F Cipriani

F Leblanc

J M Illiano

J J Berthelier

F Müller

Physcial Characterization of Molybdenum microtips Field Emitter arrays

Keywords:

We used a dedicated hemispherical energy analyzer to measure energetic and angular distributions of electrons emitted from Molybdenum microtips integrated in a 1 cm 2 Field Emitter Array designed by the CEA/LETI laboratory. Such cathodes typically deliver about 25mA at an extraction voltage of 100V, and are studied in order to replace heated wires as electron sources for space applications. We find that the energy distribution of the beam strongly depends on the extraction voltage, and is therefore expected to vary across the emission life time of the device, at a rate depending both on the alteration of the resistive structure with time, and on the fate of adsorbed contaminants at the tip surface. A semi-empirical model of the emitters is proposed and used to determine parameters of energetic and angular distributions. The energy dispersion of the beam is found to increase from 2eV±20% up to 9eV±20% eV, for extraction voltages varying from 40 to 100 V. The mean angular dispersion of the beam is found to be 42 • ±20% when a null electric field is set at the grid extraction surface.

Introduction

Most of mass spectrometers on space missions have used electron impact ion sources based on thermoionic electron guns. Thermoionic sources provide electron current up to typically 500 μA, requiring several watts of power to heat the emitting filament up to typically 2000K. High power consumption, filament burning, and outgassing of heated materials in the source are the main shortcomings of the hot filament technique. Cold electron emitters such as Field Emitter Arrays (FEA) originally developed for flat screen applications, offer in that sense interesting perspectives for the next generation of space instruments.

In the present study we test Molybdenum microtips field emitters. Such cathodes have been extensively studied in the context of vacuum microelectronics applications such as flat panel display, microwave power tubes, or pressure gauges [START_REF] Temple | Recent progress in field emitter array developpment for high performance applications[END_REF], and their performance continuously improved over the past decades [START_REF] Schlesser | Energy distribution of field emitted electrons from diamond coated molybdenum tips[END_REF]. In particular several papers describing their current-voltage characteristics, emission area, emission stability as a function of exposition to different gases and pressures, can be referred to [START_REF] Spindt | Physical properties of thin-film field emission cathodes with molybdenum cones[END_REF][START_REF] Brenac | Caractéristiques énergétiques de cathodes à micropointes à émission de champ[END_REF][START_REF] Constancias | Emission d'Electrons à partir de Micropointes pour Ecrans Plats : Simulation, Caractérisation et Confrontations Experimentales[END_REF].

Critical reliability issues have been pointed out (see for instance [START_REF] Aplin | Reliability tests of gated silicon field emitters for use in space[END_REF] for silicon field emitters) originating from various phenomena as degradation due to ion bombardment, arcing, or thermal breakdown due to electromigration. The present study was conducted with 1 cm 2 Mo tips cathodes containing 14400 tips.mm -2 and delivering a 25mA nominal current. The samples were fabricated at CEA-LETI in Grenoble, France, using lithographic technique to obtain Spindt type tip emitters [START_REF] Constancias | Emission d'Electrons à partir de Micropointes pour Ecrans Plats : Simulation, Caractérisation et Confrontations Experimentales[END_REF].

Hot filaments have the advantage of emitting electrons with very low energies that can be accelerated and focused with a simple electrostatic optics. This is not the case of FEA emitters due to the field effect which causes electron extraction. In order to design ion sources it is thus necessary to have a detailed knowledge of the properties of the electron beam such as its energy or angular distribution. The knowledge of such parameters also allows to define optimized ion sources geometries and optics potentials in order to precisely control the electrons energy and spatial distribution of the electronic charge in the source. This is essentially done to obtain large ionization efficiencies, useful when probing low density media in space. For this purpose we set up an experiment based on an hemispheric energy analyzer, allowing a simultaneous measurement of the angular and energetic distributions of the emitted electrons [START_REF] Cipriani | La couronne neutre suprathermique de Mars : modélisation et développement d'une instrumentation pour sa mesure in situ[END_REF]. This experiment also allows to study the current stability as a function of the nature and pressure of the ambient atmosphere.

Previous field emission energy distribution measurements carried out by [START_REF] Brenac | Caractéristiques énergétiques de cathodes à micropointes à émission de champ[END_REF]; [START_REF] Purcell | Nanoprotrusion model for field emission from integrated microtips[END_REF], have revealed a complex behaviour due to processes affecting the structure of the material making the microtips as well as peculiarities in the design of the component.

In the second and third sections of this paper we describe the emitter structure, as well as the experimental setup and principle of operation. In the fourth section we show energy and angular distributions measurements that are analyzed and modelled in the fifth section with the help of previous results from [START_REF] Baptist | Microtips and resistive sheet : A theoritical description of the emissive properties of the system[END_REF]; [START_REF] Purcell | Nanoprotrusion model for field emission from integrated microtips[END_REF]. We actually show that due to the large width of energy distributions of emitted electrons, such cathodes are probably not the best candidates for an efficient design of ion sources.

Description of the device

The component is a square matrix composed of 32 × 32 independent pixels (0.4 × 0.4μm 2 ) covering a total area of 11.2 × 11.2mm 2 . The extraction grids and tips (Figure 1) are polarized through two groups of Niobium tracks at right angle from each other. To ease the practical handling of the FEA, components have been glued into a Pin Grid Array (PGA) support using a UHV qualified cement. The Niobium tracks are bonded to the leads of the PGA. Molybdenum tips have an average height of 1μm and an average half cone angle of 15 • [START_REF] Constancias | Emission d'Electrons à partir de Micropointes pour Ecrans Plats : Simulation, Caractérisation et Confrontations Experimentales[END_REF].

The diameter of extraction grids apertures is 1.4μm (cf. Figure 1). Extraction grids and cathodic contacts are isolated by a 1μm silica layer, as shown on Figure 3. A 800nm silicon layer, referred to as the resistive sheet, is included in the structure between tips and cathodic contacts, in order to obtain good emission homogeneity, and to prevent destructive arcing between the tips and the extraction grids [START_REF] Constancias | Emission d'Electrons à partir de Micropointes pour Ecrans Plats : Simulation, Caractérisation et Confrontations Experimentales[END_REF].

The electrical equivalent structure is shown on Figure 3. The equivalent resistance for the silicon layer between one tip and the cathodic contact through which it is biased has a value of 300 MΩ (Muller, personal communication). This value has been shown to be current dependant, as a consequence of the variation of the equilibrium point of the voltages distribution across the tips (see [START_REF] Baptist | Microtips and resistive sheet : A theoritical description of the emissive properties of the system[END_REF] for more details), and also because of its semi-conductive nature.

However its variations actually remain within 5% of the mean value, in the 0-100V Vgc range.

The complete electrical equivalent model shown on Figure 3 can be somehow simplified when considered in continuous voltage operation.

Overview of the Experimental Setup

The experimental setup is installed in a dedicated vacuum chamber with a base vacuum level of 10 -8 mbar, achieved by turbomolecular pumping. Continuous measurements of the residual gas were carried out using a quadrupole mass spectrometer. Cathode baking was not performed before operation as we reproduced operations onboard a satellite where the large power necessary for heating would not be available. A satisfactory vacuum level of 2.3 × 10 -8 torr was obtained after two weeks of pumping. The vacuum composition is basically a mixture of the residual nitrogen used as the venting gas (41%), and residual air. The potential interference of those gases on our measurements is very weak at this pressure level. We noted that degassing of H 2 from the vacuum chamber walls remained fairly weak compared with the nitrogen level (about 25% of the nitrogen peak).

As shown on Figure 4 the main components of the experimental setup are : (1) a sample holder attached to a XY linear travel, (2) the Field Emitter Array, and (3) the hemispheric energy analyzer. Electrical connections were realized with Kapton insulated cables or with bare rods connected to feedthroughs on the chamber flanges. The PGA support is mounted on an alumina support followed by bare rods. The whole assembly can be moved in front of the analyzer entry plane along two X,Y perpendicular directions with a positioning resolution of 5μm.

Measurements of electrons characteristics were preformed with the retarding potential hemispheric analyzer, discussed in details in [START_REF] Cipriani | La couronne neutre suprathermique de Mars : modélisation et développement d'une instrumentation pour sa mesure in situ[END_REF]. In this setup (see Figure 4) the FEA is positioned in a plane parallel to the entrance plane of the analyzer, which has an entry hole 250μm in diameter. The distance between the cathode' surface and the entrance plane is 5mm. More than 99% of the emitted beam is collected on the entrance plane and the electrons passing through the entry hole enter the analyzer cavity within a 2 • aperture cone.

The inner surface of the analyzer entrance plate is insulating and supports 18 materialized equipotentials under the form of rings of width 1mm, polarized in such a way that the retarding field in the analyzer cavity is radial. Electrons entering the hemisphere are therefore decelerated by the radial retarding field along their trajectories. A potential energy diagram of the experiment is shown on Figure 2. The retarding field amplitude is set by the difference between the collector potential V coll, and the potential V d of the entrance plate of the analyzer.

Electrons are emitted from an energy level theoriticaly close to the Fermi level of the tips, which are biased to potential Vp, and are accelerated towards the extraction grids, which are biased to a more positive potential Vg. Therefore electrons energies reach a maximum value close to e(Vg-Vp)-Φ grid at the grids surface (where Φ grid is the work function of the grid). The collector is in stainless steel with a work function Φ coll of 4.85eV according to [START_REF] Benenson | Handbook of Physics[END_REF]. We estimate this value as representative of the actual work function of the collector within ± 0.5eV. The collector's potential has to be varied from a value V coll max such that

V coll max > e(V g-V p)+Φ coll , downto a value V coll min such that V coll min < e(V g-V p)+Φ coll .
The cut-off condition corresponds to equality. The collector's potential is varied in 100meV steps, resulting in a non zero electron current when the retarding potential V min is lower that the lowest energy of the electrons, and providing measurement of the energy distribution when varied from V min to V max (retarding potential preventing electrons to hit the collector). Moving the FEA in front of the entrance plane provides a at the same time measurement of the electron intensity as a function of the angle between the FEA tips axis and the central axis of the hemispherical analyzer.

