Dark-field Transmission Electron Microscopy Images Research Articles

Detection of BiGa hetero-antisites at Ga(As,Bi)/(Al,Ga)As interfaces

In this work, we show how diffraction-based chemically sensitive dark-field transmission electron microscopy (DFTEM) reveals the presence of Bi hetero-antisites (BiGa) at the interface of Ga(As,Bi)/(Al,Ga)As quantum well (QW) structures grown by molecular beam epitaxy on GaAs(001). The presence of BiGa is demonstrated by the striking appearance of “dark-lines” at the interfaces under two-beam DFTEM imaging conditions using the (002) diffraction spot. Additional analytical scanning (S)TEM procedures reveal Ga depletion and Bi accumulation at the exact position of the dark-lines, consistent with BiGa at this location. The precise location of the dark-lines agrees with the position of growth interruptions made to adjust substrate temperature and the As/Ga flux ratio and, most importantly, the realization of a Bi pre-treatment before QW growth. We believe the Bi pre-treatment may have favored formation of BiGa hetero-antisites. We validate the use of g002 DFTEM for further investigations of the intricate bismuth incorporation into the lattice and its dependence on the growth conditions. Finally, g002 DFTEM imaging is positioned as a very powerful technique for the detection of point defects in general in materials with the zinc-blende crystal structure, beyond dilute bismide alloys.

Cross-sectional transmission electron microscopy studies of Si coated heat-treated stainless steel with improved optical absorption for solar thermal applications

Heat-treatment of stainless steel substrate is reported to improve absorptance due to an increase in roughness along with the formation of surface oxides. Depositing a thin Si layer on stainless steel prior to the heat-treatment results in the formation of islands of nano-micrometer sizes with gaps between them which helps in improving absorptance. Cross-sectional Transmission Electron Microscopy studies are required to understand the development of the surface structures and improved optical properties. The cross-sectional bright- and dark-field Transmission Electron Microscopy images showed the presence of a bright Si-rich region where small nanocrystalline islands were embedded. On top of this layer, larger nanocrystalline protrusions were observed. Energy Dispersive X-ray analysis and Selected Area Diffraction data showed that the substrate elements diffuse out through the Si layer forming smaller islands of various Fe-Cr-Mn-based silicides in the Si layer and islands of Fe, Cr, and Mn-based silicides and oxides at the surface. The High-Resolution Transmission Electron Microscopy images showed that the islands are nanocrystalline in nature and are embedded in a Si-rich amorphous region. The absence of O in the Si-rich layer in SAD data was observed. Also, from the top layer EDX data, Fe peak intensity was found to be higher than Cr peak intensity. These observations indicate that the presence of the Si layer helps in reducing the inward diffusion of O into the substrate, thus preventing internal oxidation. Also, this Si layer prevents the Cr from forming a continuous Cr oxide layer and instead forms mixed oxide/silicide islands which enhance optical properties. Similar studies in the heat-treated substrate without a Si layer showed the presence of sharp-angled nanocrystalline structures of Fe-Cr-based oxides and smooth-angled nanocrystalline structures of Fe-Mn-based oxides on the surface.

Statistical Determination of Atomic-Scale Characteristics of Au Nanocrystals Based on Correlative Multiscale Transmission Electron Microscopy

Abstract Atomic-scale characteristics of individual nanocrystals (NCs), such as the crystallographic orientation of their facets and the kind and density of crystal structure defects, play a tremendous role for the functionality and performance of the whole NC population. However, these features are usually quantified only for a small number of individual particles, and thus with limited statistical relevance. In the present work, we developed the multiscale approach available in transmission electron microscopy (TEM) further, and applied it to describe features of different types of Au NCs in a statistical and scale-bridging manner. This approach combines high-resolution TEM, which is capable of describing the characteristics of NCs on atomic scale, with a semi-automatic analysis of low-magnification high-angle annular dark-field scanning TEM images, which reveals the nanoscopic morphological attributes of NCs with good statistics. The results of these complementary techniques are combined and correlated. The potential of this multiscale approach is illustrated on two examples. In the first one, the habitus of Au NCs was classified and assigned to multiply twinned nanoparticles and nanoplates. These classes were quantified and related to different stacking fault densities. The second example demonstrates the statistical determination of crystallographic orientations and configurations of facets in Au nanorods.

A modulated structure derived from the XA-type Mn2RuSn Heusler compound.

A modulated structure derived from the inverse Heusler phase (the XA-type and the disordered variant L21B-type) has been observed in rapidly quenched Mn2RuSn ribbons. The powder X-ray diffraction pattern of the quenched ribbons can be indexed as an L21B-type structure. Electron diffraction patterns of the new structure mostly resemble those of the XA-type (and the disordered variant L21B-type) structure and additional reflections with denser spacing indicate a long periodicity. Orthogonal domains of the modulated structure were revealed by a selected-area electron diffraction pattern and the corresponding dark-field transmission electron microscopy images. The structure was further studied by the crystallographic analysis of high-resolution transmission electron microscopy images. A model for the modulated structure has been proposed to interpret the experimental results.

Nanostructural Analysis of Co‐Re/γ‐Al2O3 Fischer‐Tropsch Catalyst by TEM and XRD

AbstractA rhenium promoted Fischer‐Tropsch (FT) cobalt catalyst supported on γ‐Al2O3 has been investigated by Transmission Electron Microscopy (TEM) and X‐ray Diffraction (XRD), before and after reduction. Electron diffraction, High Resolution TEM and Electron Energy Loss Spectroscopy were used to confirm the oxidation state. Cobalt aggregate, particle and crystallite sizes have been studied in detail and measured by TEM and XRD. A cobalt particle size of 10.0±2.4 nm obtained from bright field TEM images for the reduced material is consistent with the XRD analysis of the calcined catalyst. After reduction dark field TEM imaging gave a volume‐weighted crystallite size of 7.5±2.5 nm, which is close to the value obtained by XRD. The particles had lost the parallel orientation and physical continuity within the alumina pore structure that were present before reduction. The latter was confirmed by electron tomography. Lamellae identified with the presence of Hexagonal Close Packed cobalt were observed in the predominantly Face Centred Cubic particles.

Strain-Assisted Single Pt Sites on High-Curvature MoS 2 Surface for Ultrasensitive H 2 S Sensing

Open AccessCCS ChemistryRESEARCH ARTICLE7 Dec 2022Strain-Assisted Single Pt Sites on High-Curvature MoS2 Surface for Ultrasensitive H2S Sensing Zhenggang Xue, Chun Wang, Yujing Tong, Muyu Yan, Jiangwei Zhang, Xiao Han, Xun Hong, Yafei Li and Yuen Wu Zhenggang Xue NEST Laboratory, Department of Physics, Department of Chemistry, College of Science, Shanghai University, Shanghai 200444 Hefei National Laboratory for Physical Sciences at the Microscale, Collaborative Innovation Center of Chemistry for Energy Materials (iChEM), School of Chemistry and Materials Science, National Synchrotron Radiation Laboratory, University of Science and Technology of China, Hefei 230026 Google Scholar More articles by this author , Chun Wang Jiangsu Collaborative Innovation Centre of Biomedical Functional Materials, School of Chemistry and Materials Science, Nanjing Normal University, Nanjing 210023 Google Scholar More articles by this author , Yujing Tong Hefei National Laboratory for Physical Sciences at the Microscale, Collaborative Innovation Center of Chemistry for Energy Materials (iChEM), School of Chemistry and Materials Science, National Synchrotron Radiation Laboratory, University of Science and Technology of China, Hefei 230026 Google Scholar More articles by this author , Muyu Yan Hefei National Laboratory for Physical Sciences at the Microscale, Collaborative Innovation Center of Chemistry for Energy Materials (iChEM), School of Chemistry and Materials Science, National Synchrotron Radiation Laboratory, University of Science and Technology of China, Hefei 230026 Google Scholar More articles by this author , Jiangwei Zhang Dalian National Laboratory for Clean Energy, State Key Laboratory of Catalysis, Dalian Institute of Chemical Physics, Chinese Academy of Sciences, Dalian 116023 Google Scholar More articles by this author , Xiao Han Hefei National Laboratory for Physical Sciences at the Microscale, Collaborative Innovation Center of Chemistry for Energy Materials (iChEM), School of Chemistry and Materials Science, National Synchrotron Radiation Laboratory, University of Science and Technology of China, Hefei 230026 Google Scholar More articles by this author , Xun Hong Hefei National Laboratory for Physical Sciences at the Microscale, Collaborative Innovation Center of Chemistry for Energy Materials (iChEM), School of Chemistry and Materials Science, National Synchrotron Radiation Laboratory, University of Science and Technology of China, Hefei 230026 Google Scholar More articles by this author , Yafei Li Jiangsu Collaborative Innovation Centre of Biomedical Functional Materials, School of Chemistry and Materials Science, Nanjing Normal University, Nanjing 210023 Google Scholar More articles by this author and Yuen Wu *Corresponding author: E-mail Address: [email protected] Hefei National Laboratory for Physical Sciences at the Microscale, Collaborative Innovation Center of Chemistry for Energy Materials (iChEM), School of Chemistry and Materials Science, National Synchrotron Radiation Laboratory, University of Science and Technology of China, Hefei 230026 Dalian National Laboratory for Clean Energy, Dalian 116023 Google Scholar More articles by this author https://doi.org/10.31635/ccschem.022.202101628 SectionsSupplemental MaterialAboutAbstractPDF ToolsAdd to favoritesDownload CitationsTrack Citations ShareFacebookTwitterLinked InEmail Engineering the local three-dimensional structure of metal sites has an important effect of maximizing the activity and selectivity of single-atom site catalysts. Here, we engineered a strain-assisted single Pt sites structure on a highly curved MoS2 surface to enhance its H2S sensor property. By introducing N-methyl-2-pyrrolidone (NMP) as guiding molecules, a multilayer MoS2 structure with bending base planes was achieved. This bending behavior could inject not only uniform in-plane strain into the original inert MoS2 basal plane but also introduce sufficient accessible sites to anchor Pt monomers. Further experimental and theoretical results showed that the high-curvature MoS2 surface endowed 0.8% stretch strain onto the low-coordinated single Pt sites with a unique “tip” effect, which led to more accumulation of electrons around the Pt species, thereby accelerating the electric transfer process between H2S and supports. The final catalyst delivered pronouncedly enhanced H2S sensing response and response speed at room temperature. Our proposed strain-assisted strategy might create a new path to design highly active single-atom site catalysts for gas sensors. Download figure Download PowerPoint Introduction One frontier in the synthesis of single-atom site catalysts includes locating highly distributed metal atoms onto the desired positions to form a site-special geometric framework so as to improve the intrinsic activity of metal sites.1–6 For example, in the case of transition metal dichalcogenides (TMDs) materials comprised active edges and inert basal plane, selectively modifying the single metal atoms at the edge or in-plane sites is crucial for their catalytic performance due to differences in coordinated environment and steric configuration of metal sites. Many attempts have proven that metal atoms around the coordinatively unsaturated edge regions possess unique geometric and electronic configurations, and thus, exhibit superior activity during the catalytic process.7,8 However, these edge regions usually only occupy a little part over the whole TMDs supports, which severely restrict the concentrations and density of metal sites. One effective strategy used to solve this problem is to sufficiently activate the inert basal plane of TMDs materials as active regions to anchor the metal species.9,10 Some physical or chemical methods have been reported to fully utilize inert basal atoms, such as constructing vacancies,11,12 manipulating strain,13,14 creating distortion,15 inducting doped atoms,16–19 and so on. Among these, injecting lattice strain into the TMDs planes could create the active surface and induce the strain-assisted single metal sites, which might be beneficial to improve the catalytic activity. However, the strain-related single-atom site catalysts supported on active TMDs surfaces have rarely been studied. Here, we present a strain-assisted single Pt site on a highly curved MoS2 surface to enhance the H2S sensor property. In our design, the N-methyl-2-pyrrolidone (NMP) was employed as guiding solvent molecules to construct multilayer fullerene-like MoS2 configurations. This closed structure introduced uniform in-plane strain and effectively prevented strain recovery of MoS2 layers. The X-ray absorption fine structure (XAFS) characteristics and density functional theory (DFT) calculation further revealed the high-curvature MoS2 surface endowed 0.8% stretch strain onto the low-coordinated single Pt sites. The final Pt decorated bending MoS2 (denoted as Pt-B-MoS2) enhanced the sensing response of H2S gas markedly and reduced the response/recovery time at room temperature. Further calculations and simulations indicated that the special Pt sites could enhance the H2S adsorption and showed a unique tip effect to induce more accumulation of electrons, which further accelerated an electric transfer between H2S and supports. Experimental Methods Materials Ammonium tetrathiomolybdate [(NH4)2MoS4], NMP, ethanol, hydrazine hydrate (N2H4·H2O) were purchased from Shanghai Chemical Reagents (Shanghai, China). Chloroplatinic acid (H2PtCl6·6H2O) was obtained from Sigma-Aldrich (Shanghai, China). Deionized water was used throughout this study. All chemicals were used as received without further purification. Preparation of S-MoS2 and B-MoS2 The B-MoS2 was prepared by a facile solvothermal method: Typically, 0.5 mmol of ammonium tetrathiomolybdate was mixed with 80 mL of NMP solution and stirred for 0.5 h. Then 2 mL N2H4·H2O was added to the solution and heated to 80 °C in an oil bath under gentle reflux. After stirring for 12 h, the black-brown products were harvested and washed three times with water, then dried at 60 °C overnight. The as-B-MoS2 precursors obtained were heated to 800 °C at a rate of 5 °C/min and kept for 2 h under a flowing Ar2 atmosphere. The S-MoS2 was prepared by the same process except that H2O replaced the NMP solution as a corresponding solvent. Preparation of Pt-B-MoS2 and Pt-S-MoS2 In a typical procedure, 150 mg B-MoS2 (or S-MoS2) powder was dispersed in 20 mL water and ultrasonic mixed for 10 min. After stirring for 1 h, 0.375 mL of H2PtCl6 aqueous solution (Pt: 4 mg/mL) was added to the solution mixture, followed by stirring for 4 h. Then the as-formed precipitates were washed three times with water and dried at 60 °C under a vacuum. Characterization Powder X-ray diffraction (XRD) measurements were recorded on a Rigaku Miniflex-600 (Tokyo, Japan) operated at 40 kV voltage and 15 mA current using a Cu Kα radiation (λ = 0.15406 nm) at a step width of 8°/min. Scanning electron microscopy (SEM) was performed on JSM-6700F (JEOL, Tokyo, Japan). Transmission electron microscope (TEM) images were recorded on a Hitachi-7700 (Tokyo, Japan) set at 100 kV. The high-resolution TEM and high-angle annular dark-field scanning TEM (HAADF-STEM) images were recorded on an FEI Tecnai G2 F20 S-Twin high-resolution transmission electron microscope (HRTEM; Hillsboro, OR, United States) set at 200 kV, and a JEOL JEM-ARM200F TEM/STEM (Tokyo, Japan), respectively, with a spherical aberration corrector worked at 300 kV. Through-focal HAADF series were acquired at nanometer intervals, with the first image under-focused (beyond the beam exit surface) and the final image over-focused (before the beam entrance surface). Then the images were aligned manually to remove the sample drift effects. X-ray photoelectron spectroscopy (XPS) experiments were performed at the Catalysis and Surface Science End station at the BL11U beamline of National Synchrotron Radiation Laboratory (NSRL) in Hefei, China. Soft X-ray absorption spectroscopy (Soft-XAS, S L-edge) was carried out at BL12B X-ray Magnetic Circular Dichroism (XMCD) station and BL10B photoemission end-station of National Synchrotron Radiation Laboratory (NSRL, Hefei in China) in total electron yield (TEY) mode. The samples were coated on double-sided carbon tape for characterization. UV–vis absorption spectra were collected by an Agilent Cary 60 spectrophotometer (Santa Clara, CA, United States). Room-temperature electron paramagnetic resonance (EPR) spectra were obtained using a JEOL JES-FA200 spectrometer (300 K, 9.062 GHz; Tokyo, Japan). XAFS measurement and data analysis: The X-ray absorption fine structure data were collected at 1W1B station in Beijing Synchrotron Radiation Facility (BSRF). The storage rings of BSRF was operated at 2.5 GeV with a maximum current of 250 mA. The X-ray absorption near edge structure (XANES) data were recorded in fluorescence mode. All samples were pelletized as disks of 13 mm diameter using poly(1,1-difluoroethylene) powder as a binder. The acquired extended X-ray absorption fine structure (EXAFS) data were processed according to the standard procedures using the ATHENA module implemented in the IFEFFIT software packages.20 The EXAFS spectra were obtained by subtracting the post-edge background from the overall absorption and then normalizing with respect to the edge-jump step. Subsequently, χ(k) data in the k-space were Fourier transformed to real (R) space using a Hanning window (dk = 1.0 Å−1) to separate the EXAFS contributions from different coordination shells. Fabrication and response test of the gas sensor Gas sensing tests were performed on a commercial CGS-8 Gas Sensing Measurement System (Beijing Elite Tech Company Limited), and the gas sensing performances of the sensors were measured using a data acquisition system and subsequent transformation to a response. First, the products were mixed with ethanol to form a paste and coated on the sensor substrates. After drying in the air, the substrates were welded into the circuit board. All tests were implemented at room temperature (25 °C). The H2S sensitivity of the sensor was defined as S = (Rg − Ra)/Ra = ΔR/Ra, where Ra was the baseline resistance of the sensors when exposed to air, and Rg was the resistance of the sensors when exposed to gas analytes. The sensitivity of the oxidizing gas NO2 sensor was defined as S = (Ra − Rg)/Ra = ΔR/Ra. Results and Discussion The strategies used to activate the basal plane and introduce active Pt sites in MoS2 are illustrated in Figure 1a. The straight MoS2 (denoted as S-MoS2) nanosheet with normal crystal arrangement was synthesized by facilely reducing the (NH4)2MoS4 precursors through N2H4 reagent in the water solution at 80 °C (see Experimental Procedures). The TEM and HRTEM images ( Supporting Information Figure S1) showed that the S-MoS2 nanosheets possessed layer-assembled morphology, and no obvious bending structure was observed. The lattice spacing of S-MoS2 was 0.66 nm, consistent with the (002) surfaces of 2H-MoS2.9 Interestingly, when the water solution was replaced by NMP reagent, the MoS2 layer was inclined to curve and close, establishing a hollow fullerene-like nanosphere with an average shell thickness of ∼7 nm ( Supporting Information Figures S2 and S3). The HAADF-STEM image in Figure 1c indicated that the bending MoS2 (denoted as B-MoS2) was highly curved at an atomic scale, and the interconnected fullerene spheres supported the high-quality MoS2 surface. XRD patterns in Supporting Information Figure S4 suggested that both S-MoS2 and B-MoS2 were typical structures of the trigonal prismatic (2H) phase (PDF no. 37-1492).21 Based on the HAADF-STEM image in Figure 1c, the detailed lattice strain components of εxx and εyy were displayed in Supporting Information Figure S5. The average strain distribution of εxx and εyy ranged from ∼−10%to 10% and 0% to 30%, respectively, indicating that the visual curvature induced the generation of noticeable strain between the atoms. More representative zones were investigated, as shown in Supporting Information Figures S6 and S7. The hollow B-MoS2 sphere exhibited uniform layer numbers of ∼11, consistent with the TEM images in Supporting Information Figure S3. The HAADF-STEM images and the corresponding simulation configurations of straight and bending MoS2 configurations in Figure 1b clearly showed the differences in atomic array and direction. Moreover, compared with S-MoS2, slight redshifts of the Raman E 2 g 1 and A1g peaks were observed in B-MoS2 (Figure 1d), suggesting that a tensile strain was introduced into the B-MoS2 sample, in accordance with previous reports.10 Furthermore, the UV–vis spectrum ( Supporting Information Figure S8a) of B-MoS2 showed an apparent redshift of A and B excitons with respect to the S-MoS2, ascribed mainly to the structural stretch.22 The Mo 3d XPS analysis in Supporting Information Figure S8b showed that the doublet peaks of Mo in B-MoS2 shifted to lower binding energy compared with Mo in S-MoS2, suggesting a reduced oxidation state of the Mo species in B-MoS2. The Mo XANES characterization is displayed in Supporting Information Figure S9a: The pre-edge centroid of B-MoS2 shifted into lower energy compared with S-MoS2, suggesting a decreased oxidation of Mo in B-MoS2. This result agreed well with the XPS analysis. Moreover, the EXAFS spectrum of B-MoS2 ( Supporting Information Figure S9b) and S-MoS2 ( Supporting Information Figure S9c) and quantitative analysis of EXAFS fitting ( Supporting Information Table S1) showed that the average Mo-S bond length in B-MoS2 was 2.43 Å, which was longer than that of S-MoS2 (2.33 Å). Figure 1 | (a) The schematic diagram for forming process of Pt-B-MoS2. (b) The HAADF-STEM images and corresponding simulating configurations of S-MoS2 and B-MoS2. (c) The HAADF-STEM image of the B-MoS2 sample. (d) Raman spectra for the S-MoS2 and B-MoS2. Inset: the HAADF-STEM image of B-MoS2. (e) The HAADF-STEM image of Pt-B-MoS2. Download figure Download PowerPoint Further, when Pt species were introduced to the curved B-MoS2 surface, the HAADF-STEM image (Figure 1e) clearly showed that the isolated Pt atoms (circled in Figure 1e) distributed uniformly over the B-MoS2 surface (denoted as Pt-B-MoS2), with no apparent agglomeration. More representative images are shown in Supporting Information Figure S10. Energy-dispersive X-ray analysis ( Supporting Information Figure S11) exhibited homogeneous distribution of Mo, S, and Pt compositions on the bending frameworks. Moreover, we explored the chemical environment of Pt species on the supports employing XPS: For pristine B-MoS2, the S 2p XPS displayed a characteristic doublet peak at ∼163.4 and ∼162.2 eV (Figure 2a). After functionalization by the Pt atoms, the doublet peaks of S 2p in Pt-B-MoS2 shifted to higher binding energy (0.3 eV left shift) compared with S in B-MoS2. Similarly, the S L-edge XANES spectra in Supporting Information Figure S12 exhibited a lower peak intensity of Pt-B-MoS2 compared with origin B-MoS2. These changes occurred mainly because of the spontaneous reduction of Pt4+ species on MoS2 surface and the formation of Pt-S coordinate bond, resulting in electrons’ transfer from S to Pt; thus, modifying the electrical environment of S atoms.8,23 In addition, the same peak shift of Mo 3d was also observed in Supporting Information Figure S13, in which Mo in Pt-B-MoS2 shifted to higher energy compared with B-MoS2.17 The Pt 4f XPS analysis in Figure 2b showed that the Pt binding energy in Pt-B-MoS2 was higher than that of the Pt nanoparticles (Pt0, at ∼71.2 eV) but lower than PtO2 (Pt4+, at ∼74.0 eV), indicating its ionic Ptδ+ (0 < δ < 4) nature. However, for the Pt loading of S-MoS2 samples (donated as Pt-S-MoS2), the Pt exhibited an additional peak at ∼74.0 eV, suggesting the existing higher oxidation state of the Pt4+ species. This was attributed mainly to the differences in local geometric configuration and electronic environment of the Pt single atoms on special supports. Moreover, the XPS results showed that the contents of surface Pt in Pt-B-MoS2 was ∼0.83 wt %, which is higher than that in Pt-S-MoS2 (0.69 wt %) and closed to the inductively coupled plasma atomic emission spectroscopy results of ∼0.91 wt % ( Supporting Information Table S2). The EPR analysis ( Supporting Information Figure S14) revealed no detection of notable sulfur vacancy peak (at g = 2.003) in the B-MoS2 samples,11,24 demonstrating negligible defect sites over the support surface for plausible capture of dissociative metal monomers. Therefore, the Pt atoms on the bending MoS2 surface were inclined to disperse and stabilize only through the strong covalent metal-support interaction,23,25 distinct from defect-stabilized Pt atoms in Pt-S-MoS2. Figure 2 | (a) The S 2P XPS of Pt-B-MoS2 and B-MoS2. (b) The Pt 4f XPS of Pt-B-MoS2 and Pt-S-MoS2. (c) Pt L-edge XANES spectra. (d) The EXAFS spectra (e) corresponding fitting curve of Pt L3 edge. (f) The diffusion path of Pt atoms on B-MoS2 surface and the analysis of diffusion barrier. Download figure Download PowerPoint Next, the Pt L-edge XANES characterization (Figure 2c) revealed the oxidation states of Pt in Pt-B-MoS2 and Pt-S-MoS2, in which the white line intensity was situated between the PtS2 and Pt foil. The EXAFS spectrum (Figure 2d) displayed a notable peak at ∼2.2 Å, contributed by the Pt-S coordination path. No Pt–Pt characteristic peak (at ∼2.8 Å) was detected for both Pt-B-MoS2 and Pt-S-MoS2 samples, confirming the atomic dispersion of Pt on the supports. Interestingly, the intensity of Pt in Pt-B-MoS2 was weaker than that in Pt-S-MoS2, suggesting a lower coordinated number (CN) of Pt atoms in Pt-B-MoS2. Furthermore, the first-shell EXAFS fitting curve (Figure 2e) and k space fitting curve obtained ( Supporting Information Figure S15) validated that the central Pt in Pt-B-MoS2 coordinated directly with three S atoms, with an average bond length of 2.34 Å ( Supporting Information Table S3), which were 0.8% longer than that of that of Pt-S-MoS2. In contrast, the Pt in Pt-S-MoS2 showed a higher CN of ∼6.0 ( Supporting Information Figure S16 and Table S3), which might have led to an increase in valence states of Pt, in accordance with the XPS results. The best-fitting results of PtS2 nanoparticles and the corresponding Pt L-edge wavelet transform contour plots are displayed in Supporting Information Figures S17 and S18 and Table S3, respectively. Only one maximum peak at ∼5.9 Å−1 was discovered in all Pt-B-MoS2, Pt-S-MoS2 and PtS2 samples, regarded as the Pt-S coordination path. No obvious Pt–Pt bond was observed, demonstrating highly distributed single Pt atoms in Pt-B-MoS2 and Pt-S-MoS2 catalysts. To further corroborate the Pt-sites configuration on the special support, theoretical investigations were implemented employing DFT methods. As seen in Figure 2f, the curved MoS2 framework was constructed and the Pt species were introduced into the special surface. We found that the adsorbed Pt atom could stabilize on the top of the Mo atom bound to three adjacent S atoms, consistent with the EXAFS results. Meanwhile, the diffusion of Pt atom along with the bending surface required additional energy barriers of 0.8 eV to the adjacent hollow site and 0.7 eV to the top of S atom, implying excellent structural stability for this framework. Moreover, when we introduced S-vacancy or Mo-vacancy into this curved MoS2 structure, it was extremely unstable, different from the S-vacancy or Mo-vacancy stablized single atoms in straight MoS2 supports. The three-dimensional topographic maps of the proposed Pt-B-MoS2 and Pt-S-MoS2 theoretical models are shown in Supporting Information Figure S19. Notably, the Pt atom in Pt-B-MoS2 protruded from the MoS2 plane, and the heights of adjacent in-plane S atoms were inconsistent due to the bending strain. In order to examine the catalytic activity of as-obtained S-MoS2, B-MoS2, Pt-S-MoS2, and Pt-B-MoS2 samples, gas-sensing performance tests were conducted. For an intact sensing process, when the reducing gases arrived at the p-type MoS2 surface, the supports trapped electrons from the adsorbed gas molecules, leading to a decrease in hole concentration and an increase in resistance. Once the sensors were exposed in the air atmosphere again, the desorption behavior of the gases molecules resulted in resistance recovery. The introduction of a noble metal such as Pt served as the chemical and electrical sensitizer to promote the activation of adsorbed species and facilitate the electron transmission process. First, the Pt-B-MoS2 sensors were circularly exposed to H2S atmosphere ranged from 0.5 to 10 ppm at room temperature. As seen in Figure 3a, on the basis of various H2S concentrations, the Pt-B-MoS2 sensor exhibited cyclic and stable dynamic response and recovery characteristics. Also, the baseline of the Pt-B-MoS2 sensors almost recovered in the air atmosphere, in comparison with poor reversibility of B-MoS2 ( Supporting Information Figure S20). Moreover, the responses (defined as S = (Rg − Ra)/Ra = ΔR/Ra) for all four sensors increased continuously with increasing H2S concentrations (Figure 3b). In detail, the B-MoS2 sensor exhibited an enhanced H2S response compared with S-MoS2, suggesting that introducing the bending stain enabled the active MoS2 surface to strengthen the adsorption of H2S gas molecules.26 In particular, the Pt-B-MoS2 sensor (205% to 10 ppm H2S) exhibited ∼4-, 11- and 21-fold higher responses than those of Pt-S-MoS2 (52%), B-MoS2 (18%), and S-MoS2 (10%), respectively. This was ascribed mainly to the highly active Pt atoms on the activated MoS2 supports; both improved the gas adsorption and promoted the activation of adsorbed H2S species and the subsequent electron transport process. Considering the permissible occupational limitation for H2S in air environment is 10 ppm,27 such a high sensing response surpassed those of most reported H2S sensors ( Supporting Information Figure S21a and Table S4) and met the detection requirements. More gas-sensing performances based on the different MoS2 morphology are displayed in Supporting Information Table S6. Furthermore, the Pt-B-MoS2 sensor exhibited a significant response value of 18% even at an extremely low H2S concentration of 100 ppb. However, the limit of detection for Pt-S-MoS2 sensor toward H2S is only 1 ppm, while the sensing signals could not be perceived for S-MoS2 and B-MoS2 sensors once the concentration of H2S was below 5 ppm. The response and recovery times of the four samples at varying H2S concentrations were further investigated (Figures 3c and 3d): The Pt-B-MoS2 sensors showed a much faster response and recovery times (25 and 20 s) at 10 ppm H2S than those of Pt-S-MoS2 (50 and 55 s), B-MoS2 (1224 and 890 s), and S-MoS2 (1684 and 720 s). This significant increase in response and recovery speed confirmed that the site-specific Pt configuration facilitated the adsorption of the target gas and electron transport process; thus, enhancing the intrinsic activity of catalysts. In addition, the responses of S-MoS2, B-MoS2, Pt-S-MoS2, and Pt-B-MoS2 sensors toward different interfering gases at 10 ppm concentration were further examined (Figure 3e): The response of Pt-B-MoS2 sensor for H2S was at least ten times higher than other analyte gases, including nitrogen monoxide (NO), ammonia (NH3), carbon monoxide (CO), methanal (HCHO), hydrogen (H2), and nitrogen dioxide (NO2), indicating a significantly high selectivity. This sufficiently corroborated that the well-designed Pt sites with coordination-consistent configuration provided favorable binding sites to adsorb H2S molecule selectively and optimize the gas-sensing behavior. The cycle stability tests are displayed in Figure 3f, in which the response, response time (Tres), and recovery time (Trec) of Pt-B-MoS2 sensors remained almost unchanged during 19 cycles. The long-term durability measurement was performed in Supporting Information Figure S21b, suggesting a superior stability of Pt-B-MoS2 sensors. Figure 3 | (a) Time-related dynamic responses of Pt-B-MoS2. (b) The responses of S-MoS2, B-MoS2, Pt-S-MoS2, and Pt-B-MoS2 in different H2S concentrations. The (c) response and (d) recovery time of four samples at various H2S concentrations. (e) The selectivity of S-MoS2, B-MoS2, Pt-S-MoS2, and Pt-B-MoS2 in different gases. (f) The cycle stability tests of Pt-B-MoS2 sensors. Download figure Download PowerPoint Finally, DFT calculations were conducted to investigate the structure-activity relationships between special-sites stabilized Pt-B-MoS2 catalysts and gas-sensing performance. Initially, four different models of S-MoS2, B-MoS2, Pt-S-MoS2, and Pt-B-MoS2 surfaces were built based on the above experimental results (Figure 4a). In comparison, the adsorption energy Eads of B-MoS2 (−0.26 eV) was larger than that of S-MoS2 (−0.21 eV) upon interaction with H2S molecules, which demonstrated that the introduced stain by bending structure could effectively activate the MoS2 supports and promote gas adsorption, agreeing well with previous report. Furthermore, the gas adsorption was elevated further once the Pt atom was introduced into the B-MoS2 or S-MoS2 supports. Remarkably, the Pt-B-MoS2 structure exhibited the strongest adsorption ability (Eads = −1.63 eV) than other models, implying that the Pt-B-MoS2 surface was beneficial for adsorption and interacted with the H2S molecules. The electric density distribution of the Pt-B-MoS2 framework was further investigated: As displayed in Figure 4b, the extrusive Pt atom aggregated more electrons over the bending surface, and thus, revealed a tip-like center around the Pt species. The highly localized electric configuration was proven earlier to serve as active sites for assembling high-concentration reactants and facilitate the reaction process such as electrocatalytic CO2 reduction,28,29 electrocatalytic hydrogen evolution,2 and others. Therefore, the tip effect induced by curve-supported Pt sites improved the adsorption and interactions considerably with H2S molecules. To explore steeply into the interactions between supports and target H2S molecules during reaction process, the partial density of states (PDOS) analysis of H2S/Pt-S-MoS2 and H2S/Pt-B-MoS2 models were investigated. The outermost atomic orbits of Pt-5d, S-3p, and H-1s are displayed in Figure 4c. Compared with H2S/Pt-S-MoS2 model, we found that the Pt-5d orbitals and S-3p orbitals in H2S/Pt-B-MoS2 model possessed a noticeable overlap below the Fermi level, indicating a strong hybridization between these two orbitals. Therefore, the stain-assisted Pt sites might have acted as a crucial role of reactant (H2S) dragging, and thus, enhanced the electric transfer between adsorbed H2S and Pt sites. Figure 4d further depicts intuitional insigh

