Semih Ener
email: semih.ener@tu-darmstadt.de

Konstantin P Skokov

Dhanalakshmi Palanisamy

Thibaut Devillers

Johann Fischbacher

Gabriel Gomez Eslava

Fernando Maccari

Lukas Schäfer

Léopold V B Diop

Iliya Radulov

Baptiste Gault

Gino Hrkac

Nora M Dempsey

Thomas Schrefl

Dierk Raabe

Oliver Gutfleisch

Twins -A weak link in the magnetic hardening of ThMn12-type permanent magnets

Keywords: magnets, ThMn12-type compounds, coercivity, twin boundaries, microstructure

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Introduction

Nd2Fe14B-type magnets are key components in electro-mobility, wind turbines, robotics and automatization with expected strong increases in demand in all sectors. The rare-earth (RE) supply crisis of ten years ago led to a drastic increase in prices and revealed the dependency on critical RE elements [START_REF] Gutfleisch | Magnetic Materials and Devices for the 21 st Century: Stronger, Lighter, and More Energy Efficient[END_REF] for the production of high performance permanent magnets. Significant efforts are ongoing to find alternative material systems which present magnetic properties similar or superior to those the Nd2Fe14B benchmark, but which contain less critical elements [START_REF] Skokov | Heavy rare earth free, free rare earth and rare earth free magnets -Vision and reality[END_REF]. RE-lean Nd-and Sm-based alloys with ThMn12-type structures (tetragonal I4/mmm #139) have been known for more than 30 years. More recently, in 2015, Hirayama et al. reported the intrinsic magnetic properties of NdFe12Nx thin films (saturation magnetization µ0Ms ~ 1.66 T, anisotropy field µ0Ha ~ 8 T and Curie temperature Tc ~550 °C) to be superior to those of Nd2Fe14B [START_REF] Hirayama | NdFe12Nx hard-magnetic compound with high magnetization and anisotropy field[END_REF][START_REF] Hirayama | Rare-Earth Lean Hard Magnet Compound NdFe12[END_REF].

However, the ThMn12-type structure was preserved only in thin-films with thicknesses up to 360 nm. A major drawback of the NdFe12Nx system is the necessity of interstitial modification by nitrogen to obtain high saturation magnetization, anisotropy field and Tc. Decomposition of the nitride starting from around 570°C prevents conventional densification via sintering or hot pressing. Nonetheless, investigations of materials with the Nd(Fe,X)12-composition show the possibility to stabilize ThMn12 structure in the bulk state by addition of doping elements, such as Cr, V, Ti, Mo, W or Si [START_REF] Buschow | Structure and properties of some novel ternary Fe-rich rare-earth intermetallics (invited)[END_REF][START_REF] De Mooij | Some novel ternary ThMn12-type compounds[END_REF], that substitute Fe atoms. Among these elements, Ti is very effective and stabilizes the structure with minimal doping (approx. 7.7 at.%) [START_REF] Coey | Comparison of the intrinsic magnetic properties of R2Fe14B and R(Fe11Ti); R = rare earth[END_REF].

The sign of the Stevens factor αJ of Sm 3+ is opposite to that of Nd 3+ , and as a consequence Sm(Fe,Ti)12 exhibits a high anisotropy field (µ0HA ~ 10 T at room temperature [START_REF] Hu | Intrinsic magnetic properties of the iron-rich ThMn12-structure alloys R(Fe11Ti)[END_REF]) without any interstitial modification, which makes it a promising material for the production of anisotropic bulk and fully dense permanent magnets. In 2017, Hirayama et al. reported that the intrinsic magnetic properties of cobalt doped Sm(Fe1-xCox)12 thin films are superior to those of Nd2Fe14B [START_REF] Hirayama | Intrinsic hard magnetic properties of Sm(Fe1-xCox)12 compound with the ThMn12 structure[END_REF]. Other studies showed that the addition of Zr reduces the Ti concentration required to stabilize the ThMn12type structure in Sm-based bulk samples [START_REF] Kuno | .5Ti1.0-0.5 compounds as new permanent magnet materials[END_REF][START_REF] Hagiwara | Structural and magnetic properties of rapidly quenched (Sm,R)(Fe,Co)11.4Ti0.6 (R = Y, Zr)[END_REF]. Zr substitution also leads to an increase in the saturation magnetization, but the anisotropy field decreases.

Even though the intrinsic magnetic properties of both Nd and Sm based ThMn12-type thin films are reported to be superior to those of Nd2Fe14B, the production of bulk ThMn12-type structures remains challenging [START_REF] Inoue | Influence of nitrogenating conditions on magnetic properties of NdFe8Co3TiNx[END_REF][START_REF] Yang | Structure and magnetic properties of mechanically alloyed Nd(Fe,V)12Nx compounds[END_REF][START_REF] Yang | Magnetic properties and magnetic domain structures of NdFe10.5Mo1.5 and NdFe10.5Mo1.5[END_REF][START_REF] Yang | Effect of interstitial nitrogen on the structural and magnetic properties of NdFe10.5V1.5Nx[END_REF][START_REF] Wuest | Preparation of nanocrystalline Ce1-xSmx(Fe,Co)11Ti by melt spinning and mechanical alloying[END_REF][START_REF] Tozman | Intrinsic magnetic properties of Sm(Fe1-xCox)11Ti and Zr-substituted Sm1-yZry(Fe0.8Co0.2)11.5Ti0.5 compounds with ThMn12 structure toward the development of permanent magnets[END_REF][START_REF] Saito | Magnetic properties of SmFe12-based magnets produced by spark plasma sintering method[END_REF][START_REF] Neznakhin | Structure and magnetic properties of (Sm0.9Zr0.1)Fe11Ti alloys with ThMn12-type structure[END_REF][START_REF] Dirba | Development of ultra-fine grain sized SmFe12-based powders using hydrogenation disproportionation desorption recombination process[END_REF]. The hardness parameter, defined as:

= (1)
quantitatively determines the probability of resisting self-demagnetization when hard magnetic materials are fabricated in any possible shape. An empirical rule of thumb to select materials as potential candidates for application as permanent magnet [START_REF] Skomski | Magnetic anisotropy -How much is enough for a permanent magnet[END_REF] is that κ must be significantly greater than one. For Nd2Fe14B the κ value is 1.5, while it rises to a value of 2.1 for SmFe11Ti [START_REF] Hirosawa | Perspectives for high-performance permanent magnets: applications, coercivity, and new materials[END_REF]. The excellent intrinsic magnetic properties measured on single crystalline samples (K1, Ms) and the high value of the hardness parameter κ cannot guarantee the direct use of the material for permanent magnet production. Indeed, each class of hard magnetic material requires its own special microand nanostructure, which needs to be controlled by utilizing peculiarities of the phase diagram and solid-state phase transformations. More importantly, in the context of this work, defects resulting from phase formation under specific processing regimes can serve as centers where reverse magnetic domains nucleate and propagate even in weak magnetic fields, thereby drastically reducing the coercivity of the sample. Due to the above mentioned promising intrinsic properties, the SmFe11Ti material is selected as a prototypical representative for ThMn12-type materials.

