Ordered Mesoporous Titanium Nitride as a Promising Carbon-Free

Despite the extraordinary gravimetric energy densities, lithium-oxygen (Li-O2) batteries are still facing a technological challenge; limited round tri...
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Ordered Mesoporous Titanium Nitride as a Promising Carbon-Free Cathode for Aprotic Lithium-Oxygen Batteries Byung Gon Kim,†,∥ Changshin Jo,‡,§,∥ Jaeho Shin,† Yeongdong Mun,‡ Jinwoo Lee,*,‡,§ and Jang Wook Choi*,† †

Graduate School of Energy, Environment, Water, and Sustainability (EEWS) and KAIST Institute NanoCentury, Korea Advanced Institute of Science and Technology (KAIST), 291 Daehakro, Yuseong-gu, Daejeon 34141, Republic of Korea ‡ Department of Chemical Engineering and §School of Environmental Science and Engineering, Pohang University of Science and Technology (POSTECH), 77 Cheongam-Ro, Pohang, Gyeongbuk 37673, Republic of Korea S Supporting Information *

ABSTRACT: Despite the extraordinary gravimetric energy densities, lithium-oxygen (Li-O2) batteries are still facing a technological challenge; limited round trip efficiency leading to insufficient cycle life. Recently, carbonaceous electrode materials were found to be one of the primary origins of the limited cycle life, as they produce irreversible side products during discharge. A few investigations based on noncarbonaceous materials have demonstrated largely suppressed accumulation of irreversible side products, but such studies have focused mainly on the materials themselves rather than delicate morphology control. As such, here, we report the synthesis of mesoporous titanium nitride (m-TiN) with a 2D hexagonal structure and large pores (>30 nm), which was templated by a block copolymer with tunable chain lengths, and introduce it as a stable air-cathode backbone. Due to the well-aligned pore structure and decent electric conductivity of TiN, the battery reaction was quite reversible, resulting in robust cycling performance for over 100 cycles under a potential cutoff condition. Furthermore, by protecting the Li metal with a poreless polyurethane separator and engaging a lithium iodide redox mediator, the original capacity was retained for 280 cycles under a consistent capacity condition (430 mAh g−1). This study reveals that when the appropriate structure and material choice of the air-cathode are coupled with an advanced separator and an effective solution-phase redox mediator, the cycle lives of Li-O2 batteries can be enhanced dramatically. KEYWORDS: mesoporous structure, poreless separator, redox mediator, self-assembly, titanium nitride

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impairing the reversibility of each cycle and eventually leading to short cycle life. It is thus desired to find alternatives to carbon materials for the air-cathode to avoid this critical shortcoming. Along this research direction, recent studies investigating Au,13 TiC,14 Ti4O7,15 TiSi2,16 Ru,17 RuO2,18 Mo2C,19 and Co3O420,21 as air-cathode materials are remarkable. In general, an air-cathode should meet the following standards: high electric conductivity, proper porosity, chemical stability against oxygen radical attack, and electrochemical stability with the electrolyte in the given operation voltage window. In particular, we pay attention to pore size to warrant the reversible reaction of a discharging product, Li2O2. In fact, it turns out22,23 that micropores (30 nm). As for material selection, titanium nitride (TiN) was chosen because TiN has been used28−30 as catalysts or catalyst supports for various chemical reactions, which particularly take place in electrochemical energy conversion and storage systems due to its outstanding resistance against corrosion and decent electric conductivity. In spite of these conspicuous advantages, the synthesis of ordered mesostructured TiN has proven to be challenging due to the lack of a controllable sol−gel process involving titanium and nitrogen precursors, along with strict requirements of air- and water-free conditions for the conversion of TiO2 to TiN. So far, in generating pure and ordered mesoporous TiN, the nitridation of nanostructured TiO2 at a high temperature using NH3 gas is the only possible way.31,32 However, many attempts have failed due to the thermal instability of preformed mesoporous TiO2, with only composite structures that include mesoporous TiN having been reported.33−35 In addition, although direct growth of porous or tubular TiN on a substrate36 has been demonstrated, the tuning of TiN loading, scalability, and material cost are yet to be clarified. In an effort to overcome the difficulty of producing ordered meso-structured TiN, in the present investigation, we develop a synthetic scheme based on an evaporation-induced selfassembly (EISA) process using a block copolymer (BCP) with a high molecular weight. The final mesoporous TiN, namely m-TiN, has a 2D hexagonal structure bearing mesopores with a diameter of 37 nm. Furthermore, titanium oxynitride (TiOxNy) and titanium suboxide (TinO2n−1 under Magneli phase) layers with relatively high conductivity exist15,28,37 on the surface of m-TiN and retard the oxidation of TiN to TiO2, thus preserving the high electrical conductivity of TiN in the long term, in contrast with previous cases based on bulk TiN that suffer from inferior conductivity and electrochemical performance due to the formation of a surface oxide layer.14,38 Utilizing these structural and physical properties, when examined as an air-cathode, m-TiN exhibited cyclability for over 100 cycles in a potential cutoff condition (2.0−4.7 V vs Li/Li+). Moreover, the cycle life was further extended to more than 380 cycles upon protection of the Li metal anode and the addition of a redox mediator, LiI. This study indicates that m-TiN is a promising option as an aircathode and that air-cathode material selection and meticulous cell design are also crucial factors in guaranteeing the sustainable operation of Li-O2 cells.

Scheme 1. Schematic Representation of Block Copolymer Directed Evaporation-Induced Self-Assembly and Subsequent Heat-Treatment Processes for the Synthesis of Ordered m-TiN

assembly of a lab-made amphiphilic BCP, poly(ethylene oxide)block-poly(styrene) or PEO-b-PS in abbreviation (Mn = 25 kg mol−1, 20 wt % of PEO block). When the EISA process progressed in a Petri dish at 50 °C, mesoscopic phase separation between the hydrophilic (hydrolyzed titanium isopropoxide (TTIP)/PEO) blocks and the hydrophobic (PS) blocks was formed. When TTIP, a titanium-containing precursor in our case, is added to the polymer solution, hydrolyzed TTIP selectively binds with the hydrophilic PEO blocks via hydrogen bonding. The final structure of m-TiN was completed after a nitridation of mesoporous-TiO2 (m-TiO2), as described in the Methods section. See Figure S1 for additional scanning electron microscopy (SEM), scanning transmission electron microscopy (STEM), and energy-dispersive X-ray spectroscopy (EDS) results of m-TiN. It is known39 that the size of meso-phase separation can be tuned by controlling the block lengths of PEO-b-PS. In particular, compared to pluronic BCPs, PEO-b-PS can develop into mesoporous materials with thicker pore walls and larger pore sizes (>30 nm).40,41 The enlarged thickness of the pore wall would indeed be beneficial for maintaining the overall structure during high-temperature nitridation that gives rise to large stress. Note that due to stress generation, ordered m-TiN has never been produced directly from the corresponding TiO2 counterparts via nitridation.33 The density difference between both materials explains the stress generation: 3.78 g cm−3 (anatase TiO2) vs 5.22 g cm−3 (cubic TiN). The porosity of m-TiN was investigated by carrying out N2 physisorption analysis (Figure 1a). The isotherms of both mTiO2 and m-TiN belong to the type-IV with hysteresis present at P/P0 = 0.85−0.95, which is reflective of mesoporous structures containing uniform and large pores. The Barrett− Joyner−Halenda (BJH) pore size distributions calculated from the adsorption curves reveal that the main pore sizes of m-TiN and m-TiO2 are 37 and 23 nm, respectively (Figure 1a, inset). The increase in pore size after nitridation is attributed to a thermal shrinkage of the pore wall and a change in crystal structure during high-temperature treatment. While the X-ray diffraction (XRD) spectra (Figure 1b) of m-TiN and m-TiO2 confirm their crystal structures under the cubic (JCPDS no.: 87-0633) and anatase (JCPDS no.: 78-2486) symmetry, respectively, the XRD spectra also indicate that their primary crystalline sizes are similar according to the Debye−Scherrer

