The resulting samples are analyzed by X-ray diffraction (XRD) and scanning electron microscopy (SEM) including electron backscatter diffraction (EBSD). Ge-Fresnoite  is observed at the immediate surface of all samples and oriented nucleation is proven in both compositions. After a very fast kinetic selection, the crystal growth of BTG into the bulk occurs via highly oriented dendrites where the c-axes are oriented perpendicular to the surface. The growth of this oriented layer is finally blocked by dendritc BTG originating from bulk nucleation. The secondary phases BaTiGe3O9 (benitoite) and BaGe4O9 are also identified near the surface by XRD and localized by EBSD which additionally indicates orientation preferences for these phases. This behaviour is in contrast with previous reports from the Ba2TiSi2O8 as well as the Sr2TiSi2O8 systems.Two glasses of the compositions 2 BaO - TiO Barium may be partially or completely substituted by Sr to form Sr2TiSi2O8 (STS) while Si may partially or completely be replaced by Ge to form Ba2TiGe2O8 (BTG) or a solid solution. Incommensurate modulations occur in all three compositions3213Fresnoite in its original form has the composition BaAs these crystals are not ferroelectric and thus cannot be poled, the polar BTG crystals must be aligned during crystal growth if piezoelectric properties are required. Apart from growing single crystals in the Czochralski process, glass-ceramic materials containing oriented BTG crystals have been produced via the surface crystallization of glasses789101112131415161718192021222325262728302 were chosen to avoid spontaneous crystallization during cooling as well as homogeneous bulk nucleation. In order to suppress bulk nucleation789101114151617181920212223Usually, glass compositions with a certain excess of GeO2-GeO2 system has been studied22. At least the 22 glass compositions presented in 2, NiO, Fe2O3, V2O5, Sb2O3 or metallic Cu, all other compositions only vary in the concentrations of the basic components BaO, TiO2 and GeO2. Selected properties of these glasses and the resulting glass ceramics are also stated.The glass forming range in the BaO-TiO3+ is generated in the melt due to the reducing conditions at high melting temperatures3+/Ti4+ ratio should increase if the glass is melted at higher temperatures. The presence of Ti3+ leads to a decrease in viscosity and network connectivity and gives rise to a dark blue to violet coloration. This has already been shown for the case of fresnoite melts in the BTS system3+ led to higher nucleation rates and hence promoted crystallization254+ is reduced to elemental Si which in turn forms a liquid Pt-Si-alloy in the melt. As there are multiple Pt-Ge-alloys2 containing glasses considering the reducing conditions at high temperatures.An effect of the glass melting temperature on the optical properties of a glass in this system has been reported335 pyramidal units132 –3.67 GeO2  is observed after annealing at temperatures from 695 to 720 °C while it changes very little from 720 to 740 °C and decreases at 750 °C, probably due to the bulk crystallization triggered by the formation of BaTiGe3O9Crystallized BTG glass-ceramics range from transparent13151617181920212213381314151617181920213O9 (benitoite) and BaGe4O9 have been reported at the immediate surface after crystallizing glasses containing additional GeO2141616183O9 phaseWhile only BTG was detected after crystallizing a stoichiometric melt2 –3.67 GeO2 glass at 690 °C2012Specific annealing regimes enabled to solely crystallize oriented benitoite from the 2 BaO – TiO2 provided the first case where the texture detected at the immediate surface also prevailed during growth334042XRD-patterns of the surface crystallized layers showed a non-statistical orientation distribution indicating a 001 texture79131416171819202122919211617402 – 3.67 GeO2 glass has been proposed4+ ions (5 fold instead of 4 or 6) may be related to the crystallization mechanism and that benitoite appears firstA mechanism for the formation of BTG in the 2 BaO – TiO33-values from 3 to 24 × 10−12 m/V have been reported in various glass-ceramics7891011121314et al. used the Marker Fringe technique and different refractive indices in the calculations for some articles1333-values increase if the topmost layer of crystals is removed and are hence linked to the degree of crystal orientation in the glass-ceramic33-value of 10 × 10−12 m/V stated for BTG single crystals733-meter.Piezoelectric measurements have shown that the positive end of the poles points towards the growth directionDue to the properties outlined above, fresnoite based glass ceramics have been proposed for applications as hydrophones2O3·B2O3 glassThe application of electron backscatter diffraction (EBSD) to the surface crystallization of glass-ceramics containing BTS403839383940412 · 2.75GeO2 were 0.75 mol% of additional GeO2 are added to the stoichiometric BTG composition to enable the direct comparison to previous analysis featuring BTS40382 and 10 cm long2 –3.67 GeO2, has been intensively studied for its optical non-linearity  increases with the increasing GeO2 concentration. In contrast Tx(bulk) seems to have a maximum somewhere between adding 0.75 and 1.66 mol% of GeO2.3O9) and BaGe4O9 are expected to only form locally and close to the surface of the BTG glass-ceramics which poses an ambiguity problem for building the material files necessary for EBSD analysis. To improve the reliability of these material files, powders of both phases were produced via solid-state reactions. Material files were then built and optimized using the EBSD-patterns obtained from the respective powder and the Inorganic Crystal Structure Database (ICSD) files 36185 and 83734.The phases benitoite  currently contain an entry with the Wycoff-positions of Ge-benitoite which are, however, necessary for building a material file in the given software. As a rule of thumb, a pattern may be considered reliably indexed in the TSL software if the primary orientation solution receives at least 30 votes, a fit of less than 1° and a Confidence Index (CI) value larger than 0.100.ICSD-file 36185 features hexagonal benitoite of the composition BaTiSi4O9, the material file also showed very good results but a confidence index of 0.000 for almost all detected orientations due to pseudo symmetries where the first and second orientation solutions receive the same number of votes. This problem also occurred during the analysis of surface crystallizing BaAl2B2O7 glass-ceramics4O9 but with space group 143 instead of 150) led to practically identical results. Thus the CI fails as a reliability indicator in its usual form in the case of the material files built for BaGe4O9. If, however, an orientation solution receives enough votes and a sufficient fit factor in combination with a CI of exactly 0.000, this may be considered as a correct orientation but with an ambiguity of a 180° rotation around the [0001] direction.The material file of hexagonal benitoite performed very well for single patterns as well as in a scan of the embedded benitoite powder. In the case of trigonal BaGe4O9 is that the achieved material files may be used to index the EBSD-patterns of both phases, making a phase separation based on EBSD alone unreliable. This type of problem may be solved by either fine tuning the material files based on characteristic kikuchi bands, e.g. for the separation of cubic and hexagonal ZnSThe real problem for the analysis of benitoite and BaGeet al.Two entries concerning the low temperature phase of BTG can be found in the ICSD data base: file no. 39133  and file no. 281271 . While the a and b axes almost have the same values in both files, the c-axis in file 281271 is basically twice as long as in file 39133 because it was doubled in order to describe a modulated structure at room temperature4O9 powder.The EBSD-patterns presented in 4O9 is achieved. Additionally, the position of the 001-pole is constant for the first two orientation solutions for these patterns. Hence even if the CI-value drops below 0.1 there is still a good chance that the orientation of the c-axis is correct.The patterns 1–4 are clearly indexed as BTG using a material file based on the ICSD file 39133. Despite the minimal difference between the a- and b-axes in the material file, the indexing parameters indicate a very reliable orientation solution and a clear separation between BTG and benitoite/BaGe4O9. These fit values can easily be decreased by choosing a smaller interplanar angle tolerance (IAT) than the 3° used here which, however, proved to be less useful during the EBSD-scans. The CI-values of the BaGe4O9 file are always 0.000 due to the pseudo symmetry problem outlined above, making the CI-value useless as a phase differentiating factor in this case. The patterns 6 and 8 receive exactly the same number of votes from both material files and pattern 5 receives almost the same fit factor. Activating the band width ratio matching function did not lead to a satisfactory phase separation either.Apart from some fit values slightly larger than 1°, the patterns 5–8 receive reliable indexing parameters for both benitoite and BaGe4O9. The latter two phases may not be separated reliably at this point and hence only the material file of benitoite is used for further measurements as it does not suffer from the pseudo symmetry problem of the BaGe4O9 file. The three phase system in the glass-ceramic is thus reduced to the two phase system of BTG and benitoite/BaGe4O9 for further analysis.These results show that the orientation measurements are satisfactory and BTG may be reliably separated from benitoite/BaGeAs it turned out in the subsequent measurements, the grain CI-standardization clean-up procedure, which only changes CI-values but no crystal orientations, provided sufficient datasets for analysis while still using a CI-filter to remove unreliable data points originating e.g. from the residual glass.4O9. While these weak peaks are not a clear proof that these phases, also described for other BTG containing glass ceramics1416Polished samples of the glass 1 were annealed at temperatures from 690 to 970 °C in order to achieve comparability with the surface crystallization research performed on glass-ceramics containing BTG7891011121314151617181920212223403223789101112131415161718192021224O9 which cannot be reliably separated at this time as outlined above. Hence a complete layer of benitoite comparable to that proposed to occur at the surface of the glass BTG55 crystallized at 690 °C4O9. As both phases occur according to the XRD-measurements in 4O9 or both at this point.SEM-micrographs obtained from the surfaces of samples annealed at the stated temperatures are presented in A similar ring texture indicating a c-axis tilt of 58° at the surface was detected in the BTS equivalent of the current glass composition after crystallization at temperatures from 790 to 970 °C404O9. 