Employing solid state phase transformation, the Bi2Te3/In2Te3 interface density was tuned by modifying the driving force for In2Te3 precipitation. The structure-property relationship in this strongly anisotropic material is characterized thoroughly and systematically for the first time. Unexpectedly, with increasing Bi2Te3/In2Te3 interface density, an increase in electrical conductivity and a decrease in the absolute Seebeck coefficient were found. This is likely to be due to electron accumulation layers at the Bi2Te3/In2Te3 interfaces and the interplay of bipolar transport in Bi2Te3. Significantly improved thermoelectric properties of Bi2Te3-In2Te3 composites as compared to the single phase (Bi1−xInx)2Te3 solid solution are obtained.Layered (Bi Research on thermoelectrics has witnessed a renewal of interest with the drastic improvement of thermoelectric properties owing to nanoscaled structures, e.g. in thin film superlattice/quantum well multilayered structures1246782Te3 formed by the decomposition of Pb2Sb6Te11 is controlled by the temperature and time of the decomposition process2Te3-In2Te3 samples, controlling the crystal orientation of the parent solid solution and the lamellar structure that is formed by solid state precipitation. The lamellar spacing is adjusted in a wide range by varying the supersaturation of the parent solid solution. Thoroughly characterized microstructural parameters are correlated with the thermoelectric performance for the first time. Suppression of bipolar transport and enhanced interface density contributed to an improved performance of the material.Several traditional metallurgical methods have shown their potential to achieve nanoscaled structures and to generate interfaces that are beneficial for the thermoelectric properties121314151617181920121314161718202Te3  solid solutions with the {001} plane being parallel to the growth axis were obtained by a seeding zone melting technique2Te3-In2Te3 phase diagram according to previous work in the literature17Oriented 2Te3. Note that homogenization after seeding zone melting is not necessary, as a perfectly homogeneous crystal is generated. This reduces possible artifacts and the misinterpretation of measured values significantly. Considering the temperature dependent solubility of In in Bi2Te3, as illustrated by the red line in 2Te3 matrix in all samples, but a different Bi2Te3/In2Te3 interface density. After annealing for 6 days, the sample was water quenched. Annealing for a longer time of up to 12 days was also performed, but no further change in microstructure and thermoelectric properties was observed. This confirms that 6 days are long enough to reach the equilibrium. The microstructure consists of micro-/nanoscaled Bi2Te3/In2Te3 (“BTIT”) lamellae of different spacing, depending on the initial In concentration. For simplification, all the BTIT composites are termed as BTIT-c with the initial In concentration c = 4, 6, 7.5, in at%In.Samples with In concentrations of 4, 6 and 7.5 at% were vacuum sealed in quartz and isothermally annealed for the precipitation of In2O3 powder suspension, and finally polishing with 50 nm colloidal silica. The microstructure was observed using a scanning electron microscope (SEM) equipped with a backscattered electron (BSE) detector and energy dispersive X-ray (EDX) spectrometry. Thickness and volume fraction of the phases in the samples were analyzed using the image processing software Image J. X-ray diffraction (XRD) was used to characterize the constituent phase(s) of the sample. Thin sections of the specimens were examined in the transmission electron microscope . TEM samples were prepared by means of Focused Ion Beam (FIB) using an in-situ lift-out technique.The samples were prepared metallographically by grinding with a series of SiC papers up to a grit size of 4000, by polishing with 3 μm and 1 μm Al2 and ∅6 × 1 mm2 were cut along and perpendicular to the growth direction, respectively. The electrical conductivity (σ) and Seebeck coefficient (S) were measured using the thermoelectric measurement system ZEM-3 in the temperature range from room temperature to 300 °C. The equilibrium phase diagramktot) was calculated from the values of thermal diffusivity (α), density (ρ) and specific heat (CP) by the relationship k = αρCP, where α was measured using a laser flash system (Netzsch LFA-457), ρ was measured by the Archimedes method, and CP was determined by a combination of the Dulong-Petit law and a differential scanning calorimeter (Netzsch STA-449FA). The measured values fall into the same range and are also consistent with the calculated values by the Dulong-Petit law. For each sample state, 3 samples prepared by the same experimental procedure were characterized with respect to their thermoelectric properties. All samples exhibited similar or identical temperature dependent Seebeck coefficients and thermal diffusivities. A maximum variation of ±12% of the measured electrical conductivity was observed, and average values were finally used in the present work. The variation of the measured heat capacities and thermal conductivities was ±5%. The electronic contribution to the thermal conductivity (ke) is determined from the electrical conductivity utilizing the Wiedemann-Franz law, ke = LσT. From the total thermal conductivity ktot, the electronic contribution (ke) is subtracted, yielding kLB, the lattice thermal conductivity kL plus further contributions, particularly the bipolar contribution kB.Measurement of thermoelectric properties was performed for all samples. Plates of 4 × 4 × 10 mm2Te3 grains grow in a preferential orientation perpendicular to the [001] direction. In2Te3 peaks also show the anisotropy, considering that all the visible peaks are indexed (h k l) with k = l.// is nearly two orders of magnitude higher than σ⊥ for the whole temperature range  in this temperature range due to the increasing number of thermally generated holes, considering the narrow band gap of (or below) 130 meVS|⊥ is only a little smaller than |S|// and exhibits similar temperature dependence. The bright phase in 2Te3 phase, whose In concentration was measured as ~3 at%, and the black phase represents the In2Te3 phase with negligible Bi concentration. More details on the composition distribution across the two phases will be discussed in detail below.Since the samples produced by zone melting and precipitation exhibit pronounced microstructural anisotropy, also anisotropic thermoelectric properties are expected. re range . This anktot// is higher than ktot⊥, reflecting also the anisotropy of the electrical conductivity. However, after subtracting the electronic contribution (ke = LσT) from the total thermal conductivity, the remaining thermal conductivity kLB// becomes smaller than kLB⊥. The value of the Lorenz number is not only dependent on charged carrier concentration and temperature, but also can be influenced by quantum well effects, which e.g. occurs in superlattice/quantum well thin film materials302Te3 phase but also from the lamellar aligned microstructure, make the determination of the charge carrier concentration a complex issue. In the literature, Lorenz numbers in the range of 1.2…1.4 × 10−82Te3-based superlattice−82Te33334−8kLB. However, this lead to non-physical negative kLB// values of BTIT-7.5, most likely due to the complex band structure or the complex scattering mechanism in BTIT. According to published work on Bi2Te3T and thus decrease with temperature. Whereas this may approximately hold for kLB//, it is definitively not the case for kLB⊥, indicating a bipolar contribution to the thermal conductivity (see discussion below). The figure of merit, ⊥, as shown in The temperature dependence of the thermal conductivities in both directions is shown in When discussing the thermoelectric properties of BTIT, one has to keep in mind that the material is a two-phase composite material see . Effecti3738404242S) as calculated by a mixture rule follows from the Seebeck coefficients of the two phases 1 and 2 by one of the relationshipsThe total Seebeck coefficient (σ) should be a consequence of the conductivity of the individual phases followingand the total electrical conductivity  of BTIT-7.5 divert drastically from the calculated values. The measured S// values are close to the values calculated with the GEMT model . All models suggest a decreasing σ// with increasing volume fraction of In2Te3. However, the large difference between the calculated σ and σ// suggest that the Bi2Te3/In2Te3 interfaces influence the thermoelectric properties in a more complex manner.The calculation results for the properties of the BTIT composites at room temperature are plotted in irection . Due to 2Te3/In2Te3 interfaces on the transport behavior of BTIT, the interface density in BTIT was varied in a controlled way. Since the thermoelectric properties of Bi2Te3 are very sensitive to the composition, the Bi2Te3/In2Te3 interface density was modified by changing the supersaturation of the parent 2Te3 phase while maintaining the same annealing temperature, i.e. the same composition of each constituent phase. 2Te3  after precipitation annealing at 400 °C. For avoiding projection effects in the lamellar spacing measurements, microstructural analysis was performed in cross sections of the samples. The grey phase in 2Te3 phase with an identical In concentration of ~3 at% for all samples, and the black phase represents the In2Te3 phase with negligible Bi content. The composition distribution in each phase was determined in the TEM, as shown in 2Te3/In2Te3 interfaces is observed. The concentration of Te in the 2Te3 phase is a little higher than 60 at%, while the Te concentration in the In2Te3 phase is somewhat lower than 60 at%. Inhomogeneity and lattice misfit at heterophase interfaces have been demonstrated to cause transport processes across the interface and hence change the thermoelectric performance. A high resolution TEM image of a Bi2Te3/In2Te3 interface is also shown in 2Te3/In2Te3 interface is sharp and coherent. As can be clearly seen from 2Te3 increases and the lamellar spacing decreases with increasing In concentration. Simultaneously, the length of the In2Te3 plates increases. For all samples, the thickness of the In2Te3 plates is similar, in the order of ~100 to 200 nm. For BTIT-6, the lamellar spacing ranges from ~500 nm to ~3.5 μm, with the majority of lamellar spacings being in the size class ~1.2 μm. For BTIT-7.5, the lamellar spacing is smaller, ranging from ~200 nm to ~2 μm, and most of the lamellae have a spacing below 1 μm. For verifying that 6 days are enough for reaching thermodynamic equilibrium at 400 °C, we compared the measured lamellar spacing values with the diffusion length, D is the diffusion coefficient and t is the involved time scale. D is of the order of 10−12 m2 s−1, as found in previous work on the diffusion of metallic elements in Bi2Te347482Te3 precipitates for all BTIT samples. This in turn confirms that the equilibrium state at 400 °C is reached. With the In concentration increasing from 4 to 7.5 at%, the volume fraction of In2Te3 after annealing was measured as 1.5 ± 0.5%, 7.3 ± 1.5% and 13.5 ± 1.5%, respectively. These values are consistent with those calculated by the lever rule based on the solubility of In in Bi2Te3 (3at%In at 400 °C) and negligible solubility of Bi in In2Te3. In the present work, it can be seen from 2Te3 volume fraction mainly comes from the decrease of lamellar spacing rather than the increase in thickness of In2Te3.To get deeper insight into the effect of the BiS| decreases, and the electrical conductivity σ increases, indicating mixed conduction or bipolar transport behaviorS|, but the lowest σ of all samples. With increasing volume fraction of In2Te3 and decreasing lamellar spacing, |S| decreases while σ increases. Since the conductivity of In2Te3 is much smaller than that of Bi2Te3, the increase of σ with increasing In2Te3 volume fraction is not straightforward to understand.To get deeper insight into the physical transport mechanisms, the thermoelectric properties of the solid solution sample (BT-3 at%In) and the composite samples  were characterized in the temperature range from room temperature to 300 °C, see H via the equation RH = A/ne, where A is related to the anisotropy factor of the effective carrier mass and the Fermi energy levels2Te3 and n-type Bi2Te3, (i.e. the coexistence of both holes and electrons) make the accurate determination of A not straightforward. Due to a lack of detailed information, an approximation, RH = 1/ne is used in our present work. We get qualitative information concerning the increase of the electrical conductivity with increasing BT/IT interface density, as listed in 2Te3 and the BT/IT interfaces to the carrier concentration is negligible. This is partly due to the very low room temperature hole concentration of In2Te3, 7.8 × 10 9cm−3, which is nearly 10 orders of magnitude lower than that of BT-3In. The increasing σ of BTIT with increasing In2Te3 volume fraction mainly comes from the increasing carrier mobility.For qualitatively analyzing the possible reason for the increasing σ, Hall tests at room temperature were performed on all the samples. The measured Hall coefficients of BTIT at room temperature are given in ktot than BT-3, and BTIT-7.5 exhibits a comparable ktot as BT-3, as shown in 2Te3 fraction. The remaining thermal conductivity after subtracting the electronic contribution contains the lattice thermal conductivity and the bipolar contribution. The pure lattice thermal conductivity should decrease with increasing temperature due to phonon-phonon interactionLB decreases with increasing In2Te3 fraction, and for BTIT-7.5 the overall temperature dependence is inverted from increasing to decreasing 2Te3/p-In2Te3 interface is proposed in the present work, see 2Te3 and the minimum work function for Bi2Te3 are used. For p-type In2Te3, an electron affinity of 3.85 eV and a band gap of ∼1.12 eV have been reported2Te3 in BTIT in the present work. In view of the p-type behavior, we assume a Fermi level close to the valence band. Then the work function (WF) for In2Te3 ranges from 4.41 to 4.97 eV. As to n-type Bi2Te3 with 3 at% In, so far the band structure is not known and cannot be determined on the basis of the results in our present work. The recently reported band gap value, 0.08 eV, for Bi2Te3 with 4 at% In2Te3. Considering the n-type transport behavior of 2Te3, a Fermi level at the same position of the conduction band is used. Even though the Fermi level of In2Te3 is higher than that of 2Te3 at the 2Te3/In2Te3 interface, the electrons tend to flow from In2Te3 to 2Te3 to balance the Fermi level at the contact, which in turns yields band bending. 2Te3 then can be expected to act as electron acceptor, and the formation of an electron accumulation layer close to the interface to In2Te3 can be safely assumed.In the present work, the enhanced electrical conductivity . Such a metal-like surface state of Bi2Te3 may also be a possible reason for the increase of electrical conductivity with increasing BT/IT interface density. Together with the metal-like behavior at the BT/IT interfaces, a metal-like surface state of Bi2Te3 as a prominent example for a topological insulator592Te3 was reported to exhibit a high room temperature charge carrier mobility, particularly a hole mobility of 1820 cm2V−1 s−1 and electron mobility of 2890 cm2V−1 s−1. Previous work on Cu2Se-CuAgSeBand alignment at hetero-interfaces is regarded as one of the main reasons for enhanced electrical conductivity in two-phase nanocomposites, for example SiGe-Si522Te3 and In2Te3 of 3/SrTiO32Tentative Density Functional Theory (DFT) calculations with VASP .Publisher's note: Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.


This structure attains 1.4% of reflectance power losses in the operation spectral range of GaAs solar cells (440–870 nm), a 12.5% relative reduction of reflection power losses in comparison with a standard bilayer. The work presented here considers a fabrication process based on laser interference lithography and dry etching, which are compatible with solar cell devices. The experimental results are consistent with scattering matrix simulations of the fabricated structures. In a broader spectral range (400–1800 nm), the simulation estimates that the nanostructure also significantly outperforms the standard bilayer coating (3.1% vs. 4.5% reflection losses), a result of interest for multijunction tandem solar cells.The moth-eye structure has been proposed several times as an antireflective coating to replace the standard optical thin films. Here, we experimentally demonstrate the feasibility of a dielectric moth-eye structure as an antireflective coating for high-index substrates, like GaAs. The fabricated photonic crystal has Si Optical losses (reflection and parasitic absorption) in solar cells are typically minimized by using thin film dielectric multilayer coatings with high transmittance1345678910111213141516171819679111213143456789101112131415161718192)22Not every nano-lithography technique is suitable for low cost fabrication of large area devices such as solar cells. We have chosen laser interference lithography (LIL) for its capability to uniformly pattern large areas (>1 m3N4) on a high index substrate, GaAs. The design was globally optimized using 3D simulations in a previous work3N4 nano-cones in a square lattice with the following optimal dimensions: period a = 344 nm, height H = 512 nm, and radius R = 163 nm. Between the nanostructure and the substrate there are two index matching thin films, one made of Si3N4 (d = 49 nm), like the nanostructure, and one of TiO2 (d0 = 51 nm). These dimensions were used as a target for fabrication.In this work, we have fabricated a moth-Eye nanostructure made with a transparent dielectric (Si2 deposition is done using atomic layer deposition (ALD) and the Si3N4, using plasma enhanced deposition (PECVD). Both techniques are extensively used in electronics and photovoltaics252627282931323334Our structure is fabricated on a GaAs substrate and optimized for III–V tandem solar cells. These cells generally operate in optical concentration conditions, allowing to reduce the semiconductor area by up to three orders of magnitude. Still, the fabrication techniques used in this paper are fully compatible with large area solar cell production. The TiO2 using ALD. Subsequently 616 nm of Si3N4 were deposited using PECVD. Both layers were deposited at a relatively low temperature (200 °C) to ensure process compatibility with a range of photovoltaic technologies. The samples were covered with a diluted positive photoresist (S1805 + PGMEA 1:1) by spin coating at 5000 rpm for 60 seconds and baked at 115 °C for one minute before exposure. The obtained photoresist thickness was 150 nm. We did not use a bottom antireflective resist, as the sample itself is a better ARC.The initial step for the fabrication was the deposition of 51 nm of TiOμm wide lines with a period of 500 μm was defined by exposing the photoresist using conventional proximity mask UV lithography. This pattern was used to measure the height after the etching step using a profilometer. The sample was then exposed using LIL with a Lloyd mirror configuration and a 12 mW, 405 nm semiconductor diode laser on a temperature controlled mount. The light from the laser was spatially filtered using a lens (f = 4 mm) and a pin hole (5 μm) at the focal plane. A distance of 35 cm separated the pin hole from the sample and the aluminum mirror. The exposure intensity over the sample was controlled using a GaP photodiode mounted below the sample mount.An auxiliary pattern consisting of an array of 50 To obtain a square lattice (2D) the pattern must be exposed twice in two orthogonal directions. The pattern inside the unit cell will be similar to a square or a circle depending one the exposure dose3N4 layer using RIE with a 200 W, 460 V, 10 mTorr plasma with a gas flow of 25 sccm of N2 and 25 sccm of CHF33N4 and 15–20 nm/min for the photoresist. In order to avoid excessive thermal damage to the photoresist, the attack was pulsed following a periodic sequence of 30 seconds of plasma, 2 minutes of vacuum, 1 minute of N2 flux and 1 minute of vacuum. To complete the dry etching the loop was repeated 20 times, namely 10 minutes of plasma. The final height of the nanostructure was measured using a profilometer, obtaining 390 ± 10 nm. 2 plasma during 30 seconds to remove possible photoresist residues.The pattern was transferred to the Siμm diameter spot using a 20X magnification objective lens (NA = 0.4) at the lower vertical arm of the beam splitter. The signal reflected from the sample was collected by the objective and sent through the beam splitter to a lens focused on a silicon photodiode on the upper arm of the beam splitter. For measurements at wavelengths above 700 nm, a dichroic filter was used to block short wavelength light due to the second diffraction order from the monochromator. For our experimental set up we choose optical elements optimized for visible and near infrared transmission. The bounds of our experimental spectra (440 nm to 870 nm) are also limited by the position of the GaAs absorption edge and the spectral distribution of our illumination source.The reflectance was measured using the light from a halogen lamp dispersed by a 0.3 m focal length monochromator. Temporal fluctuations of the excitation beam were compensated with a reference silicon photodiode. The beam was modulated at 477 Hz with a chopper to allow for lock-in signal detection. A 200 micron core size multi-mode optical fiber was used to allow for mechanical decoupling of the monochromator and the rest of the set up. The collimated excitation beam was aligned with a Glan-Thompson polarizer at one the horizontal arms of a 45 degree beam splitter, and then focused to a 30 μm radius), as the fabricated nanostructure is not perfectly homogeneous, see a, was extracted from The reflectivity measurements were compared with the simulations results obtained using the scattering matrix method (SMM)2, d0, is 5 nm bigger than the fitted one. This small difference is most likely a result of the GaAs native oxide not being included in our simulations. An important question might arise in view of the SEM images presented in et al. indicated in ref. 45The experimental and simulated reflectance data are shown in IAMG(λ) the spectral irradiance, and R(λ) the reflectance.being λ0 = 440 nm and λ1 = 870 nm), the nanostructures lead to a 12.5% relative reduction of reflection power losses in comparison with a standard bilayer (1.4% vs. 1.6% absolute power loss). Extending the integration range to λ0 = 400 nm and λ1 = 1800 nm, our simulations predict a reflection power loss of 3.1%, while the bilayer leads to a 4.5% power loss. Thus, a 31% relative reduction of power losses is expected. As seen in In the visible and near infrared range of the spectrum (2 (nano-cones ARC) and ZnS (bilayer). The higher absorption in the bilayer results from the higher extinction coefficient of ZnS vs. TiO2. In fact, the measured extinction coefficient of our TiO2 layers .Publisher's note: Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.


