelectron spectra were measured. TiC Single Crystal Preparation and Treatment. For studying TiC interaction with Li2O2, we used a polished single crystal cut along (755) plane (MaTecK). The surface was cleaned by heating at 1150 °C for 20 min at an oxygen pressure of 10−7 mbar followed by cooling to room temperature and subsequent fast heating (flashing) to 1300 °C several times. Oxidation of the crystal surface was carried out at room temperature (pO2 = 200 mbar), at 500 °C (pO2 = 10−7 mbar), and at 1100 °C (pO2 = 10−7 mbar) for 30 min. Graphene was grown by chemical vapor deposition on clean TiC (755) surface from propylene (C3H6) at 850 °C and C3H6 pressure of 10−6 mbar during about 5 min followed by slow cooling to room temperature at a rate of 1−2 °C/min. Li2O2 was deposited on clean, oxidized, and graphene covered TiC (755) surface. To obtain Li2O2, metallic lithium was deposited from a Li dispenser (SAES) under an oxygen pressure of about 10−7 mbar at room temperature. X-ray Photoemission Studies. X-ray photoemission studies were performed using several facilities at Helmholtz Zentrum, Berlin (HZB), and Elettra, Trieste. XPS studies of TiC (755) surface reactivity were performed at the RGBL beamline (HZB). High-resolution Ti 2p, C 1s, O 1s, and Li 1s core level spectra were acquired at a photoelectron kinetic energy of 50 eV to provide ultimate surface sensitivity and at 200 eV for composition quantification using a SPECS Phoibos 150 electron energy analyzer. Photon energies were calibrated using the secondorder diffraction reflection from the plane grating. In situ photoemission observation of the solid electrolyte electrochemical cell was implemented at both ISISS beamline (HZB) and ESCAmicroscopy beamline (Elettra). For the surface of the TiC positive electrode, Ti 2p, C 1s, O 1s, and Li 1s spectra were acquired at a photoelectron kinetic energy of 200 eV under an oxygen pressure of 0.1 mbar after consequent discharge steps using a SPECS Phoibos 150 NAP electron energy analyzer (ISISS) and at a photon energy of 647.9 eV under a dynamic oxygen pressure of ∼10−2 mbar provided by capillary22 using a 48-channel hemispherical analyzer (ESCAmicroscopy). In the latter case, composition mapping possibility was provided by focusing the X-ray beam to a 120 nm spot using a Fresnel zone plate. In some cases, the electrode surface was treated with atomic oxygen to eliminate adventitious carbon and sputtered with Ar+ ions of 500 eV for 10 min.

In the search for alternatives, researchers considered a number of options including noble metals,16 and also nonoxide conducting binary compounds of transition metals, primarily titanium.17,18 In contrast to the others, the latter option, which includes titanium carbide (TiC) and nitride (TiN), seems attractive due to commercial availability of these compounds in the form of nanopowders, high conductivity, and affordable price. Reported by Thotiyl et al., porous TiC cathode demonstrated high cycle stability (>98% capacity retention after 100 cycles) with Li2O2 as the main discharge product.17 Such an outstanding performance was attributed to a thin TiO2 layer growing during battery operation and protecting the surface from nucleophilic attacks of O2− and Li2O2 and further oxidation. As it was further demonstrated by Adams et al., overall recharge performance is quite sensitive to the surface layer chemistry defined by TiC powder prehistory.18 The optimal thickness of the TiO2 protective layer, which should be thick enough to inhibit oxidation and at the same time maintain reasonable electron transport though it, was estimated to be about 2 nm for both TiC and TiN. At the same time, both TiC and TiN are subjected to natural oxidation that results not only in the formation of oxide layers composed by titanium suboxide, dioxide, and oxycarbide/ oxynitride, but also, in the case of TiC, leads to the formation of elemental carbon precipitates on the surface.19−21 An in-depth understanding of surface chemistry is required to gain control over the composition of surface layer and fine-tune the surface chemistry of TiC and other binary titanium conductive compounds, thus opening the way to reliable design of noncarbon stable electrodes for oxygen redox in aprotic media. Here, we discover the reasons of the degradation of TiC, and some other Ti binary compound electrodes, based on comprehensive studies of the surface chemistry in model chemical and electrochemical systems probed by in situ and ex situ experiments. We show that the interactions of TiC with LiORR products and intermediates lead to progressive growth of passivation layers. We believe that this knowledge is helpful for designing TiC-based materials for LAB positive electrode.