Oxide Two-Dimensional Electron Gas with High Mobility at Room-Temperature.

The prospect of 2‐dimensional electron gases (2DEGs) possessing high mobility at room temperature in wide‐bandgap perovskite stannates is enticing for oxide electronics, particularly to realize transparent and high‐electron mobility transistors. Nonetheless only a small number of studies to date report 2DEGs in BaSnO3‐based heterostructures. Here, 2DEG formation at the LaScO3/BaSnO3 (LSO/BSO) interface with a room‐temperature mobility of 60 cm2 V−1 s−1 at a carrier concentration of 1.7 × 1013 cm–2 is reported. This is an order of magnitude higher mobility at room temperature than achieved in SrTiO3‐based 2DEGs. This is achieved by combining a thick BSO buffer layer with an ex situ high‐temperature treatment, which not only reduces the dislocation density but also produces a SnO2‐terminated atomically flat surface, followed by the growth of an overlying BSO/LSO interface. Using weak beam dark‐field transmission electron microscopy imaging and in‐line electron holography technique, a reduction of the threading dislocation density is revealed, and direct evidence for the spatial confinement of a 2DEG at the BSO/LSO interface is provided. This work opens a new pathway to explore the exciting physics of stannate‐based 2DEGs at application‐relevant temperatures for oxide nanoelectronics.

Boosting the Rate Performance of Li–S Batteries via Highly Dispersed Cobalt Nanoparticles Embedded into Nitrogen-Doped Hierarchical Porous Carbon

Open AccessCCS ChemistryRESEARCH ARTICLE5 Aug 2022Boosting the Rate Performance of Li–S Batteries via Highly Dispersed Cobalt Nanoparticles Embedded into Nitrogen-Doped Hierarchical Porous Carbon Cheng Yuan, Pan Zeng, Chen Cheng, Tianran Yan, Genlin Liu, Wenmin Wang, Jun Hu, Xin Li, Junfa Zhu and Liang Zhang Cheng Yuan Institute of Functional Nano & Soft Materials (FUNSOM), Jiangsu Key Laboratory for Carbon-Based Functional Materials & Devices, Soochow University, Suzhou 215123 Google Scholar More articles by this author , Pan Zeng Institute of Functional Nano & Soft Materials (FUNSOM), Jiangsu Key Laboratory for Carbon-Based Functional Materials & Devices, Soochow University, Suzhou 215123 Google Scholar More articles by this author , Chen Cheng Institute of Functional Nano & Soft Materials (FUNSOM), Jiangsu Key Laboratory for Carbon-Based Functional Materials & Devices, Soochow University, Suzhou 215123 Google Scholar More articles by this author , Tianran Yan Institute of Functional Nano & Soft Materials (FUNSOM), Jiangsu Key Laboratory for Carbon-Based Functional Materials & Devices, Soochow University, Suzhou 215123 Google Scholar More articles by this author , Genlin Liu Institute of Functional Nano & Soft Materials (FUNSOM), Jiangsu Key Laboratory for Carbon-Based Functional Materials & Devices, Soochow University, Suzhou 215123 Google Scholar More articles by this author , Wenmin Wang *Corresponding authors: E-mail Address: [email protected] E-mail Address: [email protected]u.cn Hefei National Laboratory for Physical Sciences at the Microscale, University of Science and Technology of China, Hefei 230026 Google Scholar More articles by this author , Jun Hu National Synchrotron Radiation Laboratory, University of Science and Technology of China, Hefei 230026 Google Scholar More articles by this author , Xin Li Hefei National Laboratory for Physical Sciences at the Microscale, University of Science and Technology of China, Hefei 230026 Google Scholar More articles by this author , Junfa Zhu National Synchrotron Radiation Laboratory, University of Science and Technology of China, Hefei 230026 Google Scholar More articles by this author and Liang Zhang *Corresponding authors: E-mail Address: [email protected] E-mail Address: [email protected] Institute of Functional Nano & Soft Materials (FUNSOM), Jiangsu Key Laboratory for Carbon-Based Functional Materials & Devices, Soochow University, Suzhou 215123 Google Scholar More articles by this author https://doi.org/10.31635/ccschem.021.202101214 SectionsSupplemental MaterialAboutAbstractPDF ToolsAdd to favoritesDownload CitationsTrack Citations ShareFacebookTwitterLinked InEmail Lithium–sulfur (Li–S) batteries are one of the most promising alternatives to lithium–ion batteries because of the advantageous high energy density and low cost. However, the practical applications of Li–S batteries are hampered by a severe shuttle effect and sluggish polysulfide redox conversion. Herein, highly dispersed cobalt nanoparticles (∼0.8 wt %) embedded into nitrogen-doped hierarchical porous carbon ([email protected]) are designed as an effective electrocatalyst for Li–S batteries, which exhibit a synergistic effect of anchoring and dual-directional catalytic conversion of polysulfides. The experimental and theoretical studies reveal that [email protected] not only provides strong chemical affinity to polysulfides but also lowers the Li+ diffusion barrier and facilitates the precipitation and decomposition of Li2S, thus effectively inhibiting the shuttle effect and promoting the reaction kinetics of polysulfides. In addition, the well-dispersed Co nanoparticles in the three-dimensional carbon matrix guarantee the exposure of abundant polysulfide confining sites and catalytically active sites. Accordingly, the Li–S batteries assembled with [email protected] functional separators harvest a high rate performance (808.4 mAh g−1 at 10 C), a long-lasting cycle stability (0.055% decay per cycle over 1000 cycles at 4 C), and a superior areal capacity retention (5.78 mAh cm−2 after 100 cycles with a high sulfur loading of 7 mg cm−2). Download figure Download PowerPoint Introduction With the advancement of personal portable electronic devices and electric vehicles, there is an ever-increasing demand for high-power and fast-charging battery systems.1,2 Lithium–sulfur (Li–S) batteries have been proposed as a promising candidate for next-generation energy storage systems because of the high energy density, nature abundance, and environmental friendliness of sulfur. Nevertheless, the practical application of Li–S batteries faces the challenges of poor conductivity, shuttled polysulfides, and sluggish reaction kinetics.3 Multifarious strategies, including design of sulfur hosts, novel electrolytes and additives, and functionalized separators, have been adopted to circumvent the above-mentioned challenges.4–11 Among them, functionalizing separators with certain materials has proved effective at inhibiting the shuttle effect and improving the electrochemical performance.12,13 Initially, various carbonaceous materials, for example, one-dimensional carbon nanotubes, two-dimensional graphene, and three-dimensional porous carbon, have been utilized as functional materials to physically confine polysulfides.14–17 However, the weak interaction between nonpolar carbon materials and polar polysulfides cannot effectively suppress the polysulfide shuttle effect, resulting in a low sulfur utilization. In contrast, polar metallic compounds, such as metal oxides, metal sulfides, and metal–organic frameworks, have been introduced into carbon materials to improve the bonding with polysulfides through the Lewis acid–base interaction.18–21 However, the slow redox kinetics of polysulfides still leads to the accumulation of polysulfides and increases the electrolyte viscosity, especially under high sulfur loading and high-current rate conditions.22 Recently, catalysis has been introduced into Li–S batteries to simultaneously capture and convert polysulfides to accelerate the reaction kinetics.15,23–25 Noble metals, such as Pt and Ir, which possess excellent catalytic activity, have been utilized to improve the electrochemical performance of Li–S batteries; however, the high cost of noble metals restricts their practical applications.25–27 In contrast, 3d transition metals have also attracted extensive attention because of their low cost and decent catalytic activity.28–31 For example, Co and Ni nanoparticles combined with carbon materials have been synthesized using zeolitic imidazolate frameworks to improve the polysulfide redox kinetics.32–35 Nevertheless, the content of metals in the synthesized materials is usually high, and the rate performance and cycling stability are still unsatisfactory.36 Additionally, most of these reports only focused on one direction sulfur reaction (mainly reduction), and the fast cycling of sulfur species has not yet been finally realized. In principle, to realize the rapid polysulfide redox conversion, the designed catalyst should simultaneously possess a strong affinity for polysulfides, fast Li+/electron diffusion rate, and high catalytic activity toward polysulfide redox conversion. Here, we endow these properties to one dual-directional catalyst, namely well dispersed Co nanoparticles embedded into nitrogen-doped hierarchical porous carbon ([email protected]) with a low Co content (∼0.8 wt %), which is synthesized using a simple one-step pyrolysis strategy.37 The Li–S batteries assembled with [email protected] functional separators show a high rate performance of 808 mAh g−1 at 10 C, a low capacity fading rate of 0.055% per cycle over 1000 cycles at 4 C, as well as a high areal capacity of 5.78 mAh cm−2 after 100 cycles with a sulfur loading of 7 mg cm−2. Combining experiments and density functional theory (DFT) calculations, we demonstrate that the multiple effects of [email protected] are responsible for the improved electrochemical performance of Li–S batteries, which can not only strongly anchor polysulfides and facilitate the Li+/electron transport but also accelerate the polysulfide redox kinetics by reducing the energy barriers for Li2S deposition and decomposition. This work presents a facile strategy to develop cost-effective catalysts toward dual-directional polysulfide redox conversion for practical Li–S batteries. Experimental Methods Preparation of [email protected] First, 1.441 g urea and 0.198 g Co(NO3)2·6H2O were completely dissolved in 7.5 mL deionized water (DIW) and stirred for 0.5 h to form solution A. Then, 0.75 g polyethylene oxide-polypropylene oxide-polyethylene (P123) was completely dissolved in 7.5 mL DIW, 1 mL commercial acidic silica was added, and then the solution was stirred for 0.5 h to form solution B. After that, solution B was added to solution A and stirred for 3 h at room temperature. The mixture was gradually evaporated at 80 °C. The evaporated precursor was fully ground and moved into a porcelain boat under high-temperature pyrolysis. The precursor was heated to 500 °C for 2 h, and then heated to 900 °C for 2 h under N2 atmosphere. After cooling to room temperature, the obtained black powder was etched with 5% hydrofluoric acid (HF) (v/v in water) solution. The [email protected] was obtained by vacuum filtration and dried at 60 °C in a vacuum oven. As contrast samples, N-HPC was prepared without Co source by the same procedure, and N-C was obtained without Co source and silica sol. Polysulfide absorption experiments Sulfur and lithium sulfide (Li2S) in a 5:1 molar ratio were dissolved in 1,2-dimethoxyethane (DME) solvent with stirring at 60 °C for 48 h to produce Li2S6 solution. The concentration of Li2S6 solution was diluted to 2 mmol L−1 (mM). Then 10 mg of [email protected], N-HPC, and N-C powders were placed in 2 mL Li2S6 solution, respectively. After the solution was left standing for a designated time, the supernatant was absorbed for UV–vis spectroscopy, and the adsorbed powder was dried for X-ray photoelectron spectroscopy (XPS). Evaluation of polysulfide conversion kinetics [email protected], N-HPC, and N-C were mixed with polyvinylidene fluoride (PVDF) in N-methylpyrrolidinone (NMP) solvent with a mass ratio of 4:1, respectively. The mixed material was coated on the carbon coated aluminum foil, and then dried under vacuum at 60 °C for 12 h. The obtained electrodes were cut into disks with a diameter of 12 mm. Li2S6 (0.2 M) solution containing 1 M lithium bis(trifluoromethanesulfonyl)imide (LiTFSI) and 2.0 wt % LiNO3 in DME and 1,3-dioxolane (DOL) (1:1, v/v) solvent (conversional electrolyte) was used as electrolyte. Two identical electrodes were used to assemble a symmetrical battery with 30 μL electrolyte. Cyclic voltammetry (CV) tests were performed with a scan rate of 3 mV s−1 and a voltage range of 1 to −1 V. The nucleation and dissolution experiments of Li2S A 0.5 M Li2S8 solution was prepared by mixing a 7:1 molar ratio of sulfur and Li2S in tetraglyme solvent. The [email protected] and N-C powders were ultrasonically dispersed in ethanol, and then the dispersed solution was dropped on a carbon cloth with a diameter of 12 mm and an areal mass loading of 0.2 mg cm−2. The loaded carbon cloth was used as the cathode, and lithium metal was used as the anode; 15 μL of 0.5 M Li2S8 solution was placed on the cathode side, and 15 μL of conversional electrolyte was placed on the anode side. The assembled batteries were galvanostatically discharged to 2.06 V at a current of 0.112 mA and then potentiostatically discharged to 2.05 V until the discharged current was below 0.01 mA. Similarly, the batteries were assembled using the same electrode material and electrolyte. The batteries were first galvanostatically discharged to 1.7 V at a current of 0.1 mA and then allowed to stand for 300 s. During the standing process, the voltage moved toward the equilibrium potential higher than 1.8 V, and therefore, the batteries then were galvanostatically discharged to 1.8 V at 0.1 mA to ensure that polysulfides were fully converted into solid Li2S. Finally, the nucleation and dissolution capacity of Li2S were evaluated by calculating the integral area of the curve drawn according to Faraday’s law. Preparation of [email protected] functional separations [email protected] functional separations were prepared by a slurry casting method. [email protected], acetylene black, and PVDF were mixed in NMP solvent with a mass ratio of 7:2:1, and the mixed precursor was coated on the polypropylene separator (PP, Celgard-2500). After drying, functional separators were cut into disks with a diameter of 19 mm. The areal loading of the material was controlled at 0.4 mg cm−2. Fabrication of Li–S batteries and electrochemical tests The Li–S batteries were assembled under Ar in a glove box with the oxygen and water content below 0.1 ppm. The amount of electrolyte was maintained at an electrolyte to sulfur (E/S) ratio of 15 μL mg−1. The sulfur cathode consists of 70% sulfur powder, 20% acetylene black, and 10% PVDF. Sulfur loading was approximately 1 mg cm−2 for routine testing (E/S = 15 μL mg−1) and 7 mg cm−2 for high loading testing (E/S = 15 and 10 μL mg−1). Note that the high sulfur loading cathode was obtained by mixing 90% sulfur and 10% conductive carbon with carbon cloth as the current collector, and the weight of carbon cloth was not taken into consideration in general. The batteries were assembled with lithium metal as the anode and sulfur as the cathode with modified separators to form the standard 2025 coin-type batteries. The galvanostatic charge–discharge test was measured by the Landt CT2001A battery test system. A CHI660E electrochemical workstation was used for CV tests, and the scan rate was from 0.1 to 0.5 mV s−1 in the voltage range of 1.7 to 2.8 V. Electrochemical impedance spectroscopy (EIS) was measured in the frequency range of 10 mHz to 100 kHz. Computational Methods All the DFT simulations in this work were performed using the Vienna ab initio Simulation Package (VASP).38 The all-electron projector augmented wave (PAW) model was used for describing the electron-ion interaction.39 The Perdew–Burke–Ernzerhof (PBE) exchange and correlation functional was employed for the structural optimization and the energy calculation.40 A kinetic energy cutoff of 400 eV was used for the plane-wave expansion of the electronic wave function. The Hellmann–Feynman force and energy convergence criteria were set to 0.02 eV/Å and 10–4 eV in structural optimization, respectively. The four layers of (3 × 3) supercell of Co (111) containing 144 atoms with the bottom two layers fixed were used as the substrates for redox reaction of Li–S molecules, with periodic vacuum layers of 15 Å and Г-centered k-point sampling grid of 1 × 1 × 1.The climbing-image nudged-elastic-band (CI-NEB) approach was applied to calculate the energy barrier of single Li diffusion and Li decomposition from Li2S molecule on the monolayer nitrogen-doped graphene (N-G) and Co (111) surfaces.41 To further study the absorption effect between Li2Sx (x = 1, 6) and samples, the first-principle simulations were performed. As shown in Supporting Information Figure S9, two atom monolayers (AMs) [N-G, Co (111) surfaces] were selected as the representative to calculate the binding energy with Li2Sx (x = 1, 6), respectively. From the top view of adsorption conformations of Li2Sx (x = 1, 6) on the AMs, we can conclude that the chemical interaction is the most important interaction between Li2Sx (x = 1, 6) and AMs. The binding strengthens between Li2Sx (x = 1, 6) and AMs, and the Co (111) surface has the highest binding energy for Li2Sx (x = 1, 6) as compared to N-G surface. The Co (111) surface has the highest binding energy for Li2Sx (x = 1, 6) as compared with N-G surface. The binding energy, Eb, is computed to measure the binding strength between Li2Sx (x = 1, 6) and AMs. It is defined as the energy difference between the summation of pure Li2Sx and pure AMs and Li2Sx − AM adsorbed system (ELi2Sx + AM) and can be expressed as Eb = ELi2Sx + AM − (ELi2Sx + EAM). With this definition, a negative binding energy indicates that the binding interaction is favored. All the Eb magnitudes and atomic configurations showed below are those of the most stable adsorption cases. Results and Discussion Synthesis and characterization of [email protected] We employed a one-step pyrolysis to synthesize [email protected] (Figure 1a). First, a mixture of cobalt nitrate [Co(NO3)2], urea, P123 (surfactant), and SiO2 sol were dissolved in DIW at room temperature and evaporated at 80 °C to evenly distribute Co and SiO2. Then the mixed precursors were pyrolyzed at 900 °C and etched by HF solution to obtain the functional composite of [email protected] For comparison, two other samples, that is, N-HPC without Co and N-C without Co or SiO2 sol, were also prepared by the same method. As shown in the scanning electron microscopy (SEM) images, the as-synthesized [email protected] exhibits a hierarchical porous structure with a three-dimensional carbon skeleton (Figures 1b and 1c and Supporting Information Figures S1a and S1b). N-HPC shows a similar structure to that of [email protected], whereas N-C displays unevenly distributed carbon blocks ( Supporting Information Figures S1c–S1f), implying that the formation of a hierarchical porous structure is related to the synergistic effect of the silica template and carbon pyrolysis. Figure 1 | (a) Illustration of the [email protected] synthesis. (b and c) SEM images, (d and e) TEM images, (f and g) high-resolution HADDF-STEM images, (h) HADDF-STEM image and the corresponding element mapping of [email protected] Download figure Download PowerPoint The transmission electron microscopy (TEM) images of [email protected] (Figures 1d and 1e) further reveal the three-dimensional porous structure and uniformly distributed mesopores with an average size of 10–20 nm. Furthermore, the high-resolution high-angle annular dark-field scanning TEM (HADDF-STEM) images (Figures 1f and 1g and Supporting Information Figure S2) show well-dispersed Co nanoparticles that are surrounded by a graphitic carbon matrix. In addition, the distance between adjacent lattice fringes is 0.21 nm, corresponding to the (111) crystal plane of metallic Co. The Co content of [email protected] determined by inductively coupled plasma-optical emission spectrometry (ICP-OES) is ∼0.8 wt % ( Supporting Information Table S1). Energy-dispersive X-ray spectroscopy (EDS) mappings also demonstrate that Co, N, and C elements are homogeneously distributed and lack aggregated Co nanoparticles in [email protected] (Figure 1h and Supporting Information Figure S3). The formation of highly dispersed Co nanoparticles is related to the etching of unstable and agglomerated Co nanoparticles by HF solution. The X-ray diffraction (XRD) pattern of [email protected] (Figure 2a) clearly demonstrates the characteristic graphitic carbon peak at 26.2°, which is consistent with the TEM results. However, no such peak is observed for N-HPC and N-C, suggesting that the Co source facilitates the formation of graphitic carbon. Note that the diffraction peak of metallic Co is absent because of the low Co content in [email protected]42 Moreover, the Raman results (Figure 2b) suggest that the intensity ratio of G band (1590 cm−1) to D band (1330 cm−1) (IG/ID) is higher for [email protected] (1.02) compared with that of N-HPC (0.92) and N-C (0.87),43 implying a higher degree of graphite and therefore promoted electrical conductivity for [email protected] The N2 adsorption and desorption isotherms show the characteristics of typical IV isotherms with the pore size of 5–20 nm, indicating that [email protected] is mainly dominated by mesoporous structure (Figures 2c and 2d). [email protected] also exhibits the largest specific surface area (1248.32 m2 g−1) among the samples ( Supporting Information Table S2), which should be beneficial for the physical constraint of polysulfides and nucleation of Li2S.44 Figure 2 | (a) XRD patterns, (b) Raman spectra, (c) N2 adsorption/desorption isotherms, and (d) the corresponding pore size distributions of [email protected], N-HPC, and N-C. (e) N 1s XPS spectrum of [email protected] (f) Co K-edge XANES, (g) FT-EXAFS, and (h–j) WT-EXAFS of [email protected], Co foil, and CoO. Download figure Download PowerPoint XPS was carried out to investigate the chemical states of Co and N in [email protected] As shown in Figure 2e, the N 1s XPS spectrum displays four different N species ( Supporting Information Table S4), that is, pyridinic N (398.7 eV, 37.7%), pyrrolic N (399.9 eV, 7.2%), graphitic N (401.2 eV, 49.3%), and oxidized N (403.1 eV, 5.8%). Furthermore, the N contents of [email protected], N-HPC, and N-C samples were calculated ( Supporting Information Table S3), indicating the successful incorporation of N in the carbon matrix. It has been previously reported that an N-doped carbon surface could facilitate polysulfide redox kinetics by improving the electron transfer, resulting in enhanced sulfur utilization and cycling performance.45,46 The Co 2p XPS spectrum ( Supporting Information Figure S4) reveals that in addition to metallic Co, Co2+ and Co3+ are also observed, which may be ascribed to the slight carbonization of certain Co atoms during the pyrolysis process.33 To further probe the electronic structure and local coordination environment of Co nanoparticles in [email protected], Co K-edge X-ray adsorption near-edge structure (XANES) and extended X-ray absorption fine structure (EXAFS) measurements were performed. The overall spectral shape of [email protected] is like that of Co foil (Figure 2f), reflecting the dominant metallic property of Co in [email protected]42 Moreover, in contrast to CoO, only one dominant peak at 2.12 Å is observed for the Fourier-transformed EXAFS (FT-EXAFS) spectrum of [email protected] (Figure 2g), corresponding to the first shell scattering of the Co–Co bond. Note that the Co–Co bond length of [email protected] is slightly shorter than that of Co foil, which could be related to the size effect of Co foil and [email protected]47 The wavelet transform (WT) analysis of [email protected] displays one intensity maximum at 6.90 Å−1 that is correlated with the Co–Co bond (Figures 2h–2j), further illustrating the presence of metallic Co in [email protected] Li + diffusion kinetics We next utilized [email protected] as a functional material to modify the commercial PP separators of Li–S batteries ( Supporting Information Figures S5a–S5c). The as-prepared functional separators exhibit an excellent electrolyte wettability and a strong mechanical toughness ( Supporting Information Figures S5h–S5l). For comparison, N-HPC and N-C functional separators were also fabricated and investigated ( Supporting Information Figures S5d–S5g, S5m, and S5n). Figure 3a demonstrates the initial discharge/charge voltage profiles of Li–S batteries using [email protected], N-HPC, and N-C functional separators at 0.2 C (1 C = 1675 mA g−1). ΔH1 and ΔH2 represent the reduction from S8 to short-chain Li2S4 and then to Li2S2/Li2S during the discharge process, respectively.48 Clearly, Li–S batteries with [email protected] functional separators exhibit larger ΔH1 and ΔH2 values (Figure 3a and Supporting Information Figure S6), suggesting that more sulfur is involved in the electrochemical reduction process by converting to polysulfides and then to Li2S. Figure 3 | (a) Initial discharge/charge voltage profiles and (b) Nyquist plots of Li–S batteries assembled with [email protected], N-HPC, and N-C functional separators. (c–e) CV curves of Li–S batteries with different separators at different scan rates from 0.1 to 0.5 mV s−1 and (f–h) the corresponding linear fits of redox peak currents with IA, IB, and IC. (i and j) Calculated energy profiles of Li diffusion along the minimum energy path on the surfaces of Co (111) and monolayer N-G. The green, brown, silver, and blue balls correspond to Li, C, N, and Co atoms, respectively. Download figure Download PowerPoint The ion diffusion kinetics is a vital factor affecting polysulfide redox kinetics, which we investigated by EIS and CV measurements.49 The EIS results reveal that [email protected] functional separators show the smallest charge-transfer resistance (10.2 Ω) (Figure 3b), indicative of the fast charge transfer at the Co/polysulfides interface.50 Notably, N-HPC also has an excellent electrical conductivity that improves the interfacial charge-transfer rate.51 Figures 3c–3e show the CV measurements of Li–S batteries assembled with different separators at various scan rates from 0.1 to 0.5 mV s−1. The cathodic peaks IA and IB correspond to the reduction process of S8 to Li2Sx and Li2Sx to Li2S2/Li2S, respectively, while the anodic peak IC represents the oxidation process from Li2S2/Li2S to S8.52 Compared with N-C and N-HPC, the Li–S batteries with [email protected] functional separators show the highest peak currents with reduced voltage polarizations ( Supporting Information Figure S7), suggesting the presence of Co nanoparticles greatly accelerates the polysulfide redox kinetics.53 According to the Randles–Sevcik equation,49,54 the Li+ diffusion coefficients are proportional to the slopes of Ip/v0.5 (Figures 3f–3h), where Ip is the peak current and v is the scan rate. The calculated slope values and Li+ diffusion coefficients are summarized in Supporting Information Table S5. It reveals that the Li–S batteries with [email protected] functional separators demonstrate the largest Li+ diffusion coefficient, implying that the well dispersed Co nanoparticles in [email protected] could greatly enhance the Li+ diffusion rate, which should be beneficial for the redox kinetics of polysulfides. In addition, the redox peaks in the CV curves of Li–S batteries with [email protected] functional separators overlap well in the initial three cycles ( Supporting Information Figure S8), suggesting the excellent cycling stability and high reversibility of sulfur conversion. To better understand why well distributed Co nanoparticles could facilitate the Li+ diffusion kinetics, we have calculated the Li diffusion barriers on N-G and Co surfaces along with the Li diffusion path ( Supporting Information Figure S9 and Figures 3i and 3j). The calculated diffusion barriers are 0.45 and 0.36 eV for N-G and Co, respectively, indicating the improved Li diffusion on the Co surface. Overall, the enhanced electron/ion diffusion rate on [email protected] surface ensures the accelerated polysulfide redox kinetics during the electrochemical process, which could greatly suppress the shuttle effect and enhance the sulfur utilization. Polysulfide adsorption and catalytic conversion The adsorption and conversion of polysulfides play a decisive role in improving the electrochemical performance of Li–S batteries. To probe the bonding strength with polysulfides, adsorption experiments were performed by adding [email protected], N-HPC, N-C powders into Li2S6 solution using DME as solvent. As shown in Figure 4a, after 30 min of standing, the [email protected] and N-HPC solutions turned transparent, whereas the N-C solution remained pale yellow. The UV–vis spectra (Figure 4a) further demonstrate that [email protected] binds the strongest with polysulfides. To further explore the strong affinity between [email protected] and polysulfides, synchrotron radiation photoemission spectroscopy (SRPES) measurements are performed on Li2S6 before and after interacting with [email protected] (Figure 4b). For the S 2p SRPES spectrum (photon energy: 250 eV) of Li2S6, there are two peaks, which are assigned to bridging sulfur (SB0) and terminal sulfur (ST−1) species, respectively.55 After interacting with [email protected], both species shift to lower binding energies because of charge transfer from Co to Li2S6, confirming the strong chemical interaction between [email protected] and Li2S6. This scenario is also verified by the first-principle simulations. Figures 4c and 4d show the adsorption conformations of Li2S6 on Co and N-G surfaces, respectively. The binding energy (Eb) between Li2S6 and Co (−5.41 eV) is much stronger compared with that of N-G (−0.30 eV), similarly to the case of Li2S ( Supporting Information Figure S10). The results clearly reflect that the well dispersed Co nanoparticles in [email protected] can effectively confine polysulfides, which is the first

