Tangi Aubert

Ljiljana Palangetic

Mohammad Mohammadimasoudi

Kristiaan Neyts

Jeroen Beeckman

Christian Clasen

Zeger Hens

Large scale and electro-switchable polarized emission from semiconductor nanorods aligned in polymeric nanofibers

Keywords: quantum dots, microemulsion, electrospinning, liquid crystal, display, LED TOC

dynamics, [START_REF] Nobile | Self-assembly of Highly Fluorescent Semiconductor Nanorods into Large Scale Smectic Liquid Crystal Structures by Coffee Stain Evaporation Dynamics[END_REF] impregnation of 3D templates, [START_REF] Lutich | Macroscale Alignment of CdSe/CdS Nanorods by Porous Anodic Alumina Templates[END_REF] mechanical rubbing of spin-coated NRs layers, [START_REF] Amit | Semiconductor Nanorod Layers Aligned Through Mechanical Rubbing[END_REF] stretching of polymer films [START_REF] Peng | Shape Control of CdSe Nanocrystals[END_REF][START_REF] Talapin | Highly Emissive Colloidal CdSe/CdS Heterostructures of Mixed Dimensionality[END_REF] or electrically driven alignment. [START_REF] Ryan | Electric-Field-Assisted Assembly of Perpendicularly Oriented Nanorod Superlattices[END_REF][START_REF] Gupta | Self-Corralling" Nanorods under an Applied Electric Field[END_REF][START_REF] Hu | Electric-Field-Driven Accumulation and Alignment of CdSe and CdTe Nanorods in Nanoscale Devices[END_REF] Although in most of these techniques the NRs are aligned over areas in the micro-to millimeter range at best, some have already shown to be suitable for relatively large scale alignment. For instance, Amit et al. have obtained a polarization ratio of 0.55 (contrast ratio of 3.5) over a several cm 2 area by mechanical rubbing of CdSe/CdS NRs. [START_REF] Amit | Semiconductor Nanorod Layers Aligned Through Mechanical Rubbing[END_REF] We recently reported a novel technique for the alignment of CdSe/CdS NRs by applying an alternating electric field during a dip-coating procedure achieving a very high polarization ratio of 0.60. [START_REF] Mohammadimasoudi | Fast and Versatile Deposition of Aligned Semiconductor Nanorods by Dip-coating on a Substrate with Interdigitated Electrodes[END_REF] Nevertheless, these techniques can result in a heterogeneous distribution of the NRs over the substrate. In addition, the aligned NRs obtained by these techniques cannot be conveniently manipulated and processed, which would be a great asset for display and lighting applications. As a further requirement, the NRs should also be aligned horizontally in order to be able to collect the polarized emission perpendicularly to the substrate.

In this context, a very interesting approach has been developed for the alignment of gold NRs in polymeric nanofibers made by electrospinning. [START_REF] Roskov | Long-Range Alignment of Gold Nanorods in Electrospun Polymer Nano/Microfibers[END_REF] During electrospinning, the polymer solution undergoes a strong uniaxial extensional flow, which aligns the NRs along the nanofiber axis and prevents rotational diffusion from randomizing the particle's orientation. This technique has proven to efficiently align even NRs with low aspect ratios. [START_REF] Roskov | Long-Range Alignment of Gold Nanorods in Electrospun Polymer Nano/Microfibers[END_REF] Generally, the electrospun nanofibers are collected on a conductive, grounded plate as non-woven mats. The reason for this is the random movement of the charged jet caused by the bending instability which originates from the interactions between the electrostatic charges within the jet, and between the electrostatic charges and the applied electric field. However, using different types of collectors and/or manipulating the electric field enables ordered structures to be formed, which is an additional asset of the technique. [START_REF] Teo | A Review on Electrospinning Design and Nanofibre Assemblies[END_REF] In this work, we employ electrospinning to produce films made of aligned polyvinylpyrrolidone (PVP) nanofibers containing silica coated CdSe/CdS NRs oriented parallel to the nanofiber axis. The uniaxially aligned nanofibers were obtained by simply replacing the standard grounded plate collector with a parallel plate collector, consisting of two free-standing stripes of electrical conductors separated by a gap. When the randomly moving parts of a nanofiber get close to the edges of the conducting plates, they are "pulled" by the electrostatic forces which stretch the nanofiber across the gap and orient it perpendicularly to the parallel stripes. [START_REF] Li | Electrospinning Nanofibers as Uniaxially Aligned Arrays and Layer-by-Layer Stacked Films[END_REF][START_REF] Reddy | Metallic and Bi-Metallic Janus Nanofibers: Electrical and Self-Propulsion Properties[END_REF] Using this type of collector we produced a free standing film of 1.5 cm 2 that showed polarized emission with a large scale polarization ratio of 0.45.

The alignment of NRs in such polymeric films offers important advantages for their use in light emitting devices. With this approach the NRs are aligned within a protective, low refractive index matrix, whereas other techniques often produce rather compact films of NRs which can result in a self-quenching of their emission. [START_REF] Shirasaki | Emergence of Colloidal Quantum-Dot Light-Emitting Technologies[END_REF] In addition, such flexible and freestanding films can be manipulated and further processed for the realization of a working device which we also demonstrate here. Indeed, besides the large scale alignment of NRs we go one step further towards a realistic lighting application through the integration of this film in a liquid crystal (LC) cell. Liquid crystals are electro-optic materials that are commonly used in display applications. In such applications a voltage across the liquid crystal layer is used to change the direction of the liquid crystal, which in turn has an effect on the polarization state of light passing through. It appears that the role of the aligned nanofibers is not only limited to the alignment of the NRs, but can also be used to induce alignment of the nematic LC. With this approach we have fabricated a simple device emitting polarized light over a large area in which the initial direction of polarization can be switched over 90°.

CdSe/CdS dot-in-rods. The CdSe/CdS NRs were synthesized according to a procedure described in the literature (see the Supporting Information, paragraph 1 for the complete synthesis details). [START_REF] Carbone | Synthesis and Micrometer-Scale Assembly of Colloidal CdSe/CdS Nanorods Prepared by a Seeded Growth Approach[END_REF] The first step is the synthesis of spherical CdSe QDs with a wurtzite crystalline structure. The QDs are then used as the cores for the seeded growth of an anisotropic CdS shell, resulting in a so-called dot-in-rod heterostructure. The CdSe/CdS NRs involved in this work have an average diameter of 4.1 nm and an average length of 29.3 nm (see the Supporting Information, Figure S1 for a transmission electron microscopy (TEM) image). The NRs show an emission spectrum centered at 600 nm (see the Supporting Information, paragraph 2 and Figure S2 for the optical characterization of the NRs).

