HR) is the process whereby two DNA molecules that share high sequence similarity are able to recombine to generate hybrid DNA molecules. Throughout evolution, the ability of HR to identify highly similar DNA sequences has been adopted for numerous biological phenomena including DNA repair, meiosis, telomere maintenance, ribosomal DNA amplification and immunological diversity. Although Rad51 and Dmc1 are the key proteins that promote HR in mitotic and meiotic cells, respectively, accessory proteins that allow Rad51 and Dmc1 to effectively fulfil their functions have been identified in all examined model systems. In this Review, we discuss the roles of the highly conserved Swi5‐Sfr1 accessory complex in yeast, mice and humans, and explore similarities and differences between these species.Homologous recombination ( Moreover, the process of replicating DNA before each cell division has a propensity to generate double‐strand breaks (DSBs) in DNA, resulting in broken chromosomes. Such damage can lead to genomic instability, a hallmark of cancer. Two major DNA repair pathways protect our genomes from DSBs: homologous recombination (HR) and nonhomologous end joining (NHEJ) HR‐mediated DSB repair is dependent on the presence of an intact stretch of DNA that shares high sequence similarity with the broken DNA molecule (i.e. is homologous). In contrast, NHEJ does not involve a homologous partner and the broken ends are ligated together. NHEJ is therefore thought to be inherently more mutagenic than HR. Although first observed during the meiotic cell cycle, HR has since been shown to be involved in a variety of biological processes other than DSB repair, including the generation of immunological diversity as well as numerous aspects of chromosomal biology such as telomere length regulation and ribosomal DNA maintenance The central proteins in HR are the Rad51 and Dmc1 recombinases, which are the eukaryotic homologues of the bacterial recombinase RecA. These RecA family recombinases are defined by their ability to directly promote strand exchange in a homology‐dependent manner and differ from classical site‐specific recombinases. Although Rad51 is required for both mitotic and meiotic HR Schizosaccharomyces pombe) Swi5‐Sfr1 and the budding yeast (Saccharomyces cerevisiae) homologue Mei5‐Sae3. Furthermore, the discovery of the Swi5‐Sfr1 recombination accessory complex in mammals has provided support for the notion that proteins involved in fundamental processes such as DNA repair are highly conserved throughout evolution. In this review, we will focus on the roles of Swi5‐Sfr1 in HR. We will first describe the identification and in vivo characterisation of Swi5‐Sfr1 in yeast, before turning our attention to the biochemical characteristics of this protein complex. A mechanistic model for the role of Swi5‐Sfr1 in HR will be proposed, followed by a description of how Swi5‐Sfr1 was identified and characterised in mammals, and how it compares with its yeast counterpart. Finally, we will discuss the implications of these findings in a broader context and consider the key questions that remain to be answered.Since the last review on Swi5‐Sfr1 S. cerevisiae and S. pombe led to the characterisation of numerous genes that are required for functional mitotic and meiotic HR. These genes encode recombination accessory proteins such as Rad52 S. pombe was linked to mating‐type switching, a phenomenon that occurs in some homothallic fungi and allows the conversion of cell type through an HR‐dependent DNA rearrangement  and Rrp2 (HLTF in H. sapiens), which are both homologues of S. cerevisiae Uls1 with predicted helicase and ATPase activities rrp1− or rrp2− single mutants did not show sensitivity to DNA damage, the rad57−rrp1− and rad57−rrp2− double mutants were more sensitive to DNA damage than the rad57− single mutant. In fact, these double mutants were almost as sensitive to DNA damage as the rad57−sfr1− double mutant (and the rad51− single mutant). In contrast, the DNA damage sensitivity of the sfr1−rrp1− and sfr1−rrp2− double mutants resembled the sfr1− single mutant. Importantly, the rad51−rrp1− and rad51−rrp2− double mutants showed the same DNA damage sensitivity as the rad51− mutant, indicating that rad51− is epistatic to rrp1− and rrp2−. Taken together, these data suggested that Rrp1 and Rrp2 have a role specifically in the Swi5‐Sfr1 pathway of HR. Consistent with this model, yeast‐two hybrid data indicated that Rrp1 and Rrp2 interact not just with each other, but also with Swi5 rrp1− and rrp2− mutant phenotypes, it is highly likely that the two proteins form a functional complex in vivo to regulate Swi5‐Sfr1‐specific HR. This specialised regulation of the Swi5‐Sfr1 HR pathway by Rrp1‐Rrp2 provides further evidence for the idea that recombination accessory factors differentially influence the molecular process of HR. The role of Rrp1‐Rrp2 in HR remains to be determined, but the most recent model proposed by Dziadkowiec et al. This hypothesis was substantiated by subsequent genetic analyses involving Rrp1  experiments in vivo co‐IP assays demonstrated conclusively that Mei5 forms a complex with Dmc1 during meiosis; yeast two‐hybrid analysis indicated that this interaction required the N terminus of Mei5 Following the discovery of Swi5‐Sfr1 in in vitro three‐strand exchange assay was used to demonstrate that Rad51 from S. cerevisiae (ScRad51) facilitates pairing of circular single‐stranded DNA (ssDNA) with homologous linear dsDNA and subsequent strand exchange in an ATP‐dependent manner in vitroScRad52. Several groups independently demonstrated that ScRad52 can overcome the inhibitory effect of RPA in the three‐strand exchange reaction H. sapiens (HsRAD52) could stimulate RAD51‐mediated DNA pairing in vitro, these experiments did not involve RPA S. pombe or S. cerevisiae, stimulates HsRAD51 in vitro by overcoming the inhibitory effect of RPA HsRAD52, serves as a canonical recombination mediator in humans HsRAD52 and ScRad52, BRCA2 is unable to efficiently anneal RPA‐coated ssDNA. It is likely that BRCA2 and HsRAD52 have assumed divergent roles in mammalian cells throughout the course of evolution, with BRCA2 taking on the role of mediator and HsRAD52 taking on the role of annealer.An S. cerevisiae Rad51 paralogues Rad55 and Rad57 were purified as a complex and shown to stimulate strand exchange by overcoming the inhibitory effect of RPA in vitroS. pombe (consisting of Rlp1‐Sws1‐Rdl1), although a homologue of the Csm2 subunit has not been discovered in vitroS. pombe, as there are no reports detailing their purification.Following the characterisation of Rad52, other recombination accessory proteins were studied. The in vivo evidence suggested that Swi5‐Sfr1 is also a recombination accessory factor that directly promotes Rad51 activity  was purified and its effect on the strand exchange reaction was examined Extensive This characterisation of Rad51‐mediated strand exchange was followed by an investigation into the roles of Swi5‐Sfr1 and Rad52 in Dmc1‐mediated strand exchange SpRad52 is not required for crossover formation during meiosis in vitro and in vivo data lead to a model where Swi5‐Sfr1 takes on the role of a canonical mediator during meiosis in addition to its usual role in strand exchange activation to promote Dmc1‐dependent interhomologue recombination.What is the significance of Swi5‐Sfr1 during meiosis? Genetic observations have established that S. cerevisiae, strongly supported the notion that this complex promotes the activity of Dmc1 only. Sae3‐Mei5 was shown to overcome the inhibitory effect of RPA on Dmc1 in the D‐loop assay, which measures strand invasion et al. S. cerevisiae meiosis, they also highlight the need to determine how Swi5‐Sfr1 differentially modulates Rad51 and Dmc1 activity during meiosis in S. pombe.Results obtained from biochemical experiments with Sae3‐Mei5, the meiosis‐specific homologue of Swi5‐Sfr1 in in vitroet al. Structural studies were instigated to gain further insight into the role of Swi5‐Sfr1 in HR. Attempts to crystallise the full length Swi5‐Sfr1 complex were unsuccessful, so Swi5 in complex with a 180 residue N‐terminal deletion mutant of Sfr1 (Sfr1C) was crystallised instead How does Swi5‐Sfr1 promote Rad51 activity? As mentioned above, the presence of Swi5‐Sfr1 stabilises the Rad51 filament in an ATP‐dependent manner, preventing the displacement of Rad51 by RPA Mus musculus) in vivo and in vitro, and both proteins localised to the nucleus. Furthermore, SWI5 protein could not be detected in Sfr1−/− mouse embryonic stem (ES) cells, and SFR1 levels were diminished in Swi5−/− cells, strongly suggesting that association of SWI5 with SFR1 enhances the stability of both proteins in vivo. Consistent with an evolutionarily conserved role for SWI5‐SFR1 in HR, in vitro co‐IP experiments demonstrated that SWI5, but not SFR1, interacts with RAD51, whereas the opposite is true in S. pombein vivo HR, measured through the use of a fluorescence‐based assay, were not affected by the absence of either SWI5 or SFR1. Nonetheless, Swi5−/− and Sfr1−/− cells were more sensitive than wild‐type to chemicals that cause DNA strand breaks. In an attempt to reconcile these results, the BRC3 peptide, which is known to inhibit HR, was expressed in mouse ES cells. A 10‐fold reduction in HR was seen in wild‐type cells, but a 20‐fold reduction was seen in cells lacking SWI5 or SFR1. Taken together, these results suggest that SWI5‐SFR1 does promote HR in mice, but that its contribution to HR only becomes significant under conditions where HR is already compromised.Based on amino acid conservation, homologues of Swi5 and Sfr1 were identified in mice .Extensive H. sapiensS. pombe and mouse homologues . Interestingly, both subunits of human SWI5‐SFR1 are capable of binding to RAD51. Moreover, in contrast to what was seen with a fluorescence‐based HR assay in mouse cells, the absence of SWI5 or SFR1 in human cells led to a clear reduction in HR. Accordingly, these mutant cell lines displayed increased