Chlamydomonas reinhardtii cells under different conditions. Firstly, the impact of stress conditions (nitrogen starvation) on the cell wall structure is assessed. Secondly is described the effect of solvents, in the context of lipid extraction, on the microalga's cell, particularly its lipid droplets, in the perspective of understanding the mechanisms behind solvent extraction of lipids. Furthermore, the role of the cell wall as a barrier to the solvent extraction action is highlighted.Data are related to Confocal Laser Scanning Microscopy (CLSM) observations of lipid-enriched  Specifications tableValue of the data•The presented data provides information on the role of the cell wall as a barrier to molecules extraction; this information is of prime importance for lipid extraction from microalgae•The data will be helpful to relate the microalga's cell wall structure and the ability of a given solvent to extract lipids•The data gives significant information that might be crucial to highlight mechanisms of solvent extraction action, especially the role of the cell wall in the extraction efficiency•The data presented here might lead to further investigations on microalgae cell wall during stress conditions and solvent extraction•The data provides information to the scientific community on the effect of coupled solvents on the microalgae cells structure in the frame of the extraction of lipids from microalgae1Chlamydomonas reinhardtii are reported. Chlamydomonas reinhardtii, during stress conditions, is shown for the first time in Three experiments using confocal microscopy imaging on 7 days-stressed 22.1−2·s−1, 100 rpm agitation, a temperature of 24 °C ± 1 °C and 0.004% CO2.Lipid accumulation was induced by culturing the cells under the following conditions: nitrogen depleted TAP medium, light intensity of 150 μmol·m2.2In order to visualize the cell wall and lipid droplets using confocal laser scanning microscope  two fluorescent dyes were used: Concanavalin A conjugated with tetramethylrhodamine  for the cell wall and Bodipy 505/515 (green fluorescence) for the lipids. Chlorophyll was detected by red autofluorescence.The stock solution of ConA  was prepared by dissolving it in 0.1 M sodium bicarbonate (pH ≈ 8.3) to a final concentration of 2 mg/ml. Fluorescence intensity was analyzed after 20 min of incubation time.Bodipy 505/515 (green fluorescence) was prepared by dissolving it in DMSO to a concentration of 100 mg/L. The final concentration of bodipy in the cell sample was 1.5 µg/ml. Fluorescence intensity was analyzed after 10 min of incubation time.2.32.3.1Chlamydomonas reinhardtii cells were mixed with ethyl acetate in a total volume of 100 µL during 10 min .For this experiment, 2.3.27 cells·mL−1.The following solvent mixture was used with a ratio of 1:1: dodecane as the biocompatible solvent with dichloromethane as the highly polar solvent. The solvent/cell ratio was 1:1 and the contact time was 30 min. The cell concentration of the sample was 1 × 10


The transient reduction in adenosine triphosphate (ATP) consumption and cardiac contractility may be initially beneficial; however, sustained pH changes can be maladaptive, leading to myocardial damage and electrical arrhythmias.To demonstrate the feasibility of mapping intracellular pH within the in vivo rodent heart. Alterations in cardiac acid‐base balance can lead to acute contractile depression and alterations in CaHCO3− and CO2 resonances individually while preserving signal from the injected hyperpolarized [1‐13C]pyruvate. The large flip angle pulses were placed within a three‐dimensional (3D) imaging acquisition, which exploited CA‐mediated label exchange between HCO3− and CO2. Images at 4.5 × 4.5 × 5 mm3 resolution were obtained in the in vivo rodent heart. The technique was evaluated in healthy rodents scanned at baseline and during high cardiac workload induced by dobutamine infusion.Spectrally selective radiofrequency (RF) pulses were used to excite the The intracellular pH was measured to be 7.15 ± 0.04 at baseline, and decreased to 6.90 ± 0.06 following 15 min of continuous β‐adrenergic stimulation.13C]pyruvate. The new method is anticipated to enable assessment of stress‐inducible ischemia and potential ventricular arrythmogenic substrates within the ischemic heart. Magn Reson Med 77:1810–1817, 2017. © 2016 The Authors Magnetic Resonance in Medicine published by Wiley Periodicals, Inc. on behalf of International Society for Magnetic Resonance in Medicine. This is an open access article under the terms of the Creative Commons Attribution License, which permits use, distribution and reproduction in any medium, provided the original work is properly cited.Volumetric maps of intracellular pH can be obtained following an injection of hyperpolarized [1‐ Specifically, 13C‐MRS of CO2 and HCO3− concentrations within the heart would enable a ratiometric calculation of pH via the Henderson‐Hasselbalch formula. Unfortunately, this is infeasible using conventional  13C magnetic resonance, because of limited 13C natural abundance and the inherently low sensitivity of the MR experiment.An attractive option for noninvasive pH13C‐labeled substrates relative to their thermal equilibrium polarization in excess of 10000‐fold 13C]pyruvate, for example, provides key biochemical information regarding tricarboxylic acid (TCA) cycle flux as well as anaerobic metabolism via observation of [1‐13C]lactate 13C‐bicarbonate has previously been used to allow the noninvasive in vivo mapping of extracellular pH by exploiting CA‐mediated exchange within the blood and extracellular space 13C‐bicarbonate 13C‐bicarbonate have also been investigated, using base‐catalyzed hydrolysis in conjunction with biocompatible 13C‐enriched carbonates to reduce toxicity concerns from the presence of Cs+13C]pyruvate by pyruvate dehydrogenase (PDH) involves mitochondrial decarboxylation as an initial step before entry into the TCA cycle. This results in the production of intracellular hyperpolarized 13CO2, which is vented from the mitochondrial matrix into the cytoplasm, where a cytosolic CA isoform catalyzes the rapid exchange with H13CO3−. The ratio of these two signals can be used for determination of intracellular pH 13CO2 signal  is often at the limit of detection, which has up until now restricted this technique to whole‐heart assessment 13C‐based intracellular pH measurements to identify focal regions of myocardial ischemia, it will be necessary to develop optimized data acquisition methods that increase the measured 13CO2 signal.Recent developments in hyperpolarized magnetic resonance using the dissolution dynamic nuclear polarization (DNP) method provide transient signal increases in i mapping in the heart, using hyperpolarized 13C MRI. To this aim, we have developed an optimized data acquisition in which spectrally selective radiofrequency (RF) pulses are used to excite the HCO3− and CO2 resonances individually while preserving signal from [1‐13C]pyruvate. A variable timing scheme is used to exploit CA‐mediated label exchange between HCO3− and CO2, which enables the use of large flip angles for a 3D imaging acquisition. Images at 4.5 × 4.5 × 5 mm3 resolution are obtained in the in vivo rodent heart. The technique is evaluated in healthy rodents scanned at baseline and during high cardiac workload induced by dobutamine infusion. The new method is anticipated to enable new cardiovascular imaging examinations in which the severity of stress‐inducible ischemia can be evaluated.The aim of this study was to investigate the feasibility of performing in vivo pHAll experiments were performed on an Agilent 7 Tesla (T) MRI system . All animal investigations conformed to Home Office Guidance on the Operation of the Animals (Scientific Procedures) Act (HMSO) of 1986, to institutional guidelines, and were approved by the University of Oxford Animal Ethics Review Committee.13C]pyruvate. When required, a separate tail vein catheter was placed in the contralateral tail vein for delivery of dobutamine during the high cardiac workload experiments, described subsequently. Animals were then placed in a home‐built animal handling system Male Wistar rats  were scanned in this study. Two animals were used in initial pilot scans to determine the feasibility of the spectrally selective excitation scheme used for subsequent imaging studies. A further three animals were then scanned in two separate scan sessions consisting of rest and stress conditions to demonstrate the feasibility of the newly developed 3D imaging sequence, described subsequently. Anesthesia was induced at 2.5–3% isoflurane in oxygen and nitrous oxide . Anesthesia was maintained by means of 2% isoflurane delivered to, and scavenged from, a nose cone during the experiment. A tail vein catheter was placed for intravenous injection of hyperpolarized [1‐13C] pyruvate was administered 15 min after the start of infusion through a separate tail vein catheter. Dobutamine infusion was maintained throughout the 13C scan. The heart rate immediately prior to hyperpolarized [1‐13C]pyruvate injection was 380 ± 20 bpm at rest and 450 ± 10 bpm during stress .Scans were performed in a group of three subjects  in separate experiments under either rest or stress conditions. Cardiac workload was increased by continuous dobutamine infusion  into the tail vein using a small animal infusion pump . The volume of the injection line was filled with dobutamine, and hyperpolarized [1‐13C]pyruvic acid  was hyperpolarized in a prototype DNP hyperpolarizer at 93.979 GHz and 100 mW for 60 min 13C]pyruvate. A total of 2 mL of hyperpolarized [1‐13C]pyruvate was injected over 20 s via the tail vein.pyruvate (170 ppm) unexcited , 10−5 stopband (6.5 ppm from the center of the passband)). An open‐source RF pulse design toolbox was used to generate these pulses (http://rsl.stanford.edu/research/software.html). The RF pulse for 20° HCO3− excitation is shown in Figure 13C enriched urea . The experimental profiles show good agreement with Bloch simulation, revealing a suppression level of at least 10−3, with the remaining discrepancies likely caused by RF amplifier nonlinearities Using the Shinnar Le‐Roux (SLR) algorithm HCO3− and CO2 . The sequence alternated between excitation of all metabolites , of HCO3− only , or of CO2 only .An ECG‐gated pulse‐and‐acquire sequence was used to demonstrate the signal‐to‐noise ratio (SNR) gain from the large flip angle excitation and to determine the optimal timing window for subsequent imaging experiments of myocardial 13C]pyruvate using the different excitation pulses. These spectra show signals from a nonselective 10° hard pulse excitation in blue, as well as signals from spectrally selective excitation . The acquired spectra are localized to the heart as a result of coil placement. Figure HCO3− and 3.1‐fold for CO2. Furthermore, despite the large flip angles (70°) delivered to the CO2 resonance, rapid exchange with the HCO3− pool and new production from unaffected pyruvate regenerate this peak for subsequent measurements.Figure HCO3− and CO2 in the heart. A stack‐of‐spirals trajectory was used for spatial encoding . The eight phase encodes were acquired in a linear order across k‐space. Sequential CO2 and HCO3− volumes were acquired by switching the excitation pulse used in each block. The TR was varied between the two resonances to