Silica encapsulation of the NRs. The CdSe/CdS NRs are typically dispersed in apolar solvents, such as toluene or chloroform, which are not suitable for electrospinning due to their incompatibility with the polymers typically used for electrospinning or their fast evaporation rate. In order to transfer them to a polar solvent (ethanol), the NRs were coated with a thin silica shell through a reverse microemulsion process. This technique has already proven to be very efficient for the silica coating of initially hydrophobic nanocrystals, including CdSe/CdS dot-in-rods. [START_REF] Pietra | Synthesis of Highly Luminescent Silica-Coated CdSe/CdS Nanorods[END_REF] Here, 10 nmol of NRs were mixed with 100 mL of n-heptane (99%, VWR) and 32 mL of polyoxyethylene (4) lauryl ether (Brij L4, formerly known as Brij 30, Sigma-Aldrich). After 15 min under magnetic stirring, 5 mL of MilliQ H2O and 0.5 mL of NH4OH (28% in water, VWR) were slowly added to form the microemulsion. After 1 hour under stirring, 0.25 mL of tetraethyl orthosilicate (TEOS, 98%, Sigma-Aldrich) was added to the microemulsion. The reaction was left stirring for two days to form the silica shell. The CdSe/CdS@SiO2 NRs were then further functionalized with polyethylene glycol (PEG) to provide them with enhanced colloidal stability in polar solvents and to prevent aggregation. PEGylation of the CdSe/CdS@SiO2 NRs was done in a one-pot process, following the growth of the silica shell, by adding 120 mg of methoxy-PEG-silane (mPEG-silane, Mw ≈ 1000 g/mol, Gelest) to the microemulsion. After 3 additional days of stirring, the reaction was stopped by adding a large volume of ethanol, resulting in the destabilization of the microemulsion and precipitation of the particles which were collected by centrifugation. The PEGylated CdSe/CdS@SiO2 particles were further purified twice with n-heptane and redispersed in ethanol. Figure 1a shows the TEM image of the CdSe/CdS@SiO2 NRs evidencing that a homogenous silica coating of about 6 nm in lateral thickness is formed. The rod shape of the objects is largely preserved, albeit with a reduced aspect ratio. The initial aspect ratio of 7.1 of the CdSe/CdS NRs is reduced to 2.4 for the CdSe/CdS@SiO2 NRs. Figure 1a also shows that some of the objects actually consist of two CdSe/CdS NRs connected within the silica matrix, resulting in some cases in V shapes with an angle of 110°. The connections between rods could not be avoided in our experiments and are believed to occur already when the CdSe/CdS NRs are in their apolar solvent prior to the silica encapsulation. This assumption is supported by TEM observations of the CdSe/CdS NRs deposited on a Formvar layer prior to their silica encapsulation (see the Supporting Information, Figure S3). Electrospinning of the nanofibers. The NRs were mixed with a polymer by adding 240 mg of polyvinylpyrrolidone (PVP, Mw = 1,300,000 g/mol, Sigma-Aldrich) to a solution containing 4.65 nmol of the CdSe/CdS@SiO2 NRs in 2.81 g of ethanol. Even though the NRs were coated with silica, we avoided using water since this is a rather harsh environment that tends to deteriorate the optical properties of the semiconductor NRs. [START_REF] Pietra | Synthesis of Highly Luminescent Silica-Coated CdSe/CdS Nanorods[END_REF] The relatively high molecular weight has been chosen for the ease of nanofiber formation by electrospinning. [START_REF] Palangetic | Dispersity and Spinnability: Why Highly Polydisperse Polymer Solutions are Desirable for Electrospinning[END_REF] The nanofibers were made using a climate controlled electrospinning chamber (EC-CLI) from IME Technologies. The temperature and relative humidity were set to 25 °C and 30% respectively. The flow rate was 0.55 mL/h (controlled with a HARVARD syringe pump). The tip-to-collector distance was 15 cm, and the voltage was set to 14 kV. The nanofibers were electrospun from a needle with an inner diameter of 0.25 mm and an outer diameter of 0.52 mm. In order to obtain aligned nanofibers, a parallel plate collector was used. [START_REF] Reddy | Metallic and Bi-Metallic Janus Nanofibers: Electrical and Self-Propulsion Properties[END_REF] The collector consists of two aluminum stripes separated by a void gap of 1 cm wide and 1.5 cm long. In order to ensure that the nanofibers are deposited within the desired area and that they have support, the gap and the space around the stripes have been covered with Teflon sheets. With this type of collector, the electrostatic forces generated by the two conductive stripes induce a uniaxial alignment of the charged spun nanofibers across the insulating gap. [START_REF] Li | Electrospinning Nanofibers as Uniaxially Aligned Arrays and Layer-by-Layer Stacked Films[END_REF] The collected nanofibers, denoted as CdSe/CdS@SiO2@PVP, show a high degree of alignment as evidenced by scanning electron microscopy (SEM) observations (Figure 1b). Electrospinning was carried out up to 5 min, producing a free standing film of 1.5 cm 2 that could then be manipulated (Figure 1c). Nevertheless, the electrospinning technique used here can easily be adapted to produce even larger films by increasing the length of the stripes. The TEM image of a single nanofiber in Figure 1d shows a relatively high load of CdSe/CdS@SiO2 NRs embedded in the polymeric nanofiber. Moreover, this image already evidences that the NRs are indeed aligned along the axis of the nanofiber. Even NRs with a V shape are found to be oriented at their best in the nanofibers, resulting in a maximum angular deviation of 35° with respect to the nanofiber axis. The TEM image in Figure 1d also shows that the NRs are evenly distributed within the fiber. The NRs do not appear to aggregate (keeping in mind that TEM observations provide only a 2D projection), which confirms the good colloidal stability of the silica coated NRs in the polar polymeric solution and the good affinity of the polymer for their PEGylated surface. The absence of aggregation is an important parameter to avoid selfquenching of the NRs emission in lighting applications. 17

Large scale polarized emission

To measure polarized emission with a fluorescence microscope, the film of aligned CdSe/CdS@SiO2@PVP nanofibers was detached from the collector and placed between two glass slides. The nanofibers were excited with green light, using a xenon lamp with a band pass filter (510-560 nm), in order to avoid emission from the PVP or the LC which also absorb in the UV range (see the Supporting Information, Figure S4 for their optical characterization). The emitted light was collected by an objective (10×) after passing through a rotatable linear polarizer and was detected with an Andor CCD camera. One should note that in our experiments the excitation light is not polarized. This is preferred in order to avoid preferential absorption by the NRs that have the best alignment and allow for a better assessment of the true polarization ratio of the sample. Figure 2 gives the fluorescence microscopy images and integrated intensities depending on the polarizer angle. This angle is defined as the difference between the nanofibers long axis and the polarizer axis. As clear evidence of the alignment of the NRs inside the nanofibers and of the parallel alignment of the nanofibers, a bright image is obtained when the polarizer is parallel to the nanofibers (Figure 2a) and a dark image is obtained when the polarizer is perpendicular to the nanofibers (Figure 2b). Figure 2c further evidences that the variation in intensity, when rotating the polarizer, follows a classic cos 2 function. [START_REF] Lutich | Macroscale Alignment of CdSe/CdS Nanorods by Porous Anodic Alumina Templates[END_REF][START_REF] Amit | Semiconductor Nanorod Layers Aligned Through Mechanical Rubbing[END_REF][START_REF] Mohammadimasoudi | Fast and Versatile Deposition of Aligned Semiconductor Nanorods by Dip-coating on a Substrate with Interdigitated Electrodes[END_REF] From these results one can calculate the polarization ratio (PR) defined by: PR = (I∥ -I⊥)/(I∥ + I⊥), where I∥ and I⊥ are the integrated intensities when the polarizer is parallel and perpendicular to the nanofibers respectively, i.e., at 0° or 90°. One should keep in mind that the polarization ratio of a single CdSe/CdS NR with a spherical core is limited to 0.75, so that a polarization ratio of 1 cannot be achieved. [START_REF] Talapin | Highly Emissive Colloidal CdSe/CdS Heterostructures of Mixed Dimensionality[END_REF][START_REF] Pisanello | Dots in Rods as Polarized Single Photon Sources[END_REF][START_REF] Hadar | Polarization Properties of Semiconductor Nanorod Heterostructures: From Single Particles to the Ensemble[END_REF] From the large scale images shown in Figure 2a andb, we obtain a polarization ratio of 0.45. For measurements on individual nanofibers, the highest polarization ratio was found to be 0.55 (see the Supporting Information, Figure S5 for fluorescence microscopy images of a single fiber). This indicates that the polarization is relatively well preserved when aligning the nanofibers over a large area. The few non-aligned nanofibers, which become visible in Figure 2b when the polarizer is in the 90° position, have a negligible contribution to the overall integrated intensity. The 0.45 polarization ratio obtained here has the same order of magnitude as obtained with other techniques for the alignment of NRs at large scale. This is comparable for instance with the 0.34 obtained by Rizzo et al. through solvent evaporation at a water/air interface, [START_REF] Rizzo | Polarized Light Emitting Diode by Long-Range Nanorod Self-Assembling on a Water Surface[END_REF] or the 0.55 obtained by Amit et al. through mechanical rubbing. [START_REF] Amit | Semiconductor Nanorod Layers Aligned Through Mechanical Rubbing[END_REF] Although we have previously demonstrated that polarization ratios up to 0.6 could be obtained through electrical-driven alignment, [START_REF] Mohammadimasoudi | Fast and Versatile Deposition of Aligned Semiconductor Nanorods by Dip-coating on a Substrate with Interdigitated Electrodes[END_REF] the present approach offers important advantages with regard to scale-up and further processing as demonstrated in the following section. In addition, the polarization ratio of 0.45 obtained here can most likely be improved, especially if one can avoid the connection between NRs. It should also be noted that the excitation wavelengths used in the experiment (510-560 nm) actually correspond to absorption by the CdSe cores in the CdSe/CdS NRs (see the Supporting Information, Figure S2 for the absorption spectrum of the NRs). This excitation range was initially chosen in order to avoid emission from the PVP and LC. However, additional experiments using excitation in the UV range (330-380 nm) also resulted in a polarization ratio of 0.45 (see the Supporting Information, Figure S6 for fluorescence microscopy images under UV excitation). This observation is in line with previous reports indicating that in such heterostructures the dipole momentsresponsible for the polarized emission propertyare aligned in the same direction in the CdSe core and in the CdS shell. [START_REF] Lutich | Macroscale Alignment of CdSe/CdS Nanorods by Porous Anodic Alumina Templates[END_REF] Figure 2. Fluorescence microscopy images of the aligned CdSe/CdS@SiO2@PVP nanofibers with the polarizer oriented (a) parallel (0°) and (b) perpendicular (90°) to the nanofibers. (c) Normalized integrated intensities depending on the polarizer angle (black markers) fitted with a cos 2 function (red line).