RESULTS AND DISCUSSION Our synthetic strategy to produce m-TiN is illustrated in Scheme 1. The mesoscopic structure stems from the self1737

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Figure 1. (a) N2 physisorption isotherms (inset: pore size distributions), (b) XRD patterns, and (c) SAXS patterns of m-TiO2 and m-TiN. (d) SEM, (e) TEM, and (f) HR-TEM images of m-TiN. The insets in (f) are an HR-TEM image magnified from the red box (bottom) and an FFT pattern of the image in (f) (top).

Figure 2. (a) CV profiles of the m-TiN cells under pure Ar and O2 when measured at 0.1 mV s−1 in the voltage range of 2.0−4.5 V. (b) The voltage profiles of the OMC and m-TiN cells and (c) the cycling performance of the m-TiN, m-TiO2, OMC, and bulk-TiN cells when measured at a current density of 70 mA g−1 for 100 cycles. (d) TEM images of m-TiN and bulk-TiN. (Inset) Magnified TEM image of m-TiN. (e) Ti 2p XPS spectra of pristine bulk-TiN and m-TiN and (f) m-TiN after 150 cycles.

g−1, and 0.27 cm3 g−1, respectively. The pore properties of the porous materials used in this study are summarized in Table S1 in the Supporting Information. The structural information was further attained by employing small-angle X-ray scattering (SAXS) analyses (Figure 1c). The as-self-assembled sample before heat-treatment exhibited highorder reflections at the scattering wave vector (q) values corresponding to the (200), (210), and (300) orientations along with the primary (100) peak, indicating that pores are arranged in a 2D hexagonal array upon the self-assembly.43 The d-spacing, d*, in pore can be correlated to the peak position, q*

equation:42 16.3 nm (TiN) vs 15.9 nm (TiO2). Thus, the higher density of TiN (5.22 g cm−3) is the main origin of the shrink in wall thickness. The Brunauer−Emmett−Teller (BET) surface area and pore volume of m-TiN are 58 m2 g−1 and 0.36 cm3 g−1, respectively, whereas those of the templating m-TiO2 are 65 m2 g−1 and 0.30 cm3 g−1. Using BCP with a different length of PS blocks, another m-TiO2 with a similar BET surface area to that of m-TiN was also prepared as a control sample (Figure S2) to clarify the effect of the enhanced electric conductivity of TiN on the electrochemical performance. The pore size, BET surface area, and pore volume of this m-TiO2 are 33 nm, 57 m2 1738

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Figure 3. (a) XRD spectra of m-TiN at the pristine, first discharged, first charged states, and after 100 cycles. (b) SEM images of m-TiN at the pristine, first discharged, and first charged states when measured at a current density of 70 mA g−1. (c) Li 1s XPS spectra of m-TiN after the first discharge and charge and the fifth discharge and charge. (d) SEM images of OMC and m-TiN after the first full discharge.

value by d* = 2π /q*. Compared to that of the as-selfassembled, the (100) peak of m-TiO2 shifted to a larger q value, reflecting the decreased d-spacing in the pore arrangement (35.7 to 34.3 nm). This observation can be explained by the removal of organic components during heat treatment.43,44 After nitridation, the primary (100) peak shifted to a lower q value (d(100) = 39.5 nm), which is consistent with the increased pore size observed in the BET analysis. Also, the high-order scattering peaks disappeared, and the primary (100) peak became broadened, which indicates a slight loss in long-range ordering and/or modified local pore structure.45,46 The 2D hexagonal pore arrangement of m-TiN was reflected in its SEM and TEM images (Figure 1d−f and Figure S1a). The SEM images of m-TiN in Figure S1b,c show a highly ordered hexagonal array of pores along the [100] direction. However, the morphology of the pore wall was changed from a highly dense, monolithic structure in m-TiO2 to assembled nanoparticles (Figures S2 and S3). Nonetheless, the selfsupportive structure was maintained after the nitridation process. The high-resolution TEM (HR-TEM) image (Figure 1f) reveals the cubic structure of TiN, which is in good agreement with the XRD spectrum (Figure 1b). Therefore, structural characterizations support the first synthesis of pure and self-supportive TiN with a highly ordered mesoporous structure through soft template. Ordered mesoporous carbon (OMC), a control sample to observe the effect of TiN on electrochemical performance, was found to have a similar 2D hexagonal meso-porosity with the pore dimension centered near 36 nm (Figure S4a,b,d,e). The Raman spectrum (Figure S4f) of OMC exhibited an enhanced D-band peak at 1350 cm−1 compared to that of single wall carbon nanotubes, indicating the amorphous and defective character of the carbon in OMC. Commercial bulk-TiN (bulk-TiN) (Figure S5) was also prepared to see the effect of the porosity of m-TiN. In order to assess m-TiN as an air-cathode, coin-type cells were prepared by employing Li metal disk as a reference/ counter electrode. See the Methods section for detailed cell

preparation. Figure 2a shows cyclic voltammetry (CV) profiles of the m-TiN cell under Ar and O2 atmospheres. The m-TiN cell in the O2 atmosphere clearly showed reversible cathodic (ORR) and anodic (oxygen evolution reaction (OER)) peaks, whereas no visible peaks were observed from the Ar atmosphere, verifying that the redox reactions observed from the O2 atmosphere involve O2 and that m-TiN does not trigger any significant side reactions.47,48 Consistently, the galvanostatic discharging scan did not show any tangible capacity in the Ar atmosphere (Figure S6a). The CV results provided the first indication of better cycling stability of m-TiN over OMC (Figure S7), and while the current profiles of the OMC cell gradually decayed with cycling, the m-TiN cell was able to maintain the initial profile to a great extent. Both cells were further tested under the galvanostatic mode with a voltage cutoff condition, 2.0−4.7 V (Figure 2b). As in the CV tests, m-TiN showed more sustainable cycling behavior than OMC for 100 cycles; the specific capacity of OMC dropped upon cycling, whereas m-TiN was much better at maintaining the initial profiles, supporting once again the improved durability of m-TiN for Li-O2 cells. This cycling performance of m-TiN is remarkable, as the capacity retention is comparable or superior to those of previous cells14,15,17,19,49,50 that were tested under similar conditions (potential cutoff, noncarbonaceous air-cathode, and ether-based electrolyte). The distinct cycling performances of m-TiN, mTiO2, OMC, and bulk-TiN are comparatively displayed in Figure 2c. The m-TiN started with 390 mAh g−1, and its specific capacity rose to 830 mhA g−1 at the 20th cycle and finally ended at 623 mAh g−1 at the 100th cycle. The gradual increase of the specific capacity is attributed to a certain activation process, such as the establishment of the three-phase boundary,18,51,52 although the detailed mechanism requires further investigation for clarification. The activation process was reflected in the decreased interfacial resistance from electrochemical impedance spectroscopy (EIS) results during the corresponding cycles (Figure S8a). In contrast, OMC suffered 1739