Major Factors Influencing the Size Distribution Analysis of Cellulose Nanocrystals Imaged in Transmission Electron Microscopy.

Size distributions of cellulose nanocrystals (CNCs), extracted from softwood pulp via strong sulfuric acid hydrolysis, exhibit large variability when analyzed from transmission electron microscopy (TEM) images. In this article, the causes of this variability are studied and discussed. In order to obtain results comparable with those reported, a reference material of CNCs (CNCD-1) was used to evaluate size distribution. CNC TEM specimens were prepared as-stained and dried with a rapid-flushing staining method or hydrated and embedded in vitreous ice with the plunge-freezing method. Several sets of bright-field TEM (BF-TEM), annular dark-field scanning TEM (ADF-STEM) and cryogenic-TEM (cryo-TEM) images were acquired for size distribution analysis to study the contributing factors. The rapid-flushing staining method was found to be the most effective for contrast enhancement of CNCs, not only revealing the helical structure of single CNCs but also resolving the laterally jointed CNCs. During TEM specimen preparation, CNCs were fractionated on TEM grids driven by the coffee-ring effect, as observed from contrast variation of CNCs with a stain-depth gradient. From the edge to the center of the TEM grids, the width of CNCs increases, while the aspect ratio (length to width) decreases. This fractionated dispersion of CNCs suggests that images taken near the center of a droplet would give a larger mean width. In addition to particle fractionation driven by the coffee-ring effect, the arrangement and orientation of CNC particles on the substrate significantly affect the size measurement when CNC aggregation cannot be resolved in images. The coexistence of asymmetric cross-section CNC particles introduces a large variation in size measurement, as TEM images of CNCs are mixed projections of the width and height of particles. As a demonstration of how this contributes to inflated size measurement, twisted CNC particles, rectangular cross-section particles and end-to-end jointed CNCs were revealed in reconstructed three-dimensional (3D) micrographs by electron tomography (ET).

Electroforming and threshold switching characteristics of NbOx films with crystalline NbO2 phase

Threshold switching (TS) and negative differential resistance (NDR) characteristic of niobium oxide (NbOx) films have been actively studied for neuromorphic computing. Generally, the electroforming process is required for TS and NDR in NbOx films. However, different electroforming and TS properties have been reported for NbOx films with different crystallinities or chemical compositions. This study investigates the effect of thermal annealing on the microstructures of NbOx films and compares the electroforming, TS, and NDR characteristics of amorphous, partially crystallized, and fully crystallized films. The distributions of crystalline NbO2 phase in NbOx films annealed at various temperatures were analyzed using transmission electron microscopy dark-field imaging, and it was observed that the distribution of crystalline NbO2 phase influenced the electroforming process. Moreover, TS characteristics improved in the thermally annealed NbOx films with crystalline NbO2 phases.

Operando leaching of pre-incorporated Al and mechanism in transition-metal hybrids on carbon substrates for enhanced charge storage

•A "nano-tailoring" strategy via electrochemical leaching of Al species is proposed•Defective and highly active materials were formed via the "nano-tailoring" strategy•The integrated oxygen vacancies can boost charge storage of potassium-birnessite•The reducibility of M2+ is identified as the key descriptor for the reconstruction rate From the viewpoint of optimizing charge-storage dynamics, we for the first time propose a universal "nano-tailoring" strategy to fabricate low-crystalline and highly active materials via electrochemically dynamic leaching of pre-incorporated Al. With M2+Al hydroxides as the precursor, potassium-birnessite MnO2 with abundant defective sites and enriched edge-active sites is configured, resulting in the enhanced charge-storage capability. Soft XAS and Raman mapping techniques are adopted to finely track the dynamic reconstruction principles and intrinsic active sites. It is found that the reducibility of M2+ serves as the key descriptor for the reconstruction rate, which correspondingly is divided into two branches: oxidization-boosting type and non-oxidization type. This work can provide a novel avenue and spark new ideas to electrochemically modulate and ameliorate electrode materials for advanced energy storage and conversion applications in the future. Insufficient exposure and utilization of active sites often induces an inferior reactivity for transition-metal-based two-dimensional (2D) materials. In response, we for the first time propose a universal "nano-tailoring" strategy to incorporate abundant defects and active sites into low-crystallinity nanosheets by electrochemically leaching of Al species. With MnAl layered double hydroxides (LDHs) as a representative example, potassium-birnessite MnO2 (AK-MnO2) with oxygen vacancies and abundant edge sites is successfully produced. The oxygen vacancies are shown to help optimize the electron-transfer and ion-adsorption capability. These integrated advantages endow the AK-MnO2 with a high capacitance value of 239 F g−1 at 100 A g−1. By further combining with soft X-ray absorption spectroscopy techniques, we unravel that the reducibility of M2+ in M2+Al-LDH serves as the key descriptor for the reconstruction rate. This "nano-tailoring" strategy can provide some important implications and clues to manipulating 2D materials for efficient energy storage and conversion. Insufficient exposure and utilization of active sites often induces an inferior reactivity for transition-metal-based two-dimensional (2D) materials. In response, we for the first time propose a universal "nano-tailoring" strategy to incorporate abundant defects and active sites into low-crystallinity nanosheets by electrochemically leaching of Al species. With MnAl layered double hydroxides (LDHs) as a representative example, potassium-birnessite MnO2 (AK-MnO2) with oxygen vacancies and abundant edge sites is successfully produced. The oxygen vacancies are shown to help optimize the electron-transfer and ion-adsorption capability. These integrated advantages endow the AK-MnO2 with a high capacitance value of 239 F g−1 at 100 A g−1. By further combining with soft X-ray absorption spectroscopy techniques, we unravel that the reducibility of M2+ in M2+Al-LDH serves as the key descriptor for the reconstruction rate. This "nano-tailoring" strategy can provide some important implications and clues to manipulating 2D materials for efficient energy storage and conversion. 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To the best of our knowledge, aluminum, as a typical amphoteric metal element, displays a tendency to spontaneously dissolve in alkaline electrolytes, while this is thermodynamically unfavorable for transition-metal elements. In addition, Al-based transition-metal hybrids have a big family, including double/triple/multiple oxides, hydroxides, sulfides, and so on. Accordingly, it is expected that pre-incorporating Al into the lattice of transition-metal hybrids, followed by its in situ leaching, can generate a series of advanced materials with promoted reactivity, which remain less concerning and require more attention and detailed explorations. Herein, a universal "nano-tailoring" strategy, via the operando leaching of pre-incorporated Al driven by electrochemical effects, is first proposed for realizing the holistic activation of transition-metal hybrids. Using MnAl layered double hydroxides (LDHs) as a demonstration, we show that low-crystallinity potassium-birnessite MnO2 (AK-MnO2) can be produced quickly after the "nano-tailoring" process. With the help of soft X-ray absorption spectroscopy (sXAS), the structural evolution is finely tracked and decoupled. Interestingly, due to the irreversible leaching of Al and the intercalation of K+, the microstructure may undergo the emergence and release of internal strain, resulting in a positively tailored microstructure with enriched edge sites and oxygen vacancies. Also, density functional theory (DFT) calculation demonstrates that the introduced oxygen vacancies would be capable of promoting charge transfer and facilitating the adsorption of electrolyte ions, leading to optimized reaction kinetics for charge storage. As such, the as-formed AK-MnO2 hybrids can deliver a high capacitance value of 356 F g−1 at 1 A g−1, with a high retention rate of up to 67% at a super-large current density of 100 A g−1, indicative of a superior rate performance. More importantly, the universality of the "nano-tailoring" process to generate low-crystallinity nanosheets with enriched active sites is deeply verified and decoupled. We reveal that the leaching transformation is highly dependent on the reducibility of the adopted M2+, which acts as a prominent factor for the reconstruction rate. It is believed that the unique and universal "nano-tailoring" strategy can stimulate broad interest and inspiration of researchers to configure novel and highly efficient materials with positively tuned properties. As illustrated in Figure 1A, the "nano-tailoring" strategy involves M2+Al-based binary hydroxides as the precursor, where the molecular breakage and reconstruction happens during the electrochemical scanning process in an alkaline medium. As a typical example, during the "nano-tailoring" process, the MnAl-LDH will undergo a transformation into MnAl-based hydroxides, accompanied by partial leaching of Al, when soaked in an alkaline electrolyte for a short time (2 min), yielding Mn(Al)-OH as the intermediate. Subsequently, it will be reconstructed into low-crystallinity birnessite-type MnO2 with K+ in the interlayer after cyclic voltammetry (CV) cycling treatment, denoted as AK-MnO2. The dynamic reconstruction is ex situ tracked by scanning electron microscopy (SEM) images and the corresponding size distributions. As shown in Figures 1B–1D, the large LDH nanosheets are broken and reconstructed to much smaller ones with a higher stacking density. The average size decreases from 726 nm for the MnAl-LDH to 202 nm for the Mn(Al)-OH hybrids, and finally to 55 nm for the AK-MnO2 hybrids (about 13 times decrease in size). Notably, the small and dense nanosheets would deliver more abundant active sites and feature numerous short-distance channels for fast charge transfer, and thus are beneficial for highly efficient charge storage. The microstructure is further uncovered by transmission electron microscopy (TEM) images. As depicted in Figure S1A, the MnAl-LDH precursor displays a large size, and a lattice spacing of 0.26 nm corresponds to the (012) plane of LDH. For the Mn(Al)-OH, a broken structure with numerous in situ-formed holes can be observed, and the lattice spacing of 0.33 nm corresponds to the (100) plane of Mn(OH)2 (Figure S1B). Impressively, the well-tailored AK-MnO2 with densely stacked and ultrasmall nanosheet microstructure (Figure 1E) displays a low crystallinity, as shown in the high-resolution (HR)-TEM image (Figure 1F). It is further identified by the result of selected area electron diffraction (SAED), where no lattice patterns can be detected (Figure 1F), indicative of the generation of lattice distortion and randomly stacked internal construction. Furthermore, as presented by high-angle annular dark-field scanning TEM (HAADF-STEM) images and the corresponding elemental maps, the Mn, O, and K elements are uniformly distributed on the surface of the AK-MnO2 (Figure 1E). Also, we carried out atomic force microscopy (AFM) characterization to better study the morphology of the AK-MnO2. In Figure S2, the ultrathin and wrinkled nanosheet morphology with a thickness of about 7.7 nm is presented. In addition, the AK-MnO2 demonstrates a highly increased surface area (157 m2 g−1) as well as more abundant pore structure (Figure S3), compared with the MnAl-LDH (58 m2 g−1). The enlarged surface area, enabled by the novel "nano-tailoring" process, can help give access to more electrolyte ions at a given time and shorten the transport distance, thus accounting for the enhanced charge storage. The "nano-tailoring" process of the MnAl-LDH is further decoupled by X-ray diffraction (XRD) patterns (Figure S4). Notably, the intense peaks located at 26.4°, 43.1°, 54.6°, and 77.6° can be indexed to the characteristic peaks of carbon paper (CP). For the MnAl-LDH hybrids, the peaks at 10.1°, 20.2°, 33.1°, 36.4°, and 58.8° can be indexed to the (003), (006), (012), (015), and (110) planes of the LDH phase, respectively. Explicitly, for the Mn(Al)-OH hybrids, the peaks at 18.8°, 31.6°, 36.8°, 49.9°, and 59.5° can be indexed to the (001), (100), (101), (102), and (111) diffraction planes of Mn(OH)2 (JCPDS card no. 18-0787), respectively, indicative of the phase transformation from the MnAl-LDH to the Al-doped Mn(OH)2 phase. However, as for the AK-MnO2, it displays only one significant peak at about 12.7°, as well as an indiscernible wide band in the range of 35° to 40°. This implies the low-crystallinity/amorphous features of the [MnO6] octahedra laminate.36Lin B. Zhu X. Fang L. Liu X. Li S. Zhai T. Xue L. Guo Q. Xu J. Xia H. Birnessite nanosheet arrays with high K content as a high-capacity and ultrastable cathode for K-ion batteries.Adv. Mater. 2019; 31: 1900060Crossref PubMed Scopus (132) Google Scholar,37Dong Z.H. Lin F. Yao Y.H. Jiao L.F. Crystalline Ni(OH)2/amorphous NiMoOx mixed-catalyst with Pt-like performance for hydrogen production.Adv. Energy Mater. 2019; 9: 1902703Crossref Scopus (78) Google Scholar Notably, an increase of interlayer spacing from Mn(Al)-OH (d spacing: 4.7 Å) to AK-MnO2 (d spacing: 7.0 Å) can be attributed to a certain amount of K+ intercalated into the interlayer. To further confirm the existing species, Raman measurements were carried out. For the AK-MnO2 hybrids, the characteristic peaks at 265, 317, 396, 469, 578, and 654 cm−1 are matched with those of birnessite-type MnO2 with a little shift (Figure 1G), attributed to the varied ion radius between K+ and Na+.38Julien C. Raman spectra of birnessite manganese dioxides.Solid State Ionics. 2003; 159: 345-356Crossref Scopus (570) Google Scholar And the intense peaks at 1,350 and 1,580 cm−1 could be indexed to the D and G bands of CP. Moreover, Raman maps were recorded in the green, orange, and pink regions (as marked in Figure 1G), where the uniform distribution for the signals of the CP (orange and pink) and the AK-MnO2 (green) are presented in the maps (Figure 1H), indicative of the intimate interaction between them. The sXAS technique was used to decouple the fine structure of the as-formed hybrids and understand the dynamic phase reconstruction process. The mechanism of sXAS is illustrated in Figure 2A (refer to Yang and Devereaux and Lin et al.).39Yang W. Devereaux T.P. Anionic and cationic redox and interfaces in batteries: advances from soft X-ray absorption spectroscopy to resonant inelastic scattering.J. Power Sources. 2018; 389: 188-197Crossref Scopus (132) Google Scholar,40Lin F. Nordlund D. Markus I.M. Weng T.-C. Xin H.L. Doeff M.M. Profiling the nanoscale gradient in stoichiometric layered cathode particles for lithium-ion batteries.Energy Environ. 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Energy Mater. 2017; 7: 1602010Crossref Scopus (34) Google Scholar The Mn L-edge soft XAS spectra are shown in Figure 2B. It is worth noting that the typical peaks at 638.8, 640.2, and 641.6 eV can be assigned to Mn2+, Mn3+, and Mn4+, respectively, while the peak intensity is generally used to identify the relative content.42Henderson C.M.B. Cressey G. Redfern S.A.T. Geological applications of synchrotron radiation.Radiat. Phys. Chem. 1995; 45: 459-481Crossref Scopus (66) Google Scholar Clearly, for the MnAl-LDH, the peak at 638.8 eV is much more intensive than the others, implying that the Mn primarily exists in the +2 status. Compared with the MnAl-LDH, the Mn(Al)-OH hybrids are also dominated by Mn2+, despite the slight increase in peak intensity of Mn3+. Interestingly, the Mn4+ ions are dominant for the as-formed AK-MnO2, where a certain amount of Mn3+ remained. It is noteworthy that the charge compensation can be realized by the intercalation of K+, which is described below in detail. Moreover, compared with the 30-cycled sample (AK-MnO2), no further obvious changes in Mn L-edge are found after 50 cycles, implying that a stable status can be attained after 30 cycles. These results correspond finely to the Mn 2p and Mn 3s X-ray photoelectron spectroscopy (XPS) results (Figure S5). To gain a more complete picture of the reconstruction process, the O K-edge characterization was conducted (Figure 2C). The pre-edge region (below 533 eV) and broad-band region (over 533 eV) can be attributed to the transition of O 1s to Mn 3d and O 2p hybrid-state orbital, and O 1s to Mn 4sp and O 2p hybrid-state orbital, respectively. Generally speaking, the positive shift of the broad band can indicate an increase in Mn valence.43Oishi M. Yamanaka K. Watanabe I. Shimoda K. Matsunaga T. Arai H. Ukyo Y. Uchimoto Y. Ogumi Z. Ohta T. Direct observation of reversible oxygen anion redox reaction in Li-rich manganese oxide, Li2MnO3, studied by soft X-ray absorption spectroscopy.J. Mater. Chem. A. 2016; 4: 9293-9302Crossref Google Scholar Compared with the MnAl-LDH, there is no detectable shift of the broad band for the Mn(Al)-OH, indicative of the unobvious oxidation of Mn species. Nevertheless, the significantly positive shift can be detected for the 30-cycled sample (AK-MnO2), indicative of the oxidation of Mn species to a great degree. The 50-cycled sample demonstrates no more shift compared with the AK-MnO2, indicative of complete transformation after 30 electrochemical cycles, very consistent with the results we mentioned above. For the C K-edge spectra (Figure S6), the two new peaks at 299 and 302 eV, which correspond to the characteristic peaks of K L3- and L2-edge,44Dey A. Krishnamurthy S. Bowen J. Nordlund D. Meyyappan M. Gandhiraman R.P. Plasma jet printing and in situ reduction of highly acidic graphene oxide.ACS Nano. 2018; 12: 5473-5481Crossref PubMed Scopus (23) Google Scholar emerge for 30- and 50-cycled electrodes, indicative of the intercalation of K+. This is further proved by the presence of the intense K 2p peak signals for AK-MnO2 (Figure S7A). On top of that, the drastic reconstruction process enables the formation of internal defects to a great degree, which is studied and analyzed by the depth spectra of O 1s (Figure 2D). The peaks located at 530, 531.7, and 533 eV can be finely detected, corresponding to the bonding of metal and oxygen, oxygen vacancy, and physically/chemically adsorbed O, respectively.45Bao J. Zhang X.D. Fan B. Zhang J.J. Zhou M. Yang W.L. Hu X. Wang H. Pan B.C. Xie Y. Ultrathin spinel-structured nanosheets rich in oxygen deficiencies for enhanced electrocatalytic water oxidation.Angew. Chem. Int. Ed. 2015; 54: 7399-7404Crossref PubMed Scopus (923) Google Scholar,46Fang G.Z. Zhu C.Y. Chen M.H. Zhou J. Tang B.Y. Cao X.X. Zheng X.S. Pan A.Q. Liang S.Q. Suppressing manganese dissolution in potassium manganate with rich oxygen defects engaged high-energy-density and durable aqueous zinc-ion battery.Adv. Funct. Mater. 2019; 29: 1808375Crossref Scopus (373) Google Scholar Impressively, the intensive peak at 531.7 eV implies a sign