allow for regeneration of CO2 from the larger HCO3− resonance . The time required to acquire the two metabolic volumes was 6 s.Figure HCO3− and CO2 signal. High SNR spectra can be obtained for at least 40 s after the start of the scan.Figure sinθ. Linear nearest‐neighbor interpolation using the image magnitude was used to fill in missing HCO3− and CO2 images in the corresponding time series. pHi was estimated using the Henderson‐Hasselbalch equation (pH=pKa+ log10HCO3‐/CO2) with pKa = 6.15. pHi maps were masked by including only voxels that reached four standard deviations above the mean magnitude HCO3− noise signal in the final analysis.The spiral trajectory was predicted using a premeasured gradient impulse response function P < 0.05 level.A paired, two‐tailed Student's t‐test was used to compare pH values between rest and dobutamine stress conditions. Statistical significance was considered at the HCO3− and CO2 using the 3D pHi mapping pulse sequence in a midventricular slice. This set of images directly demonstrates the localization of HCO3− and CO2 signals to the myocardium, with no detectable signal within the blood pool. The images with dashed borders were linearly interpolated from neighboring time points. The initial missing HCO3− time point was linearly extrapolated from the first two frames of the HCO3− time series. The acquisition time for the central kz spiral readout was used as the time point for each image, and this varied between frames resulting from slight variations in heart rate over the course of the scan.Figure HCO3− and CO2 in the rodent heart, summed over the first 20 s of data acquisition at both rest and during high cardiac workload. The surface coil placement on the chest wall next to the heart and high cardiac PDH flux localize HCO3− and CO2 signals to the heart, demonstrating that any aliasing of HCO3− or CO2 from outside of the FOV is negligible. Spatial pHi maps using the ratio between HCO3− and CO2 images 20 s following the start of injection show pHi = 7.2 ± 0.2 at rest (mean ± SD) across the heart in this subject. In a separate experiment, injection of hyperpolarized [1‐13C]pyruvate during a continuous dobutamine infusion revealed a decrease in pHi to 6.8 ± 0.2 in this subject.Figure i maps (20 s after the start of injection) in multiple subjects. The HCO3− and CO2 images are localized to the myocardial tissue, as shown by the anatomical segmented GRE images. The pH maps demonstrate a uniform myocardial pHi of 7.15 ± 0.04  across the whole heart, at rest, consistent with global pHi measurements in rats and pigs i decreased to 6.90 ± 0.06 . Figure i at rest and during dobutamine stress.Figure HCO3−/CO2 buffering system to obtain a spatial three‐dimensional map of the myocardial intracellular pH. The sensitivity of the method is demonstrated by detection of a decrease in pHi following 15 min of high cardiac workload.The proper functioning of cells relies on tight control of intracellular pH. Despite the vital importance of this physiological parameter, rapid noninvasive mapping of intracellular pH within the in vivo heart is impossible with existing techniques. Here, we build upon the previous use of hyperpolarized substrates 2+ handling, and contractile function. Under certain conditions, a decline in cardiac pHi can trigger electrical arrhythmia i could allow assessment of stress‐inducible ischemia and potential ventricular arrythmogenic substrates within the ischemic heart.It is anticipated that this development will enable new experimental paradigms in which acute, focal alterations in pH can be detected in vivo. Disruption to the acid‐base balance within cells affects enzyme activity, ion channel behavior, Ca13CO2 and H13CO3− may not equilibrate on the time scale of the hyperpolarized 13C experiment. The effect of this incomplete equilibration would result in an artefactual low pHi measurement. Previous work has shown agreement between 13C and 31P‐MRS measurements in isolated, perfused hearts subjected to ischemia and subsequent reperfusion The correct use of the Henderson‐Hasselbalch equation for intracellular pH measurement requires that two necessary conditions are satisfied 2 is vented into a cytoplasmic domain with elevated CA activity 2 efflux from the myocardium into the extracellular space is likely low on the experimental timescale, as the ratio between HCO3− and CO2 reaches a plateau within 5 s in vivo HCO3− and CO2 are confined to the myocardium, which is dominated by the intracellular space. The close agreement of the hyperpolarized 13C pH measurement with the 31P measurement in the perfused heart indicates minimal contamination with the extracellular space 2 venting from the myocardium into the blood is likely low, as no HCO3− or CO2 signals were visualized within the blood pool here. The absence of any HCO3− and CO2 signal within the blood pool is consistent with dilution into a much larger volume, the fact that mature erythrocytes are anucleate and lack mitochondria, and association of 13CO2 with hemoglobin, which would shorten the T1 and rapidly remove the hyperpolarized signal. It is interesting to contrast the rapid steady state reached in rodents to measurements made in pigs, in which the measured pHi appears to rise over a 20 s period, suggesting a difference between species in the rate at which 13CO2 leaves the cytosolic compartment during the duration of the experiment The second condition is that the acid‐base pair are detected from the same cellular compartment. In rodents, mitochondrial CA activity is low, and mitochondrial CO1‐adrenergic agonist, is a positive inotrope used clinically to increase cardiac workload as a pharmacologic stress testing agent 13C]pyruvate, despite the high intracellular buffering capacity and normal perfusion of the myocardium. Similarly, increased cardiac workload reduces intracellular pH in isolated cardiomyocytes during high‐frequency electrical stimulation Dobutamine, a βHCO3− and CO2 resonances to preserve the limited reservoir of hyperpolarized [1‐13C]pyruvate magnetization HCO3− and CO2 signals, the measured pH is independent of the concentration of pyruvate or its downstream metabolites. The main requirement is that the PDH flux is high enough to enable the detection of CO2 and HCO3− within the cell. One limitation of this approach is that the arrival of pyruvate into the heart is not monitored. Although pyruvate concentration is not required for pH determination, imaging the arrival of hyperpolarized pyruvate at the heart would allow a simultaneous measurement of cardiac PDH flux, which would be useful to assess tissue viability HCO3− and CO2 images, which would be useful in a heterogeneous subject group in which the arterial input function is expected to be variable. Alternatively, co‐polarization of metabolically inert 13C urea alongside [1‐13C]pyruvate could be used to obtain the arrival time without disturbing the hyperpolarized pyruvate magnetization The optimized acquisition in this study uses spectrally selective imaging of the i determination is the relatively low SNR of the CO2 signal specifically, which is 10% of the HCO3− signal at physiological pH levels. This is consistent with an error propagation analysis of the Henderson‐Hasselbalch equation, which leads to σpH2=1/SNRbic2+1/SNRCO22, where σX denotes the standard deviation of X, and SNRX=X/σX denotes the signal‐to‐noise ratio. The SNR of the CO2 is typically ∼5 in the in vivo experiments shown here, which is consistent with the single subject standard deviation reported in Figure 2; this angle was chosen to improve SNR while avoiding excitation by flip angles above 90° due to B1+ inhomogeneity, which would decrease the CO2 signal. This is a reasonable trade‐off, as the transverse magnetization created by 70° excitations is 94% that of a 90° excitation.The limiting factor for accurate pH2 SNR may further be improved by shortening the duration of the CO2 excitation. Specifically, the 10 ms pulse used for excitation of HCO3− (170 ppm) was designed to avoid excitation of [1‐13C]pyruvate (160 ppm). CO2, which is positioned at 124.5 ppm, is three times further away in chemical shift, and a shortened pulse (∼3.5 ms) would reduce the effective echo time and reduce T2*‐related dephasing across the volume. Although this could potentially improve SNR, the difference in echo time would need to be compensated by T2* mapping within the heart. This could be done using multiecho 1H images, and converting the T2* values to 13C via multiplication by γ1H/γ13C ≈ 4.The CO31P‐MRS, suffers from contamination of the pH‐sensitive Pi resonance with 2,3‐diphosphoglycerate  found in the ventricular blood. It is possible to perform a saturation transfer experiment to remove the contribution from 2,3‐DPG, but this method remains limited to a whole‐heart measurement with long scan time due to limited SNR 13C and 31P‐based measurements agree in isolated, perfused rat hearts in which blood is not present 13C measurements described here.Currently, the validation of noninvasive methods for intracellular pH mapping within the myocardium is challenging. The gold standard for noninvasive intracellular pH measurement, 13C substrates to study human heart disease is rapidly approaching the clinic. The first application of hyperpolarized magnetic resonance in humans at the University of California, San Francisco, demonstrated the potential of the technique to stage prostate cancer and assess the response to treatment 13C]pyruvate, has a direct path for clinical translation.The use of hyperpolarized 13C]pyruvate, mitochondrial decarboxylation results in cytosolic CO2 and HCO3−. Spectrally selective excitation is combined with a rapid 3D imaging readout, and the rapid regeneration of CO2 due to carbonic anhydrase‐mediated exchange is used to obtain whole‐heart intracellular pH maps in vivo. The intracellular pH was measured to be 7.15 ± 0.04 at baseline, and decreased to 6.90 ± 0.06 following 15 min of continuous β‐adrenergic stimulation. The new method is anticipated to enable assessment of stress‐inducible ischemia and potential ventricular arrhythmogenic substrates within the ischemic heart.An optimized sequence for mapping intracellular pH within the rodent heart in under a minute is proposed. Following an injection of hyperpolarized [1‐