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Figure 4. (a) The first voltage profile of the m-TiN cell when measured at 50 mA g−1, along with schematic illustrations showing the m-TiN morphology change during discharge and charge. (b−e) TEM images at the (upper) low and (bottom) high magnifications of the m-TiN obtained at the points marked in (a). The insets in (d) show magnified TEM images from the marked boxes.

of Li2O2 (Figure 3a),10 whereas the spectra at the pristine, first charged, and 100th charged states did not consistently show those peaks, supporting the reversibility of the m-TiN cell for a great number of cycles. The reversible formation and decomposition of Li2O2 were also observed by ex situ SEM imaging (Figure 3b); at the end of discharge in the first cycle, disk-like Li2O2 was formed. However, this Li2O2 disappeared at the end of the subsequent charge, making the overall morphology consistent with that of the pristine state. From XPS analysis (Figure 3c), the spectrum of the first discharged sample can be deconvoluted to the peaks corresponding to LiF, Li2CO3, and Li2O2.10 The spectrum, however, disappeared after the first charge, and the reversible trend was maintained in the fifth cycle as well. Nonetheless, the spectrum at the fifth discharge was slightly changed compared to that at the first discharge, which is likely to be reflective of inevitable parasitic reactions associated with the Li2O2-induced electrolyte instability and the electrolyte oxidation at high OER voltages.7,8 Interestingly, the Li2O2 morphology of the m-TiN cell at the first discharge is clearly distinct from that of the OMC cell (Figure 3d); the Li2O2 of the m-TiN cell was shaped like disks, whereas the OMC cell produced toroid-like Li2O2. The different nucleation environments can explain this morphological difference; the OMC with a large number of defects on the surface prefers superoxide-to-peroxide nucleation at multiple sites, which eventually leads to a toroid-like morphology as the crystals grown from multiple sites combine with each other. By contrast, m-TiN with much less defects allows the crystals to grow from a single or a small number of nucleation points, resulting in continuous crystalline growth along the same plane to form the disk-like morphology. This rationale is consistent

from a drastic capacity decay from the early period of cycling as indicated by its EIS results (Figure S8c). The specific capacities of m-TiO2 and bulk-TiN were quite low from the beginning of cycling, presumably due to the relatively low electric conductivity of TiO2 (Figure S9) and the insufficient surface area of bulk-TiN (Table S1), respectively. Bulk-TiN also forms a TiO2 layer on its surface upon exposure to air, which additionally impairs the electric conductivity of bulk-TiN if the thickness of the formed layer exceeds several nanometers.14,28,38,53 In fact, the existence of a surface rutile-TiO2 layer was confirmed by an etching test (Figure S10) and Raman spectroscopy (Figure S11), and the thickness of this layer was determined to be approximately 5 nm via TEM analysis (Figure 2d). Due to the insufficient electric conductivity originating from the relatively thick TiO2 layer, bulk-TiN was previously deemed inappropriate for air-cathodes.14,38 Note that the mTiN, bulk-TiN, and m-TiO2 electrodes are all free of a carbon conducting agent to avoid side reactions related to carbonaceous electrodes so that the electronic conductivity can be limited by the choice of air-cathode backbone. In contrast with bulk-TiN, m-TiN showed a negligible oxide layer (Figure 2d), and its X-ray photoelectron spectroscopy (XPS) spectrum (Figure 2e) indicates that the surface of m-TiN contains titanium oxynitride (TiOxNy) and titanium suboxide (TinO2n−1: Magneli phase) with higher electric conductivity than that of TiO2.15,28,29,37 Moreover, the surface compositions of m-TiN remained unchanged for 150 cycles (Figure 2f). The reversibility of the m-TiN cell was investigated using a series of analytical tools. The XRD pattern at the first discharged state exhibited peaks at 32.8° and 34.9°, corresponding to the (100) and (101) crystalline orientations 1740

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Figure 5. (a) SEM images of m-TiN at different cycle numbers. (b) C 1s and Li 1s XPS spectra of m-TiN after 150 cycles. (c) FT-IR spectra of m-TiN at different cycle numbers. (d) 1H NMR spectra obtained from the electrolyte of the m-TiN cell after the first and 150 cycles. (e) XRD spectra and color of the Li metal anode of the m-TiN cell at the pristine state and after 90 cycles.

with the previous reports.49,51,54−58 Favorably, the disk-like Li2O2 morphology is known49,54,57 to be more suitable for the decomposition of Li2O2 during charge and is thus more likely to offer reversibility upon cycling. Importantly, the XPS (Figure 2f) and XRD (Figure 3a) spectra were largely consistent for 150 and 100 cycles, respectively, demonstrating sustainable operation with m-TiN in Li-O2 batteries. The structure of m-TiN was monitored throughout the first cycle by taking TEM (Figure 4) and SEM (Figure S12) images at different discharge and charge points marked in Figure 4a. The pristine m-TiN exhibited well-aligned channels (Figure 4b) through which the electrolyte containing Li ions and oxygen molecules can diffuse.23,59 These channels began to be filled with discharge products when the specific capacity reached 200 mAh g−1 (Figure 4c). At this point, the Li2O2 growth was insignificant from the outer surface of m-TiN according to the SEM image (Figure S12), implying that much of the newly formed Li2O2 is confined within the pores of m-TiN. Such spatial restriction of Li2O2 appears to be better at avoiding detachment of Li2O2 from the air-cathode, which is indeed crucial for long-term cyclability.24 Once fully discharged, the disk-like Li2O2 began to show up on the surface of m-TiN particles, as seen from the TEM (Figure 4d), SEM (Figure S12), and EDS (Figure S13) results. A lattice distance obtained from the orange box in Figure 4d corresponds to the (101) crystalline orientation of Li2O2 and thus verifies the presence of crystalline Li2O2,60 consistent with the XRD (Figure 3a) and XPS (Figure 3c) results. Finally, the TEM images at the fully charged state showed a well-recovered internal channel structure and a clear surface (Figure 4e), further supporting the reversibility of the m-TiN cell. In spite of the improved cycling performance of the m-TiN cell compared to the other control cells, capacity fading was observed after around 80 cycles. To find the reasons for this

fading, SEM, XPS, and Fourier transform infrared (FT-IR) analyses were carried out. The SEM images (Figure 5a) of the m-TiN air-cathode showed a gradual accumulation of parasitic reaction products with cycling, providing a clue for the capacity decaying.15,18,50,61 According to the XPS (Figure 5b) and FT-IR (Figure 5c) results, these products were identified as carbonates and carboxylates,10,62 and the corresponding peaks grew as cycling progressed, reconfirming that the accumulation of these side products is the primary reason for the capacity fading of the m-TiN cell. The increased interfacial resistance (Figure S8b) and carbon content from EDS elemental mapping (Figure S14) can be understood along the same line of logic. To further elucidate the parasitic reactions, the separator soaked with the remaining electrolyte was characterized with 1 H nuclear magnetic resonance (NMR) spectroscopy (Figure 5d). The 1H NMR spectrum at the 150th cycle exhibited peaks at 8.41 and 4.8 ppm, which are related to formate and water, respectively.63 The accumulation of the aforementioned side products including formate and the generation of water are both attributed8,9,11 to the decomposition of ether solvents by oxidative potentials during charge as well as the attack of reactive intermediate compounds during discharge. The generation of water is also contributed by the decomposition of poly(vinylidene fluoride) (PVDF) binder by the attack of reactive discharge intermediate compounds.64 In addition, LiOH was detected from the Li metal anode by XRD characterization (Figure 5e) along with other side products according to the C 1s branch in the XPS spectra (Figure S15). The LiOH formation can be explained by crossover oxygen65 and the water formed by the aforementioned mechanism.9,64 This degradation of the Li metal anode is another critical reason for the capacity fading of the m-TiN cell. As can be expected from the unstable character of the carbon electrode material, the OMC cell displayed more side products on the 1741