Figure 1 compares coercivity values from literature of benchmark Nd2Fe14B-based materials and Sm(Fe,X)12-based materials [START_REF] Dirba | Development of ultra-fine grain sized SmFe12-based powders using hydrogenation disproportionation desorption recombination process[END_REF][START_REF] Ohashi | Magnetic properties of Fe-rich rareearth intermetallic compounds with a ThMn12 structure[END_REF][START_REF] Tatsuki | Enhancement of the Magnetic Properties of Sm-Fe and Nd-Fe-N Alloys with ThMn12 Structure by Utilizing HDDR Phenomena[END_REF][START_REF] Sugimoto | Magnetic properties and microstructures of the (SmFe10V2)1-x-(Sm2Fe17)x cast alloys[END_REF][START_REF] Dirba | Anisotropic, singlecrystalline SmFe12-based microparticles with high roundness fabricated by jet-milling[END_REF][START_REF] Gabay | Mechanochemical synthesis of magnetically hard anisotropic RFe10Si2 powders with R representing combinations of Sm, Ce and Zr[END_REF][START_REF] Singleton | Magnetic and structural properties of melt-spun rare-earth transition-metal intermetallics with ThMn12 structure[END_REF][START_REF] Strzeszewski | High coercivity in Sm(FeT)12 type magnets[END_REF][START_REF] Okada | High Coercivity in Melt-Spun SmFe10(TiV)2 Ribbons[END_REF][START_REF] Sugimoto | Enhancement of Magnetic Properties of Sm(Fe,Co,Ti)12 Melt-Spun Ribbons by Refining Crystallized Grains[END_REF][START_REF] Pinkerton | Magnetic hardening of SmFe10V2 by melt-spinning[END_REF][START_REF] Wang | Magnetic and structural studies in Sm-Fe-Ti magnets[END_REF][START_REF] Saito | Magnetic properties of Sm-Fe-Ti nanocomposite magnets with a ThMn12 structure[END_REF][START_REF] Shield | Microstructures and phase formation in rapidly solidified Sm-Fe and Sm-Fe-Ti-C alloys[END_REF][60][START_REF] Sepehri-Amin | Achievement of high coercivity in Sm(Fe0.8Co0.2)12 anisotropic magnetic thin film by boron doping[END_REF], as a function of grain size. For the Nd2Fe14B system coercivity values up to 2.95 T have been reported for thin films, while coercivity around 1.95 T have been achieved for sintered bulk samples. Even though the Sm-based ThMn12-systems show similar anisotropy field, the reported coercivity values do not surpass 1.1 T in melt-spun ribbons. The highest coercivity reported for the ThMn12-type bulk system is 1.07 T for Sm8Fe75.5Ga0.5Ti8V8. Recently, coercivities of 1.2 T and 1.26 T has been reported for Sm(Fe0.8Co0.2)12B0.5 thin films and Sm0.815Zr0.185(Fe0.81Co0.19)11.19Ti0.81 mechanochemically synthesized oriented particles, respectively, as indicated by the dashed line in Figure 1. For Sm-Fe-Ti ternary compositions the reported coercivity values do not surpass 0.51 T, as indicated by a dotted line in Figure 1.

Figure 1.

Comparison of reported coercivity values as a function of grain size. Open symbols and full symbols correspond to Nd2Fe14B and Sm(Fe,X)12 systems, respectively. Different colors correspond to different sample production methods and shaded areas are guides to the eye. The highest reported coercivity of a ThMn12-type structure is shown by the dashed line [START_REF] Gabay | High-coercivity ThMn12-type monocrystalline Sm-Zr-Fe-Co-Ti particles by high-temperature reduction diffusion[END_REF]. The highest reported coercivity for a ternary Sm(FeTi)12 sample is shown by the dotted line [START_REF] Sugimoto | Enhancement of Magnetic Properties of Sm(Fe,Co,Ti)12 Melt-Spun Ribbons by Refining Crystallized Grains[END_REF].

A possible explanation for the low coercivity of SmFe11Ti-based magnets may be the absence of a non-magnetic grain boundary phase which magnetically decouples adjacent grains, as found in classical Nd2Fe14B-type sintered magnets [START_REF] Palanisamy | Grain boundary segregation, phase formation, and their influence on the coercivity of rapidly solidified SmFe11Ti hard magnetic alloys[END_REF]. Another reason might be the presence of macroscopic twin boundaries and microscopic slip-bands, which were previously reported in ThMn12-type structures [START_REF] Dirba | Anisotropic, singlecrystalline SmFe12-based microparticles with high roundness fabricated by jet-milling[END_REF][START_REF] Skokov | Magnetic properties of Gd3FexTi3 (x=34, 33, …, 24), TbFe11Ti and TbFe10Ti single crystals[END_REF][START_REF] Maccari | Rapid solidification of Nd1+XFe11Ti compounds: Phase formation and magnetic properties[END_REF]. Although the twins have been observed in different ThMn12type material systems before, up to now, their influence on the physical properties were not investigated in detail. It is important to state that there are no reports on twins in high performance magnets based on Nd2Fe14B-, SmCo5-, Sm2Co17-and Sm2Fe17Nx-type rare earth transition metal phases. In contrast and equally important, a mixture of stacking faults, anti-phase boundaries and abundant twins have been shown to occur in rare earth free L10-type hard magnetic phases. For the Fe-Pd, Fe-Pt and Co-Pt systems twins conjugate along {110}-planes, while for MnAl systems twins conjugate along {111}-planes. These twins can form during a diffusional transformation or recrystallization [START_REF] Palanisamy | Elemental segregation to twin boundaries in a MnAl ferromagnetic Heusler alloy[END_REF][START_REF] Yanar | Evolution of microstructure and defect structure in L10-ordered manganese aluminide permanent magnet alloys[END_REF][START_REF] Landuyt | Defect structure and magnetic properties of MnAl permanent magnet materials[END_REF][START_REF] Jakubovics | The effect of crystal defects on the domain structure of Mn-Al alloys[END_REF][START_REF] Bance | Role of twin and antiphase defects in MnAl permanent magnets[END_REF][START_REF] Klemmer | Magnetic hardening and coercivity mechanisms in L10 ordered FePd ferromagnets[END_REF][START_REF] Shima | Hard Magnetic Films[END_REF][START_REF] Gau | Electron Microscopy of Mn-Al-C Magnets[END_REF][START_REF] Zhang | Magnetic domains and coercivity in polytwinned ferromagnets[END_REF]. In 2001, Attané et al., reported that micro-twins observed in FePt films act as pinning centers that increase coercivity [START_REF] Attané | Domain wall pinning on strain relaxation defects in FePt(001)/Pt thin films[END_REF]. Microscopic twins and stacking faults observed in Co-based high-density magnetic recording media materials such as CoCr and CoCrPt were also credited with increasing coercivity [START_REF] Glijer | Correlation of structural defects and magnetic properties in CoCrPt/Cr and CoCrTaPt/Cr thin films for ultra high density magnetic recording media[END_REF][START_REF] Demczyk | Growth Characteristics of Co-Cr Thin Films for Magnetic Recording[END_REF]. On the other hand, twins in Mn-Al have been reported to be deleterious for coercivity, by decreasing locally the energy necessary to nucleate reversed magnetic domains [START_REF] Jia | L10 rareearth-free permanent magnets: The effects of twinning versus dislocations in Mn-Al magnets[END_REF]. The bivalent role that twins appear to play in different hard magnetic material systems highlights the importance of a deeper understanding of the link between microstructure and magnetic properties in such technologically important materials.