Integration in a liquid crystal cell

One of the advantages of this method of aligning the NRs in polymeric nanofibers is that the resulting large scale flexible film can be conveniently manipulated and further processed. Here, a film of aligned CdSe/CdS@SiO2@PVP nanofibers has been integrated in a liquid crystal cell which is illustrated in Figure 3. The cell consists of two glass substrates coated with indium-tin-oxide (ITO) layers as the electrical contacts. The bottom substratei.e. on the side where emission is collectedis also coated with a 300 nm thick layer of mechanically rubbed nylon-66. The film of aligned nanofibers was placed in the cell so as to have a 90° angle between the alignments of the nylon layer and the nanofibers. The cell was then sealed with glue containing spacer beads of 10 µm (Sekisui Chemicals) to ensure a uniform cell gap. An opening was left on the side of the cell to infiltrate it with a commercial E7 (Merck) LC. The director of the liquid crystal follows the alignment of the nanofibers and of the nylon at both sides and inbetween it makes a twist of 90°. This device was then analyzed by fluorescence microscopy under green excitation, with the emitted light passing through a rotatable polarizer. The different experimental conditions and the resulting fluorescence microscopy images are shown in Figure 3. The microscopy image in Figure 3a evidences that almost no light is collected when the polarizer is oriented parallel to the nanofibers (0° position). Since both the NRs and the polarizer are oriented in the same direction, this can only be explained by a 90° rotation of the initial polarization direction of the NR emission. The 90° rotation of the polarization direction is caused by the fact that the linear polarization direction is following the twist of the LC director. Indeed, for the parameters in this work the twist of the director is slow and the change of polarization is according to the Maugain regime as demonstrated in the Supporting Information (paragraph 7). [START_REF] Yeh | Optics of Liquid Crystal Displays[END_REF] It confirms that the nematic phase of the LC is aligned by the nanofibers. Aligned nanofibers have already been reported to orient twisted nematic LC phases as efficiently as rubbed polymer layers. [START_REF] Toan | Director Orientation of Nematic Liquid Crystal using Orientated Nanofibers Obtained by Electrospinning[END_REF] This could be explained by the shape of the nanofibers, as waveguides and edges have been reported to align LCs. [START_REF] Desmet | Modeling Nematic Liquid Crystals in the Neighborhood of Edges[END_REF] In addition, the long polymer molecules within the nanofibers undergo a coil-stretch transition in the extensional flow of the spinning process, [START_REF] Clasen | Capillary Breakup Extensional Rheometry of Semi-Dilute Polymer Solutions[END_REF][START_REF] Vadillo | Microsecond Relaxation Processes in Shear and Extensional Flows of Weakly Elastic Polymer Solutions[END_REF] of which a residual alignment along the nanofiber axis has been reported in the solid nanofiber state, [START_REF] Fong | Elastomeric Nanofibers of Styrene-Butadiene-Styrene Triblock Copolymer[END_REF][START_REF] Catalani | Evidence for Molecular Orientation and Residual Charge in the Electrospinning of Poly(butylene terephthalate) Nanofibers[END_REF] and could also contribute to the LC alignment similarly to a rubbed polymer layer. The fact that the same template matrix can be used for the alignment of the NRs and of the LC constitutes a considerable added value of this approach. In Figure 3b, the polarizer is rotated perpendicularly to the nanofibers (90° position) and therefore lets the polarized light emitted by the NRs go through. This is the "on" position. Here again a polarization ratio of 0.45 is found between the fluorescence microscopy images of Figure 3a and 3b, confirming that the large scale polarized emission is not affected by the integration of the aligned nanofibers film in the LC cell. In Figure 3c an AC electric field (sine wave, 50 V, 1 kHz) is applied over the ITO layers, resulting in a change of the orientation of the LC director perpendicular to the substrates, along the electric field. As a consequence, the polarization of the light emitted by the NRs is not rotated anymore and is blocked by the polarizer. This is the "off" position. Here, a polarization ratio of 0.43 is found, which has been calculated using I0V (Figure and I50V (Figure 3c) instead of I∥ and I⊥ respectively. This value is very close to the 0.45 of the aligned nanofibers film itself, which shows the high efficiency of the electrical switching of this polarized light emitting device. If the polarizer is rotated to the 0° positioni.e. parallel to the nanofibersthe emitted light is collected again (see the Supporting Information, Figure S7). In this case, the amount of light collected corresponds to 77% of the "on" position, which indicates a relatively good vertical alignment of the LC director in the presence of the 50 V electric field. Nevertheless, the small decrease in collected light still suggests that the vertical alignment is not perfect, which is most likely the case at the vicinity of the alignment layers (rubbed nylon on one side, fibers film on the other). This could result in a distortion of the linearly polarized light originating from the NRs, which then becomes slightly elliptical. In conclusion, we have successfully aligned silica coated CdSe/CdS NRs in PVP nanofibers by electrospinning, in spite of the relatively low aspect ratio (2.4) of the NRs. Aligning the nanofibers themselves parallel resulted in a flexible film of 1.5 cm 2 showing a polarization ratio of 0.45. Such a self-supporting film, similar to a foil, can be manipulated and combined with any kind of substrate, which is an important asset of the approach undertaken here for the realization of a device. Through its integration in a LC cell we have shown that the film of aligned nanofibers can be used to induce alignment of a nematic LC. Finally, we have demonstrated that this simple device, combining aligned nanofibers and liquid crystal, constitutes an efficient way of electrically switching on and off polarized light while fully preserving the polarization ratio over a large area. The strategy developed here can readily be applied to a large variety of colloidal objects, taking advantage of the high versatility of the electrospinning technique, and paving the road for the design of novel devices where the large scale organization of nanosized building blocks matters.

Figure 1 .

 1 Figure 1. (a) TEM image of the PEGylated CdSe/CdS@SiO2 NRs. (b) SEM image of the aligned nanofibers. (c) Photography of the whole aligned nanofibers sample (red arrow indicates the alignment direction of the nanofibers). (d) TEM image of a CdSe/CdS@SiO2@PVP nanofiber.

Figure 3 .

 3 Figure 3. LC cell with CdSe/CdS@SiO2@PVP nanofibers and corresponding fluorescence microscopy images depending on the polarizer orientation and electric field. (a) Polarizer is parallel to the nanofibers (0°) and no electric field is applied. (b) Polarizer is rotated perpendicularly to the nanofibers (90°) and no electric field is applied (ON-position). (c) Polarizer is maintained perpendicular to the nanofibers (90°) and a 50 V electric field is applied (OFF-position). Insets of the fluorescence microscopy images in (b) and (c): corresponding pictures of the LC cell using a portable UV lamp for the excitation.

Acknowledgments

The authors acknowledge BelSPo (IAP 7.35, photonics@be), FWO-Vlaanderen (Project G.0760.12; KaN 1509012N), Ghent University (BOF12/GOA/015), ERC-2007-StG starting grant 203043 NANOFIB, and the SoPPom program of the Flemish Strategic Initiative for Materials (SIM) for financial support.

Author information * Corresponding authors: tangi.aubert@ugent.be, zeger.hens@ugent.be $ Present address: Department of Materials, Eidgenössische Technische Hochschule (ETH) Zürich, 8093 Zürich, Switzerland Associated content Supporting Information Available: Synthesis details for the CdSe/CdS NRs, additional TEM and optical characterizations, additional fluorescence microscopy analyses, condition for the Maugain regime. This material is available free of charge via the Internet at http://pubs.acs.org.


Y Taibi

M Poulain

M Legouera

A Mebrek

INVESTIGATION AND PROPERTIES OF NOVEL LOW MELTING GLASSES IN THE TERNARY SYSTEM Sb 2 O 3 -PbO-Me 2 O 3

Keywords: Lead antimony glass, Thermal properties, Density, Hardness

New glasses based on Sb 2 O 3 -PbO association have been synthesized and characterized. The limits for glass formation have been investigated in ternary systems encompassing an oxide of IIIA group elements like Al 2 O 3 , Ga 2 O 3 and In 2 O 3 as the third component. Fast quenching is required to prevent melt crystallization, which results in thin samples. Chemical analysis of elements shows a rather good agreement between nominal and analyzed compositions. Glass transition temperature, Tg, ranges between 255°C and 280°C, depending on composition. Density and hardness have been measured for typical composition.