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Figure 6. (a) The first voltage profiles of the OMC, m-TiN, OMC + LiI, m-TiN + LiI, and m-TiN + LiI + PU cells. (b) The cycling performances of the same five cells. (c) SEM images and (d) C 1s XPS spectra of the m-TiN + LiI cell and OMC + LiI cell after 130 cycles. (e) XRD spectra and color of the Li metal anode (inset) of the m-TiN + LiI cells with and without PU separator after 80 cycles. All of the electrochemical tests in (c−e) were based on a current density of 70 mA g−1 with a fixed capacity of 430 mAh g−1.

characteristic of carbonaceous electrode materials for sustainable operation of Li-O2 cells. In contrast, the m-TiN and mTiN + LiI cells were able to extend the cycle life by mitigating the parasitic reactions7,8,11 involving carbonaceous electrodes and the accumulation of discharge products. This was clearly shown by the relatively latent drop in capacity at around the 102th and 130th cycle, respectively, compared to the OMC cells. The SEM images (Figure 6c) taken for the m-TiN + LiI and OMC + LiI cells after 130 cycles reflect the distinct interfacial stability between both cells. The m-TiN + LiI cell preserved the original powder morphology of m-TiN to a great extent, whereas the air-cathode of the OMC + LiI cell was fully covered with side products, burying the original powder morphology of OMC. In the case of XPS spectra (Figure 6d) taken after the same number of cycles, the m-TiN + LiI cell exhibited much more suppressed peaks related to side products, including Li2CO3 (289.7 eV) and O−CO (289 eV), compared to the OMC + LiI cell. As can be seen from the SEM images and XPS spectrum of the m-TiN cell in Figure S18, the formation of unwanted side products was markedly lessened for the m-TiN + LiI cell in comparison with the mTiN cell. However, even the m-TiN + LiI cell was able to extend the constant capacity period for only additional 28 cycles compared to the m-TiN cell. We suspected that the destabilization of the Li metal anode originating from crossover oxygen and water as well as direct reactions with oxidized LiI contributes substantially to the limited cycle life of the m-TiN + LiI cell.65,70 As evidence of such a scenario, LiOH was detected from the Li metal anode after 80 cycles by XRD analysis (Figure 6e). The spontaneous reaction70−72 of the oxidized LiI with the Li metal anode was also captured by the I 3d branch of the XPS spectrum (Figure S19a) of the Li metal anode. Furthermore, the copresence of other side products such as carbonates, carboxylates, and binder

electrode surface even after a smaller number of cycles (Figure S16),7,8 which is also reflected in its increased interfacial resistance with cycling (Figure S8c). In order to extend the cycle life, a redox mediator (RM), lithium iodide (LiI),66 was employed in the m-TiN cell because the RM is known to facilitate the decomposition of Li2O2 and reduce the charging overpotential by mediating the charge transfer during charge, although LiI may promote side reactions involving LiOH in the presence of water.67,68 In particular, when a constant capacity condition is applied, the suppressed overpotential is beneficial in extending the cycle life.4,69 Figure 6a displays the initial voltage profiles of the OMC and m-TiN cells with and without LiI. It is observed that the OMC cell showed lower overpotentials than those of the m-TiN cell throughout the first discharge, regardless of the RM addition. The distinct discharging behaviors are ascribed to the less conductive oxynitride/suboxide layer on the surface of m-TiN and the smaller surface area of m-TiN. The effect of the RM in lowering the overpotential was more prominent during charge, as both cells exhibited significantly reduced charging overpotentials upon the addition of LiI. Despite the distinct electric conductivities of m-TiN and OMC, both the m-TiN + LiI cell and the OMC + LiI cell showed almost identical charging profiles, indicating that charge transfer is more critical for efficient charging process than the electric conductivity of aircathode.3 The series of capacity retention results (Figure 6b) measured under a constant capacity condition (430 mAh g−1) reveals the effects of the electrode material and the RM more explicitly (Figure S17 for voltage profiles of all samples during cycling). The inferior cycling performance of OMC is most evident. Even with the addition of LiI, the capacity fading was more drastic than the other three cells based on m-TiN. The OMC and OMC + LiI cells started to drop in capacity at around the seventh and 12th cycle, respectively, confirming the undesirable 1742

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(Mn = 30 kg mol−1), HCl, H2O, and THF in a mass ratio of 1:0.2:0.34:18. The solution was first dried at 50 °C, then cured at 100 °C. The resulting film was calcined at 800 °C under N2 atmosphere, followed by hydrofluoric acid etching to remove SiO2. Li-O2 Battery Fabrication and Electrochemical Tests. For preparation of the air-cathode, m-TiN powder and PVDF were dispersed in NMP in a weight ratio of 95:5 without any conducting agent. The slurry was then cast onto a gas diffusion layer substrate (TGPH-090, 16 mm in diameter, Toray, Japan) serving as a current collector, followed by a drying step at 70 °C for 12 h under vacuum to remove residual solvent. The loading density of the electrode components was 0.4−0.5 mg cm−2. Based on this electrode loading, the total capacity is 0.35−0.43 mAh for the cells measured under a fixed capacity of 430 mAh g−1. The electrochemical tests were carried out using modified 2032-coin cells, whose bottom covers have 20 holes with a diameter of 1 mm to facilitate oxygen flow. Each cell was composed of a Li metal anode (thickness = 600 um, Honjo, Japan), a glass fiber membrane (GF/D, Whatman, USA) impregnated with electrolyte, and the as-prepared air-cathode. The electrolyte was prepared by dissolving 1 M LiClO4 in TEGDME and was stored in a bottle containing 3 Å molecular sieves inside an Ar-filled glovebox. LiClO4 was chosen because other fluorine-containing salts are known to be decomposed by superoxide radical attack.75 The electrochemical tests of the Li-O2 cells were conducted under CV and galvanostatic modes using a battery cycler (WBCS 3000, WonAtech, Korea) at room temperature. CV was carried out at a scan rate of 0.1 mV s−1 in the voltage range of 2.0−4.5 V vs Li/Li+. The galvanostatic charge/ discharge tests were conducted at current densities of 50, 70, and 100 mA g−1 in the voltage range of 2.0−4.7 V vs Li/Li+. The upper voltage cutoff was set to 4.7 V to avoid electrolyte decomposition (Figure S6b).9,47 For the electrochemical tests with LiI, the electrolyte was prepared by adding 0.05 M LiI in the same electrolyte, and a poreless PU separator was located between the GF and the Li metal to prevent Li metal degradation from crossover oxygen and water during cycling. The PU separator was prepared following a previous report.70 For ex situ analysis of the Li metal, a polyethylene (SK innovation, Korea) separator was added between the PU separator and the Li metal anode to prevent the PU separator from sticking to the Li metal anode. EIS analysis was carried out by using a VSP tester (Biologic, France) in the frequency range of 10 mHz to 1 MHz. All electrochemical tests were conducted inside a sealed chamber full of pure oxygen (99.999%) at 1 atm to avoid moisture contamination (Figure S22). Characterization of Synthesized Samples, Air-Cathodes, and Li Metal Anodes. For characterization of the electrodes after electrochemical cycling, the cells were disassembled in an Ar-filled glovebox, and the air-cathodes and Li metal anodes were washed with DME and dried under vacuum for 12 h in an antechamber connected to a glovebox. XRD patterns of the synthesized powders and aircathodes were obtained using an X-ray diffractometer (Bruker D8 Advanced XRD/micro XRD, Bruker/Rigaku, USA/Japan) with Cu Kα (λ = 0.15406 nm). To avoid exposure to air and moisture, a gastight XRD holder was used. SAXS analysis was performed for the structural characterization of mesoporous materials using the 4C SAXS beamline at Pohang Light Source (PLS, Korea). The morphologies of the synthesized powders and the air-cathodes were characterized by fieldemission SEM (FE-SEM, Hitachi S-4200/S-4800, Japan) and fieldemission TEM (FE-TEM, Tecnai F-30 S-Twin, FEI, USA). The elemental mapping of the synthesized powders was conducted using EDS (Sirion, FEI, USA). The chemical compositions and bonding characteristics of the samples were analyzed using FT-IR (Alpha FT-IR Spectrometer, Bruker, USA) spectroscopy, high-resolution dispersive Raman microscope (Horiba Jobin Yvon, France), and XPS (K-alpha, Thermo VG Scientific, England) with an Al−Kα line as the X-ray source. The C 1s peak at 284.8 eV was used as a reference. Nitrogen adsorption/desorption measurements were conducted using a surface area analyzer at 77 K (TriStar 3020, Micrometrics, USA). The electrolyte stability was determined by 1H NMR analysis (Agilent 400 MHz 54 mm NMR DD2, USA). The residual electrolyte in the separator was extracted by dipping the separator in DCCl3 solvent.