Measurements

Measurement sequence

We randomly selected a set of 5 pixels (out of 1024) on which measurements of the energetic and angular dispersions of the beam have been carried out. The coordinates of such pixels on the 32 columns (c) by 32 lines (l) matrix are : (c9,l14), (c9,l15), (c9,l16), (c17,l15), and (c27,l8).

During our measurements cathodic contacts were biased at 0V, extraction grids were biased at a potential Vg varying between +20V and +100V. The entrance plate was biased to a potential Vd equal to Vg+20V, such that the electric field is directed from the entrance plate towards the FEA, that is opposite to the retarding field in the analyzer. This allows to reduce the potential bubble forming near the entry hole and disturbing electron trajectories. Such biasing also prevent most of the secondary electrons emitted by the entrance plate to be lost by back scattering, ensuring a minimum error on the emitted current measurement. The potential of the protecting electrode (surrounding the emitting surface, in red in Figure 4) was set near 0V to produce a slight beam focusing effect and avoid hazardous electronic trajectories outside the analyzer volume.

Current voltage characteristic and emission stability

In order to avoid destructive arcing, extraction grids potentials were operated with a duty cycle of 10%. That is, the extraction voltage Vgc is applied during 1ms, every 10ms. The FEA is operated in a residual vacuum (P 10 -8 mbar). After stablization of the Fowler-Nordheim characteristic, the emitted current density is 25mA.cm -2 when the extraction voltage is 100V.

Figure 5 shows the typical evolution of the emitted current over a period of 30h for a subset of 32 pixels forming one column of the emitter. The average current is 65.9μA with variations less than 2% during the measurement. A stabilized current-voltage characteristic of this column is plotted on Figure 6, the emitted current reaching 80μA at 100V. The inset of Figure 6 shows a Fowler-Nordheim fit to the measured current. As can be observed, a slight deviation from the Fowler-Nordheim fit occurs at high voltage. Such deviation is typical for this type of emitter and results from the presence of a resistive sheet between cathodic contacts and tips. At high current values flowing through the resisitve layer, the resulting voltage drop at the tips lowers the emitted current compared with the Fowler-Nordheim current for the same voltage without resistive sheet. As will be shown in section 5, such resistive layer has direct consequences on the width of the energy distribution of emitted electrons. As will be shown in detail in section 5 our observations are consistent with the Total Energy Distribution measurements by [START_REF] Purcell | Nanoprotrusion model for field emission from integrated microtips[END_REF] :

Results of the energy analysis

-a shift is observed between the effective collection threshold voltage and the threshold expected if emitted electrons had energies close to eVgc. This shift increases with Vgc.

-the energy spread of the distribution increases with Vgc, reaching roughly 10eV at 100V in our case.

A comparison of collected currents for the set of 5 pixels tested when biased at 60V is shown in Figure 8. Though the amplitude of the total collected current varies by a factor of 7, the variability over pixels of the mean values of the energy distributions deduced from the curves is not larger than 10% of the average mean energy value. Such variations are representative of variations observed at any extraction voltage between 40V to 100V.

The factor of 7 in the current level at the collector reflects a variability of the total emitted current between pixels. Such variability can be related to a number of factors : structural discrepancies in the resistive material, the presence of surface contaminants, pixel individual history, variations of the number of active tips from one pixel to another. The relative importance of each one of those factors is however difficult to single out.

We monitored the temporal evolution of the energy spectra, for electrons emitted by pixels (17,15) (green and black curves) and (9,14) (red and yellow curves) over time intervals of respectively one month and 3 months. The corresponding effective emission times (involving operations at different extraction voltages and a constant pressure level) were respectively of 146 and 104 hours. Between the beginning and the end of those periods we observed a decrease of the mean energies of the distributions close to 2eV in the case of pixel (17,15), and close to 1eV in the case of pixel (9,14). This decrease is consistent with the alteration of the resistance between tips and cathodic contacts. Those measurements give indications on variations of the energetic structure of emitted electrons which are important to consider in order to optimize parameters of ion sources optics.

Results of the angular analysis

Measurements presented here have been carried out at an extraction voltage of 60V.

Electrodes potentials were set in the same configuration than described in section 4.2.

By simultaneously measuring the position of the sample holder central axis along two orthogonal X and Y axis (cf. Figure 4), and the collected current, we could derive the shape of spot due to the emitted beam on the analyzer entrance plane, as well as the angular dispersion of the beam. For that purpose the position of the emitting area with respect to the sample holder central axis has been measured and is known with a precision of ±100μm. A numerical model of the experiment and emitters (see part 5) then allows to determine a beam angular profile.

Positioning the component from one step to the following takes on average 10 seconds.

Since the temporal variability of the beam can be important with respect to this time scale, our observations are based on averaged current values at each (X,Y) position of the emitter.

Figure 9 shows a 3D profile of the beam (intensity as a function of the spatial coordinates on the entrance place) emitted by pixel (9,15). The maximum amplitude of the average collected current measured in the pulsed mode is close to 500pA. The collected current reaches a few pA on the external boundary of the beam. We define the beam boundary by setting a threshold at 5% of the maximum current value. Figures 10 and11 show normalized current profiles measured respectively along the X and Y axis, for the set of five pixels considered in this study. Each curve is an average profile of 3 passes of the emitter along the axis (each point is an average over five values successively acquired at 1s time intervals). The distance travelled by the component between measurement positions is 1mm, in order to minimize the total integration time and variations of the emitted current over the whole sequence. As illustrated with pixel (17,15) profile on Figures 10 and11, increasing the spatial frequency to 0.5 mm does not significantly affect the profiles shape.

Normalized profiles shown on Figures 10 and 11 exhibit similar slopes, while maxima are shifted along the displacement axis, consistently following the emitting pixel position.

Since electron trajectories depend on non homogeneous grids potentials and potentials of nearby electrodes, direct calculation of the angular profile is not possible from geometrical considerations only. Therefore we have used a semi-empirical model of the emitters, presented in the following section, which has been integrated into a numerical model of the experiment.

Along the Y axis, the collected current decreases very steeply when the distance from the maximum position increases, with an average 2.25 mm width at half maximum and 4.61 mm on the entry plane of the analyzer. Along the X axis, the slope is shallower with an average 3.44 mm width at half maximum and 7.95 mm on the entry plane. This indicates an asymetrical repartition of the current with respect to the analyzer axis. Those values vary by about 20% over the set of tested pixels.

Clearly a homogeneous beam density can not be expected from those kind of emitters, as demonstated by beam section imaging on phosphorous targets in [START_REF] Purcell | Nanoprotrusion model for field emission from integrated microtips[END_REF] and [START_REF] Constancias | Emission d'Electrons à partir de Micropointes pour Ecrans Plats : Simulation, Caractérisation et Confrontations Experimentales[END_REF]. Indeed, the nanograin structure of emitting centres covering the tip surface as well as charging effects near the extraction grids affect the spatial sidtribution of the emitted current. However, those effects are smoothed out as the emitter number is increased, and can not be responsible for the observed asymmetry of the beam.

Interpretation

This section provides an interpretation of the data shown on Figure 7, 10 and 11. An emphasis is put on deriving parameters useful to optimize ion sources optics, such as those allowing to model energetic and angular distributions of emitted electrons.

Energy distribution of emitted electrons

As reported in [START_REF] Baptist | Microtips and resistive sheet : A theoritical description of the emissive properties of the system[END_REF] the resistive sheet (silicon layer shown on Figure 3) plays the role of an electrical ballast able to :(1) regulate emitted currents as a function of tips locations on the sheet, (2) allow a uniform emission, (3) protect the FEA from destruction through arcing between tips and the extraction grid. The resistive sheet also acts as a coupling device between tips, since current tubes flowing from each tip to the closest cathodic contact cross the common resistive structure. Such current induces a voltage drop between the emitting tip and the extraction grid, therefore affecting the potential of every other tip. Since one tip typically emits 1μA at an extraction voltage of 100V, with a resistance value of 10 7 Ω the voltage drop per tip reaches values higher than 10V.

Another study [START_REF] Purcell | Nanoprotrusion model for field emission from integrated microtips[END_REF] reports detailed measurements of the energetic structure of the emitted beam from 1 to 4 mm 2 FEA, using a 135 • hemispherical energy analyzer with a nominal resolution of 10 meV, at four stages of operation of the FEA :(1) during the first emission, (2) during the seasonning process, (3) after the seasonning process, (4) before and after flash experiments. The base vacuum was 5.10 -11 torr and the samples were not baked before field emission, which is similar to our case. The observed energy spectra proved to be rapidly evolving with main and secondary peaks appearing and disappearing during the seasonning process. Increase of the extraction voltage caused an increase of the distance between peaks with typical widths of the whole distribution ranging from 10V up to 30V.

Such observations were interpreted with a model accounting for the presence of nanometric structures on the surface of the tip material (relevant when tips surfaces are not cleaned by heating), and mobile surface contaminants. Following this model electrons are emitted from localized energy bands associated with those structures, which are energy-shifted with respect to the Fermi level of the tip material and between each other.