4O9 or both cannot be made at this point. Due to the small grain size and the topography of the surface a validation of the phases using WDS is deemed unreliable. Larger crystals of the non-BTG phases could help solve this problem in the future.The common texture observed in 2 of the bulk crystallization in a sample annealed at 790 °C for 10 h showed multiple orientations over larger areas, i.e. bulk nucleation occurs in this system but with a low nucleation rate.Cross sections of these samples were prepared in order to analyze the crystal growth into the bulk. 1314161718224O9. The non-BTG areas basically occur parallel to the BTG growth structures, are less than 1 μm thick but frequently reach several μm in length. Assuming they are composed of benitoite the c-axes would be generally oriented parallel to the c-axes of the neighbouring BTG but with a lower degree of orientation and without an epitaxial relationship. Some data points attributed to benitoite/BaGe4O9 are observed more than 100 μm below the surface but they do not occur at all in scans covering the boundary between surface and bulk crystallization. Hence it appears that benitoite/BaGe4O9 crystallizes in minimal amounts during the initial growth of the oriented surface layer whenever the local stoichiometry is suitable.Only 1.1% of the reliably indexed data points in the surface near scan are attributed to benitoite/BaGe4O9 observed below the surface, the following measurements only feature results concerning BTG. To confirm the observed bulk nucleation and gain further insight into the microstructure of BTG in the oriented layer, part of the sample featured in Due to the minimal amounts of benitoite/BaGeWhile the contrast in the SEM-micrograph is very low, the IPF+IQ-map of the performed EBSD-scan shows that the BTG growth structures appear as lines when viewed from the side but as stars when viewed parallel to their growth direction. As the EBSD-scan was performed with a step size of 200 nm, the thermal footprints of the scan step overlap which can lead to increased temperatures in the scan area and even cause EBSD-pattern degradationThe method of EBSD-pattern degradation was applied to polished and unpolished surfaces using the same conditions stated in ref. These results are observed in all samples annealed at temperatures in the range from 750 to 970 °C for 10 to 20 h. Obvious differences between the crystallization below and above the phase transition temperature of BTG were not observed except for a decrease of the nucleation rate and the fact that higher annealing temperatures tend to enable a better EBSD-pattern acquisition from the untreated surface. Perhaps the long annealing times at high temperatures allow some of the twinning probably present in the system3g to be 674 °C and the onset of crystallization occurred at ca. 757°C. Samples were annealed at temperatures from 680 to 970 °C for times from 0.5–12 h and in a two-step process at 690 °C for 3 h with another 3 h at 720 °C in analogy to ref. The prepared glass was transparent and showed the same yellow coloration already observed in glass 1. A Differential scanning calorimetry (DSC)-measurement of the glass showed the glass transition temperature T3O9 (benitoite) observed in ref. 4O9. Annealing for only 0.5 h at 690 °C did not provide any indication of crystallization in the XRD patterns and a subsequent analysis of the sample surface using SEM/EBSD also failed to produce any indication of crystallization.The XRD-patterns presented in 4O9 were obtained from all surfaces. Despite the strong topography, EBSD-scans were performed so that {001} and {0001} textures could be calculated. Relevant information concerning the datasets used for these calculations are stated in EBSD-patterns reliably indexable as either BTG or benitoite/BaGe4O9 or both at this point. The Euler angle distributions presented below show that the c-axes of BTG are now predominantly tilted by 29° from the surface normal in contrast to the 35° observed for glass 1 in 4O9 are basically the same.In analogy to glass 1, annealing at higher temperatures increases the indexing rate significantly as well as the detected amount of BTG at the immediate surface. While the double texture observed in glass 1 occurs in almost all calculated textures, the phases in the sample annealed at 970 °C for 10 h show an additional orientation preference for crystals with their c-axes oriented parallel to the surface. This preference is also observed in the texture of BTG obtained from the sample crystallized via the two step process. As with glass 1, it is impossible to state whether these textures originate from benitoite, BaGe4O9 are detected by EBSD in the top layer while the bulk structures are composed of BTG as well as benitoite/BaGe4O9. Texture information of more representative scans covering a) 300 × 300 μm2 and b) 260 × 100 μm2 are presented to show that the top layer of mainly BTG is highly oriented with the c-axis perpendicular to the surface while no clear texture is observed in the bulk for either phase. Sample preparation is very challenging in these samples because the crystals are much finer than those observed in glass 1. Instead of the systematic pores at the growth front observed in Cross sections of various samples were prepared and analysed. 4O9 occur at the surface of almost all samples. Only crystallizing glass 2 at 680 °C for 1 h provided a one phase system of BTG according to the XRD-pattern, which is not very well suited to exclude the existence of a phase, instead of the benitoite detected in ref. The results presented above indicate that BTG and secondary phases such as benitoite and BaGe4O9 at 690 °C was observed by both groups, e.g. only 3.3% of the data points of the corresponding EBSD-scan in A complete layer of either phase at the surface is impossible in all other samples as EBSD-patterns reliably indexable as BTG were obtained from the surface of every annealed sample which is in agreement with ref. Due to the high nucleation rate, the high fraction of BTG in samples crystallized at 970 °C in both glasses and the sole occurrence of BTG in glass 2 at 680 °C, it seems acceptable to assume that BTG nucleated independently from the other phases in these experiments. This means that oriented nucleation is indicated for BTG as well as for at least one of the secondary phases. As oriented nucleation was also observed after the crystallization of two glasses in the STS system32A temperature dependent texture formation similar to that observed in the BTS systemIn both glasses, nucleation is followed by a very rapid growth selection so that ca. 10 μm below the surface only BTG-crystals with their c-axes perpendicular to the surface prevail while the secondary phases practically disappear. Hence the idea that the degree of orientation is affected by the surface roughness observed after crystallization4O9 play a role similar to that outlined in ref. As outlined above, it is basically impossible that benitoite and/or BaGeThe cracks and large number of pores observed in the BTG glass-ceramics, especially at the interface between surface and bulk crystallization in glass 1, increase the difficulty of achieving a high quality surface treatment. The pores at the growth front are very similar to the holes observed during the crystallization of glass fibers with the composition BTG55Due to the basically perfect c-axes orientation observed in the crystallized surface layer featured in the The uneven growth fronts observed in a number of articles and discussed in the context of nucleation sites13141618In summary the growth of BTG in both glasses featured here is quite similar to the growth of BTS but widely in contrast to STS in their respective versions of the glass1 composition. The novel observation of a multi-phase oriented nucleation including a triple texture further cement the fact that assuming statistically oriented nucleation in theoretical models concerning crystal nucleation in glasses is questionable. Both BTG and BTS show oriented nucleation with the c-axis tilted by 55 and 58 ±10° respectively for the equivalent compositions. In addition to the temperature dependence of oriented nucleation observed here and in the BTS2 · 2.75 GeO2 was melted from a mixture of the reagent grade raw materials BaCO3, TiO2 and GeO2 in a Pt-crucible at 1550 °C in an inductive furnace for 2 h and stirred for another 2 h at this temperature to homogenize the melt. The glass was then poured on a brass block, quenched for about 7 s with a brass stamp and subsequently transferred to a furnace preheated to a temperature of 760 °C (which is close to Tg). The glass 2 was prepared in the same inductive furnace using the same raw materials. It was melted at 1300 °C for 1 h and stirred for another 20 min. After quenching it was transferred to a furnace preheated to 660 °C. The respective furnaces were switched off allowing the glasses to cool with a rate of approximately 3 °C/min. Crystalline powders of BaTiGe3O9 and BaGe4O9 were prepared via solid state reactions using BaCO3 , TiO2  and GeO2  as raw materials. The stoichiometric mixture of BaTiGe3O9 was heat treated at 1230 °C while the stoichiometric mixture of BaGe4O9 was heat treated at 1050 °C where they were held for 20 h with three intermediate regrinding steps.Glass 1 with the composition 2 BaO · TiO−3 Pa.Cross sections and powders were embedded in Araldite CY212 for further treatment. Samples were manually polished with shrinking grain sizes down to 0.75 μm diamond suspension. A final finish of 30 min using colloidal silica was applied. Conductivity of the surface was achieved by mounting the sample using Ag-paste and applying a thin coating of carbon at about 10The glass 1 was analyzed by DTA according to Hartmann & Braun. The grain size of the powder was smaller than 70 μm and a heating rate of 10 °C/min was used. Glass 2 was analyzed by DSC using a Linseis DSC Pt 1600 using Pt/Rh10 crucibles and a heating rate of 10 K/min. X-ray diffraction (XRD) was performed using CuKα–radiation in a SIEMENS D5000 diffractometer. SEM analyses were performed using a Jeol JSM 7001F scanning electron microscope equipped with an EDAX Trident analyzing system containing a Digiview 3 EBSD-camera. EBSD-scans were performed using a voltage of 20 kV and a current of ca. 2.40 nA. The scans were captured and evaluated using the software TSL OIM Data Collection 5.31 and TSL OIM Analysis 6.2. All scans were cleaned using the “grain CI standardization” function. Afterwards, unreliable data points were removed by applying a CI filter of 0.1. No further cleanups which actually modify orientations were applied. Pole figures of textures are presented in multiples of a random distribution (MRD).How to cite this article: Wisniewski, W. et al. Oriented Nucleation of both Ge-Fresnoite and Benitoite/BaGe4O9 during the Surface Crystallisation of Glass Studied by Electron Backscatter Diffraction. Sci. Rep.6, 20125; doi: 10.1038/srep20125 (2016).