fragments (−CF2−) was implied by C 1s XPS spectrum (Figure S19b). We therefore introduced a poreless polyurethane (PU) separator70,73 (denoted as m-TiN + LiI + PU) to protect the Li metal anode from corrosion and side reactions involving crossover oxygen, water, and the derivatives of LiI. As a result of the incorporation of PU separator, the constant capacity of 430 mAh g−1 was preserved for 280 cycles (Figure 6b), corresponding to 2.15 times improved performance compared to the m-TiN + LiI cell. The protection of the Li metal anode was indicated by its XRD pattern (Figure 6e) that is free of the LiOH peaks after 80 cycles and the XPS spectrum (Figure S20). When the constant capacity was lowered to 300 mAh g−1, the cycle life was further extended to nearly 400 cycles (Figure S21). This series of cycling results verifies that the rationally designed carbon-free m-TiN can improve the cycle life compared to that of carbonaceous counterparts, especially when appropriate separators and RM are effectively employed.

CONCLUSIONS In summary, ordered mesoporous TiN with a main pore size of 37 nm was synthesized by BCP-assisted EISA and subsequent nitridation. During discharge, the sufficient pore sizes and wellaligned channel structure of m-TiN allow to first confine Li2O2 internally before the external crystal growth in the disk-like morphology. The nature of the carbon-free electrode in conjunction with conductive oxynitride/suboxide surface layers and large surface area of m-TiN jointly contribute to the improved cycle life of up to 100 cycles even under a fixed potential cutoff condition. Moreover, when integrated with a LiI RM and a poreless PU separator, the cycle life reaches >280 cycles under a fixed capacity of 430 mAh g−1. This study reveals the importance of the careful selection of air-cathode material and systematic cell design in warranting stable electrode− electrolyte interfaces at both sides of electrodes and thus extending the cycle life of Li-O2 batteries. METHODS Chemicals. TTIP, hydrochloric acid (HCl, 37 wt % in water), tetrahydrofuran (THF), bulk-TiN, tetraethylorthosilicate (TEOS), PVDF, N-methyl-2-pyrrolidone (NMP), lithium perchlorate (LiClO4), LiI, 1,2-dimethoxyethane (DME), and tetraethylene glycol dimethyl ether (TEGDME) were purchased from Sigma-Aldrich and used without further purification. Phenol-formaldehyde resin (resol) was synthesized following the basic polymerization process.74 PEO-b-PS was prepared by atomic transfer radical polymerization (ATRP).40 The molecular weight of PEO block was fixed to Mn = 5 kg mol−1 for all PEO-b-PSs. Preparation of OMC, m-TiO2, and m-TiN. For m-TiO2 with a pore size of 23 nm, 1.2 g of TTIP was added to 4.4 g of the polymer solution consisting of PEO-b-PS BCP (Mn = 25 kg mol−1), HCl, H2O, and THF in a mass ratio of 1:1.1:1.9:18. For m-TiO2 with a pore size of 33 nm, 1.1 g of TTIP was added to the same polymer solution but with the BCP of Mn = 33 kg mol−1. After stirring for 1 h, the solution was poured into a glass Petri dish at 50 °C with a cover. Once the solvent evaporated, the film was placed inside an oven at 100 °C to cure the BCP/inorganic hybrid film for 12 h. The as-made sample was calcined at 500 °C for 3 h under air with a heating rate of 1 °C min−1. After cooling down to room temperature, the resulting m-TiO2 was calcined again at 700 °C for 10 h under NH3 atmosphere with a heating rate of 1 °C min−1 to convert TiO2 to TiN, and the final mTiN was obtained. For preparation of OMC, we followed the procedure in a previous study.39 Instead of TTIP, 0.38 g of resol and 0.17 g of TEOS were added to 4 g of the polymeric solution consisting of PEO-b-PS BCP 1743