Emission features presented in this paper are consistent with observations reported both in [START_REF] Purcell | Nanoprotrusion model for field emission from integrated microtips[END_REF] and in [START_REF] Baptist | Microtips and resistive sheet : A theoritical description of the emissive properties of the system[END_REF]. Indeed, the mean energy of the beam inferred from Figure 7 is much lower than the value expected from clean tips with no resistive structure, and the energy spread is on the contrary much larger than the theoretical behaviour observed from single tips, reaching roughly 10eV at 100V in our case.

Mean Energy of emitted electrons :

Using a Gaussian distribution of the energy of emitted electrons in our simulation, we fitted the current curves shown on Figure 7. From those fits we obtain a range of energy dispersions of the beam between 1eV and 4eV, and a range of voltage drops between 5 and 25eV, when the extraction voltage varies between 40V and 100V. Since our samples are not baked and very probably contaminated (our base vacuum level is larger by 3 orders of magnitude compared to that of [START_REF] Purcell | Nanoprotrusion model for field emission from integrated microtips[END_REF]), effects of granular surface structures and contaminants certainly account for a significant part of the observed energetic spread.

On the tip scale, the resistive structure significantly influences the energy of emitted electrons. This structure is modelled on Figure 12, in which R1 represents the resistance of the silicon layer between the cathodic contact and the tip base . Since the structure is symmetrical with respect to the tip axis, Figure 12 Currents are noted as following : i e is the total emitted current, i g is the current intercepted by extraction grids, i f is the current flowing from the cathodic contact towards the silica, which divides into currents i f 1 and i f 2 . The current flowing through cathodic contact is noted i c , and is the sum of i e and i f . R2 + λR1 is the sum of the silica resistance R2, and of an unknown resistance λR1 encountered in the resistive sheet by the leakage current flowing from the extraction grid to the cathodic contact (i f 2 in the silica layer and i f in the resistive sheet).

It represents the isolation between the cathodic contact and the extraction grid. The resistance between the extraction grid and the tip is the sum of R2 and of a lateral resistance, noted μR1, corresponding to the physical volume of the resistive sheet contained in the current tube flowing from the insulator base to the tip base (corresponds to i f 1 ). Measurements that we have carried out (see below) tend to show that a non negligible leakage current between the extraction grid and the cathodic contact is present during the emission process. The equilibrium voltages of the tips will therefore also depend on the behaviour of the resistive sheet regarding the leakage current. The potential at the silica/silicon interface associated with this current may therefore be different from the nearby tip potential, hence inducing an equilibrium current i f 1 .

We adopt here a simple resistive model to describe those different effects, assuming that resistances associated with currents i f and i f 1 (which cross the resistive sheet as does i e ) can be seen as proportional to the resistance R1.

We moreover consider that extraction grids collect very few electrons (the seasonning process of the tips is achieved). Therefore we assume a leakage current i g arising partly from such electrons, and from surface currents flowing on the silica and silicon exposed to the vacuum (due for instance to carbon deposition during the emission or fabrication process (Muller, personnal communication)), being less than 10% of the emitted current. This is consistent with current values observed from the difference between the collected current at the entrance plane of the analyzer, and the emitted current measured at the cathodic contact.

In order to estimate R2 + λR1, we carried out measurements of the pixel isolation when the cathode is not active. For those measurements the emitter is polarized in such a way that the field emission process is nullified (therefore i f =i f 2 ).

Assuming that the isolation resistance remains equal when the emitter is polarized in normal mode, we consider this measurement as a good estimate of R2 + λR1. Results for a set of 10 pixels are shown on Figure 13. The pixel isolation is observed to vary as a function of the extraction voltage. This is due to both the semi-conductive nature of the silica/silicon material, and to variations of the resistive sheet equivalent resistance with the extraction voltage as reported in [START_REF] Baptist | Microtips and resistive sheet : A theoritical description of the emissive properties of the system[END_REF]. The range of observed values extends from several tens of MΩ near 0V to about 2 MΩ for V gc = 100V . This implies that a leakage current as high as 50μA per pixel is likely to occur in the cathodic contact at this voltage level.

We find that this behaviour can be modelled by a power law of the form :

R2 + λR1 = R 0 • V gc -α (1)
where α takes values between 0.4 and 0.5, depending on the quality of the isolation at the moment of the measurement. Therefore such a leakage current crossing a resistive sheet of average resistance 170kΩ (estimate of the equivalent resistance for the whole pixel structure), leads to an average 8.5V voltage drop at the tips, which is of the same order than the voltage drop due to the emitted current crossing the resistive sheet.

Indeed, in order to estimate an equivalent resistance for the pixel as seen by the emitted current we also have to consider that, due to the geometry involved, the current tube flowing from the tip base to the cathodic contact only affects the nearby area of the resistive sheet.

Moreover, although there are 1760 tips per pixel, the proportion of emitting tips is actually close to 50% [START_REF] Constancias | Emission d'Electrons à partir de Micropointes pour Ecrans Plats : Simulation, Caractérisation et Confrontations Experimentales[END_REF]. Therefore using a value for R1 which is 300MΩ (Muller, personal communication), the equivalent resistance seen by the emitted current is 340kΩ. This value causes again a mean 8.5V voltage drop assuming an average emitted current of 25μA per pixel at an extraction voltage of 100V.

We can now estimate the mean voltage drop at the tips, writing the following expression of Vg-Vp :

V gp modelled = V g -V p = (V g -V c) 1 - 1 A -k × i e B A ( 2
)
where :

A = 1 + μ + 1 + 1 + μ λ R2 R1 B = μR1 + R2 1 + μ λ (3)
Coefficient k accounts for the proportion of the emitted current collected by the extraction grids. In equation ( 2) i e represents the current measured at the analyzer entrance, we therefore assume that a current k × i e (with k > 1) is actually emitted. This has no effect on the model considered on Figure 12, but helps to figure out the influence of the intercepted current (i g = (k -1)i e ) on the potential difference (Vg-Vp).

Using measured values of R2+λR1 for pixel (17,15), we studied the influence of parameters k, λ, μ, and R1, on the potential difference (Vg-Vp). Results are summarized on Figure 14.

On this figure the green line with squares shows the potential difference (Vg-Vp) associated with mean energy values e(Vg-Vp)=e(Vg-Vc)-e(Vp-Vc) of the beam deduced from data corresponding to pixel (17,15) as shown on Figure 12. We carried out a parametric study by varying μ from 0 up to 0.5, λ from 0.5 up to 2, R1 from 170kΩ up to 680kΩ (equivalent to a proportion of emitting tips from 100% down to 25%), and k from 1.1 to 1.5 (around 10% of the emitted current in usually lost in the extraction grids during nominal operation of the FEA). On Figure 12 we plotted four sets of curves in the parameter space, illustrating the most important features. For each set of curves λ has been given 4 values which are 0.5, 1, 1.5, and 2.

Increasing λ reduces the discrepancy between the modelled and observed (Vg-Vp). Comparing the black set of curves (corresponding to R1 value of 170kΩ), with the cyan (R1 = 340kΩ), and the yellow (R1 = 680kΩ) sets of curves, we note that increasing the proportion of emitting centres (all other parameters staying equal) increases the discrepancy between observed and modelled values of (Vg-Vp). Also, increasing μ from 0 to 0.2 increases this discrepancy, as can be seen from the red and blue sets of curves (the effect however remains smaller in this case).

At last, increasing the leakage current at the grids from 10% to 50% of the emitted current decreases the discrepancy between the modelled and measured values, as can be observed from the cyan and blue sets of curves.

The main features derived from the modeled values of (Vg-Vp) (all of which are not shown here) are the following :

-the Vgp model usually overestimates the observed values : therefore the resistive structure alone can not explain the observed energy shift, -a proportion of emitters lower than or equal to 50% is required to allow parameters λ and μ to stay in a realistic range of values (that is λ values close to 1 and μ values close to 0), -increasing the current intercepted by the grids lowers the discrepancy between model and observations.

We conclude that our model gives a satisfactory qualitative description of the resistive contribution to the mean energy of the emitted beam. Based on that, in order to build a phenomenological model of the emission, we now propose a simple statistical description that accounts for the effect of nanometric protrusions at the tip surface.

The potential difference V pc between the tips and the cathodic contacts can be written as the sum of a resistive contribution V p R and a contribution due to the material granularity

V p n : V pc = V p R + V p n (4)
The V p R component describes the influence of both the resistive sheet and that of the insulating layer on the mean tip potential. From previous considerations we can write :

V p R = V gc applied -V gp modelled (5)
where V gc applied = V g -V c and V gp modelled in given by equation (2).

Equation ( 4 According to this approach we obtain that nanometric surface structures accounts for about 50% of the mean value of V pc.

Energy dispersion of emitted electrons :

While we previously worked on the pixel scale,

we now scale the study on the tip scale in order to consider dispersion of the parameters inside one pixel. The resistive model proposed on Figure 12 is also valid on the tip scale. [START_REF] Purcell | Nanoprotrusion model for field emission from integrated microtips[END_REF] demonstrated that nanometric protrusions were responsible for a discrete and shifted peaking of the energy spectra therefore leading to a large overall energetic dispersion. We choose to describe such behaviour in a practical and representative way by using two parameters, the mean energy eV p n and the dispersion σ n , associated to a gaussian envelope of the spectra noted Γ(eV p n , σ n ).