sensitivity to ionising radiation. Taken together, these data provide compelling evidence that the role of SWI5‐SFR1 as a recombination accessory factor has been conserved throughout evolution.Shortly after the discovery of SWI5‐SFR1 in mice, a homologous complex was identified in There are numerous noteworthy attributes of the Swi5‐Sfr1 complex that are conserved from yeast to mammals. In this section, we will discuss some of the more striking similarities and differences.S. pombe, the C‐terminal half of Sfr1 is required for the interaction with Swi5; since Swi5‐Sfr1C but not Sfr1N can stimulate Rad51, it is likely that Swi5‐Sfr1 complex formation is essential for its function S. pombe and mice, it is important to note that the N terminus of human SWI5 has an extension that does not exist in most other species. Regardless of this difference, in each case, the module required for the interaction contains a highly conserved coiled‐coil motif The nature of the Swi5–Sfr1 interaction is of interest. In S. pombe, it was shown that Swi5‐Sfr1, but not Swi5 alone, interacts with Rad51 S. pombeThe interaction between Swi5‐Sfr1 and Rad51 is thought is to be essential for Swi5‐Sfr1 function. In S. pombe, Swi5 has Sfr1‐independent roles in mating‐type switching swi5− and sfr1− mutants in S. pombe are indistinguishable with regard to DNA damage sensitivity swi5−/− and sfr1−/− mouse cells showed similar sensitivity to ionising radiation, sfr1−/− cells were more sensitive to camptothecin and etoposide, which are type I and type II topoisomerase poisons, respectively In addition to these structural differences, there are numerous physiological differences. In S. cerevisiae and S. pombe, Swi5‐Sfr1 plays an integral role in meiotic HR; this is reflected in its ability to serve as a canonical mediator for Dmc1 in vitroin vitroIn both S. pombe. However, even in this model organism, where Swi5‐Sfr1 is best characterised, the molecular role of Swi5‐Sfr1 in HR is not fully understood. Furthermore, the interplay between Swi5‐Sfr1 and other recombination accessory factors, such as the Rad51 paralogues Rad55‐Rad57, is poorly characterised at best. Do the two complexes physically interact? In what way do their stimulatory effects on Rad51/Dmc1 differ (e.g. strand invasion or strand exchange)? Related to this, can they synergistically stimulate Rad51/Dmc1? Why does Swi5‐Sfr1 have a specific regulatory mechanism mediated by Rrp1/Rrp2 S. cerevisiae, where Sae3‐Mei5 is meiosis‐specific, why is it that Rad51 does not require a Swi5‐Sfr1 homologue for mitotic HR? In addition to these questions, the possibility of post‐translational modifications regulating Swi5‐Sfr1/Sae3‐Mei5 has not been explored, although phosphorylation of both Swi5 and Sfr1 has been reported There is no doubt that Swi5‐Sfr1 is a critical regulator of HR in in vitro, remains an enigma. It is not known whether human SWI5‐SFR1 can exert any stimulatory effect on RAD51‐mediated strand exchange in vitro. Likewise, it is not known if SWI5‐SFR1 plays a role in mammalian gametogenesis. Although sequences of Swi5‐Sfr1 homologues from vertebrates and fungi have been deposited into the National Center for Biotechnology Information databases, to our knowledge, there are no known Swi5‐Sfr1 homologues in plants. If plants do not possess a Swi5‐Sfr1 homologue, it would be of interest to determine whether other recombination accessory proteins , or recombinases themselves, compensate for this. Future work in S. pombe should focus on illuminating the precise contribution of Swi5‐Sfr1 and other accessory factors to HR, which will be achieved by elucidating the various mechanisms of Rad51 stimulation. Furthermore, structural studies are required to identify the potential mechanism of RAD51 stimulation by mouse SWI5‐SFR1. Additionally, it would be of interest to determine whether mouse SWI5‐SFR1 acts synergistically with the RAD52/BRCA2 family of recombination mediators, as was reported for S. pombeIn comparison to the yeasts, relatively little is known about mammalian SWI5‐SFR1. Despite recent advances in the biochemical characterisation of mouse SWI5‐SFR1, the human complex, which is yet to be studied This manuscript was prepared by BA with guidance from YM and HI.


Study of room temperature PL was performed in both single-crystal and polycrystalline NGO (substrates and pellets) respectively. Several NGO pellets were prepared with varying Ce concentration and their room temperature PL was studied using 325 nm laser. It was found that the PL intensity shows a systematic quench with increasing Ce concentration. XPS measurements indicated that nearly 50% of Ce atoms are in the 4+ state. The PL quench was attributed to the novel concept of super hydrogenic dopant (SHD)”, where each Ce4+ ion contributes an electron which forms a super hydrogenic atom with an enhanced Bohr radius, due to the large dielectric constant of the host. Based on the critical Ce concentration for complete quenching this SHD radius was estimated to be within a range of 0.85 nm and 1.15 nm whereas the predicted theoretical value of SHD radius for NdGaO3 is ~1.01 nm.In this study we report the existence of novel ultraviolet (UV) and blue emission in rare-earth based perovskite NdGaO It has had a revolutionary impact in changing or introducing electronic135n or p type dopants. It has been also predicted that at very low concentrations (of about only 1%) these super hydrogenic orbitals will overlap with each other to create a band and this band may cause magnetic (ferro or antiferro) effects in the material123 (NGO), this novel approach can be also utilized to study other exotic materials like rare-earth oxides (mostly with R2O3 formula) which exhibit large band-gap, high dielectric constant with potential applications in electronic and photonic technologies. In this paper we report the origin of room-temperature ultraviolet (UV) and blue emission in NdGaO3 (NGO) bulk systems and the subsequent quench of PL by Ce4+ doping in the place of Nd3+ ions in NGO as a case study of super-hydrogenic dopant (SHD).In this study we report a novel electronic doping strategy via super-hydrogenic dopants in oxides which can be a very interesting approach for fundamental understanding as well as potential applications. As we know, the radius of hydrogen atom in free space is about 0.5 Å; but the whole situation changes when this atom is placed into a high dielectric constant material. The Rydberg radius of that atom becomes much larger - the free space radius multiplied by the dielectric constant of the system, these kind of super-hydrogenic atoms can be realized in wide band-gap oxides (with high dielectric constants) by doping them with a = 5.43 Å, b = 5.50 Å and c = 7.71 Å while its orthorhombically distorted pseudo cubic cell has a lattice constant of 3.87 Å. It is close to lattice constants of many perovskite materials and therefore generally used as substrates for colossal magneto-resistance thin film and other perovskite thin film growth. Due to the very low dielectric loss of NGO in the GHz frequency rangeNGO is a paramagnetic insulator, which undergoes a magnetic phase transition into an antiferromagnetic phase at ~1 K due to strong super-exchange interaction3Al5O12)163+ ion in NGO single crystals were studied by Orera et al.Rare earth based oxides, perovskites and their interfaces are of great interest in optical studies as rare earth ions can serve as emission centers for laser systems, such as Nd:YAG  were used to find the Ce dopant level in doped NGO poly-crystalline samples. Unfortunately, Ce elemental level measurements using RBS and PIXE were not possible as the atomic mass or Ce (Z = 57) and Nd (Z = 58) are comparable and thus the signals of Ce and Nd overlap with each other in both RBS and PIXE spectra. Followed by RBS-PIXE, we instead tried the Secondary Ion Mass Spectroscopy (SIMS) technique for quantifying the Cerium dopant in the doped NGO samples . samples . Five sa samples -inset, ieaks see , indicat2. Generally, PL spectra of NGO single crystals exhibit a very sharp UV emission peak at 388 nm. There are a few of emission peaks, which can overall be categorized into three bands. The first band around 360 and the second one around 390 nm are UV emission peaks while the third one around 420 nm pertains to blue emission. The characteristic energy level diagram of rare earth ions Ln3+, known as Dieke diagram only changes slightly in different crystal structures−1, respectively. The ground state of Nd3+ is 4I9/2. The 325 nm laser excites electrons in the ground state 4I9/2 to energy levels of ~30769 cm−1, which are 2L15/2, 4D7/2 and 2I13/2 levels [schematized in red in 4D3/2, 2P3/2 and 2P1/2. Based on the energy level diagram of Nd3+4D3/2 → 4I9/2; the emission around 390 nm corresponds to the optical transitions 4D3/2 → 4I11/2 and 2P3/2 → 4I9/2; and the emission around 420 nm can be ascribed to the optical transitions 4D3/2 → 4I13/2, 2P3/2 → 4I11/2 and 2P1/2 → 4I9/2. Such possible emission processes are schematized in Room temperature PL spectra were measured in the range of 330 nm to 700 nm. However, no PL was detected above 450 nm, which is consistent with the results reported2, which could be several orders of magnitude higher than the intensity of the tungsten lamp through a 0.5 m double monochromatorIt is worthwhile to point out that no photoemission was seen in NGO under the tungsten lamp excitation4D3/2 → 4I11/2 and/or 2P3/2 → 4I9/2 transition of Nd3+. We find that initially the PL peak height of the 390 nm line in pure NGO poly-crystalline sample was ≈10 k counts, which drastically drops to almost half i.e. ≈5.2 k at only 0.2% elemental doping of Ce4+ and finally the 390 nm PL peak intensity completely suppressed by only 5% elemental doping of Ce4+. 