To date, no detailed investigation has been reported on the nature of twins in SmFe11Ti-based compounds and their effect on the hard-magnetic properties. In this study, we address this issue by comprehensively investigating twins at different length-scales in (1) twinned bulk crystals, [START_REF] Skokov | Heavy rare earth free, free rare earth and rare earth free magnets -Vision and reality[END_REF] poly-crystalline bulk samples and (3) poly-crystalline thin films. At various stages we also refer to the benchmark Nd2Fe14B to elucidate commonalities and disparities. We reveal changes in atomic ordering across twin-like structures and a change in composition at the twin boundaries with respect to the adjacent lattice, using advanced electron microscopy and atom probe tomography.

Micromagnetic modelling support our experimental results and elucidate how twins influence local and global magnetization reversal and ultimately macroscopic hysteresis properties in the SmFe11Ti system. We also consider how twin formation is associated with the accumulation of internal stresses during solidification and grain growth during thermal treatments, as a consequence the energy of the twin formation is becoming more favoured in comparison to the formation of specific grain boundaries. Molecular dynamics shows that there is a fine balance between twin formation and certain grain boundary phases. In this work, we study in detail on all length scales how twins act as a weak link in the magnetic hardening of ThMn12-type permanent magnets.

Experimental and Simulation Details

Poly-crystalline and single-crystalline sample preparation and characterization: Polycrystalline Sm1.25Fe11Ti ingots were prepared by melting high-purity (>99.99%) starting elements in an induction furnace under a purified argon atmosphere. Using an arc-melting setup, the ingot was suction-cast into a rectangular slab of 0.5 mm thickness, to ensure fast solidification/cooling (approx. 10 3 K.s -1 ). The suction-cast slabs were wrapped in Mo foil, sealed in quartz ampules under Ar atmosphere and annealed at 1175 ℃ for 30 minutes to stabilize the Sm(Fe,Ti)12 phase and then quenched in room temperature water. It is worth mentioning that an off-stoichiometric (higher Sm concentration) composition was used to compensate for evaporation losses during the melting and annealing procedures.

Structural characterization and phase determination of the annealed samples were done using room temperature x-ray powder diffraction (XRD). The measurements were carried out on a Stoe Stadi P diffractometer with Mo Kα1 radiation, in transmission mode and an angular 2θ range from 5° to 50°. Phase matching and structural refinements were performed using FullProf software [START_REF] Rodríguez-Carvajal | Recent advances in magnetic structure determination by neutron powder diffraction[END_REF].

Microstructural analysis was carried out using a Tescan VEGA3 scanning electron microscope (SEM) with secondary electron (SE) and backscattered electron (BSE) detectors. The chemical composition of the observed phases was determined by energy-dispersive X-ray spectroscopy (EDS) measurements.

Magneto-optical Kerr effect (MOKE) measurements were performed on a Zeiss Axio Imager.D2m microscope with polarized light function equipped with an electromagnet (evico magnetics GmbH).

The MOKE measurements were done using polar contrast mode to ensure highest contrast between the magnetic domains pointing in opposite directions. To enhance the magnetic contrast of the images, the non-magnetic background image was subtracted from the measurement images by using the KerrLab software.

SmFe11Ti single crystals were produced by re-melting Sm1.25Fe11Ti polycrystalline samples and then slowly cooling them at a rate of 1 ℃/h. Re-melting was carried out in a ZrO2 crucible in a sealed quartz tube under a protective Ar atmosphere. The obtained particles, of typical size 500-1000 µm, were mechanically separated from each other and they are referred as twinned crystals.

Selected crystals were oriented using back-scattering Laue x-ray diffraction and mounted for Thin film sample preparation and characterization: Sm-based 1:12-type films of thickness 1.5 µm were deposited by triode sputtering onto thermally oxidized Si wafers at 700 ℃. 50 nm Ta buffer layers were used to prevent chemical reaction between the wafer and the magnetic layer and 10 nm Ta capping layers were added to prevent oxidation. We present results from two samples with composition Sm0.7Fe11.3Ti0.9 and Sm0.7Zr0.2Fe11.4Ti0.7, as estimated by EDS, referred to as the Sm(Fe,Ti)12 film and the (Sm,Zr)(Fe,Ti)12 film, respectively. Atomic structure and diffraction analysis of the films were carried out using transmission electron diffraction (TEM) microscopy.

TEM samples were prepared in cross-sectional configuration by mechanical polishing (down to 25 µm) followed by ion milling with a Gatan PIPS II (Precision Ion Polishing System). Ion milling of the TEM samples was performed at 3 kV for several hours until an electron transparent hole appears, followed by 5 minutes of cleaning at 0.5 kV.

High resolution microscopy on polycrystalline samples:

To locate grains with twin boundaries, electron backscatter diffraction (EBSD) analysis was performed on the polycrystalline samples using a Jeol JSM-6500F SEM operated at 15 kV. EBSD acquisition was performed using the TSL OIM EBSD system equipped with a Digiview IV EBSD detector with a step size of 0.1 µm.

Correlative electron microscopy and atom probe tomography (APT) investigations were carried out using a JEOL 2200FS aberration-corrected electron microscope (Thermo-Fischer TITAN) in scanning-transmission (STEM) mode and a local-electrode-atom-probe (LEAP 5000 XR, Cameca) instrument equipped with reflection lens, respectively. For APT, needle-shaped specimens were prepared from the twin boundary regions using a dual beam SEM/focused-ion-beam (FIB) instrument (FEI Helios Nanolab 600i) with an in-situ-lift-out procedure detailed in references [START_REF] Palanisamy | Elemental segregation to twin boundaries in a MnAl ferromagnetic Heusler alloy[END_REF][START_REF] Felfer | A reproducible method for damagefree site-specific preparation of atom probe tips from interfaces[END_REF][START_REF] Makineni | Correlative Microscopy-Novel Methods and Their Applications to Explore 3D Chemistry and Structure of Nanoscale Lattice Defects: A Case Study in Superalloys[END_REF]. The prepared lamella, containing the twin boundary region, was positioned and Pt-welded using the in-situ gas-injection system on the electro-polished posts of a halved TEM Mo-grid. The Mo-grid with the attached specimen was held in a special correlative holder designed in-house, as described in reference [START_REF] Herbig | Combining structural and chemical information at the nanometer scale by correlative transmission electron microscopy and atom probe tomography[END_REF]. The welded regions on the Mo-posts were sharpened by Ga ions accelerated at 30 kV with currents ranging from 80 pA to 0.78 nA. The sharpened specimens were checked by TEM to get the exact location of the twin boundary and further Ga-ion milling was performed such that the boundary is within 150 ~ 200 nm from the apex of the specimen. Before APT analysis, the specimens were cleaned by 5kV with 8 pA current to remove regions of the specimen severely damaged by the implantation of high-energy Ga-ions.