Introduction

The use of lead oxide as a significant glass component is older than modern chemistry and physics. In 1675 George Ravenscroft defined a new glass composition containing lead oxide, leading to larger working range in the molten state, reduced hardness and higher refractive index, which increases surface brightness [START_REF] Navarro | Constitución, fabricación, propiedades[END_REF][START_REF] Moretti | English lead crystal. The Ravenscroft recipe and points not clear on the process of development of "flint" glass formulation[END_REF].

Heavy Metal Oxide Glasses (HMOG) make a family of special glasses that has been the subject of numerous studies. They may be roughly defined as glasses free of classical vitrifiers SiO 2 , B 2 O 3 , P 2 O 5 [START_REF] Dumbaugh | [END_REF]. These glasses exhibit high refractive index, close to 2, large optical non-linearity and extended optical transmission in the infrared spectrum [4,5,6]. Among HMOG, germanate glasses appear as an intermediate group as GeO 2 is considered as a classical network former according to Zachariasen's rules [7], and also because their refractive index is somewhat lower.

Antimonite glasses based on Sb 2 O 3 have attracted growing interest, mainly because numerous glass forming systems have been identified so far, especially in polyanionic systems: oxyhalides and oxysulfides. In addition their stability against devitrification may be very large. Antimony oxide glasses have been reported in association to alkali oxides and in multicomponent systems [8,9,10].

According to previous reports, PbO-Sb 2 O 3 glasses exhibit good chemical durability and low crystallization rates [11,12]. Addition of the oxides of III A group elements like Al 2 O 3 , Ga 2 O 3 and In 2 O 3 to the PbO-Sb 2 O 3 based glass is likely to ameliorate their physical properties. The choice of these oxides was based on the occurrence of a binary glass forming range in combination to PbO. Table 1 reports the compositional limits of these binary glasses: [13]; 67 -95 mol%PbO [14] 69 -81mol%PbO [15]; 58 -82 mol% PbO [16]; 57 -82 mol% PbO [17]; ˂ 5 mol% PbO [18] The aim of the present work is to explore new glass compositions in multicomponent systems based on antimony oxide and to assess the potential of these glasses

Experimental

Materials used in this research are Sb 2 O 3 (99,9%), PbO (99,9%), In 2 O 3 (98%), Ga 2 O 3 (98%), and Al 2 O 3 (98%) from Aldrich, WWR and Metal Europe. Glass samples were prepared using the standard melt-quenching technique. Crystalline powders were intimately mixed in an agate mortar. Then batch was heated up to melting in glass crucible for 5 min. Finally the melts were cast and quenched between two brass plates to obtain a fast cooling rate close to a few hundred K/S. In these conditions, sample thickness ranges between 0.8 and 1mm.

Main physical characterizations include X ray diffraction, thermal analysis microhardness and density. X-ray diffraction patterns were obtained using Cu Kα radiation on Philips PW1830 difractometer. Differential scanning calorimetric measurements were performed on small samples using a DSC 2010 from TA Instrument in aluminum sealed pans under N 2 atmosphere. Current heating rate was 10 K.min -1 and experimental error was ±2°C for glass transition and onset of crystallization. Chemical composition was checked using Oxford Link ISIS probe in a JEOL 6400 Scanning Electron Microscope dispersive -ray spectrometer working at 20 KV. Density was measured using an ACCUPYC 1330 Micrometrics Helium Picnometer with a precision of about ±0,0005g/cm 3 . The samples were indented with Matsuzawa MXT set-up. The value is averaged from 10 indentations or more, made on each sample.

Results

Glass formation:

Our primary investigation of the binary Sb 2 O 3 -PbO based system confirms that the limits for glass formation are extended, when faster quenching rates are used. These limits lie between 80 and 10% mol Sb 2 O 3 as shown in Fig1.

Prepared glasses are stable at room atmosphere. At visual inspection, they appear transparent and yellow colored for glass samples between 80 and 40 mol % Sb 2 O 3 . As PbO content increases, glass samples appear dark. This coloration arises from lead reduction: Indeed, as shown in figure 2, a metallic phase area corresponding to pure lead was observed in a dark glass sample. In practice, melt must be kept at room atmosphere for a longer time -e.g. 8 mnto ensure oxidizing atmosphere.

In order to confirm glass composition, elemental analysis was performed for some samples with nominal composition (100-x) Sb 2 O 3 .xPbO ( x= 50,60,70,80,90). The analytical results reported in Table 2 show a rather good agreement between nominal and analyzed Sb concentration, while lead concentration is smaller than expected. Silicon content that arises from crucible increases as PbO concentration and melting time increase. For all these systems, the experimental vitreous area is very narrow. For example, In 2 O 3 content should not exceed 5 mol % while it may reach 15 mol % with Al 2 O 3 and Ga 2 O 3 . Sample color remains yellow fro the ternary glasses containing Al 2 O 3 and Ga 2 O 3 but it turns to green for the largest In 2 O 3 contents.

X Ray Diffraction

The XRD patterns confirm the amorphous character of the samples, as shown in figure 4 for typical glass compositions. These patterns show a broad main peak at 2ϴ = 27,9°. Beyond this peak, a second one is observed at 2ϴ = 54°.

Thermal Properties:

Binary system

Differential Scanning Calorimetry has been implemented between 100 and 500°C. Typical DSC scans are reported in figure 5. The evolution of the characteristic temperatures corresponding to glass transition (Tg), onset of crystallization (Tx) and exotherm maximum (Tp) versus concentration is plotted in figure 6(a) The stability factors S [19] and T=Tx-Tg [20] that give an estimate of glass forming ability appears in figure 6(b). Glass transition varies between 255 and 276°C. As one could expect, characteristic temperatures depends on Sb 2 O 3 content. These results suggest that glass stability against devitrification is maximum for high Sb 2 O 3 concentration (> 70 %) and for large lead concentration (90 % PbO). Intermediate compositions appear more prone to devitrification.

Ternary system

For comparison, thermal measurements have been carried out on a set of three samples with the same general, formula: x Sb 2 O 3 -(95-x) PbO-5Me 2 O 3 (Me = Al, Ga, In, and 10< x <80) The composition dependence of glass transition temperature (Tg) is reported in fig 7 . The curves expressed the same behavior as observed in the binary system for glass transition: Tg increases for large PbO concentrations (> 40%). In a general way, Tg increases with ionic radius (Al 3+ ˂ Ga 3+ ˂ In 3+ ). We have also prepared samples with different In 2 O 3 contents. The basic composition -50 Sb 2 O 3 -50PbO -was chosen in the range of minimum stability in order to make clear the effect of additive concentration on glass stability. As shown in figure 8 These curves suggest that Me 2 O 3 addition has a stabilizing effect by comparison to the binary glass. At first glance, glasses containing Ga content appear as the most stable ones. However additional experiments should be implemented to confirm these observations: first, no data were collected for Me 2 O 3 concentrations between 5 and 10 mol %; and second, no stable glass was obtained for In 2 O 3 concentrations larger than 5 mol % (see fig. 3). Then one would expect the sharp decrease of the stability factor beyond some threshold value that could be found experimentally.

Density

The composition dependence of the density of the binary glass samples in the Sb 2 O 3 -PbO system is shown in fig 11 .   As one could expect, density d decreases as lead oxide is replaced by antimony sesquioxide. This change is approximately linear. Another set of density measurement has been implemented in the ternary system with the following formula: xSb 2 O 3 -(95-x) PbO-5Me 2 O 3 (Me = Al, Ga, In). As shown in figure 12, density also decreases linearly with the substitution of PbO by Sb 2 O 3 .

Microhardness

Discussion

This study exemplifies the respective parts of the glass components. Antimony sesquioxide Sb 2 O 3 is clearly a vitrifying oxide. Even though the existence of pure vitreous Sb 2 O 3 has been for long a matter of discussion, the occurrence of numerous binary glasses with high Sb 2 O 3 concentration confirms its glass forming ability. It complies with Zachariasen's rules [7] insofar as it forms tetrahedra in the solid state that must be written SbO 3 (LP), where LP is a lone pair of s 2 electrons. In this way it offers similarities with phosphate groups PO 4 in which one oxygen is non-bridging. The vitreous network is constructed from these SbO 3 (LP) tetrahedra, with up to three non-bridging oxygens. Lead oxide PbO is known to have a dual role, depending on glass composition and also on the nature of the other vitrifyers. Solid state chemistry teaches that Pb 2+ commonly acts as a large cation that leads to compounds isostructural with those obtained with Sr 2+, Ba 2+ or Eu 2+ . Among other examples one may quote perovskite PbZrO 3 and fluorite-type PbF 2 . In these structures, coordination number is high: 12 in perovskite and 8 in fluorite [START_REF] Chiang | Physical Ceramics-Principles for Ceramic Science and Engineering[END_REF]. Based on this coordination and the corresponding size effect, divalent lead is considered as a modifier that fills empty space within the vitreous network. This behavior is common at low PbO content. However, at high concentrations, the structural behavior of Pb 2+ is different and compares to that of trivalent antimony. Its coordination number drops down while a lone pair of s 2 electrons completes a smaller coordination polyhedron. This polyhedron typically consists of 4 oxygens and one lone pair. This makes the basis of the PbO structure. Then this PbO 4 group can enter the vitreous network. Indeed there are numerous examples of glass compositions with a very large content of lead oxide. In the Sb 2 O 3 -PbO binary system, there are two different glass domains, the first one for large Sb 2 O 3 concentrations and the second one when lead is predominant. Not surprisingly characteristic temperatures reflect this situation (fig. 6): less stable glasses are found in the intermediate composition range.