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(11) Yao, X.; Dong, Q.; Cheng, Q.; Wang, D. Why Do Lithium− Oxygen Batteries Fail: Parasitic Chemical Reactions and Their Synergistic Effect. Angew. Chem., Int. Ed. 2016, 55, 11344−11353. (12) Li, F.; Tang, D.-M.; Jian, Z.; Liu, D.; Golberg, D.; Yamada, A.; Zhou, H. Li-O2 Battery Based on Highly Efficient Sb-Doped Tin Oxide Supported Ru Nanoparticles. Adv. Mater. 2014, 26, 4659−4664. (13) Peng, Z.; Freunberger, S. A.; Chen, Y.; Bruce, P. G. A Reversible and Higher-Rate Li-O2 Battery. Science 2012, 337, 563−566. (14) Ottakam Thotiyl, M. M.; Freunberger, S. A.; Peng, Z.; Chen, Y.; Liu, Z.; Bruce, P. G. A Stable Cathode for the Aprotic Li-O2 Battery. Nat. Mater. 2013, 12, 1050−1056. (15) Kundu, D.; Black, R.; Berg, E. J.; Nazar, L. F. A Highly Active Nanostructured Metallic Oxide Cathode for Aprotic Li-O2 Batteries. Energy Environ. Sci. 2015, 8, 1292−1298. (16) Xie, J.; Yao, X.; Madden, I. P.; Jiang, D.-E.; Chou, L.-Y.; Tsung, C.-K.; Wang, D. Selective Deposition of Ru Nanoparticles on TiSi2 Nanonet and its Utilization for Li2O2 Formation and Decomposition. J. Am. Chem. Soc. 2014, 136, 8903−8906. (17) Li, F.; Tang, D.-M.; Chen, Y.; Golberg, D.; Kitaura, H.; Zhang, T.; Yamada, A.; Zhou, H. Ru/ITO: A Carbon-Free Cathode for Nonaqueous Li-O2 Battery. Nano Lett. 2013, 13, 4702−4707. (18) Li, F.; Tang, D.-M.; Zhang, T.; Liao, K.; He, P.; Golberg, D.; Yamada, A.; Zhou, H. Superior Performance of a Li-O2 Battery with Metallic RuO2 Hollow Spheres as the Carbon-Free Cathode. Adv. Energy Mater. 2015, 5, 1500294. (19) Kundu, D.; Black, R.; Adams, B.; Harrison, K.; Zavadil, K.; Nazar, L. F. Nanostructured Metal Carbides for Aprotic Li-O2 Batteries: New Insights into Interfacial Reactions and Cathode Stability. J. Phys. Chem. Lett. 2015, 6, 2252−2258. (20) Su, D.; Dou, S.; Wang, G. Single Crystalline Co 3 O 4 Nanocrystals Exposed with Different Crystal Planes for Li-O2 Batteries. Sci. Rep. 2014, 4, 5767. (21) Lee, H.; Kim, Y.-J.; Lee, D. J.; Song, J.; Lee, Y. M.; Kim, H.-T.; Park, J.-K. Directly Grown Co3O4 Nanowire Arrays on Ni-foam: Structural Effects of Carbon-Free and Binder-Free Cathodes for Lithium-Oxygen Batteries. J. Mater. Chem. A 2014, 2, 11891−11898. (22) Zhang, Y.; Zhang, H.; Li, J.; Wang, M.; Nie, H.; Zhang, F. The Use of Mixed Carbon Materials with Improved Oxygen Transport in a Lithium-Air Battery. J. Power Sources 2013, 240, 390−396. (23) Xie, J.; Yao, X.; Cheng, Q.; Madden, I. P.; Dornath, P.; Chang, C.-C.; Fan, W.; Wang, D. Three Dimensionally Ordered Mesoporous Carbon as a Stable, High-Performance Li-O2 Battery Cathode. Angew. Chem., Int. Ed. 2015, 54, 4299−4303. (24) Kushima, A.; Koido, T.; Fujiwara, Y.; Kuriyama, N.; Kusumi, N.; Li, J. Charging/Discharging Nanomorphology Asymmetry and RateDependent Capacity Degradation in Li-Oxygen Battery. Nano Lett. 2015, 15, 8260−8265. (25) Ye, Y.; Jo, C.; Jeong, I.; Lee, J. Functional Mesoporous Materials for Energy Applications: Solar Cells, Fuel Cells, and Batteries. Nanoscale 2013, 5, 4584−4605. (26) Park, J.-B.; Lee, J.; Yoon, C. S.; Sun, Y.-K. Ordered Mesoporous Carbon Electrodes for Li-O2 Batteries. ACS Appl. Mater. Interfaces 2013, 5, 13426−13431. (27) Park, J.; Jeong, J.; Lee, S.; Jo, C.; Lee, J. Effect of Mesoporous Structured Cathode Materials on Charging Potentials and Rate Capability of Lithium-Oxygen Batteries. ChemSusChem 2015, 8, 3146−3152. (28) Avasarala, B.; Haldar, P. Electrochemical Oxidation Behavior of Titanium Nitride based Electrocatalysts under PEM Fuel Cell Conditions. Electrochim. Acta 2010, 55, 9024−9034. (29) Thotiyl, M. M. O.; Sampath, S. Electrochemical Exidation of Ethanol in Acid Media on Titanium Nitride Supported Fuel Cell Catalysts. Electrochim. Acta 2011, 56, 3549−3554. (30) Molinari, V.; Giordano, C.; Antonietti, M.; Esposito, D. Titanium Nitride-Nickel Nanocomposite as Heterogeneous Catalyst for the Hydrogenolysis of Aryl Ethers. J. Am. Chem. Soc. 2014, 136, 1758−1761.

ASSOCIATED CONTENT S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsnano.6b07635. Additional SEM, TEM, EDS, BET, XRD, XPS, Raman, and electrochemical results of the electrode materials (PDF)

AUTHOR INFORMATION Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. ORCID

Jang Wook Choi: 0000-0001-8783-0901 Author Contributions ∥

These authors contributed equally to this work.

Notes

The authors declare no competing financial interest.

ACKNOWLEDGMENTS J.L. acknowledges the support by Basic Research Laboratory program and Korea Science and Engineering Foundation (KOSEF) grant funded by the Korean Government (MEST) (NRF-2016R1A4A1010735 and NRF-2015M1A2A2056557). J.W.C. acknowledges the support by the Energy Efficiency and Resources Core Technology Program of the Korea Institute of Energy Technology Evaluation and Planning (KETEP), which is granted financial resources from the Ministry of Trade, Industry, and Energy, Republic of Korea (20152020104870). REFERENCES (1) Bruce, P. G.; Freunberger, S. A.; Hardwick, L. J.; Tarascon, J.-M. Li-O2 and Li-S Batteries with High Energy Storage. Nat. Mater. 2012, 11, 19−29. (2) Abraham, K. M.; Jiang, Z. A Polymer Electrolyte-Based Rechargeable Lithium/Oxygen Battery. J. Electrochem. Soc. 1996, 143, 1−5. (3) Choi, J. W.; Aurbach, D. Promise and Reality of Post-Lithium-Ion Batteries with High Energy Densities. Nat. Rev. Mater. 2016, 1, 16013. (4) Grande, L.; Paillard, E.; Hassoun, J.; Park, J.-B.; Lee, Y.-J.; Sun, Y.K.; Passerini, S.; Scrosati, B. The Lithium/Air Battery: Still an Emerging System or a Practical Reality? Adv. Mater. 2015, 27, 784− 800. (5) Li, F.; Zhang, T.; Zhou, H. Challenges of Non-Aqueous Li-O2 Batteries: Electrolytes, Catalysts, and Anodes. Energy Environ. Sci. 2013, 6, 1125−1141. (6) Lu, Y.-C.; Gallant, B. M.; Kwabi, D. G.; Harding, J. R.; Mitchell, R. R.; Whittingham, M. S.; Shao-Horn, Y. Lithium-Oxygen Batteries: Bridging Mechanistic Understanding and Battery Performance. Energy Environ. Sci. 2013, 6, 750−768. (7) Ottakam Thotiyl, M. M.; Freunberger, S. A.; Peng, Z.; Bruce, P. G. The Carbon Electrode in Nonaqueous Li-O2 Cells. J. Am. Chem. Soc. 2013, 135, 494−500. (8) McCloskey, B. D.; Speidel, A.; Scheffler, R.; Miller, D. C.; Viswanathan, V.; Hummelshøj, J. S.; Nørskov, J. K.; Luntz, A. C. Twin Problems of Interfacial Carbonate Formation in Nonaqueous Li-O2 Batteries. J. Phys. Chem. Lett. 2012, 3, 997−1001. (9) Freunberger, S. A.; Chen, Y.; Drewett, N. E.; Hardwick, L. J.; Bardé, F.; Bruce, P. G. The Lithium-Oxygen Battery with Ether-Based Electrolytes. Angew. Chem., Int. Ed. 2011, 50, 8609−8613. (10) Kim, B. G.; Kim, S.; Lee, H.; Choi, J. W. Wisdom from the Human Eye: A Synthetic Melanin Radical Scavenger for Improved Cycle Life of Li-O2 Battery. Chem. Mater. 2014, 26, 4757−4764. 1744