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EXPERIMENTAL SECTION

Disc Electrode Preparation. Disc electrodes were fabricated by electrical discharge machining of TiC (99.5%), TiN (99.5%), and TiO (99.9%) sputtering targets (Cathay Advanced Materials, Ltd.). Subsequently, the obtained discs were grinded using rough and then soft abrasive paper. Polishing was performed using 6, 3, and 1 μm diamond suspensions (Bio Diamant Liquids, Heraeus Kulzer GmbH) until the surface attained a mirror shine. Cell Assembly and Electrochemical Measurements. For cyclic voltammetry studies, we used a hermetic three-electrode cells designed and fabricated in-house. A 0.1 M solution of lithium perchlorate (battery grade, 99.99%, Sigma-Aldrich) in dimethyl sulfoxide (DMSO, anhydrous, ≥99.9%, Sigma-Aldrich) was used as electrolyte. Water content in the electrolyte did not exceed 30 ppm as determined by Karl Fischer titration with a 831 KF coulometer (Metrohm). Platinum wire was used as a counter electrode. Ag+/Ag Vycor frit separated electrode with 0.01 M AgNO3 (Sigma-Aldrich) and 0.1 M tetrabutylammonium perchlorate (for electrochemical analysis, ≥99.0%, Sigma-Aldrich) in acetonitrile (gradient grade for liquid chromatography, ≥99.9%, Merck Millipore) electrolyte was employed as a reference. The reference electrode potential was calibrated vs ferrocene/ferrocenium couple (98%, Aldrich). The cells were assembled and sealed in an Ar-filled glovebox (Labconco Protector, CA). Assembly and cyclic voltammetry of the cells for ex situ analysis of the TiC electrode surface evolution upon cycling in DMSO-based 8249


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Figure 1. Ti 2p (a,c,e,g) and C 1s (b,d,f,h) photoemission spectra for clean (a,b) and oxidized (c−h) TiC (755) single crystal before (gray) and after (black) Li2O2 deposition measured at ultimate surface sensitivity provided by electron kinetic energy of 50 eV. A clean vicinal (755) TiC surface was obtained by annealing the single crystal in oxygen (10−7 mbar) at 1150 °C for 20 min and further “flashing” up to 1300 °C for several times. Such a procedure allowed one to eliminate oxygen from the surface; residual oxygen intensity never exceeded 5 at. % for the layer probed with a photoelectron kinetic energy of 200 eV. Fresh Li2O2 was deposited by evaporation of metallic Li under oxygen pressure of ∼10−7 mbar in the preparation chamber of the spectrometer. Dots show experimental data, while solid lines represent fitted curves. Components are shown for spectra recorded after Li2O2 deposition. Insets demonstrate component fractions in spectra before (left bars) and after (right bars) Li2O2 deposition. Ex situ XPS analysis of polished disc electrodes before and after discharge was performed using Kratos Axis ULTRA DLD instrument equipped with a monochromatic Al Kα (1486.7 eV) source. All spectra were fitted by Gaussian/Lorentzian convolution functions using Unifit 2014 software. Spectra background was optimized using a combination of Shirley and Tougaard functions simultaneously with spectral fitting. Atomic fractions were calculated from integral peak intensities obtained at fixed kinetic energy (200 eV) normalized by photoionization cross sections23 and photon flux (see the Supporting Information for details). Layer thickness was estimated using the Hill equation18 taking into account corresponding values of inelastic mean free path calculated using TPP 2M formula24 and atomic density (see the Supporting Information).