The diffusivity of Cr, Fe and O is enhanced concurrently in the nanocrystalline surface layer, which causes the fast oxidation in the initial oxidation stage. The formation of 2O3 inner layer would inhabit fast diffusion of alloy elements in the nanocrystalline surface layer of P91 steel in the later oxidation stage, and it causes a decrease in the parabolic oxidation rate compared with conventional specimens. This study provides a novel approach to improve the oxidation resistance of heat resistant steel without changing its Cr content.An interesting change of scale sequence occurred during oxidation of nanocrystalline surface layer by means of a surface mechanical attrition treatment. The three-layer oxide structure from the surface towards the matrix is Fe Due to long-term exposures to high temperatures, the high temperature creep resistance and high temperature oxidation resistance of P91 steel should be simultaneously improved. A Cr content of around 9 wt% in martensitic heat resistant steels is required for the optimization in creep properties, and creep resistances of 9Cr heat resistant steels could be further improved by the control of precipitation behaviors1The nanocrystalline structure with a large number of grain boundaries, which can act as fast atomic diffusion channels810The chemical composition of commercial P91 steel used in the present work is 0.1 C-0.6 Mn-8.8 Cr-0.95 Mo. A polished plate specimen (Ф50 mm ×4.0 mm in size) of a tempered P91 steel with conventional grain size of 14 μm was subjected to SMAT. SMAT was performed only on one side of the specimens. The set-up and procedure were described as follows: hardened 0.8 C steel balls, with 5 mm in diameter and mirror-like surfaces, were placed in a reflecting chamber that was vibrated by a vibration generator with a frequency of 20 kHz. After being treated for 30 min, the surface roughness of SMAT-ed sample is comparable to that of the original, unpolished specimen. Other details of the experimental setup and the SMAT processing are described in the reference3 were cut by an electrical discharge machine for the nanocrystalline specimens and conventional grain specimens. The surfaces of specimens, except for the SMAT treated surface, were polished to mirror-like surface by 2.5 μm Al2O3 polishing paste, then cleaned in distilled water and then in acetone with ultrasonic agitation for 15 min prior to oxidation. Oxidation behaviors were tested in supercritical water oxidation equipments operated at 848 K and 14.1 MPa, and the total time of oxidation was up to 48 days (1152 h). Each oxidation date at different time was calculated by averaging the value from three specimens. Oxide scale were characterized by SEM, EPMA, XRD and TEM. Cross-sectional morphologies of the specimens were observed by using SEM with back-scattered electron (BSE) and Electron probe micro-analyzer (EPMA). Microstructure characterizations of nanocrystalline microstructure were also examined using TEM. Thin foil specimens for TEM observation were prepared using Focused Ion Beam instrument (FIB).In order to make comparisons of the oxidation resistance between convertional materials and SMAT-ed nanocrystalline materials, oxidation specimens with 10 × 20 × 2 mmIt is reported that the grains in the surface layer region within 100 μm thick are refined into nanometer scale, and the thermal stability of nanocrystalline microstructure is excellent until the temperature above 1033 KW is the weight gain in mg/cm2, k is the oxidation rate constant in mg/cm2/h, and t is time in hour. The oxidation rate constant k is 0.044 for the conventional specimens. However, the weight gain data (24 h) of the nanocrystalline specimens could not fitted to the parabola or a power law in 3O4 and FeCr2O4, and this will lead to an acceleration of the oxidation rate in the early stage before the formation of a more protective Cr-rich 2O3 oxide scale as shown in 2O3 inner layer on the nanocrystalline surface layer of P91 steel as shown in k from 0.044 in the conversional specimen to 0.026 in the nanocrystalline specimen 2O3 inner layer at the oxide/alloy interface. This is different from the typical two-layer scale consisting of an Fe3O4 outer layer and an FeCr2O4 inner layer on the conventional surface of P91 steel133O4 and FeCr2O4 are similar in both the conventional and nanocrystalline specimens. However, the inner oxide scale 2O3 is richer in Cr, which can be clearly seen from the cross-sectional scale morphology in 3O4, FeCr2O4 and 2O3 were all detected. Further analysis was performed using TEM to verify the morphology and phase structure of different types of oxides. The nanocrystalline specimen exposed after 576 h was selected to prepare the TEM foils because of the processing depth limitation in the FIB instrument. The selected area electron diffraction pattern (SAED) of TEM indicated that the oxide contained the spinel FeCr2O4 and the corundum 2O3 in 2O4 to 31.4 wt% in 2O3. The map-scan profiles also proved that the inner oxide scale 2O3 is richer in Cr er in Cr .2O3 scale in 9 Cr steel to form at 848 K. The scales of 10 Cr steel typically consist of an outer Fe3O4 layer and an inner layer containing 3O4 at 923 K in Ar + 50%H2O2O3, Fe3O4, and 3O4 is formed on P91 steel at 923 K in steam atmospheres, and it is still composed of Fe3O4 and 3O4 at 1073 K2O3 layer. However, the scale is composed mainly of 2O3 on AISI 430 steel with 16.21 wt% Cr content at 923 K, and a layer of Cr2O3 is formed at 1073 K owing to the higher diffusion rate of Cr at this temperature than at 923 K2O3 layer could form on the nanocrystalline surface layer in P91 steel with 8.8 wt% Cr at 848 K.Cr diffusion in the metal matrix and oxide is not fast enough for a stable corundum 2O3 layer. Before the formation of the 2O3 layer, the diffusivity of oxygen and iron sharply increase on the nanocrystalline surface layer of P91 steel. So oxidation rates of the nanocrystalline specimens are significantly accelerated in the initial stage. Increase of the oxidation rate with decreasing grain size is attributed to the larger grain-boundary area, which result in an increase in the short circuit diffusion paths2O3 layer, due to its higher density of grain boundaries in the nanocrystalline surface layer, a higher flux of Cr goes towards the 2O3-the alloy interface, while the spinel 2O3 and FeCr2O4 layer inhabits Cr and O/Fe diffusion through it. Therefore, the 2O3 scale at the alloy-oxide interface has a higher concentration of Cr. O and Fe diffusivity would also decrease due to the formation of 2O3 inner layer in the nanocrystalline surface layer of P91 steel. The enhancement of Cr diffusion guarantee the stable growth of the Cr-rich 2O3 scale, which significantly improve oxidation resistance. Even a Cr2O3 scale may be formed at higher temperatures owing to the higher diffusion rate of Cr.The oxidation process of the nanocrystalline surface layer is divided into two stages, namely before and after the formation of a protective How to cite this article: Xia, Z. X. et al. Improve oxidation resistance at high temperature by nanocrystalline surface layer. Sci. Rep.5, 13027; doi: 10.1038/srep13027 (2015).


The films were anatase TiO2, with good n-type electrical conductivities as determined via Hall effect measurements. The film doped with 2.25 at.% W showed the lowest resistivity at 0.034 Ω.cm and respectable charge carrier mobility (14.9 cm3/V.s) and concentration (×1019 cm−3). XPS indicated the presence of both W6+ and W4+ in the TiO2 matrix, with the substitutional doping of W4+ inducing an expansion of the anatase unit cell as determined by XRD. The films also showed good photocatalytic activity under UV-light illumination, with degradation of resazurin redox dye at a higher rate than with undoped TiO2.Tungsten doped titanium dioxide films with both transparent conducting oxide (TCO) and photocatalytic properties were produced  An undoped TiO2 film was made from a toluene (25 mL) solution of [Ti(OEt)4], while for the W doped TiO2 films dopant concentrations (2–20 mol.% in solution) of [W(OEt)6] were used. This is to our knowledge the first example of W doped TiO2 being formed by AACVD. All films were well adhered to the substrate passing the ScotchTM tape test. The undoped films were optically transparent and colourless where as the doped films were translucent blue, with the colour intensity increasing with increasing dopant concentration. The films were air stable, showing no change in colour or conductivity after 3 months.Thin films of undoped and W doped TiO2 (2 films deposited using both AACVD and PVD9The XRD patterns for the undoped and W doped films show reflections matching the anatase phase of TiO2 . The und6+ (0.60 Å) and Ti4+ (0.605 Å), the substitutional doping of W6+ with Ti4+ should impart almost no change in the TiO2 unit cell. However, XPS results (see below) show that doped films contain W in both the 4+ and 6+ oxidation state. Therefore the observed linear shift of the XRD peaks to lower 2θ values, thus indicating an expansion in the TiO2 unit cell, was due to the partial replacement of Ti with W4+ that has an ionic radius of 0.66 Å. This follows closely with observations made by Kafizas and Parkin for W doped TiO2 system studied via combinatorial atmospheric pressure CVD where the majority of the W was in the 5+ oxidation state (0.62 Å)Due to the similarities in the ionic radii of W2 films showed peaks expected for TiO2 anatase with tetragonal symmetry  was shifted to higher wavenumbers indicating an expansion of the unit cell. The shift in the Eg peak increased linearly with increasing W concentration and is consistent with formation of a solid solution as observed by the XRD data.Raman patterns taken for the undoped and W doped TiOsymmetry 20. Upon Wavelength dispersive X-ray spectroscopy was used to determine the amount of W present in the films. The results show a linear increase in W content before possible saturation is reached at 15 mol.% of W in the AACVD precursor solution and 4.5 at.% in the films .2 film consists of compact domes roughly 100 nm in width that transform to pyramidal features perpendicular to the substrate upon W incorporation. With increasing dopant concentration the features coalesce to form larger pyramids that are roughly 500 nm in width. Furthermore, the 4.47 W at.% and 4.65 W at.