Near-atomic-scale compositional analysis was performed using a LEAP 5000 XR instrument operated in laser pulsing mode with a pulse repetition of 200 kHz and a pulse energy of 40 pJ. The specimen was kept at a base temperature of 60 K with a target detection rate of 5 ions detected per 1000 pulses. The APT reconstruction and analysis were performed using IVAS 3.8.2 software.

Micromagnetic simulations: Magnetization reversal in a single twinned grain was computed by micromagnetic simulations. The shape of the grain was generated with the Neper software [START_REF] Quey | Large-scale 3D random polycrystals for the finite element method: Generation, meshing and remeshing[END_REF].

The twin boundary was defined with the computer aided design software Salome [START_REF] Ribes | Salomé platform component model for numerical simulation[END_REF]. The geometric model is meshed with tetrahedral finite elements using MeshGems [START_REF] Meshgems | Distene's MeshGem Software[END_REF]. In order to treat the magnetostatic open boundary problem [START_REF] Qiushi | A review of finite element open boundary techniques for static and quasi-static electromagnetic field problems[END_REF], the finite element mesh for the magnetostatic scalar potential is extended outside the magnetic grains. The mesh size at the grain surfaces and twin boundaries was 1.5 nm. We computed the demagnetization curve by minimization of the Gibbs free energy for a decreasing external field. To solve for the next local minimum, we used a preconditioned conjugate gradient method [START_REF] Exl | Preconditioned nonlinear conjugate gradient method for micromagnetic energy minimization[END_REF]. Plots of the demagnetization curves were created using matplotlib [START_REF] Hunter | Matplotlib: A 2D Graphics Environment[END_REF].

Molecular dynamics simulations:

The molecular dynamics calculations were carried out by using modified Morse potentials on supercell structures with the crystallographic interface orientation and composition taken from experiments. The different surface energies of the 1:12 phase were calculated. The twin formation was incorporated as an additional interface energy contribution.

After the total energy with and without twin formation has been derived, we saw that once the grain reaches a certain size, twin formation is deemed preferential. This result underpins the findings in the experiments.

Results and Discussion

Single crystals and intrinsic magnetic properties

Magnetic anisotropy is the key requisite for achieving a large coercivity, and in the ideal case of uniform rotation of the magnetic moments in single domain particles, the coercivity Hc is given by Hc = Ha = 2K1/µ0Ms (Ha is anisotropy field, K1 is anisotropy constant, µ0 is the permeability of free space, Ms is the spontaneous magnetization). In reality, prediction of Hc from K1 and Ms, using the Stoner Wohlfarth model [START_REF] Stoner | A mechanism of magnetic hysteresis in heterogeneous alloys[END_REF], leads to a large overestimation. This is known as Brown's paradox [START_REF] Skokov | Heavy rare earth free, free rare earth and rare earth free magnets -Vision and reality[END_REF][START_REF] Kronmüller | Analysis of the magnetic hardening mechanism in RE-FeB permanent magnets[END_REF][START_REF] Brown | Virtues and Weaknesses of the Domain Concept[END_REF], and it is attributed to local magnetic softening caused by crystallographic defects, strain, secondary phases, magnetic and chemical inhomogeneities which reduce Hc to about αHa (α is the "Kronmüller factor" , which is related to nano-structural imperfections). As-cast materials often exhibit α = 0.01, and via sophisticated processing methods, e.g. melt spinning, jet milling, or die upsetting, α can reach to 0.25 of the theoretical limit in laboratory-made Nd-Fe-B and Sm-Cobased magnets.

We start by reporting in Figure 2 the room temperature magnetization curves of a SmFe11Ti twinned crystal, with the magnetic field applied along the [100] and [001] directions. The typical behavior of easy-axis magnetocrystalline anisotropy is observed; the data measured along the 4-fold symmetry axis [001] clearly identifies this as the easy magnetization direction, leading to a µ0Ms = 1.17 T (calculated by converting the mass magnetization to volume magnetization considering the x-ray density of the SmFe11Ti). At room temperature, curves recorded along the [100] and [110] directions are practically identical (the [110] curve is not shown), reflecting a weak magnetic anisotropy within the basal plane of this system at room temperature (see supplementary materials for details of the crystal structure). The anisotropy field, µ0Ha, estimated from the slope change in the magnetic hard [100]-axis curve (identified by an arrow in Figure 2), amounts to 10.5 T at 300 K. The value of K1 is estimated to be 627 J/kg (approx. 4.8 MJ/m 3 , considering the theoretical xray density). The hard-axis magnetization curve does not reach the same maximum value as the easy-axis curve above µ0Ha, due to anisotropy of the magnetic moment. This observed difference of the easy and hard axes magnetization values above the anisotropy field indicates a possible anisotropy of the magnetization, as has been reported for different rare-earth transition-meal systems [START_REF] Karpenkov | Intrinsic magnetic properties of hydrided and non-hydrided Nd5Fe17 single crystals[END_REF][START_REF] Bolyachkin | Magnetic anisotropy of YFe3 compound[END_REF][START_REF] Kuz'min | Magnetic anisotropy of La2Co7[END_REF]. Additionally, the low field region of the hard axis measurements shows non-linear behavior which is due to the formation of twins in the "single-crystals". The dashed red line corresponds to the hard axis measurement of a twin-free SmFe11Ti single crystal, as reported in [START_REF] Diop | Determination of the crystal field parameters in SmFe11Ti[END_REF]. One can see that the spontaneous magnetization of the Nd2Fe14B single crystal is significantly higher than that of the SmFe11Ti single crystal. On the other hand, the anisotropy field of SmFe11Ti is larger than that of Nd2Fe14B.