2p and C 1s spectra of the TiC single-crystal pristine surface are presented as gray lines in Figure 1a and b, respectively. The surface composition was shown to correspond to stoichiometric TiC within our accuracy; that is, both Ti and C atoms are present at the terraces. The Ti 2p spectrum demonstrates asymmetry that can be attributed to partial oxidation of Ti atoms at surface step edges. As can be expected from Gibbs free energy calculations (see Table S1), Li2O2 deposition results in surface oxidation that is in line with recently reported DFT calculations.25 Ti 2p spectrum shows evidence for the appearance of oxycarbide species, whereas elemental carbon with certain amounts of oxygen-containing surface groups and carbonate species is produced according to the C 1s spectrum. On the basis of these data, we can suggest the following scheme of titanium carbide oxidation by Li2O2:

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RESULTS AND DISCUSSION As lithium−air battery discharge results mainly in Li2O2, we performed photoemission studies of the surface chemistry on a TiC single crystal upon lithium peroxide deposition. Typical Ti 8250


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oxidation upon Li2O2 deposition. It should be mentioned, however, that the atomic structure of such elemental carbon produced by room temperature oxidation is not quite clear yet.19 While the peak position in C 1s spectrum fits well to sp2hybridized carbon, we observed a noticeable oxidation of these carbon species upon Li2O2 deposition in contrast to the previously reported stability of sp2 carbon in graphene.15 Creating an artificial carbon protecting layer by chemical vapor deposition of graphene on a TiC surface28 also enabled suppression of the reactivity toward Li2O2. The results of this experiment are illustrated in Figure 2. Both C 1s and Ti 2p

2/aTiC + 4Li 2O2 = 2/aTiOa C1 − a + Li 2CO3 + C + 3Li 2O

(1)

Indeed, the surface chemistry of TiC nanopowders used for fabrication of air electrodes in previous works17,18 is much more complex in comparison to the freshly cleaned single crystal. Carbide powders are covered by overlayers of complex composition, which are produced due to interaction of TiC with environmental oxygen.19−21,26 At the earliest stages, the reaction between a clean (100) surface of TiC and hollowpositioned oxygen molecule is favorable only due to the enthalpy of C−C or C−O bond formation. Oxidation includes formation of elemental carbon or CO/CO2 elimination accompanied by oxygen bonding to titanium, generation of oxycarbides of various composition, and further formation of TiO2 nuclei. At room temperature, carbon predominantly segregates at the surface, while higher temperatures lead to prevailing CO (and CO2) emission with TiO2 remaining on the surface. These steps are summarized in the following reactions: TiC + n/2O2 = TiC1 − nOn + nC

(2)

TiC + nO2 = TiC1 − nOn + nCO

(3)

TiC + 3/2nO2 = TiC1 − nOn + nCO2

(4)

TiC1 − nOn + (3 − 2n)/2O2 = TiO2 + (1 − n)CO

(5)

We further evaluate the “protecting” ability of the overlayers formed on TiC due to oxidation by molecular oxygen under different conditions. We start with oxidation of TiC single crystal at high temperature forming a TiO2 layer, which is anticipated to be stable toward further oxidation by Li2O2 due to the highest oxidation state of titanium in it. To prepare such a layer, we oxidized a clean TiC surface at 500 or 1100 °C under 10−7 mbar oxygen partial pressure. At 1100 °C, we observed a TiO2 layer with an estimated effective thickness of about 0.3−0.4 nm and elemental carbon residues (Figure 1c and d), while oxidation at 500 °C results in an oxide layer with higher carbon amount as it follows from Figure 3e and f. Unfortunately, both layers obtained at 500 and 1100 °C did not prevent reactions with Li2O2 (Figures 1c−f) that can be rationalized presuming the multicomponent nature and nonuniformity of such coatings.27 In both cases, we observed progressive surface oxidation that is represented by the following reaction schemes:

Figure 2. C 1s spectra TiC (755) surface of single crystal covered with graphene layer before (in gray, top) and after (in black, bottom) Li2O2 deposition. Spectra were acquired at 50 eV photoelectron kinetic energy providing surface sensitivity. Dots denote experimental data, while solid lines show fit results. Deconvolution into components is shown for spectra, obtained after Li2O2 deposition.

