Advanced Materials

Published by Wiley
Online ISSN: 1521-4095
Discipline: Materials Science
Learn more about this page
Aims and scope

Advanced Materials has been bringing you the latest progress in materials science every week for over 30 years. Read carefully selected, top-quality Research Articles, Reviews, and Perspectives at the cutting edge of the chemistry and physics of functional materials.



Recent publications
  • Yichao TangYichao Tang
  • Mingtong LiMingtong Li
  • Tianlu WangTianlu Wang
  • [...]
  • Metin SittiMetin Sitti
Wireless miniature soft actuators are promising for various potential high‐impact applications in medical, robotic grippers, and artificial muscles. However, these miniature soft actuators are currently constrained by a small output force and low work capacity. To address such challenges, we report a miniature magnetic phase‐change soft composite actuator. This soft actuator exhibits an expanding deformation and enables up to a 70 N output force and 175.2 J/g work capacity under remote magnetic radio frequency heating, which are 106–107 times that of traditional magnetic soft actuators. To demonstrate its capabilities, we first design a wireless soft robotic device that can withstand 0.24 m/s fluid flows in an artery phantom. By integrating it with a thermally‐responsive shape‐memory polymer and bistable metamaterial sleeve, we design a wireless reversible bistable stent towards for future potential angioplasty applications. Moreover, it can additionally locomote inside and jump out of granular media. At last, the phase‐change actuator can realize programmable bending deformations when a specifically designed magnetization profile is encoded, enhancing its shape‐programming capability. Such a miniature soft actuator provides an approach to enhance the mechanical output and versatility of magnetic soft robots and devices, extending their medical and other potential applications. This article is protected by copyright. All rights reserved
Chiral metasurfaces can exhibit a strong circular dichroism, but it is limited by the complicated fabrication procedure and alignment errors. Here, we demonstrate a new type of self‐aligned suspended chiral bilayer metasurface with only one‐step electron beam lithography exposure. A significant optical chirality of 221°/μm can be realized using suspended metasurfaces with a thickness of 100 nm. Furthermore, we experimentally demonstrate that such a structure is capable of label‐free discrimination of the chiral molecules at zeptomole level, exhibiting a much higher sensitivity (orders of magnitude) compared to the conventional circular dichroism spectroscopy. The fundamental principles for chiral sensing using molecules‒metasurfaces interactions have been explored. Benefiting from the giant chiroptical response, our proposed meta‐device may offer promising applications for ultrathin circular polarizers, chiral molecular detectors, and asymmetry information processing. This article is protected by copyright. All rights reserved
Growth of Cu(In,Ga)Se2 (CIGS) absorbers under Cu‐poor conditions gives rise to incorporation of numerous defects into the bulk, whereas the same absorber grown under Cu‐rich conditions leads to a stoichiometric bulk with minimum defects. This suggests that CIGS absorbers grown under Cu‐rich conditions are more suitable for solar cell applications. However, the CIGS solar cell devices with record efficiencies have all been fabricated under Cu‐poor conditions, despite the expectations. Therefore, in the present work, we investigate both Cu‐poor and Cu‐rich CIGS cells and show that the superior properties of internal interfaces of Cu‐poor CIGS cells, such as p‐n junction and grain boundaries, makes them always the record‐efficiency devices. More precisely, by employing a correlative microscopy approach, we discover for the first time the typical fingerprints for superior properties of internal interfaces necessary for maintaining a lower recombination activity in the cell. These are a Cu‐depleted and Cd‐enriched CIGS absorber surface near the p‐n junction and the negative Cu factor (∆β) and high Na content (> 1.5 at.%) at the grain boundaries. Thus, this work provides key factors governing the device performance (efficiency), which can be considered in the design of next‐generation solar cells. This article is protected by copyright. All rights reserved
Ga2O3 and its polymorphs are attracting increasing attention. The rich structural space of polymorphic oxide systems such as Ga2O3 offers potential for electronic structure engineering, which is of particular interest for a range of applications, such as power electronics. γ‐Ga2O3 presents a particular challenge across synthesis, characterisation, and theory due to its inherent disorder and resulting complex structure – electronic structure relationship. Here, density functional theory is used in combination with a machine learning approach to screen nearly one million potential structures, thereby developing a robust atomistic model of the γ‐phase. Theoretical results are compared with surface and bulk sensitive soft and hard X‐ray photoelectron spectroscopy, X‐ray absorption spectroscopy, spectroscopic ellipsometry, and photoluminescence excitation spectroscopy experiments representative of the occupied and unoccupied states of γ‐Ga2O3. The first onset of strong absorption at room temperature is found at 5.1 eV from spectroscopic ellipsometry, which agrees well with the excitation maximum at 5.17 eV obtained by PLE spectroscopy, where the latter shifts to 5.33 eV at 5 K. This work presents a leap forward in the treatment of complex, disordered oxides and is a crucial step towards exploring how their electronic structure can be understood in terms of local coordination and overall structure. This article is protected by copyright. All rights reserved
Noncovalent macrocycle‐confined supramolecular purely organic room‐temperature phosphorescence (RTP) is the current research hotspot. Herein, we report a high‐efficiency noncovalent polymerization activated near‐infrared (NIR) emissive RTP‐harvesting system in aqueous solution based on the stepwise confinement of cucurbit[7]uril (CB[7]) and β‐cyclodextrin grafted hyaluronic acid (HACD). Compared with dodecyl chain bridged 6‐bromoisoquinoline derivative (G), the dumbbell‐shaped assembly G⊂CB[7] presents an appeared complexation‐induced RTP signal at 540 nm via the first confinement of CB[7]. Subsequently, benefitting from the stepwise confinement encapsulation of β‐cyclodextrin cavity, the subsequent noncovalent polymerization of the binary G⊂CB[7] assembly enabled by HACD can contribute to the further enhanced RTP emission intensity approximately 8 times in addition to an increased lifetime from 59.0 µs to 0.581 ms. Moreover, upon doping a small amount of two types of organic dyes, Nile blue (NiB) or Tetrakis(4‐sulfophenyl)porphyrin (TPPSS) as an acceptor into the supramolecular confinement assembly G⊂CB[7]@HACD, efficient RTP energy transfer occurs accompanied by a long‐lived NIR emitting performance (680, 710 nm) with a high donor/acceptor ratio. Intriguingly, the prepared RTP‐harvesting system is successfully applied for targeted NIR imaging of living tumor cells by utilizing the targeting ability of hyaluronic acid, which provides a new strategy to create advanced water‐soluble NIR phosphorescent materials. This article is protected by copyright. All rights reserved
  • Guiying HeGuiying He
  • Kaia R. ParentiKaia R. Parenti
  • Luis M. CamposLuis M. Campos
  • Matthew Y. SfeirMatthew Y. Sfeir
Singlet fission is commonly defined as the generation of two triplet excitons from a single absorbed photon. However, ambiguities within this definition arise due to the complexity of the various double triplet states that exist in SF chromophores and corresponding interconversion processes. To clarify this process, singlet fission is frequently depicted as sequential two‐step conversion in which a singlet exciton decays into a bound triplet pair biexciton state which dissociates into two “free” triplet excitons. However, this model discounts the potential for direct harvesting from the coupled biexciton state. Here, we demonstrate that individual triplet excitons can be extracted directly from a bound triplet pair. We demonstrate that due to the requirement for geminate triplet‐triplet annihilation in intramolecular singlet fission compounds, unique spectral and kinetic signatures can be used to quantify triplet pair harvesting yields. We achieve an internal quantum efficiency for triplet exciton transfer from the triplet pair of greater than 50%, limited only by the solubility of the compounds. The harvesting process is not dependent on the net multiplicity of the triplet pair state, suggesting that an explicit, independent dissociation step is not a requirement for using triplet pairs to do chemical or electrical work. This article is protected by copyright. All rights reserved
  • Siyoung LeeSiyoung Lee
  • Hajung RohHajung Roh
  • Junsoo KimJunsoo Kim
  • [...]
  • Kilwon ChoKilwon Cho
Auditory sensors have shortcomings with respect to not only personalization with wearability and portability but also detecting a human voice clearly in a noisy environment or when a mask covers the mouth. In this work, we exploited an electret‐powered and hole patterned polymer diaphragm into a skin‐attachable auditory sensor for the first time. The optimized charged electret diaphragm induces a voltage bias of >400 V against the counter electrode, which reduces the necessity of a bulky power source and enables the capacitive sensor to show high sensitivity (2.2 V/Pa) with incorporation of an elastomer nanodroplet seismic mass. The sophisticated capacitive structure with low mechanical damping enabled a flat frequency response (80–3000 Hz) and good linearity (50–80 dBSPL). The hole‐patterned electret diaphragms help our skin‐attachable sensor detect only neck‐skin vibration rather than dynamic air pressure, enabling a person's voice to be detected in a harsh acoustic environment. The sensor operated reliably even in the presence of surrounding noise and when the user was wearing a gas mask. Therefore, our sensor shows strong potential of a communication tool for disaster response and quarantine activities, and of diagnosis tool for vocal healthcare applications such as cough monitoring and voice dosimetry. This article is protected by copyright. All rights reserved
  • Shiyong LiShiyong Li
  • Ye WangYe Wang
  • Miaojin WuMiaojin Wu
  • [...]
  • Weibo CaiWeibo Cai
During cerebral ischemia‐reperfusion (I‐R) injury, the infiltration of monocyte/macrophages (Mo/Mφ) into the ischemic penumbra causes inflammatory damage but also regulates tissue repair in the penumbra. The regulation and balance of Mo/Mφ polarization has been considered as a potential therapeutic target for treating cerebral I‐R injury. Herein, our findings demonstrated that glabridin (Gla)‐loaded nanoparticles (i.e., NPGla‐5k) could effectively inhibit M1‐polarization and enhance M2‐polarization of Mo/Mφ. Positron emission tomography (PET) imaging showed that NPGla‐5k could selectively accumulate in the spleen following intravenous injection. Spleen‐targeted Cy5‐NPGla‐5k can colocalize with peripheral macrophages in the penumbra at 24 h after tail vein injection. Interestingly, NPGla‐5k treatment could reduce inflammatory damage, protect dying neurons and improve nervous system function. The protective effect of spleen‐targeted NPGla‐5k against cerebral I‐R injury in mice encourages an exploration of their use for clinical treatment of patients with cerebral I‐R injury. This article is protected by copyright. All rights reserved
  • Fang WangFang Wang
  • Zhiyi LiuZhiyi Liu
  • Tao ZhangTao Zhang
  • [...]
  • Weida HuWeida Hu
Room‐temperature‐operating highly sensitive mid‐wave infrared (MWIR) photodetectors are utilized in a large number of important applications including night vision, communications, and optical radar. Many previous studies have demonstrated uncooled MWIR photodetectors using two‐dimentional narrow‐bandgap semiconductors. To date, most of these works utilize atomically thin flakes, simple van der Waals (vdW) heterostructures, or atomically thin p‐n junctions as absorbers, which have difficulty meeting the requirements for state‐of‐the‐art MWIR photodetectors with a blackbody response. Here, we report a fully depleted self‐aligned MoS2‐BP‐MoS2 vdW heterostructure sandwiched between two electrodes. This new type of photodetector exhibits competitive performance, including a high blackbody peak photoresponsivity up to 0.77 A/W and low noise‐equivalent power of 2.0 × 10−14 W/Hz1/2, in the MWIR region. A peak specific detectivity of 8.61 × 1010 cmHz1/2/W under blackbody radiation was achieved at room temperature in the MWIR region. Importantly, the effective detection range of our device is twice that of state‐of‐the‐art MWIR photodetectors. Furthermore, the device presents an ultra‐fast response of ∼4 μs both in the visible and short‐wavelength infrared bands. These results provide an ideal platform for realizing broadband and highly sensitive room‐temperature MWIR photodetectors. This article is protected by copyright. All rights reserved
Fabrication and microstructure of nanoMOF aerogels. a) Schematic illustration of the formation of zeolitic imidazolate framework‐8 (ZIF‐8) particles via self‐assembly of 2‐methylimidazole (2‐MiM) linkers and tetrahedrally coordinated Zn²⁺ ions, with and without the use of cetrimonium bromide (CTAB) for nanoMOF stabilization. b) Representative scanning electron microscopy (SEM) image of micrometer‐sized ZIF‐8 particles formed without CTAB; the inset shows the partial sedimentation of the dispersion after 2 h. c) Representative SEM image of nanoZIF‐8 (nZIF‐8) particles formed with CTAB; the inset shows a stable dispersion with no sedimentation after 2 h. d) Schematic image of the hierarchical structure of wood: from timber, via fibers, to nanofibrils and cellulose polymer chains. e) Representative atomic force microscopy image of the cationic cellulose nanofibrils (CNFs). f) Illustration of the formation mechanism of anisotropically frozen hydrogels and a representative photograph of the aerogels, composed of 90 wt% nZIF‐8 and 10 wt% CNFs, achieved by lyophilization. g,h) SEM images of the aerogels imaged in the transverse (g) and longitudinal (h) direction. i) High‐magnification SEM image of a porous part of the pore wall shown in (h).
Mechanical properties of the MOF‐based aerogels and adsorption capacity of anionic dyes. a,b) Compressive stress–strain curves for AGnZIF‐8, compressed in the longitudinal (a) and transverse (b) directions of the ice growth during freeze casting. c) Specific modulus of aerogels as a function of MOF content; data from our study and the literature.[9,20] d) Cyclic compressions (five cycles) to 50% compression of the original size of an AGnZIF‐8. e) Dye adsorption capacity of AGnZIF‐8 aerogels in 0.05 mg g⁻¹ dye solutions of brilliant blue (BB) and Congo red (CR) at pH 7.4, with photographs taken immediately after submersion and after 336 h. f) SEM image of the MOF‐covered aerogel pore walls, before and after 366 h of BB sorption, showing the intact microstructure of the MOFs after dye adsorption.
Gas adsorption and separation performance and flame‐retardant properties. a) Schematic image of how the AGnZIF‐8 aerogels selectively adsorb and hinder the transmission of CO2 while allowing the passage of CH4. b) Gravimetric equilibrium adsorption isotherms of CO2 and CH4 as a function of gas pressure at 293 K. c) Cyclic equilibrium CO2 adsorption isotherms in pristine nZIF‐8 powder and in AGnZIF‐8 at 293 K. d) Cyclic equilibrium CO2 adsorption isotherms in AGnZIF‐8 at 293 K. e) CO2 column breakthrough measurement of AGnZIF‐8 at 293 K and 2 bars with a flow direction parallel to the direction of freezing. f) Horizontal flame burning test of pure CNF aerogels (AGCNF) and AGnZIF‐8. The panels show photos of the samples at the start of the flame test, 2 or 5 s after flame removal, and the residues at the end of the test. g) Photographs of AGCNF, nZIF‐8, and AGnZIF‐8 samples before the cone calorimetry test (top row) and after a heat flux exposure of 35 kW m⁻² (bottom row). h) Proposed schematic mechanism[³²] for the self‐extinguishing behavior of the AGnZIF‐8.
Carbonization characteristics and charge storage for the carbonized aerogels. a) Photographs and schematic drawings of non‐carbonized and carbonized AGnZIF‐8 aerogels (CAGs) at 900 °C (CAG⁹⁰⁰). Scale bars: 5 mm. b) SEM and c) transmission electron microscopy (TEM) images of CAG⁹⁰⁰. d) Raman spectra of AGnZIF‐8 carbonized at 800 (CAG⁸⁰⁰), 900 (CAG⁹⁰⁰), and 1000 (CAG¹⁰⁰⁰) °C. e,f) XPS C1s spectra of AGCNF (e) carbonized at 900 °C (CCNF⁹⁰⁰) and CAG⁹⁰⁰ (f), as well as g,h) N1s spectra of CCNF⁹⁰⁰ (g) and CAG⁹⁰⁰ (h). i) Cyclic compressive testing (three cycles) to 50% compression of the original size of a CAG⁹⁰⁰. j) Cyclic voltammetry curves of CAGs at a scan rate of 50 mV s⁻¹ and k) Ashby plot of the specific modulus (transverse and longitudinal) and specific capacitance of CAGs at 800, 900, and 1000 °C, compared with those of other reported conductive aerogel electrodes.[⁴⁵]
Metal‐organic frameworks (MOFs) are hybrid porous crystalline networks with tunable chemical and structural properties. However, their excellent potential is limited in practical applications due to their hard‐to‐shape powder form, making it challenging to assemble MOFs into macroscopic composites with mechanical integrity. While a binder matrix enables hybrid materials, such materials have a limited MOF content and thus limited functionality. To overcome this challenge, nanoMOFs are combined with tailored same‐charge high‐aspect‐ratio cellulose nanofibrils (CNFs) to manufacture robust, wet‐stable, and multifunctional MOF‐based aerogels with 90 wt% nanoMOF loading. The porous aerogel architectures show excellent potential for practical applications such as efficient water purification, CO2 and CH4 gas adsorption and separation, and fire‐safe insulation. Moreover, a one‐step carbonization process enables these aerogels as effective structural energy‐storage electrodes. This work exhibits the unique ability of high‐aspect‐ratio CNFs to bind large amounts of nanoMOFs in structured materials with outstanding mechanical integrity – a quality that is preserved even after carbonization. The demonstrated process is simple and fully discloses the intrinsic potential of the nanoMOFs, resulting in synergetic properties not found in the components alone, thus paving the way for MOFs in macroscopic multifunctional composites. This article is protected by copyright. All rights reserved
Demonstration of key parameters for the quantification of adsorbed proteins using DDM. A) The workflow to assess proteins adsorbed on NPs starts by mixing the NPs with labeled proteins and letting them adsorb. The suspension is then transferred into a glass capillary for fluorescence microscopy imaging. Time series of images are then analyzed following the process described in the text to extract the 2D Fourier transform and the DICF. B) Examples of DICF g(q,τ) obtained at different protein concentrations showing the values of parameters ANP(q) and B*(q) used for the quantification of the amount of protein adsorbed, where red solid curves are fits to Equation (3)  f(q,τ) = exp (−Dq²τ) for diffusive particles. C) Experimental evidence showing that the parameter Aratio, which is directly proportional to the amount of adsorbed protein, is independent of the spatial frequency q and can therefore be averaged over q. D) Experimental validation of the linear relationship between Aratio and the concentration of particles in the medium as predicted by Equation (4), represented by the red line. E) Validation of the power‐law dependence (black lines) of Aratio with the average total fluorescence intensity, independently of the frame rate used for the acquisition.