We assume that the energy dispersion of electrons in the beam is due to two terms. First, the dispersion of R1 values across the pixel (based on [START_REF] Baptist | Microtips and resistive sheet : A theoritical description of the emissive properties of the system[END_REF] and Constancias (1998) (the dispersion of field enhancement factors across the pixel is assumed to be negligible in comparison). Second, the dispersion due to nanoprotrusions, described by an ad hoc parameter σ n . The associated total dispersion can therefore be written :

σ p = σ p R + σ pn (6)
Differentiation of (5) relatively to R1 allows to derive the resistive contribution :

δV p R = 1 A δR1 R1 ki e R1 - B A (1 + μ -A) -V gc 1 + μ A -1 (7)
The influence of nanometric structures can be in principle described by using the Fowler-Nordheim energy distribution, but is more complex to evaluate. A high field enhancement indeed occurs at the surface of the nanometric objects [START_REF] Purcell | Nanoprotrusion model for field emission from integrated microtips[END_REF], such that the electric field penetrates deeply into the material, allowing electrons from lower energy levels to be extracted. Therefore the corresponding energy spread is in principle proportional to the full width at half maximum of the Fowler-Nordheim distribution function that we recall here :

dJ dE • 1 J 0 = exp /d d [1 + exp /k B T ] (8)
where the Boltzmann constant is noted k B , the temperature T, J 0 is the emitted current density at zero temperature, = E -E F the relative energy with respect to the Fermi level of the emitter, and d a function of the applied electric field F (V.m -1 ), and work function Φ(eV )

of the emitter :

d = 0.93.10 -10 F √ Φ (9)
Since the Full Width at Half Maximum of the FN distribution is related to parameter d (see [START_REF] Brenac | Caractéristiques énergétiques de cathodes à micropointes à émission de champ[END_REF] for instance) the σ pn parameter derived from our measurements could in principle be used to estimate an equivalent electric field enhancement factor describing the influence of nanoprotrusions on the emission. Such calculations are however beyond the scope of the present study, but interesting to mention. Following the same approach than for deriving mean energy values, we calculate values of the parameter σ pn by adjusting σ p to the energetic dispersion derived from the numerical simulation of the experiment, assuming a gaussian distribution of the energy of the emitted electrons.

We assume a value of 10% for δR1 R1 across the tips in one pixel. Figure 16 shows the simulated mean energy and energy dispersion of electrons for each extraction voltage, constrained by our measurements. The energy dispersion increases with the emitted current, consistently with [START_REF] Purcell | Nanoprotrusion model for field emission from integrated microtips[END_REF].

Angular distribution of emitted electrons

Null potentials are applied to all extraction grids except those located just above the emitting tips, therefore, the emitted beam is focalised along the Y axis, and defocalised along the X axis. This situation can be seen on Figures 10 and11, which are respectively cuts along the X and Y axis, associated to the above configuration, for an idealized axisymetrical beam. The analysis of simulated electrons trajectories reveals that some electrons are deflected backward and will return to the cathode, most of the time charging the isolating layer (the grid to silica area ratio is about 15%).

Simulation results shown on Figures 10 and11, have been obtained with the assumption of a 45 • half angle dispersion of the beam. Over the set of five pixels tested we find an average half angle of 42 • ±20%. We think that this value is representative of the actual angular dispersion of the tested sample. In our model the number of emitted electrons decreases as a cosine function of the emitted angle. This assumption is in good agreement with the observed profiles.

Conclusions

We used a hemispheric retarding field energy analyzer [START_REF] Cipriani | La couronne neutre suprathermique de Mars : modélisation et développement d'une instrumentation pour sa mesure in situ[END_REF] We derived the energy distribution of the electrons as a function of the extraction voltage with a resolution of 100meV. Those measurements proved to be stable and repeatable. The mean energy and FWHM of the electrons are found to increase with the extraction voltage.

The FWHM varies from 2eV at 40V to 9eV at 100V, consistently with observations carried out in [START_REF] Purcell | Nanoprotrusion model for field emission from integrated microtips[END_REF]. As a consequence the mean energy is expected to vary across the emission life time of the device, at a rate depending both on the alteration of the resistive structure with time, and on the fate of adsorbed contaminants at the tip surface. We also found a mean angular dispersion of the beam of 42 • ± 20%, over the set of tested pixels. We used those observations to derive a simple phenomenological model of the emission accounting for the resistive structure of the component, as well as for the influence of nanoprotrusions on the emitted energy distribution, and for the mean angular dispersion of the beam. Such model can be used in order to design and optimize ion sources integrating future space instruments.

Space applications using FEA as electron guns suppose working in vacuum levels of the order of 10 -6 mbar. The gaseous environment surrounding some planets and satellites in the Solar System, such as Mars or Europa, contains variable amounts of oxygen, carbon based molecules, and water vapor, which is also the main component outgassed from vessels structures. Some of those gases have been shown to be critical for the lifetime of FEA, including those based on microtips. For those reasons emission performances of such electron emitters, whatever their nature, have to be monitored in representative gaseous environments in order to infer their applicability to space applications. for pixel :(17,15) green, (9,14) blue, (27,8) red, (9,16) magenta, (9,15) cyan.

Figure 7

 7 Figure7shows the evolution of the Collector current as a function of the collector voltage

  represents half of the electrode structure encompassing each tip. However resistance values indicated on this Figure refer to the complete structure.

  ) is constrained by the mean energies of the beam fitting the measurements shown on Figure 7. Those mean energy values are respectively 35.03eV (for Vgc=40V), 50.70eV (for Vgc=60V), 64.20eV (for Vgc=80V), and 75.40eV (for Vgc=100V), and correspond to mean voltage drops at the tips shown on Figure 15. On that figure yellow triangles represent the observed values of the mean energy, and are adjusted by a second order polynomial plotted on the graph. The curve with squares shows corresponding Vgp values, obtained by subtracting the polynomial values from Vgc. The green curve with disks shows values of V gp modelled , obtained with realistic values of our parameters such that k = 1.1, μ = 0.1, λ = 1.0, and R1 = 340kΩ. Adjusted values of Vpc and V p n are respectively represented by the blue curve with diamonds, and by the red curve with stars.

  to characterize the electron beam emitted by Mo microtips Field Emission Arrays made by the CEA/LETI laboratory. Those components integrate extraction grids and a resistive sheet. Measurements were carried out as a function of the extraction voltage, at normal incidence of the beam with respect to the analyzer entrance, without sample baking.

Fig. 1 -

 1 Fig. 1 -. SEM image showing extraction grid holes (φ = 1.4μm), and underlying Mo tips. The mean distance between two holes centers is 3μm.

Fig. 2 -

 2 Fig. 2 -. Potential energy diagram of the experiment. Electrons having an energy E 0 with respect to the Fermi level of the tips material leave the tips at potential Vp, and are accelerated towards the grids at potential Vg. They have an energy close to the grids equal to e(V g -V p) -Φ grid where Φ grid is the work function of the grid material. Electrons leaving the grids are further accelerated towards the entry plane of the analyzer at potential Vd. They enter the retarding field controlled by(Vcoll-Vd). Their energy close to the collector surface is e(V coll -V p) -E 0 -Φ coll , where Φ coll is the collector's work function.

Fig. 3 -

 3 Fig. 3 -. Emitting unit structure and dimensions. A molybdenum tip is represented (in green), two cathodic contacts and extraction grids made of Niobium (orange), the silica insulating layer (yellow), and the resistive sheet made of silicion (grey). The equivalent electrical structure per tip is also shown.

Fig. 4 -

 4 Fig.4-. Experimental setup overview. The FEA is bonded in a PGA cavity (blue part on the drawing), plugged into an alumina support (green part). The emitting surface is parallel to the entrance plane of the analyzer (red line, named Diaphragm), which has a small entry hole (diameter 250μm). An alumina support with conducting concentric rings is placed at the back of the entrance plane, in front of the hemispheric collector. The on-axis part of the collector (centering collector) is a 1 mm 2 area insulated from the main collector part, allowing centering of the emitter with respect to the analyzer entry hole.

Fig. 5 -

 5 Fig. 5 -. Variations over 30h of the emitted current (one column of 32 pixels), for an extraction potential of 90V.

Fig. 6 -

 6 Fig. 6 -. Typical courant-voltage plot for one stabilized 32 pixels column. The inset shows the associated Fowler-Nordheim plot.

Fig. 8 -

 8 Fig. 8 -. Comparison of collected currents at an extraction voltage of 60V for different pixels :(17,15) 1-green,(17,15) 2-black, (9,14) 1-red, (9,14) 2-yellow, (27,8) cyan, (9,16) blue, (9,15) magenta.

Fig. 9 -

 9 Fig. 9 -. 3D profil of the collected current for pixel(9,15), whith V gc = 60V, V d = 80V . The current reported on the Z axis is a mean value (duty cycle of 10%).

Fig. 10 -

 10 Fig. 10 -. Normalized collected current as a function of the emitter position along the X axis,for pixel :(17,15) green, (9,14) blue, (27,8) red, (9,16) magenta, (9,15) cyan.

Fig. 12 -

 12 Fig. 12 -. Resistive structure of the emitting unit (a detailed description is given in the text).

Fig. 13 -

 13 Fig. 13 -. Evolution of the isolation between the grids and the cathodic contacts as a function of an inverse extraction voltage.

Fig. 14 -

 14 Fig. 14 -. Potential difference between the extraction grid and tips (Vg-Vp) noted Vgp, as a function of the applied voltage Vgc. The green curve represents values of Vgp deduced from measurements of Figure 12. Details on curve families are given in the text.