spectra practically do not change after Li2O2 deposition with the exception of the appearance of a tiny amount of carbonate that can be a result of incomplete graphene coverage. Being disregarded in previous works,17,18 the carbon layer demonstrates a reasonable protective ability against Li2O2. Its reactivity with superoxides,14 which are produced electrochemically during battery discharge, however, makes such coating not effective enough. We illustrate this by in situ spectroscopy of an operating electrochemical cell with solid-state lithium-conducting electrolyte, shown in Figure 3a. TiC nanopowder was deposited onto a solid glass-ceramic electrolyte disc and served as the working electrode. A piece of lithium foil was employed as the counter electrode. We used near ambient pressure X-ray photoelectron spectroscopy (NAP XPS) to probe the reactions between TiC and ORR products and intermediates, while the influence of side reactions with liquid electrolytes was avoided. As expected, galvanostatic discharge of the cell inside the XPS chamber leads to a gradual increase in the Li 1s and O 1s intensities, which can be explained by the formation of a discharge product layer on the surface (see Figure S1). According to our spectromicroscopy data shown in Figure 3g−i, the discharge products grow as monolayer-thin islands gradually covering the electrode surface. The surface of TiC nanopowder used as a positive electrode in our in situ cell (see the Supporting Information for more details on carbide nanopowder characterization) is covered by oxides and significantly contaminated by elemental carbon originating from natural oxidation and/or from the synthesis

2/bTiOx Cy + 4Li 2O2 = 2/bTiOx+b + 2/bCy − b + Li 2CO3 + C + 3Li 2O

TiOx + (2 − x)Li 2O2 = TiO2 + (2 − x)Li 2O

(6) (7)

Much lower reactivity with lithium peroxide was observed for a TiC surface preliminary exposed to near-ambient oxygen pressure (200 mbar) for 30 min at room temperature emulating natural oxidation. Such a procedure resulted in a surface layer consisting mainly of titanium oxycarbide phases and elemental carbon that can be seen from the Ti 2p and C 1s spectra in Figure 1g and h. In contrast to the clean surface and TiC, exposed to oxygen at high temperatures, both showing strong reactivity with Li2O2, crystal with carbon/oxycarbide overlayer demonstrated reasonable stability that can be ruled out by comparing the spectra in Figure 1c and d: the peak related to Li2CO3 is absent, and titanium does not undergo further 8251


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Figure 3. (a) Schematic of operando experiment with a two-electrode electrochemical cell consisting of lithium counter electrode, solid lithium conducting NASICON-type electrolyte, and TiC nanopowder, deposited onto electrolyte rough surface and served as a working electrode. The cell was assembled in Ar atmosphere, transferred to the spectrometer without contact to air, and connected to potentiostat via electrical feedthrough of the spectrometer. Experimental C 1s spectra of the TiC electrode before (b) and after (d) the cell discharge are shown by dots; solid lines represent fitting results. Components are colored filled. (c) Evolution of component fractions in Ti 2p spectrum during cell discharge in galvanostatic mode. (e) C 1s spectrum of TiC nanopowder electrode after Ar+ sputtering. The evolution of Ti 2p spectrum of such a “cleaned” electrode is shown in (f), while maps (g)−(i) illustrate effective thickness of the discharge product layer at different discharge depths for the cell with Ar+ sputtered TiC electrode. The thickness was calculated using attenuation of Ti 2p intensity (see the Supporting Information for details).