The optical quality of the glass samples is very sensitive to the redox conditions. There are several redox couples that interact in glass melt: Sb 3+ /Sb 5+ , Pb 2+ /Pb 4+ , Pb 0 /Pb 2+ . Equilibrium constants vary with temperature. In order to limit the formation of reduced species, it is necessary to maintain an oxidizing atmosphere that can be provided by room atmosphere. However, in some cases, longer meltingand fining-time are required. When silica crucibles are used, this results in a larger contamination by silica. This drawback is avoided with other crucible materials such as alumina.

The influence of the third component on thermal properties is unexpected. In classical oxide glasses -silicates, phosphates, borates-glass transition temperature decreases as the ionic radius of the modifying element increases: Tg is smaller when sodium is replaced by potassium, rubidium or cesium. When network formers are incorporated, Tg variation is ruled by the relative bond strength of the new element. For example, the Si/Ge substitution leads to lower Tg's. In these glasses, the bond strength of aluminum vs. oxygen is much larger than that of gallium and indium. However, for the same nominal composition, larger Tg is observed with Ga 2 O 3 and In 2 O 3 (fig. 9). Then, it appears that bond strength is not the major parameter to account for this observed variation. Rather changes in glass structure should be considered. While these three cations -Al,Ga,Inare likely to enter the vitreous network, their coordination number is expected to be 6. As a result, the number of non-bridging oxygens is reduced, and consequently, glass transition temperature is lowered. This mechanism could also explain the increase of microhardness with the incorporation of the Me 2 O 3 oxides in the Sb 2 O 3 -PbO binary glass (fig. 13).

Conclusion

The incorporation of trivalent oxides in the Sb 2 O 3 -PbO binary glass has been studied. Characteristic temperatures, density and microhardness have been measured for binary glasses. The limits for glass formation have been defined by a set of systematic experiments in the Sb 2 O 3 -PbO-Me 2 O 3 (Me=Al,Ga,In) ternary systems. Up to 10 mol % Me 2 O 3 oxide could enter glass composition. Thermal measurements show that characteristic temperatures increase with oxide addition. In the mean time, glass is less prone to devitrification. Glass transition temperature increases according to the sequence: Tg(Al)< Tg(Ga)< Tg(In). Microhardness variation is similar, and maximum values are observed for indium containing glasses. It is suggested that this unexpected variation occurs from the decrease of the number of non-bridging oxygens. This decrease could result from the 6 coordination number of the trivalent cations that take place in the vitreous network.

  , a first result concerns Tg that increases with In 2 O 3 concentration. Another set of glass samples was prepared with the general formula: 50 Sb 2 O 3 -(50-x) PbO-x Me 2 O 3 (Me= Al, Ga, In). Glass transition Tg increases quasi-linearly with Me 2 O 3 addition. Once again, larger Tg values are observed for cations with larger ionic radius.

Figure 10

 10 Figure 10 reports the variation of the stability factor S as a function of Me 2 O 3 concentration for the 50 Sb 2 O 3 -(50-x) PbO-x Me 2 O 3 glasses.These curves suggest that Me 2 O 3 addition has a stabilizing effect by comparison to the binary glass. At first glance, glasses containing Ga content appear as the most stable ones. However additional experiments should be implemented to confirm these observations: first, no data were collected for Me 2 O 3 concentrations between 5 and 10 mol %; and second, no stable glass was obtained for In 2 O 3 concentrations larger than 5 mol % (see fig.3). Then one would expect the sharp decrease of the stability factor beyond some threshold value that could be found experimentally.

Figure 13

 13 Figure 13 shows the microhardness variation of the 50Sb 2 O 3 -(50-x) PbO-xMe 2 O 3 (Me = Al, Ga, In) glasses as a function of Me 2 O 3 concentration. Microhardness increases with Me 2 O 3

Figures

  Figures of the manuscript :

Table 1 M 2

 12

	O 3	Maximum reported glass forming range (mol % PbO)
	Al 2 O 3	50-95 mol% PbO
	Ga 2 O 3
	In 2 O 3

Table 2 :

 2 Glass formation has been investigated in variousSb 2 O 3 -PbO-Me 2 O 3 ternary systems, where Me = Al, In, Ga. Corresponding vitreous areas in nominal compositions are reported in fig3.

	Acronym	mol %		Nominal composition	Analyzed composition
		Sb 2 O 3 PbO	Sb	Pb	O	Sb	Pb	O	Si
	50Sb 2 O 3 50PbO	50	50	28.57	17.14	57.14	29.5	13.4	55.8	1.3
	40Sb 2 O 3 60PbO	40	60	25	18.75	56.25	21.7	17.2	54.8	0.9
	30Sb 2 O 3 70PbO	30	70	20.68	24.13	48.27	21.5	22.3	52.6	3.5
	20Sb 2 O 3 80PbO	20	80	15.38	30.76	53.84	15.6	27.1	54.0	3.3
	10Sb 2 O 3 90PbO	10	90	8.69	39.13	52.17	8.2	31.8	55.6	4.4

Acknowledgment

The authors would like to thanks PNE (Algerian program) for the financial support.


F Ortega

J Fettar

S Chatterjee

V Bansal

J.-M Tonnerre

E Mossang

L Ortega

F Fettar

J Chatterjee

S Auffret

I.-L Prejbeanu

B Dieny

Depth-resolved magnetization profile of MgO/CoFeB/W perpendicular half magnetic tunnel junctions

Depth-resolved magnetization profile of MgO/CoFeB/W perpendicular half magnetic tunnel junctions

I. INTRODUCTION

Magnetic tunnel junctions have been a keen focus in spintronics known for their usage in nonvolatile spin-transfer-torque (STT) magnetic random access memories [START_REF] Ohno | Magnetic tunnel junction for nonvolatile CMOS logic[END_REF], for which MgO/CoFeB/Ta trilayer systems with perpendicular magnetic anisotropy (PMA) have been proven to be a prime candidate [START_REF] Ikeda | [END_REF]. W insertion between CoFeB and Ta has been shown to improve PMA, thermal robustness and tunnel magnetoresistance (TMR) ratios [3,4]. Moreover, TMR and STT excitation processes are interfacesensitive quantities [5]. Probing magnetic buried interfaces in a layer exhibiting perpendicular magnetization, which limits the thicknesses to the nanometer, is a major challenge. In particular, better understanding of the magnetic interfaces in basic stacks for functional devices can help move beyond the ideal layers and perfect interfaces used for theoretical transport models and understand the difference between predictions and experiments. In this paper, the magnetization profile of a MgO/CoFeB/W free layer exhibiting PMA after annealing at 400°C is studied by X-ray resonant magnetic reflectivity (XRMR) in the soft X-ray range. Analysis of the angular dependence of the magnetic asymmetry reveals a perpendicular magnetization gradient with a significant increase in Fe and Co magnetization at the MgO interface and a 50% reduction at the W interface. In addition, the evolution of the magnetization profile, exposed to an external inplane (ip) magnetic field, is discussed. This is the author's peer reviewed, accepted manuscript. However, the online version of record will be different from this version once it has been copyedited and typeset.