Boosting the Energy Migration Upconversion through Inter-Shell Energy Transfer in Tb 3+ -Doped Sandwich Structured Nanocrystals

Open AccessCCS ChemistryRESEARCH ARTICLE2 Jul 2021Boosting the Energy Migration Upconversion through Inter-Shell Energy Transfer in Tb3+-Doped Sandwich Structured Nanocrystals Dan Xu, Jin Xu, Xiaoying Shang, Shaohua Yu, Wei Zheng, Datao Tu, Renfu Li and Xueyuan Chen Dan Xu CAS Key Laboratory of Design and Assembly of Functional Nanostructures, Fujian Key Laboratory of Nanomaterials, and State Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fuzhou, Fujian 350002 University of Chinese Academy of Sciences, Beijing 100049 Google Scholar More articles by this author , Jin Xu *Corresponding authors: E-mail Address: [email protected] E-mail Address: [email protected] CAS Key Laboratory of Design and Assembly of Functional Nanostructures, Fujian Key Laboratory of Nanomaterials, and State Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fuzhou, Fujian 350002 Fujian Science and Technology Innovation Laboratory for Optoelectronic Information of China, Fuzhou, Fujian 350108 Google Scholar More articles by this author , Xiaoying Shang CAS Key Laboratory of Design and Assembly of Functional Nanostructures, Fujian Key Laboratory of Nanomaterials, and State Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fuzhou, Fujian 350002 Google Scholar More articles by this author , Shaohua Yu CAS Key Laboratory of Design and Assembly of Functional Nanostructures, Fujian Key Laboratory of Nanomaterials, and State Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fuzhou, Fujian 350002 Google Scholar More articles by this author , Wei Zheng CAS Key Laboratory of Design and Assembly of Functional Nanostructures, Fujian Key Laboratory of Nanomaterials, and State Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fuzhou, Fujian 350002 Fujian Science and Technology Innovation Laboratory for Optoelectronic Information of China, Fuzhou, Fujian 350108 Google Scholar More articles by this author , Datao Tu CAS Key Laboratory of Design and Assembly of Functional Nanostructures, Fujian Key Laboratory of Nanomaterials, and State Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fuzhou, Fujian 350002 Fujian Science and Technology Innovation Laboratory for Optoelectronic Information of China, Fuzhou, Fujian 350108 Google Scholar More articles by this author , Renfu Li CAS Key Laboratory of Design and Assembly of Functional Nanostructures, Fujian Key Laboratory of Nanomaterials, and State Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fuzhou, Fujian 350002 Fujian Science and Technology Innovation Laboratory for Optoelectronic Information of China, Fuzhou, Fujian 350108 Google Scholar More articles by this author and Xueyuan Chen *Corresponding authors: E-mail Address: [email protected] E-mail Address: [email protected] CAS Key Laboratory of Design and Assembly of Functional Nanostructures, Fujian Key Laboratory of Nanomaterials, and State Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fuzhou, Fujian 350002 University of Chinese Academy of Sciences, Beijing 100049 Fujian Science and Technology Innovation Laboratory for Optoelectronic Information of China, Fuzhou, Fujian 350108 Google Scholar More articles by this author https://doi.org/10.31635/ccschem.021.202101047 SectionsSupplemental MaterialAboutAbstractPDF ToolsAdd to favoritesTrack Citations ShareFacebookTwitterLinked InEmail It remained challenging to fabricate Tb3+-doped lanthanide nanocrystals (NCs) to simultaneously acquire strong energy migration upconversion (EMU) emissions of Tb3+ while suppressing the Tm3+ UV upconversion emissions that cause background biofluorescence issues in bioapplications based on Tb3+-doped EMU NCs. Herein, we report a novel sandwich structured [email protected]@shell scheme for the design of EMU NCs, for example, NaLuF4∶Yb/[email protected]4∶[email protected]4∶Tb NCs, wherein Yb3+, Tm3+, and Tb3+ are incorporated separately into the inner core, middle shell, and outer shell, respectively. We found that in the sandwich structured NCs, the effective inter-shell energy transfer from Gd3+ in the middle shell to Tb3+ in the outer shell accelerated the Yb3+–Tm3+ five-photon upconversion and the subsequent Tm3+ to Gd3+ energy transfer processes, which could eventually lead to almost complete inhibition of Tm3+ UV upconversion emissions, concurrent with the strong EMU emissions of Tb3+. Our findings might stimulate new concepts for manipulating upconversion emissions of lanthanide NCs. Download figure Download PowerPoint Introduction Lanthanide-doped upconversion nanocrystals (UCNCs) have become ideal candidates for applications in biological labeling and imaging fields owing to their attractions such as sharp emission peaks, long luminescence lifetimes, high photostability, low toxicity, deep light penetration depth, and others.1–4 Notably, the most common activators applied to achieve efficient upconversion luminescence (UCL) are restricted to Er3+, Tm3+, and Ho3+ due to their ladder-like energy levels that ensure effective sensitization by Yb3+ ions.5 Nonetheless, a series of lanthanide ions (e.g., Tb3+, Eu3+) without long-lived intermediate states can also be utilized as activators to accomplish their own intrinsic UCL through the energy migration upconversion (EMU) mechanism, which was first realized in lanthanide-doped NCs by Wang et al.6 In 2011, Wang’s group6 designed a [email protected] NaGdF4@NaGdF4 NC with the sensitizer/accumulator (Yb3+/Tm3+) and activator X3+ (X3+ = Tb3+, Eu3+, Dy3+, or Sm3+) ions separately incorporated into the core and shell of NC, respectively. In the EMU NC, the near-infrared (NIR) excitation energy was accumulated in the core by the Yb3+–Tm3+ five-photon upconversion process, followed by energy transfer from Tm3+ (1I6) to Gd3+ (6P7/2).7 Then the energy hopped randomly between Gd3+ ions, and finally was captured by the activator X3+ doped in the shell to achieve the EMU emissions.8–11 Among the aforementioned lanthanide activators, Tb3+ ion possesses a relatively higher quantum efficiency of EMU emissions, mainly due to the large energy gap (∼16,000 cm−1) between the first excited state (5D4) and the highest ground state (7F0) of Tb3+ that generates very weak phonon relaxation in the 5D4 excited state.12,13 In other words, the 5D4 → 7FJ transitions of Tb3+ are resistant to the deleterious high-frequency vibration modes of molecules or ligands in complex biological media that promote nonradiative phonon relaxation.14,15 In fact, benefiting from such a unique characteristic of Tb3+, the luminescence intensity or lifetime of 5D4 → 7FJ transitions of Tb3+ can work as a stable and reliable detection signal to guarantee a high accuracy of biological labeling and imaging, employing the Tb3+-doped NCs as nano-bioprobes or Tb3+ complex as a molecular probe.16,17 For instance, the sensitive and high-accuracy quantification of lysosomal nitric oxide (NO) with both ratiometric and luminescence lifetime detection modes was successfully established by Dai et al.18 based on intramolecular luminescence resonance energy transfer (LRET) from the 5D4 excited state of luminescent Tb3+ complex to the first excited singlet state of rhodamine. Nevertheless, the UV excitation of the Tb3+ complex inevitably causes issues of poor tissue penetration of the excitation light, as well as background biofluorescence from endogenous chromophores and proteins during in vivo biodetection.19–26 Unfortunately, even if substituting the luminescent Tb3+ complex with the EMU NCs (e.g., NaGdF4∶Yb/[email protected]4∶Tb), excitable by 980 nm NIR light can help in the light penetration depth,27–30 the issue of background biofluorescence, induced by the UV upconversion emissions of Tm3+ would still exist. This is because the Tm3+ UV emissions of considerable magnitude always coexist with the 5D4 → 7FJ emissions of Tb3+, as reported previously for Tb3+-doped EMU NCs.31 As such, in vivo ratiometric and luminescence lifetime detection/imaging approaches based on LRET from 5D4 of Tb3+ doped in EMU NCs to a custom-designed energy acceptor would be confronted with the Tm3+ UV emissions that trigger an issue of background biofluorescence. Likewise, the same dilemma might be encountered in the EMU NCs-mediated super-resolution imaging of cytoskeleton protein or organelle of cells based on the stimulated emission depletion (STED) nanoscopy technique.32 The Tm3+ UV emissions could compromise the spatial resolution of STED nanoscopy by imposing background biofluorescence in proximity to the EMU NCs. Herein, in this context, we proposed a sandwich structured33–36 [email protected]@shell scheme to design EMU NCs. As a proof-of-concept experiment, we synthesized the sandwich structured NaLuF4∶Yb/[email protected]4∶[email protected]4∶Tb NCs with a sensitizer (Yb3+), an accumulator (Tm3+), and an activator (Tb3+), separately incorporated into the inner core, middle shell, and outer shell of NC, respectively. The analyses, based on steady-state and transient photoluminescence (PL) spectroscopy, revealed that the Yb3+–Tm3+ five-photon upconversion and subsequent Tm3+ to Gd3+ energy transfer processes could be enhanced by the effective Gd3+ → Tb3+ inter-shell energy transfer. Such an enhancing effect led to almost complete suppression of Tm3+ UV upconversion emissions, with respect to the strong EMU emissions of Tb3+ (Figure 1) under 980 nm excitation at a power density above the threshold of several tens of W/cm2. Figure 1 | Schematic illustration of the sandwich structured [email protected]@shell design strategy of EMU NCs for acquiring nearly completely inhibited Tm3+ UV upconversion emissions in parallel with the strong EMU emissions of Tb3+. Download figure Download PowerPoint Experimental Methods Chemicals and materials Ln2O3 (Ln = Yb, Tm, Gd, Tb, and Lu) (99.99%), trifluoroacetic acid, cyclohexane, and ethanol were purchased from Sinopharm Chemical Reagent Co. (Shanghai, China). Oleic acid (OA), oleylamine (OM), and 1-ocatedecence (ODE) were purchased from Sigma-Aldrich (Shanghai, China). All the chemical reagents were used as received without further purification. Ln(CF3COO)3 was prepared as reported in the literature.37 Synthesis of core NaLuF4∶Yb/Gd NCs Monodisperse NaLuF4∶Yb/Gd was synthesized via a modified thermal decomposition process.38 In a typical reaction, Ln(CF3COO)3 (0.5 mmol, 33% mol of Gd, 49% mol of Yb, 18% mol of Lu) and Na(CF3COO) (0.75 mmol) were added to a 50 mL flask containing OA (5.9356 g), OM (2.9627 g), and ODE (3.0197 g) at room temperature (RT). Then the mixture was heated at 120 °C for 15 min under magnetic stirring in a N2 atmosphere to dissolve the trifluoroacetate precursors and remove residual water and oxygen. Subsequently, the resulting transparent solution was heated to 310 °C under N2 flow with vigorous stirring for 60 min and then cooled to RT naturally. The resulting NaLuF4∶Yb/Gd NCs were precipitated by the addition of ethanol, collected via centrifugation (12,000 rpm, 5 min), washed twice with ethanol, and redispersed in cyclohexane. Synthesis of [email protected] NaLuF4∶Yb/[email protected]4∶Tm NCs The [email protected] NCs were fabricated using the seed-mediated method. In a typical synthetic procedure, NaLuF4∶Yb/Gd NCs were added to a 50 mL flask with OA (6.2954 g) and ODE (5.5626 g). The solution was stirred at 90 °C in a N2 atmosphere to remove cyclohexane. Then the reaction mixture was cooled to RT. Thereafter, Ln(CF3COO)3 (0.425 mmol, 98% mol of Gd, 2% mol of Tm) and Na(CF3COO) (0.425 mmol) were added, followed by heating the mixture at 120 °C for 15 min under magnetic stirring in N2 atmosphere to dissolve the trifluoroacetate precursors and remove residual water and oxygen. Subsequently, the resulting transparent solution was heated to 300 °C under N2 flow with vigorous stirring for 60 min and then cooled to RT naturally. The resulting NaLuF4∶Yb/[email protected]4∶Tm NCs were precipitated by the addition of ethanol, collected via centrifugation (vide supra), as indicated above, washed twice with ethanol, and redispersed in cyclohexane. Synthesis of [email protected]@shell NaLuF4∶Yb/[email protected]4∶[email protected]4∶Tb NCs The [email protected]@shell NCs were fabricated by using the presynthesized [email protected] NCs as seeds. In a typical synthesis process, NaLuF4∶Yb/[email protected]4∶Tm NCs were added to a 50 mL flask with OA (6.2954 g) and ODE (5.5626 g). The solution was stirred at 90 °C in a N2 atmosphere to remove cyclohexane. Then the reaction mixture was cooled to RT. Thereafter, Ln(CF3COO)3 (0.5 mmol, 85% mol of Lu, 15% mol of Tb) and Na(CF3COO) (0.5 mmol) were added; the mixture was heated at 120 °C for 15 min under magnetic stirring in N2 atmosphere to dissolve the trifluoroacetate precursors and remove residual water and oxygen. Subsequently, the resulting transparent solution was heated to 300 °C under N2 flow with vigorous stirring for 60 min and then cooled to RT naturally. The resulting NaLuF4∶Yb/[email protected]4∶[email protected]4∶Tb NCs were precipitated by the addition of ethanol, collected via centrifugation, as indicated earlier, washed twice with ethanol, and redispersed in cyclohexane. The synthetic procedures for [email protected]@shell NaLuF4∶Yb/Gd/[email protected]4@NaLuF4∶Tb and NaLuF4∶Yb/[email protected]4∶[email protected]4 NCs were identical to that for NaLuF4∶Yb/[email protected]4∶[email protected]4∶Tb NCs, except for Tm(CF3COO)3 was added as a core precursor for NaLuF4∶Yb/Gd/[email protected]4@NaLuF4∶Tb NCs; Tb(CF3COO)3 was not added as a shell precursor for NaLuF4∶Yb/[email protected]4∶[email protected]4 NCs. Characterization Powder X-ray diffraction (XRD) patterns of the samples were collected on an X-ray diffractometer (MiniFlex 600; Rigaku, Tokyo, Japan) with Cu Kα1 radiation (λ = 0.154187 nm). Transmission electron microscopy (TEM) measurements, including high-angle annular dark-field scanning TEM (HAADF-STEM) image and energy-dispersive X-ray spectroscopy (EDS) element mapping, were performed on an FEI Tecnai F20 TEM (Hillsboro, OR). Upconversion PL spectra were collected under 980 nm laser excitation provided by the corresponding continuous-wave laser diode. PL decays were measured with a customized UV to mid-infrared steady-state and phosphorescence lifetime spectrometer (FSP920-C; Edinburgh Instruments Ltd., Livingston, UK) equipped with a digital oscilloscope (TDS3052B; Tektronix, Beaverton, OR) and a tunable mid-band optical parametric oscillator (OPO) pulse laser as the excitation source [410–2400 nm, 10 Hz, and a pulse width of ∼5 ns (Vibrant 355II; OPOTEK, Carlsbad, CA)]. The absolute UCL quantum yield (QY) of sandwich structured NaLuF4∶Yb/[email protected]4∶[email protected]4 NCs was measured with a customized UCL spectroscopy system (FLS920; Edinburgh Instruments Ltd., Livingston, UK) at RT upon a 980 nm diode laser excitation at a power density of ∼120 W/cm2. All the spectral data were corrected according to the spectral response of the spectrometer. Results and Discussion The sandwich structured [email protected]@shell NaLuF4∶Yb/[email protected]4∶[email protected]4∶Tb EMU NCs were synthesized through a seed-mediated method,39–41 which involves the growth of NaLuF4∶Yb/Gd(49/33%) core NCs via the thermo-decomposition of lanthanide trifluoroacetates,42,43 followed by the successive epitaxial deposition of NaGdF4∶Tm(2%) and NaLuF4∶Tb(15%) shell layers on the core NCs serving as seeds. The typical low-magnification TEM images of NCs showed that the core NaLuF4∶Yb/Gd, [email protected] NaLuF4∶Yb/[email protected]4∶Tm, and [email protected]@shell NaLuF4∶Yb/[email protected]4∶[email protected]4∶Tb NCs were almost spherical with average diameters of ∼18.3 ± 1.6, 22.8 ± 1.9, and 25.4 ± 2.4 nm, respectively (Figures 2a–2c), as corroborated by the incrementally narrowed powder XRD peaks, which were well-indexed as the hexagonal phase44–46 of NaLuF4 ( Supporting Information Figure S1). The single crystalline property with hexagonal structure was further validated using high-resolution TEM (HRTEM) observation47–49 of an individual [email protected]@shell EMU NC and the corresponding selected area electron diffraction (SAED) pattern (Figures 2d and 2e).50,51 The HAADF-STEM image in Figure 2f confirmed visually, the inner core area and outer shell layers of EMU NCs, and the EDS mappings conducted on several randomly selected EMU NCs (Figures 2f and Supporting Information Figure S2) verified the distribution of Yb in the core, Gd, Tm in the middle shell, and Tb in the outer shell.52–55 Figure 2 | TEM images of (a) core NaLuF4∶Yb/Gd NCs, (b) [email protected] NaLuF4∶Yb/[email protected]4∶Tm NCs, (c) [email protected]@shelll NaLuF4∶Yb/[email protected]4∶[email protected] NaLuF4∶Tb NCs. The scale bar is 50 nm. (d) HRTEM image of a single [email protected]@shell NC and its corresponding fast Fourier transform (FFT) pattern (inset). (e) SAED pattern of NaLuF4∶Yb/[email protected]4∶[email protected]4∶Tb NCs assemblies. (f) HAADF-STEM image and corresponding EDS element mappings of Yb3+, Gd3+, and Tb3+ ions for several randomly selected [email protected]@shell NCs. Download figure Download PowerPoint Figure 3a shows the characteristic EMU emissions of Tb3+ doped in the sandwich structured NaLuF4∶Yb/[email protected]4∶[email protected]4∶Tb NCs deposited on the quartz glass substrate with film thickness in the order of 0.1 mm, upon 980 nm excitation at a power density of ∼200 W/cm2. The strong emission bands of Tb3+ recorded at 489, 542, 584, and 619 nm are from its optical transitions of 5D4 → 7F6, 7F5, 7F4, 7F3, respectively56 [the absolute QY of Tb3+ EMU emissions ≈ 0.4 ± 0.01%]. Concomitantly, weak UV upconversion emissions of Gd3+ (6P7/2 → 8S7/2) and Tm3+ (1I6 → 3H6, 1I6 → 3F4, and 1D2 → 3H6) were observed at 314, 289, 344, and 361 nm, respectively.6,57 For the radiative transitions from 1G4, 1D2, 1I6 of Tm3+ and 5D4 of Tb3+ in the sandwich structured NCs, the smaller double-log intensity-power slopes in the high-power regime than in the low-power regime were observed ( Supporting Information Figure S3). Typically, such behavior is ascribed to the absorption saturation of lanthanide ions. Owing to the unique ladder-like energy levels of Tm3+, a high excitation power density promoted the population of higher energy levels of Tm3+ (from 1G4 to 1D2 to 1I6), affiliating the energy transfer from the Tm3+1I6 to Gd3+6P7/2, to Tb3+5D0 levels. As a result, the Tb3+/Tm3+ intensity ratio should be dependent on the excitation power density. Particularly, the intensity ratio between the strongest emission band of Tb3+ (5D4 → 7F5) and the maximum UV emission of Tm3+ (1D2 → 3H6) was calculated to be ∼58.9 at an excitation power density of ∼200 W/cm2 (Figure 3a); even if the power density was decreased to the level of ∼70 W/cm2, the intensity ratio was still as large as ∼25.0 ( Supporting Information Figure S4). Obviously, the Tm3+ UV upconversion emissions could be inhibited to a negligible magnitude with respect to the strong EMU emissions of Tb3+ in the sandwich structured NCs when the excitation power density was above a threshold of several tens of W/cm2. From this aspect, it is expected that by using such sandwich structured EMU NCs as nano-bioprobes for the cellular biological detection/imaging or super-resolution imaging of cells,18,32 the background biofluorescence issue, triggered by Tm3+ UV emissions would be circumvented. Subsequently, in a control experiment, we synthesized a series of non-sandwich structures, NaLuF4∶Yb/Gd/Tm(49/33/x%, x = 0.5, 1, 2)@NaGdF4@NaLuF4∶Tb(15%) with Yb3+ and Tm3+ co-doped into the inner core of NC ( Supporting Information Figure S5). Upon 980 nm excitation at a power density of ∼200 W/cm2, Tm3+ UV emissions of considerable magnitude, concomitant with the intense 5D4 → 7FJ (J = 3, 4, 5, 6) emissions of Tb3+, were observed in the non-sandwich structured NCs (Figure 3b), resembling the emission pattern of Tb3+-doped EMU NCs reported previously ( Supporting Information Table S1).6,8,10 The calculated intensity ratios between the strongest emission band of Tb3+ (5D4 → 7F5) and the maximum UV emission of Tm3+ (1D2 → 3H6 or 1I6 → 3F4, depending on specific circumstance) were ∼0.2, 1.1, and 0.6 for the non-sandwich structured NCs doped with 0.5%, 1%, and 2% Tm3+, respectively (Figure 3b), contrasting the large value of ∼58.9 observed in the sandwich structured NCs at an identical excitation power density. Meanwhile, as listed in Supporting Information Table S2, the QY of Tb3+ EMU emissions in the sandwich structured NCs was even larger than that in the non-sandwich structured NCs. Such a remarkable difference in emission property between the sandwich and non-sandwich structured NCs could well be explained by comparing their respective microscopic EMU processes. Figure 3 | UCL spectra of (a) sandwich structured NaLuF4∶Yb/[email protected]4∶[email protected]4∶Tb NCs and (b) non-sandwich structured NaLuF4∶Yb/Gd/Tm (49/33/x%, x = 0.5, 1, 2)@NaGdF4@NaLuF4∶Tb(15%) NCs. Inset is the photograph of a thin powder sample deposited on a quartz glass substrate upon 980 nm excitation at a power density of ∼200 W/cm2. (c) Left panel: energy level diagrams of Yb3+, Tm3+, Gd3+, and Tb3+ ions, and the proposed EMU process of Tb3+ ion. Right panel: proposed energy transfer pathways in the core area and shell layers of non-sandwich structured (upper) and sandwich structured (lower) NCs accounting for the EMU process of Tb3+. (d) UCL decays from 6P7/2 of Gd3+ in non-sandwich structured (denoted as Yb3+/Tm3+) NCs doped with 1% Tm3+ and sandwich structured (denoted as Yb3+@Tm3+) NCs, fitted by using double- and single-exponential functions, respectively: I = A1exp[−(t − t0)/τ1] + A2exp[−(t − t0)/τ2] + I0; I = Aexp[−(t − t0)/τ] + I0. Download figure Download PowerPoint Specifically, for the non-sandwich structured NaLuF4∶Yb/Gd/[email protected]4@NaLuF4∶Tb NCs, the excitation energy accumulated by the Yb3+-Tm3+ five-photon upconversion process initially transferred from Tm3+ (1I6) to Gd3+ (6P7/2) within the inner core. Afterward, the energy randomly hopped in a circuitous way between Gd3+ ions accommodated both in the inner core and middle shell. Finally, the energy transferred directly from Gd3+ in the middle shell to Tb3+ in the outer shell (i.e., inner-shell energy transfer) (left panel of Figure 3c). Here, it should be noted that the random energy hopping between Gd3+ ions could also transfer energy to the quenching defects or trap states58–62 in the crystal lattice, or transfer energy back to Tm3+ ions, which was deleterious to the EMU process of Tb3+ (upper right panel of Figure 3c). By contrast, for the sandwich structured NaLuF4∶Yb/[email protected]4∶[email protected]4∶Tb NCs, although the energy hopping between Gd3+ ions in the middle shell still existed, the circuitous random hopping in the inner core was alleviated effectively, compared with the non-sandwich structured NCs (lower right panel Figure 3c). Meanwhile, for the energy hopping between Gd3+ ions in the confined middle shell (∼ 2.3 nm), the length of the average hopping path was markedly constrained because the energy could be preferentially transferred from Gd3+ to neighboring Tb3+ in the thin outer shell (∼1.3 nm) through the direct inter-shell energy transfer.10 Consequently, as evidenced in Figure 3d, Gd3+ in the non-sandwich structured NCs presents a double-exponential trend with a fast decay (τ1 = 1.07 ms), followed by a long-lived component (τ2 = 4.60 ms). The long-lived component derived from the circuitous random energy hopping between Gd3+ ions because the energy hopping decreased the nonradiative rates related to defect-induced quenching and back energy transfer from Gd3+ to Tm3+. In fact, the measured lifetime of the long-lived component related to the random energy hopping between Gd3+ ions was shorter than, but could asymptotically approach the lifetime of Gd3+ ions under ideal conditions, where the only radiative process of Gd3+ ions and ion–ion interactions existed. Unlike the non-sandwich structured NCs, Gd3+ in the sandwich structured NCs exhibited a simple single-exponential decay with a lifetime of ∼0.72 ms, probably due to the effectively alleviated random energy hopping. The alleviated random energy hopping in sandwich structured NCs was further validated by the observation (vide infra) that, Gd3+ in the NaLuF4∶Yb/[email protected]4∶[email protected]4 NCs (with no Tb3+ doping in the outer shell) also presented a double-exponential trend similar to the non-sandwich structured NCs, due to the absence of the above-mentioned Gd3+ → Tb3+ inter-shell energy transfer, beneficial to inhibition of energy hopping. Altogether, the depressed random energy hopping between Gd3+ ions combined with the effective Gd3+ → Tb3+ inter-shell energy transfer was responsible for the more remarkable EMU emissions of Tb3+ in the sandwich structured NCs relative to the non-sandwich structured counterparts (Figures 3a and 3b and Supporting Information Table S1). To further investigate the detailed EMU process of Tb3+ in the sandwich structured NaLuF4∶Yb/[email protected]4∶[email protected]4∶Tb NCs, we elaborately synthesized a blank sample, that is, the NaLuF4∶Yb/[email protected]4∶[email protected]4 NCs with no Tb3+ doping in the outer shell, under otherwise identical thermodynamic conditions to those of sandwich structured EMU NCs. As validated by TEM observation, the blank sample maintained the same crystal morphology and structures as the sandwich structured NCs (Figures 4a and 4b and Supporting Information Figure S6). Consequently, in the following experiments, we could unveil the variation in steady-state populations63 of lanthanide excited states before and after the occurrence of Gd3+ → Tb3+ inter-shell energy transfer, respectively, by comparing their corresponding UCL intensities of lanthanide ions. To quantify the variations in steady-state populations of lanthanide excited states accurately, the UCL intensities of colloidal solutions of blank sample and the sandwich structured NCs that were well-dispersed in cyclohexane at the same concentration (∼1.3 mg/mL) (Figures 4c and 4d) were recorded under identical excitation conditions with effective excitation power density in the order of several W/cm2 (upper panel of Supporting Information Figure S7 and Figure 4e). Here, the NCs colloidal solution sample placed in a quartz cuvette has great superiority over the thin film sample fabricated by depositing NCs on a quartz glass substrate to analyze variations in the steady-state populations of lanthanide excited states. This is because the collected emission intensity of the thin-film sample was dependent on the stacking density of NCs, while the film sample showed poor repeatability of NCs stacking density. Additionally, it should be mentioned that the effective excitation power density on sandwich structured NaLuF4∶Yb/[email protected]4∶[email protected]4∶Tb NCs differed markedly between the thin film and colloidal solution samples (∼ 200 vs. several W/cm2) under different experimental setups for the acquisition of UCL spectra ( Supporting Information Figure S7), thus, resulting in the different UCL patterns, as shown in Figures 3a and 4e. Figure 4 | TEM images of (a) NaLuF4∶Yb/[email protected]4∶[email protected]4 NCs (blank sample, denoted as without Tb3+) and (b) sandwich structured NaLuF4∶Yb/[email protected] NaGdF4∶[email protected]4∶Tb NCs; denoted as with Tb3+). The scale bar is 200 nm. Insets of (a and b) are corresponding fast Fourier transform (FFT) patterns. Photographs of colloidal solutions of (c) sandwich structured and (d) blank sample NCs, well-dispersed in cyclohexane under 980 nm excitation. (e) Comparison of steady-state UCL spectra of blank sample and sandwich structured NCs colloidal solutions. (f) Simplified model depicting the proposed EMU process of Tb3+. Download figure Download PowerPoint As compared in Figure 4e, decreases in the steady-state populations of excited states 6P7/2 of Gd3+, 1I6, 1D2, and 3H4 of Tm3+, along with the increase in that of 1G4 of Tm3+, appeared in the sandwich structured NCs relative to the blank sample. The decrease in the population of Gd3+6P7/2 excited state was consistent with the