procedure as can be seen from the C 1s spectrum in Figure 3b. The effective thickness of the TiO2 layer was estimated to be about 0.6 nm. At the same time, no further oxidation of TiC particle core was observed (Figure 3c). Carbon, however, is slowly oxidized that we figure out by comparing the initial C 1s spectrum with that recorded after the cell discharge (Figure 3d) and revealing a significant amount of oxygenated groups and some carbonate on the surface upon discharge. This fact along with the observation of nearly no changes in the Ti chemistry suggests the conclusion that the carbon “protecting” layer degrades due to previously reported reactions with ORR intermediates.14,15 Again, elimination of surface layers from the TiC nanopowder promotes oxidation of the carbide itself. We used Ar+ bombardment for thinning the carbon layer. Such treatment resulted in a residual monolayer thickness of TiO2 coating on the nanoparticle surface. Figure 3e demonstrates the resulting C 1s spectrum of such nanopowder after Ar+ bombardment. Although not all of the elemental carbon was removed, discharge of the cell with an electrode prepared from such TiC powder resulted in a progressive increase of the TiO 2 component intensity during discharge (Figures 3f and S3). The evolution of the C 1s spectrum during discharge (Figure S4) gives also evidence for carbon oxidation. Together with Li2O2, the byproducts form a passivating layer covering TiC nanoparticles. While lithium peroxide is removed upon cell charge (see Figure S5), TiO2 stays at the surface. We further observe the same situation for mirror-polished polycrystalline TiC disc electrodes (see Figure S6) cycled in oxygen-saturated DMSObased electrolyte. Cyclic voltammograms in Figure 4a reveal a significant electrode passivation. Indeed, a freshly polished TiC electrode has a natural oxide layer on its surface, which is known to increase the work function by at least 1.3 eV.20

Nevertheless, the electrode demonstrated a pronounced cathodic peak associated with ORR on the TiC surface. Current density for the first cathodic sweep is comparable to that for glassy carbon (Figure S7); however, Figure 4a clearly shows the continuous decrease of the cathodic peak current for further cycles. The peak position shifts toward more negative potentials. We believe that the observed fade in the peak current density along with peak position shift are caused by electrode passivation with an insulating layer inhibiting electron transport, while mass transport conditions remain unchanged from cycle to cycle. Even faster passivation was observed for TiN and TiO, both possessing metallic conductivity (∼1−10 kS/cm)29 similar to TiC and reasonable specific weight. Even at the first cycle, ORR occurred at very high overpotentials (see Figure 4b and c). Growth of the passivating layers upon cycling of polycrystalline TiC was further confirmed by ex situ XPS analysis revealing progressive oxidation of TiC surface upon cycling. As can be seen from the Ti 2p spectrum in Figure 5a, the surface of freshly polished electrode is covered by a natural oxide/ oxycarbide film and adventitious carbon.30,31 TiOxCy is represented by two components (II and III), which are well resolved in high-resolution spectra recorded for the TiC single crystal (Figure 1). Analysis of the angular dependence of Ti 2p component relative intensities shown in Figure S8 reveals that the surface is covered by a TiO2 layer, while oxycarbide is distributed across the surface film and a subsurface bulk. The spectral changes after exposure of the electrode to electrolyte were found to be negligible (see Figure S9). In contrast, fast increase in TiO2related intensity is observed upon cycling of the electrode in oxygen-saturated electrolyte, thus giving a hint that TiC side reactions with Li2O2 and superoxide intermediate species generated during cathodic sweeps are the reason for surface 8252


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Figure 4. Cyclic voltammograms of polished TiC (a), TiN (b), and TiO (c) dense polycrystalline electrodes recorded for freshly polished samples. Voltammograms were recorded in 0.1 M LiClO4 solution in oxygen-saturated DMSO at a sweep rate of 100 mV/s. Inset in (a) shows the picture of TiC, TiN, and TiO (left to right) electrodes after polishing.

Figure 5. (a) Ti 2p spectra (obtained ex situ, photon energy 1486.6 eV) and component fractions distribution for pristine dense polycrystalline TiC electrode after exposure to the electrolyte and the same electrode, cycled for 1, 4, and 30 voltammetry cycles. The electrode was taken from the cell after anodic sweeps, so the unreacted discharge products were removed from the surface by oxidation. The electrode was washed with pure DMSO and transferred to the ultrahigh vacuum (UHV) environment of the spectrometer without exposure to air. (b) Dependencies of produced Li2O2 layer thickness, effective thickness of oxidized Li2O2 (calculated from cyclic voltammograms), and TiO2 passivation layer thickness on cycles number (calculated from XPS data). Corresponding C 1s spectra are shown in Figure S10.