DOI: 10.1021/acsnano.6b07635 ACS Nano 2017, 11, 1736−1746

Article

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(50) Xu, C.; Gallant, B. M.; Wunderlich, P. U.; Lohmann, T.; Greer, J. R. Three-Dimensional Au Microlattices as Positive Electrodes for LiO2 Batteries. ACS Nano 2015, 9, 5876−5883. (51) Liao, K.; Wang, X.; Sun, Y.; Tang, D.; Han, M.; He, P.; Jiang, X.; Zhang, T.; Zhou, H. An Oxygen Cathode with Stable Full DischargeCharge Capability based on 2D Conducting Oxide. Energy Environ. Sci. 2015, 8, 1992−1997. (52) Shang, C.; Dong, S.; Hu, P.; Guan, J.; Xiao, D.; Chen, X.; Zhang, L.; Gu, L.; Cui, G.; Chen, L. Compatible Interface Design of CoObased Li-O2 Battery Cathodes with Long-Cycling Stability. Sci. Rep. 2015, 5, 8335. (53) Viswanathan, V.; Pickrahn, K. L.; Luntz, A. C.; Bent, S. F.; Nørskov, J. K. Nanoscale Limitations in Metal Oxide Electrocatalysts for Oxygen Evolution. Nano Lett. 2014, 14, 5853−5857. (54) Xu, J.-J.; Wang, Z.-L.; Xu, D.; Zhang, L.-L.; Zhang, X.-B. Tailoring Deposition and Morphology of Discharge Products towards High-Rate and Long-Life Lithium-Oxygen Batteries. Nat. Commun. 2013, 4, 2438. (55) Zhang, P.; Sun, D.; He, M.; Lang, J.; Xu, S.; Yan, X. Synthesis of Porous δ-MnO2 Submicron Tubes as Highly Efficient Electrocatalyst for Rechargeable Li-O2 Batteries. ChemSusChem 2015, 8, 1972−1979. (56) Xu, S.-M.; Zhu, Q.-C.; Long, J.; Wang, H.-H.; Xie, X.-F.; Wang, K.-X.; Chen, J.-S. Low-Overpotential Li-O2 Batteries Based on TFSI Intercalated Co-Ti Layered Double Oxides. Adv. Funct. Mater. 2016, 26, 1365−1374. (57) Huang, L.; Mao, Y.; Wang, G.; Xia, X.; Xie, J.; Zhang, S.; Du, G.; Cao, G.; Zhao, X. Ru-decorated Knitted Co3O4 Nanowires as a Robust Carbon/Binder-Free Catalytic Cathode for Lithium-Oxygen Batteries. New J. Chem. 2016, 40, 6812−6818. (58) Liao, K.; Zhang, T.; Wang, Y.; Li, F.; Jian, Z.; Yu, H.; Zhou, H. Nanoporous Ru as a Carbon- and Binder-Free Cathode for Li-O2 Batteries. ChemSusChem 2015, 8, 1429−1434. (59) Guo, Z.; Zhou, D.; Dong, X.; Qiu, Z.; Wang, Y.; Xia, Y. Ordered Hierarchical Mesoporous/Macroporous Carbon: A High-Performance Catalyst for Rechargeable Li-O2 Batteries. Adv. Mater. 2013, 25, 5668− 5672. (60) Zhu, Y. G.; Jia, C.; Yang, J.; Pan, F.; Huang, Q.; Wang, Q. Dual Redox Catalysts for Oxygen Reduction and Evolution Reactions: Towards a Redox Flow Li-O2 Battery. Chem. Commun. 2015, 51, 9451−9454. (61) Kim, B. G.; Lee, J.-N.; Lee, D. J.; Park, J.-K.; Choi, J. W. Robust Cycling of Li-O2 Batteries through the Synergistic Effect of Blended Electrolytes. ChemSusChem 2013, 6, 443−448. (62) Yang, C.; Wong, R. A.; Hong, M.; Yamanaka, K.; Ohta, T.; Byon, H. R. Unexpected Li2O2 Film Growth on Carbon Nanotube Electrodes with CeO2 Nanoparticles in Li-O2 Batteries. Nano Lett. 2016, 16, 2969−2974. (63) Yilmaz, E.; Yogi, C.; Yamanaka, K.; Ohta, T.; Byon, H. R. Promoting Formation of Noncrystalline Li2O2 in the Li-O2 Battery with RuO2 Nanoparticles. Nano Lett. 2013, 13, 4679−4684. (64) Black, R.; Oh, S. H.; Lee, J.-H.; Yim, T.; Adams, B.; Nazar, L. F. Screening for Superoxide Reactivity in Li-O2 Batteries: Effect on Li2O2/LiOH Crystallization. J. Am. Chem. Soc. 2012, 134, 2902−2905. (65) Assary, R. S.; Lu, J.; Du, P.; Luo, X.; Zhang, X.; Ren, Y.; Curtiss, L. A.; Amine, K. The Effect of Oxygen Crossover on the Anode of a Li-O 2 Battery using an Ether-Based Solvent: Insights from Experimental and Computational Studies. ChemSusChem 2013, 6, 51−55. (66) Lim, H.-D.; Song, H.; Kim, J.; Gwon, H.; Bae, Y.; Park, K.-Y.; Hong, J.; Kim, H.; Kim, T.; Kim, Y. H.; Lepró, X.; Ovalle-Robles, R.; Baughman, R. H.; Kang, K. Superior Rechargeability and Efficiency of Lithium-Oxygen Batteries: Hierarchical Air Electrode Architecture Combined with a Soluble Catalyst. Angew. Chem., Int. Ed. 2014, 53, 3926−3931. (67) Liu, T.; Leskes, M.; Yu, W.; Moore, A. J.; Zhou, L.; Bayley, P. M.; Kim, G.; Grey, C. P. Cycling Li-O2 Batteries via LiOH Formation and Decomposition. Science 2015, 350, 530−533. (68) Kwak, W.-J.; Hirshberg, D.; Sharon, D.; Shin, H.-J.; Afri, M.; Park, J.-B.; Garsuch, A.; Chesneau, F. F.; Frimer, A. A.; Aurbach, D.;