Higher-voltage asymmetric-electrolyte metal-air batteries

Asymmetric-electrolyte metal-air batteries (AMABs) with high operating voltage and energy density are highly appealing, but some challenges remain. In the April 7, 2021 issue of Matter, Jung-Ho Lee and co-workers first reported a novel separator and an atomically dispersed catalyst to achieve high voltage and long stability of AMABs. Asymmetric-electrolyte metal-air batteries (AMABs) with high operating voltage and energy density are highly appealing, but some challenges remain. In the April 7, 2021 issue of Matter, Jung-Ho Lee and co-workers first reported a novel separator and an atomically dispersed catalyst to achieve high voltage and long stability of AMABs. The increasing concerns about the increasing energy demand and environmental crisis have led to an extensive search for a new and efficient energy conversion device. Among the developed energy storage devices, aqueous metal-air batteries (MABs) are particularly promising due to high theoretical energy density, high safety, eco-friendliness, as well as high tolerance to moisture. They show considerable potential in portable and large-scale power supply and storage applications.1Liu Q. Pan Z. Wang E. An L. Sun G. Aqueous metal-air batteries: Fundamentals and applications.Energy Stor. Mater. 2020; 27: 478-505Crossref Scopus (86) Google Scholar,2Wang H.-F. Xu Q. Materials Design for Rechargeable Metal-Air Batteries.Matter. 2019; 1: 565-595Abstract Full Text Full Text PDF Scopus (188) Google Scholar However, the major obstacles to the large-scale implementation of MABs are CO2 poisoning issues and sluggish air electrode reactions, namely oxygen reduction reaction (ORR) and oxygen evolution reaction (OER). In view of the above tough issues, Yu and Manthiram3Yu X. Manthiram A. Electrochemical Energy Storage with Mediator-Ion Solid Electrolytes.Joule. 2017; 1: 453-462Abstract Full Text Full Text PDF Scopus (17) Google Scholar first proposed the concept of asymmetric-electrolyte MABs (AMABs) in 2017, which fundamentally solved the problem of continuous CO2 ingression on the side of air electrode. More importantly, the higher ORR potential (1.23 V versus SHE) delivered in the cathodic acidic electrolyte (e.g., H2SO4) of AMABs increases the output voltage and power output of practical devices. Although significant progress has been made in improving the electrochemical performance and long-term stability of AMABs in recent years,4Yu J. Li B.-Q. Zhao C.-X. Liu J.-N. Zhang Q. Asymmetric Air Cathode Design for Enhanced Interfacial Electrocatalytic Reactions in High-Performance Zinc-Air Batteries.Adv. Mater. 2020; 32: e1908488Crossref PubMed Scopus (50) Google Scholar, 5Li L. Manthiram A. Long-Life, High-Voltage Acidic Zn-Air Batteries.Adv. Energy Mater. 2016; 6: 1502054Crossref Scopus (77) Google Scholar, 6Yu X. Gross M.M. Wang S. Manthiram A. Aqueous Electrochemical Energy Storage with a Mediator-Ion Solid Electrolyte.Adv. Energy Mater. 2017; 7: 1602454Crossref Scopus (18) Google Scholar the emerging AMABs still face several key obstacles and drawbacks that still need to be overcome. The first obstacle is the high demand for cost-effective separator: the currently used bipolar polymer films and ceramic solid-state membrane show ultrahigh cost ($3,200 m−2 for bipolar polymer membrane and $6,000/kg for ceramic solid-state membrane),7Cai P. Li Y. Chen J. Jia J. Wang G. Wen Z. An Asymmetric-Electrolyte Zn-Air Battery with Ultrahigh Power Density and Energy Density.ChemElectroChem. 2018; 5: 589Crossref Scopus (32) Google Scholar limited ionic conductivity, and the unavoidable crossing between H+ and OH−. Another obstacle is the lack of low-cost, highly efficient, and stable bifunctional ORR/OER catalysts in the acidic electrolyte.8Zhang P. Zhao Y. Zhang X. Functional and stability orientation synthesis of materials and structures in aprotic Li-O2 batteries.Chem. Soc. Rev. 2018; 47: 2921-3004Crossref PubMed Google Scholar To address the above bottleneck issues, in the April 7, 2021 issue of Matter, the new materials design reported by Lee and coworkers,9Lin C. Kim S.-H. Xu Q. Kim D.-H. Ali G. Shinde S.S. Yang S. Yang Y. Li X. Jiang Z. et al.High-voltage asymmetric metal-air batteries based on polymeric single-Zn2+-ion conductor.Matter. 2021; 4: 1287-1304Abstract Full Text Full Text PDF Scopus (8) Google Scholar includes two important breakthroughs: the novel separator fabrication and atomically dispersed catalyst design to achieve the high-voltage and long-term stability AMABs. The first, it is the first reported the cost-effective polyacrylonitrile (PAN)-based single-Zn2+-ion conductor to selectively transport Zn2+ ions and inhibit the crossover of H+ and OH– ions. Specifically, Zn2+ ion-selective transport membrane (ZnSTM) (Figure 1A) was formed from a mixture of Zn(OAc)2 and PAN heated at 140°C in the air by cross-linking polymerization between adjacent molecules. As-obtained ZnSTM possesses good mechanical stability and flexibility, and its tensile strength is 0.8 Ma. X-ray diffraction (XRD), Fourier transform infrared spectroscopy and X-ray photoelectron spectroscopy clearly reveal that the typical Zn2+ anchoring ladder structure in ZnSTM is formed by a cyclization reaction, which transforms the C≡N in the PAN into C=N (Figure 1B). The microstructure of ZnSTM was further investigated using transmission electron microscopy (TEM) (Figure 1C). The microphase separation structure presented in ZnSTM consists of a hydrophobic polymer backbone (bright region) and a hydrophilic domain (dark region) generated by Zn2+ ions aggregated around the N-related groups at the edge of the carbon backbone. The hydrophilic domains with a Zn2+ ionic conductivity of 1.9 × 10−4 S cm−1 at 313 K provided a fast transport channel for Zn2+ ions. The Arrhenius plot shows that the calculated activation energy of ZnSTM is about 0.2 eV (<0.4 eV), indicating that the migration of Zn2+ in ZnSTM follows the Grotthuss jump mechanism (Figure 1D). According to the Bruce-Vincent-Evans equation, the migration number of Zn2+ ion is determined to be 0.8, suggesting that the ion conduction mainly comes from the mobility of the cation (Zn2+) in ZnSTM. In addition, the self-made diffusion cell experiment confirmed that ZnSTM can effectively inhibit the cross of H+ and OH−. Computational calculation reveals that the attractive force between Zn2+ sites and guest ions (H +/OH− ions) inhibits the transport of H + and OH− ions. On the contrary, the repulsive force between Zn2+ binding sites and guest Zn2+ ions is conducive to the migration of Zn2+ ions.Figure 1Structure and properties of ZnSTM, Co/NS, and the assembled AZnAB based on ZnSTM separator and Co/NS catalystShow full caption(A) Optical images of ZnSTM.(B) XRD pattern of ZnSTM. Inset shows its chemical structure.(C) TEM image and corresponding SAED pattern (inset) of ZnSTM.(D) Arrhenius plots of ZnSTM.(E) ORR polarization curves of NS, Co/NS, Co-NPs/NS, and Pt/C in O2-saturated 0.1 M HClO4.(F) OER polarization curves of NS and Co/NS in O2-saturated 0.1 M HClO4 with rotating rate of 1,600 rpm.(G) Comparison of in situ EXAFS obtained on Co/NS before and after in situ exposure to periodic changed ORR and OER potentials.(H) Schematic of the theoretical voltage and configuration of aqueous AZnAB.(I) Comparison of energy densities of the AZnAB and previously reported ZnAB at various current densities.View Large Image Figure ViewerDownload Hi-res image Download (PPT) (A) Optical images of ZnSTM. (B) XRD pattern of ZnSTM. Inset shows its chemical structure. (C) TEM image and corresponding SAED pattern (inset) of ZnSTM. (D) Arrhenius plots of ZnSTM. (E) ORR polarization curves of NS, Co/NS, Co-NPs/NS, and Pt/C in O2-saturated 0.1 M HClO4. (F) OER polarization curves of NS and Co/NS in O2-saturated 0.1 M HClO4 with rotating rate of 1,600 rpm. (G) Comparison of in situ EXAFS obtained on Co/NS before and after in situ exposure to periodic changed ORR and OER potentials. (H) Schematic of the theoretical voltage and configuration of aqueous AZnAB. (I) Comparison of energy densities of the AZnAB and previously reported ZnAB at various current densities. The second, an atomically dispersed Co electrocatalyst supported on nitrogen-doped carbon nanosheet (Co/NS) was developed as the cathodic catalyst for AMABs, which has high activity and stability for both ORR and OER in acidic medium. In this work, Co/NS was prepared by pyrolysis of imidazolate frameworks of bimetallic ZnCo zeolitic encapsulated with sodium chloride crystals ([email protected]). The coexistence of nanocluster with an average diameter of ~1 nm, lattice fringes of 0.2 nm, and single atoms in the synthesized Co/NS sample were confirmed by the aberration-corrected atomic-resolution high-angle annular dark-field scanning TEM (AC-HAADF-STEM) image. In line with the AC-HAADF-STEM, X-ray absorption spectrum also revealed co-existence of the nanoclusters with Co-Co coordination number of 2 and Co species of Co-Nx (x = 1–4) in the Co/NS sample. Through the test of catalytic performance, the atomically dispersed Co/NS catalyst shows excellent ORR and OER performance in O2-saturated 0.1 M HClO4. The results showed that the half-wave potential of ORR of as-prepared Co/NS was 0.71 V, and the Tafel slope was 82.1 mV dec−1 (Figure 1E). The overpotential of OER for Co/NS at 10 mAcm–2 was 605 mV after 2,000 consecutive cyclic voltammetry cycles, and no active degradation was observed after 20 h of electrocatalysis at 10 mAcm−2 (Figure 1F). In situ extended X-ray absorption fine structure (EXAFS) data (Figure 1G) confirmed that the bond length of Co–N varies significantly between ORR and OER potentials, and periodically decreases first and then increases, suggesting that the active sites of ORR and OER are Co–Nx species. To demonstrate the practical potential, the asymmetric Zn-air battery (AZnAB) was assembled on the basis of the prepared novel ZnSTM separator and atomically dispersed Co/NS catalysts (Figure 1H). Remarkably, the AZnAB delivered an operating voltage of 1.96 V, an ultrahigh specific energy density of 1,354 Wh kgZn−1, and long-term stability (~100 h) with a round-trip efficiency of 76.6%. The specific energy density of 1,354 Wh kgZn−1 is higher than those of recently reported conventional ZnABs (Figure 1I) and even approaches the theoretical value of the alkaline electrolyte-based ZnABs. And after 300 cycles over 100 h, an increased round-trip efficiency of 70.8% can be obtained, which is comparable with the state-of-the-art Pt/C+RuO2-based AZnAB (75.1%). Moreover, the smart combination strategy is also applicable to the other type AMABs, including ASiAB and ASnAB. Impressively, both ASiAB and ASnAB have demonstrated enhanced operating voltage, power density, and excellent rate performance over the recently reported SiABs and SnABs. This work not only opens up a new way for the design of new separators that can selectively transport heavy multivalent ions in AMABs but also makes an important step for the practical application of AMABs. However, AMABs still have room for improvement in the formation of dendrites in the anode in alkaline electrolytes, the internal resistance of the cathode, and the material science and engineering of the electrode. In addition, this work provides a general strategy for AMABs, which is also applicable to the other-type assembled asymmetric MABs (metals, Si, Sn), and has guiding significance for the design of next generation metal-air batteries with large current and long-term energy storage applications. High-voltage asymmetric metal–air batteries based on polymeric single-Zn2+-ion conductorLin et al.MatterJanuary 29, 2021In BriefThe development of asymmetric-electrolyte metal–air batteries (AMABs) has been hindered by lack of polymer separator that can selectively transport metal ions and cost-effective bifunctional electrocatalyst. Here, we report breakthroughs in both key components by constructing a polyacrylonitrile membrane that can selectively transport Zn2+ ions, and an atomically dispersed Co electrocatalyst that can stably catalyze ORR and OER in acid. The assembled AMABs (metals Zn, Si, Sn) showed stable battery performance with enhanced operating voltage, power densities, and excellent rate performance that surpass those of recently reported MABs. Full-Text PDF Open Archive

High-Performance Ru 2 P Anodic Catalyst for Alkaline Polymer Electrolyte Fuel Cells