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II. X-RAY RESONANT MAGNETIC REFLECTIVITY

XRMR is a powerful tool to probe the depth magnetization profile of a specific magnetic element in thin films, with sub-nanometer spatial resolution [6,7]. It exploits the dependence of the atomic scattering factor (ASF) with respect to the scattering vector q, incident photon energy, the polarization state of the incident and reflected beam as well as amplitude and orientation of the magnetization [8]. When the X-ray photon energy is tuned to an absorption edge of a magnetic element, the energydependent correction terms f'(E)+if''(E) to the Thomson scattering f 0 (q) and energy-dependent magnetism-sensitive scattering term f m (E) = m'(E)+im''(E) can be enhanced. The effect is particularly strong at the L 23 edges of 3d element in the soft X-ray range, where it is possible to use Xray circular magnetic dichroism (XMCD) effects [9]. In addition, thanks to the longer wavelength, angle-dependent specular reflectivity can be measured up to large scattering angles, providing sensitivity to oop magnetization [10,11]. The XRMR measurements are performed at room temperature using the RESOXS end station [12] at SEXTANTS beamline [13] at Synchrotron SOLEIL (France). The current magnetization device allows the application of a magnetic field in all directions of the sample plane but not perpendicular to the sample. Therefore, only thin films exhibiting 100% oop spontaneous magnetization can be investigated. The magnetic contrast is derived from the measurement of two intensities I p and I m , obtained either by using right and left circularly polarized incident light, keeping the direction of the applied magnetic field unchanged or measuring at remanence, or by reversing the direction of an applied magnetic field while keeping the X-ray helicity unchanged. The magnetic asymmetry is defined as

R = [(I p -I m )/(I p +I m )].
The sample is first magnetized by a permanent perpendicular magnet of µ o H = 0.4 T brought close to the sample surface followed by measurements at remanence [11] and, second, is subjected to an ip magnetic field of 0.17 T (maximum ip field currently available) at the intersection of the diffraction plane and the sample plane.

III. RESULTS AND DISCUSSION

We study here a half magnetic tunnel junction [MgO(~1)/Co 8 Fe 72 B 20 (1.6)/W (2)/Ta (1)/Pt (3)], part of a series of samples prepared with Ta and W/Ta cap layers of different thicknesses (nominal thickness in nm). They are deposited on [Si/SiOx//Ta (3)/CoFeB (0.4)]. Details on the preparation of the stack are given in [3]. Figure 1 shows the hysteresis loops obtained by Extraordinary Hall Effects for the perpendicular magnetic field, where the signal is proportional to the perpendicular component of the magnetization, and by SQUID for the parallel magnetic field. All measurements were performed at 300K and 1T. The M(H perp ) curve shows that 100% of the saturation magnetization (M s ) can be maintained at the zero field. The shape of the M(H par ) curve indicates that before saturation, a canted magnetic configuration could appear [14]. M s = 1500125 emu/cm 3 is properly determined after a calibration using a Fe sample of similar size and the same parameters for the magnetic measurements and by considering the actual rather than nominal thickness of CoFeB. This is the author's peer reviewed, accepted manuscript. However, the online version of record will be different from this version once it has been copyedited and typeset. The study of the magnetization profile relies on the refinement of R, derived from angle-dependent reflectivity, which requires the determination of the structural parameters of the heterostructure. In order to allow a more unequivocal analysis of the X-ray reflectivity [15], the data are collected at various incident photon energies 670 eV and 760 eV, far from both the Fe L 3 (706.8 eV) and Co L 3 (778.1 eV) edges and at a few energies close to the edges. Figure 2 shows the non-magnetic intensities I ave = (I p +I m )/2 and the fitting curves as a function of the scattering vector q for energies close to the Fe L 3 edge. The variations in the shape of the intensities are linked to the energy dependence of f 1 = Z+f' (f 0  Z in the soft X-ray range) and f 2 = f" for Fe in Co 8 Fe 72 B 20 (shown in the inset). They are well reproduced (also at the edge Co), leading to consistent parameters. The density (ρ in mol/cm 3 ), thickness (t in nm) and the interfacial roughness (σ in nm) for the main layers are listed here: Pt (0.0910.003 / 2.920.24 / 0.240.01), Ta (0.0830.004 / 1.090.18 / 0.560.2), W (0.0890.005 / 1.930.24 / 0.290.19), CoFeB (0.1290.015 / 1.360.10 / 0.50.04) and MgO (0.1590.017 / 1.360.08 / 0.310.02). This is the author's peer reviewed, accepted manuscript. However, the online version of record will be different from this version once it has been copyedited and typeset. We now turn to the analysis of the R measured in the remnant condition. The reflectivity curves are collected up to an incident angle of 80° corresponding to q max 7 nm -1 at the Fe L 3 edge. The dichroic effect in I p and I m is illustrated in Figure 3a for 706.8 eV. Together with the absence of splitting at low angles, a clear splitting at larger angles is a strong indication of a net perpendicular magnetization, as expected. Such features are present in the data for all the energies at both edges. Figure 3(b,c,d) displays the q-dependence of R derived at 706.8 eV, 709 eV and at 778.1 eV. They exhibit changes related to m'(E) and m"(E) (not shown) that are derived from reference samples [16]. The refinement is performed by considering an increasing number of sub-layers in CoFeB and by adjusting for each one a weighting factor , modifying the f m (E) amplitude and the thickness, while the other structural parameters are kept fixed. This corresponds to modifying the magnetization, averaged laterally over a certain thickness across the layer. Its amplitude is linked to the XMCD of the reference samples. A value  = 1 implies an amplitude of the 3d magnetic moment of 2.1 µ B for Fe and 1.7 µ B for Co, which results from the application of the sum rules [16]. A homogeneous oop magnetization profile across the CoFeB layer did not permit to fit the data and a minimum of four sub-layers was required to get the best fits at all energies, as illustrated in Fig. 3. The adjustment is constrained by simultaneously refining R and I ave . The CoFeB layer exhibits an oop magnetization gradient as depicted by the  changes from 1.16  0.13 at the MgO/CoFeB interface to 0.52  0.17 at the CoFeB/W one, in the inset of Fig. 3a for Fe and similar for Co (Fig. 3d). This corresponds to a gradual decrease of 2.40.3 µ B to 1.10.4 µ B . Another depth profile model, based on a dead layer 0.6 nm thick [3,4] at the upper interface and homogeneous magnetization over the rest of the layer, was tested and cannot reproduce the data over the entire angular range (Fig. 3b). This is the author's peer reviewed, accepted manuscript. However, the online version of record will be different from this version once it has been copyedited and typeset. We finally present the study of the magnetization profile under a 0.17 T ip magnetic field applied parallel to the diffraction plane. Referring to the parallel hysteresis loop (Fig. 1), this value is insufficient to saturate the magnetization but appears sufficient to achieve a single domain configuration. Figure 4a shows I p and I m reflectivity curves at 706.8 eV that now exhibit a splitting at low angles indicating the presence of a net ip magnetization component. At the same time, a net oop contribution persists as can be seen at larger angles in the 706.8 eV and 709 eV magnetic asymmetries (Figures 4b and4c). This behavior observed for the five Fe resonant energies (no data at the Co L 3 edge). This is in agreement with a canted configuration [17]. The fits of R lead to a similar magnetization gradient (inset) with a tilt angle of  = 343 degrees for the four sub-layers. We did not try to probe the distribution of tilt angles across the layer [18] since a constant value for a 1.36 nm thick layer provides a very good fit for the data. This is the author's peer reviewed, accepted manuscript. However, the online version of record will be different from this version once it has been copyedited and typeset.

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IV. CONCLUSION

A first notable characteristic of the Fe magnetization profile within the annealed CoFeB layer is a high 3d magnetization in the first 0.3  0.09 nm thick sublayers (about 2 ML) at the MgO interface. Such an enhancement of the magnetization has also been observed for Fe on MgO studied by XMCD [19] and for Fe in the 2-3 ML near both MgO interfaces in a MgO/Fe (12 nm)/MgO (001) trilayer by using SXRMR [20] and theoretically discussed [21]. We found a similar magnetization increase for Co. A second feature is the lowering in magnetization towards the W interface, which is more pronounced after 0.7 nm from the MgO interface (thickness of the two sublayers) . This decrease is ascribed to a localized reduction of M S and T C due to alloying of CoFeB with W , as proposed earlier [22]. This is supported by a 0.5 nm roughness at the CoFeB/W interface. The averaged  value from the Fe magnetization profile amounts to 0.910.13, corresponding to 1.90.3 µ B , which leads to 1480234 emu/cm 3 in reasonably good agreement with the value derived from SQUID data. We also characterize the magnetization profile for various values of the applied magnetic field from 0.17 T to 0.02 T. The modification concerns mainly the tilt angle that varies linearly from 34 to 4 degrees, indicating the magnetization gradient is a robust property of the CoFeB layer in the investigated sample. This is the author's peer reviewed, accepted manuscript. However, the online version of record will be different from this version once it has been copyedited and typeset.

PLEASE CITE THIS ARTICLE AS DOI:10.1063/9.0000343

This investigation is part of a larger investigation dedicated to the comparison of the magnetization profile between as-deposited samples and annealed samples with a Ta (1 nm) cap layer and a W (2 nm)/Ta (1 nm) cap layer. The full results will be published elsewhere.

  DOI:10.1063/9.0000343

FIG. 1 .