Metal‐intercalated iron selenides are a class of superconductors that have received much attention but are less understood in comparison with their FeAs‐based counterparts. Here, the controversial issues such as Fe vacancy, the real phase responsible for superconductivity, and lattice stability have been addressed based on first‐principles calculations. New insights into the distinct features in terms of carrier doping have been revealed. The origin of this issue is largely due to the phase separation that inevitably occurs in these systems at 500–578 K,2Fe4Se5) and the SC phase.xFe2−ySe2 film grown on SrTiO3 (001) substrate, an Fe‐vacancy‐free phase KFe2Se2 with √2 × √5 charge ordering was observed in the SC region.Tc) of 44 K in the K‐Fe‐Se system was also reported.One of the most controversial issues is whether the SC phase is Fe vacancy free or not, i.e., the FeSe layers in ATc from 30 to 46 K, which are often mixed in a single sample. For instance, apart from a dominated 30 K phase, a 44 K phase of trace amount sometimes can be observed in a KxFe2−ySe2 sample obtained by a high‐temperature route.x = 0.6–0.7), but its real compositions are unknown. The results in ref.0.3Fe2Se2 with a Tc = 44 K and K0.6Fe2Se2 with a Tc = 30 K. Superconductivity with Tc of 30 K and 43 K was also observed in KxFe2−ySe2 samples with √2 × √2 × 1 superstructure due to K vacancy ordering, but the phases for different Tc have not been specified.Tc of 44 K.Another puzzling issue is the existence of several SC phases with different xFe2−ySe2 as an example using first‐principles calculations. Two competing factors are found to dominate the formation of the phases and the structural evolution. One is due to the energy increase caused by the accumulation of the negative charge in FeSe layers. The other is due to the Coulomb attraction between the K ion layers and the negatively charged FeSe layers. Then we show that the intercalated K content at 0.25 ≤ x ≤ 0.6 can stabilize the body centered tetragonal (bct) lattice and Fe is favored to fully occupy its site. At x > 0.6, the structure is stabilized by creating Fe vacancies in FeSe layers. At low intercalated K level, x < 0.25, the structure will collapse due to the lattice instability. Then a schematic phase diagram is constructed accordingly whereby we speculate the SC phases are KxFe2−ySe2 with 0.25 ≤ x ≤ 0.6 and y = 0. Our finding sheds some light on understanding this distinct SC family and should also be applicable to other metal‐intercalated iron selenides besides KxFe2−ySe2.So far, there has been no consensus on the answers to the above issues. Moreover, the underlying mechanism of the lattice stability and vacancy is also poorly understood, in particular, in terms of the origin for multifarious structures. To address these issues, here, we first study the energetic change and structural evolution as a function of intercalated metal content by taking KTo begin with, we study the formation energy for the process of K intercalation. The process can be expressed as a chemical reactionEI = EKxFe2Se2 −EK0Fe2Se2 − xEK, where EKxFe2Se2 and EK0Fe2Se2 are the total energies of KxFe2Se2 and an assumed K0Fe2Se2 with a similar bct structure but without any K ion, and EK is the energy of elemental K. The variation of ΔEI versus K content is shown in FigureThe formation energy per unit cell can be described as ΔEI in terms of energy from the structural constituent units and their interactions. For the K intercalation process, as shown in Figure xFe2Se2 with a bct structure. Hence, the following contributions to the total formation energy are considered: (1)EK ion layers = E(K ion layers)x+ − xEKThe formation of K ion layers: Δ(2)Ee‐ doping =E(FeSe layers)x− −EFeSe layersThe electron doping in FeSe layers: Δ(3)EFeSe layer deformation = EFeSe layers −EK0Fe2Se2The deformation of FeSe layers: Δ(4)EC =EKxFe2Se2 −E(FeSe layers)x− − E(K ion layers)x+The Coulomb attraction between K ion layers and FeSe layers: ΔThen we break down ΔE(K ion layers)x+ is the energy of K ion layers, E(FeSe layers)x− the energy of charged FeSe layers, and EFeSe layers the energy of FeSe layers in KxFe2Se2. Therefore, the total formation energy can be written as the sum of these four energy changes: ΔEI = ΔEK ion layers + ΔEe‐doping + ΔEFeSe layer deformation + ΔEC.And eI when they are negatively charged with more and more electrons. This is easily understood since charging a neutral FeSe sheet will incur additional energy like charging a capacitor. Inversely, with the increasing K content, more charges will come into effect in the Coulomb attraction between K ion layers and the charged FeSe layers and thus greatly enhance |ΔEC|, which will sufficiently offset the increase of energy induced by electron doping into the FeSe layer. In contrast with these two contributions, the other two energy contributions due to the formation of K ion layers and the deformation of FeSe layers are much smaller. Therefore, we conclude that the former two contributions dominate the energy change during the K intercalation.The variations of these four energy contributions as a function of the K content are calculated and the results are shown in Figure EFe vacancy  by removal of an Fe atom from the unit lattice of KxFe2Se2 and its dependence on the K content. It is found that ΔEFe vacancy only fluctuates above and below zero when x ≤ 0.6 as shown in Figurex > 0.6, ΔEFe vacancy rapidly drops, revealing that Fe vacancies are favored at high levels of K intercalation. This is easily understood since K has a smaller electronegativity than Fe, when the doped electrons due to K are over a limit that Fe—Se bond can accommodate, the surplus electrons will prefer to transfer into Fe2+. In this way, Fe is repelled out of the lattice in the form of an element. The final compound with Fe vacancy is more energetically a favored state. The limit is around at x = 0.6. The K2Fe4Se5 phase contains 20% Fe vacancy corresponding to x = 0.8 is a manifestation of this argument. This tendency of appearance of Fe vacancy with K content is supported by the experimental results.To understand the tendency for appearance of Fe vacancy in the K‐intercalated iron selenides, we go further to consider the energy change Δ0.2Fe2Se2, where negative frequencies appear, meaning the structure is unstable. Partial DOPS further indicate that it is mainly because of the considerable amount of highly unstable Se atoms, which are unbonded due to the absence of K nearby. The phonon density, however, exhibits no imaginary mode for all other compounds with x ≥ 0.25. As examples, Figure 0.25Fe2Se2 and K0.5Fe2Se2. The above results demonstrate that the phases of KxFe2Se2 free of Fe vacancy are both energetically and dynamically favored in the K intercalation range of 0.25 ≤x ≤ 0.6.Besides the above considerations on the formation energy, the influence of lattice dynamics on the structural stability should also be accounted for. This is, in particular, true for the compounds with low levels of intercalated K, which can induce lattice instability. Figure x ≤ 0.6. Both electronic and size effects of K will be taken into account. In order to better understand the role of electron doping in the structural change, we charge K0Fe2Se2 with various electron concentrations, which allows us to explore the effect of electron doping while excluding the size influence of K. As shown in Figurea, the Fe—Se bond, and the Se—Fe—Se angle clearly indicate that FeSe layers are stretched along the ab plane upon electron doping alone.  For a more realistic case, other factors such as the lattice mismatch and the Coulomb attraction must be included. Figure a, stretches FeSe layers, and reduces the lattice constant c. For instance, the lattice constant a of K0.25Fe2Se2 and K0.5Fe2Se2 expands from 3.65 Å to 3.69 Å and c shrinks from 14.42 Å to 13.92 Å. The predicted lattice constant c of K0.25Fe2Se2, 14.42 Å agrees with 14.28(4) Å of K0.3Fe2Se2 having Tc of 44 K.a and c for K0.5Fe2Se2, 3.69 Å and 13.92 Å are consistent with the periods observed using scanning tunneling microscope along [110] and [001], 5.5 Å (√2a) and 14.1 Å,2Si2 structures, such as KFe2As22As2.Now we are focused on the trend of structural change upon K intercalation in the range of 0.25 ≤0.25Fe2Se2 and K0.5Fe2Se2. Both structures can survive for at least 1 picosecond using 2√2 × 2√2 × 1 super cells at temperatures up to 500 K, indicating that they are stable at these temperatures.0.25Fe2Se2 are larger than those of K0.5Fe2Se2, suggesting that K0.25Fe2Se2 is less stable. The variations of free energy with temperature for the two phases are also calculated and shown in Figure 0.5Fe2Se2 at temperatures above 268 K, consistent with the result obtained by MD. Despite their variations in stability, both of them have the electronic structures similar to KFe2Se2 , which suggests that they should have similar properties to KFe2Se2. Considering stoichiometry, formation energy, stability, and electronic structures, K0.25Fe2Se2 and K0.5Fe2Se2  are proposed to be responsible for the observed superconductivity at 44 and 30 K,Furthermore, we explore the temperature‐dependent stability of structures in the K region of interest. To this end, we perform the molecular dynamics (MD) simulations on KxFe2–ySe2 system in Figurex > 0.6, phases with Fe vacancies tend to exist, agreeing well with the observed antiferromagnetic K2Fe4Se5 phase. As for the low level K‐intercalated compounds (x < 0.25), there has been no report on the synthesis of free‐standing KxFe2Se2. We note that the two identified SC phases K0.3Fe2Se2 and K0.6Fe2Se2 lie in the region of 0.25 ≤ x ≤ 0.6.Tc is 44 K for the former phase and 30 K for the latter one, suggesting that Tc is dependent on the K content or the doped electron concentrations. Although experimentally observed K contents in SC phases by far are discrete, recent reports about carrier concentration tuning of Tc from 30 to above 40 K in FeSe thin flakesTc with carrier concentration can be continuous, similar to the cases in other high‐temperature superconductors. Further improvement of Tc can be expected considering the experimental progresses achieved in single‐layer FeSe films.3, NH2, or C2H4(NH2)2 along with alkali metal can stabilize the structures.Tc of bulk iron‐selenide‐based superconductors by controlling carrier doping while stabilizing the structures by intercalations or effect of substrates. At the moment, tremendous efforts are needed toward this goal. Because of the calculation limit and accuracy, we do not consider the slight off‐stoichiometric cases, say, the Fe vacancy concentration is less than 3 at%. It should be pointed out that such slight off‐stoichiometry is tolerated in KxFe2Se2 just like many other materials which may be the reason that superconductivity was observed in previous reports of refs. Based on the results presented above, we schematize a phase diagram for the KxFe2–ySe2 as a function of intercalated K content using first‐principles calculations. Two factors dominating the formation of the phases and the structural evolution are confirmed. One is due to the accumulation of negative charge in FeSe layers, the other is due to Coulomb attraction between K ion layers and negatively charged FeSe layers. The structural evolution of this series of phases is summarized: at 0.25 ≤ x ≤ 0.6, the bct lattice is stable and Fe is favored to full occupy its site; at x > 0.6, FeSe layers tend to exclude Fe atoms and create Fe vacancies; and at x < 0.25, the structure will collapse for the dynamic instability. A schematic phase diagram is constructed accordingly and the possible route to further improve Tc is suggested. The phases responding to the observed superconductivity are proposed to be K0.25Fe2Se2 and K0.5Fe2Se2 in terms of stoichiometry, formation energy, stability, and electronic structures. Though based on the study of KxFe2−ySe2, our results should be meaningful to understand the SC and its related phases in metal‐intercalated iron selenides and other similar SC systems.In conclusion, we carefully investigate the energetic change and structural evolution of KAs a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