Fig. 15 -

 15 Fig. 15 -. Comparison of the modelled and measured energy shift as a function of the extraction voltage Vgc.

Table 1

 1 Adjusted values of parameters describing the emission. Mean energy values and energy dispersion deduced from measurements are respectively noted E moy and ΔE HW HM .

	shows adjusted


Jean Marie Le Breton
email: jean-marie.lebreton@univ-rouen.fr

Rodrigue Lardé

H Chiron

V Pop

Dominique Givord

Olivier Isnard

I Chicinas

A structural investigation of SmCo 5 /Fe nanostructured alloys obtained by high energy ball milling and subsequent annealing

come

A structural investigation of SmCo5/Fe nanostructured alloys obtained by high-energy ball milling and subsequent annealing

Introduction

Nanocrystalline exchange-spring "hard-soft" magnets are composite magnetic materials consisting of two magnetic phases, one soft and one hard, mixed in a nanostructure in order to combine the high coercivity of the hard magnetic phase with the high magnetisation of the soft magnetic phase. These composite materials exhibit magnetic properties which are interesting both from a fundamental point of view and for applications [1][2][3][4][5]. The exchange-spring behaviour can be understood on the basis of the intrinsic parameters of the hard and soft magnetic phases which are coupled by exchange [6,7]. However, the role of the microstructure in the spring mechanism is not well understood [7][8][9][START_REF] Skomski | Permanent Magnetism[END_REF]. In previous investigations, we established the determinant influence of the microstructure on the hardsoft exchange coupling in SmCo 5 /α-Fe powders obtained by mechanical alloying [START_REF] Pop | [END_REF][12][13][14]. Mechanical alloying is a widely used preparation technique to obtain nanocrystalline structures exhibiting original magnetic properties including exchange spring [1,[15][16][17][18][19][20][21][22][23][24]. It involves the synthesis of materials by high-energy ball milling, in which pure or pre-alloyed elements are milled to achieve alloys or composite materials formation. We concentrate here on the mechanical milling technique to blend two different magnetic phases. We applied mechanical milling to synthesize SmCo 5 /α-Fe nanocomposites starting from a mixture of SmCo 5 powder (milled for two hours) and 20 wt% of elemental iron. The coercivity and the remanence were found to be dependent upon the process conditions and they could be improved by adjusting the milling and/or heat treatment conditions (time and temperature) [START_REF] Pop | [END_REF][12][13][14]. Microstructural investigations of as-milled and annealed powders showed that the initial magnetic phases are destroyed upon milling, and that annealing the as-milled samples permit recovering the initial phases [START_REF] Pop | [END_REF]13,14]. However, our previous Mössbauer investigations suggested that Fe/Co interdiffusion occur during the process [START_REF] Pop | [END_REF]. Since then, this result has been confirmed by other authors [25][26][27].

In the present study, we proceed to a systematic study of the microstructure of SmCo 5 /α-Fe alloys by Mössbauer spectrometry with the aim to follow accurately the evolution of the Fe-containing phases upon milling and subsequent annealing. The results are conforted by tomographic atom probe analysis and interpreted in relation with the magnetic properties of the materials, by considering the influence of the Fe/Co interdiffusion on the exchange coupling phenomenon.

Experimental

An alloy ingot with the SmCo 5 nominal composition was prepared by induction melting under argon atmosphere. The purity of the elements was 99.9%. The ingot was crushed into small pieces and it was subsequently mechanically milled for two hours in a high energy planetary mill under argon atmosphere. The SmCo 5 powder thus obtained was mixed with an iron powder (grain size below 40 µm) in a ratio of 80 (SmCo 5 )/ 20 (Fe) weight percent, corresponding to a ratio of 75 (SmCo 5 )/ 25 (Fe) atomic percent. The mixture was mechanically milled under argon atmosphere using the high-energy planetary mill. Because the rare earths are very sensitive to oxygen, for the preparation of rare-earth permanent magnetic materials it is necessary to mill in inert atmosphere in order to avoid rare earth oxidation. Several milling times were used ranging from 2 to 8 hours.

In order to investigate the influence of the annealing on the evolution of the SmCo 5 /α-Fe exchange coupling, samples made of the as-milled powder were sealed inside evacuated silica tubes for further heat treatments. These treatments were performed at temperatures ranging from 450 to 650 °C for 0.5 to 10 hours. The maximum annealing temperature of 650 °C had been chosen in order to avoid the crystallite growth. 57 Fe Mössbauer spectrometry experiments were performed at room temperature in transmission geometry, using a 57 Co source in a rhodium matrix. The Mössbauer spectra were fitted according to a least squares method [28]. The isomer shift (relative to metallic α-Fe at room temperature) and hyperfine field are denoted δ and B respectively. Estimated errors for the hyperfine parameters originate from the statistical errors σ given by the fitting program, taking 3σ. X-ray diffraction (XRD) patterns were recorded in the angular range 2θ = 25 -100°. For these experiments a Brucker D8 powder diffractometer with the Co(Kα 1 ) radiation ( = 0,1788970 nm), was used.

The powder milled for 8h was analysed by Laser Assisted Wide Angle Tomographic Atom Probe (LAWATAP). The LAWATAP is a 3D analytical high resolution microscope which provides the spatial distribution of atoms of the analysed specimen . The principle of atom probe is based on the field evaporation of surface atoms as ions and their identification by time of-flight mass spectrometry. The high electric field required (a few tens V/nm) is obtained by applying a high voltage to the specimen prepared in the form of a sharply pointed needle (tip radius less than 50 nm). The specimen preparation involves a two-step procedure [29]. In the first step, a single powder particle between ~1 and ~50µm is mounted onto a stainless steel support tip using an ex-situ micro-manipulator and an optical microscope. In the second step, the particle is annularly milled into a very sharp tip by Focused Ion Beam. The tip is biased at a high positive DC voltage V o in the range 2-10 kV and cooled to low temperatures (20-80 K). The vacuum in the analysis chamber is usually in the range 10 -8 Pa. Surface atoms are then field-evaporated by means of high-frequency (100 kHz) laser pulses superimposed on the DC voltage V o [30]. The evaporated atoms are collected by a time-resolved position detector placed in front of the specimen. This allows measuring the time of flight of each ion and recording its impact position. Both informations allow the chemical nature of the evaporated ions to be determined and the position from which the atoms originate at the tip surface to be derived. The atom lateral positions at the surface of the tip sample are obtained from a reverse projection of ion impacts coordinates on the position sensitive detector. The depth position is deduced from the order of detection of atoms [31]. The sample is field-evaporated atomic layer by atomic layer, and a typical volume of analysis of 10 4 nm 3 is collected in a few hours. After 3D reconstruction of the analysed volume, the spatial distribution of atoms is observable at the atomic scale in the real space. From the data set of the 3D reconstruction, the chemical composition or the concentration depth profiles can be calculated everywhere in the analysed volume. More details about tomographic atom probe can be found in references 32 and 33.

Results

Structural analysis of annealed powders

The XRD patterns of the powder milled for 8 h and subsequently annealed are shown in figure 1. The pattern of the as-milled powder shows only a broad and low intense peak corresponding to α-Fe. As the annealing temperature increases, the width of the α-Fe peak decreases, and the intensity of the peaks corresponding to the Sm-Co phase increases. This is in agreement with the recrystallization process of the hard and soft phases induced by the annealing treatment and reported before [START_REF] Pop | [END_REF]13,14]. Additionally, the main peaks of the Sm 2 O 3 oxide are observed, indicating that some oxidation occured during the annealing treatment. In the pattern of the powder annealed 30' at 650°C, the position of the most intense α-Fe peak corresponds to a lattice parameter a = 0.2857 ± 0.0005 nm. This value is lower than that of the pure α-Fe phase (0.2866 nm). Considering that the Co atomic radius is lower than the Fe atomic radius, this reveals the presence of Co in the α-Fe phase.

The Sm-Co peaks can be indexed in the hexagonal P6/mmm system with lattice parameters a = 0.4900 ± 0.0005 nm and c = 0.4100 ± 0.0005 nm. These lattice parameters are respectively lower and higher, than those of the pure SmCo 5 phase (0.4995 nm, 0.3978 nm). Such variations in the lattice parameters are typical of the introduction of transition metal dumbbells in the structure [34,35]. This classically occurs when Fe atoms are introduced in the SmCo 5 phase, in agreement with the fact that the pure SmFe 5 phase does not exist [START_REF] Burzo | Magnetic properties of metals[END_REF]. Note that in the following we refer to the formation of the Sm(Co-Fe) 5 alloy whereas, as Fe enters into the phase, the stochiometry is expected to become enriched in transition metal as a result of dumbbell formation.

The grain size has been estimated for both soft and hard magnetic phases, according to Scherrer's formula. A size less than 30 nm is obtained for the samples milled for 6h or more. For the samples milled for 8h, the α-Fe grain size is derived to be of about 18 nm.

Mössbauer analysis: Fitting procedure

All the Mössbauer spectra were recorded at room temperature. They were fitted consistently throughout the whole series of samples, with the following contributions:

-a magnetic contribution with a hyperfine field equal or close to that of α-Fe (33.0 T), -a second magnetic contribution with a hyperfine field of about 27.5 T, -a paramagnetic contribution, to account for the excess of absorption at the centre of the spectra, when detected. The magnetic contribution of the α-Fe phase has been systematically considered. To take into account its large width in different spectra, this contribution has been fitted with several components, that is a distribution of magnetic sextets. The second magnetic contribution is characterized by an hyperfine field of about 27.5 ± 0.3 T, which is typical of a rare-earth / iron compound. This value is in agreement with the room temperature mean hyperfine field of dilute 57 Fe nuclei in SmCo 5 compound [START_REF] Forker | [END_REF]. Consequently, this magnetic component has been attributed to Fe atoms present in Sm-Co-rich regions. This is in agreement with the XRD analysis, thus confirming that the Sm-Co regions destroyed upon milling recrystallize as Sm(Co,Fe) 5 In the following, the second magnetic contribution will be denoted as a Sm(Co,Fe) 5 . Preliminary fittings were performed on the spectra where the relative intensity of this contribution is the most important of the series (figure 2). They showed that the Sm(Co,Fe) 5 contribution can be fitted with three components, with hyperfine fields in the range 26.2-29.5 T (these three components are displayed in figure 2, with the resultant contribution). This has been used for all the spectra. In some spectra, a paramagnetic contribution has been fitted with a weak relative intensity, which is attributed to Fe-containing paramagnetic phases such as oxides or nanoparticles which are superparamagnetic at room temperature.