Quantification of LYZ on NPs, aggregation and fractal dimension. A) Adsorption isotherm and particle size measurement of PS NPs and LYZ mixtures. In the strong aggregation regime highlighted in gray, aggregates fractal dimension was quantified simultaneously by DDM after 20–30 min of incubation time. Measurements of protein adsorption (nb = number) by DDM are compared with the adsorption measurement from the quantification of free proteins by fluorescence spectroscopy after centrifugation. B) Protein adsorption triggers NPs aggregation. The aggregates fractal dimension, df, can be monitored by DDM. Curves in (i) are fits to the Fisher–Burford model from Equation (8). C) Quantification of adsorbed protein amounts by DDM shows excellent agreement with PALS providing an additional validation of the quantification method of DDM. The red line is a guide to the eye. D) Adsorption isotherms of LYZ on PS NPs at different particle concentrations were well described by the Hill adsorption model with a Hill coefficient n = 1 (dashed curves). E) Adsorption isotherms shown in D normalized by the concentration of NPs collapse into a single master curve and fitted by a modified Langmuir isotherm Γ=ΓmaxCp/CNPKD+Cp/CNP$\Gamma = \frac{{{\Gamma _{\max }}{C_{\rm{p}}}{\rm{/}}{C_{{\rm{NP}}}}}}{{{K_{\rm{D}}} + {C_{\rm{p}}}{\rm{/}}{C_{{\rm{NP}}}}}}$. F) Dilution of proteins/NPs mixture shows that the amount of adsorbed proteins after dilution is always consistent with the amount expected by an adsorption isotherm performed at the same dilution. Dashed lines are the fitted isotherms from (D). DDM data in (A–F) are averages of at least five acquisitions and the error bars represent one standard deviation.
Quantification of adsorbed serum proteins. A) Adsorption isotherm, particle size, and fractal dimension of PS NPs/serum mixtures show multiple saturation plateaus reminiscent of significant variations in the PC composition along the isotherm. “Cst” represents the parameter kept constant in the adsorption experiment. B) The BSA adsorption isotherm can be described by the Hill isotherm and exhibits strong aggregation as all the other proteins tested. The 4‐most dilute experiments (100 > CP/CNP) were tested at CNP = 425 pm, whereas all other experiments were tested at CNP = 42.5 pm. C) Superposition of the serum and BSA adsorption isotherms shows strong overlapping in the high concentration region. In that region, the serum PC is expected to be largely composed of BSA. The lines in the upper panels are fits to a sum of two expressions of Equation (9) in (A) and to Equation (9) in (B) that were also re‐used in (C). The lines in the lower panels for size measurements are guide to the eye. DDM data in the figure are averages of at least five videos and error bars represent one standard deviation.
Kinetics of the PC formation in vitro. A) Protein exchange at the PC interface can be performed by DDM in different configurations to test the reversibility of protein adsorption. Changes in adsorbed labeled protein is modeled using Γ(t) = Γeq + ΔΓexp(−kt), with free parameters Γeq and ΔΓ are the adsorbed amount at equilibrium and the change of the adsorbed amount of labeled proteins, respectively. B) Monitoring of protein exchange at the PC interface in presence of lysozyme at different concentrations of labeled proteins. C) Evolution of the adsorption rate constant with the protein content allows to extract adsorption constants. The red line is a linear fit to the experimental data points, and the error bars are the parameters’ fitting standard errors. DDM data in (A) and (B) represent one measurement and the error bars represent the standard deviation of the 〈Aratio〉q.
In vivo evaluation of the PC kinetics in zebrafish larvae. A) Imaging of NPs in the zebrafish larva (48 hpf) post injection show rapid and stable immobilization of the NPs on the lining of the blood capillaries of the subcaudal region (i). In the right panels, the normalized average fluorescence intensity captured by the camera (ii) and the time evolution of the PC composition of NPs immobilized in the CVP area of the zebrafish larvae in vivo (iii). Data in (ii) and (iii) are cluster averages over n = 3 independent experiments (see Figure S10, Supporting Information) and the error bars their associated standard deviation. B) Dynamics of the immobilized NPs show a single process on the whole range of spatial frequency (i) where its normalization by τq² leads to a master curve demonstrating the diffusive nature of the motion of captured NPs (ii). The extracted coefficient of the immobilized NPs for many measurements (shown in the right panel) is consistent with the NPs being adsorbed at the cells surface but not internalized (iii). Error bars denote three standard errors to the data distribution.
Here, we propose a new theoretical framework that enables the use of Differential Dynamic Microscopy (DDM) in fluorescence imaging mode to quantify in situ protein adsorption onto nanoparticles (NP) while simultaneously monitoring for NP aggregation. We use this methodology to elucidate the thermodynamic and kinetic properties of the protein corona (PC) in vitro and in vivo. Our results show that protein adsorption triggers particle aggregation over a wide concentration range and that the formed aggregate structures can be quantified using the proposed methodology. Protein affinity for polystyrene (PS) NPs was observed to be dependent on particle concentration. For complex protein mixtures, our methodology identifies that the PC composition changes with the dilution of serum proteins, demonstrating a Vroman effect never quantitatively assessed in situ on NPs. Finally, DDM allows monitoring of the evolution of the PC in vivo. Our results show that the PC composition evolves significantly over time in zebrafish larvae, confirming the inherently dynamic nature of the PC. The performance of the developed methodology allowed to obtain quantitative insights into nano‐bio interactions in a vast array of physiologically relevant conditions that will serve to further improve the design of nanomedicine. This article is protected by copyright. All rights reserved
Nanostructure engineering is a key strategy for tailoring properties in the fields of batteries, solar cells, thermoelectrics, and so on. Limited by grain coarsening, however, the nanostructure effect gradually degrades during the materials’ manufacturing and in‐service period. Herein, we develop a strategy of cleavage‐fracture for grain shrinking in the Pb0.98Sb0.02Te sample during the sintering, and the grain size remains stable after repeated tests. Moreover, the initial grain boundary is filled by fractured slender grains and enriched by dislocations, evolving into a hierarchical grain‐boundary structure. The lattice thermal conductivity (κlat) is greatly reduced to approach the amorphous limit. As a result, a record‐high ZT value of about 1.9 is obtained at 815 K in the n‐type Pb0.98Sb0.02Te sample and a decent efficiency of 6.7% in thermoelectric device. This strategy for grain shrinking will shed light on the application of nanostructure engineering under high temperature and extreme conditions in other material systems. This article is protected by copyright. All rights reserved
A scalable synthetic route to colloidal atoms has significantly advanced over the past two decades. Recently, colloidal clusters with DNA‐coated cores called “patchy colloidal clusters” have been developed, providing a directional bonding with specific angle of rotation due to the shape complementarity between colloidal clusters. Through a DNA‐mediated interlocking process, they were directly assembled into low‐coordination colloidal structures, such as cubic diamond lattices. This review details the significant progress in recent years in the synthesis of patchy colloidal clusters and their assembly in experiments and simulations. Furthermore, an outlook is given on the emerging approaches to the patchy colloidal clusters and their potential applications in photonic crystals, metamaterials, topological photonic insulators, and separation membranes. This article is protected by copyright. All rights reserved
a) Analytical phenomenological modeling‐based temperature–strain–strontium content (x) 3D‐phase diagram illustrating the stability regimes for paraelectric, c, c/a, and a1/a2 domain structures for Pb1−xSrxTiO3 thin films. The mixed‐phase domain region where c/a and a1/a2 domains coexist is calculated by phase‐field simulations and is shown in the yellow color at x = 0.0, 0.1, 0.2, 0.3, 0.4 and 0.5. Strain position of the DyScO3 (110) is schematically shown as a plane. b) Analytical phenomenological modeling‐based temperature–strontium content (x) phase diagram of a Pb1−xSrxTiO3 thin film on a DyScO3 (110) substrate. The mixed‐phase region was calculated using phase‐field simulations and is shown in yellow. c–f) The mixed‐phase region shows a gradual transition in the domain architecture as a function of strontium content (x) such as from small droplets of a1/a2 domains inside a c/a domain matrix at x = 0.15 (c), to a connected‐labyrinth structure of c/a domains at x = 0.19 (d), to a disconnected‐labyrinth structure at x = 0.21 (e), and finally to c/a droplets in an a1/a2 matrix at x = 0.275 (f).
a) θ–2θ X‐ray diffraction scans for 100 nm‐thick Pb1−xSrxTiO3 thin films (with, from top to bottom, x = 0, 0.06, 0.17, 0.20, 0.23, 0.25, 0.27, and 0.30) grown on DyScO3 (110) substrates. b–e) AFM topography images of the Pb1−xSrxTiO3 thin films showing a1/a2 droplets in a c/a matrix for x = 0.06 heterostructures (b), a connected‐labyrinth structure for x = 0.17 heterostructures (c), a disconnected‐labyrinth structure for x = 0.23 heterostructures (d), and c/a droplets in an a1/a2 matrix for x = 0.27 heterostructures (e). f) RSM about the 220‐diffraction peak of the DyScO3 (110) substrate revealing 002‐ and 200‐diffraction peaks for the Pb1−xSrxTiO3 thin films corresponding to the tilted c and tilted/untilted a domain variants, respectively, for an x = 0.20 heterostructure. g) Schematic of the domain tilting phenomenon observed in the c/a domain stripes with characteristic angles defined. h) Measured tilt angles for the c and a domains and the net tilt angle between them as a function of film composition.
a–d) From left to right, a binarized image used for analysis based on the experimental AFM data, the AFM topography image, the corresponding phase‐field simulation, and a binarized image used for analysis based on the phase‐field simulation for Pb1−xSrxTiO3 thin films with x = 0.06 (a), x = 0.17 (b), x = 0.23 (c), and x = 0.27 (d) (experimental x values noted here, but corresponding phase‐field x values are noted on the figure). e) Calculated normalized Euler characteristic (χN) versus strontium content (x) for the experimental part (red circles) and phase‐field generated (blue squares). f) Derivative of the piecewise fitted curve of the normalized Euler characteristic (χN) for the experimental and simulated data. Extrema of the derivative of the experimental curve (red line) are used to demarcate various topological regions observed in the Pb1−xSrxTiO3 thin films.
a–c) AFM topography images of 100 nm‐thick Pb0.83Sr0.17TiO3 thin films grown on DyScO3 (110) substrates in the as‐grown state (a) and in the same area after annealing at 270 °C for 24 h (b) and 190 h (c). d–f) Phase‐field simulations showing similar domain alignment phenomenon in the connected‐labyrinth structure for the as‐grown state (d) and after annealing for 300 000 (e) and 600 000 (f) timesteps. g–n) Various types of defects are observed in the Pb0.83Sr0.17TiO3 thin films such as threefold junctions (g), stripe end‐points (h), fourfold junctions (saddle points) (i), handle defects (j), disconnected stripes (k), droplets (l), target skyrmions (m), and dislocations (n). o) The defect densities of three‐ and four‐fold junctions (Dtype1) and the stripe end‐points (Dtype2) as a function of annealing time. The insets show the dislocation and fourfold defect types.
a) Dielectric permittivity (εr) as a function of strontium content (x). b) Shared perimeter between c/a and a1/a2 domains normalized by majority phase as a function of strontium content (x). c) Dielectric permittivity as a function frequency measured at different applied background DC electric fields for x = 0.00 heterostructure. d) Intercepts obtained from a linear logarithmic fit of the dielectric permittivity−frequency plot at various applied background DC electric fields are plotted as a function of the applied DC electric fields for x = 0.00 heterostructure. e) Zero‐field permittivity intercept of Intrinsic‐only, combined intrinsic–extrinsic contribution, and the difference between them (extrinsic‐only contribution) is plotted as a function of strontium content (x). f) Phase‐field simulations for permittivity coefficient (k33) for mixed‐phase Pb1−xSrxTiO3 thin film with x = 0.1875.
The potential for creating hierarchical domain structures, or mixtures of energetically degenerate phases with distinct patterns that can be modified continually, in ferroelectric thin films offers a pathway to control their mesoscale structure beyond lattice-mismatch strain with a substrate. Here, it is demonstrated that varying the strontium content provides deterministic strain-driven control of hierarchical domain structures in Pb1-x Srx TiO3 solid solution thin films wherein two types, c/a and a1 /a2 , of nanodomains can coexist. Combining phase-field simulations, epitaxial thin-film growth, and detailed structural, domain, and physical-property characterization, it is observed that the system undergoes a gradual transformation (with increasing strontium content) from droplet-like a1 /a2 domains in a c/a domain matrix, to a connected-labyrinth geometry of c/a domains, to a disconnected labyrinth structure of the same, and, finally, to droplet-like c/a domains in an a1 /a2 domain matrix. A relationship between the different mixed-phase modulation patterns and its topological nature is established. Annealing the connected-labyrinth structure leads to domain coarsening forming distinctive regions of parallel c/a and a1 /a2 domain stripes, offering additional design flexibility. Finally, it is found that the connected-labyrinth domain patterns exhibit the highest dielectric permittivity. This article is protected by copyright. All rights reserved.
a) Diagram of anatase TiO2 nanosheet with dominant {001} facets. b) Atomic structure of (001) surface of the top view. c) Standard diffraction pattern of anatase TiO2 in [001] direction. d) Diagram of anatase TiO2 nanosheet with dominant {111} facets. e) Atomic structure of (111) surface of the top view. f) Standard diffraction pattern of anatase TiO2 in [111] direction.
Structural characterization of anatase TiO2 nanosheets. a) HRTEM image of nanosheets in [001] orientation. b) FFT pattern from (a). c) HAADF‐STEM image of nanosheets with exposed {001} facets. d) Filtered atomic‐resolution image of {001} facets. e) HRTEM image of nanosheet in [111] orientation. f) FFT pattern from (e). g) HAADF‐STEM image of nanosheets with exposed {111} facets. h) Filtered atomic‐resolution image of {111} facets.
Orientation determination of anatase TiO2 nanosheets by means of NBD. a,c) HAADF‐STEM images of anatase TiO2. b,d) Diffraction patterns at different positions.
In situ photoreduction deposition of platinum on the different anatase TiO2 nanosheets. a–c) HRTEM images of Pt/TiO2‐001 prepared under different irradiation times. d–f) HRTEM images of Pt/TiO2‐111 prepared under different irradiation times.
Structural characterization and photocatalytic performance evaluation of different anatase TiO2 nanosheets. a) XRD patterns of the as‐prepared Anatase‐{001} and Anatase‐{111} samples. b) TEM images of anatase TiO2 with dominant {001} facets. c) TEM images of anatase TiO2 with dominant {111} facets. d) UV–visible diffuse reflectance spectra of the Anatase‐{001} and Anatase‐{111}, insert is the plots of (ahv)1/2 versus photon energy (hv). e–g) Mott–Schottky plots and scheme of electronic band structures for Anatase‐{001} and Anatase‐{111}. h) Time course of photocatalytic hydrogen evolution on the Anatase‐{001} and Anatase‐{111} using MeOH as the hole acceptor. i) Photocatalytic oxygen evolution on the Anatase‐{001} and Anatase‐{111} using Fe(NO3)3 as electron acceptors. Reaction conditions: cat. 100 mg, methanol solution (10 vol%) for hydrogen production, 10 mm Fe(NO3)3 aqueous solution for oxygen evolution, 300 W Xe lamp, Pt was deposited as a cocatalyst for hydrogen production.
Titanium dioxide (TiO2) nanocrystals have attracted great attention in heterogeneous photocatalysis and photoelectricity fields for decades. However, contradictious conclusions on the reporting of crystallographic orientation and exposed facets of TiO2 nanocrystals frequently appeared in the literature. Herein, using anatase TiO2 nanocrystals with highly exposed {001} facets as a model, we clarified the misleading conclusions that exist on anatase nanocrystals. Although the TiO2‐001 nanocrystals are recognized to be dominated by {001} facets, in fact, anatase nanocrystals with both dominating {001} and {111} facets are always co‐existing due to the similarities in lattice fringes and intersection angles between two types of facets (0.38 nm and 90° in [001] direction, 0.35 nm and 82° in [111] direction). We also give a paradigm for determining the crystallographic orientation and exposed facets based on transmission electron microscope (TEM) analysis, which provides a universal methodology to nanomaterials for determining the orientation and exposed facets. This article is protected by copyright. All rights reserved
Wireless stand‐alone interactive display (WiSID). a) Conceptual illustration and schematic of WiSID composed of three layers with an AC power unit. b) A cross‐sectional SEM image of ZnS:Cu/PVDF‐TrFE (above the dotted line) on a PEDOT:PSS electrode (under the dotted line). c) A photograph of a mechanically flexible WiSID on an AC power unit. d) The schematic of the cross section and equivalent circuit of WiSID with load resistor in place of the stimuli‐responsive layer. e) Induced voltages and f) power supplied to a display layer, load, and WiSID as a function of load impedance. g,h) FEA results of AC field analysis of a WiSID without stimuli (g) and with stimuli (h) on the stimuli‐responsive layer. i) Induced voltage between two parallel PEDOT:PSS electrodes of a WiSID without the stimuli‐responsive layer as a function of the distance between the AC power unit and the electrodes of the WiSID with an AC input voltage and frequency of 525 V and 10 kHz, respectively, in air. j) Induced voltage as a function of area of a transmitting electrode of AC power and a receiving electrode of the WiSID with an AC input voltage and frequency of 525 V and 10 kHz, respectively, in air. k) Induced voltage as a function of the relative permittivity of the medium (power transfer area) with an AC input voltage, frequency, and operating distance of 525 V, 10 kHz, and 1 mm, respectively.
Pressure‐ and temperature‐interactive WiSIDs. a) The schematic of a stimuli‐interactive WiSID with an AC power unit. b) SEM image of a topographically pyramid‐shaped micropatterned ionic gel layer as a pressure‐responsive layer is shown in the top image. The bottom schemes show schematic illustration of an ion transport mechanism in the PEO/LiTFSI/PEGDME composite as a temperature‐responsive layer at low (left) and high (right) temperature. c) Photographs of characteristic EL from a pressure‐interactive WiSID under the pressure of 20 kPa (top) and from a temperature‐interactive WiSID operated at 80 °C on the heating stage (bottom). d) Plots of impedance change and sensitivity of the pressure‐interactive WiSID as a function of pressure. Sensitivity is defined as Sp = δ(∆Z/Z0)/δp, where p is applied pressure and Z and Z0 are the impedance with and without applied pressure, respectively. e) EL spectra of the pressure‐interactive WiSID as a function of pressure in the range of 0–20 kPa. f) Luminance of the device as a function of pressure. The series of photographs of the device captured at different pressures are shown in the inset. g) Plots of impedance change of the temperature‐interactive WiSID as a function of temperature upon cooling and heating. h) EL spectra of the temperature‐interactive WiSID versus temperature in the range of 30–80 °C. i) Luminance of the temperature‐interactive WiSID as a function of temperature in the range of 30–80 °C. The series of photographs of the device are captured at different temperatures shown in the inset. j) Heating/cooling cycle endurance of the temperature‐interactive WiSID in both impedance and EL upon repetitive temperature change from 35 to 45 °C. k) SPL spectra of the temperature‐interactive WiSID versus frequency when the temperature‐responsive layer was heated from 30 to 80 °C. l) SPL of the temperature‐interactive WiSID as a function of temperature in the range of 30–80 °C upon cooling and heating.
Liquid‐interactive WiSID for dynamic monitoring of medical fluids. a) The schematic of a liquid‐interactive WiSID with an AC power unit. b) The photographs of a liquid‐interactive WiSID operated without deionized water (top) and with deionized water (bottom). c) A plot of impedance changes of the device with five different medical fluids. d) Luminance and SPL of the device with five different medical fluids. e) The schematic of the tube‐type liquid‐interactive WiSID with a mechanically flexible AC power unit wrapping the WiSID. A cross‐sectional scheme of the tube‐type WiSID with an AC power unit is shown on the right. f) A photograph of a test patient for intravenous therapy with a tube‐type WiSID connected to an infusion extension medical tube. The diameter of the tube was ≈3 mm. g) Photographs of the tube‐type WiSID (left) and the flexible AC power unit (right). h,i) When the medical fluid was filled through the tube‐type WiSID, both AC field‐induced EL (bottom photograph) (h) and inverse piezoelectric sound (i) were developed with source AC frequency and voltage of 10 kHz and 525 V. j) Photographs of a tube‐type WiSID connected to a three‐way medical connector for IV therapy. NS and TPN medical fluids were injected with the four‐step programmed prescription. k) The impedance variation of the device during the four‐step. l) Field‐induced EL and m) SPL variation of the device during the four‐step IV therapy.