The results show that the slight change of interlayer geometry can significantly modulate the charge transfer time from 100 fs to 1 ps scale. Detailed analysis further reveals that the transfer rate in MoS2/WS2 bilayers is governed by the electronic coupling between specific interlayer states, rather than the interlayer distances, and follows a universal dependence on the state‐coupling strength. The results establish the interlayer stacking as an effective freedom to control ultrafast charge transfer dynamics in 2D heterostructures and facilitate their future applications in optoelectronics and light harvesting.Light‐induced interlayer ultrafast charge transfer in 2D heterostructures provides a new platform for optoelectronic and photovoltaic applications. The charge separation process is generally hypothesized to be dependent on the interlayer stackings and interactions, however, the quantitative characteristic and detailed mechanism remain elusive. Here, a systematical study on the interlayer charge transfer in model MoS The MX2 materials share similar crystalline structures and symmetries, but possess distinct electronic properties in bandgaps, photoabsorption, and spin–orbit coupling strength.2 are of particular interests because many of them form type II heterojunctions,Since the discovery of graphene and the rise of MoS2 heterostructures is of central importance in their photoresponse, which determines both the speed and efficiency of the charge separation.2/WS2 bilayer takes place in an ultrafast time scale.2 heterostructures, the quantitative characteristic and detailed mechanism for the interlayer stacking/interaction dependence still remain elusive.The interlayer charge transfer in MX2/WS2 bilayer system with variable stacking configurations. Our results demonstrate that the interlayer twisting, translation, or spacing can significantly modulate interlayer charge transfer time from 100 to 1000 fs scale, a huge modulation that has not been realized before. Further analysis reveals that the transfer rate in MoS2/WS2 bilayers is governed by the coupling between specific interlayer states  and follows a universal exponential dependence on their dipole coupling matrix values. This work establishes a firm correlation between charge transfer dynamics and interlayer stacking/interaction in 2D heterostructures on the atomic level, and thus facilitate their future applications in, for example, high speed optoelectronics and new generations of light harvesting.Here we utilize the state‐of‐the‐art time‐dependent density functional theory method (TDDFT) and demonstrate a systematical study on the interlayer charge transfer in model MoS2/WS2 as a model system to study the interlayer charge transfer process in 2D heterostructures is mainly initiated by the ultrafast experiment from Wang and co‐workers at UC Berkeley.2/WS2 bilayers have a type II band alignment: the conduction band minimum and valence band maximum at K point reside at MoS2 and WS2 layers, respectively [Figure2 has an energy bandgap smaller than WS2, we can selectively excite the MoS2 by purposely choosing the excitation light with photon energy in between the bandgap of MoS2 and WS2. After photoexcitation, holes in MoS2 will have the chance to transfer into the WS2 because its valence band maximum is lower than that of WS2; while the electrons will stay in MoS2 since its conduction band minimum is also lower than WS2  methods,2 nanowire2/WS2 bilayer is profoundly determined by the direct excitation of electron–hole pair in the individual layer since the interlayer interactions will only disturb the optical transition slightly.2/WS2 bilayer , they are referred to as interlayer state. It should be noted that |−2> and |−1> states at K point are not sensitive to interlayer coupling, and they do not change much with the number of layers, while the states at Γ point is quite allergic to the interlayer interactions. We note that direct interlayer excitation has a negligible strength as compared to intralayer excitation due to the minimal wavefunction overlap. Once this occurs, the interlayer exciton is formed right away upon excitation, and no charge transfer takes place in a short time <1 ps. By choosing the photon energy to be that of MoS2 bandgap, it will mainly excite the electron from |−2> to |1> state and leaves the holes in MoS2 layer. The holes will transfer into WS2 layers afterward.MoSer Figure2a, at 2/WS2 bilayer. At the beginning, the energy of hole states (orange curve) is 0.48 eV below the valence band maximum. As the energy levels evolve in real time accompanied by the ionic motion, the hole states are getting closer to the valence band maximum. At t = 150 fs the difference becomes as small as 0.13 eV and the excited hole states in MoS2 and original occupied states in WS2 get largely mixed, or the hole transfer takes place. In Figure 2 . At t = 60 fs, ≈25% of the holes have transferred into WS2, and clearly at t = 150 fs, about 50% of the holes are localized in WS2. It means that the hole transfer between MoS2 and WS2 is taking place in an ultrafast way.In the following, we look into the detailed hole transfer dynamics. We simulate the time evolution of interlayer charge transfer by simultaneously solving the time‐dependent Kohn–Sham equation and the Newtonian motion of ions .A1g mode was found to be critical to the photoexcited hole dynamics. Here by doing Fourier transformation of dipole moment along the vertical direction of the heterojunction, we find that the A1g phonon mode indeed plays an important role in the process of ultrafast hole transfer (see the Supporting Information for details). The formation of intralayer exciton prior to interlayer charge transfer and the collective motion of excitons is secondary for the dynamics of photoexcited charge, because the main driving force in such an ultrafast process is the specific state coupling as detailed below.To gain quantitative information on charge transfer dynamics, we integrate the hole (electron) density (χ) on the WS1‐2H, AB2‐2H, AB3‐2H, AA1‐3R, and AA2‐3R in MoS2/WS2 bilayers. For AB ones, the M‐X bond directions in the two layers are opposite; while for AA ones, the directions are the same. So between AA and AB ones, the bilayers rotate by an angle of π; and between AA1‐3R (AB1‐2H) and AA2‐3R (AB2‐2H), the two layers translate in‐plane by 1.1 Å. From our calculations  and hole dynamics. Thus, the hole dynamics of AB3‐2H stacking mode is presented for comparison in Figure S6 (Supporting Information). Since the charge transfer is related to the interlayer electronic coupling, naively one would expect that the shorter the interlayer distance is, the faster the interlayer charge transfer will be, or τAB1 ≈ τAB2 ≈ τAA1 ≪ τAA2.Now we turn to evaluate the dependence of charge transfer dynamics on interlayer stackings and interactions. There are five typical stacking configurations of ABons Table1 and at/τ), where τ is the charge transfer lifetime. As marked in FigureAB1 or τAA2 is around 100 fs timescale and τAB2 or τAA1 is around 1000 fs scale (or τAB1 ≈ τAA2 ≪ τAB2 ≈ τAA1). It seems that the charger transfer time has no obvious correlation to interlayer coupling strength, which is quite counter‐intuitive. The timescale is obtained by fitting to the simulated trends of ultrafast charge transfer upon photoexcitation. Direct simulation with trajectories longer than picoseconds is not only computationally unaffordable, but also blocked by the divergence of orbital propagation and break‐down of mean field approach. We found that simulation in a relatively short timescale already gives a very reasonable fitting result of the charge transfer timescale .In our results, we fit the time evolution data with an exponential equation χ = a + b*exp (−M) to evaluate the coupling strength between the two states asZ^ is the position operator along the vertical direction normal to the MX2 plane. Quite interestingly, we found that MAB1 > MAA2 ≫ MAB2 > MAA1, which is positively correlated to 1/τ. Detailed analysis further revealed that 1/τ is exponentially dependent on M, or 1/τ ∝ Me . Again, all the data seat around the same 1/τ versus Me curve. So, the interlayer coupling between the |−2> and |−1> states is the universal factor that determines the charge transfer dynamics in MoS2/WS2 bilayers. To clarify why the specific coupling strength between |−2> and |−1> states at K point is crucial, we have checked all relative states (occupied or unoccupied states) and found only M= <−2|Z^|−1>, is allergic to the photoexcited hole dynamics . The couplings between other pairs of states are much smaller . The dependence of charge transfer rate on the interlayer‐state‐coupling M justifies the importance of M as a measure for interface ultrafast dynamics.To further check whether this exponential dependence is universal in determining the interlayer charge transfer dynamics, we artificially tune the interlayer spacing from 6.0 to 7.3 Å in ABM. We use AB1‐2H and AA1‐3R bilayers as an example for illustration, which are both stable with the interlayer spacing of 6.3 Å. We directly draw the spatial distribution of |−2> .Now we try to give out a physical picture why slight change of interlayer stacking configuration will give out dramatically different > Figure b and |−1> Figure  d in bot2/WS2 bilayers, where the interlayer stacking configurations are not well controlled.2 bilayers with defined interlayer geometry can be directly grown by chemical vapor deposition methods in several groups.2/WS2 bilayers with different stackings predicted here.In the original experiment by Hong et al., the authors used film transfer method to prepare MoS2/WS2 bilayers. Our study reveals that slight interlayer geometry modulation of twisting, translation, or spacing can tune charge transfer dynamics very effectively, resulting in a change of transfer lifetime spanning from 100 to 1000 fs timescale. A universal exponential relationship between the charge transfer rate and the coupling strength between specific interlayer electronic states is established, based on detailed analysis of stacking‐mode and layer‐spacing modulation on the charge transfer dynamic processes. Based on these findings, one could utilize physical or chemical methods to control the interlayer geometry and therefore to control the charge transfer quantum dynamics, thus facilitating future applications of 2D heterostructures in novel optoelectronic and light harvesting devices.In summary, we employ ab initio TDDFT‐MD methods to investigate the ultrafast interlayer charge transfer dynamics in MoS2 vertical heterostructures were performed using density functional theory implemented in the Vienna ab initio simulation package−2 eV Å−1 and the total energy variation is less than 10−4 eV. The Brillouin zone was sampled by a set of 25 × 25 × 1 k‐mesh with an energy cutoff of 400 eV for plane waves. The thickness of vacuum layer was set to be larger than 15 Å, so that interactions between repeated images are avoided. Formation energy (E) of the MoS2/WS2 heterostructure is defined as E=(EMoS2/WS2− EMoS2− EWS2)/N, where EMoS2/WS2, EMoS2, EWS2 are, respectively, the total energies of MoS2/WS2 bilayers, MoS2 and WS2 layers, and N is the number of atoms in the supercells.First principle calculations of MX2 and MoS2 in experiment (0.7 eV) were correctly reproduced by the PBE functional (0.5 eV).The excited‐state real‐time TDDFT simulations were carried out with the time‐dependent ab initio package TDAPThe authors declare no conflict of interest.SupplementaryClick here for additional data file.