4+ replaces Nd3+, it donates an extra electron to the system. As this electron doping forms a deep level it does not contribute to the electrical conductivity of the material. Using the doping-dependent drop in the peak height shown in 22/cm3. We find that the total PL almost quenches near 5% Ce doping – so that means effective Bohr radii of 5% Ce atoms envelops almost all the Nd3+ ion centers which are clearly the source of PL in this study. This assumes that all Ce ions are in the 4+ state. To determine the valence of the Ce we refer to 20/cm3 related to the quenching, so the volume occupied by each Ce dopant electron comes out to be ~2.62 × 10−21 cm3. This value further leads to an effective Ce4+ electronic radius of ~8.5 Å if we assume all PL quench happened precisely at 5% Ce doping level.Using Avogadro number, molecular weight and density of NGO we can calculate the number of NGO molecules per unit volume that comes out to be ~3.46 × 10The simple calculation procedure described above can be represented by the mathematical equation shown below:wherex – Dopant concentration required to completely quench the PL signal.N – The Avogadro number;3);M – The molecular weight of the material  to 1.15 nm (for 2% doping). Using the free space Rydberg radius value of 0.5 Å and considering the bulk dielectric constant of NGO of ~20.2 we estimate a SHD radius of 1.01 nm which lies exactly in the middle of our experimental range of 0.85–1.15 nm.Now, if we assume that the precise doping level  lies somewhere between 2 to 5% of Ce3+ ion. Following the observation of NGO PL emission we have seen a systematic quench of one sharp UV line (390 nm) with very small amount of Ce4+ doping in the place of Nd3+. This PL quench has been explained by the novel idea of a super-hydrogenic dopant where the SHD radius is enhanced by the large dielectric constant of the host. The experimentally derived range of SHD radius comes out to be 0.85 nm to 1.15 nm where the predicted radius of 1.01 nm (from the bulk dielectric constant of the host) in the middle of that range. Thus using PL quenching we may able to estimate the Bohr radius of other super hydrogenic dopants in wide band-gap solids with large dielectric constants with the possibility of new functionalities.In this study we have reported the PL properties of NGO single crystals and polycrystalline bulk samples where UV (~360 and ~390 nm) and blue emission (~420 nm) lines were observed. The observed UV and blue emission can be understood based on the energy level diagram of the Nd3, NdAlO3 or Lu2O3, Nd2O3, Gd2O3 etc. to study SHD using the sharp photo luminesces from the rare earth atoms. This sharp atomic PL lines and their subsequent quench (with doping) are very crucial to study, understand and quantify the effect of SHD and its dimension. Other perovskite systems (widely used in oxide-electronics research) like SrTiO3, KTaO3 etc. may not be the ideal systems to study SHD phenomena as they usually have broad PL emission originating from band-edge or defect states.In final remark, we would like to add that although SHD means super-hydrogenic dopant – it clearly has nothing to do with doping using hydrogen atom/ions. Any element which can act as an electronic dopant can be a SHD in a host with large dielectric constant. However, in terms of the host material we need rare-earth based oxides and perovskite materials like NdGaO3) from CrysTec GmbH, Germany, for the single crystal study of the NGO pure system. Further, we have prepared one pure NGO poly-crystalline pellet and five different Ce doped NGO poly-crystalline pellets. For the pure NGO bulk sample we have used the high purity (99.995% or higher) powders of Nd2O3 and Ga2O3. Moreover, for other five doped NGO bulk samples high purity CeO2 powder was used along with Nd2O3 and Ga2O3 powders. At first, the powders are mixed for about 1 h by hand using mortar and pestle and the mixed powder was calcined at 1000 °C for ≈12 h in an oven at a closed environment in atmospheric pressure. The calcined powder is then further ground in mortar and pestle again and then subsequently pressed into 1-in diameter pellets by hydraulic press. Finally, the pellets were sintered at 1350 °C under the same atmospheric pressure for ≈16 h.In this study we have initially used (1 1 0) oriented NGO single crystal (5 × 5 × 0.5 mm2) and recorded using a JY-Horiba LabRAM HR spectrometer. Further elemental composition, valence and structural analyses were done using Rutherford backscattering spectroscopy ((using 3.5 Mev singletron accelerator)), secondary ion mass spectroscopy (using ToF-SIMS-IV from ION-TOF GmbH), x-ray photoelectron spectroscopy  and x-ray diffraction .PL properties of NGO single crystals and thin films were studied by exciting with the 325 nm line of a He-Cd laser with a spot size of ~800 nm and a power of 10 mW, .Publisher’s note: Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.


There is growing interest in replacing conventional TCOs such as indium tin oxide with lower cost, earth abundant materials. In the current study, we dope Zr into thin ZnO films grown by atomic layer deposition (ALD) to target properties of an efficient TCO. The effects of doping (0–10 at.% Zr) were investigated for ~100 nm thick films and the effect of thickness on the properties was investigated for 50–250 nm thick films. The addition of Zr4+ ions acting as electron donors showed reduced resistivity (1.44 × 10−3 Ω·cm), increased carrier density (3.81 × 1020 cm−3), and increased optical gap (3.5 eV) with 4.8 at.% doping. The increase of film thickness to 250 nm reduced the electron carrier/photon scattering leading to a further reduction of resistivity to 7.5 × 10−4 Ω·cm and an average optical transparency in the visible/near infrared (IR) range up to 91%. The improved n-type properties of ZnO: Zr films are promising for TCO applications after reaching the targets for high carrier density (>1020 cm−3), low resistivity in the order of 10−4 Ω·cm and high optical transparency (≥85%).Transparent conducting oxides (TCOs), with high optical transparency (≥85%) and low electrical resistivity (10 The dopants used in ZnO should be shallow donors that provide extra ionized electrons. Dopants such as B uch as B ,12, Ge −3 Ω·cm),uch as B  have beeuch as B ,17 and Guch as B ,19,20. Tncy ≥85% .2+ [4+ compared to Zn2+ [i.e., 0.745 Å for Zr and 0.740 Å for Zn [4+. An additional advantage of using Zr as a dopant is that it does not readily bond with Zn atoms, hence is unlikely to form secondary intermetallic phases [Zirconium was chosen as the dopant in the current work due to its abundance, comparable ionic size to Zn and because it can potentially act as double donor providing up to two extra free electrons per ion when substituted for Zn2+ . The clo to Zn2+   should Å for Zn . Al can Å for Zn , which cc phases . A numbec phases ,26, low c phases , sol-gelc phases ,28,29,30c phases ,32,33, pc phases , and ALDc phases . et al. [ALD was used in the current study to produce modulated (delta) doped ZnO during the deposition itself, which is considered an advantageous method for accurate control of the carrier concentration. The self-limiting nature of ALD also provides excellent control over film thickness, good uniformity and conformality, while the relatively low growth temperature permits the use of temperature sensitive substrates such as polymers. A comprehensive report is presented in this paper about the limits in electrical properties of Zr doping in ZnO thin films, and how this affects the optical properties. Lin et al.  reported2O sources held at room temperature, while the TEMAZ was heated at 95 °C. Films were deposited on sodalime glass microscope slides (cleaned with isopropanol and dried with nitrogen), and on virgin test grade Si (100) wafers. The growth rates of ZnO and ZrO2 individually at 200 °C were 1.87 and 0.65 Å/cycle respectively. Zr doped ZnO with target doping ratios between 0 and ~10 at.% were deposited using an ALD delta doping methodology similar to the one reported by Chalker et al. [2. The doping percentage was calculated from the number of doping cycles and the number of Zr atoms per ZrO2 monolayer. The film compositions presented in this study are nominal. The overall thickness of the layer was then controlled by repeating this “master” cycle multiple times. To improve the quality of the films, an ultra-thin (~2 nm) buffer layer of Al2O3 was deposited on the substrates by ALD  prior to the deposition of the doped ZnO. The Al2O3 process promotes uniform deposition of the target films (probably due to self-cleaning), and is also believed to enhance the optical performance, since it provides transverse optical confinement with minimum absorption loss due to its wide energy gap [The deposition was carried out in an Oxford Instruments OpAL ALD reactor  at 200 °C using diethylzinc (DEZ) and tetrakis-ethylmethylamino zirconium (TEMAZ) as Zn and Zr sources respectively (supplied by SAFC Hitech), with water vapour as a co-reactant. Each precursor was delivered via vapour draw, with DEZ and Hr et al. . As illuergy gap . Films grown on glass were used for all subsequent optical analysis, while films on Si (100) were used for microstructural analysis. Film thickness was assessed using a Rudolph Research Auto-EL-IV Ellipsometer  operating at 633 nm. The microstructure was analysed by atomic force microscopy (AFM)  using a Bruker Multimode 8. The microstructure was calculated from the X-ray diffraction (XRD) patterns obtained using a Rigaku Miniflex Diffractometer  in the Bragg-Brentano geometry and using a Cu Kα X-ray source. The chemical state of the films was determined by X-ray photoelectron spectroscopy (XPS) using a FISONS VG Escalab MKII Scientific XPS , with Al Kα radiation. The electrical properties of the films were assessed by four-point-probe (4PP) and by Hall effect measurements. The 4PP measurements were carried out using a homebuilt system comprising of a Lucas Signatone Corp SP4 Probe head  and a Keithley 2400 series sourcemeter , while the Hall measurements used the same sourcemeter attached to a BioRad Hall probe station with a 0.3 T fixed magnet. Hall measurements were carried out using van-der-Pauw sample configuration with indium contacts. The optical properties of the films were assessed by UV-Vis transmission using a Shimadzu SolidSpec-3700DUV dual beam UV-Vis spectrophotometer  and by photoluminescence (PL) using a Horiba JY LabRam HR HR800 confocal Raman microscope  fitted with a He-Cd UV laser (325 nm).20 to 3.81 × 1020 cm−3 as the doping is increased from zero to 4.8 at.