Haptic WiSTID‐on‐finger for wireless human‐interactive display. a) The conceptual illustration of a H‐WiSTID‐on‐finger. The device architecture and equivalent circuit of a H‐WiSTID‐on‐finger are also shown. AC frequency‐dependent EL, inverse piezoelectric sound, and tactile vibration acquired from a H‐WiSTID‐on‐finger is illustrated on the right‐hand side. b) The photographs of the H‐WiSTID‐on‐finger before (top) and after (bottom) touch. The characteristic AC EL is apparent from the device when touched. c) Photographs of the H‐WiSTID‐on‐finger as a function of distance between the device and power unit. d) EL intensity and e) SPL of the device versus distance between the AC power unit and H‐WiSTID‐on‐finger. f) A plot of vibration amplitude of the H‐WiSTID‐on‐finger as a function of AC frequency. The threshold vibration amplitude of Pacinian corpuscles is also shown as a function of AC frequency. g) A plot of tactile sensing rate versus AC frequency. The tactile vibration developed in the H‐WiSTID‐on‐finger was quantified as a function of frequency with a sensing rate from 0 to 10. The results were obtained with an AC voltage of 220 V at a distance of 0.2 mm between the AC power unit and the device. h) The variation of the tactile sensing rate as a function of distance between the AC power unit and H‐WiSTID‐on‐finger. The applied AC voltage and frequency were 220 V and 100 Hz, respectively.
Haptic WiSTID‐on‐finger arrays for three‐mode electronic braille system. a) The conceptual illustration of the H‐WiSTID‐on‐finger with the three‐way non‐contact finger interactive outputs of AC EL, tactile vibration, and inverse piezoelectric sound using the designed AC power source. b) The photograph of the H‐WiSTID‐on‐finger emitting AC EL. The characteristic AC EL is apparent from the device when touched with the low and high frequencies of 200 Hz and 10 kHz with an AC voltage of 220 V. c) Tactile sensing rate of the H‐WiSTID‐on‐finger. Tactile vibration is generated from the device when the H‐WiSTID‐on‐finger is touched on the power unit with an AC voltage of 220 V. d) SPL of the device versus frequency with the low and high frequencies of 200 Hz and 10 kHz with AC voltage of 220 V. SPL peak at the frequency of 10 kHz is much greater than that at 200 Hz. e) The schematic (left) and photograph (right) of the 3 × 2 arrays of H‐WiSTID‐on‐fingers. f) The schematic of the 3 × 2 arrays of AC power units. g) The photograph of the operation of the arrays of H‐WiSTID‐on‐fingers when brought onto the arrays of the AC power units. h) The braille cell of the alphabet “N.” i) The photograph, j) the tactile vibration rate, and k) SPL of the arrays of H‐WiSTID‐on‐fingers displaying the braille alphabet “N.” l) Photographs of the arrays of H‐WiSTID‐on‐fingers showing the braille cells of a word “HELLO.”
With recent advances in interactive displays, the development of a stand‐alone interactive display with no electrical interconnection is of great interest. Here, we present a wireless stand‐alone interactive display (WiSID), enabled by direct capacitive coupling, consisting of three layers: two in‐plane metal electrodes separated by a gap, a composite layer for field‐induced electroluminescence (EL) and inverse piezoelectric sound, and a stimuli‐responsive layer, from bottom to top. Alternating current (AC) power necessary for field‐induced EL and inverse piezoelectric sound is wirelessly transferred from a power unit, with two in‐plane electrodes remotely separated from the WiSID. The unique in‐plane power transfer through the stimuli‐sensitive polar bridge allows stand‐alone operation of the WiSID, making it suitable for the wireless dynamic monitoring of medical fluids. Moreover, we demonstrate a haptic wireless stand‐alone trimodal interactive display mounted on a human finger, whereby touch is wirelessly displayed in various outputs of EL, inverse piezoelectric sound, and tactile vibration, making it suitable for a wireless three‐mode smart braille display. This article is protected by copyright. All rights reserved
a) Simulated AA stacking mode in different views and the theoretical pore size distributions for USTB‐6. b) Experimental, Pawley refined, and simulated PXRD patterns as well as the difference plot for USTB‐6. c) FTIR curves of PTO‐NH2, HATN‐CHO, and USTB‐6. d) Solid‐state ¹³C NMR spectrum of USTB‐6. e) N2 adsorption (solid) and desorption (hollow) curves at 77 K. f) Vertical view of the deformation charge density images of USTB‐6@G (the turquoise regions on the graphene and the yellow regions on the USTB‐6 indicate charge transfer from the graphene to the COF via conjugation of π–π interactions).
a–f) SEM, TEM, HR‐TEM, for USTB‐6 (a–c) and USTB‐6@G (d–f). g–j) AFM topography images and AFM height profiles of graphene (g,h) and USTB‐6@G (i,j).
a) CV curves of first five cycles of USTB‐6@G in the range of 1.2–3.9 V (scan rate: 0.2 mV s⁻¹). b) Nyquist plots of USTB‐6, USTB‐6/G, and USTB‐6@G. c,d) Cycling performances at 2 C and rate performances for USTB‐6, USTB‐6/G, and USTB‐6@G. e) Charge/discharge profiles of USTB‐6@G at different current density. f) Long‐term cycling stability of USTB‐6@G at 5 C. g) Comparison of the performance to other COF cathodes in LIBs. For figures (c,d,f), the charge and discharge capacities are represented by the colored and gray symbols, respectively.
a) CV profiles of USTB‐6@G collected at different scan rates. b) The b values of different peaks. c) Capacitive‐controlled contribution at the scan rate of 0.2 mV s⁻¹. d) Capacitive‐controlled contribution ratios at different scan rates. e) Charge and discharge curves of USTB‐6@G during the first cycle at a current density of 0.2 C. f,g) High‐resolution ex situ XPS spectra of N 1s (f) and O 1s (g) about the corresponding charge/discharge states of the USTB‐6@G electrode. h) The Li⁺ storage behavior of USTB‐6.
Scheme for the synthesis of USTB‐6 and USTB‐6@G.
Poor electronic and ionic conductivities of covalent organic frameworks (COFs) severely restrict the development of COF‐based electrodes for practical rechargeable batteries, therefore inspiring more research interests from the direction of both material synthesis and technology. Herein, a dual‐porous COF, USTB‐6, with good crystallinity and rich redox‐active sites has been conceived and reticulated by the polymerization of 2,3,8,9,14,15‐hexa(4‐formylphenyl)diquinoxalino [2,3‐a:2,3‐c]phenazine and 2,7‐diaminopyrene‐4,5,9,10‐tetraone. In particular, the graphene‐participated heterogeneous polymerization of the same starting materials affords the uniformly dispersed COF nanosheets with the thickness of 8.3 nm on the conductive carbon substrate, effectively enhancing the electronic conductivity of COF‐based electrode. Such graphene‐supported USTB‐6 nanosheets cathode used in lithium‐ion battery exhibits a specific capacity of 285 mA h g−1 at current density of 0.2 C and excellent rate performance with the still prominent capacity of 188 mA h g−1 at 10 C. More importantly, a capacity of 170 mA h g−1 is retained by USTB‐6 nanosheets cathode after 6000 cycles change and discharge measurement at 5 C. This article is protected by copyright. All rights reserved
A) Schematic of hydroxypropyl cellulose (HPC) thermal nanoimprint lithography on poly(dimethylsiloxane) (PDMS). B) Reflectance spectra allowing the estimation of the film thickness on silicon: 30 nm (solid purple), 125 nm (solid blue), 192 nm (solid orange), and 330 nm (solid red) for the HPC aqueous solutions at different concentrations. The dotted lines represent the corresponding simulated interference generated by differential refraction for thin films with heights of 30 nm (dotted purple), 125 nm (dotted blue), 192 nm (dotted orange), and 330 nm (dotted red). C) Photograph of the ≈30 nm patterned HPC stencil with periodicities of 600 nm (Λ) resulting from the 21 mg mL⁻¹ aqueous HPC solution (patterned area is 0.49 cm², scale bar: 0.7 cm). D–F) SEM images of different areas of the thinnest (30 nm) patterned HPC layer, with (F) showing a region at the edge of the film where holes are clearly visible. G) The PDMS substrate after removal of the HPC with water (indentation can be clearly observed), and H) the cross‐section of a PDMS substrate with the imprinted HPC. I) Schematic showing imprinting of an HPC film with volume lower than the negative volume of the stamp on a hard (top) and soft (bottom) substrate. Details of sample preparation can be found in Scheme S3, Supporting Information.
A) Schematic of the etching of the residual layer with UV–ozone, followed by inking with poly(methylhydrosiloxane), removal of the hydroxypropyl cellulose (HPC) with water, and introduction of the substrate into a growth solution to form gold nanoparticles (AuNPs). B) Schematic depicting competition between secondary solution nucleation and surface growth. C) Maximum extinction of growth solutions containing different water sources (Milli‐Q or high‐performance liquid chromatography, HPLC, water) and surfactant/capping ligand conditions (cetyltrimethylammonium bromide or chloride, CTAB/C, respectively), and all applying ascorbic acid as a mild reductant, over 15 min. D) Maximum extinction of growth solutions for substrates with different inking conditions. E) SEM image obtained under environmental conditions (60 Pa, low vacuum in water atmosphere) showing the patterned array after 10 min growth. F–I) SEM images showing a mixture of products suggesting different crystal structures and twinning: F) isotropic products, G) triangles, H) rods, and I) platelets/truncated triangles.
A–D) SEM images of patterned nanoparticle substrates after 1 min (A), 2 min (B), 5 min (C), and 10 min (D) growth with a growth solution containing cetyltrimethylammonium bromide (capping ligand/surfactant), gold salt, and ascorbic acid as a mild reducing agent. E–H) Size distributions for the 1 min (E), 2 min (F), 5 min (G), and 10 min (H) samples (150 nanoparticles each, n = 10). I,J) Representative SEM images showing pattern yield and common defects: multiple particles per area (pink circles), vacancies (purple circles), and particles outside the pattern (green areas). Patterned regions with single particles are indicated with a white circle (n = 10).
A) Schematic of overgrowth process into nanostars on poly(dimethylsiloxane) (PDMS). B) Normalized and smoothed (Savitzky–Golay) transmission spectra comparing the plasmon band position before (5 min growth in cetyltrimethylammonium bromide solution with ascorbic acid added as a weak reducing agent) and after stars overgrowth (5 min in a solution containing capping ligand/surfactant laurylsulfobetaine, HCl and ascorbic acid). C–E) SEM images of the nanostar arrays at different magnifications.
A–C) SEM images showing patterned hydroxypropyl cellulose (HPC) films with lattice parameter Λ = 400 nm (A) and the resulting patterned nanoparticle substrates (B,C). D–F) SEM images showing the Λ = 500 nm stencil (D) and the corresponding patterned nanoparticles (E,F) (SEM data for the Λ = 600 nm HPC stencils and resulting gold nanoparticle arrays are shown in Figures 1 and 3). G) Smoothed (Savitzky–Golay) and normalized transmission spectra of the substrates following refractive index matching. The dashed lines represent the theoretical position of Rayleigh–Wood anomalies (nPDMS = 1.4). H–J) Contour plots showing measured angular dependence of the optical response optima for the Λ = 400 nm (H), Λ = 500 nm (I), and Λ = 600 nm (J) lattices. The dashed lines represent the predicted positions and displacement of the Rayleigh anomalies, depending on the azimuthal angle, using nPDMS = 1.4. Additional SEM images can be found in Figures S13–S15, Supporting Information. All arrays were fabricated after 5 min of growth in the cetyltrimethylammonium bromide/ascorbic acid solution optimized from Figure 2 to limit secondary nucleation.
Precise arrangements of plasmonic nanoparticles on substrates are important for designing optoelectronics, sensors, and metamaterials with rational electronic, optical, and magnetic properties. Bottom‐up synthesis offers unmatched control over morphology and optical response of individual plasmonic building blocks. Usually, the incorporation of nanoparticles made by bottom‐up wet chemistry starts from batch synthesis of colloids, which requires time‐consuming and hard‐to‐scale steps like ligand exchange and self‐assembly. Herein, we develop an unconventional bottom‐up wet‐chemical synthetic approach for producing gold nanoparticle ordered arrays. Water‐processable hydroxypropyl cellulose stencils facilitate the patterning of a reductant chemical ink on which nanoparticle growth selectively occurs. Arrays exhibiting lattice plasmon resonances in the visible region and near infrared (quality factors >20) were produced following a rapid synthetic step (<10 min), all without cleanroom fabrication, specialized equipment, or self‐assembly, constituting a major step forward establishing in situ growth approaches. We further demonstrate the technical capabilities of this method through modulation of the particle size, shape, and array spacings directly on the substrate. Ultimately, establishing a fundamental understanding of in situ growth has the potential to inform the fabrication of plasmonic materials, opening the door for in situ growth fabrication of waveguides, lasing platforms, and plasmonic sensors. This article is protected by copyright. All rights reserved
In cancer radiotherapy, the lack of fixed DNA damages by oxygen in hypoxic microenvironment of solid tumors often lead to severe radioresistance. Nitric oxide (NO) is a potent radiosensitizer that acts in two ways. It can directly react with the radical DNA thus fixes the damage. It also normalizes the abnormal tumor vessels, thereby increasing blood perfusion and oxygen supply. To achieve these functions, the dosage and duration of NO treatment need to be carefully controlled, otherwise it would lead to the exact opposite outcomes. However, a delivery method that fulfills both requirements is still lacking. We designed a NO depot for the control of NO releasing both over quantity and duration for hypoxic tumor vessel normalization and radiosensitization. In B16 tumor‐bearing mice, the depot could provide low dosage NO continuously and release large amount of NO immediately before irradiation for a short period of time. These two modes of treatment worked in synergy to reverse the radioresistance of B16 tumor more efficiently than releasing at single dosage. This article is protected by copyright. All rights reserved
a) The scheme of the Cd‐doped and anion‐exchange QDs using the hot injection method. b) The FTIR spectra of the pristine, doped, and exchanged QDs. c,d) XPS spectra of Cd 3d (c) and Pb 4f (d) for pristine, doped, and exchanged QDs. e) The crystal model and f) calculated crystal model of the bromine ions filling into halogen vacancies based on the Cd‐CsPb(BrxCl1−x)3 QDs using DFT calculations. g) The scheme of vacancy defects in the pristine CsPb(BrxCl1−x)3 QDs. h) Mechanism description of the anion‐exchange: ① Bromine ions replace chloride ions; ② The passivation of vacancy defects.
a–c) TEM images of pristine QDs (a), doped QDs (b), and exchanged QDs (c). d) XRD patterns, e) absorption and photoluminescent spectra, and f) time‐resolved spectroscopy of the pristine, doped, and exchanged QDs.
Temperature‐dependent PL measurement. a,d) PL spectra from 80 to 290 K for the films of pristine (a) and exchanged (d) QDs. b,e) PL decay curves from 80 to 290 K of the films of pristine QDs (b) and exchanged QDs (e). c,f) The integrated PL intensity (Arrhenius fit) versus of the films of pristine QDs (c) and exchanged (f) QDs. g,h) PL spectra from 300 to 360 K for pristine QDs (g) and exchanged QDs (h).
a,b) AFM images of pristine (a) and exchanged (b) QD films. c,d) Current density–voltage curves of the hole‐only device (c) and electron‐only device (d) under dark condition. The inset shows the scheme of the device structure.
a) The architecture of device. b) Electroluminescence spectra driven by voltages from 3 to 7 V. Inset image show the luminescence of device. c) The curve of the voltage–current density luminance of the device. d) The curve of the current efficiency of the device. e) EQE of the devices based on the exchange QDs. f) CIE coordinate of the device.
Halogen vacancies are of great concern in blue‐emitting perovskite quantum dot light‐emitting diodes because they affect their efficiency and spectral shift. Here, the enriched‐bromine surface state is realised using a facile strategy that employs a PbBr2 stock solution for anion exchange based on Cd‐doped perovskite quantum dots. We find that the doped Cd ions are expected to reduce the formation energy of halogen vacancies filled by the external bromine ions, and the excess free bromine ions in solution are enriched in the surface by anchoring with halogen vacancies as sites, accompanied with the shedding of surface long‐chain ligands during the anion exchange process, resulting in a Br‐rich and “neat” surface. Moreover, the surface state exhibits good passivation of the surface defects of the controlled perovskite QDs and simultaneously increases the exciton binding energy, leading to excellent optical properties and stability. Finally, the sky‐blue emitting perovskite QLEDs (490 nm) are conducted with a record external quantum efficiency of 14.6% and current efficiency of 19.9 cd/A. Meanwhile, the electroluminescence spectra exhibit great stability with negligible shifts under a constant operating voltage from 3 to 7 V. This strategy paves the way for improving the efficiency and stability of perovskite QLEDs. This article is protected by copyright. All rights reserved
Synthesis of PQDs using different approaches. a,b) Schematic illustrations of binary‐precursor (a) and ternary‐precursor (b) approaches for the synthesis of PQDs. c–f) TEM images of control PQDs (c), PEAI‐based PQDs (d), TMSI‐based PQDs (e), and TMSI·TOP‐based PQDs (f). g) Light absorption and PL spectra of the PQDs synthesized using different approaches.
Characterization of PQDs synthesized using different approaches. a) XRD patterns and b) FTIR spectra of control, TMSI‐ and TMSI·TOP‐based PQDs. c) Pb 4f, d) I 3d and e) Si 2p XPS spectra of control, TMSI‐ and TMSI·TOP‐based PQDs. f,g) The calculated surface structure of the TMSI‐based PQD before (f) and after (g) the CST. The corresponding ELF images are also included in the figures.
Charge‐carrier dynamics and orientation of PQD solids. a–c) 2D pseudocolor fs‐TA plots of control (a), TMSI‐based (b) and TMSI·TOP‐based (c) PQD solid films. d–f) GIWAXS patterns of control (d), TMSI‐based (e), and TMSI·TOP‐based (f) PQD solid films. g) Integrated 1D profiles of the GIWAXS patterns within the PQD solid films. h) Schematical illustration of the orientation of PQD solids prepared using the PQDs synthesized using different approaches.
Photovoltaic performance of PQDSCs. a) Cross‐sectional SEM image of the PQDSC. b) J–V curves of control, TMSI‐ and TMSI·TOP‐based PQDSCs. The inset shows the photovoltaic parameters of these devices. c) IPCE spectrum and integrated photocurrent density of the TMSI·TOP‐based PQDSC. d) Stabilized current density and power output of the TMSI·TOP‐based PQDSC. e) PCE statistics of control, TMSI‐ and TMSI·TOP‐based PQDSCs. f) Stability of PQDSCs storage under ambient conditions.