Open AccessCCS ChemistryRESEARCH ARTICLE1 May 2022High-Performance Ru2P Anodic Catalyst for Alkaline Polymer Electrolyte Fuel Cells Yuanmeng Zhao†, Fulin Yang†, Wei Zhang†, Qihao Li†, Xuewei Wang, Lixin Su, Xuemei Hu, Yan Wang, Zizhun Wang, Lin Zhuang, Shengli Chen and Wei Luo Yuanmeng Zhao† Hubei Electrochemical Power Sources Key Laboratory, College of Chemistry and Molecular Sciences, Wuhan University, Wuhan, Hubei 430072 †Y. Zhao, F. Yang, W. Zhang, and Q. Li contributed equally to this work.Google Scholar More articles by this author , Fulin Yang† Hubei Electrochemical Power Sources Key Laboratory, College of Chemistry and Molecular Sciences, Wuhan University, Wuhan, Hubei 430072 †Y. Zhao, F. Yang, W. Zhang, and Q. Li contributed equally to this work.Google Scholar More articles by this author , Wei Zhang† Key Laboratory of Automobile Materials MOE, School of Materials Science & Engineering, Electron Microscopy Center, and International Center of Future Science, Jilin University, Changchun 130012 †Y. Zhao, F. Yang, W. Zhang, and Q. Li contributed equally to this work.Google Scholar More articles by this author , Qihao Li† Hubei Electrochemical Power Sources Key Laboratory, College of Chemistry and Molecular Sciences, Wuhan University, Wuhan, Hubei 430072 †Y. Zhao, F. Yang, W. Zhang, and Q. Li contributed equally to this work.Google Scholar More articles by this author , Xuewei Wang Hubei Electrochemical Power Sources Key Laboratory, College of Chemistry and Molecular Sciences, Wuhan University, Wuhan, Hubei 430072 Google Scholar More articles by this author , Lixin Su Hubei Electrochemical Power Sources Key Laboratory, College of Chemistry and Molecular Sciences, Wuhan University, Wuhan, Hubei 430072 Google Scholar More articles by this author , Xuemei Hu Hubei Electrochemical Power Sources Key Laboratory, College of Chemistry and Molecular Sciences, Wuhan University, Wuhan, Hubei 430072 Google Scholar More articles by this author , Yan Wang Key Laboratory of Automobile Materials MOE, School of Materials Science & Engineering, Electron Microscopy Center, and International Center of Future Science, Jilin University, Changchun 130012 Google Scholar More articles by this author , Zizhun Wang Key Laboratory of Automobile Materials MOE, School of Materials Science & Engineering, Electron Microscopy Center, and International Center of Future Science, Jilin University, Changchun 130012 Google Scholar More articles by this author , Lin Zhuang Hubei Electrochemical Power Sources Key Laboratory, College of Chemistry and Molecular Sciences, Wuhan University, Wuhan, Hubei 430072 Google Scholar More articles by this author , Shengli Chen Hubei Electrochemical Power Sources Key Laboratory, College of Chemistry and Molecular Sciences, Wuhan University, Wuhan, Hubei 430072 Google Scholar More articles by this author and Wei Luo *Corresponding author: E-mail Address: [email protected] Hubei Electrochemical Power Sources Key Laboratory, College of Chemistry and Molecular Sciences, Wuhan University, Wuhan, Hubei 430072 Suzhou Institute of Wuhan University, Suzhou, Jiangsu 215123 Google Scholar More articles by this author https://doi.org/10.31635/ccschem.021.202100810 SectionsSupplemental MaterialAboutAbstractPDF ToolsAdd to favoritesDownload CitationsTrack Citations ShareFacebookTwitterLinked InEmail Exploring efficient and economical electrocatalysts and understanding the mechanism for alkaline hydrogen oxidation reaction (HOR) are crucial to facilitate the development of alkaline polymer electrolyte fuel cells (APEFCs). Herein, Ru2P was synthesized and used as an anodic HOR electrocatalyst for APEFC, achieving a peak power density of 1.3 W cm−2, the highest value among Pt-free anode electrocatalysts reported under the same conditions. From the density functional theory (DFT) calculations and experimental results, it was found that besides the optimized hydrogen binding energy, the enhanced adsorption strength of oxygenated species (OH*) and the reduced work function of Ru2P contributed to the enhanced HOR performance. The normalized exchange current densities of Ru2P/C were 0.37 mA cmECSA−2 and 0.27 mA μgRu−1, respectively, both approximately three times higher than those of Ru when conducted in the rotating disk electrode (RDE) system. Our work provides a new pathway for exploring highly active Pt-free HOR electrocatalysts and expanding the family of anodic electrocatalysts for APEFCs. Download figure Download PowerPoint Introduction Proton-exchange membrane fuel cells (PEMFCs) are regarded as the most significant renewable energy conversion technology for efficiently accomplishing transformation from hydrogen to electrical energy, which involves anodic the hydrogen oxidation reaction (HOR), as well as the cathodic oxygen reduction reactions (ORRs).1–3 Currently, Pt remains the most advanced ORR electrocatalyst for PEMFCs, but its use is limited by its high cost and scarcity.4,5 Due to their harsh acidic environment, most nonnoble metals and their alloys are labile and hence lead to low durability in PEMFCs.6 As an alternative to PEMFCs, alkaline polymer electrolyte fuel cells (APEFCs) have received considerable attention by virtue of the development of nonnoble-metal-based electrocatalysts with Pt-like stability and high activity in cathode for ORR in alkaline electrolytes.4,7,8 However, the kinetics of anodic HOR in alkaline solution, even for the most advanced Pt-based catalysts, are approximately two to three orders of magnitude more sluggish than in acidic solution,9,10 which means that more Pt-based catalysts are required to overcome the sluggish anodic HOR in APEFCs. Consequently, developing Pt-free materials as highly active and stable anodic catalysts for APEFCs is highly desirable. Currently, the mechanism for alkaline HOR to guide the catalyst design is still in debate. The hydrogen binding energy (HBE) theory is widely accepted as a valid descriptor to assess the HOR performance by optimizing the strength of adsorbed H (H*) on the surface of electrocatalysts.11–14 However, the most controversial issue is whether HBE is the sole descriptor in alkaline conditions. In fact, Markovic et al.15–18 proposed the bifunctional theory that the strength of both the adsorbed H and adsorbed OH (OH*, a hydroxyl species adsorbed on the active site), are decisive. Furthermore, it has recently been highlighted that the HOR performance of Pt in alkaline media is not only dependent on these thermodynamic characteristics, but also relies on the kinetic factors at electrode/electrolyte interface, for example, the potential of the zero (free) charge (PZC).19,20 Despite the fact that some Ni-based nonprecious metal electrocatalysts have been reported for catalyzing alkaline HOR, their catalytic performances are still far below the practical application in APEFCs.4,6–8,21,22 Most of the APEFCs with Ni-based materials as the anodes only show the peak power density around 0.1 W cm−2.22–24 As the cheapest metal in the platinum group of metals (PGM), Ru-based electrocatalysts have received considerable attention. Ohyama et al.25 reported a Ru/C catalyst applied as the anodic catalyst for APEFC. The peak power density of 0.23 W cm−2 at 323 K was obtained, which is superior to that of their home-made Pt/C (0.20 W cm−2), but still far inferior to that of most commercial Pt/C (0.40–0.82 W cm−2).26–30 According to the electrochemical tests conducted in a rotating disk electrode (RDE) system, the exchange current densities (one of the most vital parameters to assess a catalyst’s intrinsic catalytic activity) of Ru-based catalysts were much lower than that of Pt.31–33 Recently, ruthenium phosphides, such as Ru2P, have been reported as effective electrocatalysts for the hydrogen evolution reaction (HER) due to the optimized HBE induced by the electron transfer between Ru and P atoms.34–36 Some precious transition-metal-based phosphides (TMPs), like palladium phosphides (Pd5P2 and PdP2) and iridium diphosphides (IrP2), have been reported as catalysts for HOR in acidic media.37,38 However, to the best of our knowledge, TMP-based catalysts toward alkaline HOR, especially tested in APEFCs, have so far rarely been reported. Furthermore, in view of the same intermediate during the reaction process (the adsorbed hydrogen, H*) for hydrogen electrode reaction (HER and HOR), we thus expect that the HER-active Ru2P may be a potential candidate for catalyzing alkaline HOR efficiently. In this work, Ru2P nanoparticles loaded on carbon supports (Ru2P/C) have been prepared. When using Ru2P/C as the anodic electrocatalyst for APEFC, a peak power density of 1.3 W cm−2 was obtained, which is the highest record among Pt-free anodic electrocatalysts under the same condition, and even higher than commercial Pt/C in the small-polarization region. Furthermore, the electrochemical tests under RDE system revealed that Ru2P/C exhibits Pt-like HOR performance under alkaline media with ECSA and mass normalized exchange current densities of 0.37 mA cmECSA−2 and 0.27 mA μgRu−1, respectively, both approximately three times higher than those of Ru/C. Experimental Methods Chemicals and materials Ruthenium acetylacetonate (Ru(acac)3), rhodium acetylacetonate (Rh(acac)3), palladium acetylacetonate (Pd(acac)2), and ruthenium chloride hydrate (RuCl3·nH2O) were obtained from Wuhan Changcheng Chemical Co. Ltd (Wuhan, China). tri-n-octylphosphine (TOP), oleylamine (OAm), oleic acid (OA), 1-octadecene (ODE), dodecyl amine (DDA), and tri-n-octylphosphine oxide (TOPO) were obtained from Aladdin Chemical Ltd. (Shanghai, China) Nickel aceylacetonate (Ni(acac)2), cobalt acetylacetonate (Co(acac)2), potassium hydroxide (KOH), perchloric acid (HClO4), isopropyl alcohol, hexane, and ethanol were bought from Sinopharm Chemical Reagent Co. Ltd. (Shanghai, China). Commercial Pt/C (20 wt %) as well as Nafion (5 wt %) were acquired from Sigma-Aldrich (United States). These chemicals were utilized directly without other treatment. Ultrapure water was used during our whole experimental process. Synthesis of Ru2P nanoparticles 0.1 mmol Ru(acac)3 and 1 g TOPO were added into a 50 mL three-neck flask. This system was vigorously magnetically stirred and heated to 120 °C and maintained for 30 min under vacuum to remove residual low-boiling solvents and oxygen. After degassing, the system was heated up to 320 °C, and 0.5 mL TOP was injected. The mixture solution was held at 320 °C for 2 h. Then the reaction was slowly cooled to room temperature (r.t.). The resultant Ru2P nanoparticles were cleaned by adding ethanol to the reaction mixture, followed by centrifugation at 9000 rpm for 5 min, and thereafter washed with ethanol as well as hexane several times to get rid of organic solvent. Eventually, the product obtained was dried at r.t. under vacuum. Synthesis of Ru nanoparticles The synthesis of Ru nanoparticles was similar to what was reported in the previous literature with slight modification.39 Typically, 0.1037 g RuCl3·nH2O, 0.5 mL OA, and 3 mL DDA were added into a 50 mL three-neck flask. The system was vigorously magnetically stirred and heated to 120 °C, then maintained for 30 min under vacuum to remove residual low-boiling solvents and oxygen. After degassing, the reaction system was heated to 320 °C and maintained for 1 h. The postprocess was the same as the synthesis of Ru2P nanoparticles. Synthesis of Rh2P nanoparticles 3 mL OAm, 0.08 mmol Rh(acac)3, and 3 mL ODE were added into a 50 mL round-bottom flask. This system was vigorously magnetically stirred and heated to 120 °C and maintained for 30 min under vacuum to remove residual low-boiling solvents and oxygen. After degassing, the system was heated up to 240 °C and maintained for 60 min, and then 0.4 mL TOP was injected. Thereafter, the system was heated up to 320 °C and held at 320 °C for 2 h. Then the reaction was slowly cooled to r.t. The postprocess was the same as the synthesis of Ru2P nanoparticles. Synthesis of Pd3P nanoparticles 0.1 mmol Pd(acac)2 and 5 mL OAm were added into a 50 mL three-neck flask. This system was vigorously magnetically stirred and heated to 120 °C and maintained for 30 min under vacuum for removing residual low-boiling solvents and oxygen. After degassing, the system was heated up to 320 °C, and 0.5 mL TOP was injected. This mixture solution was held at 320 °C for 2 h. Then this reaction was slowly cooled to r.t. The postprocess was the same as the synthesis of Ru2P nanoparticles. Synthesis of Co2P nanoparticles 0.18 g Co(acac)2 and 5 mL OAm were added into a 50 mL three-neck flask. This system was vigorously magnetically stirred and heated to 120 °C and maintained for 30 min under vacuum for removing residual low-boiling solvents and oxygen. After degassing, the reaction system was heated up to 300 °C, and 0.9 mL TOP was injected. This mixture solution was held at 300 °C for 2 h. The postprocess was the same as the synthesis of Ru2P nanoparticles. Synthesis of Ni2P nanoparticles The synthesis method of Ni2P nanoparticles was analogous to Co2P, except that 0.18 g Ni(acac)2 instead of Co(acac)2 was used, with other experimental parameters unchanged. Materials characterization Powder X-ray diffraction (XRD) patterns were acquired using a Bruker D8-Advance X-ray diffractometer (Bruker, Germany). Catalyst morphologies and sizes were obtained using a Tecnai G20 U-Twin (FEI Co., United States) transmission electron microscopy (TEM) equipped with an energy-dispersive X-ray spectroscopy (EDX). High-angle annular dark-field scanning TEM (HAADF-STEM) image and EDX elemental mapping were acquired by using JEM-ARM300F GRAND ARM (JEOL, Japan) and JEM-2100F (JEOL, Japan). X-ray photoelectron spectroscopy (XPS) was carried out with a Thermo Fischer ESCALAB 250 Xi spectrophotometer (Thermo Fisher Scientific, United States). And inductively coupled plasma atomic emission spectroscopy (ICP-AES) was carried out on a Thermo IRIS Intrepid II XSP atomic emission spectrometer (Thermo Elemental, United States). The H2 temperature programmed desorption (H2-TPD) was conducted with a Micromeritics ChemiSorb 2720 (Micromeritics, United States). Catalyst preparation 5 mg Ru2P nanoparticles and 20 mg XC-72 carbon were added into 20 mL hexane and stirred in a N2 flow for 12 h at r.t. The precipitation was obtained by centrifugation. Then the Ru2P/C was dried under vacuum. Once dry, the Ru2P/C was annealed at 600 °C under 5% H2/N2 to remove surfactants and obtain high crystallinity. Finally, the Ru2P/C catalyst was obtained. The Ru/C catalyst was prepared and postprocessed in the same way. Fifty percentage of loaded catalysts were prepared with 30 mg Ru2P or Ru and 30 mg XC-72 carbon, and postprocessed in the same way for use in single-cell tests. Preparation of working electrodes 4 mg of catalyst was added to 2 mL of mixture solvent containing 5% Nafion/isopropyl alcohol (v/v = 1∶99) with ultrasonic treatment to form a homogeneous solution. Glassy carbon electrode (GCE, 5 mm in diameter) was polished with 0.05 μm gamma alumina powders for a neat surface and then washed. Once the GCE was dried, 5 μL catalyst-ink was drop-coated on the surface of the GCE (loading: ∼0.01 mgPGM cm−2). Electrochemical measurements The electrocatalytic properties were assessed in a three-electrode system using a CHI 760E electrochemical workstation (CH Instruments, Inc., Bee Cave, United States) at 30 °C. The GCE coated with catalyst samples, a graphite rod, and Hg/HgO electrode were applied as the working, counter, and reference electrode, respectively. And the electrolyte was 0.1 M KOH. Before testing the HOR, cyclic voltammetry (CV) was recorded in an Ar-saturated solution to obtain the steady voltammetry curves. Then, polarization curves were obtained by applying a rotation disk electrode (RDE) with a rotation speed of 1600 rpm at a scan rate of 10 mV s−1 in H2-saturated electrolyte. Electrochemical impedance spectra (EIS) tests were carried out after each RDE measurement with the alternating-current (AC) impedance spectra from 200 to 0.1 kHz and a voltage perturbation of 10 mV. All the potential data in this paper were obtained iR-free. Membrane-electrode assembly and single-cell test The as-synthesized Ru2P/C or Ru/C was used as the anode catalyst while the Pt/C (Johnson–Matthey) was used as the cathode catalyst. The home-made quaternary ammonia poly(N-methyl-piperidine-co-p-terphenyl) (QAPPT) was used for the ionomer and the membrane in electrodes. More details of QAPPT can be found in previous literature40,41 ( Supporting Information Table S1). The ink was prepared by mixing Ru2P/C, Ru/C, and Pt/C with QAPPT ionomer solution (20 mg mL−1) in n-propanol, which contained 80 wt % of catalyst and 20 wt % ionomer. Then the ink was ultrasonic for 40 min and then sprayed on the QAPPT APEs (25 ± 3 μm in dry state) to form a catalyst-coated membrane (CCM) with an electrode area of 4 cm−2. The metal loading was controlled to be 0.4 mg cm−2 for anode and 0.4 mg cm−2 for the cathode. Next, the prepared CCM was soaked in 2 M KOH at 80 °C for 24 h for the purpose of changing the anion into OH−, and then washed with distilled water to remove the excess KOH. To make the membrane electrode assembly (MEA) in situ, the CCM was placed between two pieces of carbon paper (AvCarb GDS3250), and the hot-pressing was not used. H2/O2 single-cell APEFCs were tested by an 850e Multi-Range Fuel Cell Test Station (Scribner Associates, Inc., Southern Pines, NC) in a galvanic mode at 80 °C. H2 and O2 were humidified at 80 °C (100% RH) and fed with a flow rate of 1000 sccm with a backpressure of 0.1 MPa symmetrically on both sides. The fuel cell was activated at a constant current for a moment, and then the cell voltage at every current density was recorded. Computational details and theoretical models All the DFT calculations applied in this paper were performed by the PWscf code of the Quantum ESPRESSO package. And the generalized gradient approximation (GGA) method with the Perdew–Burke–Ernzerh (PBE) functional was applied to illustrate the exchange and correlation interactions.42 The interaction between valence electrons and ionic cores was described by Ultrasoft pseudopotential.43 The kinetic energy cutoff was 33 Ry while the charge density cutoff was 330 Ry. The Brillouin zone was sampled using the Monkhorst-Pack method, and the number of k points was set according to this model, of which 2 × 4 × 1 was set for Ru2P geometric optimization and 6 × 6 × 1 for Ru. The convergence threshold was set as 1.0e−6 Ry. In this study, we constructed correlative theoretical models to simulate Ru2P, as well as Ru and Pt for comparison. According to the TEM image (see below), the Ru2P (112) facet was exposed. And the (112) facet showed the highest intensity among all the facets as indicated by the XRD results of Ru2P. Thus, we chose (112) surface as the model of Ru2P for calculation, and Pt (111) and Ru (001) for comparison. For all structure models, four layers were used for calculations, and the top two layers and adsorbate were fully relaxed, with a vacuum gap of ∼15 Å. A slab of (2 × 1 × 1) was used for Ru2P, and slabs of (3 × 3 × 1) were used for Pt and Ru. The calculation of adsorption free energies The adsorption free energy of H (ΔGH*) is usually considered as a descriptor to assess the HER/HOR performance (a catalyst with ΔGH* ≈ 0 could be an excellent candidate for HER).44 The different adsorption free energies of H were determined by the following formula ΔGH* = ΔEH* + 0.24 eV, where ΔEH* represents the binding energy of H* adsorption. The calculation of OH binding energies ΔEOH* = E(OH*) − E(*) − (E(H2O) − 1/2E(H2)),45 for which E(OH*) and E(*) are the ground-state energies of the surfaces with OH* and the clean surface, respectively. E(H2O) and E(H2) are the calculated DFT energies of H2O and H2 molecules in the gas phase. Results and Discussion Amorphorous Ru2P nanoparticles (NPs) were initially prepared by a colloidal method, in which ruthenium acetylacetonate [Ru(acac)3], tri-n-octylphosphine oxide (TOPO), and tri-n-octylphosphine (TOP) were used as Ru precursor, solvent, and P precursor, respectively. The obtained NPs were supported on a certain amount of XC-72 carbon and annealed at 600 °C under 5% H2/N2 to remove the residual surfactants and obtain high crystallinity (denoted as Ru2P/C). X-ray powder diffraction (XRD) was carried out to investigate the crystal structure of the product. As expected, distinct characteristic diffraction peaks corresponding to Ru2P (PDF#89-3031) were observed, which further confirmed the successful formation of Ru2P/C (Figure 1a). Ru nanoparticles were prepared through a similar approach reported previously with some modifications39 and following the same high-temperature annealing process (denoted as Ru/C). Characteristic XRD diffraction peaks of Ru (PDF#88-1734) further confirmed the successful synthesis of Ru/C (Figure 1a). Figure 1 | (a) XRD patterns of Ru2P/C and Ru/C. (b) Cell voltage and power density versus current density curves obtained from APEFC with Ru2P/C (red), Ru/C (blue), and Pt/C (black) as the anode catalysts, respectively. Pt/C was used as the cathode catalyst. All of the metal or metal phosphide loadings in the anodes and the cathodes are controlled to be 0.4 mg cm−2. Single-cells test was conducted at 80 °C. (c) The expanded region of the cell voltage versus current density curves for Ru2P/C and Pt/C from (b). Download figure Download PowerPoint The APEFC performances of Ru2P/C, Ru/C, and commercial Pt/C were first conducted. Figure 1b shows the corresponding plot of cell voltage and power density versus current density. The exact loadings of Ru in Ru2P and Ru were confirmed by ICP-AES with the contents of 43.76 and 48.12 wt %, respectively ( Supporting Information Table S2). When tested with an anode metal (or metal phosphide) loading of 0.4 mg cm−2, Ru2P/C achieved a peak power density of 1.3 W cm−2 (at the current density of 3.0 A cm−2) at 80 °C with 0.1 MPa backpressure, which is the highest value among the reported Pt-free anode catalysts under this condition ( Supporting Information Table S3). For comparison, Ru/C only achieved a peak power density of 0.7 W cm−2 (at the current density of 1.4 A cm−2). Even though the peak power density of Pt/C (1.4 W cm−2, at the current density of 3.2 A cm−2) was a bit higher than Ru2P/C, the current density of Ru2P/C was larger than Pt/C till the cell voltage decrease to 0.87 V, as shown in Figure 1c, suggesting the higher apparent activity of Ru2P/C than Pt/C in the small-polarization region. Due to the same cathodic catalysts, we could confirm that Ru2P/C is a better HOR catalyst, superior to Ru/C, and even comparable with the commercial Pt/C for APEFC. The durability of the Ru2P-catalyzed APEFC single cell was tested in H2/air at 80 °C with 0.2 MPa backpressure. The cell voltage could be maintained at above 0.7 V for 100 h, indicating the stability of the Ru2P-catalyzed APEFC single cell ( Supporting Information Figure S1). To evaluate the alkaline HOR catalytic performance in detail, RDE measurements were conducted in 0.1 M KOH, and the catalyst film-modified glass carbon electrode (GCE) was invoked as the working electrode with a total loading of ∼50 μgcat. cm−2. The exact loadings of Ru in Ru2P and Ru were acquired by ICP-AES with the content of 19.11 and 20.26 wt %, respectively ( Supporting Information Table S4). Commercial Pt/C (20 wt %) was also studied for comparison ( Supporting Information Figure S2). CVs were obtained in an Ar-saturated electrolyte (Figure 2a). The CV peaks of Ru2P/C and Ru/C at about 0.1 V stemmed from the stripping process of the underpotentially deposited hydrogen (H-UPD), accompanied by the oxidation of the Ru surface.33,46–48 And the polarization curves of the three catalysts with iR-compensation tested in the H2-saturated electrolyte are shown in Figure 2b. Similar phenomena can be observed that show that the anodic currents for Ru-based catalysts rise initially from the potential at 0 to around 0.1 V and decline afterward. The decreased currents were illustrated by the excessive oxidation of Ru-species that led to the inhibited behaviors of H adsorption and H2 oxidation.33,46 The HOR/HER current density of Ru2P/C increased much faster than that of Pt/C and Ru/C at around 0 V, suggesting the highest apparent HOR activity of Ru2P/C near the equilibrium potential. It should be noted that due to the surface oxidation, the HOR current density of Ru2P/C at the potential above 0.1 V was inferior to that of Pt/C, which was in accordance with the single-cell test (Figure 1c). Figure 2 | HOR performances of Ru2P/C, Ru/C, and Pt/C. (a) CVs curves of Ru2P/C, Ru/C, and commercial Pt/C. (b) HOR polarization curves and (c) linear-current potential region around the equilibrium potential of HOR/HER of Ru2P/C, Ru/C, and commercial Pt/C after iR-compensation. (d) MA and SA of Ru-based catalysts by rotating disk electrode in 0.1 M KOH. The data of different colors come from different references (dark gray, orange, violet, green, magenta, blue, and red represent refs 33, 49–53 and this work, respectively). Download figure Download PowerPoint Exchange current density (j0) was assessed by the linear-current potential region (−5 to 5 mV) (Figure 2c), using the approximate Butler–Volmer equation, j = j0(ηF/RT),31,33 where j, η, F, R, and T represent the current density, overpotential, Faraday’s constant (96,485 C mol−1), universal gas constant (8.314 J mol−1 K−1), and temperature on the Kelvin scale, respectively. Obviously, the apparent j0 of Ru2P/C was higher than that of Ru/C and Pt/C. In addition, to eliminate interruptions from the currents resulting from the oxidation of Ru or P, the steady-state polarization curve was acquired by the chronoamperometry test ( Supporting Information Figure S3). In addition, similar polarization behaviors between the steady-state polarization curve obtained from the chronoamperometry test and the transient polarization curve acquired from linear scanning voltammetry (LSV) at the micropolarization regions were observed, suggesting that the employment of a transient polarization curve to calculate the j0 does not lead to overestimation. After normalizing by the weight of the PGM ( Supporting Information Table S5), the obtained mass-specific exchange current density (j0,m) of Ru2P/C (0.27 mA μg−1) was a little higher than that of commercial Pt/C (0.21 mA μg−1), and about three times higher than that of Ru/C (0.10 mA μg−1). To further clarify the intrinsic activity of the identified active sites, we tried to estimate the electrochemically active surface areas (ECSAs) of the three samples by the Cu underpotential deposition (Cu-UPD) stripping method ( Supporting Information Figure S4 and Table S5).33 The ECSA-normalized exchange current density (j0,s, specific activity) of Ru2P/C was measured to be 0.37 mA cm−2, which is about three times higher than that of Ru/C (0.12 mA cm−2) and is comparable with that of commercial Pt/C (0.41 mA cm−2). Figure 2d and Supporting Information Tables S5 and S6 show the mass activity (MA) as well as the specific activity (SA) of our samples and other Ru-based catalysts reported previously.33,49–53 It is clear that Ru2P exhibits remarkable HOR performance among the reported Ru-based catalysts, close to the benchmarked Pt-Ru system. The long-term durability of Ru2P/C was probed via accelerated degradation tests (ADTs) ( Supporting Information Figure S5). The CV curves ( Supporting Information Figure S5a) and HOR polarization curves ( Supporting Information Figure S5b) of Ru2P/C were conducted before and after 1000 CVs between 0 and 0.72 V versus RHE. Then the exchange current densities were evaluated (inset of Supporting Information Figure S5b) by means of the approximate Butler–Volmer equation. And after the ADTs, we can see that the apparent exchange current density descended a bit after 1000 cycles but was still higher than the fresh Ru/C, indicating its good HOR stability. Figure 3 | (a) HRTEM image of a single Ru2P nanoparticle of Ru2P/C. Inset shows the SAED pattern of Ru2P/C. (b) HAADF-STEM image and EDX elemental mapping of Ru and P in Ru2P/C. Ru 3p (c) and P 2p (d) core-level XPS spectra of Ru2P/C and Ru/C. Download figure Download PowerPoint The size effect of Ru2P nanoparticles for HOR performances was also examined. Typically, Ru2P/C with average sizes at about 3.3, 4.5, and 6.6 nm were obtained by annealing at 500, 600, and 700 °C, respectively ( Supporting Information Figures S6–S8). The XRD patterns sho

Detecting minute amounts of nitrogen in GaNAs thin films using STEM and CBED

Nitrogen (N) is a common element added to GaAs for band gap engineering and strain compensation. However, detection of small amounts of N is difficult for electron microscopy as well as for other chemical analysis techniques. In this work, N in GaAs is examined by using different transmission electron microscopy (TEM) techniques. While both dark-field TEM imaging using the composition sensitive (002) reflections and selected area diffraction reveal a significant difference between the doped thin-film and the GaAs substrate, spectroscopy techniques such as electron energy loss and energy dispersive X-ray spectroscopy are not able to detect N. To quantify the N content, quantitative convergent beam electron diffraction (QCBED) is used, which gives a direct evidence of N substitution and As vacancies. The measurements are enabled by the electron energy-filtered scanning CBED technique. These results demonstrate a sensitive method for composition analysis based on quantitative electron diffraction.

An efficient combination strategy for high-performance asymmetric-electrolyte metal–air batteries

Asymmetric-electrolyte metal–air batteries (AMABs) are the most promising technologies for realizing efficient energy conversion. However, the issues in cathode and separator parts seriously impede their practical applications. This article previews the latest finding on high-voltage AMABs with long cycle stability. Impressively, the generality of the developed approach is demonstrated by the AZnABs, ASiABs and ASnABs, exhibiting the enhanced operating voltage, power densities, and excellent rate performance. Asymmetric-electrolyte metal–air batteries (AMABs) are the most promising technologies for realizing efficient energy conversion. However, the issues in cathode and separator parts seriously impede their practical applications. This article previews the latest finding on high-voltage AMABs with long cycle stability. Impressively, the generality of the developed approach is demonstrated by the AZnABs, ASiABs and ASnABs, exhibiting the enhanced operating voltage, power densities, and excellent rate performance. The development of advanced energy storage technologies is a promising approach to cope with the increment of energy and environment challenges. Among various renewable energy systems, aqueous metal–air batteries (MABs) have attracted intensive attention owing to their high theoretical energy density, high safety, eco-friendliness, as well as high tolerance to moisture.1Liu Q.F. Pan Z.F. Wang E.D. An L. Sun G.Q. Aqueous metal–air batteries: fundamentals and applications.Energy Storage Mater. 2020; 27: 478-505Crossref Scopus (86) Google Scholar,2Pei Z. Yuan Z. Wang C. Zhao S. Fei J. Wei L. Chen J. Wang C. Qi R. Liu Z. Chen Y. A flexible rechargeable Zinc–air battery with excellent low-temperature adaptability.Angew. Chem. Int. Ed. Engl. 2020; 59: 4793-4799Crossref PubMed Scopus (90) Google Scholar However, the low redox potential and CO2 poisoning issues greatly limit the practical energy density and durability of MABs.3Fu J. Cano Z.P. Park M.G. Yu A. Fowler M. Chen Z. Electrically rechargeable zinc–air batteries: Progress, challenges, and perspectives..Adv. Mater. 2017; 29: 1604685Crossref Scopus (804) Google Scholar To address the above problems, an attractive concept of asymmetric-electrolyte MABs (AMABs) was proposed in 2017.4Yu X. Manthiram A. Electrochemical energy storage with mediator-ion solid electrolytes.Joule. 2017; 1: 453-462Abstract Full Text Full Text PDF Scopus (17) Google Scholar In AMABs, an ion-conductive solid-electrolyte separator is generally used to make the acidic catholyte and alkaline anolyte separated to avoid the persistent CO2 ingression at the side of the air electrode. Moreover, the delivered oxygen reduction reaction (ORR) potential (1.23 V versus SHE) in an acidic medium will be higher, which increases the output voltage and energy density of the practical device.5Koper M.T.M. Theory of multiple proton-electron transfer reactions and its implications for electrocatalysis.Chem. Sci. (Camb.). 2013; 4: 2710-2723Crossref Scopus (376) Google Scholar Unfortunately, the emerging AMABs still can not compete with the demand of the large-scale application due to the following challenges currently faced: 1) the commercially available separators, no matter the bipolar polymer membrane or ceramic solid-state membrane, show low ion conductivity, unavoidable crossover of H+ and OH– ions and ultrahigh cost (e.g., $ 3200 m-2 for bipolar polymer membrane);6Cai P.W. Li Y. Chen J.X. Jia J.C. Wang G.X. Wen Z.H. An asymmetric-electrolyte Zn–air battery with ultrahigh power density and energy density.ChemElectroChem. 2018; 5: 589-592Crossref Scopus (32) Google Scholar 2) the lack of cost-effective and highly efficient bifunctional catalysts in the challenging acidic electrolyte.7Yang Y.C. Yang Y.W. Pei Z.X. Wu K.H. Tan C.H. Wang H.Z. Wei L. Mahmood A. Yan C. Dong J.C. et al.Recent progress of carbon-supported single-atom catalysts for energy conversion and storage.Matter. 2020; 3: 1442-1476Abstract Full Text Full Text PDF Scopus (95) Google Scholar In this issue of Matter, Lee et al. reported a combination strategy that includes the novel separator fabrication and atomically dispersed catalyst design to realize the high-voltage and long-term stable AMABs.8Lin C. Kim S.-H. Xu Q. Kim D.-H. Ali G. Shinde S.S. Yang S. Yang Y. Li X. Jiang Z. Lee J.-H.o. High-voltage asymmetric metal–air batteries based on polymeric single-Zn2+-ion conductor.Matter. 2021; 4 (this issue): 1287-1304Abstract Full Text Full Text PDF Scopus (8) Google Scholar The cost-effective polyacrylonitrile (PAN)-based single-Zn2+-ion conductor is first reported to selectively transport Zn2+ ions and inhibit the crossover of H+ and OH– ions. In addition, an atomically dispersed Co electrocatalyst supported on nitrogen-doped carbon nanosheet (Co/ NS ) with high activity and stability for both ORR and oxygen evolution reaction ( OER ) in acidic medium is prepared as the cathodic catalyst for AMABs. Remarkably, the as-fabricated asymmetric Zn–air battery (AZnAB) delivers an ultrahigh specific energy density of 1,354 Wh kgZn-1 (approaching the theoretical value of the alkaline ZnAB), and long-term stability (100 h) with a round-trip efficiency of 76.6%. Moreover, the generality of the developed strategy is successfully verified in the other-type assembled asymmetric MABs (metals, Si, Sn). The Zn2+ ion-selective transport membrane (ZnSTM) is synthesized by heating a mixture of Zn(OAc)2 and PAN at 140 °C in air condition. The chemical structure is investigated by a series of characterizations including X-ray diffraction (XRD), Fourier transform infrared spectroscopy (FTIR) and transmission electron microscopy (TEM) techniques. The Zn2+ anchored ladder structure (Figure 1A) is formed through a cyclization reaction from C≡N to C = N bond. The crosslinking reactions between the adjacent ladder polymeric molecular chains enable mechanically stable and flexible polymer membrane (tensile strength of 0.8 Ma). TEM image confirms the microphase separation structure composed of the hydrophobic polymeric backbones and hydrophilic domains generated by Zn2+ ions aggregated around the N-related groups. The hydrophilic domains with a Zn2+ ionic conductivity of 1.9 × 10−4 S cm-1 at 313 K could offer the rapid Zn2+ ions transport channel (Figure 1B). The Arrhenius plot shows the calculated activation energy of ZnSTM is ca. 0.2 eV (< 0.4 eV) (Figure 1C), indicating the Zn2+ migration follows the Grotthuss hopping mechanism. Computational calculation unveils that the attractive force between skeleton coordinated Zn2+ sites and H+/ OH– ions inhibits the transport of H+ and OH– ions, while repulsive force between skeleton coordinated Zn2+ sites and guest Zn2+ ions benefits for the migration of Zn2+ ions. The efficient cathodic electrocatalysts for ORR and OER in acidic media are highly desirable but challenging.9Wang X.X. Swihart M.T. Wu G. Achievements, challenges and perspectives on cathode catalysts in proton exchange membrane fuel cells for transportation.Nat. Catal. 2019; 2: 578-589Crossref Scopus (413) Google Scholar,10Reier T. Nong H.N. Teschner D. Schlögl R. Strasser P. Electrocatalytic oxygen evolution reaction in acidic environments-reaction mechanisms and catalysts.Adv. Energy Mater. 2016; 7: 1601275Crossref Scopus (530) Google Scholar In this work, the Co/NS is prepared by pyrolysis of bimetallic ZnCo zeolitic imidazolate frameworks, and further used as bifunctional catalysts for ORR and OER in O2-saturated 0.1 M HClO4. Aberration-corrected atomic-resolution high-angle annular dark-field scanning TEM (AC-HAADF-STEM) image indicates the average diameter of nanocluster is ∼1 nm with a lattice fringe of 0.2 nm. And, the local coordination environment is presented by the X-ray absorption spectrum, revealing that the Co species exist in the form of Co-Nx (x = 1-4). As a result, the as-prepared Co/NS exhibits a half-wave potential of 0.71 V for ORR in 0.1 M HClO4 and a small Tafel slope of 82.1 mV dec-1 (Figure 1D). The over-potential of the Co/NS at 10 mA cm-2 is 605 mV along with no observed activity degradation after 20 h electrocatalysis (Figure 1E). In situ extended X-ray absorption fine structure (EXAFS) data (Figure 1F) show the bond length of Co-N periodically reduces and then increases between ORR and OER potential, demonstrating the active site for ORR and OER is Co-Nx species. As a demo of the real applications, the AZnAB is assembled using the developed ZnSTM and atomically dispersed Co/NS catalysts (Figure 1G). Notably, the aqueous AZnAB delivers an operating voltage of 1.96 V, an ultrahigh specific energy density of 1,354 Wh kgZn-1, as well as long-term stability (∼100 h) with a round-trip efficiency of 76.6%. The specific energy density of 1,354 Wh kgZn-1 approaches the theoretical limit of alkaline electrolyte-based ZnAB, higher than that of previously reported conventional ZnABs (Figure 1H). And, the round-trip efficiency after 300 cycles is up to 70.8%, which is comparable to Pt/C+RuO2-based AZnAB (75.1%) under identical conditions. In addition, the developed combination strategy is extended to the other type AMABs including ASiAB and ASnAB. Both ASiAB and ASnAB display the enhanced operating voltage, power densities, and rate performance than those of recently reported SiABs and SnABs. Significant progress in AMABs has been made through the design of ions conductor and electrocatalyst. More research is recommended to be carried out to further improve the overpotential of the oxygen electrode in the acidic media, especially for the OER. In addition, the internal resistance from the membrane and cathode still needs to be optimized to promote the output current density of the fabricated AMABs. Apart from the nanoengineering in separator and cathode, the issues of anode such as the formation of Zn dendrites in the alkaline electrolyte are urgently addressed. The strategy from system thinking plays a key role in the commercialization process of the AMABs. High-voltage asymmetric metal–air batteries based on polymeric single-Zn2+-ion conductorLin et al.MatterJanuary 29, 2021In BriefThe development of asymmetric-electrolyte metal–air batteries (AMABs) has been hindered by lack of polymer separator that can selectively transport metal ions and cost-effective bifunctional electrocatalyst. Here, we report breakthroughs in both key components by constructing a polyacrylonitrile membrane that can selectively transport Zn2+ ions, and an atomically dispersed Co electrocatalyst that can stably catalyze ORR and OER in acid. The assembled AMABs (metals Zn, Si, Sn) showed stable battery performance with enhanced operating voltage, power densities, and excellent rate performance that surpass those of recently reported MABs. Full-Text PDF Open Archive