%, leading to a decrease in resistivity from 3.02 × 10−3 to 1.44 × 10−3 Ω·cm. 4+ ions replace Zn2+ ions, thus donating two extra electrons to the system [Zn) in the bulk of the film rather than locating at the surface as is the case of Cu and Ag [Zn has lower formation energy than interstitial (Zri) and oxygen vacancy (ZrO) positions [Zn), which indicates that Zr is in its ionic state (Zr4+) in the Zr-O structure. The binding energies of Zr3d5/2 and Zr3d3/2 are 182.4 and 184.7 eV respectively , hydrogen acting as donor [4+). The hydrogen concentration is not expected to change significantly as Zr is added and hence is not expected to contribute significantly to the increase in carrier density due to doping. The main increase in carrier density with doping is attributed to Zr4+ acting as an electron donor, potentially providing up to two extra carriers per ion. If each Zr4+ dopant ion were acting as a double donor, then the addition of 4.8 at.% Zr to the ZnO would theoretically lead to an increase of ~4.1 × 1021 cm−3 in the carrier density. In reality, Hall effect measurements indicate that the addition of 4.8 at.% Zr causes the carrier density to rise by only ~2.8 × 1020 cm−3, which possibly suggests the formation of neutral defects during doping and segregation of Zr at grain boundaries.The ionic bonding between Znas donor  and the 2, which inhibits effective doping of the ZnO and hence leads to a reduction in carrier concentration. The ZrO2 tends to cluster at grain boundaries, and this has been shown to result in suppression of grain growth and hence smaller grains [As the doping level increases beyond 4.8 at.%, the carrier concentration begins to decrease again, resulting in an increase of resistivity. Previous studies ,37 have r grains ,37. AFM Zn). The PL data (Zn concentration for the heavier doped films. The PL shows two main emission features for the doped films attributed to the recombination of carries from the conduction band minimum (CBM) to the VZn0 state (3.14–3.15 eV) [Zn0 emission feature is the shoulder on the high energy side of the main peak and is related to the zinc vacancy concentration in the films. The integrated area of this PL feature was found by a nonlinear curve fitting (Loranzian) using two peaks for all films and the data are shown in the inset graph. The intensity of the high energy peak was zero for the two lower doped films (0–1 at.%) and thus is not shown in the graph. The integrated area of the CBM to VZn0 emission shows a slight reduction as the doping level increases from 3 to 5 at.% followed by an increase as the doping level increases up to 9.1 at.%. Zinc vacancies act as intrinsic acceptors and are easily formed in n-type materials by means of their formation energy decrease as the Fermi level increases [Zn is proportional to the grain boundary areas. The grain boundary areas increase at high doping level would be expected to correlate to an increase of the density of zinc vacancies per unit volume. The combination of VZn and Zr4+ segregated at the boundaries is likely to lead to the formation of neutral defects and hence to the reduction of free electrons. The increase of the defect concentration observed in the PL data is consistent with the carrier density reduction for the heavier doped films.In addition to dopant segregation having a direct effect on the effectiveness of the doping, the suppressed grain size due to the doping is believed to have a secondary effect on the carrier density by reducing zinc vacancies (V PL data  showed i3.15 eV)  and the ncreases . Grain bncreases , hence, 2/Vs It may be postulated that this behavior is consistent with an interface carrier scattering mechanism, with a higher value observed (for the thicker film) being consistent with that of the bulk material. The carrier density reaches 4.6 × 1020 cm−3, and it is largely independent of film thickness (for films >80 nm thick) as expected. The drop-off effect for very thin films (<100 nm) is attributed to the strain sensitivity of very thin films before reaching the bulk thickness.The carrier mobility  decreaseThe tensile strain measured from  [−4 Ω·cm) [The lowest resistivity value achieved in this work was 7.5 × 100−3 Ω·cm ,37, and [−4 Ω·cm . It is a−4 Ω·cm)  and Ga di.e., before normalisation) of ~85 nm doped ZnO films with various fractions of Zr. Films with 1 to 4.8 at.% Zr have a high normalised transparency, averaging 87%–88% in the visible (380–780 nm) and 89% in the near IR (780–1400 nm) range . The normalized transparency in the visible spectrum was at its maximum (91%) for a thickness of 161 nm, and in the near IR region reached 91% with a thinner film (50 nm). The high transparency and low resistivity results in high values of the figure of merit calculated by the formula derived by Haacke [T is the average transmittance in the visible region and sR is the sheet resistance. The value for the thicker doped film (250 nm) was 0.93 × 10−2 Ω−1 while for a thinner one (83 nm) was at 0.14 × 10−2 Ω−1. For comparison ITO typically has 1.19 × 10−2 Ω−1 [TCF = 0.10 × 10−2 Ω−1) [TCF = 0.12 × 10−2 Ω−1) [TCF = 0.12 × 10−2 Ω−1) [TCF = 0.36 × 10−2 Ω−1) [y Haacke  as expre10−2 Ω−1 . The res0−2 Ω−1) , 180 nm 0−2 Ω−1) , 100 nm 0−2 Ω−1)  and 200 r IR 780–00 nm ranΔE = 0.26 eV) which are in close agreement with PL values of 3.21 and 3.52 eV (ΔE = 0.31 eV). The slight differences are comparable with the expected difference between absorption and emission processes. As expected, the long wavelength plasma edge of the transmission spectra is affected by doping as the plasma frequency is proportional to the square root of the carrier concentration. The optical gap was measured experimentally by both optical transmission and by PL. In transmission the values for undoped and 4.8 at.% doped samples were 3.27 and 3.53 eV  due to many body effects from the increased carrier concentration. To confirm those effects, mathematical models were applied to calculate the effective masses of parabolic   effect d019 cm−3 ) and ban 0.28 m0 ) and non 0.28 m0  and to e 0.28 m0  and band 0.28 m0  as expreThe experimental results agree reasonably well with the non-parabolic band model with both effects. The non-parabolic CB was due to the band deformation by the high number of free electrons. Hence, by modulated Zr doping the optical gap can be tuned in a predictable way.4+ to Zn2+. At heavier doping the carrier concentration was reduced due to Zr segregation to grain boundaries consistent with the suppression of grain growth, and the formation of neutral defects at the boundaries (increase of zinc vacancies). The resistivity was decreased as thickness increasing (7.5 × 10−4 Ω·cm) leading to scattering reduction. The remaining decrease was accompanied by a carrier density increase (4.6 × 1020 cm−3) influenced by the tensile strain. The optical gap of the degenerate films could be enlarged up to 3.5 eV due to the net effects of the Burstein-Moss shift and bandgap renormalisation. The transparency of the doped films was up to 91% in the visible and near IR regions. The properties resulted in a high figure of merit (0.93 × 10−2 Ω−1) that is close to figures of merit for ITO films (1.19 × 10−2 Ω−1 [−2 Ω−1) [−2 Ω−1) [20 cm−3), low resistivity in the order of 10−4 Ω·cm and high optical transparency (≥85%).ALD grown ZnO: Zr films showed that doping offers control over resistivity reduction up to 4.8 at.% due to the extra ions offered by the substitution of Zr10−2 Ω−1 ) which i[−2 Ω−1)  and Hf-d[−2 Ω−1) . As a re


IV‐based secondary building units (SBUs), as in the UiO‐66 series, are receiving widespread research interest due to their enhanced chemical and mechanical stabilities. We report the synthesis and extensive characterisation, as both bulk microcrystalline and single crystal forms, of extended UiO‐66 (Zr and Hf) series MOFs containing integral unsaturated alkene, alkyne and butadiyne units, which serve as reactive sites for postsynthetic modification (PSM) by halogenation. The water stability of a Zr–stilbene MOF allows the dual insertion of both −OH and −Br groups in a single, aqueous bromohydrination step. Quantitative bromination of alkyne‐ and butadiyne‐containing MOFs is demonstrated to be stereoselective, as a consequence of the linker geometry when bound in the MOFs, while the inherent change in hybridisation and geometry of integral linker atoms is facilitated by the high mechanical stabilities of the MOFs, allowing bromination to be characterised in a single‐crystal to single‐crystal (SCSC) manner. The facile addition of bromine across the unsaturated C−C bonds in the MOFs in solution is extended to irreversible iodine sequestration in the vapour phase. A large‐pore interpenetrated Zr MOF demonstrates an I2 storage capacity of 279 % w/w, through a combination of chemisorption and physisorption, which is comparable to the highest reported capacities of benchmark iodine storage materials for radioactive I2 sequestration. We expect this facile PSM process to not only allow trapping of toxic vapours, but also modulate the mechanical properties of the MOFs.Metal–organic frameworks (MOFs) containing Zr IV MOFs, in particular the so‐called UiO‐66 series6O4(OH)4 clusters, demonstrate increased chemicalMetal–organic frameworks (MOFs)1 and 2, respectively, , whereby chemical transformations are performed on pre‐synthesised MOFs whilst maintaining crystallinity.3, constructed from a butadiyne‐containing organic ligand, alongside a pair of analogous interpenetrated isoreticular Zr and Hf MOFs,4 and 5, respectively, containing an extended 4,4′‐‐dibenzoate (peb2−) linker 4(sdc)6]n (1) to [Zr6O4(OH)4(meso‐sdc‐Br2)6]n (21‐Br), as both bulk microcrystalline and single‐crystal material , we decided to investigate the use of N‐bromosuccinimide (NBS) as a milder, alternative brominating agent  resulted in degradation of the MOF.00 °C res1. Addition of N,N′‐diphenylthiourea (DPT), a known catalyst for activation of NBS, gave small amounts of the bromohydrinated product, 1‐Br‐OH, alongside the brominated material 21‐Br (entry 3), presumably resulting from excipient water in the reagent or solvent. Incorporation of hydroxyl units within MOFs allows pore hydrophilicity to be tuned, yet direct synthesis of such MOFs is not always possible due to coordination of the hydroxyl group to the metal ions.1‐Br‐OH, with increased amounts of NBS resulting in slightly better conversion (entry 4), but the yield of 1‐Br‐OH remained low. Deliberate addition of water (entry 5) gave significantly more bromohydrinated product, while the combined presence of both water and DPT gave full conversion and up to 67 % bromohydrination (entries 6–8) versus bromination. Reduced quantities of NBS (entries 9 and 10) gave similar product distributions; around 60–65 % of 1‐Br‐OH and 35–40 % of 21‐Br. Similar product distributions were obtained when the reaction was performed on a larger scale (entry 11).Addition of 15 equivalents of NBS (entry 2) was necessary to quantitatively brominate 1, which results exclusively in meso‐sdc‐Br22−,R,S)/‐isomers are obtained . The fact that 1 remains crystalline during the transformation highlights its high mechanical stability, as it is able to withstand a change in geometry of the central carbon atoms as their hybridisation changes from sp2 to sp3. The chemical stability of 1 is also apparent as it retains its framework structure in the presence of significant amounts of water. The incorporation of hydroxyl groups generates a platform for subsequent functionalisation of MOFs, with alkoxide formation1, enabling the insertion of orthogonal functionality into the MOF in one‐step, we have chosen Br2 as a favoured brominating agent for Zr‐MOFs bearing integral unsaturated functionalities.Considering the similarity of bromohydrination to bromination of 2) for incorporation into a UiO‐66 series MOF. Bulk microcrystalline and single‐crystal samples of [Zr6O4(OH)4(bdb)6]n (3) were synthesised in the presence of 30 equivalents of benzoic acid .3 adopts the typical UiO‐66 topology, with Zr6O4(OH)4 clusters connected 12‐fold by bdb2− ligands  Å. N2 adsorption isotherms performed on bulk samples at 77 K prove the permanent porosity of 3, which displays a type I isotherm with stepwise adsorption observed at low partial pressures, characteristic of filling of the smaller tetrahedral and larger octahedral pores of UiO‐66 type MOFs  in CHCl3, with the transformation followed using a number of experimental techniques .With 43‐Br was evidenced by PXRD .trans‐dibromoalkene carbon atoms in 43‐Br compared to the analogous MOF 22‐Br  suggests that the single species is the trans,trans‐bdb‐Br4‐H2 product; that is, the bromination of 3 occurs stereoselectively. In contrast, the liquid‐phase bromination of the dimethyl ester of the ligand (bdb‐Me2), using conditions similar to those employed for the bromination of 3 , resulted in a 1H NMR spectrum containing resonances for three species, which we have assigned to be the trans,trans, cis,trans and cis,cis geometric isomers of bdb‐Br4‐Me2 .The transformation to D Figure  c; excel4 Figure  d provid−1, representative of the butadiyne functionality, combined with the emergence of a distinct shoulder on the alkene peak at about 1600 cm−1, attributable to the formation of the tetrabromodialkene moiety as 3 is brominated to 43‐Br. Excluding initial solvent loss (<200 °C), thermogravimetric analysis (TGA) in air reveals that 3 is thermally stable to about 460 °C, typical of UiO‐66 type MOFs . The TGA profile of 43‐Br displays an additional mass loss event that we expect to correspond to debromination, but which could not be fully deconvoluted from the overall decomposition. The total mass loss between 200–550 °C (82.8 %) corresponds very closely to the theoretical mass loss (82.9 %) expected upon decomposition of 43‐Br to ZrO2, in contrast to the 70.7 % mass loss for 3 over the same temperature range (expected 69.3 %), which, along with bromine analysis  and the spectroscopic data, suggests quantitative bromination.Raman spectroscopy Figure  e also s3 were suspended in CHCl3 and exposed to Br2 for four days, thereafter the CHCl3 was replenished several times before finally resolvating in DMF. In contrast to our solution of the crystal structures of 21‐Br and 22‐Br, it was only possible to collect unit cell parameters for 43‐Br. We would expect the lowest energy arrangement of the bdb‐Br42− linker to result in the tetrabromodialkene unit being geometrically frustrated both horizontally and vertically as a consequence of the steric bulk of the bromo substituents; this conformation has been observed in the crystal structure of a tetrabromodialkene analogue4‐H2  Å) and 43‐Br (a=32.7864(7) Å) clearly shows a mechanical contraction associated with the bdb2− linker shortening upon bromination, consistent with the brominations of 1 and 2.43‐Br is not porous and so cannot examine the mechanical modification by gas uptake. We hypothesise that the large mechanical contraction combined with the significant ligand disorder brought on by bromination may result in pore collapse.Single crystals of 2−), were reported in 2011.6O4(OH)4(peb)6]n (4) and the Hf derivative [Hf6O4(OH)4(peb)6]n (5) .l‐proline modulation were smaller (50 μm) than those from conventional benzoic acid modulated syntheses (100 μm), and so benzoic acid modulated crystals were analysed by SCXRD. Both 4  Å and a=39.806(5) Å, respectively, and are structurally similar to their functionalised analogues.2− ligands is evident, as they bow in and out of the linear plane separating adjacent M6O4(OH)4 (M=Zr or Hf) clusters. Experimental PXRD patterns of bulk samples of 4 and 5 prepared by l‐proline modulation show excellent agreement ]‐dibenzoate  obtained . Raman spectra of the MOFs . The TGA profiles of 4 and 44‐Br are clearly different  represents debromination, there is excellent agreement between the observed mass loss of approximately 39.1 % and the theoretical Br content of 40.1 %, which, in concert with a bromine content of 41.9 % by elemental analysis, suggests quantitative bromination has been achieved. N2 adsorption isotherms prove the permanent porosity of 4 which has a BET surface area of 2650 m2 g−1, decreasing to 1440 m2 g−1 upon bromination to 44‐Br , suggesting that the transformation from alkyne to dibromoalkene units results in a mechanical contraction of the ligand and thus the MOF.Excluding solvent loss (<200 °C), TGA reveals that t Figure  e, with 4 and 5 were suspended in CHCl3 and exposed to Br2 for four days, thereafter the CHCl3 was replenished several times before finally resolvating in DMF . Both MOFs were successfully brominated in a single‐crystal to single‐crystal manner with data collections possible, resulting in an accurate description of the structural alterations introduced during bromination. The brominated MOFs were found to retain the same Fd3‾m space group as their parent structures, with bromination resulting in a mechanical contraction of the MOFs, as observed through reductions in the unit cell edges from 39.8116(7) Å to 39.067(7) Å as 4 is brominated to 44‐Br and from 39.806(5) Å to 39.0451(3) Å as 5 is brominated to 45‐Br. The solid‐state structures of 44‐Br and 45‐Br exhibit significant disorder in their linker components, as would be expected due to the number of frustrated positions the dissymmetric linker could adopt along the linear vector.Single crystals of both 45‐Br, it is possible, despite the linker disorder resulting in multiple bromine positions, to visualise the trans,trans‐configuration of the peb‐Br42− linker , indicates that the central benzene ring of the linker is free to rotate. Comparison of the experimental PXRD pattern with the pattern calculated from our structure model of 45‐Br shows good agreement , with the total iodine uptakes measured through gravimetric investigations, alongside 1H NMR spectroscopic analysis to establish the percentage chemisorption , allowing us to crystallographically characterise trans‐edb‐I2‐H2,trans,trans‐bdb‐I4‐H2 and trans,trans‐peb‐I4‐H2, from the iodination of 2, 3 and 4, in turn .The stereoselectivities of these vapour phase iodinations are analogous to the solution phase brominations, as can be seen by the n Figure  b‐5d. Th2 vapours by the MOFs, combined with their high porosity, makes them excellent candidates for sequestration of radioactive I2, and we are currently examining the scope of substrates which can be irreversibly captured by the MOFs in this manner. In addition, the change in hybridisation of integral linker atoms upon halogenation will modulate the mechanical properties of the MOFs, and we are investigating these effects in detail.In conclusion, we have demonstrated that the combined chemical and mechanical stabilities of Zr and Hf MOFs enable their postsynthetic halogenation across integral unsaturated linker functionalities. Both solution‐phase bromination and vapour‐phase iodination proceed stereoselectively and in excellent yields across alkenes, alkynes, butadiynes and in interpenetrated frameworks, while the water stability of the Zr MOFs allows aqueous bromohydrination of alkene units. Bromination is shown to be reversible by abstraction with pyrrolidine to restore the parent material, making it an unusual example of reversible covalent postsynthetic modification, and we are currently investigating the potential for multiple cycles of sequestration and regeneration of the MOFs. The chemisorption of ICCDC 1443195–1443201 contain the supplementary crystallographic data. These data can be obtained free of charge by The Cambridge Crystallographic Data Centre.Experimental details for this work can be found in the Supporting Information. 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.