For elucidating and controlling its topological quantum state, high-quality Cd3As2 thin films have been highly desired. Here we report the development of an elaborate growth technique of high-crystallinity and high-mobility Cd3As2 films with controlled thicknesses and the observation of quantum Hall effect dependent on the film thickness. With decreasing the film thickness to 10 nm, the quantum Hall states exhibit variations such as a change in the spin degeneracy reflecting the Dirac dispersion with a large Fermi velocity. Details of the electronic structure including subband splitting and gap opening are identified from the quantum transport depending on the confinement thickness, suggesting the presence of a two-dimensional topological insulating phase. The demonstration of quantum Hall states in our high-quality Cd3As2 films paves a road to study quantum transport and device application in topological Dirac semimetal and its derivative phases.A well known semiconductor Cd 3As2, the high-quality Cd3As2 films are still rare. Here, Uchida et al. grow high-crystallinity and high-mobility Cd3As2 thin films and observe quantum Hall states dependent on the confinement thickness.Despite many achievements in the topological semimetal Cd Among them, the topological Dirac semimetal (DSM) has attracted burgeoning attention as emergence of Dirac fermions in three-dimensional (3D) materials14. The 3D DSM state is particularly intriguing as a parent phase of exotic topological phases such as 3D topological insulator3, Weyl semimetal4, and two-dimensional (2D) topological insulator6, which are realized by symmetry breaking in DSM7. As crystalline materials of 3D DSM, Cd3As2 and Na3Bi have been theoretically suggested8, and their key electronic structures have been directly confirmed by angle-resolved photoemission and scanning tunneling spectroscopy12. A classification scheme of DSM in terms of the crystal point group symmetry has also been developed13.Topological materials, which are characterized by a non-trivial electronic band topology, have great potential for unprecedented quantum transport phenomena3As2 has been known as a stable II–V type semiconductor from the early period, its film quality has been limited due to the necessity of low-temperature growth for resolving its high volatility  and valence bands (VB) with inverted orbital character, touching to form a 3D Dirac dispersion centered at the Dirac points ±kD  and substrate (SrTiO3) significantly improves the crystallinity and the electron mobility of the films. The film triangular lattice is epitaxially grown on the substrate square lattice with aligned in-plane axes, owing to a good match of their projected lattice distances. As confirmed in the transmission electron microscopy image  reported for Cd3As2 single-crystals20.High-quality Cdage Fig. , the Cd ern Fig. , Bragg pn = 1 × 1018 cm−3) and different thicknesses . Shubnikov-de Haas (SdH) oscillations and corresponding plateau-like structures are resolved from a few teslas in longitudinal resistance Rxx and Hall resistance Ryx. As the field increases, integer quantum Hall states clearly emerge down to the quantum limit with filling factor ν = 2. Rxx is further suppressed and finally becomes zero. Simultaneously, Ryx exhibits quantized values over wide field ranges, which are expressed as 1/Ryx = −ν(e2/h) = −sn(e2/h), with the degeneracy factor s and a non-negative integer n. A swell around the quantized values confirmed for thinner films is probably an artifact typically appearing in pulsed field measurements of such high resistance samples27. Although such deformations of transport data taken in pulsed fields were corrected by calculating the effective current through the sample as detailed in Methods and Supplementary Fig. s shows a dramatic change depending on the film thickness, governing the appearance of the quantum Hall effect. It is altered from s = 2 to 1 when the thickness increases only by 2 nm from t = 14 to 16 nm. For clarifying the origin of this change, we analyze the quantum transport in detail.Figure 3As2 films, in order to extract effective masses and also quantum scattering times using the Dingle expression28. For the 12 nm film, as a typical example, the oscillation amplitude gradually decreases with elevating temperature but remains finite up to about 100 K BF to the primary oscillation frequency BF,1, the Fermi surface area AF is calculated to be AF = 3.3 × 10−3 Å−2 for the 12 nm film, for example. The dimensional change is also reflected in a clear field-angle dependence of the oscillation and magnetoresistance, as shown in Supplementary Figs. EF, measured from the Dirac points, is estimated to be EF = 116 meV by using −1 and vF = ħkF/m* = 8.9 × 105 m/s in the following reported hyperbolic dispersion with an onset energy of E0 = 50 meV and an energy difference between the conduction band bottom and the Dirac points of ECB = 35 meV12.From the Fourier transformation of the SdH oscillations, further information about the 2D Fermi surface can be extracted Fig. . By applBF,2 is detected at lower frequencies, which is ascribed to the subband splitting due to the quantum confinement. The subband electronic structure is also evaluated assuming the same Fermi velocity and onset energy.For the thicker 16 and 23 nm films, another peak 3As2 films can be comprehensively explained by considering a confinement effect on the original Dirac dispersion as schematised in Fig. k⊥[112]) and the film normal direction (k||[112]) are summarized in Fig. kF and EF determined from the above analysis of the SdH oscillations using the dispersion relationship are plotted along the in-plane momentum direction. Here the band edge positions are interpolated from previous calculations17. The band gap Eg is also almost consistent within the error bars with estimations from the g factor change (for details see Methods).The various quantum Hall states appearing in the Cd31,t is the film thickness, nc is an integer numbering the confined subband, and δ(E) is the total phase shift at the interfaces (for details see Methods). In the structural plot expressing this relationship in Fig. k||[112] (E)) and the quantization condition curves ((2πnc − δ(E))/2t) determines energy and momentum of the bottom of the subbands for each thickness. This agrees rather well with the experimental trends including the appearance of the second subband (nc = 2) above 16 nm. Reflecting the large Fermi velocity of the Dirac dispersion, the band gap sharply opens when the confinement thickness decreases below 23 nm, giving rise to the dramatic g factor change observed in the quantum Hall effect. The band character inversion, which occurs when crossing projected kD, is also confirmed between t = 12 and 23 nm, as denoted by the CB character change from blue (Cd 5s) to pink (As 4p). In the case of the thick films where the subband is located inside the projected kD, the gap energy and g factor become nearly unchanged.The quantum confinement condition along the film normal direction is given by the following formula3As2 films. As illustrated in Fig. 3As26, should be induced by the confinement as long as the system remains 2D. Surprisingly, the two dimensionality is maintained up to 23 nm, where the orbital character only of the first subband (nc = 1) is inverted, suggesting the presence of a 2D topological insulating phase at this thickness. Below this thickness, another topological phase transition to a trivial insulating phase occurs associated with the sharp change in the g factor as confirmed in the thickness-dependent quantum Hall states. Since a magnetic field destroys the 2D topological insulating state by breaking time-reversal symmetry, nonlocal transport and scanning probe microscopy measurements are highly desirable for its further investigation. Applying electric gating, heterostructure fabrication, and chemical doping to such high-quality Cd3As2 films will open possibilities for further studying quantum transport and device application by tuning Fermi level, hybridization gap, and magnetic interaction in this system.Novel topological phases derived from the 3D DSM state can be expected for the high-quality Cd3As2 has been known as a high-mobility semiconductor over half a century32, its high-quality thin film growth has been quite challenging. In the 1970s and 80s, rather thicker films than 1 μm were grown by evaporating bulk Cd3As2 with a heater36 or a high-repetition-rate laser37, as plotted in Supplementary Fig. 6, a more elaborate approach using molecular beam epitaxy (MBE) has realized the epitaxial growth of single-crystalline thin films38. Compared to bulk Cd3As2, however, their crystallinity and mobility are still limited due to the low-temperature growth. In this situation, we have developed and improved a high-temperature annealing technique with combining pulsed laser deposition (PLD), and obtained high-crystallinity and high-mobility epitaxial thin films comparable to bulk quality. The mobility reaches maximum of μ = 3 × 104 cm2/Vs for n = 1 × 1018 cm−3, while it intrinsically decreases with reducing dimensions (thickness) from three (≳80 nm) to two (≲40 nm).While Cd3As2 films and TiO2/Si3N4 capping layers were deposited using KrF excimer laser. Optimization of the capping materials and their combinations has enabled high-temperature annealing of the film, while the idea of adopting a protective layer was already tried for Cd3As2 growth15. A Cd3As2 polycrystalline target was prepared by mixing 6N5 Cd and 7N5 As shots at a ratio of 3:2, keeping the mixture at 950 °C for 48 h in a vacuum-sealed silica tube, grinding and pelletizing the compound, and then resintering it at 250 °C for 30 h. The three layers were successively deposited on (001) SrTiO3 single-crystalline substrates, at room temperature and below a base pressure of 10−7 Torr. Typical laser conditions  were , , and , for Cd3As2, TiO2, and Si3N4, respectively. The shape of the Hall bar, as shown in the inset of Fig. 3As2 and TiO2 layers, while the Si3N4 layer was then deposited on the entire substrate to cover the Hall bar edges. After annealing the sample at 600 °C in air, an (112)-oriented Cd3As2 film is formed through epitaxial crystallization.Cd3As2 (14 nm) and TiO2 (30 nm)/Si3N4 (200 nm) layers shows only the SrTiO3 substrate peaks before annealing. After annealing the sample at 600 °C in air, in stark contrast, (112)-oriented Cd3As2 film peaks become clearly evident through epitaxial crystallization. Three different periodic components from the respective layers remain observed in both the Kiessig fringes. Many combinations of other substrates  and capping materials  were also tested, but they resulted in chemical reaction to substrates, cracking of capping layers, or poor crystallinity of the films. As shown in Supplementary Fig. 2/Si3N4 capping layer conduction, which is probably due to slightly oxygen-deficient TiO2 deposited under the high vacuum, is more dominant at high temperatures above about 100 K, particularly for thinner Cd3As2 films with higher resistance. At low temperatures, on the other hand, the capping layers become highly insulating and only the Cd3As2 film conduction remains, ensuring intrinsic quantum transport measurements of the Cd3As2 films.Supplementary Fig. 3As2 single-crystals20. A ϕ scan with twelve-fold symmetry reveals that the in-plane [13As2 layer (3.88 Å) has a good match with the SrTiO3 one (3.91 Å). The domain size is estimated at about a few tens of microns from the STEM observations on various areas, which are comparable to the channel length scale (~60 μm) but much larger than the magnetic length Detailed XRD data of the obtained epitaxial film, confirming its high crystallinity and flatness, are summarized in Supplementary Fig. 3As2 single-crystalline film are displayed in Supplementary Fig. I41cd)39 or the recently corrected (I41/acd)40 structure is formed in the film, while both have the similar electronic structure.Atomic-scale images of the CdRxx and Hall resistance Ryx were measured on the 60 μm-width multi-terminal Hall bar with flowing a DC current of I = 5 μA. In this Hall bar configuration, unexpected effects on the transport such as the current jetting effect in high-mobility semimetals are avoided. Aluminum electrode wires were connected to the Hall bar edges by using an ultrasonic bonding machine and then their connections were reinforced by applying silver paste. Small deformations of transport data taken in the pulsed magnetic fields were corrected based on a simple classic model27. In general, when measuring a resistive sample in pulsed fields, a small capacitive component connected parallel to the sample is slightly charged or discharged depending on the resistance change, leading to non-negligible time variation of an effective current through the sample. With the increase in the sample resistance Rxx, the effective current ix deviates more from the original set current I, and the deformations become more serious. By numerically solving the following differential equation detailed in ref.27 with a capacitance of C = 4–9 nF, we could calculate