As-milled powders

The Mössbauer spectra of the as-milled powders are shown in figure 3. The Mössbauer spectrum of the sample before milling consists uniquely in a magnetic sextet typical for α-Fe. In the spectra of the as-milled powders, an asymmetric broadening of the lines is observed, as revealed by the shoulders on the external lines of the sextet (see figure 3). In the analysis, this broadening is taken in account by considering magnetic sextets with hyperfine fields between 34 and 36 T, higher than the hyperfine field of pure α-Fe (33.0 T). The fact that both the mean hyperfine field <B Fe > and the mean isomer shift <δ Fe > of this magnetic contribution are larger than α-Fe associated values (see figures 4a and4b) is characteristic of the presence of Co atoms in the α-Fe regions, with a proportion that increases with the milling time [38]. It is worth to note that similar observations were reported for Fe 50 Co 50 mechanically milled powders [39]. As the milling time is increased, another magnetic component develops (displayed in figure 3), with a mean hyperfine field of 27.5 ± 0.3 T, characterising Fe atoms having entered in the SmCo 5 phase. The relative intensity of this contribution increases with milling time (figure 4c). The relative intensity of this contribution being proportional to the Fe content in the Sm(Co,Fe) 5 regions, its increase is in agreement with an increase in the proportion of Fe atoms in the Sm-Co-rich regions.

(2)

Velocity ( mm / s ) 0 -11 +11 Velocity ( mm / s ) 0 -11 +11 0.99 1.00 Absorption ( % ) (1) (3)
(2) The presence of Co in Fe and that of Fe in Sm-Co-rich regions is confirmed by LAWATAP analysis of the powder milled for 8h. A slice taken from the 3D reconstruction is shown in figure 5 and presents the spatial distribution of Sm, Co and Fe atoms. The figure evidences the nanostructure of the sample and reveals that the distribution of Sm is not uniform. Sm atoms are only detected in Sm-Corich regions. The regions where Sm atoms are not detected correspond to the Fe-rich regions, which are initially occupied by the α-Fe phase. Co atoms are detected in both the Sm-rich and the Fe-rich regions, with a concentration that is higher in the Sm-rich regions. Finally, Fe atoms are detected in all regions, i.e. they are also present in the Sm-rich regions. These results clearly show that Co/Fe interdiffusion occurred during milling. A more detailed LAWATAP analysis will be published elsewhere [40]. These results clearly show that high energy ball milling of SmCo 5 and α-Fe leads to a partial mixing down to the atomic scale, leading to the diffusion of Fe atoms in Sm-Co-rich regions and to the diffusion of Co atoms in Fe-rich regions. The Co/Fe interdiffusion increases with the milling time.

Annealed powders

The influence of the annealing time on the structure of the powders was investigated on samples initially milled for different times, upon annealing at 450°C. For example, the Mössbauer spectra of the powder milled for 6h and annealed for times up to 10 h is shown in figure 6. As compared with the spectrum of the as-milled powder, the spectra of the annealed powders reveal an increase of the relative intensity of the Sm(Co,Fe) 5 contribution with the annealing time. Moreover, the mean hyperfine field <B Fe > of the α-Fe(Co) contribution changes with the annealing time. The same observations are made for the other samples. The evolution of the Sm(Co,Fe) 5 relative intensity and that of the α-Fe(Co) mean hyperfine field are gathered in figures 7a and 7b, respectively. For all the samples, an important increase of the Sm(Co,Fe) 5 relative intensity occurs for short time annealing, followed by a plateau. This suggests that annealing promotes an increase of the Fe content in the Sm-Co-rich regions, up to a maximum value. The evolution of the mean hyperfine field of the α-Fe(Co) contribution shows a similar behaviour, except for the sample milled for 8 h, for which the annealing leads to a decrease of <B Fe > before reaching the plateau. This does not mean that the Co content in Fe-rich regions decreases, if one considers that, according to [38], in Fe 1-x Co x alloys with x > 0.30, the mean hyperfine field decreases as x increases. Accordingly, this suggests that for milling times lower than 6 h, the Co content in the α-Fe(Co) regions is lower than about 25 at.% in the asmilled samples and increases with the annealing time, while for 8 h milling, the Co content in the α-Fe(Co) regions is higher than about 25 at.% in the as-milled samples and increases with the annealing time as well. The influence of the annealing temperature on the phases involved is investigated for the powders milled for 8 h. The Mössbauer spectra of the powder before and after annealing for 0.5 h at temperatures up to 650°C are shown in figure 8. As the temperature increases, the relative intensity of the Sm(Co,Fe) 5 contribution increases noticeably to reach about 60% at 650°C. This shows that the Fe content in the Sm-Co-rich phase increases with the annealing temperature. This is accompanied by a decrease of the mean hyperfine field <B Fe > of the α-Fe(Co) contribution, related to a concomitant increase of the Co content in the α-Fe(Co) regions, and characteristic of a Co content higher than about 25 at.% as explained above [38]. These results show that annealing up to 650°C leads to an important increase of the Co/Fe interdiffusion. The relative intensity of the Sm(Co,Fe) 5 contribution after annealing at 650°C being 61 % indicates that more than 60% of the Fe atoms are present in the Sm-Co-rich regions. Thus, the tendency observed for long time annealing at 450°C is also observed for short time annealing at higher temperatures. Diffusion of Fe atoms in the Sm-Co-rich regions and diffusion of Co atoms in the α-Fe regions occur simultaneously leading to the fact that both the Fe content in Sm(Co,Fe) 5 and the Co content in α-Fe(Co) drastically increase with the annealing temperature. Such a structural modification of the phases involved is fully consistent with the evolution of the magnetic properties previously described in this type of alloys [START_REF] Pop | [END_REF][12][13][14].

Co/Fe interdiffusion and magnetic properties

It is well known that the mean hyperfine field <B Fe > of the α-Fe(Co) contribution is related to the Fe content in the α-Fe(Co) phase [38]. As the intensity of the Sm(Co,Fe) 5 contribution, denoted as % SmCoFe , is proportional to the Fe content in the Sm(Co,Fe) 5 phase, the Co/Fe interdifussion process during both milling and annealing can be represented by plotting the evolution of <B Fe > as a function of % SmCoFe (figure 9). The experimental points reveal a behaviour that is in full agreement with the known evolution of the hyperfine field of Fe 1-x Co x alloys with x, as shown by the insert to figure 9. According to the nominal volume fractions of the SmCo 5 and α-Fe phases and neglecting the slight difference in density of these phases, the parameter % SmCoFe corresponds exactly to the Co content in the α-Fe(Co) phase. This is confirmed by tomographic atom probe analysis as well. For example, in the powder annealed at 550°C for 0.5 h the measured Co content of the α-Fe(Co) phase (40±8 at.%) is fully consistent with the Mössbauer relative intensity of the Sm(Co,Fe) 5 contribution (39 %) [40]. Consequently, the evolution of <B Fe > with % SmCoFe reflects the evolution of <B Fe > with the Co content in Fe 1-x Co x alloys. Comparing the <B Fe >=f(% SmCoFe ) curve (figure 9) with the <B Fe >=f(x) curve (figure 9, insert), one can notice that the values of <B Fe > are lower than those expected for Fe 1-x Co x . Such a difference could be related to the small size of the α-Fe(Co) particles and to the presence of defects. Grouping all results, the various mechanisms involved in Fe vs Co interdiffusion can be summarized as follows: the increase of <B Fe > with % SmCoFe for low values of % SmCoFe corresponds to the diffusion of Fe atoms in the Sm-Co-rich regions. In this regime, the fact that <B Fe > increases with the milling time shows that the diffusion process is mechanically activated. Then, <B Fe > reaches a plateau for % SmCoFe values of about 20 to 30%, corresponding to a Co content in the α-Fe(Co) phase of 20-30 at.%. This is obtained for the powders milled for 6 and 8 h and annealed up to 10 h at 450°C, which are characterised by further Co/Fe interdiffusion. The decrease of <B Fe > observed for values of % SmCoFe higher than 30% corresponds to a further increase of the Co content in the α-Fe(Co) phase. This is obtained for the powders milled for 8 h and annealed at temperatures between 550 and 650°C. For complete Co/Fe mixing, not observed experimentally here, the relative proportions of transition metals in each phase would correspond to their relative proportions in the initial mixture, that is 71 at.% Co and 29 at.% Fe. Thus, % SmCoFe would be equal to 71% and, according to [38], <B Fe > = 33.2 T. The corresponding point reported in figure 9 is roughly aligned with the experimental points. It remains to be demonstrated whether increasing further the annealing temperature or the annealing time would lead to complete Co/Fe intermixing. Previous investigations of the magnetic properties of as-milled and annealed SmCo 5 /Fe powders revealed that the best magnetic properties are obtained for powders milled for 8 h and annealed at 550-600°C [START_REF] Pop | [END_REF]. This corresponds to a relative intensity of the Sm(Co,Fe) 5 contribution, and thus to a Co content in the α-Fe(Co) phase, comprised between 39 and 55 %. Annealing at such temperatures improves the exchange interactions [START_REF] Pop | [END_REF]. Considering the present investigation, such an improvement has to be related to the presence of both Co and Fe in all the magnetic phases. Strong Co/Fe interdiffusion may lead to better structural coupling between the Sm-(Co,Fe) and α-Fe(Co) phases, thus reinforcing the exchange interactions between the magnetically hard and soft phases. Such a positive effect of the Co/Fe intermixing on the exchange coupling has been reported in the case of SmCo 5 /Fe exchange spring multilayer films [41,42]. The evolution of the remanence in the present system is in rather good agreement with the evolution of the Co content in the α-Fe(Co) phase. The highest remanence is obtained for the powder milled for 8 h and annealed for 1.5 h at 500°C, which corresponds to a Co content of about 34 at.%, close to the maximum value of the magnetization of Fe 1-x Co x alloys [START_REF]Magnetism and Metallurgy[END_REF]. Annealing the powder for shorter times at the same temperature or at lower temperatures leads to lower remanence, as expected for Co contents below 28 at.%, and thus to a smaller value of the magnetization of Fe 1-x Co x phase. Annealing the powder at higher temperatures induces a remanence decrease, in relation with a decrease of the magnetization of the Fe 1-x Co x phase, its Co content being larger than 39%. The positive influence of an annealing treatment on the coercivity has to be related to structural modifications (grain size refinement and increase of the crystallinity), but it can be partly due as well to Co/Fe interdiffusion, through reinforcement of the intergranular exchange interactions. However, annealing at too high temperatures probably leads to an increase of the grain size, thus permitting soft phase magnetisation reversal independently of hard phase reversal. This explains the appearance of a shoulder on the demagnetization curves as shown in ref. [START_REF] Pop | [END_REF].