If this state were to become involved in superconductivity, under certain conditions a topological superconducting state could be formed, which is of high interest due to the possibility of creating Majorana fermionic states. We report directional point-contact spectroscopy data on the novel Bi2Te3/Fe1+yTe interfacial superconductor for a Bi2Te3 thickness of 9 quintuple layers, bonded by van der Waals epitaxy to a Fe1+yTe film at an atomically sharp interface. Our data show highly unconventional superconductivity, which appears as complex as in the cuprate high temperature superconductors. A very large superconducting twin-gap structure is replaced by a pseudogap above ~12 K which persists up to 40 K. While the larger gap shows unconventional order parameter symmetry and is attributed to a thin FeTe layer in proximity to the interface, the smaller gap is associated with superconductivity induced via the proximity effect in the topological insulator Bi2Te3.In the interfacial superconductor Bi Despite the exact mechanism for superconductivity in Bi2Te3/Fe1+yTe remaining unknown, it has been shown previously that any doping effect by O, Bi or Te impurities can be excluded2Te3 layer is required2Te3 quintuple layers (QL) increases, Tc rises from ~1.2 K (1 QL) to 12 K (5 QL) and then saturates. This correlates with the 5 QL critical thickness required to form a fully-developed surface state in Bi2Te382Te3 layer, and a highly complex superconducting mechanism is likely involved in the interfacial superconductivity.Tailoring a topological superconductor by combining the topologically protected surface states of a TI with a superconductor via the proximity effect is of enormous theoretical and technological interest, principally due to the possibility of finding the Majorana fermionic states which are predicted to exist in the vortex cores of topological superconductors12Te3 layer. We observe a pronounced twin gap structure with a large gap of unconventional pairing symmetry, corresponding to the interfacial superconductivity, and a smaller gap which opens below the characteristic BKT temperature and is associated with proximity-induced superconductivity in the topologically insulating Bi2Te3 layer. In addition, a pseudogap is observed which extends up to 40 K.In this article, we report directional point-contact data measured for current injection (a) parallel to the interface into the edge of the heterostructure, simultaneously probing both layers and (b) perpendicular to the interface into the top Bi2Te3(9QLs)/Fe1+yTe, chosen for their high Tc = 12 K. The Van-der-Waals bonding between the materials results in extremely high quality atomically-sharp interfaces.Point-contact spectroscopy is an energy-resolved technique directly probing the amplitude, symmetry and temperature-dependence of the superconducting gap2Te3 surface. An insulator-to-metal transition is visible in the form of a broad maximum at 76 K, which we associate with the antiferromagnetic transition of the bulk Fe1+yTe layer. At lower temperatures, the resistance passes through a minimum at 24 K before increasing steeply. The overall resistive behavior (including this increase) is typical for FeTe with a high content of interstitial excess Fey = 0.15 ± 0.02 in our Fe1+yTe bulk layer. It should be noted that FeTe does not exist in stoichiometric form and always has some excess Fe in the form of individual interstitial ions2Te3/Fe1+yTe interface reveals itself by a rapid resistance decrease below Tc = 12 K until zero resistance is reached at T0 = 8 K. From 2Te3 layer, the bulk Fe1+yTe layer and the intermediate thin interfacial layer. The resistance drop below 12 K is associated with the diverging superconducting correlation length in the interfacial layer upon approaching the BKT transition. However, 12 K is not necessarily the onset of the superconducting transition: the rapidly increasing resistance of the interface will cause the current to gradually move away from the interfacial region into the 140 nm thick bulk Fe1+yTe layer, which then begins to shunt the interface resistance. Therefore, the normal state resistance of the interface will be largely hidden due to the large Fe1+yTe thickness (140 nm) relative to that of the thin superconducting interfacial layer.Vb = 0, gradually deepening as the temperature falls. Until 15 K, the gap is rounded at low energy, but at 12 K the conductance flattens around Vb = 0 prior to the emergence of a zero bias conductance peak (ZBCP) below ~10 K. Concurrently, shoulder-like structures develop at ~10 mV and ~5 mV, which as we will now demonstrate correspond to a phase-coherent superconducting twin-gap structureZ. In such a highly two-dimensional superconductor, fluctuations are expected to significantly reduce the quasiparticle lifetime; we account for this in our model with an energy-dependent Dynes parameterΓexp[(│Vb│−Δ)/W] where Γ and W are free parameters. The shoulders at ~5 and ~10 mV are modeled by an anisotropic two-band s-wave order parameter Δ(α+(1 − α) cos θ)1,2 from a fit at our lowest achievable temperature 0.27 K, then fix α1,2 at these values for all other temperatures.To fit the temperature dependence of our data, we primarily employ a modified Blonder-Tinkham-Klapwijk (BTK) model for finite-transparency tunnel junctionsZ1 ≥ 0.35 for Δ1 and Z2 ~ 1000 for Δ2), we also attempt to model our data using a two-band Dynes model . Both models yield similar results: at 0.27 K, Δ1 = 6 meV and Δ2 = 12 meV from the BTK model, while Δ1 = 4 meV and Δ2 = 13 meV from the Dynes fit. Δ1 exhibits a pronounced anisotropy α ~ 0.7, whereas Δ2 is approximately isotropic (α = 1). In each case, Δ2 provides the dominant contribution to the spectral weight: 60 ± 5% versus 40 ± 5% for Δ1 in the BTK model, compared with a Δ1:Δ2:metallic ratio of 8 ± 0.5%: 46 ± 4% : 46 ± 4% in the Dynes model. The smaller values for Δ1 and its spectral weight from the Dynes fit are due to non-negligible Andreev reflections, which cannot be perfectly simulated by the metallic component within this model. To account for the presence of the ZBCP, we also attempted to reproduce this data using a two-gap d-wave model, but no improvement of the fit or significant variation of the gap values were observed compared to the anisotropic s-wave case. Δ1 closes at 8 K while Δ2 appears to close at 40 K. The magnitude of Δ2 abruptly increases above 8 K. As we will demonstrate later, Δ2 actually consists of a large superconducting gap which transforms continuously into a pseudogap in the temperature range between 8 and 12 K. The pseudogap then closes at 40 K.The fitting results are shown in 2Te3 surface. The current was injected perpendicular to the film and hence only probes the Bi2Te3 layer. A small 5.4 meV gap dominates the conductance at low temperature, in excellent agreement with the small gap Δ1 observed upon injecting the current into the edge of the interface. The conductance saturates for Vb > 6 mV and no signature of the larger gap Δ2 is observed. This suggests that the Bi2Te3 layer becomes superconducting and is responsible for Δ1, similar to what has been observed in Bi2Sr2CaCu2O8+δ/Bi2Se3172Sr2CaCu2O8+δ/Bi2Te32Sr2CaCu2O8+δ/Bi2Te2Se2Sr2CaCu2O8+δ/Bi2Se3 junctions21In 1,2 do not shrink significantly. This resilience of the overall gap structure demonstrates that the Cooper pairing strength is almost impervious to strong fields, in direct contrast with the critical field Hc2 = 17 T which has been estimated from resistance dataHc2 merely corresponds to a field-induced loss of phase coherence. We observe reductions in the ZBCP height of 20% and 30% in 15 T applied in-plane and perpendicular to the interface, respectively, measured with respect to the minimum in the dI/dV curve. The ZBCP does not experience any splitting, regardless of the magnetic field orientation. In In 2 ≥ 12 meV is a striking feature of our Bi2Te3/Fe1+yTe heterostructure. It exceeds the superconducting gap in bulk FeSe1−xTexTc of our heterostructures being comparable to Tc in these bulk materials. Our data constitute a demonstration of the potential for strong-coupling superconductivity which could persist up to far higher temperatures than the critical temperatures observed in the bulk iron-chalcogenides. The origin of this gap enhancement is unclear, but since the presence of the Bi2Te3 layer is essential for the appearance of superconductivity2Te3 play a certain role.The large size of Δ2Te3 layers in parallel, the twin-gap feature could also originate from two spatially separated regions each with its corresponding electronic bands involved. Superconductivity induced by the proximity effect in the Bi2Te3 layer could therefore play a role in the opening of the second gap. A proximity-induced small gap has previously been observed in devices of Bi2Se3 and Bi2Te3 which were mechanically bonded to the cuprate high-Tc Bi2Sr2CaCu2O8+δ: these showed a similar reduction of the gap in the cuprate when the TI became superconductingMultiband superconductivity with various gaps has been reported in various Fe-based superconductors10251 is approximately constant; upon increasing the temperature its magnitude decreases rapidly, reaching zero at ~8 K. In contrast, the larger gap Δ2 is almost temperature-independent between 10 K and 25 K, gradually decreasing towards zero when the temperature is increased further up to 40 K. Between 10 K and 8 K, Δ2 is slightly reduced, which clearly correlates with the opening of Δ1. Below 8 K, the sample exhibits zero resistancedI/dV close to Δ1,2. Above 8 K, global phase coherence is lost, the resistance gradually increases until the normal state is reached at 12 K and Δ1 vanishes, while Δ2 becomes a pseudogap which persists up to 40 K. The Dynes parameter Γ2 describing the quasiparticle lifetime in Δ2 also rises steeply above 8 K  origin. A pronounced pseudogap state above the superconducting critical temperature is most famous in the cuprate high temperature superconductorsTc10242829303132The edge contact data does not allow us to judge whether Δ1+yTe layer as seen in the resistivity , at which thermally-induced vortices and anti-vortices are bound into pairs45TBKT, creating a pseudogap in the density of states. The 10 meV magnitude of this superconducting gap indicates a potential for strong coupling superconductivity with stable Cooper pairing at temperatures well above 12 K. Superconducting fluctuations have been observed at temperatures many times higher than Tc in strongly-underdoped layered cuprate HTSCs36Tc2Te3(9QLs)/Fe1+yTe is supported by the fact that the pseudogap is partially suppressed by a magnetic field of 9 T in temperatures up to 40 K )38T0 is the phase coherence temperature, represented by the establishment of zero resistivity at 8 K. From However, in a pseudogap which is caused entirely by fluctuations of the superconducting order parameter, it is expected that the zero-bias conductance G2Te3 layer, interfacial contact with the Fe1+yTe induces superconductivity in Bi2Te3, which is therefore a potential candidate to host a topological superconducting state41421 increasing from zero to 5 meV within this temperature range. The ZBCP is therefore of intrinsic origin. The proximity effect of an s-wave superconductor on a TI2Te3(9QLs)/Fe1+yTe remains unclear, two possible mechanisms exist for the formation of Majorana edge states. One possibility is that the contact with FeTe drives the Bi2Te3 layer to become an intrinsic topological superconductor with Majorana surface states4041452Te3 would be required to confirm this possibility. The alternative mechanism requires nodal dx2−y2 superconductivity (associated with the Fe1+yTe) combined with Rashba spin-orbit coupling d-wave order parameter alone could create a fermionic ZBCP at the sample edge, a ZBCP composed of the fermionic edge states and Majorana fermions should form in the presence of strong spin-orbit coupling at zero field. It has been predicted that an in-plane applied field will split and shift the fermionic states to finite energy via the Zeeman effect, while the Majorana state remains at zero energyd-wave scenario unlikely. Furthermore, the ZBCP visible in our top contact  superconductors causes pronounced local in-gap states, which do not split, shift or vanish in applied magnetic fields2Te3 layer (in which Fe impurities are absent) discourages this interpretation and rather points to an unconventional pairing mechanism as its origin.A more conventional explanation for the ZBCP is related to the large Fe excess in the Fe2Te2Se junctions with a very similar temperature and magnetic field dependence as in our heterostructure2Te3/Fe1+yTe heterostructures.A similar ZBCP was observed in Bi2212/Bi2Te3(9QLs)/Fe1+yTe heterostructures reveal a highly unusual superconducting state with an extraordinarily large superconducting gap, a pronounced pseudogap and compelling evidence for proximity-induced superconductivity in the topological insulating Bi2Te3 top layer. Our experiments alone are not able to prove the topological nature of this superconductivity, nor provide an indisputable explanation of the origin of the interfacial superconductivity. Nevertheless, it is clear that Bi2Te3(9QLs)/Fe1+yTe interfaces display a similarly rich behavior to the cuprate superconductors, which remain one of the major unsolved mysteries in physics.Our Bi2Te3 film. Detailed information about the quality and characterization of this interface, including scanning transmission microscope micrographs and resistivity data may be found in Ref. 2Te3(9QLs)/Fe1+yTe, chosen for their high Tc = 12 K.The heterostructure studied in this work was synthesized by a VG-V80H MBE system. A 50 nm ZnSe buffer was first grown on a GaAs(100) semi-insulating substrate. A 140 nm thick FeTe layer is then deposited on the buffer layer, followed by a 9QL thick Bi2Te3 surface. The edge of the sample was finely polished, instantly covered by a thin layer of Au, and an isolated 100 nm wide Au strip was separated using a focused-ion-beam. The maximum contact area is ~100 × 149 nm2 . A schematic drawing of the contact configuration is shown in To fabricate a point-contact device on the edge of the bilayer, a thin slab was attached to a silicon substrate with one edge facing upwards. Ordinary low-Ohmic contacts were prepared by RS 186–3593 silver conductive paint on the BiZZ = 0 corresponds to pure Andreev reflection, while larger values represent the spectroscopic tunneling regime. In our contacts we consistently achieve Z ≥ 0.35 and our nanoscale contact area ensures that our experimental tunneling regime is ballistic and not thermal or diffusive, i.e. the applied bias voltage Vb corresponds to the electron injection energy. This is confirmed by the temperature-independent value of the normal-state contact resistance. A point contact with similar properties was established on the Bi2Te3 surface for perpendicular current injection with the help of a scanning probe device, by gently approaching a tungsten tip to the Bi2Te3 surface. The differential conductance dI/dV vs Vb was measured at temperatures from 0.27 K to 70 K in magnetic fields up to 15 T with a quasi-four-probe method, using a Keithley 6221 AC/DC current source to generate a small, constant-amplitude (10 nA) AC current IAC with frequency 5 Hz, superposed on a ramped DC bias current. A standard lock-in technique in combination with a DC multimeter was used to measure dI/dV and Vb = VDC across the junction.The edge contacts typically had resistances in the kΩ range and the data reproducibility was verified on 4 different devices. The point-contact spectral shape is strongly dependent on the tunnel barrier height parameter How to cite this article: He, M. Q. et al. Pseudogap and proximity effect in the Bi2Te3/Fe1+yTe interfacial superconductor. Sci. Rep.6, 32508; doi: 10.1038/srep32508 (2016).