Next it is instructive to compare the magnetic domain structures imaged by Kerr microscopy of 0.5-1 mm-sized single crystals of Nd2Fe14B (Figure 3a,b) and SmFe11Ti (Figure 3c,d). One can see that for the Nd2Fe14B single crystal, the branched domain pattern on the basal plane (Figure 3a, the c axis and magnetization are perpendicular to the crystal surface) and stripe-like domains on the prismatic plane (Figure 3b, the c axis and magnetization lie in the plane of the image) are rather periodic and without significant distortions. In contrast, the domain structures of the SmFe11Ti twinned crystal (Figure 3c) are locally distorted (highlighted by the arrow), and this is more pronounced when the magnetization lies in the plane of the image (Figure 3d). Such step-like features observed on stripe magnetic domains are a signature of the presence of a crystallographic twinning structure, and this phenomenon was well-studied in other materials, such as magnetic shape memory materials [START_REF] Venkateswaran | Magnetic domain memory in multiferroic Ni2MnGa[END_REF][START_REF] Ge | Direct optical observation of magnetic domains in Ni-Mn-Ga martensite[END_REF], and materials showing giant magnetostriction [START_REF] Lord | Magnetic force microscopy of Terfenol-D fracture surfaces[END_REF][START_REF] Dooley | Induction mapping of magnetostrictive materials[END_REF]. In contrast to the Nd2Fe14B single crystal, the microstructure of the SmFe11Ti crystal comprises a significant number of crystallographic twins, which leads to the appearance of micrometer-sized blocks of changing orientation. As a result, the magnetization curve measured in the hard direction of the SmFe11Ti crystal (Figure 2) is not linear but exhibits a kink at low fields -i.e. a "softmagnetic like" curvature, which is not observable in the Nd2Fe14B single crystal.

Polycrystalline samples

Figure 4 shows the magnetic domain structures of Nd-Fe-B and Sm-Fe-Ti polycrystalline bulk samples. The domain patterns were observed on arbitrarily oriented polished surfaces of isotropic samples. In Nd2Fe14B grains, Figure 4a, the domain structure is rather uniform and the domain pattern changes only across grain boundaries, i.e. due to the different crystallographic orientation of adjacent grains. In contrast, the domain structure of the SmFe11Ti grains, Figure 4b, is often divided in two (or more) well distinguishable patterns by a twin boundary, some of which are marked by red arrows. Figure 5a shows an electron back scattered diffraction (EBSD) map of the SmFe11Ti polycrystalline bulk sample evidencing grains with different orientations. The EBSD map was indexed as Sm(Fe,Ti)12 (1:12 phase) matrix phase having a tetragonal crystal structure (space group I4/mmm) with lattice parameters a = b = 0.8589 nm and c = 0.4807 nm. Some grains contain unique boundaries corresponding to twins, such as the region within the black-dashed rectangle in Figure 5a. The twin orientation relation between the twin variants is approximately 57°, as shown Figure 5b. We also show a schematic view of the orientation with respect to the shared twin boundary, and found that the twinning boundary plane corresponds to {011}-type planes in the unit cell of the Sm(Fe,Ti)12 phase. Figure 5c shows a back scattered electron (BSE) micrograph for a similar twinned region with a twin boundary inside a grain. We confirmed the crystallographic relationship of two variants of a twin by using electron diffraction analysis and the results can be found in Supplementary Materials.

We prepared an atom probe tomography (APT) specimen specifically from the location marked by the dashed red outline in Figure 5c. Figure 6a is a low magnification high-annular-angle-dark-field (HAADF) image of the specimen with the twin boundary. HAADF provides Z-contrast, i.e. heavier atoms appear brighter. Figure 6b is a high-resolution (HR) HAADF image taken along the [100] zone axis from the location marked by a black-dashed square in Figure 6a. Figure 6c is a close-up on the unit cell viewed along the [100] zone axis, which can be compared with the simulated unit cell shown in Figure 6d with Sm in red and Fe/Ti in green. As expected, atomic columns containing Sm appear bright. We can also clearly see the (011) planes, marked at black dashed lines, which exactly resemble the zig-zag twin boundary, schematically shown with a black dashed line in Figure 6b. Figure 7 details the results from the APT analysis of the same specimen. Figure 7a is the brightfield electron micrograph of the specimen and Figure 7b shows the corresponding APT reconstruction with the distribution of Fe, Ti, Sm atoms and a Sm 9 at.% iso-composition surface highlighting the location of the twin boundary. Figure 7c shows a 2D Sm compositional map evidencing a local increase in the Sm content at the twin boundary. The measured matrix composition was found to be Sm7.7Fe84.3Ti7.9, matching the 1:12 stoichiometry. The composition profile taken across the boundary, Figure 7d, shows a Sm enrichment by approx. 2 at.% and a depletion of Ti by approx. 2.7 at.% with respect to the matrix phase. A slight enrichment of Fe, by approx. 0.7 at.%, is also observed. The composition at the twin boundary is close to the stoichiometry of the 3:29-phase i.e. Sm3(Fe,Ti)29 phase. The electron diffraction measurements (see Figure S2) show no additional spots corresponding to the 3:29-phase, most probably due to the early nucleation stage of the twin boundary phase. To complement the above reported studies of twins in bulk Sm-Fe-Ti samples, we also studied twins in polycrystalline Sm-Fe-Ti and Sm-Fe-Ti-Zr sputtered films (the latter were studied to probe the eventual influence of Zr additions in films). Both films contain a first layer of isotropically oriented equiaxed grains of maximum diameter 200 nm, at the interface with the Ta buffer, and elongated grains of width ~200 nm that traverse the rest of the film thickness. The elongated grains, with lattice parameters corresponding to the Sm(Fe,Ti)12 structure, are highly textured, with a dominant {101} texture. Multiple planar defects oriented parallel to the growth direction, which are attributed to twins in the Sm(Fe,Ti)12 structure, were identified in the elongated grains of both films. A high-resolution micrograph (see Supplementary Materials) shows that along the particular zone axis, the atomic crystal structure on both sides of the twinning boundaries look identical, because they are probably seen under a similar zone axis (e.g. [1 01] and [1 01 ] with 59.8° angle between planes). High-resolution TEM micrographs confirm the formation of twins in polycrystalline film samples, as in their bulk counterparts. For further details and crystallographic phase analysis of the investigated polycrystalline sample please see the Supplementary Materials.

2.3.Micromagnetic simulations and molecular dynamics calculations

In an isotropic granular magnet, a grain may be oriented at various directions with respect to the applied external field. For a single domain grain with uniaxial anisotropy and no defect, the demagnetization curves may be very close to the Stoner-Wohlfarth prediction, depending on the direction of the external field. The curve with the minimum switching field will occur when the applied field is oriented at 45 degrees with respect to the anisotropy axis [START_REF] Stoner | A mechanism of magnetic hysteresis in heterogeneous alloys[END_REF]. In this case the switching field will be half of the anisotropy field. Before irreversible switching, the magnetization will rotate reversibly out of the anisotropy axis. In micromagnetic simulations we expect a slight reduction of those switching fields, owing to the self-demagnetizing field of the grain.

We calculated the switching fields of a single twinned SmFe11Ti grain as a function of the orientation of the applied external field and compared the results to those of a single uniaxial Nd2Fe14B grain of the same geometric shape. In the ideal microstructure of a permanent magnet, the grains are separated by a paramagnetic grain boundary phase. Thus, by comparing the single grain behavior of SmFe11Ti and Nd2Fe14B, we can derive the upper limits of coercivity for both systems.