 1 FIG. 1.Normalized magnetic loops at room temperature for perpendicular and planar magnetic fields. Inset shows a close-up of the small magnetic field region.

FIG. 2 .

 2 FIG. 2. X-ray reflectivity data (open circles) for energies near the Fe L3 edge (inset) and at 670 eV with limited resonant effects. The curves are shifted vertically for the sake of clarity and the best fits are plotted in continuous lines.

FIG. 3 .

 3 FIG. 3. Reflectivity curves for circular right and left polarized incident beam(a). Magnetic asymmetries for three resonant energies and their best fits (b, c, d). For 706.8 eV, a model with a dead layer is shown in a dashed line for comparison.

FIG. 4 .

 4 FIG. 4.Reflectivity for circular right and left polarized incident beam with H par = 0.17 T(a). Magnetic asymmetries for two resonant energies and their best fits (b, c).

V. ACKNOWLEDGMENTS

Financial support from the Agence Nationale de la Recherche, France, under Grant Agreement No. ANR-16-CE09-0009 (WAVENEXT) is acknowledged.


Raj Kumar Ramamoorthy

Sébastien Teychené

Isaac Rodriguez-Ruiz

Jean-Philippe Torré
email: jean-philippe.torre@ensiacet.fr

Insights on the formation and dissociation mechanisms of cyclopentane hydrate obtained by using calorimetry and optical microscopy

Keywords: Gas hydrates, calorimetry, optical microscopy, in-situ crystallization, cyclopentane

Identifying the characteristic properties of gas hydrates including thermo-physical properties are essential for determining their formation condition and in turn, controlling their yield/growth-rate for/in various applications/environments. Gathering these properties, during in-situ crystallization of gas hydrates, are rather challenging due to the harsh conditions they form and limited by the lack of co-existence of multiple characterization techniques. Owing to these limitations, crystallization process of some SII type gas hydrates and their characteristic properties are ill-defined. Therefore to unveil the complete picture of crystallization process of SII type gas hydrates, in this study, in-situ formation of cyclopentane hydrate was monitored using a customized DSC sample holder with a transparent window on top. To this end, in-situ crystallization of cyclopentane hydrate was simultaneously followed by coupling optical microscopy along with scanning calorimetry. Further, we show that the cyclopentane hydrate formed by consuming a part of fresh water appearing after the fusion of ice instead of at a -20 °C.

Introduction:

Understanding the crystallization process of gas hydrates is essential for a wide range of areas, including identifying their inhibiting conditions to avoid blockage of gas/oil pipelines [START_REF] Hammerschmidt | Formation of gas hydrates in natural gas transmission lines[END_REF][START_REF] Heng-Joo | Hydrate formation and inhibition in gas condensate and hydrocarbon liquid systems[END_REF], and controlling their yield for secondary refrigeration [START_REF] Skovborg | A mass transport limited model for the growth of methane and ethane gas hydrates[END_REF]Galfré et al., 2014;[START_REF] Khurana | Fundamentals and applications of gas hydrates[END_REF], and gas storage applications [START_REF] Khurana | Fundamentals and applications of gas hydrates[END_REF][START_REF] Karanjkar | Calorimetric investigation of cyclopentane hydrate formation in an emulsion[END_REF]. Generally, hydrates are ice-like crystals formed at low temperatures and high pressures [START_REF] Khurana | Fundamentals and applications of gas hydrates[END_REF][START_REF] Khurana | Fundamentals and applications of gas hydrates[END_REF] by combining water (host) and gas (guest) molecules with an appropriate stoichiometric ratio. Owing to these harsh conditions (Karanjkar et al., 2012a) and complex phenomena at water/gas interface [START_REF] Warrier | Overview: Nucleation of clathrate hydrates[END_REF], capturing such crystallization process, using common lab facilities, is quite challenging. Nevertheless, to unveil such a complex process, a very few in-situ studies of hydrate formation in suitable sample environments have been performed using X-ray diffraction (XRD) [START_REF] Fouconnier | Study of CCl3F hydrate formation and dissociation in W/O emulsion by differential scanning calorimetry and X-ray diffraction[END_REF], optical microscopy [START_REF] Sun | Morphology investigation on cyclopentane hydrate formation/dissociation in a sub-millimeter-sized capillary[END_REF], Raman spectroscopy [START_REF] Stern | Opticalcell evidence for superheated ice under gas-hydrate-forming conditions[END_REF] and differential scanning calorimetry (DSC) (Davies et al., 2010a;Davies et al., 2010b).

In this context, trichlorofluoromethane (CCl3F) and cyclopentane (CP) hydrates were found to appear after the beginning of ice melting using XRD [START_REF] Fouconnier | Study of CCl3F hydrate formation and dissociation in W/O emulsion by differential scanning calorimetry and X-ray diffraction[END_REF] and optical microscopy [START_REF] Sun | Morphology investigation on cyclopentane hydrate formation/dissociation in a sub-millimeter-sized capillary[END_REF][START_REF] Stern | Opticalcell evidence for superheated ice under gas-hydrate-forming conditions[END_REF], respectively. In another study, using optical microscopy [START_REF] Delroisse | Determination of thermophysical properties of cyclopentane hydrate using a stirred calorimetric cell[END_REF], CP hydrate was found to emerge above the equilibrium temperature of ice, and their growth rates were found to be in relation with the super-cooling temperature. Taking advantage of these findings, both Raman spectroscopy (Davies et al., 2010a) and DSC studies (Davies et al., 2010b) were performed to determine the mass transfer of guest (methane) and water molecules during the crystallization of methane hydrate. Subsequently, water molecules were found to be highly mobile in the hydrate layer and its growth rate was found to be controlled by the experimental sub-cooling temperatures. In continuation, thermophysical properties such as equilibrium temperature and enthalpy have been successfully estimated, for various gas hydrates, using DSC (Davies et al., 2010b;[START_REF] Delroisse | Determination of thermophysical properties of cyclopentane hydrate using a stirred calorimetric cell[END_REF][START_REF] Torré | A novel stirred micro calorimetric cell for DSC measurements applied to the study of ice slurries and clathrate hydrates[END_REF][START_REF] Martinez | Thermodynamic Properties of THF + CO2 hydrates in relation with refrigeration applications[END_REF].

Seemingly, most of the thermo-physical properties of gas hydrates were derived by considering that hydrates are forming during or after the beginning of ice melting.

However, this assumption is not precise and may not be sufficient to describe the complete picture of crystallization process, involving different guest molecules, for the following reasons: For the CP/water system, CP hydrate forms immediately after the emergence of water from ice melt whereas CCl3F hydrate forms at the far end of ice melting. Therefore, overlooking such pivotal information would result in an ill-defined crystallization process and in turn, in a lack of information on their inhibiting conditions and thus a poor control on the final yield of gas hydrates.

Therefore, to unveil the formation mechanism of gas hydrates, in particular, naturally forming ones, identifying a model system is essential. Among many gas hydrates, CP hydrate, formed by the stoichiometric ratio of 17 H2O molecules and one CP molecule, is a suitable candidate for understanding the crystallization processes occurring in petroleum pipelines and in sea beds, due to its mild formation condition (below 7 °C and ambient pressure: Nakajima et al., 2008) and a SII crystal structure (each unit cell is composed of 16 small cavities forming pentagonal dodecahedrons and 8 large cavities forming hexadecahedrons: [START_REF] Sloan | Clathrate hydrates of natural gases. 3rd edition[END_REF]. Thus, monitoring in-situ crystallization of using optical microscopy in coupling with either structural or thermal experiments would be an ideal combination to provide new insights on the crystallization process of SII type gas hydrates.

To the best of our knowledge, for the first time, in this study, in-situ crystallization of CP hydrate was simultaneously captured by using optical microscopy in coupling with single-cell scanning calorimetry (SCSC). These complementary and simultaneous observations show that the CP hydrate forms immediately after the beginning of ice melting and its yield is found to be controlled by the heating rate (around ice melting)

and the number of experimental cycles.

Materials and Methods:

, purchased from Sigma-Aldrich, and deionized water were used for all experiments. CP was used without further purification. All experiments were performed in quiescent conditions. In this study, SCSC (DSC Linkam 600: attached with a cryostat device) along with optical microscopy were used to follow the CP hydrate crystallization (please see supplementary information (Fig. S1) for the coupled configuration). To couple both SCSC and optical microscopy, first, standard DSC lids were pierced and sealed by transparent polymer: NOA 88. Second, for each experiment, ca. 25-28 µl of water and 15 µl of CP were loaded into a DSC pan and hermetically sealed by a customized DSC lid with the transparent NOA 88 window.