Charge‐carrier recombination and collection of PQDSCs. a) TPV and b) TPC curves of control and TMSI·TOP‐based PQDSCs. c,d) Light‐intensity dependent photovoltage (c) and photocurrent density (d) of the control and TMSI·TOP‐based PQDSCs. e,f) SCLC curves of the control (e) and TMSI·TOP‐based (f) PQDSCs. The inset in (e) shows the device structure for the SCLC measurement.
Perovskite quantum dots (PQDs) emerge as competitive optoelectronic materials for photovoltaic applications due to their ideal bandgap energy, high defect tolerance and solution processability. However, the highly dynamic surface and unperfect cubic structure of PQDs generally result in unfavorable charge carrier transport within the PQD solids and serious nonradiative recombination. Herein, the highly orientated PQD solid is demonstrated using precursor engineering accompanied by a chemical stripping treatment (CST). A combination of systematically experimental studies and theoretical calculations is conducted to fundamentally understand the resurfacing of PQDs using the CST approach. The results reveal that highly ordered PQDs could result in a high orientation of PQD solids, significantly promoting the charge carrier transport within the PQD solids. Meanwhile, the ideal cubic‐structured PQD with an iodine‐rich surface dramatically decreases surface trap states, thereby substantially diminishing trap‐assisted nonradiative recombination. Consequently, the inorganic PQD solar cell (PQDSC) delivers a power conversion efficiency of up to 16.25%. This work provides a feasible avenue to construct highly orientated PQD solids with improved photophysical properties for high‐performance optoelectronic devices. This article is protected by copyright. All rights reserved
Schematic illustration of the synthesis of Hv‐Ni3Mn0.7Fe0.3‐LDH and structural analysis. a) Illustration of the structural change of LDH after introducing H vacancies and its promoted Zn‐ion storage. b) XRD patterns of Ni3Mn0.7Fe0.3‐LDH and Hv‐Ni3Mn0.7Fe0.3‐LDH. c) SEM image of Hv‐Ni3Mn0.7Fe0.3‐LDH. d) HR‐TEM image of Hv‐Ni3Mn0.7Fe0.3‐LDH and e) corresponding SAED pattern. f–i) High‐resolution XPS spectra of Ni 2p (f), Mn 2p (g), Fe 2p (h), and O 1s (i) of Ni3Mn0.7Fe0.3‐LDH and Hv‐Ni3Mn0.7Fe0.3‐LDH.
The electrochemical performance of Ni3Mn0.7Fe0.3‐LDH and Hv‐Ni3Mn0.7Fe0.3‐LDH electrodes. a) Cyclic voltammetry curves of Ni3Mn0.7Fe0.3‐LDH and Hv‐Ni3Mn0.7Fe0.3‐LDH at a scan rate of 0.2 mV S⁻¹. b) Galvanostatic charge/discharge curves of Hv‐Ni3Mn0.7Fe0.3‐LDH at a current density of 50 mA g⁻¹. c) Galvanostatic charge/discharge curves of Ni3Mn0.7Fe0.3‐LDH at a current density of 50 mA g⁻¹. d) Rate performance and e) cycling performance of Ni3Mn0.7Fe0.3‐LDH and Hv‐Ni3Mn0.7Fe0.3‐LDH electrode at 100 mA g⁻¹. f) The Ragone plots of average output voltage versus specific capacity for Hv‐Ni3Mn0.7Fe0.3‐LDH and some state‐of‐the‐art cathode materials for aqueous ZIBs. g) Long term cycling stability of Ni3Mn0.7Fe0.3‐LDH and Hv‐Ni3Mn0.7Fe0.3‐LDH electrode at 1 A g⁻¹.
Mechanism analysis by ex situ XRD, XPS, and SEM observation. a) Typical galvanostatic charge/discharge curves. b) Ex situ XRD patterns collected during the whole charge/discharge process of Hv‐Ni3Mn0.7Fe0.3‐LDH electrode. c) Partial enlargement for the XRD diffraction of (003) plane. d–f) SEM images of Hv‐Ni3Mn0.7Fe0.3‐LDH electrode at different discharge/charge states: d) the first charge to 1.8 V, e) the first discharge to 0.8 V, and f) the second charge to 1.8 V. g–i) High‐resolution XPS spectra for Hv‐Ni3Mn0.7Fe0.3‐LDH: g) Mn 2p, h) Fe 2p, and i) Ni 2p, at various charged/discharged states. j) Normalized oxygen K‐edge soft X‐ray absorption spectra of Hv‐Ni3Mn0.7Fe0.3‐LDH in a whole battery cycling process.
First principle calculation and kinetic analysis. a) Schematic illustration of Zn²⁺ ion migration in Hv‐Ni3Mn0.7Fe0.3‐LDH. b) Diffusion barrier profiles for Zn²⁺ ion migration along the path TC→TE→TC in Ni3Mn0.7Fe0.3‐LDH and Hv‐Ni3Mn0.7Fe0.3‐LDH. c) Cyclic voltammetry curves of Hv‐Ni3Mn0.7Fe0.3‐LDH electrode at different scan rates. d) The percentage of surface capacitive contribution in Hv‐Ni3Mn0.7Fe0.3‐LDH electrode. e) GITT curves and the corresponding Zn²⁺ ions diffusion coefficients at charged state and discharged state. f,g) The variation of MnO bond length at discharged state for Ni3Mn‐LDH, Ni3Mn0.7Fe0.3‐LDH, and Hv‐Ni3Mn0.7Fe0.3‐LDH. h–k) Schematic illustration of MnO6 octahedra and structural Jahn–Teller distortion in different LDH.
Advanced cathode materials play an important role in promoting aqueous battery technology for safe energy storage. Transition metal double hydroxides are usually elusive as a stable cathode for aqueous zinc ion batteries (AZIBs) due to their unstable crystal structure, sluggish ion transportation and insufficient active sites for zinc ion storage. Here we report a trinary layered double hydroxide with hydrogen vacancies (Ni3Mn0.7Fe0.3‐LDH) as a new cathode material for AZIBs. A reversible high capacity up to 328 mAh g−1 can be obtained and cycle stably over 500 cycles with a capacity retention of 85%. Experimental and theoretical studies reveal that the hydrogen vacancies in LDH could expose lattice oxygen atoms as active sites for zinc ion storage and accelerate ion diffusion by reducing the electrostatic interactions between zinc ions and the host structure. Besides, the synergy of trinary transitional metal cations can suppress the Jahn‐Teller distortion of manganese (III) oxide octahedron and enable long cycle stability. This work provides not only a series of high performance cathode materials for aqueous zinc ion batteries but also a novel materials design strategy that can be extended to other multi‐valence metal ion batteries. This article is protected by copyright. All rights reserved
Materials and strategies used to develop the adaptive hydrogels. A) Molecular structure of the azide‐functionalized polyisocyanide (PIC), the short spacer (Dd) with a dual purpose (in the ionic cross‐linking approach and in the particle cross‐linking approach as functional groups grafted on the FeNR surface). B) Chain bundling network formation attributed to hydrophobic interactions of the polyisocyanide backbone and the ethylene glycol tails. Ionic cross‐linking via the SPAAC reaction and metal‐coordination interactions in the mixture of PIC, Dd, and iron ions results in the inside bundling cross‐linking process without changing the morphology of the gel network. Nanoparticle cross‐linking in the mixture of PIC and Dd‐functionalized FeNRs results in between bundling cross‐linking with the formation of a polymer–particle network.
Mechanical properties of the cross‐linked composite gels as a function of the cross‐link density. A) Shear modulus G′ in the LVE as a function of temperature T measured for samples with different cross‐linking efficiency, obtained in the reheating curves. B) Strain stiffening data: differential modulus K′ as a function of applied stress σ for the same samples, showing little variation for the Fe³⁺‐cross‐linked gels and a decrease from K′ ∝ σ1.5 to K′ ∝ σ0.6 at high FeNP cross‐linking densities. C) K′ versus σ with the modulus and stress normalized with G0 and σc, respectively, clearly showing the reduced stiffening response for the FeNP‐cross‐linked gels. D–G) Data extraction from the nonlinear measurements: critical stress σc (D), critical strain γc = σc/G′ (E), stiffening index m (F), and stiffening factor K′/G0 (G) plotted against the stiffening ratio R. Data are averages of n = 2–4 experiments; shading in (A–C) or error bars in (D–G) represent standard deviation. The solid lines in (D–G) are power‐law fits. For all samples: [PIC] = 1 mg mL⁻¹, [Fe³⁺] = 667 µm or [FeNRs] = 1 mg mL⁻¹, only the molar ratio of linkers varies.
Mechanical properties of gels at different pH using ionic cross‐linking and nanoparticle cross‐linking approaches. A,B) Schematic overview of the effects of pH in Fe³⁺‐ (A) and FeNP‐ (B) cross‐linked gels. C) Shear modulus G′ as a function of temperature T of PICFe3+‐2 gels (open squares), PICFeNP‐4 gels (filled circles), PIC (black) at different pH. D,E) Differential modulus K′ as a function of applied stress σ of the same gels (D), rescaled with G0 and σc to emphasize the different responses to shear stress. F–H) Stiffening parameters: stiffening index m (F), critical strain γc (G), and stiffening factor K′/G0 at σ/σc = 18 (H) as determined from the rheology curves as a function of the pH for PICFe3+‐2 gels (open), PICFeNP‐4 (filled), and PIC (black). Overall, the parameters show small variations for the Fe³⁺‐cross‐linked gels and for the FeNP‐cross‐linked gel, a strong decrease in m, γc, and K′/G0 with increasing pH, showing that pH change is a very powerful strategy to also affect the nonlinear mechanical properties of these gels. Data are averages of n = 2–4 experiments; shading in (A–C) or error bars in (D–G) represent standard deviation. Significance was determined through a paired t‐test. Significant differences: *, p ≤ 0.05; **, p ≤ 0.01; ***, p ≤ 0.001. I) Schematic overview of the rheology setup for in situ dynamic experiments. The hydrogel (pink) is formed between the plates in a larger cup. The cup allows for replacement of the aqueous phase with buffer solutions of different pH (green and blue) during the rheology experiment. J) G′ of PICFeNP‐4 as a function of time, showing a dynamic pH responsiveness. Experiment: gel was prepared and cross‐linked at pH = 7, cooled, returned to 37 °C, and equilibrated for 10 min; pH 9 buffer was placed in the cup and G′ is monitored while the pH of the gel slowly increases through diffusion. After 75 min, the pH 9 buffer was replaced by a pH 7 buffer, 55 min later by a pH 5 buffer, and later again by a pH 7 buffer, showing good reversibility of the mechanical properties. Conditions for all panels: For PICFe3+‐2 gels: [PIC] = 1 mg mL⁻¹, [linker] = 350 µm, and [Fe³⁺] = 667 µm; for PICFeNP‐4 gels, [PIC] = 1 mg mL⁻¹, [linker] = 372 µm, and [FeNR] = 1 mg mL⁻¹. T = 37 °C (D–H, J).
Self‐healing properties of differently cross‐linked PIC gels. A–C) Photos of the self‐healing properties of PIC (A), ionic‐cross‐linked gels (B), and nanoparticle‐cross‐linked gels (C). The red line in (C) indicates the fractured interface. Movies of the experiment are available in the Supporting Information. D–F) Strain amplitude measurements at with strains cycling between 1% (blue background) and 300% (yellow) strain of PIC gels (no cross‐linking) (D), ionic‐cross‐linked gels (E), and nanoparticle‐cross‐linked gels (F). Conditions: PIC control: [PIC] = 6 mg mL⁻¹; PICFe3+‐2 gel: [PIC] = 6 mg mL⁻¹, [linker] = 2100 µm, and [Fe³⁺] = 1000 µm; for PICFeNP‐4 gels, [PIC] = 6 mg mL⁻¹, [linker] = 1448 µm, and [FeNR] = 4 mg mL⁻¹; all samples: T = 37 °C. Note, the linker/PIC ratio is the same as in the PICFe3+‐2 gel, ferric ions are in excess, and that FeNP concentration matches the PICFeNP‐4 gel. G) Schematic illustration of the healing concept for the dynamic metal‐coordination interactions in the different hydrogels.
Biocompatibility of the hydrogels. A) Confocal fluorescence images of fibroblasts (HFFs) encapsulated in PIC gels and composites at day 5 after seeding. Living cells (green) and dead cells (red) stained with calcein‐AM and TOTO‐3, respectively. B) Bright‐field images of fibroblasts encapsulated in the different hydrogel groups at day 5 after seeding, showing good strong cell spreading in 3D. Scale bars in (A/B): 100 µm. C) Quantitative analysis of confocal images of living cells (green) and dead cells (red). Data are presented as percentage of live cells, n = 29; significance was determined through a one‐way ANOVA. D) CCK8 assay results for cell viability at day 1 (striped columns) and day 5 (solid columns); Statistics n = 3; Student t‐test; p‐values > 0.05, are considered not significant (ns); significant differences: *, p ≤ 0.05. All data represented as mean ± standard deviation. The color coding in panels (C) and (D) follows that used in Figure 2.
The material's properties of biological tissues are unique. Nature is able to spatially and temporally manipulate (mechanical) properties while maintaining responsiveness towards a variety of cues; all without majorly changing the material's composition. Artificial mimics, synthetic or biomaterial‐based are far less advanced and poorly reproduce the natural cell microenvironment. A viable strategy to generate materials with advanced properties combines different materials into nanocomposites. This work describes nanocomposites of a synthetic fibrous hydrogel, based on polyisocyanides (PIC), that is noncovalently linked to a responsive crosslinker. The introduction of the crosslinker transforms the PIC gel from a static fibrous extracellular matrix mimic to a highly dynamic material that maintains biocompatibility, as demonstrated by in situ modification of the (non)linear mechanical properties and efficient self‐healing properties. Key in the materials design is crosslinking at the fibrillar level using nanoparticles, which, simultaneously may be used to introduce more advanced properties. This article is protected by copyright. All rights reserved
a) Schematic illustration of the effect of CeCl3 additive on the Zn deposition process. b,c) In situ microscopy images of Zn plating process in the ZSO electrolyte and ZSO/Ce electrolyte. d,e) AFM images of the cycled Zn in the ZSO electrolyte and ZSO/Ce electrolyte.
a) The ATR‐FTIR of ZSO and ZSO/Ce electrolyte. b) Raman spectra of ZSO and ZSO/Ce electrolyte. c) Linear polarization curves, and d) EIS of Zn–Zn symmetrical cells in the ZSO electrolyte with different CeCl3 concentrations. e) The linear sweep voltammetry curves of the Zn–Ti half cells tested at a scan rate of 1 mV s⁻¹ in the ZSO electrolyte and ZSO/Ce electrolyte. f) In situ DEMS pattern displaying released H2 gas during cycling for Zn–Cu cells in ZSO and ZSO/Ce electrolyte.
a,b) The XRD pattern of pure Zn and 2 g L⁻¹ CeCl3 modified Zn anode. c) The crystal structure of metal Zn. d,e) Depth profiles and 3D visualization of surfaces on the Zn anode cycled with ZSO/Ce electrolyte with Cs⁺ sputtering; the inset in (d) shows the TOF‐SIMS spectra for Ce³⁺ and Zn²⁺ integrated over 500 s. f) The binding energy of H2O, Zn²⁺, and Ce³⁺ on Zn (002) surface. g) The initial state for the Zn anode with ZSO electrolyte or ZSO/Ce electrolyte. h,i) The potential and current distribution of Zn deposition at simulation time of 2 min in ZSO electrolyte (h) and ZSO/Ce electrolyte (i); the gray lines with the arrows and the black lines at the bottom represent the current and the initial surface of the Zn anode.
a,b) Galvanostatic voltage profiles of a Zn–Zn symmetric cell cycled in ZSO electrolyte and ZSO/Ce electrolyte for plating/stripping 1 mAh cm⁻² per cycle at current densities of 2 mA cm⁻² (a) and 10 mAh cm⁻² (b) per cycle at a current density of 40 mA cm⁻². c) The rate performance of a Zn symmetric cell using the ZSO/Ce electrolyte. d) The CPC comparison of the Zn–Zn cell between this work and other reports. The detailed references corresponding to the point number are listed in Table S1, Supporting Information. e) The CE of Cu–Zn cells using ZSO/Ce electrolyte and ZSO electrolyte. f) The voltage profiles of Cu–Zn cell cycled in (blue line) ZSO/Ce electrolyte and (red line) ZSO electrolyte at the selected cycles. g) Comparison of the cumulative areal capacity plated, CE, and current density in recent reports. The detailed references corresponding to the point numbers are listed in Table S2, Supporting Information.
a) CV curves at various scan rates. b) log(i)–log(v) plots at different peak currents of Zn–LFP cells with ZSO/Ce electrolyte. c) Rate performance of Zn–LFP cells using ZSO electrolyte and ZSO/Ce electrolyte. d) GCD curves of Zn–LFP based on ZSO/Ce electrolyte at various current densities. e) Long‐term cycling performance of Zn–LFP cell at current density of 2 C together with corresponding CEs. f) Digital photo of open‐circuit voltage of the pouch Zn–LFP cell and g) two pouch Zn–LFP cells in series. h) Digital photo showing the working states of pouch Zn–LFP cells in series to power a LED indicator.