The shape of the model is shown as an inset of Figure 8a. The volume of the grain was equal to that of a cube with 295 nm sides. For Nd2Fe14B calculations we assumed a magnetocrystalline easyaxis parallel to the z direction of the global coordinate system. The intrinsic magnetic properties used were: anisotropy constant K1 = 4.9 MJ/m 3 , magnetization µ0Ms = 1.61 T, and exchange constant A = 8 pJ/m (page 401 of [100]). The anisotropy field of this phase is μ0Ha = 7.65 T. To model the effect of a twin in a SmFe11Ti grain, we introduced a 6 nm thick twin boundary (2.4 percent of the total volume) parallel to the xz-plane which separates the twin variants of the grain into two equally sized halves. The orientation of the two easy-axes of the two variants of the twin are shown as white double-headed arrows and are parallel to the yz-plane. The easy-axes enclose an angle of 180 -58 = 122 degrees, which is consistent with results from EBSD measurements and was checked with the Twiny software [101] which gives 121.5° for (011) twin planes of SmFe11Ti.

In this model we ignored the experimentally found zig-zag shape of the twin boundary. Test calculations revealed that due to the low amplitude (1-2 nm) of the pattern, which is smaller than the domain wall width, the zig-zag shape had no effect on our micromagnetic simulations. We used the following intrinsic magnetic properties for SmFe11Ti: K1 = 4.8 MJ/m 3 and µ0Ms = 1.17 T which are both equal to the measured values for single crystals shown above. This yields an anisotropy field of µ0Ha = 10.31 T. For the exchange constant we assumed A = 10 pJ/m. For the twin boundary phase, as in agreement with the APT results, we assume a monoclinic ordered Sm3(Fe,Ti)29 phase. For a monoclinic system the magneto-crystalline anisotropy energy density K1sin²θ+K2sin²θ cos2φ depends not only on the polar angle θ, but also on the azimuthal angle φ. For Sm3(Fe,Ti)29 the anisotropy constants are K1 = 1.8 MJ/m 3 , K2 = 0.4 MJ/m 3 and the magnetization is µ0Ms = 1.13 T The two easy-axes of the two variants are shown as white double-headed arrows and are parallel to the yz-plane. The angle between the two easy-axes is 122°. The direction of the external field is given in the first bracket of the legend entries by spherical coordinates (θ,φ) where θ is the polar angle (z->x) and φ is the azimuthal angle (x->y). The second bracket gives where αi is the angle enclosed by the external field direction and one of the two easy-axes of the twin variants.

is proportional to the magnetocrystalline anisotropy energy density. Color changes from green to purple in 10° steps of the average α. Solid, dashed and dotted lines group similar rotation directions of the external field from z to y, z to x, and x to y, respectively. Entries in the legend are sorted by coercive field. a) shows the easy-axis-like demagnetization curves where average α < 61°, b) shows the hard-axis-like curves with average α >= 61°. The dotted black line on the left-hand side marks 40 percent of the anisotropy field of SmFe11Ti.

The orientation of the external field is given in spherical coordinates (θ,φ) where θ is the polar angle (z→x) and φ is the azimuthal angle (x→y). The angle between the applied field and the magnetocrystalline easy-axis is given by αi where i refers to one of the two variants of the SmFe11Ti phase. α is the average angle given by = ( + ) 2

⁄ . The anisotropy energy density is given by = . The switching field is defined as the first kink in the demagnetization curve when the behavior changes from reversible rotation to irreversible nucleation of reversed domains.

The micromagnetic simulations for a Nd2Fe14B grain without a defect (graphs not shown) give switching field (µ0Hsw) values that vary between -5.42 T for θ = 1° and -3.23 T for θ = 45°. This corresponds to 71 percent or 42 percent of the anisotropy field of Nd2Fe14B, respectively. The switching field of all other computed orientations of the external field were between those values, consistent with Stoner-Wohlfarth results mentioned above, and also considering selfdemagnetizing effects.

In contrast, despite the higher anisotropy field of the SmFe11Ti phase, much lower switching fields were found for SmFe11Ti, due to the twin defect. Figure 8 shows demagnetization curves for various orientations of the external field computed for a twinned SmFe11Ti grain split by a single 6 nm thick Sm3(Fe,Ti)29 twin boundary. Irreversible switching always starts at fields below 40 percent of the anisotropy field of SmFe11Ti. Switching fields of between -3.94 T and -0.55 T were found for external field directions of (90°, 90°) and (0°, 0°), respectively. This corresponds to 38 percent and 5 percent of the anisotropy field of SmFe11Ti, respectively. Please note that the curvature seen in the hard axis loop of Figure 2 of the single crystal measurements could not be reproduced by our model, see (90°, 0°) in Figure 8. Similarly, we did not observe any curvature in the hard axis loop for a reference calculation with zero magneto-crystalline anisotropy (K1 = 0) in the twin boundary phase. The contribution of the twin boundary phase to the total magnetization is negligible as its volume fraction is only 2.4 percent of the total volume. A possible reason for the curvature might be canting between the Fe and the Sm sublattices. The switching field strongly depends on the field angle. The magnetocrystalline anisotropy energy density is proportional to sin . The lowest switching fields are found if the average α is greater than 61°.

Due to the small thickness of the Sm3(Fe,Ti)29 phase at the twin boundary, it is not clear if the assumption of a uniaxial magnetocrystalline anisotropy is valid. Thus, we also considered the twin boundary phase to be a soft defect with K1 = 0. The switching fields were then in the range of -3.56

T to -1.3 T. In this case, the spread of the switching field values with field angle was reduced from 35 percent to 13 percent of the anisotropy field of SmFe11Ti.

If we reuse the geometric model of the twin and change the setup of the Nd2Fe14B single grain simulations to consider a 6 nm wide Nd20Fe80 ferromagnetic grain boundary phase (K1 = 0, μ0Ms = 0.93 T, A = 9.99 pJ/m [106]) but keep the easy-axes of both grains next to the boundary parallel to z, the switching field is also reduced to a range of 35 percent to 26 percent of the Nd2Fe14B anisotropy field. As mentioned before, in Nd2Fe14B permanent magnets, Nd-rich grain boundary phases with low or no magnetization are believed to magnetically decouple grains and, therefore, increase coercivity in an isotropic granular bulk magnet. Even if a similar grain boundary phase was found for SmFe11Ti, the defect due to twinning within the grains remains, and will act as a weak spot for nucleation of reversed domains. The computation of the magnetization distribution at the saddle point for thermal switching identified the twin boundary region as the location where the nucleus of reversed magnetization forms.