PM2.5 and carbonaceous particle concentrations have been declining appreciably since 2010/11 due to the use of diesel particle filters, but little change is seen in nitrogen dioxide over the period from 1995 to 2015. It is shown that the effect of NO2 from road traffic upon premature mortality was ten-fold greater than that of PM2.5 even before the widespread use of diesel particle filters, and is now considerably larger. The overwhelming contribution of diesel compared to gasoline-fuelled vehicles to emissions of both PM2.5 and NO2 emphasises the importance of further controls on emissions from diesels.Road traffic is a major source of urban air pollution responsible for substantial premature mortality. Until recently, attention has focussed primarily on exhaust emissions of particulate matter from traffic as a causal factor. From analysis of air quality measurement data from the UK and France, we demonstrate that road traffic exhaust has a far greater impact on concentrations of nitrogen dioxide than of PM Catalytic converters fitted to gasoline-powered vehicles have led to huge reductions in airborne concentrations of carbon monoxide and benzene in cities throughout the developed world7, but despite their well demonstrated adverse effects on human health, particulate matter9 and nitrogen dioxide11 have yet to be adequately controlled.The adverse effects of air pollution exposure on human health are well established12 which has required the fitting of diesel particle filters. In the case of oxides of nitrogen, there has been a gradual tightening of the type approval standard, from 0.5 g km−1 for Euro 3 light duty vehicles (in 2000) to 0.18 g km−1 for Euro 5a in September 2009 and a further reduction to 0.08 g km−1 for Euro 6 introduced in September 2014. In the case of heavy duty vehicles, Euro III required 5 g kWh−1 for NOx in October 2000, reducing to 2.0 g kWh−1 in October 2008 and to 0.4 g kWh−1 with Euro VI in January 2013. As highlighted by the Volkswagen emissions scandal, the type approval standards set by the European Union may not be representative of real-world driving emissions and various studies have shown a huge discrepancy between on-the-road emissions and the type approval limits tested in the laboratory14.In Europe, vehicle emissions are regulated through the Euro standards, and the Euro 5b standard, introduced in September 2011 for light duty vehicles, and Euro VI introduced in January 2013 for heavy duty vehicles, placed limits upon not only the mass of particles emitted per kilometre, but also regulated the particle number15 link changes in air quality at London sites to emission trends, but typically pollutants have multiple sources, and consequently changes in ambient concentrations cannot easily be attributed to an individual source category. In the case of road traffic, this problem can be overcome through the use of paired site studies in which a heavily trafficked roadside air sampling site is paired with a background monitoring site, the difference between the two representing the road traffic increment caused by emissions on the road in question16. In this study, we have used routine ambient air quality measurements downloaded from the UK Automatic Urban and Rural Network17 and Airparif18 for such paired sites in London, Glasgow and Paris in order to evaluate the change in vehicle emissions over recent years since the introduction of diesel particle filters and the implementation of the stricter Euro regulation for NOx emissions.The only truly representative way in which to evaluate the efficacy of emission controls is through ambient air measurements. Font and Fuller20. Traffic flows, and the vehicle mix have not changed appreciably over the period of the measurements  and organic carbon (OC) appear in Fig. 10 mass between 2009–11 (average) and 2015 is 8.7 µg m−3 (66.6%), and for PM2.5 is 3.0 µg m−3 (37.5%). Since exhaust emissions of particulate matter are made up predominantly from EC and organic matter, an approximate mass can be estimated from the sum of EC + 1.2 OC21. Data are available only since 2010 and the difference between 2010–11 and 2015 for EC is 2.44 µg m−3 (43.4% decrease) and for 1.2 OC is 0.80 µg m−3 giving a total reduction in particulate matter mass of 3.2 µg m−3 (35.5% decrease) which is very close to that for PM2.5  giving a clear indication that this is predominantly responsible for the reduction of the fine particle mass over this period.The paired sites of London Marylebone Road (roadside) and London North Kensington (urban background) have been used extensively in the past for evaluation of road traffic incremental pollution−3 in 2010–11 to 11657 cm−3 in 2015, an overall percentage reduction of 48.3%. The particle number count is dominated by the ultrafine particle fraction whose emission is sensitive to the sulphur content of motor fuel22. As fuel sulphur levels have been fairly stable over the period of these measurements, the reduction in particle number count would appear to be attributable to the introduction of DPFs which gives some reassurance the reduction in particle mass was accompanied by a commensurate reduction in the concentration of ultrafine particles.The trends in particle number count show a reduction in the roadside increment from 22549 cm2 and NOx for the entire period, but a substantial recent fall in ΔPM2.5 and a possible upward trend in ΔPM10. Further data on particulate matter mass are available from paired sites in Paris, Boulevard Périphérique Auteuil and Vitry-Sur-Seine, separated by 12.5 km. In this case, taking the same time period of 2009–11 to 2015, PM10 reduced in concentration by 4.0 µg m−3 (17.5%) and PM2.5 by 7.3 µg m−3 (55.1%) , required since 2011, is already having a very significant benefit for air quality through a reduction of around 36% in vehicle exhaust emissions of PM2.5 at Marylebone Road was 8.7 µg m−3, for nitrogen dioxide was 61.6 µg m−3, and for NOx was 222 µg m−3. As that NOx disperses in the urban atmosphere, the NO, which predominates in the street canyon, is largely converted by oxidation to NO2 and the NO2/NOx ratio at North Kensington in 2010/11 was 36.4/56.5 = 0.64. Consequently, the incremental NOx has an equivalent concentration of 222 × 0.64 = 143 µg m−3 of NO2 at a background urban atmospheric NO2/NOx ratio. This implies that in 2010/11 the ratio of contributions from traffic to background urban concentrations of NO2/PM2.5 was 143/8.7 or approximately 16.4 in the background urban atmosphere where most human exposure occursIn 2010/11 before the introduction of diesel particle filters, the incremental concentration of PM2.5 and NO2 can be judged crudely from their effect on all-cause mortality. This is not straightforward as PM2.5 and NO2 concentrations are normally highly correlated and under such circumstances the results of two pollutant models carry substantial uncertainty in the attribution of health effects to each pollutant. The WHO HRAPIE project3 recommended a Hazard Ratio coefficient for PM2.5 of 1.062 per 10 µg m−3. HRAPIE3 also recommend a coefficient linking NO2 and all-cause mortality of 1.055 per 10 µg m−3 but recommended reducing this by up to 33% to allow for double counting of effects of particulate matter, due to its high correlation with NO2. Subtracting the maximum of 33% from the NO2 coefficient yields a Hazard Ratio of 1.037. HRAPIE recommended application of the coefficient for concentrations >20 µg m−3. However, new evidence and in particular that from the ESCAPE study23 has shown linearity to much lower concentrations (the lowest concentration of NO2 in the ESCAPE cohorts was 1.5 µg m−3) and we therefore feel confident in applying this coefficient to the annual mean concentrations in London, which in any case well exceed 20 µg m−3. HRAPIE did not recommend attenuation of the PM2.5 coefficient for the effects of NO2, although in some studies in the meta-analysis upon which it is based, NO2 may also have exerted an effect. They did however note the possibility of double counting. HRAPIE3 noted that studies with a better quality of exposure estimate yielded on average higher hazard ratios and suggested that the result of the meta-analysis which they used  could result in an underestimation of the exposure-response slope. This being the case, we have chosen not to attenuate the coefficient for PM2.5. Had we done so, the relative proportion of health impact attributed to nitrogen dioxide would have been greater. This then gives a ratio of impact per unit mass for NO2/PM2.5 of 0.037/0.062 = 0.60. However, if the relative impacts of emission of the two pollutants from road traffic is to be estimated, the relative concentrations of the two pollutants in the background urban atmosphere (16.4 in 2010/11 and 35 in 2015) need to be taken into account, giving a ratio of 9.8 in 2010/11 and 21 in 2015 for the mortality due to traffic emitted nitrogen dioxide relative to PM2.5 from traffic. The implication is that the premature mortality impact of nitrogen dioxide from road traffic in 2010/11 exceeded that due to PM2.5 from this source by almost ten-fold and in 2015 by more than twenty-fold. There are uncertainties around this ratio. The specificity of the effects which correlate with NO2 concentrations has not been firmly established, and there may be a contribution from other, highly correlated pollutants such as ultrafine particles. On the other hand, the coefficient used for PM2.5 has not been attenuated to allow for the effect of NO2, and may over-estimate the impacts of PM2.5.The relative public health impacts of PM2.5 and NOx taken from the National Atmospheric Emissions Inventory24 disaggregated according to the vehicle type and fuel. This shows that diesel vehicles in urban areas are responsible for 97% of the PM2.5 exhaust emissions from road traffic and 90% of the NOx emissions. In the case of NOx, which arguably is responsible for a larger burden of disease than the PM2.5 from road traffic, there is a clear and urgent need to reduce the volume of diesel traffic on urban roads or to ensure that new vehicles entering the road emit far less NOx per kilometre than the current diesel fleet. The National Atmospheric Emissions Inventory24 predicts a reduction in NOx from road transport from 300 kt in 2014 to 174 kt in 2020 and 128 kt in 2025. It remains to be seen whether such reductions will in fact be achieved as this will depend critically upon the ability of motor manufacturers to install effective abatement devices on the vehicles which they produce, and that those devices remain fully effective over the lifetime of the vehicle. Past history gives little confidence that this will be achieved.Table Supplementary Information