Although aqueous Zn batteries have become a more sustainable alternative to lithium‐ion batteries owing to their intrinsic security, their practical applications are limited by dendrite formation and hydrogen reactions. We present the first application of a rare earth metal type addition to Zn batteries, cerium chloride (CeCl3), as an effective, low‐cost, and green electrolyte additive that facilitates the formation of dynamic electrostatic shielding layer around the Zn protuberance to induce the uniform Zn deposition. After introducing CeCl3 additives, the electrochemical characterizations, in‐situ optical microscopy observation, in‐situ differential electrochemical mass spectrometry, along with density functional theory calculations and finite element method simulations reveal resisted Zn dendritic growth and enhanced electrolyte stability. As a result, the Zn‐Zn cells using CeCl3 additive exhibit a long cycling stability of 2600 h at 2 mA cm–2, an impressive cumulative areal capacity of 3.6 Ah cm–2 at 40 mA cm–2 and a high coulombic efficiency of ∼99.7%. The fact that the Zn‐LiFePO4 cells with proposed electrolyte retain capacity significantly better than the additive‐free case is even more exciting. This article is protected by copyright. All rights reserved
Polymers are usually considered thermal insulators; however, significant enhancements in thermal conductivity (k) have been observed in oriented fibers and films. Despite being advantageous in real‐world applications, extending the linear thermal‐transport advantage of polymers into the three‐dimensional space in bulk materials is still limited due to the spatially interfacial phonon‐conduction barriers. Herein, inspired by the structure of tropocollagen, we discovered that weaving hierarchically arranged poly(p‐phenylene benzobisoxazole) (PBO) fibers with a spiral configuration into an epoxy matrix can yield a three‐dimensionally continuous thermal pathway. This achieves both a through‐plane k of 10.85 W/m K and an in‐plane k of 7.15 W/m K. Theoretical molecular simulations in combination with classical nonlinear modeling attribute the above spatially thermally conductive achievement to not only the hierarchical molecular, spiral and weaving structure of PBO, but also the noncrystalline chains that carry overlapping phonon density of states, thus thermally bridging adjacent high‐k crystals in the PBO fiber. Consequently, the interfacial thermal resistance among high‐k PBO crystals is suppressed to be on the order of 10–10 m2 K/W in both the through‐plane and in‐plane directions. Other advantages include being lightweight, mechanically strong, flexible, and non‐combustible. This material creates opportunities for organic polymers in high‐performance thermal management applications. This article is protected by copyright. All rights reserved
As the world steps into the era of Internet of Things (IoT), numerous miniaturized electronic devices requiring autonomous micropower sources will be connected to the internet. All‐solid‐state thin‐film lithium/lithium‐ion microbatteries (TFBs) combining solid‐state battery architecture and thin‐film manufacturing are regarded as ideal on‐chip power sources for IoT‐enabled microelectronic devices. However, unlike commercialized lithium‐ion batteries, TFBs are still in the immature state, and new advances in materials, manufacturing, and structure are required to improve their performance. In this review, we discuss the current status and existing challenges of TFBs for practical application in internet‐connected devices for the IoT. Recent progress in thin‐film deposition, electrode and electrolyte materials, interface modification, and 3D architecture design is comprehensively summarized and discussed, with emphasis on state‐of‐the‐art strategies to improve the areal capacity and cycling stability of TFBs. Moreover, to be suitable power sources for IoT devices, the design of next‐generation TFBs should consider multiple functionalities, including wide working temperature range, good flexibility, high transparency, and integration with energy‐harvesting systems. Perspectives on designing practically accessible TFBs are provided, which may guide the future development of reliable power sources for IoT devices . This article is protected by copyright. All rights reserved
In recent years, traditional antibiotic efficacy has rapidly diminished due to the advent of multidrug‐resistant (MDR) bacteria which pose severe threats to human life and globalized healthcare. Currently, the development cycle of new antibiotics cannot match the ongoing MDR infection crisis. Therefore, novel strategies are required to resensitize MDR bacteria to existing antibiotics. In this study, novel cationic polysaccharide conjugates Dextran‐graft‐Poly(5‐(1,2‐dithiolan‐3‐yl)‐N‐(2‐guanidinoethyl)pentanamide) (Dex‐g‐PSSn) were synthesized using disulfide exchange polymerization. Critically, bacterial membranes and efflux pumps were disrupted by a sub‐inhibitory concentration of Dex‐g‐PSS30, which enhanced rifampicin (RIF) accumulation inside bacteria and restored its efficacy. Combined Dex‐g‐PSS30 and RIF prevented bacterial resistance in bacteria cultured over 30 generations. Furthermore, Dex‐g‐PSS30 restored RIF effectiveness, reduced inflammatory reactions in a pneumonia‐induced mouse model, and exhibited excellent in vivo biological absorption and degradation capabilities. As an antibiotic adjuvant, Dex‐g‐PSS30 provides a novel resensitizing strategy for RIF against MDR bacteria and bacterial resistance. Our Dex‐g‐PSS30 research provides a solid platform for future MDR applications. This article is protected by copyright. All rights reserved
Rationales of material construction for unifying HSo and DeP in water electrolysis. a,b) The theoretical process insights of HSo in the HER (a) and b) DeP in the OER (b); herein the serial numbers indicate: 1, water dissociation; 2, proton concentration; 3, hydrogen spillover; 4, H2 release; 5, adsorbates’ evolution; 6, DeP; 7, proton transport; 8, O2 release; and 9, water generation. c) MoS2/NiPS3 heterostructure model. d) DOS of NiPS3 and MoS2/NiPS3. e) Planar average potential along the Z‐direction. f,g) Electronic potentials of the MoS2/NiPS3 base plane. h) Mechanism diagram of dual DeP‐enhanced OER for MoS2/NiPS3 system based on IPF.
Fabrication process and structural characterization of MoS2/NiPS3 heterostructure. a) The synthetic scheme. b,c) SEM images of bulk NiPS3 (b) and EE‐NiPS3 (c) (inset, EE‐NiPS3 in DMF exhibits the Tyndall effect). d) SEM image, e) TEM image, f) AC‐TEM image, g) SAED pattern, and h) EDX mapping of the MN‐1 heterostructure.
Phase and surface chemical characteristics of MoS2/NiPS3 heterostructures. a) XRD patterns, b) Raman spectra, and c) FT‐IR spectra of bulk NiPS3, EE‐NiPS3, and MN‐1. d,e) Ni 2p XPS spectra of EE‐NiPS3 (d) and MN‐1 (e), f,g) Mo 3d XPS spectra of MoS2 (f) and MN‐1 (g), and h,i) S 2p XPS spectra of MoS2 (h) and MN‐1 (i).
EXAFS spectra and analyses. a) Ni K‐edge XANES of EE‐NiPS3 and MN‐1, b) Mo K‐edge XANES of MoS2 and MN‐1, c) Ni K‐edge FT‐EXAFS of EE‐NiPS3 and MN‐1, d) Mo K‐edge FT‐EXAFS of MoS2 and MN‐1, e) k‐space spectra of EE‐NiPS3 and MN‐1, and f) k‐space spectra of MoS2 and MN‐1. g–l) WT for the k³‐weighted EXAFS signal of the EE‐NiPS3 Ni K‐edge (g), MN‐1 Ni K‐edge (h), and NiO Ni K‐edge (i), and the MoS2 Mo K‐edge (j), MN‐1 Mo K‐edge (k), and MoO3 Mo K‐edge (l).
Electrocatalytic performance and mechanism studies. a–f) For different electrocatalysts in 1 m KOH, a–c) HER performance: a) LSV curves, b) Tafel plots, and c) current–potential curves with different scan rates; d–f) OER performance: d) LSV curves, e) Tafel plots, and f) EIS data. g) Polarization curves of overall water splitting based on MN‐1||MN‐1 couple, and schematic diagram of overall water splitting (inset). h) Photograph of the device for overall water splitting driven by a solar cell (≈1.65 V). i) Humidity sensing response of EE‐NiPS3, MoS2, and MN‐1 when cycled between dry air and 50% RH. j) Energy barriers for HER on different sites. k) Energy barriers for four‐electron‐step OER process (at the equilibrium potential, U = 0.401 V vs SHE) and l) the reaction energy barrier for DeP on different sites for different electrocatalysts.
Hydrogen spillover has emerged to upgrade hydrogen evolution reaction (HER) activity of Pt‐support electrocatalysts, but it is not applicable to the deprotonated oxygen evolution reaction (OER). Non‐precious catalysts that can perform well in both hydrogen spillover and deprotonation are extremely desirable for sustainable hydrogen economy. Herein, an affordable MoS2/NiPS3 vertical heterostructure catalyst is presented to synergize hydrogen spillover and deprotonation for efficient water electrolysis. Internal polarization field (IPF) is clarified as the driving force nature of hydrogen spillover in HER electrocatalysis. The hydrogen spillover from MoS2 edge to NiPS3 can activate the NiPS3 basal plane to boost the HER activity of MoS2/NiPS3 heterostructure (112 mV versus RHE at 10 mA cm–2). While for OER, the IPF in the heterostructure can facilitate the hydroxyl diffusion and render the NiPS3‐to‐MoS2/P‐to‐S dual‐pathways for deprotonation. Resultantly, the stacking of OER‐inactive MoS2 on the NiPS3 surface still brings intriguing OER enhancements. Serving them as electrode couples, the overall water splitting is attested stably with a cell voltage of 1.64 V at 10 mA cm–2. This research puts forward IPF as the criterion in the rational design of hydrogen spillover/deprotonation‐unified non‐precious catalysts for efficient water electrolysis. This article is protected by copyright. All rights reserved
Resource‐abundant metal (e.g., zinc) batteries feature intrinsic advantages of safety and sustainability. Their practical feasibility, however, is impeded by the poor reversibility of metal anode, typically caused by the uncontrollable dendrite enlargement. Significant effort has been exerted to completely prevent dendrites from forming, but this seems less effective at high current densities. Herein, we present an alternative dendrite regulation strategy of forming tiny, homogeneously distributed, and identical zinc dendrites by facet matching, which effectively avoids undesirable dendrite enlargement. Confirmed by multiscale theoretical screening and characterization, the regularly exposed Cu(111) facets at the ridges of a copper nanowire are capable of such dendrite regulation by forming a low‐mismatched Zn(002)/Cu(111) interface. Consequently, reversible zinc electroplating/stripping has been achieved at an unprecedentedly high rate of 100 mA cm−2 for over 30,000 cycles, corresponding to an accumulative areal capacity up to 30 Ah cm−2. A full cell using this anode shows a high capacity of 308.3 mAh g−1 and a high capacity retention of 91.4% after 800 cycles. This strategy is also viable for magnesium and aluminum anodes, thus opening up a promising and universal avenue towards long life and high rate metal anodes. This article is protected by copyright. All rights reserved
Zinc‐ion capacitors (ZICs) is promising technology for large‐scale energy storage by integrating the attributes of supercapacitors and zinc‐ion batteries. Unfortunately, the insufficient Zn2+ storage active sites of carbonaceous cathode materials and the mismatch of pore sizes with charge carriers led to unsatisfactory Zn2+ storage capability. Herein, we report new insights for boosting Zn2+ storage capability of activated nitrogen‐doped hierarchical porous carbon materials (ANHPC‐x) by effectively eliminating micropore confinement effect and synchronously elevate the utilization of active sites. Therefore, the best‐performed ANHPC‐2 delivers impressive electrochemical properties for ZICs in terms of excellent capacity (199.1 mAh g−1), energy density (155.2 Wh kg−1), and durability (65000 cycles). Systematic ex situ characterizations together with in situ electrochemical quartz crystal microbalance and Raman spectra measurements manifest that the remarkable electrochemical performance is assigned to the synergism of Zn2+, H+, and SO42− co‐adsorption mechanism and reversible chemical adsorption. Furthermore, the ANHPC‐2‐based quasi‐solid‐state ZIC demonstrates excellent electrochemical capability with ultralong lifespan up to 100 000 cycles. This work not only provides a promising strategy to improve the Zn2+ storage capability of carbonaceous materials but also sheds lights on charge storge mechanism and advanced electrode materials design for ZICs toward practical applications. This article is protected by copyright. All rights reserved
Rechargeable sodium ion micro‐batteries (NIMBs) constructed using low‐cost and abundant raw materials in planar configuration with both cathode and anode on the same substrate, hold promises for powering coplanar microelectronics, but are hindered by the low areal capacity owing to thin microelectrodes. Here, a prototype of planar and flexible 3D‐printed NIMBs is demonstrated with three‐dimensionally interconnected conductive thick microelectrodes for ultrahigh areal capacity and boosted rate capability. Rationally optimized 3D printable inks with appropriate viscosities and high conductivity allowed the multi‐layer printing of NIMB electrodes reaching a very high thickness of ∼1200 μm while maintaining effective ion and electron transfer pathways in them. Consequently, the 3D‐printed NIMBs deliver superior areal capacity of 4.5 mAh cm–2 (2 mA cm–2), outperforming the state‐of‐the‐art printed micro‐batteries. The NIMBs showed enhanced rate capability with 3.6 mAh cm–2 at 40 mA cm–2 and robust long‐term cycle life up to 6000 cycles. Furthermore, the planar NIMB microelectrodes despite the large thickness exhibit decent mechanical flexibility under various bending conditions. This work opens a new avenue for construction of high‐performance NIMBs with thick microelectrodes capable of powering flexible microelectronics. This article is protected by copyright. All rights reserved
We report a new type of an atomically thin synaptic network on van der Waals (vdW) heterostructures, where each ultra‐small cell (∼ 2 nm thick) built with trilayer WS2 semiconductor acts as a gate‐tunable photoactive synapse, i.e., a photo‐memtransistor. A train of ultraviolet (UV) pulses onto the WS2 memristor generates dopants in atomic‐level precision by direct light‐lattice interactions, along with the gate‐tunability, leading to the accurate modulation of the channel conductance for potentiation and depression of the synaptic cells. Such synaptic dynamics can be explained by a parallel atomistic resistor network model. In addition, we show that such device scheme can generally be realized in other two‐dimensional vdW semiconductors, such as MoS2, MoSe2, MoTe2 and WSe2. Demonstration of our atomically thin photo‐memtransistor arrays, where the synaptic weights can be tuned for the atomistic defect density, provides implications for a new type of artificial neural networks for parallel matrix computations with an ultra‐high integration density. This article is protected by copyright. All rights reserved
Recently, ferromagnetic heterostructure spintronic terahertz (THz) emitters have been recognized as one of the most promising candidates for the next‐generation THz sources, owing to their peculiarities of high efficiency, high stability, low cost, ultrabroad bandwidth, controllable polarization, and high scalability. Despite the substantial efforts, they rely on external magnetic fields to initiate the spin‐to‐charge conversion, which hitherto greatly limits its proliferation as practical devices. Here, we innovate a unique antiferromagnetic‐ferromagnetic (IrMn3|CoFeB) heterostructure and demonstrate that it can efficiently generate THz radiation without any external magnetic field. We assign it to the exchange bias or interfacial exchange coupling effect and enhanced anisotropy. By precisely balancing the exchange bias effect and the enhanced THz radiation efficiency, an optimized 5.6‐nm‐thick IrMn3|CoFeB|W tri‐layer heterostructure is successfully realized, yielding an intensity surpassing that of Pt|CoFeB|W. Moreover, the intensity of THz emission is further boosted by togethering the tri‐layer sample and bi‐layer sample Besides, the THz polarization may be flexibly controlled by rotating the sample azimuthal angle, manifesting sophisticated active THz field manipulation capability. The field‐free coherent THz emission we demonstrate here shines light on the development of spintronic THz optoelectronic devices. This article is protected by copyright. All rights reserved
Strain engineering is a promising way to tune the electrical, electrochemical, magnetic, and optical properties of two‐dimensional (2D) materials, with the potential to achieve high‐performance 2D‐material‐based devices ultimately. This review discusses the experimental and theoretical results from recent advances in the strain engineering of 2D materials. We summarize some novel methods to induce strain and then highlight the tunable electrical, and optical/optoelectronic properties of 2D materials via strain engineering including particularly the previously less discussed strain tuning of superconducting, magnetic, and electrochemical properties. Also, the future perspectives of strain engineering are given for its potential applications in functional devices. The state of the survey presents the ever‐increasing advantages and popularity of strain engineering for tuning properties of 2D materials. It provides suggestions and insights for further research and applications in optical, electronic, and spintronic devices. This article is protected by copyright. All rights reserved
Fabrication and morphological characterization of tin perovskite films. a) Fabrication procedure of tin perovskite films by using an antisolvent‐based normal spin‐coating or vapor‐assisted spin‐coating (VASC) method. The VASC method is conducted by adding 200 µL of mixed solvent to the spin‐coater to form solvent vapor before depositing perovskites. b–d) SEM images of different tin perovskite films. Control (b) refers to the film prepared with the normal spin‐coating method in the absence of the PH additive. VASC (c) refers to the film prepared with the VASC method in the absence of a PH additive. VASC+PH (d) refers to the film prepared with the VASC method in the presence of a PH additive. All scale bars represent 1 µm.
Crystallinity and optical properties of tin perovskite films prepared via various methods. a) X‐ray diffraction patterns of tin perovskite films on PEDOT:PSS substrates. b) Absorption and photoluminescence spectra. c) Excitation‐intensity‐dependent PLQE. d) Time‐resolved photoluminescence decay curves (excitation: 445 nm, 5.4 nJ cm–2).
Understanding the effect of VASC method on the crystallization of tin perovskite. a,b) The in situ absorption and PL evolution of the control and VASC films measured during the spin‐coating process. The antisolvent was dropped at 30 s after starting spin‐coating. c) Schematic crystallization pathways of tin perovskite films prepared with the control and VASC methods.
Architecture and device characteristics of tin‐based perovskite LEDs. a) Schematic of the LED device structure. b) Cross‐sectional SEM image of the VASC+PH device. Scale bar: 100 nm. c) Electroluminescence spectra. d) Dependence of the current density and radiance on the voltage. e) EQE values versus current density. A peak EQE of 5.3% was achieved for the VASC+PH device. f) EQE histogram.
Tin‐based perovskites are promising candidates to replace their toxic lead‐based counterparts in optoelectronic applications, such as light‐emitting diodes (LEDs). However, the development of tin perovskite LEDs is slow due to the challenge of obtaining high quality tin perovskite films. Here, a vapor‐assisted spin‐coating (VASC) method is developed to achieve high quality tin perovskites and high efficiency LEDs. It is revealed that solvent vapor can lead to in situ recrystallization of tin perovskites during the film formation process, thus significantly improving the crystalline quality with reduced defects. An antioxidant additive is further introduced to suppress the oxidation of Sn2+ and increase the photoluminescence quantum efficiency up to ∼30%, which is an ∼4‐fold enhancement in comparison with that of the control method. As a result, efficient tin perovskite LEDs are achieved with a peak external quantum efficiency of 5.3%, which is among the highest efficiency of lead‐free perovskite LEDs. This article is protected by copyright. All rights reserved
Atmospheric water harvesting (AWH) enabled by PAM–LiCl. a) Schematic illustration of AWH process via LiCl and PAM–LiCl. b) Photographs of PAM–LiCl before (top) and after moisture capture (bottom). PAM–LiCl exhibits clear swelling after water uptake, which supports the emergence of hydrogel network–water interactions. Scale bar: 5 mm. c) XRD patterns of PAM–LiCl before and after sorption, as well as after desorption.
Sorption/desorption performance of PAM–LiCl. a) Water sorption and desorption isotherm at 25 °C. b) Water sorption and desorption kinetics of PAM–LiCl and LiCl. Sorption condition: 20% RH, 25 °C; desorption condition: 10% RH, 70 °C. The desorption water vapor pressure equals the saturated water vapor pressure at 25 °C to mimic the actual working condition. c) Sorption/desorption performance of PAM–LiCl at 10%, 20%, and 30% RH at 25 °C. d) Water desorption curves of PAM–LiCl at 60, 70, 80, and 90 °C. The inset shows the change in water uptake between the first 10 to 50 min. At all the desorption temperatures, PAM–LiCl samples were dried at the same water vapor pressure of 3.17 kPa, equivalent to the saturated vapor pressure at 25 °C. RH at each temperature are: 16% RH, 60 °C; 10% RH, 70 °C; 7% RH, 80 °C; 4% RH, 90 °C.
Facilitated water desorption in PAM–LiCl. a) Heat flow curves of LiCl, PAM–LiCl, and PAM during water release measured by DSC. b) The heat of desorption of LiCl, PAM–LiCl, and PAM. c) Heat flow curves showing melting behavior of LiCl, PAM–LiCl at the same water content of 3 g g⁻¹. d) Molecular dynamic simulation snapshots of the PAM–LiCl and LiCl at 1.5 g g⁻¹ water content. Purple: Li, green: Cl, red: water molecular, blue: polymer chain. e) Percentage of water's hydrogen bond sites occupied by polymer (PW), ion (IW), and other water molecule (WW) interactions in the LiCl and PAM–LiCl simulations. f) Calculated diffusion coefficient of water molecules in the LiCl and PAM–LiCl simulations.