(31) Li, J.; Gao, L.; Sun, J.; Zhang, Q.; Guo, J.; Yan, D. Synthesis of Nanocrystalline Titanium Nitride Powders by Direct Nitridation of Titanium Oxide. J. Am. Ceram. Soc. 2001, 84, 3045−3047. (32) Peelamedu, R. D.; Fleming, M.; Agrawal, D. K.; Roy, R. Preparation of Titanium Nitride: Microwave-Induced Carbothermal Reaction of Titanium Dioxide. J. Am. Ceram. Soc. 2002, 85, 117−122. (33) Hsueh, H.-S.; Yang, C.-T.; Zink, J. I.; Huang, M. H. Formation of Titanium Nitride Nanoparticles within Mesoporous Silica SBA-15. J. Phys. Chem. B 2005, 109, 4404−4409. (34) Jun, Y.-S.; Hong, W. H.; Antonietti, M.; Thomas, A. Mesoporous, 2D Hexagonal Carbon Nitride and Titanium Nitride/ Carbon Composites. Adv. Mater. 2009, 21, 4270−4274. (35) Ramasamy, E.; Jo, C.; Anthonysamy, A.; Jeong, I.; Kim, J. K.; Lee, J. Soft-Template Simple Synthesis of Ordered Mesoporous Titanium Nitride-Carbon Nanocomposite for High Performance DyeSensitized Solar Cell Counter Electrodes. Chem. Mater. 2012, 24, 1575−1582. (36) Chang, Y.; Dong, S.; Ju, Y.; Xiao, D.; Zhou, X.; Zhang, L.; Chen, X.; Shang, C.; Gu, L.; Peng, Z.; Cui, G. A Carbon- and Binder-free Nanostructured Cathode for High-Performance Nonaqueous Li-O2 Battery. Adv. Sci. 2015, 2, 1500092. (37) Ioroi, T.; Siroma, Z.; Fujiwara, N.; Yamazaki, S.-i.; Yasuda, K. Sub-stoichiometric Titanium Oxide-Supported Platinum Electrocatalyst for Polymer Electrolyte Fuel Cells. Electrochem. Commun. 2005, 7, 183−188. (38) Adams, B. D.; Black, R.; Radtke, C.; Williams, Z.; Mehdi, B. L.; Browning, N. D.; Nazar, L. F. The Importance of Nanometric Passivating Films on Cathodes for Li-Air Batteries. ACS Nano 2014, 8, 12483−12493. (39) Jo, C.; Park, Y.; Jeong, J.; Lee, K. T.; Lee, J. Structural Effect on Electrochemical Performance of Ordered Porous Carbon Electrodes for Na-Ion Batteries. ACS Appl. Mater. Interfaces 2015, 7, 11748− 11754. (40) Hwang, J.; Jo, C.; Hur, K.; Lim, J.; Kim, S.; Lee, J. Direct Access to Hierarchically Porous Inorganic Oxide Materials with ThreeDimensionally Interconnected Networks. J. Am. Chem. Soc. 2014, 136, 16066−16072. (41) Jo, C.; Kim, Y.; Hwang, J.; Shim, J.; Chun, J.; Lee, J. Block Copolymer Directed Ordered Mesostructured TiNb2O7 Multimetallic Oxide Constructed of Nanocrystals as High Power Li-Ion Battery Anodes. Chem. Mater. 2014, 26, 3508−3514. (42) Krawitz, A. D. Introduction to Diffraction in Materials Science and Engineering. Wiley-VCH: New York, 2001; pp 424. (43) Meng, Y.; Gu, D.; Zhang, F.; Shi, Y.; Yang, H.; Li, Z.; Yu, C.; Tu, B.; Zhao, D. Ordered Mesoporous Polymers and Homologous Carbon Frameworks: Amphiphilic Surfactant Templating and Direct Transformation. Angew. Chem., Int. Ed. 2005, 44, 7053−7059. (44) Lee, J.; Christopher Orilall, M.; Warren, S. C.; Kamperman, M.; DiSalvo, F. J.; Wiesner, U. Direct Access to Thermally Stable and Highly Crystalline Mesoporous Transition-metal Oxides with Uniform Pores. Nat. Mater. 2008, 7, 222−228. (45) Tanev, P. T.; Pinnavaia, T. J. A Neutral Templating Route to Mesoporous Molecular Sieves. Science 1995, 267, 865−867. (46) Liu, R.; Ren, Y.; Shi, Y.; Zhang, F.; Zhang, L.; Tu, B.; Zhao, D. Controlled Synthesis of Ordered Mesoporous C-TiO2 Nanocomposites with Crystalline Titania Frameworks from Organic-InorganicAmphiphilic Coassembly. Chem. Mater. 2008, 20, 1140−1146. (47) Kim, B. G.; Kim, H.-J.; Back, S.; Nam, K. W.; Jung, Y.; Han, Y.K.; Choi, J. W. Improved Reversibility in Lithium-Oxygen Battery: Understanding Elementary Reactions and Surface Charge Engineering of Metal Alloy Catalyst. Sci. Rep. 2014, 4, 4225. (48) Yoo, Y.; Kim, B. G.; Pak, K.; Han, S. J.; Song, H.-S.; Choi, J. W.; Im, S. G. Initiated Chemical Vapor Deposition (iCVD) of Highly Cross-Linked Polymer Films for Advanced Lithium-Ion Battery Separators. ACS Appl. Mater. Interfaces 2015, 7, 18849−18855. (49) Luo, W.-B.; Gao, X.-W.; Shi, D.-Q.; Chou, S.-L.; Wang, J.-Z.; Liu, H.-K. Binder-Free and Carbon-Free 3D Porous Air Electrode for Li-O2 Batteries with High Efficiency, High Capacity, and Long Life. Small 2016, 12, 3031−3038. 1745

DOI: 10.1021/acsnano.6b07635 ACS Nano 2017, 11, 1736−1746

Article

ACS Nano Sun, Y.-K. Understanding the Behavior of Li-Oxygen Cells Containing LiI. J. Mater. Chem. A 2015, 3, 8855−8864. (69) Chen, Y.; Freunberger, S. A.; Peng, Z.; Fontaine, O.; Bruce, P. G. Charging a Li-O2 Battery Using a Redox Mediator. Nat. Chem. 2013, 5, 489−494. (70) Kim, B. G.; Kim, J.-S.; Min, J.; Lee, Y.-H.; Choi, J. H.; Jang, M. C.; Freunberger, S. A.; Choi, J. W. A Moisture- and OxygenImpermeable Separator for Aprotic Li-O2 Batteries. Adv. Funct. Mater. 2016, 26, 1747−1756. (71) Lee, D. J.; Lee, H.; Kim, Y.-J.; Park, J.-K.; Kim, H.-T. Sustainable Redox Mediation for Lithium-Oxygen Batteries by a Composite Protective Layer on the Lithium-Metal Anode. Adv. Mater. 2016, 28, 857−863. (72) Zhao, Q.; Lu, Y.; Zhu, Z.; Tao, Z.; Chen, J. Rechargeable Lithium-Iodine Batteries with Iodine/Nanoporous Carbon Cathode. Nano Lett. 2015, 15, 5982−5987. (73) Kim, J.-S.; Yoo, D.-J.; Min, J.; Shakoor, R. A.; Kahraman, R.; Choi, J. W. Poreless Separator and Electrolyte Additive for Lithium− Sulfur Batteries with High Areal Energy Densities. ChemNanoMat 2015, 1, 240−245. (74) Liu, R.; Shi, Y.; Wan, Y.; Meng, Y.; Zhang, F.; Gu, D.; Chen, Z.; Tu, B.; Zhao, D. Triconstituent Co-Assembly to Ordered Mesostructured Polymer-Silica and Carbon−Silica Nanocomposites and LargePore Mesoporous Carbons with High Surface Areas. J. Am. Chem. Soc. 2006, 128, 11652−11662. (75) Nasybulin, E.; Xu, W.; Engelhard, M. H.; Nie, Z.; Burton, S. D.; Cosimbescu, L.; Gross, M. E.; Zhang, J.-G. Effects of Electrolyte Salts on the Performance of Li-O2 Batteries. J. Phys. Chem. C 2013, 117, 2635−2645.

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