Conclusion

Milling a mixture of SmCo 5 and α-Fe magnetic phases leads to Co/Fe intermixing: Co atoms diffuse in the SmCo 5 phase and Fe atoms diffuse in the α-Fe phase. This interdiffusion process is mechanically activated. Annealing the powders at temperatures up to 650°C leads to the formation of α-Fe(Co) and Sm(Co,Fe) 5 phases. Increasing the annealing temperature leads to the increase in both the Co content in the α-Fe(Co) phase and the Fe content in the Sm(Co,Fe) 5 phase. These results help in interpreting previous magnetic measurements on this system. In particular, the strong Co/Fe interdiffusion evidenced in the present study may reinforce the exchange interactions between the magnetically hard Sm(Co,Fe) 5 and soft α-Fe(Co) phases.

Figure 1 .

 1 Figure 1. XRD patterns of the powder as-milled for 8h and annealed at the indicated temperatures and times. The peaks of the Sm(Co,Fe) 5 phase are indexed in the hexagonal P6/mmm system. ( Co(Kα1) = 0,1788970 nm)

Figure 2 .

 2 Figure 2. Room temperature Mössbauer spectrum of the SmCo 5 /Fe powder milled for 8h and annealed for 30 minutes at 650°C. The components of the Sm(Co,Fe) 5 contribution are displayed in blue, their hyperfine fields being 29.4, 28.5 and 26.8 T for components (1), (2) and (3) respectively. The resultant Sm(Co,Fe) 5 contribution is displayed in red.

Figure 3 .Figure 4 .

 34 Figure 3. Room temperature Mössbauer spectra of the SmCo 5 /Fe powder before milling and after milling for the indicated times. The contribution of the Sm(Co,Fe) 5 phase is displayed.

Figure 5 .

 5 Figure 5. 2D projections of elemental mappings of Sm, Co and Fe atoms in an analysed volume of the powder milled for 8 h, obtained by LAWATAP.

Figure 6 .

 6 Figure 6. Room temperature Mössbauer spectra of the SmCo 5 /Fe powder as-milled for 6 h and annealed at 450°C for the indicated times. The contribution of the Sm(Co,Fe) 5 phase is displayed.

Figure 7 .

 7 Figure 7. Evolution with the annealing time at 450°C of (a) the relative intensity of the Sm(Co,Fe) 5 contribution, and (b) the mean hyperfine field of the α-Fe(Co) contribution. The lines are guides for the eye.

Figure 8 .

 8 Figure 8. Room temperature Mössbauer spectra of the powder as-milled for 8h and annealed for 30 minutes at the indicated temperatures. The contribution of the Sm(Co,Fe) 5 phase is displayed.

Figure 9 .

 9 Figure 9. Mean hyperfine field of the α-Fe(Co) contribution as a function of the relative intensity of the Sm(Co,Fe) 5 contribution, for all the samples investigated. Insert: hyperfine field of Fe 1-x Co x alloys as a function of the Co content x in the alloy [38].


Anthony B Parmentier

Philippe F Smet
email: philippe.smet@ugent.be

Frank Bertram

Jürgen Christen

Dirk Poelman

Structure and luminescence of (Ca,Sr)2SiS4:Eu2+ phosphors

Keywords: thiosilicates, europium, light-emitting diodes

come

Introduction

The possible energy savings associated with the replacement of incandescent light bulbs (and even fluorescent tubes) by light emitting diodes (LEDs) have triggered new efforts in phosphor research [1].

The appropriate combination of a blue or near-UV LED with a single phosphor or a mixture of phosphors should lead to energy efficient lighting devices, with good colour rendering and wide colour temperature range. Europium doped sulfide phosphors have gained renewed attention as alternatives for oxide phosphors, because they are more suited to obtain saturated red emission [2]. The thiosilicate materials Ba 2 SiS 4 [3] and Ca 2 SiS 4 [3,4] were recently investigated for use as phosphor hosts in white phosphor Structure and luminescence of Eu-doped (Ca,Sr) 2 SiS 4 phosphors converted LEDs (pcLEDs). A third phosphor in this series, Sr 2 SiS 4 , has not been studied since the work on cathodoluminescence of thiosilicate materials by Avella [5] and Olivier-Fourcade [6] in the early 19seventies. Because the recent studies on the other thiosilicates shed a substantially new light on the photoluminescent (PL) behaviour of these materials, it seems suited to reconsider Sr 2 SiS 4 as well. As the results on (Ca,Eu)-thiosilicates [4] point out the very interesting possibility of tuning the PL characteristics by changing the composition, in the current work, a series of europium-doped (Ca,Sr)-thiosilicates are prepared and investigated. Phase compositions and lattice parameters are investigated and optimized using Rietveld refinement. Photoluminescence excitation and emission spectra are recorded, and sample homogeneity is studied using cathodoluminescence spectroscopy in a scanning electron microscope (SEM-CL). Finally, the thermal quenching behaviour of the luminescence is investigated, as this is an important criterion when phosphors are to be used as wavelength convertors for white light LEDs.

Experimental

All powders are synthesized by sintering a homogeneous mixture of stoichiometric amounts of SrS (Alfa Aesar, 99.9%), CaS (CERAC 99.99%), EuF 3 (Alfa Aesar, 99.5%) and Si (Alfa Aesar, 99.9%) in a continuous H 2 S-flow at 875°C for 1 hour. Hydrogen sulfide is used as a source of extra sulfur needed to form the thiosilicate. As suggested in [4], a slight surplus of Si (5%) is used. After cooling, the samples are manually ground with a mortar and pestle in order to get a fine powder. Except for the sintering, the processing of the samples is done in ambient air. Samples are prepared with a composition ranging from Sr 2 SiS 4 to Ca 2 SiS 4 , in steps of 10% Sr 2+ -substitution by Ca 2+ . Unless stated otherwise, all samples are doped with 4% Eu 2+ , meaning that 4% of the alkaline earth ions are replaced by Eu 2+ -ions. E.g., a 20% substitution of Sr 2+ by Ca 2+ is written as Ca 0.4 Sr 1.6 SiS 4 :Eu 2+ , implying a doping with Eu 2+ of 4%.

X-ray diffraction (XRD) measurements are collected with a Bruker D5000 θ-2θ diffractometer, using Cu-Kα radiation. Lattice constants are estimated using a spreadsheet solver and refined with the Rietveld structure refinement program FullProf [7]. The goodness of fit was evaluated graphically [8]. Neither the atomic positions nor the thermal expansion factor were further refined, as this typically requires higher quality XRD spectra. The initial parameters were based on the data available for the similar compounds Ca 2 GeS 4 [9] and Eu 2 SiS 4 [10].

Photoluminescence spectra are collected using a FS920 fluorescence spectrometer (Edinburgh Instruments). Variation of the measurement temperature (70 K -475 K) is obtained by an Optistat CF helium cryostat (Oxford Instruments).

Spectrally resolved SEM-CL (cathodoluminescence in a scanning electron microscope) measurements at Structure and luminescence of Eu-doped (Ca,Sr) 2 SiS 4 phosphors CL emission spectra were not corrected for the detector sensitivity. Details of the SEM-CL setup can be found here [11].

Results

Sr 2 SiS 4

Sr 2 SiS 4 was found to be monoclinic with space group P2 1 /m (11) by Dumail [12]. XRD-measurements (Figure 1) on our undoped samples are consistent with this space group, and the lattice parameters found are in good agreement with the ones recorded by Dumail [12] (See table 1 for an overview of lattice ). As the reflections observed are also compatible with space group P2 1 (4), a refinement of the XRD results using this model was performed, and an equally acceptable fit was found.

Thus, from the data collected, the possibility of space group P2 1 (4) cannot be excluded. Nevertheless, it seems unlikely that Sr 2 SiS 4 has a different space group than the group P2 1 /m (11) reported by Dumail [12].

Johrendt reports that there is no temperature dependent P2 1 /m (11) to P2 1 (4) transition in Eu 2 SiS 4 [13].

Given that Eu 2+ is similar in size to Sr 2+ , the same behaviour is probably the case for Sr 2 SiS 4 . Furthermore

Tampier states that the similar compound Sr 2 GeS 4 is also P2 1 /m (11) [14].