Single-phase FeMnNiAl compositionally complex alloy

FeMnNiAl, a compositionally complex alloy, in close to having an equimolar composition, was prepared in bulk and thin-film forms. Arc melting under argon was used for bulk form preparation. The thin-film alloy was produced by Ion Beam Sputter Deposition on a silicon substrate at 620 K. X-Ray Diffraction (XRD) and Transmission Electron Microscopy (TEM) were used to assess the material crystallinity. Both alloy forms display a single-phase and are in crystalline state. The material has a BCC lattice with the lattice constant at 4.08 Å. Dark field TEM image reveals material nanotexturing. The potentiodynamic polarisation in 0.6 moles NaCl and H2SO4 shows that open circuit potential (against saturated calomel electrode) at room temperature is approximately - 0.21 V. The alloy is demonstrated to be nobler than 304 SS (by approximately 0.1 V) and develops passivation layer in the corrosive solutions. Nanoindentation test shows that FeMnNiAl alloy thin film has high hardness at 11.7 GPa. It is anticipated that high hardness combined with the passivation ability can make the FeMnNiAl alloy coating a viable candidate for protection of metallic constructions in erosive environments

Metal–Organic Framework-Derived CuS Nanocages for Selective CO 2 Electroreduction to Formate

Open AccessCCS ChemistryRESEARCH ARTICLE1 Dec 2021Metal–Organic Framework-Derived CuS Nanocages for Selective CO2 Electroreduction to Formate Xing Zhang, Rongjian Sa, Feng Zhou, Yuan Rui, Ruixia Liu, Zhenhai Wen and Ruihu Wang Xing Zhang State Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fujian 350002 University of Chinese Academy of Sciences, Beijing 100049 Google Scholar More articles by this author , Rongjian Sa State Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fujian 350002 Institute of Oceanography, Ocean College, Minjiang University, Fujian 350108 Google Scholar More articles by this author , Feng Zhou State Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fujian 350002 University of Chinese Academy of Sciences, Beijing 100049 Google Scholar More articles by this author , Yuan Rui State Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fujian 350002 Google Scholar More articles by this author , Ruixia Liu Innovation Academy for Green Manufacture, Chinese Academy of Sciences, Beijing 100190 Google Scholar More articles by this author , Zhenhai Wen State Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fujian 350002 Google Scholar More articles by this author and Ruihu Wang *Corresponding author: E-mail Address: [email protected] State Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fujian 350002 University of Chinese Academy of Sciences, Beijing 100049 Innovation Academy for Green Manufacture, Chinese Academy of Sciences, Beijing 100190 Google Scholar More articles by this author https://doi.org/10.31635/ccschem.020.202000589 SectionsSupplemental MaterialAboutAbstractPDF ToolsAdd to favoritesDownload CitationsTrack Citations ShareFacebookTwitterLinked InEmail Electroreduction of CO2 to target products with high activity and selectivity has techno-economic importance for renewable energy storage and CO2 utilization. Herein, we report a hierarchical CuS hollow polyhedron (CuS-HP) for electrocatalytic CO2 reduction (E-CO2R) in neutral pH aqueous media. Under E-CO2R conditions, CuS-HP undergoes structural reconstruction into sulfur-doped metallic Cu catalyst, which promotes formate production with Faradaic efficiency >90% in a wide potential range. The durability of the in situ evolved catalyst has been demonstrated by stable operation for 36 h at a formate partial current density of ∼16 mA cm−2 at −0.6 V versus reversible hydrogen electrode. Theoretical thermodynamic analysis reveals that the sulfur-doped Cu(111) facet is responsible for high formate selectivity by promoting formate production and suppressing hydrogen evolution. Download figure Download PowerPoint Introduction Electrocatalytic CO2 reduction (E-CO2R) powered by renewable electricity is a promising route to mitigate net CO2 emissions and simultaneously synthesize value-added products.1 CO2 reduction is a highly complex process as it involves multiple reaction paths and products, such as CO, formate, ethylene, ethanol, and propanol.2–6 Therefore, active and product-selective catalysts are urgently required for E-CO2R to enhance energy efficiency and reduce the cost of product separation. Formate is an important commodity chemical, which can be directly used as fuel or feedstock in fine chemical synthesis.7,8 Catalysts based on main group metals, including indium (In), tin (Sn), lead (Pb), and bismuth (Bi), have been experimentally identified as highly selective for the conversion of CO2 into formate in neutral pH aqueous media.9–19 However, many of the reported catalysts promote formate production with Faradaic efficiency >80% only at very negative potentials [<−0.8 V vs reversible hydrogen electrode (RHE)]. Recently, transition-metal cobalt (Co)-based catalysts with several atomic-layer thickness have been shown to be active for carbon dioxide-to-formate conversion at low overpotentials.20–22 Unfortunately, this performance merit could only be achieved at very low cathodic current density (<10 mA cm−2). Metallic copper (Cu) is an appealing candidate for the construction of E-CO2R catalysts because its moderate CO binding strength enables the reduction of CO2 into multicarbon products at potentials below −1.0 V versus RHE.23–25 At low overpotentials, Cu-based catalysts mainly catalyze the reduction of CO2 into CO and formate with relatively low Faradaic efficiency.8,26,27 Controlling the morphology, oxidation state, and exposed facets of Cu can tune the binding strength of the intermediates and thus the selectivity of target products.28–32 For example, sulfur (S)-doped Cu catalysts derived from CuSx precursors have exhibited remarkably enhanced formate selectivity under E-CO2R conditions.33–37 However, Faradaic efficiencies toward formate remain <80% and a mechanistic understanding of the actual active sites is still elusive. Structural reconstruction of catalyst materials under real working conditions is a common phenomenon.3,25,38 The initial morphology, composition, and coordination environment of Cu-based catalysts significantly affect structural evolution of the active catalysts under E-CO2R conditions.39–43 Therefore, it is essential to engineer the composition, morphology, and structure of Cu-based catalyst precursors or precatalysts and track their structural evolution under working conditions for understanding reaction mechanisms and improving E-CO2R performance. Herein, we present a hierarchical CuS hollow polyhedron (CuS-HP), which was synthesized via a metal–organic framework self-sacrificial template method. Using operando and postcatalysis spectroscopic characterizations, we found that the CuS phase in CuS-HP completely disappeared and a S-doped metallic Cu catalyst [Cu(S)] was generated under E-CO2R conditions. The in situ evolved Cu(S) affords formate Faradaic efficiency of 94% at −0.5 V versus RHE and partial current density of 57 mA cm−2 at −0.8 V versus RHE. Theoretical calculations identified that the S-doped Cu(111) surface is necessary for high formate selectivity due to its lower thermodynamic energy barrier for formate production and weaker hydrogen binding strength compared with those on Cu(111). Experimental Methods Synthesis of CuS-HP HKUST-1 was initially prepared per the modified literature method.44 Using ultrasonication for 3 min, 60 mg of HKUST-1 was dispersed in 30 mL of ethanol. After 2 mL of 0.5 M thioacetamide ethanol solution was added, the mixture was sealed in a gas-tight glass vessel (48 mL) and heated with stirring at 90 °C for 1 h. The black precipitate was collected by centrifugation and washed with ethanol three times. Synthesis of CuS nanosheets To 6 mL of 0.5 M thioacetamide ethanol solution was added 25 mL of 0.2 M Cu(NO3)2 ethanol solution . The mixture was then sealed in a gas-tight glass vessel (48 mL) and heated with stirring at 90 °C for 3 h. The black precipitate was collected by centrifugation and washed with ethanol three times. Synthesis of Cu nanoparticles 1-Hexadecylamine (200 mg) and Cu(OAc)2 (30 mg) were dissolved in 30 mL of ethanol. After 80 mg of l-ascorbic acid was added, the mixture was sealed in a gas-tight glass vessel (48 mL) and heated with stirring at 90 °C for 0.5 h. The reddish precipitate was collected by centrifugation and washed with ethanol three times. Other details related to material characterization, electrochemical measurements, and computational methods are available in the Supporting Information. Results and Discussion CuS-HP was readily synthesized from the HKUST-1 template via the solvothermal sulfidation method. As illustrated in Figure 1a, the octahedron-shaped HKUST-1 was initially prepared by coordination assembly of 1,3,5-benzenetricarboxylic acid and Cu2+ in methanol ( Supporting Information Figure S1).44 The formation of CuS-HP follows the Kirkendall-type mechanism. Specifically, thioacetamide releases S2−, which etches the surface of HKUST-1 to form a thin CuS seed layer. The dissolution of metastable HKUST-1 results in fast outward diffusion of Cu2+, and subsequent reaction with S2− enables the resultant CuS to continuously deposit on the outer surface, which generates hierarchical HP consisting of CuS nanosheets. The powder X-ray diffraction (XRD) pattern of CuS-HP fit well with the hexagonal CuS phase (Figure 1b). Transmission electron microscopy (TEM) and high-angle annular dark-field scanning TEM (HAADF-STEM) images show that CuS-HP has a mean size of ∼1.3 μm and is composed of crystalline CuS nanosheets with a thickness of ∼3 nm (Figures 1c–1f and Supporting Information Figure S2). The HAADF-STEM image and coupled energy-dispersive X-ray spectroscopy (EDS) elemental maps show the uniform distribution of Cu and S over CuS-HP (Figures 1g–1j). The Raman spectrum of CuS-HP exhibits a strong resonance peak centered at 470 cm−1, which is assigned to the S–S bond stretching vibration in CuSx ( Supporting Information Figure S3a).45,46 X-ray photoelectron spectroscopy (XPS) characterizations confirm that the binding energies of Cu 2p and S 2p correspond to those of CuSx ( Supporting Information Figures S3b–S3d).37 The Cu:S ratio quantified by XPS is 1.06, which conforms to the stoichiometry of CuS. Figure 1 | Structural characterization of CuS-HP. (a) Schematic illustration of the synthetic process. (b) XRD pattern. (c–f) TEM images. (g–j) HAADF-STEM image and corresponding EDS elemental mapping images. Download figure Download PowerPoint The E-CO2R performance of CuS-HP was evaluated in an H-type electrolyzer filled with CO2-saturated 0.5 M K2SO4 aqueous solution as the electrolyte ( Supporting Information Figure S4). The catalyst electrode was prepared by pasting CuS-HP on carbon fiber paper (CFP) with a mass loading of 1 mg cm−2. After the catalyst electrode was activated by 10 cyclic voltammetry (CV) scans from 0.2 to −0.8 V versus RHE at a scan rate of 20 mV s−1, repeatable CV curves were observed. The polarization curve at 2 mV s−1 shows a rapid increase of cathodic current as the potential is more negative than −0.3 V versus RHE, indicative of high electrode activity (Figure 2a). The formate Faradaic efficiency exceeds 90% in the surveyed potential region from −0.5 to −0.8 V versus RHE, while the CO production rate is negligible (Figure 2b). The corresponding formate cathodic energy efficiencies in the surveyed potential region are >50% ( Supporting Information Figure S5). For comparison, commercially available CuS nanoparticles (NPs) with mean size of 50 nm were tested under the same conditions ( Supporting Information Figure S6). The Brunauer–Emmett–Teller (BET) specific surface area of CuS-HP (20.2 m2 g−1) is two times higher that of CuS NPs (9.8 m2 g−1; Supporting Information Figure S7). However, the electrochemical active surface area (ECSA) of CuS-HP is nearly three times larger than that of CuS NPs ( Supporting Information Figures S8a–S8c). The difference in BET surface area and ECSA can be attributed to different surface wettability to gas and liquid or structural changes under electrochemical conditions. It was found that both the electrode activity and formate selectivity of CuS NPs are significantly inferior to those of CuS-HP (Figures 2a and 2b). By normalizing to their ECSA, the specific activity of CuS-HP is slightly lower than that of CuS NPs ( Supporting Information Figure S8d). Thus, this suggests that there is a significant difference in catalytically active structures derived from CuS-HP and CuS NPs under E-CO2R conditions. The robust catalytic durability of CuS-HP was further corroborated by continuous operation at −0.6 V versus RHE over 36 h (Figure 2c). When evaluated in a gas-diffusion flow cell, which reduces mass transport limitation, CuS-HP electrode affords a formate Faradaic efficiency of 95% at a current density of −300 mA cm−2 with an applied cathodic potential of −0.91 V ( Supporting Information Figure S9). The excellent catalytic performance ranks CuS-HP in the first class of reported nonprecious metal-based formate-selective E-CO2R catalysts (Figure 2d and Supporting Information Table S1). Figure 2 | E-CO2R performance of CuS-HP. (a) Polarization curves and (b) product selectivity of CuS-HP and CuS NPs. (c) Stability test and formate Faradaic efficiencies of CuS-HP at −0.6 V vs RHE. (d) Comparison of potential-dependent formate partial current densities with those of NTD-Bi,10 S-In,12 SELF-CAT-Pb,13 Sn(S)/Au,15 SnO2 QWs,17 and AC-CuSx.35 Download figure Download PowerPoint Structural reconstruction of Cu-based catalysts under E-CO2R conditions is common and the derived active structure generally correlates with the initial morphology and structure of precatalyst.39–41 The operando Raman spectroscopic characterization was performed to track structural evolution of CuS-HP under E-CO2R conditions. During the measurements, the potential applied on the electrode was reduced by steps from open circuit potential (OCP) to −0.2 V versus RHE. Each potential was held for 15 min before recording the Raman spectra. The results show that the characteristic S–S bond stretching vibration peak totally disappeared once the applied potential was more negative than −0.15 V versus RHE, indicative of the destruction of the CuS phase (Figure 3a). The CV curves of fresh CuS-HP/CFP electrode show a cathodic reduction peak started at −0.15 V versus RHE ( Supporting Information Figure S10). This cathodic reduction process is ascribed to the transformation of CuS into Cu.34 Figure 3 | Characterizations of the structural evolution of CuS-HP under E-CO2R conditions. (a) Operando Raman spectra at different applied potentials. (b) XRD patterns, (c) S 2p, and (d) Cu 2p XPS spectra before and after constant-potential electrolysis for 1 h. Download figure Download PowerPoint Ex situ XRD and XPS were performed to investigate structural and compositional changes of CuS-HP after constant-potential E-CO2R for 1 h at −0.6 V versus RHE. The XRD pattern of CuS-HP (−0.6 V) reveals the appearance of metallic Cu phase and complete disappearance of CuS phase (Figure 3b). High-resolution S 2p XPS spectra of CuS-HP (−0.6 V) show the survival of minimal Sδ− (0 ≤ δ ≤ 2) in the derived catalysts (Figure 3c). It is noted that the signals at binding energy >164 eV in S 2p spectra are ascribed to Nafion binder or adsorbed sulfate ions from electrolyte ( Supporting Information Figure S11), which will not be involved in subsequent discussion.34 Based on the XPS analyses, the content of Sδ− is 14.1 atom % in CuS-HP (−0.6 V). When the electrolysis potential was further decreased to −0.8 V versus RHE, the sulfur content remained 12.3 atom %, suggesting high stability of sulfur dopants in Cu(S) in the surveyed potential region. The Cu 2p XPS spectra and Cu LMM auger spectra show that the mean oxidation state of Cu in both CuS-HP (−0.6 V) and CuS-HP (−0.8 V) is higher than those in CuS and metallic Cu, which could be ascribed to extensive oxidation of superficial Cu atoms of in situ evolved Cu(S) upon exposure to air (Figure 3d and Supporting Information Figure S12). Thus, these operando and postcatalysis characterization results demonstrate that under E-CO2R conditions CuS-HP is reconstructed into Cu(S), which is the active species for highly selective conversion of CO2 to formate. Sulfur-free HKUST-1 and Cu NPs were also tested under the same conditions. They mainly promote H2 generation and the formate Faradaic efficiencies are <40% over the surveyed potential region ( Supporting Information Figures S13 and S14). These observations demonstrate that S-doping in Cu is important for high formate selectivity. CuS nanosheets with disorderly stacked morphology were also synthesized by direct treatment of Cu(NO3)2 with thioacetamide. The as-synthesized CuS nanosheets have similar microstructure to the assembly unit of CuS-HP ( Supporting Information Figure S15). However, the E-CO2R activity of CuS nanosheets is inferior to that of CuS-HP ( Supporting Information Figure S16), suggesting the advantages of the regularly assembled hollow morphology of CuS-HP derived from HKUST-1. It has been widely recognized that metal–organic frameworks are one kind of promising precursor for constructing functional nanomaterials with well-defined composition, morphology, and microstructure.47,48 To better understand the effects of S-doping on the formate selectivity of Cu(S), density functional theory (DFT) calculations were performed to study reaction thermodynamics of E-CO2R and competitive hydrogen evolution reaction (HER) on Cu, Cu0.96S0.04, and Cu0.87S0.13. The Cu0.87S0.13 model slab was chosen due to its similar composition to that of CuS-HP-derived Cu(S). The Cu0.96S0.04 slab was also chosen as a model for the sample with a low S-doping amount. First, we performed the calculations on the Cu(111) surface based on two reasons: (1) Cu(111) is the most stable facet in polycrystalline Cu due to its lowest surface energy among various Cu facets.49 (2) The Cu(111) facet is selective for C1 product.50 Formate production on catalyst surfaces can proceed through either COOH* or OCHO* intermediates, whereas CO and H2 evolution occur through COOH* and H* intermediates, respectively ( Supporting Information Figures S17 and S18).15,51 On the (111) surface of all model slabs, free-energy diagrams show that COOH* is the relevant intermediate for formate production and the potential-limiting step is the second proton-coupled electron-transfer process, which results in the formation of HCOOH from COOH* (Figures 4a–4c). The theoretical overpotential for formate production is 0.22 V, which is much smaller than that of Cu (0.35 V) and Cu0.96S0.04 (0.3 V). CO evolution is suppressed due to the strong binding strength of CO* on all surfaces. With the increase of S-doping amount, H* binding strength on the (111) surface becomes weaker, and Cu0.87S0.13 shows much larger theoretical HER overpotential (0.25 V) than that of pure Cu (0.17 V) (Figure 4d). We also performed DFT calculations on the (100) surface of these model slabs ( Supporting Information Figures S19 and S20). The results show that the formation of both formate and CO are unfavorable ( Supporting Information Figures S21a–S21c). Like the case of the (111) surface, the incorporation of S atom in Cu significantly weakens H* binding strength on the (100) surface ( Supporting Information Figure S21d). Therefore, the Cu(S)(111) surface is responsible for high formate selectivity of Cu(S). Figure 4 | Theoretical thermodynamic calculations for E-CO2R on Cu(111) and Cu(S)(111) surfaces. Free-energy diagrams for HCOOH and CO on (a) Cu, (b) Cu0.96S0.04, and (c) Cu0.87S0.13. (d) Comparison of thermodynamic energy barrier for HER on Cu, Cu0.96S0.04, and Cu0.87S0.13. Download figure Download PowerPoint Conclusions We have synthesized CuS-HP consisting of crystalline CuS nanosheets and explored its electrocatalytic performance for CO2 reduction. We have demonstrated that potential-induced reconstruction of CuS-HP under E-CO2R conditions generates Cu(S), which exhibits the highest formate selectivity among the reported Cu-base catalysts. Theoretical studies have revealed that S-doping in metallic Cu does not switch the reaction paths, but lowers the energy barriers for formate production and simultaneously suppresses HER. This work highlights that engineering the initial morphology and structure of precatalysts is one effective pathway to tailor in situ evolution of catalytically active structure and thus product selectivity of E-CO2R. Conflicts of Interest The authors declare no competing financial interests. Supporting Information Supporting Information is available, including additional experimental methods, Figures S1–S21, and Table S1. Acknowledgments This work was financially supported by the National Natural Science Foundation of China (nos. 22002168, 21975259, and 21673241), the Innovation Academy for Green Manufacture of the Chinese Academy of Sciences (no. IAGM2020C17), and the Strategic Priority Research Program of the Chinese Academy of Sciences (no. XDB20000000). The authors thank Prof. Y. Liang from Southern University of Science and Technology for the kind use of TEM and Raman instruments.

Phase-Controlled 1T Transition-Metal Dichalcogenide-Based Multidimensional Hybrid Nanostructures