Atmospheric water harvesting (AWH) performance of PAM–LiCl. a) Cycling stability test of PAM–LiCl by DVS. Water sorption was conducted at 20% RH, 25 °C for 180 min, and desorption was conducted at 10% RH, 70 °C for 80 min. b) DVS sorption/desorption curves of PAM–LiCl at cycle 1 and cycle 15. c) Photograph showing the AWH setup is composed of a thermometer, an AWH chamber, and a power source. d) Schematic illustration of the AWH chamber. e) Cumulative water release and water collection varying with time. Inset photograph showing the produced water.
The ubiquitous nature of atmospheric moisture makes it a significant water resource available at any geographical location. Atmospheric water harvesting (AWH) technology, which extracts moisture from ambient air to generate clean water, is a promising strategy to realize decentralized water production. The high water uptake exhibited by salt‐based sorbents makes them attractive for AWH, especially in low relative humidity (RH) environments. Salt‐based sorbents often have relatively high desorption heat, rendering water release an energy‐intensive process. We proposed a hygroscopic gel, PAM hydrogel controlled incorporated with LiCl, capable of effective moisture harvesting from arid environments. The interactions between the hydrophilic hydrogel network and the captured water enable the PAM‐LiCl to accumulate more free and weakly‐bonded water molecules, significantly lowering the desorption heat compared with conventional neat salt sorbents. Benefiting from the affinity for swelling of the polymer backbones, the developed PAM‐LiCl achieves a high water uptake of ca. 1.1 g/g at 20% RH with fast sorption kinetics of ca. 0.008 g g–1 min–1 and further demonstrates a daily water yield up to ca. 7 g/g at this condition. These findings provide a new pathway for synthesis of materials with efficient water absorption/desorption properties, to reach energy‐efficient water release for AWH in arid climates. This article is protected by copyright. All rights reserved
Effect of K values on corrosion currents and overpotentials. a) Molecular models of ligands with different K values. b) Galvanostatic voltage profiles of Cu//Zn half cells with different ligands at 1 mA cm⁻². c) Potentiodynamic polarization curves of the Zn anode with different ligands. d) Relationship between logK and corrosion current as well as the overpotential. e) Cyclic plating/stripping process of symmetric cells. f) CE stability of Cu//Zn half cells with different ligands.
The solvation between Zn ions and ligands. a) The optimized molecular structures of Zn–ligand complexes and their chelated energies including Ac‐Zn, Ox‐Zn, Gly‐Zn, NTA‐Zn, and EDTA‐Zn. Gray, blue, red, pink, and purple balls denoted carbon, nitrogen, oxygen, hydrogen, and zinc atoms, respectively. b) Corrosion reaction pathways with different Zn–ligand complexes in a ZnSO4 electrolyte. c) Relationship between log K value and the chelation energy as well as the corrosion energy barrier. d) Energy intervals between the HOMO and LUMO of EDTA and EDTA‐Zn. e) 3D and f) 2D charge‐transfer patterns between EDTA and Zn ions. g) UV–vis spectrum of electrolytes with different EDTA‐Zn molar concentrations.
Molecular‐level investigation of solvated Zn ions on Zn metal. a) In situ Raman spectra of the anode surface with EDTA‐Zn additive during plating. b) Raman spectra of the Zn surface after Zn‐metal plating for 5, 10, 15, 20, 25, and 30 min. c) Raman spectra of the Zn surface with or without the EDTA‐Zn additive. d) OH stretch vibration. e) Dynamic evolution of an EDTA‐Zn molecule during Zn plating. f) Binding energy of EDTA or a H2O molecule adsorbed on a Zn (101) surface. g) PDOS patterns of a selected O atom in EDTA‐Zn before and after binding with Zn (101). h) PDOS patterns of a selected Zn atom in an EDTA–Zn complex at its initial and final positions.
Suppressing anode corrosion by Zn‐ligand solvation. a) XRD patterns of Zn foils after immersion for one week in electrolyte with and without EDTA‐Zn. b) Potentiodynamic polarization curves of the Zn anode with and without EDTA‐Zn. c) Electrochemical impedance spectroscopy (EIS) plots of Zn//Zn symmetric cells with and without EDTA‐Zn after immersion in a 2 m ZnSO4 electrolyte for 60 min. d,g) SEM images of a Zn foil after immersion in EDTA‐based and pristine 2 m ZnSO4 electrolytes, and e,h) their corresponding element distribution on the Zn surface, and f,i) contents of Zn, S, and O.
Inhibition of dendrite growth by EDTA‐Zn. a,b) SEM images of a Zn anode with an areal capacity of 1 mAh cm⁻²: a) with EDTA‐Zn and b) without EDTA‐Zn. c,d) Optical microscopy images of the Zn electrode during Zn plating: c) in the EDTA‐Zn‐added electrolyte; d) in the ZnSO4 electrolyte. e) Cyclic plating/stripping of symmetric cells with (red curve) and without (blue curve) EDTA‐Zn at a current density of 5 mA cm⁻² with a deposition capacity of 1 mAh cm⁻². f,g) Morphology of the Zn anode after 300 h plating/stripping at 20 mA cm⁻² without (f) and with (g) EDTA‐Zn. h) Cycling performance of NVO//Zn full batteries at 2 A g⁻¹.
Changing the solvation sheath of hydrated Zn ions is an effective strategy to stabilize Zn anodes to obtain a practical aqueous Zn‐ion battery. However, key points related to the rational design remain unclear including how the properties of the solvent molecules intrinsically regulate the solvated structure of the Zn ions. We propose using a stability constant (K), namely the equilibrium constant of the complexation reaction, as a universal standard to make an accurate selection of ligands in the electrolyte to improve the anode stability. It is found that K greatly impacts the corrosion current density and nucleation overpotential. Following this, ethylene diamine tetraacetic acid with a superhigh K effectively suppresses Zn corrosion and induces uniform Zn‐ion deposition. As a result, the anode has an excellent stability of over 3000 h. This work presents a general principle to stabilize anodes by regulating the solvation chemistry, guiding the development of novel electrolytes for sustainable aqueous batteries. This article is protected by copyright. All rights reserved
Morphology and elemental characterization of anisotropic CsPbBr3 single crystals. A,B) SEM image of as‐synthesized CsPbBr3 (111) tetrahedrons. Inset, Schematic illustrations of CsPbBr3 tetrahedrons with (111) exposed facet. C,D) SEM image of as‐synthesized CsPbBr3 (100) cubes. Inset: Schematic illustrations of CsPbBr3 (100) cubes. E–H) SEM–EDS analysis of a single CsPbBr3 (111) tetrahedron. I–L) SEM–EDS analysis of a single CsPbBr3 (100) tetrahedron. The above SEM–EDS elemental mapping shows Cs (blue), Pb (green) and Br (red) exhibited a relatively homogeneous distribution inside both CsPbBr3 cube and tetrahedron with stoichiometric ratio of ≈1:1:3.
Atomic‐resolution crystal structure analysis of anisotropic CsPbBr3 single crystals. A) XRD patterns of CsPbBr3 (111) tetrahedrons and (100) cubes. B) Schematic illustration of ultrathin CsPbBr3 single‐crystal TEM slice prepared by focused ion beam (FIB). C) Low‐magnification SEM image of a single FIB slice for CsPbBr3 (100) cube from [210] zone axes. Inset: FIB cutting position for characterization. D) Cross‐section SEM image of FIB slice for CsPbBr3 (100) cube from [100] zone axes. Inset: FIB cutting position for characterization. E) Cross‐section SEM image of FIB slice for CsPbBr3 (111) tetrahedrons from [111] zone axes. Inset: FIB cutting position for characterization. F–H) AC‐HRTEM images taken from the selected areas in (C–E). Insets: Atomic‐resolution HRTEM and the corresponding overlaid atom model. I) SAED image taken from FIB slice of CsPbBr3 (100) cube.
Optical characterization of anisotropic CsPbBr3 single crystals. A,B) Photoemission cutoff and fermi edge UPS spectra of CsPbBr3 (111) tetrahedrons and (100) cubes. C,D) Absorption spectra of CsPbBr3 (111) tetrahedrons and (100) cubes. Insets of (C) and (D), absorbance versus photon energy and the determined bandgap Eg. E) Photoluminescence (PL) spectra of CsPbBr3 (111) tetrahedrons and (100) cubes. F) PL time decay trace on CsPbBr3 (111) tetrahedrons and (100) cubes.
Electronic device and electrical characterization of anisotropic CsPbBr3 single crystals. A) Schematic device structure of the anisotropic CsPbBr3 single crystal device. B) SEM image of CsPbBr3 (111) tetrahedrons and (100) cubes with Pt nanowire as metal electrodes. Scale bar: 2 µm. C,D) Dark I–V and logarithmic I–V plot characteristics of the devices. E,F) I−V trace of the devices for space‐charge‐limited current (SCLC) analysis. G,H) Anisotropic migration of carrier (electron and hole) on (100) and (111) CsPbBr3 crystal planes calculated by first‐principles. I) Surface defect formation energy of Cs, Pb, Br on (100) and (111) CsPbBr3 crystal planes calculated by first‐principles.
Engineering surface structure can precisely and effectively tune optoelectronic properties of halide perovskites, but are incredibly challenging. Herein, we report the design and fabrication of uniform all–inorganic CsPbBr3 cubes/tetrahedrons single‐crystals with precise control of (100) and (111) surface anisotropy, respectively. By combining theoretical calculations, we demonstrated that the preferred (100) surface engineering of CsPbBr3 single‐crystals enables a lowest surface bandgap energy (2.33 eV) and high‐rate carrier mobility up to 241 μm2 V–1 s–1, inherently boosting their light‐harvesting and carrier transport capability. Whereas, the polar (111) surface induces ∼0.16 eV upward surface‐band bending and ultrahigh surface defect density of 1.49 × 1015 cm–3, which is beneficial for enhancing surface defects‐catalyzed reactions. Our work highlights the anisotropic surface engineering for boosting perovskite optoelectronic devices and beyond. This article is protected by copyright. All rights reserved
a–c) STEM images of Cu‐Bi24O31Br10‐3, d) intensity profile corresponding to the dark cyan arrow in (c) and (g). e,f) EDS elemental mapping of Cu‐Bi24O31Br10‐3. g) Simulated STEM image with one Cu replacing Bi. h,i) STEM images of Cu‐Bi24O31Br10‐4.
a) Synchrotron radiation Cu K‐edge XANES, b) Cu 2p XPS, c) XAFS Bi L1‐edge, d) EXAFS spectra of Bi L1‐edge. e,f) Electrostatic potential of Bi24O31Br10 (e) and Cu‐Bi24O31Br10 (f). g,h) Electron localization function of Bi24O31Br10 (g) and Cu‐Bi24O31Br10 (h).
a) Photocatalytic NH3 production over Cu‐Bi24O31Br10 materials in pure water. b) Mass spectra of the products formed via reaction with indophenol following photocatalysis in different reaction atmospheres. c) Performance comparisons. d) Cycling test data over Cu‐Bi24O31Br10‐3. e) UV–vis diffuse reflectance spectra. f) Ultrafast transient absorption spectroscopy of Cu‐Bi24O31Br10‐3. g–i) The atomic charges distribution of Bi24O31Br10 (g) and Cu‐Bi24O31Br10 (h). i) In situ FTIR spectra from Cu‐Bi24O31Br10‐3.
a–d) Calculated Gibbs free energy profiles of the nitrogen reduction process at: a,b) Bi24O31Br10 with the associative distal pathway (a) and the alternating pathway (b), and c,d) Cu‐Bi24O31Br10 with the associative alternating pathway (c) and the distal pathway (d). e,f,h,i) Isosurface maps of IRI = 1.0 between N2 and Bi24O31Br10 (e), *NNH and Bi24O31Br10 (f), N2 and Cu‐Bi24O31Br10 (h) and *NNH and Cu‐Bi24O31Br10 (i). Envelopes of g,j) Laplacian of electron density (iso = 0.8 a.u.) of Bi24O31Br10 (g) and Cu‐Bi24O31Br10 (j). IRI is defined as a real space function that can be used to clearly exhibit chemical bonds, weak interactions, and repulsive interactions. The blue, green, and red color in Isosurface maps of IRI indicate strong attractive interactions (chemical bonds), weak attractive interactions (vdW force), and repulsive interactions, respectively, as shown in the color bar.
A universal atomic layer confined doping strategy is developed to prepare isolated Cu atoms incorporated Bi24O31Br10 materials. The local polarization can be created along Cu‐O‐Bi atomic interface, which enables better electron delocalization for effective N2 activation. The optimized Cu‐Bi24O31Br10 atomic layers show 5.3 and 88.2 times improved photocatalytic nitrogen fixation activity than Bi24O31Br10 atomic layer and bulk Bi24O31Br10, respectively, with the NH3 generation rate arrive 291.1 μmol g−1 h−1 in pure water. The polarized Cu‐Bi site pairs can increase the non‐covalent interaction between catalyst's surface and N2 molecule, then further weaken the covalent bond order in N‐N. As a result, the hydrogenation pathways can be altered from associative distal pathway for Bi24O31Br10 to the alternating pathway for Cu‐Bi24O31Br10. This strategy provides an accessible pathway for designing polarized metal site pairs or tune the non‐covalent interaction and covalent bond order. This article is protected by copyright. All rights reserved
Preparation and structural characterization of Ru/h‐BN catalyst. a) Schematic of preparation of Ru/h‐BN catalyst. b) TEM image of synthesized Ru/h‐BN catalyst. c) High‐resolution TEM image of Ru nanoparticles supported on h‐BN. d) Fast Fourier‐transforms of boxed regions shown in (c). The hexagonal patterns from Ru{100} and h‐BN{100} lattices are marked with circles and dotted lines, respectively. e) Atomistic models of Ru/h‐BN observed in (c). f) HAADF‐STEM image of Ru/h‐BN and B5 site, viewed in [001] direction of h‐BN and Ru. The left panel shows the entire shape of the Ru nanoparticle. The straight (100) edges of the Ru nanoparticle are marked with dotted lines. The top‐right panel shows low‐pass filtered, magnified image of the boxed region in the left panel. The bottom‐right panel shows intensity profiles along the lines presented in the top‐right panel. The stronger contrast at location B in profiles 1 and 3 indicates the presence of step sites at the Ru nanoparticle edge as depicted in the atomistic model. g) Cross‐sectional HAADF‐STEM image of Ru/h‐BN (left). The top‐right panel shows low‐pass filtered, magnified image of the boxed region in the left panel. The bottom‐right panel shows atomistic model of the B5 site observed in the top‐right panel. h) Schematic of the model Ru nanoparticle used for B5 site density estimation (Supporting Information). Locations where B5 sites can exist are marked with orange lines. The inset shows atomic model of B5 site. i) Dependence of B5 site density on the height and radius of curvature in the model Ru nanoparticle. j) B5 site density as a function of radius of curvature in an average Ru nanoparticle observed in our analyses. The result corresponds to the dotted line marked in (i).
Activation of Ru/h‐BN catalyst under reaction conditions. a) Comparison of stability of Ru/h‐BN, Ru/Al2O3, and Ru/SiO2 catalysts for ammonia decomposition at a GHSV of 60 000 mLNH3${\rm{m}}{{\rm{L}}_{{\rm{N}}{{\rm{H}}_{\rm{3}}}}}$ gcat⁻¹ h⁻¹ and a temperature of 450 °C over 80 h. b–d) Size distributions of Ru particles in the Ru/h‐BN, Ru/Al2O3, and Ru/SiO2 catalysts, respectively, before and after 80 h of reaction. The size d of the elongated Ru nanoparticle represents the distance along its long axis. The mean 〈d〉 and standard deviation σd of the nanoparticle size d were measured using at least 150 nanoparticles. Here, the Ru particle size in Ru/SiO2 after the 80 h reaction is not provided because of severe sintering of the Ru particles, exceeding the plot range. e) HAADF‐STEM images of Ru particles supported on h‐BN after 0, 12, and 80 h of reaction. f) Adsorption energies of Ru nanoparticles composed of 60 (left), 53 (middle) and 41 (right) Ru atoms on h‐BN(001) surface with varying numbers of Ru atoms at the interfaces. g) Time‐series TEM images of Ru particle on h‐BN sheet, observed at 450 °C under vacuum. The edge that becomes faceted over the reaction time is marked with red arrows. h) Contours of the Ru particle observed in (g) at 0 and 50 min from the onset of observation. i) H2‐TPR profiles of Ru/h‐BN, Ru/Al2O3, and Ru/SiO2. j) Ru 3p (left) and 3d (right) XPS profiles of unactivated and 12 h‐activated Ru/h‐BN catalysts. For the Ru 3p XPS spectra, the fitted peak locations of the Ru species are shown in parentheses.
Evaluation of catalytic performance. a) Ammonia conversion over Ru/h‐BN before activation (orange symbols), Ru/h‐BN after 12 h of activation (red symbols), Ru/Al2O3 (green symbols), and Ru/SiO2 (blue symbols), measured at GHSV of 60 000 mLNH3${\rm{m}}{{\rm{L}}_{{\rm{N}}{{\rm{H}}_{\rm{3}}}}}$ gcat⁻¹ h⁻¹ and different temperatures (350, 375, 400, 425, and 450 °C). Error bars smaller than the symbols are omitted. b) H2 production rates over Ru/h‐BN before activation (orange), Ru/h‐BN after 12 h of activation (red), Ru/Al2O3 (green), and Ru/SiO2 (blue), measured at a GHSV of 60 000 mLNH3${\rm{m}}{{\rm{L}}_{{\rm{N}}{{\rm{H}}_{\rm{3}}}}}$ gcat⁻¹ h⁻¹ and a temperature of 450 °C. c) Calculation of apparent activation energies of Ru/h‐BN before activation (orange symbols), Ru/h‐BN after 12 h of activation (red symbols), Ru/Al2O3 (green symbols), and Ru/SiO2 (blue symbols) using the Arrhenius equation. d) N2‐TPD spectra of Ru/h‐BN activated for different times. Desorption peaks attributed to step and terrace sites are deconvoluted and marked in red and gray, respectively. e) Mass activity and f) specific activity of Ru/h‐BN for ammonia dehydrogenation, measured at a GHSV of 60 000 mLNH3${\rm{m}}{{\rm{L}}_{{\rm{N}}{{\rm{H}}_{\rm{3}}}}}$ gcat⁻¹ h⁻¹ and a temperature of 450 °C, compared to those of reported catalysts.[36,44–49]
Ruthenium is one of the most active catalysts for ammonia dehydrogenation and is essential for the use of ammonia as a hydrogen storage material. The B5‐type site on the surface of ruthenium is expected to exhibit the highest catalytic activity for ammonia dehydrogenation, but the number of these sites is typically low. Here, we synthesize a B5‐site‐rich ruthenium catalyst by exploiting the crystal symmetry of a hexagonal boron nitride support. In the prepared ruthenium catalyst, ruthenium nanoparticles are formed epitaxially on hexagonal boron nitride sheets with hexagonal planar morphologies, in which the B5 sites predominate along the nanoparticle edges. By activating the catalyst under the reaction condition, the population of B5 sites further increases as the facets of the ruthenium nanoparticles develop. The electron density of the Ru nanoparticles also increases during catalyst activation. The synthesized catalyst shows superior catalytic activity for ammonia dehydrogenation compared to previously reported catalysts. This work demonstrates that morphology control of a catalyst via support‐driven heteroepitaxy can be exploited for synthesizing highly active heterogeneous catalysts with tailored atomic structures. This article is protected by copyright. All rights reserved
Schematic illustration of the construction of high‐speed solid‐state Li transport in the graphite anode based on the LixCu6Sn5 intermetallic crystal structure. a) The crystal structure model of η‐Cu6Sn5 ([001] projection). The arrows denote the Sn migration direction during lithium insertion. b) The crystal structure model of lithiated LixCu6Sn5 (0 < x < 13) ([110] projection). A 3D Li transport channel is formed in the LixCu6Sn5 within the zinc blende‐type crystal structure. c) Solid lithium transport channel construction based on the Cu@CuSn NWs network through the electrode to largely reduce the lithium‐ion concentration gradients and the corresponding polarization in the graphite anode.