Either two or three temperature cycles were programmed for each experiment. Typically, (i) the cooling cycle was performed at the rate of 5 °C/min from room temperature to ~ -22 °C and (ii) the heating cycle was performed by using three different heating rates:

from ~ -22 °C to -2 °C at the rate of 5 °C/min, from -2 °C to 3 °C at the rate between 0.06 and 0.18 °C/min (at the step of 0.06 °C/min) and from 3 °C to the dissociation of CP hydrate at the rate of 1 °C/min. Henceforth, experiments will be named by the heating rate set during the ice melting: CPGH1 (0.06 °C/min), CPGH2 (0.12 °C/min), and CPGH3 (0.18 °C/min). Simultaneously, an optical image of water/CP system was recorded every 0.552 s (exposure time: 0.3 s + time interval: 0.252 s) using a highspeed Imager sCMOS from LAVISION. The enthalpy of fusion of ice and dissociation of hydrate, in standard units (J/mol), were derived with the help of a calibration constant, which was estimated using standard reference samples (alkanes), as described in the reference [START_REF] Skovborg | A mass transport limited model for the growth of methane and ethane gas hydrates[END_REF].

Results and Discussion:

3.1. In-situ physical changes of CP/water system using optical microscopy in coupling with SCSC:

To unveil the crystallization process of CP hydrate, physical changes of the CP/water system were followed as a function of experimental (cooling and heating) cycles. For simplicity, selected optical images along with the evolution of calorimetric signal of systems (CPGH1 and CPGH2), during their second experimental cycle, are displayed in Fig. 1 and2, respectively. Please refer to high-speed videos (supplementary information: Video S1 and Video S2), for the continuous evolution of systems: CPGH1 and CPGH2, respectively. Irrespective of the experimental conditions, ice formation and melting, and hydrate formation and dissociation are directly visualized and traced using optical microscopy and SCSC, respectively. During the cooling cycle, even after crossing the equilibrium temperature of ice i.e. 0 °C, the CP/water system remains unperturbed.

Once the system reaches ca. -18 °C, the beginning of crystallization is directly visualized by the optical recording, and the same is confirmed by the sudden increase of calorimetric signal (the measured exotherm is proportional to the heat released by the crystallization process) using SCSC. This transition completes around -20 °C, which is evidenced by the single exothermic peak. The homogenous contrast of solid phase, along with a few droplets (for CPGH1), confirms that the solid phase is made up of H2O, as observed by [START_REF] Sun | Morphology investigation on cyclopentane hydrate formation/dissociation in a sub-millimeter-sized capillary[END_REF], instead of both H2O and H2O/CP. During the heating cycle, after holding at -22 °C, the physical appearance of solid-phase remains unmodified. Once the system approaches 0 °C, ice melts into liquid water.

With further increase in temperature and time, a new solid-phase emerges all over the recording region by consuming a part of the fresh water coming from ice melt. This process cannot be captured by SCSC due to the dominant endothermic peak of ice fusion. Within the focal plane, we observe that this phase appears close to the equilibrium temperature of ice i.e. very close to the beginning of ice melting but not until the completion of ice melting: between 2 and 3 °C. This observation is consistent with a previous study of CP hydrate [START_REF] Sun | Morphology investigation on cyclopentane hydrate formation/dissociation in a sub-millimeter-sized capillary[END_REF] but in contradiction to the formation of CCl3F [START_REF] Fouconnier | Study of CCl3F hydrate formation and dissociation in W/O emulsion by differential scanning calorimetry and X-ray diffraction[END_REF] and THF [START_REF] Muro | Formation and growth of tetrahydrofuran hydrate at the ice/hexane interface[END_REF] hydrates, wherein hydrates form at the far end of ice melting and below the equilibrium temperature of ice, respectively. With the further temperature increase, the endothermic peak of ice to water transition completes along with the appearance of a new endothermic peak welling optical observation during the second endothermic peak displays the disappearance of the newly formed phase, along with the appearance of a large number of tiny bubbles. The onset of this dissociation begins ca. 7.0 ± 1.0 °C, which confirms that the newly formed phase during the ice melting is CP hydrate, as previously reported [START_REF] Fouconnier | Study of CCl3F hydrate formation and dissociation in W/O emulsion by differential scanning calorimetry and X-ray diffraction[END_REF][START_REF] Sun | Morphology investigation on cyclopentane hydrate formation/dissociation in a sub-millimeter-sized capillary[END_REF].

Enthalpy of CP hydrate dissociation and water to CP hydrate conversion rate SCSC:

After determining the relevance of experimental conditions on hydrate formation, a series of experiments were performed to quantify the hydrate formation in relation with the heating rate (close to the temperature of ice fusion), and the number of thermal cycles, using SCSC. The resulted calorimetric and temperature records are plotted as a function of time in Fig. 3. The amount of water involved in the hydrate formation was estimated after subtracting the weight of water involved in the first endothermic peak from its initial weight. Then, the water to hydrate conversion rate, as a function of heating rate and experimental cycles, was estimated by following the procedure described by N. Gaikwad and coworkers [START_REF] Gaikwad | Clathrate hydrate equilibrium data for the gas mixture of carbon dioxide and nitrogen in the presence of an emulsion of cyclopentane in water[END_REF]. Their values, as a function of heating rate and experimental cycles, are summarized in the Table . 1.

Figure 3 .

 3 Figure 3. Calorimetric and temperature traces in relation with cooling and heating cycles: (a) CPGH1, (b) CPGH2, and (c) CPGH3.

Table 1 :

 1 Total weight of CP hydrate and water to hydrate conversion after the completion of each experimental cycle

		Weight of CP hydrate: CPGH1 (mg)	Water to hydrate conversion %	Weight of CP hydrate: CPGH2 (mg)	Water to hydrate conversion %	Weight of CP hydrate: CPGH3 (mg)	Water to hydrate conversion %
	First Cycle	0.56	1.65	0.32	0.94	0.02	0.05
	Second Cycle	0.58	1.71	0.46	1.36	0.09	0.29
	Third Cycle			0.63	1.86	0.11	0.36

Acknowledgments: This work was supported by the French Agence Nationale de la Recherche (ANR) [grant number ANR-18-CE05-0028], and undertaken in the frame of the US Partnership for International Research and Education program (National Science Foundation Award Number 1743794) and of the French R -2026 Hydrates de

Further, the authors would like to acknowledge all the people involved in the ANR MUSCOFI project for their valuable scientific input and discussions.

Funding sources :

This work was funded by the French Agence Nationale de la Recherche (ANR) [grant number ANR-18-CE05-0028].

Consequently, we found that the yield of hydrate formation increases and decreases with the number of experimental cycles and increasing heating rate, respectively, as observed for CP hydrates in a water-in-oil emulsion (Karanjkar et al., 2012b). This confirms that even with several thermal cycles hydrate formation performed in quiescent condition resulting in a static interface between water and CP is not favorable to achieve high water to hydrate conversion rates. Further, irrespective of experimental conditions, the onset, and enthalpy of hydrate dissociation are found to be ca. 7.0 ± 1.0 °C and 115,600 ± 4,500 J/mol, respectively, which is in agreement with the result of [START_REF] Delroisse | Determination of thermophysical properties of cyclopentane hydrate using a stirred calorimetric cell[END_REF]but 3% off from Nakajima et al., 2008;[START_REF] Torré | A novel stirred micro calorimetric cell for DSC measurements applied to the study of ice slurries and clathrate hydrates[END_REF] Therefore, in the future, the crystallization process of CP hydrate in presence of various additives such as surfactants i.e. sodium dodecyl sulphate would be followed using optical microscopy in coupling with scanning calorimetry. Such investigations would be very useful to deeply understand how such additives modify the crystal morphology, the formation kinetics, and the water to hydrate conversion.

Conclusions:

In-situ formation and dissociation of CP hydrate, as a function of heating rates (close to the equilibrium temperature of ice) and experimental cycles, were followed using optical microscopy in coupling with SCSC. Using these in-situ studies, we demonstrate that CP hydrate forms close to the equilibrium temperature of ice, especially, after the appearance of fresh water from ice fusion instead of either at the far end of ice melting, as observed for CCl3F or below the equilibrium temperature of ice, as observed for THF hydrate. Further, CP hydrate formation was found to be enhanced with a slow heating rate (close to the equilibrium temperature of ice) and the number of experimental cycles.

Nevertheless, even after the third experimental cycle, the inadequate quiescent condition leads to an extremely low conversion rate of CP hydrate from water. We believe that the presented experimental approach would pave a way for more hybrid experiments helping to clarify and describe the crystallization process of hydrates and in turn, provide a lever to control the final yield of gas hydrates for various applications including gas storage and transportation.

Appendix A:

Supplementary data associated with this article can be found in the Supporting Information.

Declaration of Competing Interest:

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.