Open AccessCCS ChemistryRESEARCH ARTICLE1 Nov 2021Phase-Controlled 1T Transition-Metal Dichalcogenide-Based Multidimensional Hybrid Nanostructures Hou-Ming Xu†, Chao Gu†, Xiao-Long Zhang†, Lei Shi, Qiang Gao, Shaojin Hu, Shi-Kui Han, Xiao Zheng, Min-Rui Gao and Shu-Hong Yu Hou-Ming Xu† Key Laboratory of Advanced Catalytic Materials and Reaction Engineering, School of Chemistry and Chemical Engineering, Hefei University of Technology, Hefei 230009 †H.-M. Xu, C. Gu, and X.-L. Zhang contributed equally to this work.Google Scholar More articles by this author , Chao Gu† Hefei National Laboratory for Physical Sciences at the Microscale, Division of Nanomaterials and Chemistry, Department of Chemistry, CAS Centre for Excellence in Nanoscience, Hefei Science Centre of CAS, Hefei Comprehensive National Science Centre, Institute of Energy, University of Science and Technology of China, Hefei 230026 †H.-M. Xu, C. Gu, and X.-L. Zhang contributed equally to this work.Google Scholar More articles by this author , Xiao-Long Zhang† Hefei National Laboratory for Physical Sciences at the Microscale, Division of Nanomaterials and Chemistry, Department of Chemistry, CAS Centre for Excellence in Nanoscience, Hefei Science Centre of CAS, Hefei Comprehensive National Science Centre, Institute of Energy, University of Science and Technology of China, Hefei 230026 †H.-M. Xu, C. Gu, and X.-L. Zhang contributed equally to this work.Google Scholar More articles by this author , Lei Shi Hefei National Laboratory for Physical Sciences at the Microscale, Division of Nanomaterials and Chemistry, Department of Chemistry, CAS Centre for Excellence in Nanoscience, Hefei Science Centre of CAS, Hefei Comprehensive National Science Centre, Institute of Energy, University of Science and Technology of China, Hefei 230026 Google Scholar More articles by this author , Qiang Gao Hefei National Laboratory for Physical Sciences at the Microscale, Division of Nanomaterials and Chemistry, Department of Chemistry, CAS Centre for Excellence in Nanoscience, Hefei Science Centre of CAS, Hefei Comprehensive National Science Centre, Institute of Energy, University of Science and Technology of China, Hefei 230026 Google Scholar More articles by this author , Shaojin Hu Hefei National Laboratory for Physical Sciences at the Microscale, Division of Theoretical and Computational Sciences, CAS Centre for Excellence and Synergetic Innovation Centre in Quantum Information and Quantum Physics, University of Science and Technology of China, Hefei 230026 Google Scholar More articles by this author , Shi-Kui Han *Corresponding authors: E-mail Address: [email protected] E-mail Address: [email protected] E-mail Address: [email protected] Key Laboratory of Advanced Catalytic Materials and Reaction Engineering, School of Chemistry and Chemical Engineering, Hefei University of Technology, Hefei 230009 Google Scholar More articles by this author , Xiao Zheng Hefei National Laboratory for Physical Sciences at the Microscale, Division of Theoretical and Computational Sciences, CAS Centre for Excellence and Synergetic Innovation Centre in Quantum Information and Quantum Physics, University of Science and Technology of China, Hefei 230026 Google Scholar More articles by this author , Min-Rui Gao *Corresponding authors: E-mail Address: [email protected] E-mail Address: [email protected] E-mail Address: [email protected] Hefei National Laboratory for Physical Sciences at the Microscale, Division of Nanomaterials and Chemistry, Department of Chemistry, CAS Centre for Excellence in Nanoscience, Hefei Science Centre of CAS, Hefei Comprehensive National Science Centre, Institute of Energy, University of Science and Technology of China, Hefei 230026 Google Scholar More articles by this author and Shu-Hong Yu *Corresponding authors: E-mail Address: [email protected] E-mail Address: [email protected] E-mail Address: [email protected] Hefei National Laboratory for Physical Sciences at the Microscale, Division of Nanomaterials and Chemistry, Department of Chemistry, CAS Centre for Excellence in Nanoscience, Hefei Science Centre of CAS, Hefei Comprehensive National Science Centre, Institute of Energy, University of Science and Technology of China, Hefei 230026 Google Scholar More articles by this author https://doi.org/10.31635/ccschem.021.202000578 SectionsSupplemental MaterialAboutAbstractPDF ToolsAdd to favoritesDownload CitationsTrack Citations ShareFacebookTwitterLinked InEmail Metallic-phase transition-metal dichalcogenides (TMDCs) exhibit unusual physicochemical properties compared with their semiconducting counterparts. However, they are thermodynamically unstable to access and it is even more challenging to construct their metastable-phase heterostructures. Herein, we demonstrate a general solution protocol for phase-controlled synthesis of distorted octahedral 1T WS2-based (1T structure denotes an octahedral coordination for W atom) multidimensional hybrid nanostructures from two-dimensional (2D), one-dimensional (1D), and zero-dimensional (0D) templates. This is realized by tuning the reactivity of tungsten precursor and the interaction between crystal surface and ligands. As a conceptual study on crystal phase- and dimensionality-dependent applications, we find that the three-dimensional (3D) hierarchical architectures achieved, comprising 1T WS2 and 2D Ni3S4, are very active and stable for catalyzing hydrogen evolution. Our results open up a new way to rationally design phase-controlled nanostructures with increased complexity and more elaborate functionalities. Download figure Download PowerPoint Introduction Crystal phase engineering of inorganic materials plays a key role in extending the spectrum of reachable materials and exploring their uncovered physicochemical properties.1–5 In particular, group-VI transition-metal dichalcogenides (TMDCs; the transition metals are Mo and W, and the chalcogens are S, Se, and Te) have been extensively studied due to their unique metallic octahedral or distorted octahedral 1T phases (1T structure denotes an octahedral coordination for W atom).6,7 These 1T TMDCs have been reported to exhibit superior performance for catalytic hydrogen evolution reaction (HER) compared with their corresponding semiconducting prismatic 2H phases (2H structure denotes a prismatic coordination for W atom), owing to the dramatically reduced charge-transfer resistance and increased active sites of the basal plane.8,9 Previous synthetic approaches mainly focused on solid-state procedures10 or chemical vapor deposition (CVD) growth,11 whereas the direct wet-chemical protocols are quite rare.12–14 The wet-chemical synthesis methods hold promise for scale-up production and provide a versatile ground for chemical functionalization and hybridization with other materials, but the subtle chemistry and reactivity in solution-phase synthesis (e.g., precursor, surfactant, and temperature) are difficult to control.15,16 Until now, it remains a big challenge to synthesize 1T TMDCs-based nanostructures with high-purity crystal phases and well-controlled sizes and morphologies. Precise control over thre composition, spatial distribution of each domain, and interfaces between different components in hybrid nanostructures (HNs) has become the research focus recently, which facilitates advances in synthesis and assembly methodologies, as well as fundamental insights into the structure–property–function relationships.17–21 However, a systematic design of proper synthetic conditions for fabricating HNs with controlled physical parameters remains challenging. For instance, integrating materials of different dimensionalities (e.g., one-dimensional [1D]/two-dimensional [2D],22,23 2D/1D,24 and 2D/2D25 HNs) can generate high surface area and new electron states, but little progress has been made up till now, considering the limitation of materials choice with appropriate compatibility and the difficulty in combination of different growth modes. Moreover, it is even more challenging to construct the metastable-phase heterostructures because of their thermodynamically unstable nature, especially for 1T TMDCs, which easily undergo phase transition to more stable 2H phases under external disturbance.26 Thus, a general synthetic approach with delicate control over crystal phase and dimensionality in HNs is highly desirable, which will provide a controllable platform for exploring novel, optimized, or enhanced properties. Herein, we report a general solution protocol for phase-controlled synthesis of distorted octahedral 1T WS2-based multidimensional HNs from 2D, 1D, and zero-dimensional (0D) Ni3S4 templates (denoted as 1T WS2-2D/1D/0D Ni3S4 HNs), which is realized by tuning the reactivity of the tungsten precursor and the interaction between crystal surface and ligands. By changing the reaction temperature, the corresponding prismatic 2H WS2-based HNs (denoted as 2H WS2-2D/1D/0D Ni3S4 HNs) can also be achieved. This unique toolbox for the synthesis and formation process of monodispersed, uniform, and phase-controlled multidimensional 1T WS2-2D/1D/0D Ni3S4 HNs is demonstrated in Figure 1. As a proof-of-concept application, the crystal phase- and dimensionality-dependent electrocatalysis is demonstrated by evaluating their HER performances. Remarkably, the 1T WS2-2D Ni3S4 HNs exhibit a low overpotential of 109 mV at 10 mA cm−2, and show no evidence of deactivation after long-term operation. Our studies offer new perspectives for designing advanced materials based on crystal phase and dimensionality engineering strategies. Figure 1 | Schematic illustration of the general solution protocol for phase-controlled synthesis of 1T WS2-based multidimensional HNs from 2D, 1D, and 0D Ni3S4 templates. Download figure Download PowerPoint Experimental Methods Materials synthesis Ammonium tetrathiotungstate [(NH4)2WS4] was purchased from Alfa Aesar (Haverhill, MA). 1,2-Hexadecanediol (HDD, 98%) was purchased from Aladdin (Shanghai, China). Nickel(II) acetylacetonate [Ni(acac)2] (95%), oleylamine (OAm, 70%), oleic acid (OA, 90%), 1-octadecene (ODE, 90%), and 1-dodecanethiol (DDT, 98%) were purchased from Sigma-Aldrich (St. Louis, MO). NiCl2·6H2O, acetic acid, hexane, and ethanol were purchased from Sinopharm Chemical Reagent (Shanghai, China). All chemicals were used as received without further purification. In a typical procedure, 35 mg (NH4)2WS4 was added into 10 mL OAm in a 25 mL three-neck flask and heated to 120 °C for 20 min under vacuum. Then the flask was purged with N2, and 25 mg Ni(acac)2 dissolved in 5 mL hot OAm was injected. Then the solution was maintained at 150 °C for 30 min (formation of 2D Ni3S4) and slowly heated to 210 °C for 60 min. The mixture was cooled to room temperature. Finally, the obtained product was collected by centrifugation (8000 rpm, 3 min) and washed several times with hexane and ethanol for further characterization. As for 2H WS2-2D Ni3S4 HNs, the reaction temperature was increased to 310 °C. The 1D Ni3S4 and 0D Ni3S4 were synthesized according to reported methods.27 In a typical procedure for 1T WS2-1D/0D Ni3S4 HNs, 35 mg (NH4)2WS4 and 10 mg as-synthesized 1D/0D Ni3S4 were added into 10 mL OAm in a 25 mL three-neck flask and heated to 120 °C for 20 min under vacuum. Then the solution was slowly heated to 280 °C for 30 min under N2 atmosphere. Finally, the mixture was cooled to room temperature. As for 2H WS2-1D/0D Ni3S4 HNs, the reaction temperature was increased to 340 °C. The 1T and 2H WS2 were synthesized by modification of reported methods.28 In a typical procedure for 1T WS2, 35 mg (NH4)2WS4 was added into 10 mL OAm in a 25 mL three-neck flask and heated to 120 °C for 20 min under vacuum. Then the solution was slowly heated to 280 °C for 30 min under N2 atmosphere. Finally, the mixture was cooled to room temperature. As for 2H WS2, the reaction temperature was increased to 340 °C. Characterizations Powder X-ray diffraction (PXRD) was performed on a SmartLab 9kW (Rigaku, Tokyo, Japan)X-ray diffractometer equipped with graphite monochromaticized Cu Kα radiation (λ = 1.54056 Å). Transmission electron microscopy (TEM) images were employed on HT-7700 (Hitachi, Tokyo, Japan) with an acceleration voltage of 100 kV. High-resolution TEM (HRTEM) images and high-angle annular dark-field scanning TEM (HAADF-STEM) images were carried out on JEM-ARM 200F (JEOL, Tokyo, Japan). STEM energy-dispersive X-ray spectroscopy (STEM-EDS) analysis were conducted on Talos F200X (FEI, Hillsboro, OR). X-ray photoelectron spectra (XPS) were recorded on an ESCALab 250Xi (Thermo Scientific, Waltham, MA) X-ray photoelectron spectrometer using Al Kα radiation exciting source. Raman scattering spectra were recorded with a LabRAM HR Evolution (HORIBA Jobin Yvon, Paris, France) spectrometer using the 532 nm line. Fourier transform infrared (FT-IR) spectra were measured on a Nicolet 6700 (Thermo Nicolet, Wisconsin) FT-IR spectrometer. Optical absorption spectra were measured at room temperature using a DUV-3700 UV–vis–near-infrared (NIR) spectrometer (Shimadzu, Kyoto, Japan). Electrocatalytic measurements Electrochemical measurements were performed at room temperature using a rotating disk working electrode made of glassy carbon (GC) (PINE, 5 mm diameter, 0.196 cm2) connected to a Multipotentiostat (IMPex, ZAHNER-Elektrik, Germany). Graphite rod and Ag/AgCl (3.5 M KCl) were used as counter and reference electrodes, respectively. All reported potentials were converted to the reversible hydrogen electrode (RHE) through RHE calibration in 1 M KOH, E(RHE) = E(Ag/AgCl) + 1.03 V. The preparation method of the working electrodes containing various catalysts can be found as follows. In short, 5 mg of catalyst powder was dispersed in 1 mL of ethanol with 20 μL of Nafion solution (5 wt %; Sigma-Aldrich). Then the mixture was ultrasonicated for ∼30 min to generate a homogeneous ink. Next, 8 µL of the dispersion was transferred onto the GC disk. Finally, the as-prepared catalyst film was dried at room temperature. For comparison, a bare GC electrode that had been polished and cleaned was also dried for electrochemical measurement. Before electrochemical measurements, the electrolyte (1 M KOH) was bubbled with pure Ar for at least 30 min. The polarization curves were obtained by sweeping the potential from −0.8 to 0.2 V versus RHE at room temperature and 1600 rpm with a sweep rate of 5 mV s−1. All polarization curves arising from the solution resistance were iR-corrected. The electrochemical impedance spectroscopy (EIS) measurement was performed in the same configuration at 250 mV overpotential over a frequency range from 100 kHz to 100 mHz at the amplitude of the sinusoidal voltage of 5 mV. Cyclic voltammograms performed at different sweep rates used to estimate the double-layer capacitance (Cdl) were recorded in the potential region of 0.1–0.2 V versus RHE at room temperature. For the stability evaluation, electrode potential was cycled from −0.3 to 0.2 V versus RHE with a sweep rate of 100 mV s−1. At the end of the cycling experiment, the polarization curve was obtained with a sweep rate of 5 mV s−1. Chronoamperometry data were collected on various catalyst-coated carbon fiber papers (1 cm2, catalyst loading: 1 mg) at 250 mV overpotential. The polarization curves were replotted as η (η) versus log current (log j) to get Tafel plots. By fitting the linear portion of the Tafel plots to the Tafel equation (η = b log (j) + a), the Tafel slope (b) can be obtained. Density functional theory calculations We performed the density functional theory (DFT) calculations using the Vienna ab initio simulation package (VASP)29,30 program with the projector augmented wave (PAW)31,32 method. The Perdew–Burke–Ernzerhof (PBE)33 generalized gradient approximation (GGA) exchange–correlation functional was used throughout. A 500 eV plane-wave kinetic energy cutoff was chosen, and a 4 × 2 × 1 Monhorst–Pack k-point sampling was adopted for the structure relaxation. The convergence criterion for the electronic self-consistent iteration was set to be 10–4 eV. A residual force threshold of 0.02 eV Å−1 was set for geometry optimizations. The calculations were conducted on (0 0 1) surface of 1T and 2H WS2, and (0 1 1) surface of Ni3S4 slab models. A vacuum layer 15 Å thick was used to ensure the separation between slabs. The key reaction steps in alkaline HER: (1) H 2 O + e − + * → * H + OH − (Volmer step) (2) 2 * H → H 2 (Tafel step) (3) * H + H 2 O + e − → * + O H − + H 2 (Heyrovsky step) Here, * denotes the corresponding surface models. The free energy for steps (1) and (3) should be the same at the equilibrium potential of HER. Under this assumption, one can avoid computation of the exact free energy of OH− in solutions by using computational hydrogen electrodes.34 Herein, four main stages are considered: initial state, activated water adsorption, H* intermediates formation, and H2 desorption.35 The free energies are calculated as: G 0 = G ( * ) + G ( H 2 O ) G 1 = G ( * H − OH ) G 2 = G ( * H ) + G ( OH − ) G 3 = G ( * ) + 1 / 2 G ( H 2 ) + G ( OH − ) And G 0 = G 3 The Gibbs free energy changes are calculated as follows: Δ G 1 = G 1 − G 0 Δ G 2 = G 2 − G 1 Δ G 3 = G 3 − G 2 The free energy was calculated using the equation: G = E + E ZPE − T S Here, E, EZPE, T, and S are the total energy from DFT calculations, zero-point energy (ZPE), temperature (298 K), and the entropy, respectively. ZPE values are calculated from the frequency calculations. The entropic contributions from H2 and H2O molecules are taken from Chemical Rubber Company Press (CRC) Handbook. Results and Discussion Preparation and structural analyses of 1T WS2-2D Ni3S4 HNs We achieved the vertical growth of distorted octahedral 1T WS2 nanosheets on 2D Ni3S4 nanoplates (denoted as 1T WS2-2D Ni3S4 HNs) through successive thermolysis of Ni(acac)2 and (NH4)2WS4 metal precursors in OAm solution at 210 °C for 1 h. Simply increasing the reaction temperature to 310 °C leads to the formation of corresponding prismatic 2H WS2-based counterparts (denoted as 2H WS2-2D Ni3S4 HNs) (see Experimental Methods for details). This is caused by the different reactivity of the precursor and the interaction between crystal surface and ligands: a lower temperature benefits the deactivation of (NH4)2WS4, giving rise to the metastable 1T phase.12 Moreover, 1T-WS2 consists of negatively charged WS2− surfaces that stabilized by positively charged oleylammonium ligands, hence a lower temperature could strengthen their coordination effect, and consequently the stability.36 The low- and high-magnified transmission electron microscopy (TEM) images (Figures 2a and 2b) show uniform 2D/2D hierarchical architectures with the branch length of ∼40 nm, as well as the plate length and thickness of ∼60 and ∼8 nm respectively, inherited well from the original templates ( Supporting Information Figure S1). High-resolution TEM (HRTEM) images (Figures 2c and 2d) taken from the white solid squares (1 and 2 in Figure 2b) confirm the structural relationship of the interfaces between two distinct components, indicating the formation of crystalline polydymite Ni3S4 with resolved lattice fringes of (111) and (113) planes on lateral and basal sides, respectively. This analysis agrees with the X-ray diffraction (XRD) patterns ( Supporting Information Figure S2), in which cubic Ni3S4 with spinel structure (JCPDS no. 43-1469) can be observed.37 However, diffraction peaks from WS2 are relatively weak, probably owing to its ultrathin nature (white dotted lines and dotted circles in Figures 2c and 2d). Note that the main peak in the 32° region was observed to shift toward lower angles by ∼1.2°, indicating the formation of the unusual distorted octahedral 1T phase,13 which agrees with the XRD results of the as-synthesized 1T and 2H WS2 nanosheets ( Supporting Information Figure S3). HR-HAADF-STEM images (Figures 2e and 2f) exhibit typical 1T WS2 on Ni3S4 template in achieved 1T WS2-2D Ni3S4 HNs according to their zigzag chains of W atoms. By contrast, the 2H WS2-2D Ni3S4 HNs show typical 2H WS2 with hexagonal packing crystal structures. Moreover, the corresponding intensity profile analysis (Figure 2g) along the white solid rectangles (3 and 4 in Figures 2e and 2f) reveals the spatial location of W and S atoms, further confirming the formation of 1T and 2H phases.38 These results also match well with the HR-HAADF-STEM analysis of the as-synthesized 1T and 2H WS2 nanosheets ( Supporting Information Figure S4). STEM-EDS elemental mapping images (Figure 2h) of a typical 1T WS2-2D Ni3S4 HN show W-rich branches and Ni-rich plates with S enrichment in the whole structure. By comparison, the 2H WS2-2D Ni3S4 HNs show almost the same size, morphology, crystal phases of Ni3S4, and the distribution of Ni, S, and W elements ( Supporting Information Figure S5). Together, these results support that we have succeeded in phase-controllable synthesizing the new, uniform, and high-purity 1T/2H WS2-2D Ni3S4 HNs. Figure 2 | (a) Low- and (b) high-magnified TEM images of 1T WS2-2D Ni3S4 HNs. (c) Side and (d) top view HRTEM images of the white solid squares 1 and 2 in (b). (e and f) HR-HAADF-STEM images of interfaces between WS2 and Ni3S4 of 1T/2H WS2-2D Ni3S4 HNs, respectively. Inset: crystal models of 1T/2H WS2. Color scheme: W, blue; S yellow. (g) Corresponding intensity profile along the white solid rectangles 3 and 4 in (e and f), respectively. (h) STEM-EDS elemental mapping images of a typical 1T WS2-2D Ni3S4 HN, showing the distribution of Ni (green), S (yellow), and W (red). Download figure Download PowerPoint Preparation and structural analyses of 1T WS2-1D/0D Ni3S4 HNs We also achieved the growth of 1T WS2 nanosheets on 1D Ni3S4 nanorods and 0D Ni3S4 nanoparticles (denoted as 1T WS2-1D/0D Ni3S4 HNs) by gradual thermolysis of (NH4)2WS4 in OAm solution comprising the presynthesized 1D or 0D Ni3S4 templates ( Supporting Information Figures S6 and S7) at 280 °C for 30 min. Similarly, increasing the reaction temperature to 340 °C leads to the formation of corresponding 2H WS2-based counterparts (denoted as 2H WS2-1D/0D Ni3S4 HNs) (see Experimental Methods for details). TEM ( Supporting Information Figure S8) and HAADF-STEM images (Figures 3a and 3b) of 1T WS2-1D/0D Ni3S4 HNs show that the length and diameter of the nanorods were ∼40 and ∼10 nm, respectively, and the average length of the nanoparticles was ∼15 nm. HRTEM images demonstrate good crystallinity of a typical polydymite Ni3S4 nanorod and nanoparticle, with resolved lattice fringes of (113) and (022) planes, respectively. The white dotted circles suggest the ultrathin 1T WS2 interfaced with the Ni3S4 component (insets in Figures 3a and 3b). XRD patterns further confirm the diffraction peaks are indexed to the same cubic spinel-structured Ni3S4 (JCPDS no. 43-1469). But the almost negligible diffraction peaks from WS2 in the 32° region result from its low degree of crystallinity ( Supporting Information Figure S9). HR-HAADF-STEM images (Figures 3c and 3d) of 1T WS2-1D/0D Ni3S4 HNs exhibit 1T WS2 on both 1D and 0D Ni3S4 templates with zigzag-chained W atom arrangement, whereas 2H WS2-1D/0D Ni3S4 HNs display obvious 2H WS2 with uniform crystal structures of hexagonal packing ( Supporting Information Figure S10). Furthermore, the corresponding intensity profile analysis (Figure 3e) along the white solid rectangles (1 and 2 in Figures 3c and 3d) demonstrates the spatial location of W and S atoms, confirming the formation of 1T and 2H phases. STEM-EDS elemental mapping images (Figures 3f and 3g) of 1T WS2-1D/0D Ni3S4 HNs show W-rich layers and Ni-rich rods/particles with S enrichment in the whole structure. By comparison, the 2H WS2-1D/0D Ni3S4 HNs also show almost the same size, morphology, crystal phases of Ni3S4, and distribution of Ni, S, and W elements ( Supporting Information Figures S11–S13). Figure 3 | (a and b) HAADF-STEM images of 1T WS2-1D/0D Ni3S4 HNs, respectively. Inset: HRTEM images of typical 1T WS2-1D/0D Ni3S4 HNs, respectively. White dotted circles suggest the interfaces between 1T WS2 and Ni3S4. (c and d) HR-HAADF-STEM images of 1T WS2-1D/0D Ni3S4 HNs, respectively. Inset: crystal models of 1T WS2. Color scheme: W, blue; S, yellow. (e) Corresponding intensity profile along the white solid rectangles 1 and 2 in (c and d), respectively. (f and g) STEM-EDS elemental mapping images of typical 1T WS2-1D/0D Ni3S4 HNs, showing the distribution of Ni (green), S (yellow), and W (red). Download figure Download PowerPoint Formation mechanism To gain an insight into the formation mechanism of the 1T WS2-2D Ni3S4 HNs, a series of control experiments were performed. The ratio of precursors greatly affects their final morphologies ( Supporting Information Figures S14 and S15). The nucleation and growth of 1T WS2 are controlled through reaction temperature and time. The reaction between 150 and 180 °C quickly enables the Ni precursors to be confined in the soft organic–inorganic intermediates of OAm and tungsten precursors, giving rise to uniform 2D platforms for subsequent growth.39 Increasing reaction temperature to 210 °C leads to the formation and vertical growth of high-density starting nuclei into sheet-like structures via consumption of the tungsten precursors on preformed templates, which yields uniform three-dimensional (3D) hierarchical architectures composed of 2D 1T WS2 nanosheets and 2D Ni3S4 nanoplates ( Supporting Information Figure S16), similar to the previous reported 1T MoS2-2D NiS2 structures.40 As for 1T WS2-1D Ni3S4 HNs, the reaction between 220 and 250 °C indicates no sign for the formation of ultrathin sheet-like structures on surfaces of preformed templates. After the reaction temperature is increased to 280 °C, the formation and growth of starting nuclei into sheet-like structures were observed, which yielded uniform ultrathin curled layers composed of 2D 1T WS2 nanosheets and 1D Ni3S4 nanorods ( Supporting Information Figure S17). Crystal structures and characterization of 1T/2H WS2-2D Ni3S4 HNs Figure 4a illustrates the side (up) and top (bottom) views of the atomic crystal structures of 1T/2H WS2-Ni3S4 slab models, where the symmetry differences between the distorted octahedral coordinated 1T WS2 and the prismatic coordinated 2H WS2 (red dotted circles) can be clearly observed. We used XPS to study the chemical states of the 1T/2H WS2-2D Ni3S4 HNs. As shown in Figure 4b, W 4f peaks of 1T WS2-2D Ni3S4 HNs were located at 32 eV (4f7/2) and 34.1 eV (4f5/2), which both show a shift of ∼0.9 eV to lower binding energy compared with the 2H WS2-2D Ni3S4 HNs. This binding energy shift could be originated from the change in the metal coordination and the slightly increased W–S bond length. Meanwhile, the minor fractions of the 2H phase (32.7 and 34.8 eV) and traces of oxidized tungsten (35.6 and 37.8 eV) exist. We attributed this to the H2O generated from the condensation of OAm reacting with (NH4)2WS4, which might oxidize 1T WS2 or even damage the metastable 1T phase structure. Figure 4c shows that S 2p peaks of 1T WS2-2D Ni3S4 HNs were located at 161.4 (2p3/2) and 162.6 eV (2p1/2), which both show a shift of ∼1 eV to lower binding energy compared with the 2H WS2-2D Ni3S4 HNs. These W 4f and S 2p results are consistent with the XPS analysis for pure 1T and 2H WS2 ( Supporting Information Figure S18). Moreover, the binding energy of Ni 2p peaks (Figure 4d) for 1T WS2-2D Ni3S4 HNs markedly increases by ∼0.7 eV versus the 2H WS2-2D Ni3S4 HNs. This chemical shift observed for Ni 2p peaks is remarkable, which is probably due to strong charge transfer from Ni3S4 to 1T WS2 at the nanoscale interfaces.41 Figure 4 | (a) Side and top view atomic crystal structures of typical 1T/2H WS2-Ni3S4 slab models. Inset: Red dotted circles suggest the coordination environment of single W atom in 1T/2H WS2 structures, respectively. (b) W 4f, (c) S 2p, and (d) Ni 2p XPS spectra of 1T WS2-2D Ni3S4 (top) and 2H WS2-2D Ni3S4 (bottom) HNs. Doublets of (b and c) corresponding to 1T WS2 are given in blue, 2H WS2 in green, and oxidized tungsten in pink. Doublets of (d) corresponding to Ni2+ of 1T WS2-2D Ni3S4 HNs are given in blue, Ni2+ of 2H WS2-2D Ni3S4 HNs in green; dark cyan and purple doublets are assigned to Ni3+ and concomitant shakeup satellites. (e) Raman and (f) UV–vis absorption spectra of 1T/2H WS2-2D Ni3S4 HNs, respectively. Download figure Download PowerPoint Raman spectrum of 1T WS2-2D Ni3S4 HNs shows a clearly distinct set of vibration m

Morphology and structure characteristics of nanoscale carbon materials containing graphene

A comprehensive study of the nanostructured powders (graphite GSM-2; Taunit-M; thermally expanded graphite (TEG)) by methods of transmission electron microscopy (TEM), scanning electron microscopy (SEM), X-ray diffractometry (XRD), reflection high-energy electron diffraction (RHEED), Raman spectroscopy, was carried out. The experimental XRD halo was interpreted by superimposing theoretical diffraction maxima, and an X-ray amorphous graphite phase was revealed. It was found that the X-ray amorphous phase is characterized by the limiting degree of graphite nanostructuring. From the width of the diffraction rings, the maximum sizes of graphite nanocrystals were estimated, which do not exceed 5 and 10 nm in the [0001] and [ ] directions, respectively. Carbon nanotubes and plates of turbostratic graphene were revealed. The structural and morphological parameters of the nanostructured material “Taunit-M” have been established – multi-walled nanotubes with a diameter of up to 10 nm are combined through an interlayer of X-ray amorphous carbon into flat ribbons up to 40 nm wide. Dark-field TEM images (in reflections of ) revealed moiré patterns that appear on overlapping graphene sheets due to double diffraction of the electron beam. It was found that in thermally expanded graphite, the rotation of graphene sheets ranges from 3 to 4°. Within the graphene sheets, complete dislocations with the Burgers vector b = 1/2 were revealed [1010]. The Fourier analysis of moiré images made it possible to determine the mutual orientation of graphene sheets, to reveal regions of multilayer graphene, and to identify turbostratic graphene. It is shown that the combination of RHEED, TEM, and Fourier transformations of periodic contrast of electron microscopic images is a promising approach to the analysis of the substructure and morphology of nanoscale carbon materials containing graphene and other allotropic modifications of carbon.