the exact current ix and obtain data showing negligibly small hysteresis between forward and backward field sweeps, as exemplified in Supplementary Fig. Transport measurements up to 55 T were performed using a nondestructive pulsed magnet with a pulse duration of 37 ms at the International MegaGauss Science Laboratory at the Institute for Solid State Physics of the University of Tokyo. Longitudinal resistance Rxx for the 12 nm Cd3As2 film, measured from 1.4 to 50 K in the pulsed high fields. Rxx minima at the ν = 2 quantum Hall state slowly increase from zero with elevating temperature, which can be well fitted with the standard Arrhenius plot. Obtained high activation energy of Δ = 19 K is ascribed to the unusually high Fermi velocity in Cd3As2.Supplementary Fig. Rxx and Ryx were measured using a Quantum Design Physical Properties Measurement System cryostat equipped with a 9 or 14 T superconducting magnet. Supplementary Fig. s in the quantization formula can be extracted from the increment of the plateau values. Change of the degeneracy from s = 2–1 is observed for the 23 nm film at low fields, indicating that spin splitting of oscillations occurs above about 12 T. Apparent degeneracy of s = 4 observed for the 16 nm (from ν = 16–12) and 23 nm (from ν = 20–16) films is ascribable to subband crossing. Corresponding beating pattern due to the existence of another subband can be also confirmed in Rxx for the 16 and 23 nm films. In contrast to the films below 23 nm, the Hall plateaus in Ryx become much less pronounced in the 37 nm film and almost completely disappear for the 100 nm film. This suggests that the system gradually changes from two-dimensional (2D) to three-dimensional (3D) around 40 nm.At low fields, m*) and quantum scattering time (τq) following the Dingle expression28.Rxx/R0 is the oscillation amplitude normalized by the zero-field resistance and ωc = eB/m* is the cyclotron frequency. To investigate the main conduction band, the analysis is performed assuming a single band, although the oscillation amplitude at each Landau index is affected by the existence of other subbands for the 16 and 23 nm films. The effective mass is slightly decreased with decrease of the confinement thickness, probably due to the dispersion curvature change associated with the gap opening.Supplementary Fig. Rxx, measured also with a conventional superconducting magnet. When the applied magnetic field is tilted from out-of-plane (θ = 0°) to in-plane (90°) direction in the 23 nm film, the oscillation period as well as the amplitude is substantially reduced. A gradual dimensional change from a cylindrical (2D) Fermi surface to a spherical (3D) Fermi surface is observed between 23 and 100 nm, consistent with the thickness dependence of Hall plateaus in Supplementary Fig. BF,2 than the 16 nm one, eliminating the possibility of the neck orbit as a cause of BF,2. Along with this dimensional change, considerably large negative magnetoresistance probably due to so called chiral anomaly14 is also observed for the B || I configuration (θ = 90°) on the Hall bar.Supplementary Fig. ϕB can be estimated from the intercept γ − δ in the fan diagram, on the basis of the following expression of the oscillating term in ΔRxx28.BF is the SdH oscillation frequency, γ is the phase factor expressed as γ = 1/2 − ϕB/2π, and δ is the phase shift being zero in two dimensions. For the 12 and 14 nm films, the intercept is about −0.3, which corresponds to the non-trivial Berry’s phase of ϕB ~ 0.4π. In Cd3As2, the Berry’s phase extracted from the intercept has been highly scattered and controversial43. According to the recent theoretical calculation43, when the Fermi energy is located above the saddle point of the two Dirac dispersions as in the cases of the previously reported carrier densities, the non-trivial phases at ±kD are canceled out Landau-level fan diagrams magnified around the origins are shown in Supplementary Fig. Rxx and Ryx taken at 40 mK for the 12 nm film are compared to 2 K ones, showing no major difference between them. So far, no more fine structures such as fractional quantum Hall states are confirmed in the present samples even at this ultra-low temperature up to 14 T.Further low-temperature quantum transport was also measured using a dilution refrigerator. In Supplementary Fig. Eg is also estimated from the g factor change in the following Roth equation, which is derived in the second order of the k ⋅ p perturbation theory30.Δ is the spin–orbit splitting energy (0.27 eV)29 and Ep is the energy equivalent of the principal interband momentum matrix element (−0.68 eV). Eg is estimated at ≳110, ≳110, ~55, and ~30 meV for the 12, 14, 16 and 23 nm films, by using the above parameters and respective g factors . These electronic structures parameters are summarized in Table The band gap mentclass2pt{minimEB by assigning band gap values Eg and chemical potential differences Δϕ of SrTiO3, TiO2, and Cd3As2 to the following empirical formula44.The total phase shift at both interfaces was estimated as a function of the binding energy The data supporting the plots within the paper and its Supplementary Information File are available from the corresponding author upon reasonable request.Supplementary information


However, up until now, only the NpCp3 solid state structure has been studied. Here we report on the solid state structures of UCp3 and PuCp3 which are isostructural to the Np analogue. The structural models are supported by theoretical calculations and compared to their lanthanide analogues. The observed trends in changes of bond lengths might be indicator for an increased covalency in the bonding in the tris‐cyclopentadienide actinide(III) complexes (AnCp3) compared to their lanthanide homologues.The organometallic tris‐cyclopentadienide actinide(III) (AnCp The organometallic actinide chemistry with cyclopentadienyl ligands was developed in Karlsruhe and Munich by the pioneering work of E. O. Fischer, F. Baumgärtner, and B. Kanellakopulos together with P. Laubereau, then of the National Laboratories at Oak Ridge.3 were reported 50 years ago,4 complexes.3 (Ln: lanthanide) complexes not stabilized by Lewis base adduct formation have been previously studied.3 complex, the synthesis and solid‐state structure of NpCp3 has only recently been published.III complex Pu(Cp(TMS)2)3 and its reduced PuII analogueIV organometallic plutonocene derivative.The oxidation state +III is not the most stable for all actinides. Nevertheless the solvent free non‐stabilized tris‐cyclopentadienide actinide(III) complexes AnCp3 or PuCp3 complexes are still unknown. This is because even in the case of the adduct free LnCp3 complexes, high quality single crystals are not easily obtained. Indeed different forms are sometimes observed depending on the crystallization conditions.Fifty years after the first reports, the structures of the UCp3 . Comparing them to the structures of NpCp3 and related LnCp3 complexes offers the opportunity to gain a more detailed insight in the bonding. This is important for the understanding of 4f or 5f electron behaviour and differences therein.Here, we close the knowledge gap on the solid‐state structures of AnCp3 was prepared by reductive elimination of chloride from UCp3Cl with sodium amalgam in diethylether. PuCp3 was obtained from the direct reaction of PuCl3 with a slight excess of KCp. Both were purified by filtration and evaporation of the solvent followed by extraction with pentane or pentane/Et2O mixtures. The IR spectroscopic data reveal a fingerprint consistent with that previously reported for UCp3 and PuCp3.1H NMR spectra of UCp3 show one single resonance at δH=−15.60 ppm ([D8]THF) or −13.62 ppm ([D3]MeCN) for the formed adducts under these conditions, which are in agreement with literature‐known values.13C frequency resulting in an overall comparable situation as observed in the bis‐TMS substituted uranocene derivative in3 are the 4th example of a Pu organometallic complex for which a proton resonance is reported and the 2nd complex on which multi‐dimensional NMR spectroscopy was performed.6]benzene there is one resonance observed for [PuCp3(thf)] at 11.59 ppm  giving rise to a cross‐peak in the CH correlated spectrum at 81.4 ppm . This is a sign that the Cp rings are in equilibrium due to fast chemical exchange in the sample. It seems that in all Pu organometallic complexes reported up to now the chemical shifts observed for the proton as well as for the 13C resonances appear in the same range independently on the oxidation state of the metal being +II, +III, or +IV.UCp3 and PuCp3, form crystals that are isomorphic to the NpCp3 analogue.3 complexes3 complexes with comparable cell parameters, maybe containing one axis doubled but also with an identical reduced cell.1 with a≈14.15, b≈8.70, and c≈9.60 Å, which corresponds to a monoclinic reduced cell of a≈8.30, b≈9.60, and c≈8.30 Å with β ≈116.5° . The monoclinic cell has been used before to describe LaCp333 structure which has been deposited at the CCDCBy extraction single crystals are obtained suitable for X‐ray diffraction analyses  the description in the orthorhombic space group Cmc21 is best, as in the monoclinic reduced cell for the refinement a disorder must be introduced which is not the case in the orthorhombic cell. This leads for the monoclinic case in the refinement with identical crystallographic independent cell volume to nearly double the refined parameters but higher R values. As the two compounds UCp3 and PuCp3 form the same structure, only PuCp3 is depicted representatively in Figure We are now convinced that at least in the cases for the three actinide complexes AnCp5‐coordination. The Lewis acidity of the actinides causes the formation of one additional η1‐coordination to one Cp ring of a neighboured AnCp3 residue; this Cp ring is μ‐η5,η1‐coordinated  pm) showing that Cp in KCp is a better Lewis base than in NpCp3. Lewis base adduct formation like in [UCp3(thf)] or in [UCp′3(quinuclidine)] produces a similar situation with symmetrical η5‐coordination of the Cp rings with a closer interaction to the donor atom of the Lewis base involved than observed here for the μ‐η1‐coordinated C‐atom.A coordination environment of four Cp rings three establishing η in KNpCp the KCp 5‐coordination in AnCp3  is not symmetrical: one of the rings (not the one involved in the bridging mode) in all the three structures, is localized closer to the central AnIII ion than the other two (Table 2)3.The bonding of the three Cp rings in ηwo Table . This is3 are 264.4 to 272.0 and 276.9 to 291.5 pm, respectively. Accordingly the distances between metal ions and the centres of the Cp rings  from 278(2) for UCp3 over 281(2) for NpCp3 to 283(1) pm for PuCp3 − with the central Cp ring in the bridging position . The geometry optimisations reproduced the η5,η1‐coordination of the bridging Cp ring, confirming that this unique interaction belongs to the basic bonding properties of the complexes and is not enforced by the packing effects. Similarly, the competitive nature of η5,η1‐interactions are confirmed by the calculations, the results reflecting the already described changes in the M−C distances. During the geometry optimisations we observed that the system is very flexible; it exhibits a flat potential energy hyperface. Hence slight changes in force can cause significant changes in the structure in the η1‐M−C distances. Another significant clue on the bonding was the verified importance of the 4f subshell for the Ln−Cp donor–acceptor interactions, calculations using the 4f‐in‐core Ln pseudopotentials failed to reproduce the characteristic change of the η1‐Ln−Cp distances. On the other hand, the experimentally suggested gradual change in the M−C bond lengths for η5‐ and η1‐coordinated Cp rings across the 4f/5f rows were only partially reproduced by the calculations. The probable reason lies in the already mentioned flat potential energy surface and the dimeric model structure  being unable to account for long‐range cooperative or solid‐state effects.This is because the outer orbitals of the actinide ions in AnCp3 and UCp3. They indicate that covalency in AnCp3 is higher than in LnCp3 (at least for the here reported minor actinide complexes), which is in agreement with theoretical considerations.Our experimental results described here close the knowledge gap on the solid state structure of the long known complexes PuCp3(CNC6H11) to those of the corresponding lanthanides. The IR CN‐stretching vibration of the isonitrile ligand is an excellent sensor on its binding mode and forces which enables the detection of differences between the lanthanides and actinides in their complexes MCp3(CNC6H11).With this background it seems promising to compare as well the cyclohexylisonitrile adducts AnCpCCDC 570389 (PuCp3) and 1570390 (UCp3), contain the supplementary crystallographic data for this paper. These data are provided free of charge by The Cambridge Crystallographic Data Centre. For further information, please see the Supporting Information.The authors declare no conflict of interest.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.