% films seem to show the features coalescing even further. This has been observed previously when other transition metals such as Nb and Ta have been substituted into a TiO2 matrix9via side-on SEM, show that in general there is an increase in the thickness with increasing amount of dopant. This coupled with the increase in film roughness contributes to a decrease in film transparency.Analysis of the film morphology using scanning electron microscopy (SEM) reveal that the surface of the AACVD grown films become more structured upon doping with W . The mor3/2 peak corresponding to Ti in the 4+ oxidation state was observed at 458.3 eV that matches to literature values212 samples also showed this peak at 458.2 eV for Ti4+, but an additional 2p3/2 peak at 457.0 eV corresponding to Ti3+ was also observed9133+ peak increased with increasing W concentration. Thus indicating that the charge compensation by electrons for the substitutional doping of W 4+ and/or 6+ species for Ti4+ results in the reduction of the titanium to the 3+ oxidation state.X-ray photoelectron spectroscopy (XPS) was performed on the films to determine the oxidation state of the Ti and W species . For the7/2 peaks corresponding to W6+ was seen at 35.2 eV as well as 4f7/2 peaks for W4+ at 33.8 eV246+ compared to W4+ in all the films. In a previous study where Kafizas and Parkin studied W doping of TiO2via a combinatorial APCVD method, XPS studies showed the presence of mainly W5+ and some W4+, lacking the desired 6+ oxidation that enables enhanced electrical propertiesPeaks corresponding to W 4f overlap with Ti 3s peaks in XPS, therefore W was only detectable for the 2.25 at.% and above doped films . In thesn-type conductivity (−1 and 3.4 kΩ respectively. For the 0.63 at.% doped sample the resistivity was 0.63 Ω.cm and the charge carrier concentration was to the order of 1017 cm−3 enabling a carrier mobility of 13 cm2/V.s. The resistivity (0.29 Ω.cm), charge carrier concentration (1018 cm3) and mobility (13 cm2/V.s) were also similar for the 1.65 at.% W doped film. A further increase in dopant concentration to 2.25 at.% resulted in the film with the best electrical properties for this study. There was a decrease in sheet resistance and resistivity to 212 Ω.sq−1 and 0.034 Ω.cm respectively. This is comparable to the results obtained by Xu and Jin et al. for W doped TiO2 films grown by rf sputtering. The carrier mobility, at 14.9 cm2/V.s, was slightly better than the 0.63 at.% and 1.65 at.% films while the carrier concentration was an order of magnitude higher. At 4.47 at.% doping of W in TiO2, the sheet resistance was reduced to 78 Ω.sq−1 with a resistivity of 0.042 Ω.cm and an increased carrier concentration in the order of 1020 cm−3. However, the carrier mobility was diminished to 0.32 cm2/V.s  light source with a photon flux of 1.09 × 1015 photons/cm2/s. The reduction in the concentration of the dye with time was monitored using UV-Vis spectroscopy. The formal quantum efficiency (FQE) and formal quantum yield (FQY) were calculated to enable easy comparison with literature.The UV light induced photocatalytic properties of the undoped and W doped TiO2 film was able to degrade resasurin at a rate of 1.4 × 1011 dye molecules/s/cm2, this was inferior to the 0.63 at.%, 1.65 at.% and 2.25 at.% W doped TiO2 films that had a degradation rate of 3.2 × 1011, 3.5 × 1011 and 2.4 × 1011 dye molecules/s/cm2 respectively. However at 4.47 at.% and 4.65 at.% doping levels the rates of degradation (7.9 × 1010 dye molecules/s/cm2 for both) was inferior to the undoped TiO2 film. Furthermore, compared to SGG BiocleanTM self-cleaning glass, that was able to degrade resazurin dye at 1.2 × 1011 dye molecules/s/cm2 all the films excluding the 4.47 at.% and 4.65 at.% doped films were superior.The undoped TiO2 films does indeed enhance the photocatalytic properties but excess dopants can also be detrimental. The initial increase in the photocatalytic properties of the films with doping was most likely due to the increase in surface microstructure  due to larger, better shaped particles as observed from the SEM images. This facilitates a high surface area and more photoactive sites for the dye to be degraded during UVA illumination. With excess dopant concentrations (4.47 at% and 4.65 at.%) an increase in the surface area is off set by the presence of more photoinduced charge carrier recombination sites that in total reduce the photoactivity.This shows that the presence of W in the TiO−4 dye molecules per incident photon and FQY of 3.26 × 10−4 dye molecules per absorbed photon for the 0.63 at.% doped film and 2.99 × 10−4 dye molecules per incident photon (FQE) and 3.14 × 10−4 dye molecules per absorbed photon (FQY) for the 1.65 at.% doped film. This was more than double what was observed for the undoped TiO2 film grown under the same conditions. The principle reason for this was the more structured surface morphology observed for the doped films compared to the undoped. Furthermore, the FQEs of the 0.63 at.% and 1.65 at.% W doped films were also both more than double what was observed (1.2 × 10−4 dye molecules per incident photon) for SGG BiocleanTM. The FQY of SGG BiocleanTM was however far superior with results an order of magnitude higher than the best preforming W doped TiO2 film.The formal quantum efficiency (FQE) and yield (FQY) for the films were calculated using the dye degradation rates, the UVA photon flux and UVA photon absorption for each of the films . The 0.62 films studied by combinatorial APCVDTM self-cleaning glass by Mills et al. shows a FQE of 2.5 × 10−4 dye molecules/incident photon2 samples in this study, indicating the films are on par with industry standards.Compared to literature, the W doped films display the same order of photocatalytic activity (FQE and FQY) relative to W doped TiO2 films were synthesisied via a solution based CVD technique (AACVD). The structured morphology of the doped films combined with the good conductivites (optimum of 0.034 Ω.cm) and optical transparency in the visible region make them suitable candidates for electrodes in photovoltaic devices. Furthermore, the doped films displayed enhanced photocatalytic activity compared to the undoped TiO2 sample in the destruction of resazurin redox dye when illuminated by UVA radiation. In conclusion, this paper presents the single step formation of W doped TiO2 films with four desirable functional properties:Transparent in the visible regionPhotocatalytically activeColouredElectrically conductingTransparent, electrically conductive and photocatalytically active thin films of W doped TiODepositions were carried out under nitrogen (99.99% from BOC). Precursors were placed in a glass bubbler and an aerosol mist was created using a piezoelectric device. All chemicals were procured from Aldrich and were utilised as recieved.4]  was dissolved in toluene (25 ml) and [W(OEt)5] was added in dopant amounts (0 mol.% – 20 mol.%). The resulting solution was stirred for 30 minutes and then atomised. The precursor flow was kept at 1 L.min−1. The substrate temperature was kept at 500 °C. Deposition time was 45 minutes. After the deposition the bubblers were closed and the substrates were cooled under a flow of nitrogen.[Ti(OEt)ca. 1 cm × 1 cm squares for subsequent analysis.At the end of the deposition the nitrogen flow through the aerosol was diverted and only nitrogen passed over the substrate. The glass substrate was allowed to cool with the graphite block to less than 100 °C before it was removed. Coated substrates were handled and stored in air. The coated glass substrate was cut into 1 and CuKα2 radiation of wavelengths 1.54056 Å and 1.54439 Å respectively, emitted with an intensity ratio of 2:1, a voltage of 40 kV and current of 30 mA. The incident beam angle was kept at 1° and the angular range of the patterns collected was 20° < 2θ < 66° with a step size of 0.05° counted at 0.5 s/sep.Powder X-ray diffraction (PXRD) was used to analyse the samples in a modified Bruker-Axs D8 diffractometer with parallel beam optics equipped with a PSD LynxEye silicon strip detector to collect diffracted X-ray photons. This instrument uses a Cu source for X-ray generation with CuKαScanning Electron Microscopy (SEM) was performed to determine surface morphology and film thickness using a JEOL JSM-6301F Field Emission SEM at an accelerating voltage of 5 kV.2) of each position analysed was taken.The W:Ti at.% composition was investigated via wavelength dispersive X-ray (WDX) analysis on a Phillips instrument. The average W:Ti % over four area scans (200 mmα radiation. Survey scans were collected in the range 0–1100 eV (binding energy) at a pass energy of 160 eV. Higher resolution scans were recorded for the principal peaks of Ti (2p), W (4f), O (1s), C (1s) and Si (2p) at a pass energy of 50 eV. Peak positions were calibrated to carbon (284.5 eV) and plotted using the CasaXPS software.X-ray photoelectron spectroscopy (XPS) was performed using a Thermo Scientific K-alpha photoelectron spectrometer using monochromatic Al-KUV/Visible/near IR spectra were taken using a Perkin Elmer Fourier transform Lambda 950 UV/Vis spectrometer over a wavelength range of 300 nm to 2500 nm in transmission mode. The transmission spectra were taken against an air background.2 samples were then subjected to an input current of 1 mA and a calibrated magnetic field of 0.58 T. The transverse voltage was then measured. The measurement was repeated by reversing the direction of the magnetic field and the current.Hall effect measurements were carried out using the van Der Pauw method to determine the sheet resistance, free carrier concentration (N) and mobility (μ). A square array of ohmic contacts arranged on 1 cmet al.). The intelligent ink consisted of resazurin (40 mg) redox dye in an aqueous solution (40 mL) with glycerol (3 g) and hydroxyl-ethyl cellulose (0.45 g). The mixture was aged for 24 hours in a fridge (2–5 °C) and thoroughly mixed before use.Prior to the photocatalysis study, the samples were washed with distilled water, rinsed in isopropanol and irradiated for 30 minutes with UVA light to clean and activate the surface. The photocatalysis test involved spray coating the samples with an even layer of the resazurin based ‘intelligent ink’  was used to induce the photoreduction of the resazurin redox dye on the surface of the films and a UV-visible spectrometer was employed to monitor the degradation of the dye. Formal quantum efficiency (FQE) was calculated by dividing the rate of dye molecules destroyed (cm−2 s−1) by the photon flux (1.01 × 1014 photons cm−2 s−1). The formal quantum yield (FQY) was calculated by dividing the rate of dye molecules destroyed (cm−2 s−1) by the number of photons absorbed (cm−2 s−1) by the film. The photon flux and photon absorption for each film was determined using a UVX radiometer with a detector for 365 nm radiation attached.UVA light .