Characterizations of the synthesized Cu@CuSn NWs. a) PXRD patterns of the synthesized Cu@CuSn NWs. b) High‐resolution TEM image of the Cu@CuSn NW showing the lattice fringes of (202) planes of intermetallic Cu6Sn5. c) Typical charge–discharge curve of Cu@CuSn NW membrane. d) PXRD of Cu@CuSn NW membrane at lithiation state after 30 cycles. The charge–discharge current density is 100 mA g⁻¹. e) The schematic to show that Cu3Sn buffer layer in the Cu@CuSn can enhance the adhesion between the Cu NW core and the LixCu6Sn5 shell. f) Power‐law dependence of current on the sweep rate of 0.2 V, b = 0.82. g) Schematic illustration of a symmetric cell geometry for testing diffusion coefficients. h,i) Plots of the relaxation phase of a steady‐state polarization experiment in the cell with Cu NWs (h) and Cu@CuSn NWs (i) membrane, respectively.
XPS chemical state mappings of Li 1s in the cross section of lithiated graphite anodes. a–c) The cross section SEM and the corresponding Cu and Sn elementary mapping images of G‐SLTC, respectively. d–f) The Li 1s XPS mapping image of G‐SLTC. The X‐ray induced secondary‐electron images (SXI) of the cross section (d), Li 1s overlapped chemical state mapping (e), and the corresponding XPS spectra of selected zones in Figure 3e (f). g–i) The Li 1s XPS mapping image of graphite anode. The SXI of the cross section (g), Li 1s overlapped chemical state mapping (h), and the corresponding XPS spectra of selected zones in Figure 3h (i).
Fast charging performance of full cell with the G‐SLTC anode and traditional graphite anode. a) Typical voltage profiles of full cells at the rate of 0.1 C. b) The improved CE and specific capacity of the cell with G‐SLTC and graphite anode after pre‐lithiation. c) The charging rate performance of the cells. The discharging rate is fixed at 0.5 C and the charging rate is changed from 0.2 C to 6 C. d) Typical charging voltage curves at high rate of 2 C and 3 C with constant current charging and constant voltage charging processes. e) The current rate versus capacity curve of full cells at the constant voltage process. f) Time required to charge to 140 mAh g⁻¹ at the rate of 2 C and 3 C.
Electrochemical stability of full cell at high charging rate. a) Long‐term cycling performance of full cell at a high charging rate of 2 C. The discharging rate is fixed at 0.5 C. b) The battery energy utilization of full cells in the long‐term cycles. The battery energy utilization is defined as the ratio of the charging energy to the discharging energy. c) Irreversible increase in the thickness of the electrode at delithiated state for 20 cycles.
Superior fast charging is a desirable capability of lithium‐ion batteries, which can make electric vehicles a strong competition to traditional fuel vehicles. However, the slow transport of solvated lithium ions in liquid electrolytes is a limiting factor. Here, we report a LixCu6Sn5 intermetallic network to address this issue. Based on electrochemical analysis and X‐ray photoelectron spectroscopy mapping, we demonstrate that the reported intermetallic network can form a high‐speed solid‐state lithium transport matrix throughout the electrode, which largely reduces the polarization effect in the graphite anode. Employing this design, we fabricated superior fast‐charging graphite/lithium cobalt oxide full cells and tested them under strict electrode conditions. At the charging rate of 6 C, the fabricated full cells showed a capacity of 145 mAh g−1 with an extraordinary capacity retention of 96.6%. In addition, the full cell also exhibits good electrochemical stability at a high charging rate of 2 C over 100 cycles (96.0% of capacity retention) in comparison to traditional graphite anode based cell (86.1% of capacity retention). This work presents a new strategy for fast charging lithium ion batteries on basis of high speed solid state lithium transport in intermetallic alloy hosts. This article is protected by copyright. All rights reserved
Fabrication and properties of Cu3(HHTT)2 thin films. a) Synthetic route and crystal structure of Cu3(HHTT)2. b) Fabrication procedure of Cu3(HHTT)2 thin films. c) AFM image of a Cu3(HHTT)2 thin film with a thickness of 132 nm. d) Out‐of‐plane (left) and in‐plane (right) GIXRD patterns of a 132 nm thick Cu3(HHTT)2 thin film. e) HRTEM image of a Cu3(HHTT)2 film and the corresponding FFT image (inset). f) Tauc plot of the absorption spectrum of a 132 nm thick Cu3(HHTT)2 thin film.
Temperature‐dependent transport property and derived schematic band diagrams. a) σ versus 1000/T relationship in two regions with different thermal activation energy (300–90 K). b) ln(σ) versus 1/T1/4 relationship at low T (90–10 K), showing the Mott variable range hopping behavior (55–10 K). c) Double‐logarithm plot of I–V curves at various temperatures. The red and blue regions correspond to Ohmic and space‐charge‐limited current (SCLC) conductions, respectively. d) SCLC mobility versus 1000/T relationship (255–120 K). e) Schematic band diagram used to explain the temperature‐dependent conductivity. f) Illustration of different transport mechanisms. Red, blue, and orange dashed lines correspond to carrier transport between tail states by nearest‐neighbor hopping, nearest neighbor hopping between mid‐gap states, and variable‐range hopping between mid‐gap states near Fermi level, respectively.
Device performance of photodetectors based on Cu3(HHTT)2 thin films. a) Schematic diagram of the device structure. b) I–V curves of the devices with different cycle numbers. The films with 4, 6, 8, 12 cycles have the thicknesses of 37, 62, 76, and 132 nm, respectively. c) Responsivity as a function of light intensity for three different wavelengths. d) Spectral response for different light wavelengths from 370 to 3400 nm. The error bars were calculated from four devices. e) Transient response for a photodetector under three 650 nm laser light on–off cycles.
Device performance of flexible photodetectors based on Cu3(HHTT)2 thin films. a) Responsivity as a function of light intensity for 685 nm light wavelength. b) Transient response for a flexible photodetector under three 650 nm laser light on–off cycles. c) Photocurrent of a device as a function of bending cycle under bending tests. Error bars are calculated from four devices. Inset: the photographs demonstrate the device subjected to bending tests.
Device performance of an optical synapse device based on Cu3(HHTT)2 thin films. a) Schematic illustration of a biological synapse. b) Gradual decrease of ∆PSC after triggered by an optical signal (685 nm, 840 μW cm⁻²). c) PPF behavior under stimulation of two consecutive optical signals. d) LTP/LTD characteristics under continuous light on/off cycles (685 nm, 370 μW cm⁻²). On/off times of 8 s/2 s and 2 s/8 s are used to mimic LTP and LTD side of the curve, respectively. e) Schematic illustration of the three‐layer neural network for data type analysis. f) Data type recognition accuracy as a function of training epoch.
Cu3(HHTT)2 (HHTT: 2,3,7,8,12,13‐hexahydroxytetraazanaphthotetraphene) is a novel two‐dimensional conjugated metal‐organic framework (2D c‐MOF) with efficient in‐plane d‐π conjugations and strong interlayer π‐π interactions while the growth of Cu3(HHTT)2 thin films has never been reported until now. Here, we present the successful fabrication of highly oriented wafer‐scale Cu3(HHTT)2 thin films with a layer‐by‐layer growth method on various substrates. Its semiconducting behavior and carrier transport mechanisms are clarified through temperature and frequency dependent conductivity measurements. Flexible photodetectors based on Cu3(HHTT)2 thin films exhibit reliable photo‐responses at room temperature in a wavelength region from ultraviolet (UV) to mid‐infrared (MIR), which is much broader than those of solution‐processed broadband photodetectors reported previously. Moreover, the photodetectors can show a typical synaptic behavior and excellent data recognition accuracy in artificial neural networks. This work opens a window for the exploration of high‐performance and multi‐functional optoelectronic devices based on 2D c‐MOFs. This article is protected by copyright. All rights reserved
Temperature‐dependent conductivities of neat and hybrid samples probed through DRS. a) The conductivities of poly(DOL) (1 × 10⁻³ m Al(OTf)3, green open circles and 10 × 10⁻³ m Al(OTf)3, blue open stars), self‐suspended PEG–SiO2 HNPs (ϕc = 15 vol%, orange closed circles), and hybrid system composed of both (ϕc = 2.7 vol%, red open diamonds). All samples include 2 m LiTFSI salt. b) Temperature‐dependent DC conductivity of hybrid system (ϕc = 2.7 vol%) polymerized with 10 × 10⁻³ m Al(OTf)3 for the 1st and 5th hour (as insets, shown with respect to time sweep result from Figure 2c), 24th hour, 29th hour, 48th hour, and 10th day of polymerization. Conductivity values shown in plots are room temperature conductivity and Ea values were fit either by VFT fit or Arrhenius depending on the best fit. Some error bars are smaller than the size of the scatter plot data points. Each DRS measurement was done over two samples and three iterations.
Small‐angle X‐ray scattering (SAXS) profiles to determine the structure of PEG–SiO2 HNPs/poly(DOL) hybrid material. a) Intensity profile and b) structure factor of HNPs/poly(DOL) with ϕc = 2.7 vol% with initiator Al(OTf)3 content varying from 10 to 50 × 10⁻³ m. d) Structure factor analysis gives interparticle distance dp−p (red data points) and distance between tethers (green data points) as described in the diagram in (c), as well as structure factor at q = 0. The dashed lines in (d) represent the result for self‐suspended HNPs with ϕc = 11 vol%. Each measurement was done over two samples and three iterations.
Time‐dependent dynamic shear flow measurements are used to study changes in polymerization kinetics induced by PEG–SiO2 HNPs in a poly(DOL) SPE. a) Time‐sweep measurement of DOL monomers polymerized with 10 × 10⁻³ m Al(OTf)3 with the addition of: b) 10 × 10⁻³ m LiNO3, c) PEG–SiO2 HNPs (ϕc = 2.7 vol%). d) DOL polymerized with 1 × 10⁻³ m Al(OTf)3 with PEG–SiO2 HNPs (ϕc = 2.7 vol%). This sample was measured using a 25 mm cone and plate geometry to ensure torque values above the instrument's lower limit. The inset in (d) shows a hybrid 1 × 10⁻³ m initiator system with the addition of 2 m LiTFSI. The inset in (a) represents a strain sweep at 30 min time‐point during polymerization. The schematics of polymerization process going from monomers to amorphous polymer to polymer crystals were displayed in (a). The dashed line in (b) shows data point from (a) as a comparison between poly(DOL) polymerization with and without LiNO3. The inset in (c) shows clearer peaks in G′′ and G′ at early time. All room‐temperature time‐sweep measurements were carried out at an angular frequency of ω = 10 rad s⁻¹ and strain of γ = 5%. Strain sweep was measured also at RT and ω = 10 rad s⁻¹. Al(OTf)3 contents of 1, 20–50 × 10⁻³ m for each type of sample are shown in Figures S5 and S6, Supporting Information.
Electrochemical performance of HNPs/poly(DOL) hybrid electrolyte. a) Schematic of in situ polymerization of DOL in the presence of HNPs. Slurry of HNPs, monomer, and initiator was added, ensuring good wettability before the mixture was let sit for at least 24 h of polymerization process. b) CE of various poly(DOL) electrolyte and hybrids containing PEG–SiO2 HNPs (ϕc = 2.7%) with and without 0.5 m LiNO3 at 0.1 mA cm⁻² in Li||Cu cells. c) Galvanostatic cycling profile for Li||sPAN cell with hybrid electrolyte containing ϕc = 2.7%, 2 m LiTFSI, and 0.5 m LiNO3, with d) discharge capacity and CE over cycle. e) Rate performances for the Li||sPAN cells at different c‐rates with 1.0C is current density of 1 mA cm⁻². All electrolytes include 2 m LiTFSI salt and all PolyDOL electrolytes were polymerized with 1 × 10⁻³ m Al(OTf)3. Galvanostatic cycling profile for Li||LFP cell with the same electrolyte along with its discharge capacity profile and CE is presented in Figure S16, Supporting Information.
Solid‐state electrolytes (SSEs) formed inside an electrochemical cell by polymerization of a liquid precursor provide a promising strategy for overcoming problems with electrolyte wetting in solid‐state batteries. Hybrid solid‐state polymer electrolytes (HSPEs) created by in situ polymerization of a conventional liquid precursor containing electrochemically inert nanostructures are of particular interest because they offer a mechanism for selectively reinforcing or adding new functionalities to the electrolyte—removing the need for high degrees of polymerization. The synthesis, structure, chemical kinetics, ion‐transport properties and electrochemical characteristics of HSPEs created by Al(OTf)3‐initiated polymerization of 1,3‐dioxolane (DOL) containing hairy, nano‐sized SiO2 particles are reported. Small‐angle X‐ray scattering reveals the particles are well‐dispersed in liquid DOL. Strong interaction between poly(ethylene glycol) molecules tethered to the SiO2 particles and poly(DOL) lead to co‐crystallization—anchoring the nanoparticles in their host It also enables polymerization–depolymerization processes in DOL to be studied and controlled. The utility of the in‐situ‐formed HSPE, is demonstrated first in Li|HSPE|Cu half cells, which manifest Coulombic efficiencies (CE) values approaching 99%. HSPEs are also demonstrated in solid‐state lithium–sulfur–polyacrylonitrile (SPAN) composite full‐cell batteries. The in‐situ‐formed Li|HSPE|SPAN cells show good cycling stability and thus provide a promising path toward all‐solid‐state batteries.
Developing advanced electrocatalysts with exceptional two electron (2e–) selectivity, activity and stability are crucial for driving oxygen reduction reaction (ORR) to produce hydrogen peroxide (H2O2). Herein, a composition engineering strategy has been adapted to flexibly regulate the intrinsic activity of amorphous nickel boride nanoarchitectures for efficient 2e– ORR by oriented reduction of Ni2+ with different amounts of BH4–. Among borides, the amorphous NiB2 delivers the 2e– selectivity close to 99% at 0.4 V and over 93% in a wide potential range, together with a negligible activity decay under prolonged time. Notably, an ultrahigh H2O2 production rate of 4.753 mol gcat−1 h−1 has been achieved upon assembling NiB2 in the practical gas diffusion electrode. The combination of X‐ray absorption and in situ Raman spectroscopy, as well as transient photovoltage measurements with density functional theory unequivocally reveal that the atomic ratio between Ni and B induces the local electronic structure diversity, allowing optimization of the adsorption energy of Ni towards *OOH and reducing the interfacial charge transfer kinetics to inhibit the formation of O‐O. This article is protected by copyright. All rights reserved
Overview of the manufacturing workflow and rheological characterization of salt‐containing inks. a) Schematics depicting the layer‐by‐layer stereolithographic printing of a resin containing NaCl particles, surfactant, photocurable monomers, photo‐absorber, photoinitiator, and a non‐reactive diluent. The printed body is calcined to remove the organics and further sintered to obtain a dense, binderless NaCl body. This mold is infiltrated by a desired material and finally leached to obtain the positive complex‐shaped body. b–d) The influence of the b) surfactant, c) salt, and d) diluent concentrations on the yield stress and apparent viscosity of the inks. No yielding is observed for inks with surfactant concentrations below 0.1 wt% (with respect to NaCl), which makes them unsuitable for the printing process. Orange encircled data indicates the optimal composition that was identified.
Photo‐polymerization behavior and printing fidelity of salt‐containing resins. a–c) The influence of the applied light dose (Dmax) on the cure depth (zp) of inks containing initiator concentrations in the range 1–5 wt% with respect to monomer. Fixed photo‐absorber contents of a) 0 wt%, b) 0.05 wt%, and c) 0.125 wt% are used in these experiments. The corresponding light penetration depth (ha) and critical dose (Dc) obtained by fitting the modified Beer‐Lambert law to the experimental data are indicated in the insets. The results are summarized in d) to highlight the effect of the photo‐absorber and photo‐initiator concentrations on ha and Dc. e) 2D sketch defining the parameters analyzed and illustrating the region of the ink that is illuminated with a dose equal or higher than Dc (dark orange). f,g) Printing fidelity of inks with 0, 0.05, and 0.125 wt% photo‐absorber for f) negative and g) positive features printed at a cure depth (ha) of 80 µm. The insets depict the average standard deviation of all size measurements (0–2 mm) for different photo‐absorber concentrations. Optical microscopy images show representative examples of the 2 mm designed features, whereas the dashed lines indicate the theoretical size thereof.
Shape complexity and cracking inhibition in salt‐containing printed inks. a) Complex‐shaped objects printed using optimized ink formulations. b) Measured cracked area in printed cubes of different side lengths after heat treatment at 200 °C (dark color) or 690 °C (bright color). Optical microscopy images in light transmission mode (top) illustrate the effect of 30 wt% camphor on the formation of cracks (bright areas). Scale bars: 1 mm. c) Shape changes upon drying of model layers printed using salt‐containing resins with or without camphor. All model layers are illuminated on the bottom side, and a salt‐free resin (top row) is used as control. d) Proposed mechanism for crack inhibition in inks containing camphor and salt particles. The drawings indicate the distinct crosslinking densities and amount of residual monomer expected at the bottom (B,D) and top (A,C) of a single printed layer. e) Schematics showing the presence of a percolating network of salt particles resulting from the shrinkage of the polymer continuous phase upon heating of the printed material to 200 °C. Shrinkage of the polymer phase (blue) beyond particle jamming results in detachment from the particles and the formation of interstitial pores, as indicated in the false‐colored SEM image. Scale bar: 2 µm. f) Pore size distribution and total surface area (S. Area, BET Analysis, Inset) of printed inks without (blue) or with (orange) camphor after drying at 30 and 200 °C. g) Storage (G′) and loss (G′′) moduli of printed bars subjected to DMA in torsion mode. Samples were heated to 230 °C, followed by an isothermal hold at 230 °C for 20 min.
Complex‐shaped structures made through infiltration and leaching of salt molds. a) Comparison of gyroid digital model (1) with microcomputed tomography (microCT) analysis of a printed (2), sintered salt template (3), and molded and leached silicone scaffold (4). X and Y denote specific cross‐sections, as indicated in (1). The linear shrinkage along different directions is indicated (in blue) as a percentage of the initial digital model. The scale bar of 5 mm is valid for (1–4). b) Cell viability analysis of gyroid scaffolds prepared by leaching the 3D printed NaCl templates from infiltrated samples (either silicone or PCL). Left: Quantitative analysis of cell viability from live/dead assay on three different scaffold types made from silicone or PCL (n = 5–10). The cell viability is >94% 2 days after seeding for all three scaffolds. Right: Representative confocal laser scanning microscopy image showing live/dead assay results of MC3T3‐E1 pre‐osteoblasts that have been seeded on a fibronectin‐coated silicone scaffold with a pore size of 150 µm (green: live, red: dead). c) Examples of sintered salt molds and the corresponding complex‐shaped structures of a range of materials obtained after infiltration and leaching steps. Top left: tracheal stent made of bioresorbable poly(DLLA‐co‐CL) copolymer. Top right: ultralightweight octagon lattice made from an Al‐Si 12.6% metallic alloy. Bottom left: edible bunny made from dark chocolate. Bottom right: hollow reinforced tube made by covering a salt core with a carbon fiber composite.