oxidation. As expected, the C 1s spectra of the electrode shown in Figure S10 reveal the elemental carbon accumulation on the surface during cycling due to the discussed reactions of carbide oxidation. Being typical for stoichiometric titanium dioxide after 30 cycles (Figure 5a), the related Ti 2p component (component IV) is initially located at slightly lower binding energies that can be explained by the presence of oxygen vacancies in the TiO2 passivating layer.32 Oxycarbides also remain on the surface, demonstrating no growth of the corresponding peak relative intensity in both Ti 2p and C 1s spectra (Figures 5a and S10). Indeed, passivation by TiO2 layer leads to inhibition of oxygen redox processes. Reported earlier for OER,18 TiO2 passivation layer of course hinders ORR also. Assuming Li2O2 to be a main discharge product, we use cathodic charge to calculate the effective thickness of the peroxide film grown on the electrode surface and to show that it correlates with the TiO2 layer thickness estimated using the Hill equation and our XPS data (Figure 5b). After about 10 cycles, the TiO2 thickness stabilizes at nearly 4 nm that is close to a value estimated as critical for maintaining electron tunneling sufficient for a current density of ∼1 mA/cm2 (see Figure S11).33 Although in our experiments the observed current density was even little lower, we observed a 2-fold decrease in cathodic charge passed and overpotential increase by ca. 0.1 V, while the transfer coefficient α stays nearly at the same value of about 0.5 (Figure S12). Thus, these results suggest that such oxidation leads to

excessively thick passivating film. Anodic charge spent primarily for Li2O2 oxidation on each cycle is constantly lower, which may indicate a partial loss of ORR products due to formation of soluble LiO2 intermediates,11 that diffuse away from the electrode, or side reactivity of oxygen reduced species with the electrode or electrolyte.34 Coloumbic efficiency is estimated to be ca. 65−70% and does not noticeably change with the number of cycles.

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CONCLUSIONS Titanium carbide, which is now considered as one of the candidates for lithium−air positive electrode material, reveals a rich surface chemistry in a TiC−Li−O system. TiC shows an intrinsic tendency to be easily oxidized by both oxygen and lithium peroxide; however, it can be stabilized by surface oxide layer playing the role of a protective layer. Previous investigations of TiC positive electrodes for lithium−air batteries17,18 did not take into account elemental carbon segregating at the surface during room-temperature reactions 8253


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with O2 or Li2O2, while we demonstrate that this is an important component preventing progressive oxidation of the TiC surface. Unfortunately, such an approach cannot be employed in practice as carbon is known to be oxidized in the presence of ORR intermediates.14,15 In contrast, a TiO2 layer can resolve this issue. At the same time, in stoichiometric form it exhibits a weakly O-binding surface35 that is beneficial for Li2O2 formation.36 TiO2, however, is a wide band gap semiconductor (with a gap of about 3 eV) that will unavoidably lead to ORR kinetic limitations. So, the layer thickness should be optimized to provide both reliable surface protection and a reasonable electron transport rate. Although our attempts of obtaining a uniformly dense thin film were unsuccessful so far and Li2O2 deposition still results in TiC surface degradation, we believe that further synthetic efforts can enable surface protection by dense carbon-free titanium dioxide. In turn, it will open new possibilities for TiC in lithium−air batteries.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.6b03195. Thermodynamic calculations, characterization of disc polycrystalline electrodes (Raman, SEM, XRD), additional electrochemical and XPS data, along with detailed descriptions of data processing procedures (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Author Contributions ■

A.Y.K. and E.Y.K. contributed equally.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS

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REFERENCES

We are grateful to HZB and ELETTRA for beamtimes granted. A.Y.K., E.Y.K., L.V.Y., and D.M.I. thank the Russian German laboratory at HZB for support provided. We are grateful to Prof. O. Rader who provided TiC single crystal for the experiments. We acknowledge K. I. Maslakov for his kind assistance with experiments involving laboratory X-ray photoemission spectrometer provided thanks to Lomonosov Moscow State University Program of Development. This work was financially supported by the Russian Ministry of Science and Education (RFMEFI61614X0007) and Bundesministerium für Bildung and Forschung (project no. 05K2014) in the framework of the joint Russian-German research project “SYnchrotron and NEutron STudies for Energy Storage (SYNESTESia)”.

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