By epitaxially growing twined 2O3 thin films, of which the c axis and spins of the Cr atoms lie in the film plane, we demonstrate that the exchange bias between Cr2O3 and an adjacent permalloy layer is tuned to in-plane from out-of-plane that has been observed in 2O3 films. This is owing to the collinear exchange coupling between the spins of the Cr atoms and the adjacent FM layer. Such a highly anisotropic exchange bias phenomenon is not possible in polycrystalline films.Exchange bias is one of the most extensively studied phenomena in magnetism, since it exerts a unidirectional anisotropy to a ferromagnet (FM) when coupled to an antiferromagnet (AFM) and the control of the exchange bias is therefore very important for technological applications, such as magnetic random access memory and giant magnetoresistance sensors. In this letter, we report the crystal structure manipulation of the exchange bias in epitaxial hcp Cr The perpendicular exchange bias was previously shown in (0001) oriented Cr2O3 films234715In this letter, we report the manipulation of the exchange bias by controlling the surface spin configuration via crystal orientation design. By epitaxially growing lms on 00 oriented2O3 films are grown on the 2 substrates via laser molecular beam epitaxy (LMBE) with a base pressure of 2 × 10−8 mbar (see methods for details). in-situ reflection high-energy electron diffraction (RHEED) characterization of the 2 substrate’s surface viewed from the 2O3, the RHEED pattern of the TiO2 substrate disappears and that of Cr2O3 starts to appear, as shown in 2O3 film increases, its RHEED pattern becomes brighter. 2O3 films, respectively. To be noted, four satellite RHEED spots are observed around each main diffraction spot viewed from the 2 substrates (The bstrates .2O3 films are further characterized by x-ray diffraction (XRD). The θ–2θ scans of the rutile TiO2 substrate, 10 nm, 20 nm, and 27 nm Cr2O3 films are shown in 2 substrates. For the 10 nm Cr2O3 film, a peak at 2θ of ~65 degrees is observed, which corresponds to the 2O3 crystal. As the thickness of Cr2O3 increases, the intensity of the peak at ~65 degrees becomes stronger. We note that for the 27 nm Cr2O3 film, only 2O3 thin films. The surface morphology is characterized by atomic force microscopy (AFM). The root-mean-square (RMS) roughness is 0.09 nm for the rutile TiO2 substrate after annealing in the chamber , we deposit 10 nm Py on top of the 2O3 films and measure the magnetic hysteresis loops by Magnetic Properties Measurement System  with both in-plane and out-of-plane magnetic fields at various temperatures  is defined by the mean value of the H1 and H2, i.e. Interestingly, with the ubstrate  from 10 2O3 films. The highly anisotropic exchange bias phenomenon can be attributed to the crystalline orientation difference of the Cr2O3 films. In hcp structures, the c-axis direction dictates the spin orientations of the magnetic atoms. In the 2O3 films, the spin orientations of the Cr atoms are perpendicular to the films, whereas in the 2O3 films, the spins of the Cr atoms lie in the film plane. These results further indicate the direct collinear exchange coupling between the spins of the Cr atoms and the adjacent Py layer.To characterize the anisotropy of the exchange bias effect, magnetic hysteresis loops are measured with a magnetic field perpendicular to the films . The samHC), where 2O3 films and 10 nm Py as a function of the temperature. As the temperature increases from 10 to 60 K, the in-plane exchange bias field (TB) of ~60 K. An abrupt increase in Hc below 60 K is another property of exchange biased Py, which is due to the formation of the AFM order in this 13 nm Cr2O3 thin film2O3, as indicated in zones 2TB of 13 nm Cr2O3 is ~60 K, which is much lower compared to the value reported on (0001) oriented bulk Cr2O3 single crystalsTB is highly related to Neel temperature (TN), and is usually slightly lower than the TN. Both TB and TN increase as the AFM thickness increases due to finite-size effects21223TB as a function of the thicknesses of the Cr2O3 thin films, the in-plane exchange bias for the samples consisting of 7, 10, 20 and 27 nm Cr2O3 films and 10 nm Py bilayer films are also measured. 2O3 (7 nm)/Py (10 nm) and Cr2O3 (27 nm)/Py (10 nm), respectively. The blocking temperatures of these two structures are determined to be 40 K and 100 K. The blocking temperature increases as the thickness of the Cr2O3 films increases, as shown in 2O3 film, the blocking temperature is only ~100 K, which is far below the TB of bulk Cr2O3. One possible reason is the non-trivial finite size effects arising from the grain boundaries or oxygen defects in the Cr2O3The measured 2O3 thin films by controlling the surface spin orientations of the Cr atoms via crystal orientation design. For the epitaxial growth of 2O3 films, the spin configurations of Cr atoms give rise to only in-plane exchange bias at the interface between Py and the Cr2O3 thin films, while no perpendicular exchange bias is observed. Our results along with previous studies on (0001) oriented Cr2O3 films indicate the collinear exchange coupling between the spins of the Cr atoms and the adjacent FM layer.In summary, we have demonstrated the manipulation of the exchange bias of the Cr2O3 films are grown on the (001) oriented rutile TiO2 substrates via laser molecular beam epitaxy (LMBE) with a base pressure of 2 × 10−8 mbar. Prior to the Cr2O3 growth, the substrate temperature is increased to 350 °C with a rate of 20 °C/min in the chamber with an oxygen partial pressure of 0.08 mbar. Then the Cr2O3 film is deposited from a Cr2O3 target with a laser power of (8.0 ± 0.2) mJ and a frequency of 2.0 Hz. The thickness of the Cr2O3 thin film (t) is determined from the cross section high resolution transmission electron microscopy.The 2O3 films by RF magnetron sputtering with a growth rate of 0.02 Å/s. A capping layer of 20 nm aluminum is deposited prior to taking the samples out of this sputtering chamber to prevent oxidization of Py. Magnetic Properties Measurement System  is used to measure the magnetic hysteresis loops to determine the exchange bias.A 10 nm Py is grown on top of the How to cite this article: Yuan, W. et al. Crystal Structure Manipulation of the Exchange Bias in an Antiferromagnetic Film. Sci. Rep.6, 28397; doi: 10.1038/srep28397 (2016).


A Solomon link, colloquially termed a “Solomon knot” (a  notation), is a t notation Such dou notation–5 with o notation,5 to dat notation using Sa notation of gener notation,d cyclic notation can prov notation  and investigated its coordination chemistry with primary amines and FeII salts  revealed that 2 was a metal–ligand tetramer with the formula [Fe4L4](PF6)8][1 with two molecules of n-hexylamine). 1H NMR spectroscopy  helicate topology shown in Scheme 2 was increased to 71 % (yield of isolated product) when employing 4.4 equivalents of the iron(II) salt .The reaction of 4](PF6)8]  salt (Scheme 4)2 and Fe(ClO4)2 also produced 2, although in significantly lower yields  and contaminated with other polymeric and oligomeric by-products. When FeBr2 was employed as the iron source, a different main product was obtained (Scheme 3L3]6+) 3 by 1H NMR spectroscopy  5  circular helicate 4 after 24 hours, whereas starting with a concentration of 0.55 mm of 1, the reaction produced more than 85 % of the lower nuclearity  linear helicate 5 over the same time period .[n-hexylamine-derived cyclic tetramer 2 was essentially invariant over this concentration range and no linear triple helicate was observed, illustrating the influence that subtle changes in the ligands can have over the outcomes of the self-assembly reactions.Substituting s Figure , identifure S15). In contr1 with the diamine and FeCl2 in DMSO for 24 hours, with subsequent anion exchange with aqueous KPF6, generated the Solomon link 6 in 75 % yield of isolated product bis(ethylamine), a diamine that is stereoelectronically predisposed to adopt low-energy turns. The reac1H NMR spectrum  confirmed that 6 had a structural formula consistent with a Solomon link. Single crystals of 6 suitable for X-ray crystallography were grown by slow diffusion of diethyl ether into a nitromethane solution of 6, and the structure was confirmed by X-ray crystallography . Two PF6− counter ions are positioned directly above and below the center of the helicate  dications .The e Figure , includi Database  cations, four bis and four bis(amine) building blocks to generate two interwoven 68-membered-ring macrocycles with four crossings in 75 % isolated yield. The assembly process for the tetrameric cyclic double helicate forms the basis for the Solomon link synthesis and is sensitive to structural changes in the amine, the concentration and the anion used (even though the reaction product is not the result of an anion-template mechanism). The synthesis of Solomon link 6 and the earlier pentafoil knot[The one-pot synthesis of molecular Solomon link foil knot show tha


The ability to noninvasive image through turbid media has long been a major scientific and technological goal in many disciplines. A breakthrough has been made to observe objects that were completely hidden behind an opaque scattering layer. However, such approach needs not only to scan both illumination light and detector but further off-line procedures to numerically retrieve the image of the objects. Here, we report a distant invisibility-based noninvasive method that can hide scattering layers and allows to directly image objects behind. By recording holograms of the objects through a ground glass and then using the holograms produced time-reversal lights to re-illuminate the objects, we implemented to observe objects with feature size ranging from 39 μm to 80 μm that were hidden behind a 3 mm thick ground glass. Of importance, our approach opens a door towards real-time, high speed biomedical imaging and in-site inspection of integrated devices. Imaging through turbid media is an essential observation and diagnostic tool in many fields such as astronomical observations, microscopic imaging in tissues and defects testing of microelectronic devices. Great progress has been made by using methods such as ghost imaging1346789101112131415161718192021123456791012131416171819202123678910111516171819201213123456789101112131415161720212223Here, we describe a noninvasive and robust optical approach to directly observe amplitude–only modulated objects with feature size much smaller than 100 μm (39 μm – 80 μm) deeply behind a turbid medium several millimeters (3 mm in our case) based upon the principle of complementary medium-based distant invisibility272829l, no absorption is assumed), reflection angular spectrum (AS) of the object through the turbid layer in front can be read as1H(l) and 2H(l) denote the transfer functions (TFs) of turbid medium as illumination light and reflection light of object passing through it, respectively), U is the complex amplitude of the object, xf and yf are the spatial frequencies of the object in the x axis and y axis, respectively. We see that the information of turbid medium is involved in the reflection signal of object and hence may make the object behind the turbid medium invisible . From this formula we can find that the inverse Fourier transform of above transmission light holds an angular spectrumU of object is a real number and correspondingly, AS of the object is Hermitian κ is a real coefficient. We see that the reflection signal of the object no longer includes scattering information of the turbid medium. As a result, by inverse Fourier transform .In methods for optically producing time-reversed light, both linear10x = 675 μm is shown in the right hand panel of the figure, where the notches and peaks correspond to the stripes and spaces between the stripes of the resolution target, respectively. The turbid medium is a piece of 3 mm thick acidized ground glass  shows that the intensity notches and peaks corresponding to the stripes and spaces between the stripes of the resolution target, respectively, are clearly discriminated.x = 675 μm. We can see that all the stripes are still clearly visible. When the transversal deviation is larger than −5 μm away from its original position in the In holographic conjugation technology, deviations of the re-placed hologram from its original position and orientation etc. in the recording process may seriously destroy the function of the technology. We then check the effect of position and orientation deviations on the image quality in our approach.  