Summarizing at this point, even if a similar grain boundary phase was found for SmFe11Ti, the defect due to twinning within the grains remains, and will act as a weak spot for nucleation of reversed domains. The computation of the magnetization distribution at the saddle point for thermal switching identified the twin boundary region as the location where the nucleus of reversed magnetization forms for the geometry shown in Figure 8. It is well known that ferromagnetic exchange coupling between crystallites with different orientations of the easy axes reduces the coercive field [107]. This holds for misoriented, exchange coupled grains as well as for twinned grains with a well-defined orientation relationship between the easy axes of two variants of the twinned grains. From a micromagnetic point of view, in both cases the formation of a partial domain wall causes the nucleation of reversed domains. Table 1 summarizes the micromagnetic simulations giving the hard magnetic phase, the model setup, and the range of the switching field due to the orientation of the external field.

Table 1. Range of the switching field due to the orientation of the external field for different geometries obtained by micromagnetic simulations. The columns refer to the main hard magnetic phase, the micromagnetic model setup and the coercivity range relative to the anisotropy field of the hard magnetic phases. The volume of the polyhedral grain was equal to that of a cube with an edge length of 295 nm. the interface energies play a significant role in twin formation.

Material

Summary & Conclusions

We carried out detailed structural characterization of twins in "single crystal" and polycrystalline bulk samples of composition SmFe11Ti as well as ternary Sm(Fe,Ti)12 and Zr-substituted (Sm,Zr)(Fe,Ti)12 thin films. Twins were observed at micro-and nano-length scales, and it was found that the twin orientation in all sample types is approximately 57±2 • . Thus twins appear to be a generic feature of the Sm(Fe,Ti)12-system. The twin boundary plane corresponds to a {011}plane family in the unit cell of Sm(Fe,Ti)12. The significantly higher density of twins in the thin films compared to the bulk samples may be due to substrate mediated stress. High resolution electron microscopy studies on polycrystalline bulk samples show an enrichment of Sm and a depletion of Ti at the twin boundary. The composition of the twin boundary is determined to be Sm9.6Fe84.9Ti5.5, matching the composition of the Sm3(Fe,Ti)29 phase, which has a lower anisotropy field. Modification of the intrinsic magnetic properties of the Sm3(Fe,Ti)29 phase may lead to an enhancement of the coercivity of the Sm(Fe,Ti)12-systems.

Micromagnetic simulations were carried out with a microstructure reproducing that of bulk samples, and the results show that the formation of twins reduces the estimated coercivity values to 38% of the expected switching field. The simulation results account for the observation of low coercivity values in twinned SmFe11Ti systems which confirms the detrimental effect of twins.

Additionally, the fixed angle between neighboring twin variants would hinder obtaining perfect texture which also would reduce the remanent magnetization significantly.

Formation of twin boundaries in grains larger than 1 micrometre can be attributed to the accumulation of internal stresses during solidification, where twin formation is favoured over the formation of specific grain boundaries. The specific size, morphology and interface energy dependency and the resulting detailed model of the proposed mechanism will be an object of a further study/ publication. However, tuning hard magnetic alloys with the intent to render them less prone to twin formation is challenging. Three approaches, having different probabilities of a successful outcome, are conceivable in that context. The first one lies in substitutional compositional modification of the alloy with the aim to increase the twin formation energy which in turn would reduce the tendency for twin formation. However, most of the known reported compositional variants of these materials show twins, irrespective of substitutional variations. Also, twin interface energies are generally so low compared to those of less coherent grain boundaries (i.e. about a factor of 10 smaller) that such effects do not sufficiently alter this relation to reduce their frequency.

The second option lies in exposing the material to heat treatment under stress. This measure bears the risk though that imposing stresses on systems that are already prone to twin formation could lead to more rather than less twins. However, deformation driven twinning is typically reduced at elevated temperatures, where energy dissipation by dislocation activity is often promoted over twin formation, so that this option might be worth exploring.

A third and rather promising pathway is grain size reduction as this change of volume-to-surface energy ratio decreases the need for stress relief via twin formation. For this reason, high coercivities can possibly be achieved in nanocrystalline systems like hot-pressed and die-upset ThMn12-type bulk permanent magnets but are unlikely to be observed in their sintered microcrystalline counterparts.

  magnetic characterization along three different directions: [100], [110] and [001]. Magnetic measurements were performed at room temperature using a PPMS-VSM (Quantum Design PPMS-14) under external applied magnetic fields of up to 14 T. For investigation of magnetic domain structures by the MOKE technique, two twinned single crystals were used. The observation surface of one crystal was polished perpendicular to the [100] direction (in-plane magnetization) and that of the other perpendicular to the [001] direction (out-of-plane magnetization).

Figure 2 .

 2 Figure 2. Comparison of room temperature magnetic measurements of Nd2Fe14B (blue) and SmFe11Ti (red) twinned "single-crystals" along easy and hard magnetization directions. The dashed red line indicates the magnetization behavior of a twin-free single crystal along the [100] direction [95].

Figure 3 .

 3 Figure 3. Comparison of room temperature Kerr microscopy images of (a,b) Nd2Fe14B and (c,d) SmFe11Ti twinned single-crystals along easy and hard directions.

Figure 4 .

 4 Figure 4. Comparison of room temperature Kerr microscopy images of polycrystalline (a) Nd2Fe14B and (b) SmFe11Ti bulk samples. Red arrows indicate some of the twins observed in the SmFe11Ti sample.

Figure 5 .

 5 Figure 5. (a) Inverse pole figure image of the investigated polycrystalline SmFe11Ti bulk sample derived from EBSD SEM data. (b) Measured angle between two neighboring twin variants and corresponding schematic representation of the unit cells. (c) A BSE SEM image of a twinned grain with the location of the APT specimen identified by a red triangle.

Figure 6 .

 6 Figure 6. (a) Low magnification HAADF image of SmFe11Ti at the twin boundary region. (b) High-resolution HAADF image of the interface region of the twin along the [100] zone axis. (c) High magnification of the twin boundary region and (d) schematic representation of the 011 twinning planes.

Figure 7 .

 7 Figure 7. (a) Bright-field image of one of the investigated APT specimens. (b) Atom probe tomography reconstruction of the elemental distribution of Sm, Fe and Ti. (c) 2D compositional map of Sm and (d) chemical composition distribution at the twin boundary region.

[

  102]. For our simulations we ignored the contribution of K2 and assumed that the Sm3(Fe,Ti)29 easy-axis is aligned with one of the easy-axes of two variants of the main phase [102-105].

Figure 8 .

 8 Figure 8. Demagnetization curves as a function of the direction of the external field computed for a single twinned SmFe11Ti grain where a 6 nm thick Sm3(Fe,Ti)29 twin boundary separates the two

Acknowledgments

This work is partly based on results obtained within the future pioneering program Development of magnetic material technology for high-efficiency motors (MagHEM) commissioned by the New Energy and Industrial Technology Development Organization (NEDO) of Japan. This work was also supported by the Deutsche Forschungsgemeinschaft (DFG, German Research Foundation) -Project-ID 405553726 -TRR 270.