Graphene oxide‐wrapped amorphous copper vanadium oxide is fabricated through a template‐engaged redox reaction followed by vacuum dehydration. This material exhibits high reversible capacity, excellent rate capability, and out standing high‐rate cyclability. The outstanding performance is attributed to the fast capacitive charge storage and the in situ formed copper with enhanced electrical conductivity. It has been reported that amorphous MoO2 and vanadium oxides with structural defects are able to deliver very stable cycling performance.Vanadium belongs to the early transition metal elements. It is generally accepted that vanadium‐based oxides react with lithium through the insertion reaction,Compared with binary metal oxides, ternary metal oxides have been intensively studied due to the increased electronic conductivity owing to the possible hopping processes and/or defect effect mechanisms.Figure−1 after 100 cycles at 100 mA g−1. Even at a high current density of 20 A g−1, a capacity of 672.5 mA h g−1 is retained after 7000 cycles, suggesting the excellent high‐rate capability and cyclability.Herein, we fabricate an amorphous copper vanadaium oxide decorated with graphene oxide (denoted as a‐CVO‐GO) through a facile and environmental‐friendly template‐engaged redox reaction followed by heat treatment in vacuum Figure1. The 2O spheres. The as‐prepared Cu2O template is composed of monodispersed nanospheres with a diameter of >>300 nm . The mixing of NH4VO3 with Cu2O colloidal solution at room temperature leads to the formation of hierarchical Cu3V2O7(OH)2·H2O nanoflowers with sheet‐like building blocks . During this process, the Cu2O nanospheres act as both the template and the Cu precursor; O2 is also involved in the reaction (6 Cu2O + 8 VO3− + 3 O2 + 8 H2O = 4 Cu3V2O7(OH)2 + 8 OH−). The reaction begins at the interfacial region between the Cu2O sphere and NH4VO3 solution. With the inward diffusion of VO3− and O2, the reaction propagates gradually from the outer surface to the inner core . Cu3V2O7(OH)2 hierarchical nanoflowers can be formed within 10 h.A typical synthesis procedure Figure  starts f3V2O7(OH)2. H2O nanoflowers at 250 °C. X‐ray diffraction (XRD) pattern of the a‐CVO‐GO shows no obvious peaks, indicating the poor‐crystalline or amorphous nature 2. H2O precursor. Moreover, the EDX measurements prove that the a‐CVO is composed of Cu, V and O and these three element distributed homogeneously in the selected area . The thickness of the nanosheets is estimated to be around 20 nm. The hierarchical nanoflower morphology is further confirmed by transmission electron microscopic (TEM) images . The thermogravimetry (TG) curve reveals that the content of graphene oxide in the a‐CVO‐GO is 8.06 wt% .The amorphous copper vanadium oxide nanoflowers can be obtained by vacuum dehydration of the Cure Figure2A. To 3V2O8 (JCPDS card no. 00–049–0689). The c‐CVO (Cu3V2O8) shares the same Cu/V ratio with a‐CVO‐GO . The morphology of c‐CVO is characterized by SEM , which shows nanoparticle‐constructed spheres instead of nanosheet‐constructed nanoflowers. No graphene oxide nanosheet is observed on the surface, which indicates that the graphene oxide is removed during the calcination in air. In addition, due to the calcination induced crystallization, the c‐CVO shows a significantly reduced surface area than that of a‐CVO‐GO .For the purpose of comparison, crystalline copper vanadium oxide (denoted as c‐CVO) is prepared by annealing the precursor in air at 450 °C. The XRD pattern  of the c‐CVO matches well with that of CuFigure−1. Generally speaking, the CV curves in the initial three cycles almost overlap, indicating the excellent reversibility. The charge–discharge curves at a current density of 100 mA g−1 are presented in Figure −1, respectively. The initial columbic efficiency reaches 109%, which is much higher than that of other transition metal oxide anodes based on conversion reaction mechanism. Similar anomalously high columbic efficiency has been observed in graphite anode by Dahn et al.+ insertion.−1 . At an even higher current density of 50 A g−1 , the a‐CVO‐GO delivers a capacity of 505 mA h g−1 after the first cycle, which is well above the theoretical capacity of commercialized graphite (372 mA h g−1). A capacity of 380.9 mA h g−1 can be retained after 100 cycles, corresponding to a capacity retention of 75%. However, fast capacity decay is observed after 140 cycles and this may be associated with the decomposition of electrolyte . To determine the function of graphene oxide, the electrochemical performance of a‐CVO and a‐CVO‐GO is compared in Figure S14, Supporting Information. The a‐CVO shows inferior high‐rate cycling stability. This may attributed to the reason that the graphene oxide nanosheet is able to offer increased electron conductivity especially when the structure of the electrode is destroyed.Figure −1 were performed  in Figure v1/2. Clearly, the capacitor behavior (current ∝ v) is not observed, eloquently demonstrating that lithium ions/electrons are stored in the bulk of a‐CVO‐GO rather than at the surface at the potentials of 0.01 and 3 V.To explain the extremely high rate performance of a‐CVO‐GO, the CV curves at different scan rates from 1 to 10 mV sed Figure4A. Thek1) and diffusion (k2) controlled processes Next, we used the analysis proposed by Dunn and co‐workersFigure−1. In contrast, with the increase of sweeping rate from 0.1 to 10 mV s−1, the diffusion controlled charge storage decreases from 831 to 71 mA h g−1. As a result, the contribution to the total capacity decreases from 65.3% to 16.0%. It is believed that the charge storage is mainly diffusion controlled in conventional lithium‐ion battery electrode materials. However, it is observed that the capacitive charge storage dominates the total capacity at sweeping rates above 0.5 mV s−1 in the a‐CVO‐GO. A typical separation of the capacitive and diffusion currents at a scan rate of 2 mV s−1 is shown in Figure The amount of charge stored due to both diffusion and surface limited processes at different scan rates is presented in Rct) of the electrode; while the inclined line stands for the Warburg impedance (ZW). The equivalent circuit model is provided in Figure RΩ represents the Ohmic resistance of the electrode system, including the electrolyte and the cell components. RSEI and Rct represent the resistance related to solid electrolyte interface (SEI) and the charge transfer, respectively. CPE1, CPE2, and Zw are the capacitance related to SEI, the double layer, and the Warburg impedance, respectively.−1, the charge transfer resistance of a‐CVO‐GO becomes even lower  in the low frequency range 2 aqueous solution under rapid magnetic stirring. N2H4 solution (17 μL) was then introduced into the above solution. The orange colloidal solution was formed immediately. Under continual stirring for 40 min, Cu2O spheres were collected by centrifugation and washed with water and ethanol for several times.Preparation of a‐CVO‐GO and c‐CVO: The graphene oxide was prepared by a modified Hummer's method. The fresh‐made Cu2O spheres were dispersed in 200 mL of distilled water. 50 mL of 0.02 m NH4VO3 aqueous solution was added to the Cu2O colloidal solution dropwise. Then 4 mL of the graphene oxide suspention (1 mg mL‐1) was added and ultrasoniccally dispersed. After stirring for 10 h, the color of the colloidal solution turned from orange to olive. Then the particles were collected by centrifugation and washed with distilled water and ethanol for three times. The precursor was desiccated at 70 °C in an oven. For the preparation of a‐CVO‐GO, the precursor was then transferred to the vacuum oven at 250 °C to remove the crystalline water. As a control experiment, the c‐CVO was synthesized by annealing the precursor at 450 °C in the air.Material Characterizations: XRD patterns of the samples were collected with a D8 Advance X‐ray diffractometer with area detector, using Cu Kα radiation (λ = 1.5418 Å). The microstructures were observed by field‐emission scanning electron microscopy (JEOL‐7100F), transmission electron microscopy, and high‐resolution transmission electron microscopy (HRTEM) (JEM‐2100F). Brunauer Emmet‐Teller surface area was measured by using Tristar II 3020 instrument. Thermogravimetry is carried out on a STA449c/3/G (NETZSCH).Electrochemical Characterizations: The electrochemical properties were characterized by assembing 2016‐type coin cells with lithium foil as the anode in a glovebox filled with pure argon. The cathode electrodes were composed of 60% active material, 30% acetylene black, and 10% sodium alginate binder. After coating onto copper foil, the cathode was cut into round slice with >>0.36 cm2 in area and >>0.1 mm in thickness. The mass loading of active material is 2.5–2.8 mg cm−2. A solution (1 m) of LiPF6 in EC/DMC (1:1 vol/vol) was used as the electrolyte. The cells were aged for 12 h before charge/discharge processes to ensure full penetration of the electrolyte into the electrodes. Galvanostatic charge/discharge measurements were performed by a multichannel battery testing system (LAND CT2001A). Cyclic voltammetry (0.01–3 V) was performed using an electrochemical workstation (CHI 660), electrochemical impedance spectroscopy (EIS) were tested with an Autolab Potentiostat Galvanostat (PGSTAT302N). All the measurements were carried out at room temperature. In order to perform the SEM images after cycling, the cells were disassembled and the cathode was dispersed in the alcohol for 72 h to remove the binder and the electrolyte.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.