to blur .position .For the effect of orientation deviation of hologram on image quality, we find that a little deviation from its original orientation will seriously destroy the images of the resolution target. Our experimental results show that when orientation deviation of the hologram is as small as 0.1° away from its original orientation around the Further, we check the reliability of our method in imaging objects with a little more complicated structure, such as the numbers “1, 2, 3, 4, 5, 6” with 80 μm linewidth . The holAll experimental results –6 agree It should be pointed out that the present method is based on the assumption that the objected we want to observe is amplitude-only modulated while the scattering layer has to produce a phase-only modulation. Hence our method is limited to focus on some special fields such as the observations of cancer cells in biological tissues and bright stars through atmospheric turbulence, where the ill cells and stars are generally treated as absorbing objects or luminous objects with amplitude-only modulation1234567891011121314151617181920212223589101145On the other hand, the scattering layer should remain unchanged when the hologram is calculated or fabricated as well as the measurement is implemented, which is a quite strong requirement for the technique aiming to noninvasive, and hence to in vivo imaging. To do so, one has to create a conjugation light immediately for imaging the objects behind a scattering layer, which greatly challenges the current materials and technologies. For example, Fe-doped LiNbO3 crystal, one of the typical photorefractive materials for widely using in real-time holography, shows a response time of about 2 min, and has been used to produce optical phase conjugation for turbidity suppression in biological samples21To conclude, we described both theoretically and experimentally a noninvasive imaging method to directly observe amplitude-only modulated objects behind a turbid medium by using holography-based time-reversal principle. Compared to the recent breakthrough made in noninvasive imaging approach to observe a fluorescent object that was completely hidden behind an opaque scattering layer252. Recording medium is a silver halide holographic recording plate . To record the holograms, we put the holographic recording plates in the “hologram” plane for exposure  for 25 minutes at 20°C. To evaluate the performance of holograms for hiding scattering in imaging objects through a ground glass, a step motor with 2.5 μm/step is used to control the movement and orientation of the holograms. A CCD camera is used to record the image and monitor the dynamics of the objects through the ground glass in front .Choosing the object .Calculating the object wavefront after a turbid medium (described as a random phase matrix of 512 × 512) by Fast Fourier Transform (FFT) algorithm.Encoding the calculated wavefront and reference light into a computer generated hologram (with 512 × 512 pixels) by Fresnel diffraction algorithm (containing a FFT).Optical decoding of the transmittance of hologram by using a reconstruction light (conjugated to the reference light) through Fresnel diffraction.Optical reconstruction of the image behind turbid medium by inverse FFT algorithm. The simulated results are shown in K.D.W. performed the theoretical analysis and experiments, Q.L.C. performed the experiments and simulations, Y.L.S. assisted the experiments, H.W. and G.P.W. designed and conducted the experiments, G.P.W. conceived the idea and supervised the project. All authors contributed to the final version of the manuscript.Supplementary Video S1Supplementary Video S2Supplementary Video S3Supplementary Video S4Supplementary Video S5Supplementary Video S6SUPPLEMENTARY INFO


Through spectroelectrochemical analysis of a range of non‐aqueous solvents, we describe the dependence of these processes on the electrolyte solvent and subsequent interactions formed between Na+ and O2−. The solvents ability to form and remove [Na+‐O2−]ads based on Gutmann donor number influences the final discharge product and mechanism of the cell. Utilizing surface‐enhanced Raman spectroscopy and electrochemical techniques, we demonstrate an analysis of the response of Na‐O2 cell chemistry with sulfoxide, amide, ether, and nitrile electrolyte solvents.The reduction of dioxygen in the presence of sodium cations can be tuned to give either sodium superoxide or sodium peroxide discharge products at the electrode surface. Control of the mechanistic direction of these processes may enhance the ability to tailor the energy density of sodium–oxygen batteries (NaO The possible reduced oxygen discharge products for M‐O2 electrochemistry include superoxide (MO2), peroxide (M2O2), and oxide  species.2,2 is Li2O2.2 batteries, both Na2O2⋅2 H2O222 batteries.Intensive research into lithium–oxygen Li‐O batterie2,2 electrolytes have thus far been limited to carbonate, ether, and the ionic liquid (IL) N‐methyl‐N‐propylpiperidinium bis(trifluorosulfonyl) imide (PP13TFSI)‐based electrolytes.2.2O2 as the discharge product, however sodium carbonate and carboxylates have now been detected more recently, and thought to be the major discharge product.2 cells cycled in organic carbonates as the main solvent.2, Na2O2, and Na2O2⋅H2O products.2 battery enquiries is the inducement of a solution‐based mechanism to enhance discharge capacity.Although well documented for Li‐O2O22.2 has established that the electrolyte acidic proton content promotes the formation of HO2 radicals as a phase transfer catalyst.2 enables the removal of superoxide from the surface above 5 ppm H2O/H+ content and a subsequent solution metathesis reaction with Na+ creates NaO2 nuclei that precipitate on the electrode surface. Similarly, a recent study by Jirkovsky et al.2, which promotes the formation of Li2O2 and the resulting partial dissociation of H2O to HO2 and OH−.2, along with phenol and ethanol in Li‐O2, providing additional evidence for this phenomenon.Solvation of superoxide through controlled water content increases discharge product size, including toroidal Li2 may undergo a second electron addition or a chemical disproportionation reaction to form Li2O2. These observations were qualitatively compared to Guttmann donor number, whereby high donor number solvents, including dimethylsulfoxide (DMSO), are able to better support a solution‐based mechanism enhancing battery capacity.In very dry electrolytes (≤10 ppm), control of the Lewis acidity of the alkali metal cation through electrolyte solvation allows alkali‐metal–superoxide ion pairs to form and react within the double layer to enhance discharge capacity.xOy species on planar roughened Au electrodes. Upon discharge, the high donor number solvents, DMSO and DMA, produced signals in the region for O2− and NaO2  and 1156 cm−1 , denoting varying interactions of O2− and cation between different solvents.A detailed electrochemical study is presented within the Supporting Information . Variations in the CV response were observed that were dependent both on the electrode substrate and the solvent; however, limited mechanistic insights could be directly acquired. SERS data provides insight into surface species and intermediates. By applying this technique to each electrolyte system, it is apparent that solvent choice can strongly affect the identity of Na+, as shown in Equations Within these systems, superoxide initially forms in the presence of Na2 leads to the formation of superoxide that subsequently interacts with Na+ at the interface. This interaction is heavily dependent on the solvation of both O2− at the surface and the solvation of Na+. These interactions control the acidity and basicity of the Na+ and O2−. The solvation shell in these cases will consist of anions (O2− and OTf−), cation (Na+), and solvent molecules.− presence is based on the solubility of the salt within each solvent, which induces the formation of contact, solvent‐separated, and free ion pairs in solution.−1 (Table S5) splits at a potential of 2.0 V versus Na+/Na denoting ion pair formation, but does not observably affect the NaxOy peaks within the spectra.The addition of an electron to the surface‐adsorbed O+‐O2−] is highly favorable, allowing the interaction to be ascribed to an ion pair, which corresponds to the detection of the band at 1107 cm−1. This soluble species is easily removed from the surface, which explains the quasi‐reversible nature of O2 and Na+ electrochemistry in DMSO. Furthermore, the detection of the signal for NaO2 at 1156 cm−1 is likely due to the aggregation and precipitation of NaO2 on the surface as the reductive potential increases to 2.0 V versus Na+/Na. Multiple CV scans within 0.1 m NaOTf showed that the quasi‐reversible process breaks down, and revealed the formation of two oxidation peaks within the initial cyclovoltammetric peak . This corroborated the initial ion pair formation, and the subsequent aggregation and precipitation of NaO2 on the surface. As the number of CV sweeps increases, more time has been allowed for NaO2 to precipitate, which leads to the growth of the second oxidation peak at 2.75 V versus Na+/Na.The interaction between DMSO‐solvated [Na+‐O2−] to NaO2 (1119 cm−1 to 1161 cm−1) within the spectra. The mechanism here is considered to be as stated in Equations (1) and (2). This follows the removal of NaO2 from the surface, subsequent electrolytic saturation within the double layer, and aggregation and precipitation of NaO2 , DEGDME‐based electrolyte required that the Raman analysis be carried out at 1 m NaOTf in order to reduce significant overpotentials . The assigned Na2O2 signals here were shifted from our Raman standard. This is a similar case for DEGDME . In the same spectral region as the higher Na2O2 peak in the SERS data, there was a corresponding OTf− anion peak at 760 cm−1 at OCV, however there are no peaks that appear around 710 cm−1. The OTf− band does not change in intensity under potential control, and so therefore this feature arises from Na2O2 due to O2 species. The presence of interfacial OTf− can be explained by the formation of ion pairs or aggregation of ion pairs at the interface. The doublet feature assigned to Na2O2 is shifted to a lower wavenumber than expected. If this shift is due to H2O, then it would be expected to increase the signal to a higher wavenumber.2O2 and the OTf− anion, which could of explain the observation of the red‐shifted bands (by ca. 30 cm−1) of Na2O2.To confirm that the shifted bands assigned as Na−1. This region denotes the νas SO3, which is considered the group within the anion that interacts with alkali metal cations.+ within Na2O2 at the surface may cause this observed shift. These bands relating to ion‐pair formation are detected in DMSO and DEGDME, suggesting that ion solvation may influence the discharge product.The formation of ion pairs was revealed by the appearance of a peak at 1040 cm2− in low donor number solvents increases the proximity of O2−, allowing for a second electron reduction to form a thin passivating film of Na2O2 [Eq. Therefore, the SERS data provides spectroscopic evidence that the lack of solvation of Oa2O2 Eq.  , which is removed from the surface and then aggregates and precipitates out later as NaO2 in the discharge process. Solvents with a lower ability to control the Lewis acidity of Na+ do not form an ion pair interaction with O2− and proceeds through a surface mechanism where, upon further oxygen reduction, Na2O2 is preferentially formed at the interface.In conclusion, in situ SERS investigations have shown that solvent choice can influence the overall surface discharge product of Na‐OAs 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.