Three‐dimensional printing is a powerful manufacturing technology for shaping materials into complex structures. While the palette of printable materials continues to expand, the rheological and chemical requisites for printing are not always easy to fulfill. Here, we report a universal manufacturing platform for shaping materials into intricate geometries without the need for their printability, but instead using light‐based printed salt structures as leachable molds. The salt structures are printed using photocurable resins loaded with NaCl particles. The printing, debinding and sintering steps involved in the process are systematically investigated to identify ink formulations enabling the preparation of crack‐free salt templates. Our experiments reveal that the formation of a load‐bearing network of salt particles is essential to prevent cracking of the mold during the process. By infiltrating the sintered salt molds and leaching the template in water, we create complex‐shaped architectures from diverse compositions such as biomedical silicone, chocolate, light metals, degradable elastomers and fiber composites, thus demonstrating the universal, cost‐effective, and sustainable nature of this new manufacturing platform. This article is protected by copyright. All rights reserved
Conducting, conical domain walls in LiNbO3. a) Piezoresponse force microscopy (PFM) amplitude and b) conducting atomic force microscopy (cAFM) maps of domains obtained on the top surface of 500 nm‐thick partially switched lithium niobate. Domains were created using a rastered AFM tip as a top electrode. Scale bar: 700 nm. c) PFM amplitude from the top surface of a similar LiNbO3 (LNO) thin film sample after partial switching using a mesoscale liquid top electrode. Scale bar: 150 nm. d) Cross‐sectional high‐angle annular dark‐field scanning transmission electron microscopy (HAADF‐STEM) image of the domains in LNO. The overlaid lines highlight the inclination of the walls. Scale bar: 200 nm. e) A schematic of the conducting domain wall cones in LNO, as inferred from the top surface PFM in (a) and (c), and the cross‐sectional HAADF‐STEM in (d). The dark gray and yellow rings highlight the locus of the electrical contacts made with the top and bottom electrodes, respectively.
Domain walls as Corbino disks. a) Experimental setup for the geometric magnetoresistance measurement. b) Schematic depicting conducting domain wall in LNO, along with its 2D projection, which takes the form of the Corbino disk. The black arrow shows the direction of the magnetic field used in the geometric MR measurement depicted in (a), and the various dotted arrows show the current components found through an iterative description of the current density. c) The typical Corbino disk geometry, employed for geometric magnetoresistance measurements. It consists of a sample under investigation (blue annulus) and concentric inner and outer electrodes (black and yellow rings). Electronic motion in the presence and absence of a perpendicular magnetic field is indicated by full green and dashed black arrows, respectively.
Geometric magnetoresistance (MR) measurement at conducting LNO domain walls. a) The MR response. The red “plus” motifs show the applied external magnetic field and the black circles show the normalized current response, averaged over 20 cycles. The blue dashed lines are parabolic fits (Equation (9)). Note that all fits are weighted by the uncertainty on each data point. b) Discrete fast Fourier transforms of the measured raw current and magnetic field profiles, shown for two datasets. c) A plot of the magnetic field versus the MR. The blue dashed line is a fitted parabola with B² coefficient (6.1 ± 0.5) × 10⁻³. d) A plot of log(B) versus log(MR), with a linear fit of coefficient 1.9 ± 0.1, illustrating the quadratic nature of the B–MR relationship, as expected from our geometric MR treatment.
Further MR measurements. a) Magnetic field versus MR for LiNbO3 domain wall conduction, as a function of various applied voltages (and therefore as a function of current density). b) MR versus magnetic field for various orientations of the magnetic field. At 0°, the field is in the same orientation as illustrated in Figure 2, and at 90°, the field is in the x–y plane.
Recently, electrically conducting heterointerfaces between dissimilar band‐insulators (such as lanthanum aluminate and strontium titanate) have attracted considerable research interest. Charge transport has been thoroughly explored and fundamental aspects of conduction firmly established. Perhaps surprisingly, similar insights into conceptually much simpler conducting homointerfaces, such as the domain walls that separate regions of different orientations of electrical polarisation within the same ferroelectric band‐insulator, are not nearly so well‐developed. Addressing this disparity, we herein report magnetoresistance in approximately conical 180° charged domain walls, which occur in partially switched ferroelectric thin film single crystal lithium niobate. This system is ideal for such measurements: firstly, the conductivity difference between domains and domain walls is extremely and unusually large (a factor of at least 1013) and hence currents driven through the thin film, between planar top and bottom electrodes, are overwhelmingly channelled along the walls; secondly, when electrical contact is made to the top and bottom of the domain walls and a magnetic field is applied along their cone axes (perpendicular to the thin film surface), then the test geometry mirrors that of a Corbino disc, which is a textbook arrangement for geometric magnetoresistance measurement. Our data imply carriers at the domain walls with extremely high room temperature Hall mobilities of up to ∼ 3,700cm2V–1s–1. This is an unparalleled value for oxide interfaces (and for bulk oxides too) and is most comparable to mobilities in other systems typically seen at cryogenic, rather than at room, temperature. This article is protected by copyright. All rights reserved
a) ESP map of FSA zwitterion. b) UPS spectra of SnO2 and FSASnO2 films. c) Schematic illustration of the energy diagram. d) AFM images and e) c‐AFM images of SnO2 and FSASnO2 films. f) Charge density difference of FSA on SnO2 (110) surface at an isovalue of 5 × 10–3 electrons Å⁻³. Yellow indicates electron accumulation and blue indicates electron depletion.
a) 2D GIWAXS patterns of the control and FSAPbI2 films. b) The intensity azimuthal graph of (001) diffraction of PbI2 prepared under both conditions. c) 2D GIWAXS patterns of the control and FSAFAPbI3 perovskite films. d) The intensity azimuthal graph of (100) diffraction of FAPbI3 perovskite prepared under both conditions. e) Top‐view SEM images (scale bar: 1 µm), and f) cross‐sectional SEM images (scale bar: 500 nm) of control and FSAFAPbI3 perovskite films.
a) ToF‐SIMS depth profile and b) the corresponding ion distribution of PbI3–, S–, and SnO2– in FSAFAPbI3 perovskite film. c) UV–vis absorption and steady‐state PL spectra, and d) TRPL spectra of control and FSAFAPbI3 perovskite films. e) TPC and f) TPV curves of control and FSAFAPbI3‐based devices. g) Nyquist plots of control and FSAFAPbI3‐based devices under dark condition at 0.8 V. Inset: the fitting circuit model. h) Dark J–V measurement of control and FSAFAPbI3‐based electron‐only devices. Inset: schematic illustration of the device structure.
a) Defect formation energy of PbI antisite before and after FSA adsorption on FAPbI3 (001) surface. b) Top view of theoretical models of FSA on FAPbI3 (001) surface without and with Pb–I antisite defect. c) Charge density difference of FSA on FAPbI3 (001) surface without and with PbI antisite defect at an isovalue of 5 × 10–3 electrons Å⁻³. Yellow indicates electron accumulation and blue indicates electron depletion.
a) Schematic structure of the FSAFAPbI3‐based device and FSA‐induced dipole moment at the SnO2/perovskite interface. b) J–V curves of target and control FAPbI3 PSCs under both reverse and forward scan directions. c) Statistical distribution of the PCEs of the target and control FAPbI3 PSCs. d) EQE spectra and the corresponding integrated JSC of the target and control FAPbI3 PSCs. e) J–V curves of target FAPbI3 PSCs with a 1 cm² active area. f) EQEEL and current density of FSAFAPbI3 PSCs under a bias voltage from 0 to 1.7 V. Inset: photograph of the luminescence of FSAFAPbI3 PSCs under bias voltage. g) MPP tracking measured with the target and control FAPbI3 PSCs under full solar illumination (AM 1.5G, 100 mW cm⁻² in N2 condition at 50–60 °C). Inset: photographs of the target and control FAPbI3 devices taken from the fluorine‐doped tin oxide (FTO) side after the MPP tracking.
Black phase formamidinium lead iodide (FAPbI3) with narrow band gap and high thermal stability has emerged as the most promising candidate for highly efficient and stable perovskite photovoltaics. In order to overcome the intrinsic difficulty of black phase crystallization and to eliminate the PbI2 residue, most sequential deposition methods of FAPbI3‐based perovskite would introduce external ions like methylammonium (MA+), cesium (Cs+), and bromide (Br–) ions to the perovskite structure of the light absorbing layer. Here we introduce a zwitterion‐functionalized SnO2 as the electron transport layer (ETL) to induce the crystallization of high quality black phase FAPbI3 on such SnO2 substrate. The SnO2 ETL treated with the zwitterion, formamidine sulfinic acid (FSA), can help rearrange the stack direction, orientation and distribution of residual PbI2 in perovskite layer, which reduces the side effect of the residual PbI2 to the perovskite structure. Besides, the FSA functionalization also modifies SnO2 ETL to suppress the deep‐level defects at the perovskite/SnO2 interface. As a result, the FSA‐FAPbI3 based perovskite solar cells (PSCs) exhibit an excellent power conversion efficiency up to 24.1% with 1000 h long‐term operational stability, which is among the highest values for FAPbI3 PSCs fabricated from sequential deposition. Our findings provide a new interface engineering strategy on the sequential fabrication of black phase FAPbI3 PSCs with improved optoelectronic performance. This article is protected by copyright. All rights reserved
Optical control of the domain patterns in BFO films. a) The OOP phase image of the as‐grown domain patterns, inset is the IP phase image showing the [100] directional polarization component. b) The OOP phase image of the domain structures during light on; the inset shows the corresponding IP phase image. c) The measured size changes of yellow‐colored domains depending on different light states. The red circles and black rectangles suggest the fractions of downward domains under laser on and laser off, respectively. d,e) The magnified cyan rectangles in (a) and (b), respectively. The green and blue dashed lines in (d) and (e) denote the 109° DWs and 180° DWs, respectively. The orange and cyan dashed lines in (e) denote the 71° DWs and charged 109° DWs, respectively. f) The comparison between OOP phase image (left) and IP image (right). The IP phase image displays the [010] directional polarization component. g,h) Schematic illustrations of the domain structures for the as‐grown state and photoinduced state, respectively. Two neighboring domain bundles in (g) and (h) are separated to display the polarization distribution clearly. i) Optically induced polarization switching from pristine ferroelectric variant r4⁺ to the resultant r2⁻.
Mechanisms of the selective polarization switching induced by light illumination. a,b) Local PFM phase and amplitude hysteresis loops of BFO films under different light states, respectively. c,d) Diagrams of the domains with upward and downward polarization, respectively. Eint, PS and Elas indicate the built‐in electric field, spontaneous polarization and photoinduced electric field, respectively. Note that the direction of Eint is determined by the vector sum of the electric fields induced by interfacial potential barriers between films and substrates, depolarization fields and the surface screening fields. e,f) Schematics of the domain patterns before and after ferroelastic switching. Blue, red, green and cyan lines denote the neutral 109°, 180°, 71° and charged 109° domain walls, respectively. g,h) Schematics showing the switching path of polarization variants selected by the substrate monoclinic distortion.
The switching and back‐switching of polarization induced by the coupling between electric field and light illumination. a) Writing domains poled with +30 V voltage over 25% region and −30 V voltage over 75% region. b,c) The OOP phase images of the domain patterns after electrical writing with the light off and then light on, respectively. d) Writing domains poled with +30 V voltage over 50% region and −30 V voltage over 50% region. e,f) The OOP phase images of the domain patterns after electrical writing with the light off and light on, respectively. g) Writing domains poled with +30 V voltage over 75% region and −30 V voltage over 25% region. h,i) The OOP phase images of the domain patterns after electrical writing with the light off and light on, respectively. j) The experimental downward domain evolution as a function of the electrical writing areas with +30 V voltages. Two cases for electrical writing and then light illumination are compared. k) The downward domain evolution for the regular domain patterns. Inset displays the schematic illustration of the regular domain patterns. The red circles and black rectangles in (j) and (k) suggest the fractions of downward domains under laser on and laser off, respectively.
Schematic illustrations of the controllable ferroelastic switching in BFO films under the combined stimulation of electric field and light. a) The initial regular domain patterns. b) The domain structures under light illumination. c) The domain structures poled with negative DC voltages (about −30 V). d) The domain structures obtained by external stimuli through two pathways, the one under light on/off cycle, another under electrical writing and then the light illumination. The dashed rectangles highlight the different polarization states for one logic unit consisting of two neighboring domains. e) The OOP polarization switching pathways stimulated by light and negative voltages. The higher boxes in (b) and (c) aim to highlighting the switched domains.
The magneto–electric–optical coupling in BFCO films. a) XRD reciprocal space map of the (002) diffraction for the BFCO film. b) IP magnetization‐magnetic field (M–H) curve of the BFCO film measured at 300 K. c) OOP M–H curve of the BFCO film measured at 300 K. d) Topography of BFCO film grown on KTO (001) substrates. e,f) The OOP PFM phase images of the BFCO film displaying the ferroelectric domain structures at the initial state and under light illumination, respectively. g) The MFM phase images of the BFCO film displaying the as‐grown magnetic domain structures of BFCO films. h,i) The MFM phase images of the magnetic domains after electrical writing with the light off and then light on, respectively.
Manipulating ferroic orders and realizing their coupling in multiferroics at room temperature are promising for designing future multifunctional devices. Single external stimulation has been extensively proved to demonstrate the ability of ferroelastic switching in multiferroic oxides, which is crucial to bridge the ferroelectricity and magnetism. However, it is still challenging to directly realize multi‐field driven magnetoelectric coupling in multiferroic oxides as potential multifunctional electrical devices. Here we show novel magneto‐electric‐optical coupling in multiferroic BiFeO3‐based thin films at room temperature mediated by deterministic ferroelastic switching using piezoresponse/magnetic force microscopy and aberration‐corrected transmission electron microscopy. Reversible photoinduced ferroelastic switching exhibiting magnetoelectric responses is confirmed in BiFeO3‐based films, which works at flexible strain states. This work directly demonstrates the room temperature magneto‐electric‐optical coupling in multiferroic films, which provides a framework for designing potential multi‐field driven magnetoelectric devices such as energy conservation memories. This article is protected by copyright. All rights reserved
a,b) Energy‐level diagrams and related natural transition orbitals (NTOs) for the singlet and triplet excited states of ν‐DABNA (a) and ν‐DABNA‐CN‐Me (b). Transition energies for S1 and Tn (n = 1, 2) and the spin–orbit coupling (SOC) matrix elements were calculated at the ADC(2)/def2TZVP//B3LYP/TZP and B3LYP/TZP//B3LYP/TZP levels of theory, respectively.
Photophysical properties of ν‐DABNA‐CN‐Me. a) Absorption (black) and fluorescence (blue) spectra at 300 K in toluene (1.0 × 10⁻⁵ m). b) Fluorescence spectra at 77 K with (red) and without (green) a delay time of 25 ms in PMMA (1 wt%‐doped film). c,d) Transient PL decay curve for prompt (c) and delayed (d)emission at 300 K in toluene (1.0 × 10⁻⁵ m). The red curve represents the single exponential fitting data (background = (c) 24, (d) 1). τF/τTADF: lifetimes of prompt and delayed components, respectively.
OLED performance. a) Device structure, ionization potentials (Ip), and electron affinities (Ea) (in eV) for each component. The Ip and Ea of ν‐DABNA‐CN‐Me were estimated from those of ν‐DABNA and their HOMO/LUMO energy levels. b) Normalized EL spectra. Inset: Photograph of the device in operation. c) CIE (x,y) coordinates. d) Current density–voltage–luminance characteristics. e) EQE–luminance characteristics. f) Current‐ and power‐efficiency–luminance characteristics.
Synthesis of ν‐DABNA‐CN‐Me. Yields were determined by ¹H NMR analysis.
Multiple resonance (MR) effect‐induced thermally activated delayed fluorescence (TADF) materials have garnered significant attention because they can achieve both high color purity and high external quantum efficiency (EQE). However, the reported green‐emitting MR‐TADF materials exhibit broader emission compared to those of blue‐emitting ones and suffer from severe efficiency roll‐off due to insufficient rate constants of reverse intersystem crossing process (kRISC). Herein, we report a pure green MR‐TADF material (ν‐DABNA‐CN‐Me) with high kRISC of 105 s–1. The key to success is introduction of cyano groups into a blue‐emitting MR‐TADF material (ν‐DABNA), which causes remarkable bathochromic shift without a loss of color purity. The organic light‐emitting diode employing it as an emitter exhibits green emission at 504 nm with a small full‐width at half‐maximum of 23 nm, corresponding to the Commission Internationale d’Éclairage coordinates of (0.13, 0.65). The device achieves a high maximum EQE of 31.9% and successfully suppresses the efficiency roll‐off at a high luminance. This article is protected by copyright. All rights reserved
Organic solar cells (OSCs) have experienced rapid progress with the innovation of near‐infrared (NIR)‐absorbing small‐molecular acceptors (SMAs), while the unique electronic properties of the SMAs raise new challenges in relation to cathode engineering for effective electron collection. To address this issue, we synthesized two fluorinated perylene‐diimides (PDIs), PDINN‐F and PDINN‐2F by a simple fluorination method, for the application as cathode interlayer (CIL) materials. The two bay‐fluorinated PDI based CILs possess lower the lowest unoccupied molecular orbital energy level of ca. −4.0 eV, which improves the energy level alignment at the NIR‐SMAs (such as BTP‐eC9)/cathode interface for a favorable electron extraction efficiency. The mono‐fluorinated PDINN‐F shows higher electron mobility and better improved interfacial compatibility. The PDINN‐F based OSCs with PM6: BTP‐eC9 as active layer exhibit an enhanced fill factor and larger short‐circuit current density, leading to a high power conversion efficiency (PCE) exceeding 18%. The devices with PDINN‐F CIL retained more than 80% of its initial PCE after operating at the maximum power point under continuous illumination for 750 hours. This work prescribes a facile, cost‐effective and scalable method for the preparation of stable, high‐performance fluorinated CILs, and instilling promise for the NIR‐SMAs‐based OSCs moving forward. This article is protected by copyright. All rights reserved
Journal metrics
6 days
Submission to first decision
Acceptance rate
$5,100 / £3,850 / €4,450
32.086 (2021)
Journal Impact Factor™
47.7 (2021)
Top-cited authors
Zhong Lin Wang
  • Georgia Institute of Technology
Jianhui Hou
  • Chinese Academy of Sciences
Hui-Ming Cheng
  • Shenyang National Laboratory for Materials Sciences, Institute of Metal Research, Chinese Academy of Sciences
Christoph J. Brabec
  • Friedrich-Alexander-University of Erlangen-Nürnberg
Yury Gogotsi
  • Drexel University