Monika Kurela
email: monikakurela@gmail.com

Jean-Michel Le Quere

Krzysztof Kołodziejczyk

Agnes Gacel

Sophie Guilois-Dubois

Sylvain Guyot

Effect of phenolics concentration on apple juices and French ciders colors

INTRODUCTION

The color of apple juices and ciders is an important criterion that may influence choice and acceptance of the consumers. This color is the result of polyphenol oxidation that occurs during fruit processing. It is catalyzed by polyphenoloxidase (PPO) in the presence of oxygen mainly during the crushing and the pressing of the raw material. Interestingly, apple varieties show a great diversity regarding their polyphenol profiles 1 . In this context, the purpose of the study was to investigate the impact of different proportions and concentrations of the major classes of polyphenols of apple juices in the formation of the colored and noncolored oxidation products. This investigation will help the producers to control and predict the color of the final product (apple juice or cider) depending on the composition of the raw material they used for juice and cider making. Regarding literature data, the main hypotheses were: 1) the color will decrease with the increase of procyanidins concentration, and 2) the color will increase with the increase of chlorogenic acid, epicatechin and phloridzin concentrations 2,3.

1. Sylvain Guyot, Nathalie Marnet, Philippe Sanoner, Jean-Francois Drilleau. Variability of the Polyphenolic Composition of Cider Apple (Malus domestica) Fruits and Juices. J. Agric. Food Chem., 2003,51 (21),

MATERIALS METHOD

RESULTS AND DISCUSSION

BIBLIOGRAPHY 1) Epicatechin had no marked effect on the formation of the color; 2) Procyanidins have the highest negative impact on each of the parameters (C,L,h) and appeared to be the main inhibitors of the color formation; 3) Phloridzin is important for creation of the color and shows an interaction with chlorogenic acid. 4) The intensity of the colour is increasing with the time

Conclusions

This work was a part of the regional project CICHROM and we are grateful to Region Bretagne and Région Pays de la Loire for financial support and to P2M2 analytical platform for polyphenol analysis.

AKNOWLEDGEMENT

Research were carried out using Central Composite Design method which the first stage were Factorial Design, with considering 4 factors (procyanidins, chlorogenic acid, epicatechins and phloridzin). Each factor was coded in five levels (-2, -1, 0, 1, 2) of concentrations that were chosen to cover the range of French cider apple varieties. Finally, 36 model solutions were prepared and the incubation was started by incorporating a crude concentrated PPO extract obtained from a cider apple variety (i.e. Kermerrien). Three incubation times were chosen (0, 10, 120 min.) stopped with NaF solution. At the end of each incubation time, aliquots were withdrawn and filtered. The 3.

Variation of Chroma Variation of Hue angle

Variation of Luminance

Change of the colour intensity in time

  color of the solutions was characterized by measuring the CIE L a b parameters with a spectrophotometer (SPECTROstar Nano, BMG Labtech, Ortenberg, Germany -photo on the right). Data were processed by "Statgraphic" software. Results have shown the correlation between factors and colour parameters. a) A variety of model solutions with polyphenols made of 4 mother solutions (1. procyanidins, 2. chlorogenic acid, 3. epicatechins and 4. phloridzin). b) Malate buffer : 402 mg of malic acid per 100 ml ultra-pure water c) PPO solution: 275 µl of crude extract per 5 ml of buffer (without O 2 ) d) NaF (inhibitor of PPO solution): 1 g/L; 200 µl-1 ml per tube, freezed


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Synthesis of N-phenyltrifluoroacetimidoyl chloride (CAS: 61881-19-4) via chlorination of 2,2,2-trifluoro-N-phenylacetamide using triphenylphosphine oxide/oxalyl chloride combination system

Yu Tang

To cite this version:

Yu Tang. Synthesis of N-phenyltrifluoroacetimidoyl chloride (CAS: 61881-19-4) via chlorination of 2,2,2-trifluoro-N-phenylacetamide using triphenylphosphine oxide/oxalyl chloride combination system.

hal-03325829

Synthesis of N-phenyltrifluoroacetimidoyl chloride (CAS: 61881-19-4) via chlorination of 2,2,2-trifluoro-N-phenylacetamide using triphenylphosphine oxide/oxalyl chloride combination system N-phenyltrifluoroacetimidoyl chloride (CAS: 61881-19-4) is a versatile fluorinecontaining reagent that have find wide applications, especially in the synthesis of glycans 1 . A recent report of its synthesis employing dichlorotriphenylphosphane instead of CCl 4 /PPh 3 combination as the chlorinating reagent 2 attracted our attention. Inspired by this work, we further attempted using triphenylphosphine oxide/oxalyl chloride combination 3 for the in-situ generation of dichlorotriphenylphosphane that serve as the chlorinating reagent. Under this condition, the chlorination reaction proceeds smoothly and N-phenyltrifluoroacetimidoyl chloride was formed in high yield as determined by 19 F-NMR analysis. After the reaction, triphenylphosphine oxide can be quantitatively recovered from the reaction mixture.

Experimental Procedure

A 50 mL, three-necked, round-bottomed flask equipped with a teflon-coated magnetic stir bar and a reflux condenser was charged with triphenylphosphine oxide (2.00 g, 7.19 mmol) and a solution of 1 (1.00 g, 5.29 mmol in 10 mL MeCN) was slowly added under argon atmosphere. The mixture was stirred at RT and oxalyl chloride (1.00 mL, 11.8 mmol) was added dropwise. After the addition, the mixture was stirred for 10 min, then Et 3 N (1.60 mL, 11 mmol) was added. The mixture was heated at reflux for 4 h, 19 F NMR of the reaction mixture revealed that the reaction was complete and 2 was formed in high yield (-77.1 ppm in MeCN). The mixture was cooled to RT, filtered, washed with PE (20 mL) and evaporated under reduced pressure to afford a brown solid, which was extracted with PE (20 mL×3) and filtered. The filtrate was evaporated under reduced pressure to afford crude product, which was further purified by silica gel flash column chromatography (PE) to afford 2 as a colorless oil (280 mg, 25%). The NMR spectral data of 2 was identical with the literature data 1 . The relatively lower yield was most likely due to the relatively low boiling point of the product.

In order to recover triphenylphosphine oxide from the reaction mixture, the above solid obtained after the filtration process was suspended in water (500 mL), filtered and dried to afford crude triphenylphosphine oxide (2.08 g) in essentially quantitative yield.

Acknowledgement

The author would like to thank Prof. Biao Yu at SIOC for his helpful guidance.

Financial support from the Youth Innovation Promotion Association of CAS (2021251) is acknowledged.

L'

Author Information

Yu Tang E-mail: tangyu@sioc.ac.cn https://orcid.org/0000-0002-4272-2234

State Key Laboratory of Bioorganic and Natural Products Chemistry, Shanghai Institute of Organic Chemistry, Chinese Academy of Sciences, 345 Lingling Road, Shanghai 200032, China