N2′). This dehydrative allylation procedure is very mild and atom economical, producing only water as the by-product and avoiding any unnecessary waste/steps associated with installing a leaving or activating group on the substrate. Computational studies are presented to gain insight into the mechanism of the reaction. Calculations indicate that the regioselectivity is under equilibrium control and is ultimately dictated by the thermodynamic stability of the products.A gold(I)-catalysed direct thioetherification reaction between allylic alcohols and thiols is presented. The reaction is generally highly regioselective (S This change of S-nucleophile class initially caused several problems. Firstly, the reactions are a lot slower, despite alkyl thiols being more nucleophilic. We have recently ascribed this to the fact that the higher Lewis basicity of alkyl thiols (vs. thiophenols) will push the equilibrium towards the deactivated + model catalyst, 5′, and initially considered the reaction of PhSH with Me2C(OH)CH=CH2, 4′, a simple model of allylic alcohol 4 with the n-hexyl groups replaced by methyls. Subsequent calculations with the full Echavarren catalyst, 5, provided very similar energetics, details of which are given in the Supporting Information. We report free energies derived from gas-phase optimisations with the BP86 functional, corrected for chloroform solvent (PCM method) and dispersion effects  level).Density functional theory (DFT) calculations have been performed to probe these selectivity issues. The calcN2′ and SN2 products that can be formed  through the reaction of 4′ with PhSH . The profile starts from 4′IntI, formed from 5′ by displacement of MeCN by a π-bound allylic alcohol which is itself H bonded to the external PhSH. The latter is located anti to the metal centre and so is set up for external attack of the thiol at the γ-position  with a modest barrier of only 6.8 kcal mol−1. Subsequent conformational searching revealed several additional forms of 4′IntI, the most stable of which has G=−2.5 kcal mol−1. This species, 4′bIntI, is related to 4′IntI by rotation about the C2–C3 bond which then places the {HOLHSPh} moiety over the Au centre. The overall barrier for the SN2′ process, relative to this low energy precursor, is therefore 11.6 kcal mol−1 and so remains readily accessible (assuming facile interconversion of 4′bIntI and 4′IntI).[4′IntII the SN2′ product 3 a′ can then either be released, or alternatively a second SN2′ step with PhSH can lead to the net formal SN2 product 28. Thus, displacement of H2O from 4′IntII with PhSH gives 4′IntIII . Thiol attack via 4′TS(III-IV)  then generates 4′IntIV  in which the net SN2 product 28 is now π-bound to the gold centre.2A computed mechanism for the formation of r centre. The proc IntI4′)., 28 From4′TS(I-II) and 4′TS(III-IV) are given in Figure 4′TS(I-II) the main geometric feature is a slipping of the alkene moiety induced by the approach of the thiol, with an elongation of the Au–C1 distance to 2.61 Å compared with 2.30 Å in 4′IntI. At this stage there is relatively little elongation of either the S1–H1 (1.39 Å compared with 1.37 Å in 4′IntI) or the C3–O bonds (1.47 Å compared with 1.46 Å in 4′IntI) and in these terms the transition state exhibits an early geometry. In contrast 4′TS(III-IV), equating to nucleophilic attack of a second PhSH at C3, has a late geometry with the new C3–S2 bond nearly fully formed (1.94 Å compared with 1.92 Å in 4′IntIV), extension of the C1LS1 distance to 2.37 Å and slippage of the gold away from the C2–C3 bond such that it begins to interact with C1 (AuLC1=2.67 Å).3The computed structures of N2′ steps in Figure 4′IntII, the immediate precursor to the observed SN2′ thioether product (here modeled by 3 a′), is the most stable of these. This is also consistent with the reaction conditions, which require relatively long reaction times to achieve equilibrium, but without significant heating. The experimentally observed selectivity for the SN2′ product reflects therefore the greater thermodynamic stability of this species over the alternative formal SN2 product 28, rather than any kinetic preference for the SN2′ product.The low barriers and moderate exergonicities associated with the two SN2′ thioetherification products formed between PhSH and substrates 9, 10 and 11 , being 4.0 kcal mol−1 more stable than the alternative SN2′ product 28  at +9.9 kcal mol−1 to give 24IntII in which 28 is bound to the Au centre. A second SN2′ step with PhSH then occurs via 24TS(III-IV)  to give 24IntIV at −4.8 kcal mol−1 from which 3 a′ can be released. The overall shape of the free energy surface in Figure 24IntIV) again corresponding to the most stable intermediate. In this case we also considered whether the formal SN2 product 3 a′ could be formed by a direct SN2 reaction of PhSH at an O-bound precursor, 24IntV  of 26.3 kcal mol−1 indicates this process is disfavoured kinetically relative to the facile sequential SN2′ steps shown in Figure C→D Scheme −1. We also tested the formation of the analogous allylic cation from substrate 17 and found that, while this is somewhat more accessible, it does remain strongly endergonic, by +23.9 kcal mol−1. Allylic cation formation therefore cannot account for the lack of chirality transfer observed with enantioenriched substrate 17, despite the presence of the potentially stabilising phenyl substituent.5Due to the different regioselectivity in this case we also confirmed that the formal Ss Figure . Thus, n4 Figure . IntV24 N2′ or SN2 regioselectivities) and that any regioselectivity is ultimately dictated by the thermodynamic stability of the products. With this in mind, we postulated that the apparent failure of enantioenriched (S)-18, (R)-21 and (R)-20 to transfer chirality to its respective products is due to the subsequent racemisation of any enantioenriched product that might be formed -3 ab in 80:20 e.r. and rac-29 in a 2:1 ratio. This 2:1 mixture was then subjected to the gold(I) reaction conditions (Scheme R)-3 ab completely racemised but the ratio of 3 ab/29 also increases from 2:1 to 4:1, suggesting equilibration towards the more stable isomer.[The computational studies point towards accessible isomerisations between the two possible allylic thioether products  and stereoselective (E) gold(I)-catalysed direct thioetherification of allylic alcohols has been developed. The reaction has a wide substrate scope including tertiary, secondary and certain primary allylic alcohols, and tolerates a wide range of substituents on the thiol nucleophile. Unlike the related etherification reaction employing alcohol nucleophiles, the thiol nucleophile does not need to be in excess; an approximately 1:1 ratio of substrates is sufficient with no observed self-reaction of the allylic alcohol substrate. Furthermore, the SN2′ regioselectivity is generally excellent. The exception to this is when the γ-position of the allylic alcohol is fully substituted. When alkyl thiols are used as nucleophiles, the resulting products are prone to oxidation to the sulfone during silica gel chromatography, but this is successfully avoided by employing a gold-scavenging resin Reaxa QuadraPure™ MPA prior to purification. Computational studies on selected substrates suggest that both the formal SN2′ and SN2 products can be readily and reversibly formed under the reaction conditions. The observed regioselectivities are therefore dictated by the thermodynamic stabilities of the products. The observation of any formal SN2 products will be the result of two sequential SN2′ steps, this route being kinetically favoured over a direct SN2 reaction. The lack of chirality transfer in enantioenriched substrates can also be accounted for by the racemisation of products.In conclusion, a mild, regioselective  and catalyst 5 (5 mol %) were added to a solution of allylic alcohol (1 equiv) in chloroform (0.4 m). The resulting mixture was allowed to stir at 35–60 °C for 24–72 h as required. The mixture was then filtered through a plug of silica using diethyl ether as eluent. The filtrate was concentrated under reduced pressure and the crude material was purified by flash column chromatography. Full experimental procedures, characterisation for all new compounds and copies of 1H and 13C NMR spectra are provided in the Supporting Information.Thiol


Quantum computation offers potential advantages in solving a number of interesting and difficult problems. Several controlled logic gates, the elemental building blocks of quantum computer, have been realized with various physical systems. A general technique was recently proposed that significantly reduces the realization complexity of multiple-control logic gates by harnessing multi-level information carriers. We present implementations of a key quantum circuit: the three-qubit Toffoli gate. By exploring the optical selection rules of one-sided optical microcavities, a Toffoli gate may be realized on all combinations of photon and quantum spins in the QD-cavity. The three general controlled-NOT gates are involved using an auxiliary photon with two degrees of freedom. Our results show that photons and quantum spins may be used alternatively in quantum information processing. Quantum computing is an active area of research because of its ability to efficiently solve difficult problems without efficient classical algorithms123123568S) can be classically simulated efficiently, so they are probably not sufficiently universal for quantum computation. These gates, together with other one-qubit gates, not generated by the gates in the Clifford group, form a universal set of gates for quantum computation81519Because classical computing is designed around irreversible gates, it is impossible to directly translate this expertise into the quantum world. The Gottesman-Knill Theorem says that Clifford gates . These DOFs may be regarded as high-dimensional systems. In fact, high-dimensional systems may provide different quantum correlations and may be useful in quantum information processing21222324252627282124Motivated by their scheme232930333423323337383940414243441314151617192034293044454647484950515253di denotes the spatial modes of one photon. In the following, we also denote The Toffoli gate is an important three-qubit entangling gate in quantum logic gates111210112930λ-cavity between two GaAs/Al(Ga)As distributed Bragg reflectors. The QD is located in the center of the cavity to achieve maximal light-matter coupling. If the QD is neutral, optical excitation generates a neutral exciton. If the QD is singly charged, i.e., a single excess electron is injected, optical excitation can create a negatively-charged exciton (X−), which consists of two electrons bound to one hole3738X− in the state X− in the state Consider a singly charged GaAs/InAs quantum dot (QD) inside a micropillar cavity3738under ideal conditions. In the following, this ideal spin-cavity unit is used to realize the Toffoli gate on photons and quantum-dot spins for efficient quantum information processing. Then, the experimental spin-cavity unit will be discussed in the last section.A, B and C in the statesConsider three linearly circularly polarized photons Our goal is to realize the Toffoli gate with the following formA and B are the controlling qubits while the photon C is the target photon. The detailed circuit is shown in ei in the state D in the state TAB,C is completed with the following three controlled gates.where the photons S1 shown in A as an input pulse passes through the cPS1, cavity Cy1, cPS2, sequentially. Then W1 is performed on the spin e1. Now, the pulse D from the spatial mode d1 passes through the H1, cPS3, cavity Cy1, cPS4, H2, sequentially. After these operations, the joint system consisting of the photons A and D, and the spin e1 is changed from First, from the subcircuit e1 under the basis A for the measurement outcome A and the polarization DOF of the auxiliary photon D, which is different from previous CNOT gate on the same type of input system.after the measurement of the electron spin S2 shown in B passes through the cPS5, cavity Cy2, cPS6, sequentially. Then W2 is performed on the spin e2. Now, the photon D passes through the BS1, cPS7, X1, cavity Cy2, X2, cPS8, BS2, sequentially. After these operations, the joint system consisting of the photons A, B and D, and the spin e2 is changed from Second, from the subcircuit e2 under the basis B for the measurement outcome B and the spatial mode DOF of the auxiliary photon D.after the measurement of the electron spin S3 shown in D from the spatial mode d2 passes through the cPS9, cavity Cy3, cPS10, sequentially. Then W3 is performed on the spin e3 Now, the photon C passes through the H3, cPS11, cavity Cy3, cPS12, H2, sequentially. After these operations, the joint system consisting of the photons A, B, C and D, and the spin e3 is changed from Third, from the subcircuit where a Pauli flip e3 under the basis D from the spatial mode a2 for the measurement outcome CNOTD,C performed on the auxiliary photon D and the input photon C as followsafter the measurement of the spin which is an essential three-qubit operation.D under the basis 3, two circularly polarizing beamsplitters cPS13 and cPS14, two half waveplates H5 and H6, and four single photon detectors Finally, by performing the single qubit measurements on the photon TAB,C shown in equation CNOTD,C.Thus, the Toffoli gate ei in the statesConsider three electron spins This section is to realize the Toffoli gatee1 and e2 are the controlling qubits, while the electron spin e3 is the target qubit. The detailed circuit is shown in D in the state where the electron spins D from the spatial mode d1 passes through the half waveplate H1 to H2 sequentially. The joint system consisting of the photon D and the electron spin e1 changes from First, the auxiliary photon S4) has realized the controlled-NOT gate e1 and the polarization DOF of the auxiliary photon D under the joint basis This subcircuit  has realized the controlled-NOT gate e2 and the spatial mode DOF of the photon D under the joint basis This subcircuit  has realized the controlled-NOT gate D and the input spin e3 under the joint basis where the Pauli flip Finally, the joint system D under the basis 3, two circularly polarizing beamsplitters cPS7 and cPS8, two half waveplates H3 and H4, and four single photon detectors by measuring the auxiliary photon The present Toffoli gate on a three-photon system shown in the A and B jointly control an electron spin e; their initial states are S1, the joint system consisting of three qubits and an auxiliary photon D changes from First, let two photons S2, Moreover, from the CNOT gate realized with the subcircuit S6, the joint system Furthermore, from the CNOT gate realized with the subciruit which may collapse intoMD on the photon D under the basis by performing the single qubit measurement e1 and e2 in the states A in the state S4 and S5 in D changes from Second, consider two electron spins S3 in Moreover, from the CNOT realized with the subcircuit which may collapse intoMD of the photon D under the basis after performing the measurement A in the state e in the state B in the state S1 in S5 in S3 in D changes from Third, consider one photon which may collapse intoMD of the photon D under the basis e. Thus, a Toffoli gate has been realized on two electron spins and one photon.after the measurement A in the state e1 in the state e2 in the state S1 in S5 in S6 in D changes from Finally, consider one photon which may collapse intoMD of the photon D under the basis by performing the measurement ωx = Δωc = 0, if one neglects the cavity side leakage κs ≈ 0, it easily follows that |r0| → 1 and |r| → 1 when the cooperativity parameter g2/(κγ) of cavity QED is large enough. Thus, our six Toffoli gates are deterministic and faithful. However, the side leakage from the cavity is unavoidable in the experiment444547484950515253κ. Every other unwanted photon loss, such as cavity absorption and scattering, are characterized by the overall loss rate κs. Taking into account the coupling through the cavity decay channel and neglecting the spatial dependence, the relation of the input field operator The optical selection rules of a QD-cavity system shown in equation ωc and Δωx are the frequency detunings of the cavity mode and dipole transition, respectively, in relation to the input probe light (See Method). When the quantum dot is uncoupled from the cavity (g = 0), r(ω) is reduced to373839where Δωc = Δωx = 0, the reflection coefficients |r| and |r0| are evaluated in θ and θ0 are evaluated in κs/κ and the cooperativity parameter C = g2/(κγ) of cavity QED56r| ≈ 1 and |r0| ≈ 1 when C ≫ 10 and κs/κ → 0, and these additional conditions are not required for relative phase shifts θ0 = π and θ = π because r and r0 are real under the resonant conditions Δωc = Δωx = 0. Hence, the real reflection coefficients r and r0 will be considered under the resonant conditions.These complex coefficients indicate that the reflected light may experience a phase shift32333435364445In fact, the ideal optical selection rules shown in equation in the experiment. Based on these general optical selection rules, one can also complete the Toffoli gate from our schemes.e1, e2 and e3 in the state A, B, C, and D are involved; the success of this protocol is heralded by the instance in which the detector Pj is a successful reflection probability of the j-th photon from a micropillar cavity375054Fpp,p, as evaluated in e1, e2 and e3 are involved, and one photon D is used; its success is determined by the photon D, which is detected at the detector Ess,s is evaluated in Fpp,p is evaluated in D detected at the detector For our first Toffoli gate on the three photons shown in g ≫  is necessary394041484950515253g being larger than the geometric mean of the atomic and cavity line widths. In general, the system can be parameterized in terms of two dimensionless parameters, namely, the ratios g/κ and g/γ in the cavity QED description or, in the classical description, the cooperativity parameter C and the line width ratio κ/γ. The cavity QED strong-coupling condition 2g >  corresponds to a normal-mode splitting that is much larger than the line widths of the normal modes. The cooperativity parameter of cavity QED is shown to play a central role and is given a geometrical interpretation. The cooperativity has been realized up to 27EPP,P and ESS,S are greater than 91.24% for κs/κ ≈ 0.248FPP,P and FSS,S are greater than 93.47% for κs/κ ≈ 0.248−3 on a two-dimensional lattice of qubitsκs/κ < 0.04 and the cooperativity should be C > 28 for a photonic Toffoli gate, whereas the cavity leakage ratio should be κs/κ < 0.03 and the cooperativity should be C > 34 for a Toffoli gate on a three-spin system. When the fault tolerance threshold is reduced to 1 × 10−3 using controlled phase gates based on dipole-induced transparencyC > 38 for a photonic Toffoli gate, whereas the cavity leakage ratio should be reduced to 0.015 and the relative coupling strength should be improved to 4.2 for a Toffoli gate on a three-spin system. κs/κ = 0.05 has been reported, which could be achieved by taking a pillar microcavity with the quality factor of Q = 165000 demonstrated in ref. Q = 9000, which is still in the strong-coupling regimeTypically, the cavity side leakage may greatly affect the efficiency and fidelity of the Toffoli gate. As shown in the X− stays in the ground state for the most time. With a first-order approximation, we can adiabatically eliminate E10 = E1 − E0. If the dephasing is considered for the atomic system, it may be modeled by introducing phenomenological decay terms or noise operators If the experimental electron spin decoherence and trion dephasing4144505744In conclusion, we have investigated the possibility of hybrid quantum computation assisted by the quantum spins and photons with two DOFs. Six deterministic Toffoli gates are realized on the joint system of all combinations of the photon or the quantum spin systems. Compared with previous Toffoli gates13141516171918343547484950515253A singly charged GaAs/InAs QD32333738394041424344H = H1 + H2 + H3. ρ is an arbitrary system operator. σ− and σ+ are the Pauli raising and lowering operators respectively. ωc is the resonant frequency of the dipole, and σz is the Pauli operator for the population inversion. κ is the decay rate of the cavity field due to ohmic losses in the metal. κs is the decay rate of the cavity side leakage mode due to scattering into free-space modes. The scattering rate κs may be calculated classically from the Larmor formula. where ωc = ωc − ω and X− stays in the ground state for the most time39404142where Δr(ω) is defined in equation g = 0), r(ω) is reduced to r0(ω) as shown in equation g ≫ , one can get |r| ≈ 1 and |r0| ≈ 1 under resonant conditions by adjusting ω, ωx and ωc. Thus, if the excess electron spin lies in the spin state θ = arg[r(ω)](θ0 = arg[r0(ω)]) by passing through the cavity. Conversely, if the excess electron spin lies in the spin state θ = arg[r(ω)](θ0 = arg[r0(ω)]) by passing through the cavity. Thus, two phase shifts may be obtained as373839where θ0 = π and θ = 0323344When the side leakage and cavity loss are ignored, the optical selection rules shown in equation e. The joint state is first represented by one Hadamard transformation W, i.e., e can be faithfully disentangled. The experimental performances depend on the experimental optical selection rules shown in equation To complete our Toffoli gates, the entangled excess electron spins have to be measured under the basis How to cite this article: Luo, M.-X. et al. Hybrid Toffoli gate on photons and quantum spins. Sci. Rep.5, 16716; doi: 10.1038/srep16716 (2015).