As the concentration of Eu 2+ in the samples is very small (typically 1% to 4% substitution) it is not expected to have a large impact on the crystallographic structure. To verify this, the lattice constants were determined for a sample of Sr 2 SiS 4 with 4% Eu 2+ . As expected, this substitution affected the lattice parameters only slightly (numerical values are also given in table 1). In the following discussion on the structure of (Ca,Sr) 2 SiS 4 we will therefore not mention the Eu 2+ -doping explicitly.

Avella [5] as well as Olivier-Fourcade [6] thoroughly, a low temperature measurement (figure 2b, 75K,) was carried out, yielding two emission peaks: a main emission peak at 556 nm, and an additional emission peak at 486 nm. These values were found by fitting a superposition of two Gaussian distributions to the measured emission curve for Sr 2 SiS 4 when drawn on an energy scale.

The occurrence of two emission peaks can be explained by pointing out that there are two distinct Sr 2+sites which can be occupied by Eu 2+ , as is the case for the similar compounds Eu 2 SiS 4 [13] and Sr 2 GeS 4 [9]. Although the coordination in both sites is rather similar, the orientation of the 5d orbitals of the substituting Eu 2+ -ion can be different, leading to two clearly different emission energies. The same effect was observed by Smet et al. [4] in Ca 2 SiS 4 :Eu 2+ . The intensity ratio between both emission bands can then be influenced by energy transfer from the center with higher emission energy to the lower one. The study of this process, which is affected by the dopant concentration, is the subject of future research. We determined the external quantum efficiency of Sr 2 SiS 4 at about 30%, This is slightly lower than the quantum efficiency of Ca 2 SiS 4 [4]. Although this is lower than other values reported for some other conversion phosphors in the yellow-to-red region of the spectrum [15], there is still significant room for improvement by fine-tuning the synthesis conditions.

3.2 (Ca,Sr) 2 SiS 4

Structure

Upon substitution of Sr 2+ by Ca 2+ , the samples remain monoclinic up to 40% Sr 2+ substituted by Ca 2+ , with reflections monotonically shifting towards larger angles (figure 3). This suggests the presence of a Structure and luminescence of Eu-doped (Ca,Sr) 2 SiS 4 phosphors For Ca 2+ -concentrations between 50% and 90%, additional reflections appear in the XRD pattern, e.g. 24-25°, 29-30°, 38-39°(three smaller boxes in figure 3), which are typical for an orthorhombic phase, isostructural with Ca 2 SiS 4 . This suggests that for Ca 2+ concentrations between 50% and 90%, the samples contain two distinct phases, one monoclinic and isostructural with Sr 2 SiS 4 (which is called the S-phase hereafter) and the other orthorhombic and isostructural with Ca 2 SiS 4 (the C-phase).

It is interesting to notice that for Ca 2+ -concentrations between 50% and 90%, the reflections typical for the monoclinic phase do not shift as outspokenly on adding more Ca 2+ as for Ca 2+ -concentrations up to 50%.

This means that for Ca 2+ -concentrations above 50%, the monoclinic lattice hardly shrinks upon adding more Ca 2+ , suggesting that once half of the Sr 2+ -ions is substituted by Ca 2+ -ions, it becomes difficult to force additional Ca 2+ into the monoclinic phase. Not coincidentally, at this Ca 2+ -concentration, the orthorhombic C-phase starts to appear, accommodating the surplus of Ca 2+ -ions that cannot be fitted into the monoclinic phase. Further research could show if the suggestion -supported by the PL characteristicsthat the Ca 2+ -ions only substitute substantially for one of the two Sr 2+ -sites, is confirmed.

Structure and luminescence of Eu-doped (Ca,Sr) 2 SiS 4 phosphors

The results of the Rietveld refinement for the lattice parameters of the (Ca,Sr) 2 SiS 4 -series are given in table 2. One clearly notices a shrinking c-axis (0% to 40% Ca 2+ in table 2) for the monoclinic S-phase, while the a-and b-axis are relatively unaffected. An explanation for the anisotropic shrinking could be found in the position of one in two Sr 2+ -atoms in Sr 2 SiS 4 , which form a plane with the c-axis perpendicular to it (similar to the case of Eu 2 SiS 4 , as explained by Hartenbach [10]). As a Ca 2+ substitutes for one of these Sr 2+ , the dimensions will shrink along the c-axis, but as the dimensions along the a-and b-axis are mainly determined by atoms other than the alkaline earths, these will not change much.

Structure and luminescence of Eu-doped (Ca,Sr) 2 SiS 4 phosphors Like most sulfide materials, the (Ca,Sr)-thiosilicate phosphors suffer from decomposition upon contact to ambient air due to their hygroscopic nature. After exposure to ambient air for several weeks, some degradation of the phosphor material can be observed. Although the stability is slightly less than that for the binary sulfides like CaS and SrS, the (Ca,Sr)-thiosilicate phosphors are much more stable than the similar thioaluminate compounds, like CaAl 2 S 4 or SrAl 2 S 4 [16,17].

PL -Emission Room Temperature

The body colour of the samples with 4% Eu 2+ ranges from grass green (0-10% Ca 2+ ) over whitish-yellow (20-30% Ca 2+ ) and orange-yellow (40-80% Ca 2+ ) to orange (90-100% Ca 2+ ). Small changes in Eu 2+concentration (1-4% Eu 2+ ) do not affect the body colour of the samples significantly. The body colour of a sample without any Eu 2+ is grayish-white and does not exhibit any luminescence, which confirms that only the Eu 2+ -ion is activating the luminescence. Structure and luminescence of Eu-doped (Ca,Sr) 2 SiS 4 phosphors

Figure 4 shows the emission spectra (excitation at 400 nm) for (Ca,Sr) 2 SiS 4 :Eu 2+ as a function of Ca 2+concentration. The main emission peak at 548 nm for pure Sr 2 SiS 4 shifts towards longer wavelengths as more Ca 2+ -ions are introduced into the initial mixture. This can be related to the structural properties of the S-phase with an increasing fraction of Ca 2+ : the shrinking crystal lattice causes a stronger crystal field and thus a longer emission wavelength. At around 50% Sr 2+ -substitution by Ca 2+ , the main emission peak shifts back towards the shorter wavelengths. This is explained by the contribution to the emission spectrum of the C-phase that emerges at Ca 2+ -concentrations of 50% or more. Depending on the Ca 2+concentration, the second emission peak of the S-phase at shorter wavelength (495 nm) and the second emission peak of the C-phase at longer wavelength (650 nm) are less or more noticeable when excited at 400nm. In particular, the sample with 70% Ca 2+ -concentration, is clearly a mixture of two phases (confirmed by the XRD results), as it exhibits both the 495 nm peak characteristic for a S-phase and the 650 nm peak characteristic for the C-phase. The series ends with the emission spectrum of pure Ca 2 SiS 4 .

The hypothesis of two phases (S-phase and C-phase), each causing two distinct emission peaks, will be amplified in both the section on SEM-CL results and the section on low temperature photoluminescence measurements. We will use the following abbreviations: the S1-peak (short wavelength) and S2-peak (longer wavelength) for the emission maxima of the S-phase; the C1-peak (short wavelength) and the C2peak (longer wavelength) for the emission maxima of the C-phase.

The percentage Eu 2+ that substitutes for (Ca,Sr) 2+ does not heavily influence the structure of the phosphor, but it does have a substantial impact on the photoluminescence of the material. A limited number of samples were prepared with different Eu 2+ -concentration. One finds that the Eu 2+ concentration does not affect the position of the emission peaks much, but that it does have an impact on the relative intensity of the peaks (figure 5). A deeper investigation of the (Sr,Eu) 2 SiS 4 -series is necessary to confirm if this behaviour is analogous to what was found for (Ca,Eu) 2 SiS 4 in [4].

PL -Excitation

All (Ca,Sr) 2 SiS 4 -compositions have a broad and relatively featureless excitation spectrum (figure 6) at room temperature. The main excitation band is situated from 350 to 475 nm, which means that all phosphors can be used in combination with both the standard near-UV-LEDs (390 to 400nm) and the standard blue LEDs (460 to 470nm) commercially available. Host excitation, which is situated at 4.9eV

for Ca 2 SiS 4 [18] is not very efficient.

Structure and luminescence of Eu-doped (Ca,Sr) 2 SiS 4 phosphors

SEM-CL results

To explain the luminescence behaviour of the (Ca,Sr distinct sets of spectra are obtained, one set with emission bands peaking at 485 and 585nm and one peaking at 560 and 660nm. Based on this luminescence signature and on the X-ray diffraction results, the former set can be related to the S-phase, with about 50% of the Sr 2+ -ions replaced by Ca 2+ . The latter one Structure and luminescence of Eu-doped (Ca,Sr) 2 SiS 4 phosphors can be related to the C-phase, with limited substitution of Ca 2+ by Sr 2+ . The brighter area in the centre of intensity map D could be related to a trace of CaS:Eu 2+ , or to a Ca 2 SiS 4 :Eu 2+ -like phase with higher europium concentration.

PL -Emission Low Temperature

At room temperature, the emission spectra of several (Ca,Sr) 2 SiS 4 -compositions show interesting substructure, but due to thermal broadening this cannot easily be resolved. Therefore, in this paragraph, the results of paragraph 3.2.2 (PL at room temperature) are completed with the results of low temperature measurements on the samples with 0%, 20%, 50%, 70%, 80% and 100% Sr 2+ substituted by Ca 2+ . Also, all PL emission spectra are interpreted in terms of the model developed in the paragraphs on structure and on SEM-CL results. An overview of the experimental results can be found in table 3. For the S-phase (0-40% Ca 2+ ) the emission spectrum can be modeled as a sum of two Gaussian peaks, the S1-peak and the S2-peak. The S2-peak shifts from 550 nm to about 590 nm as the Ca 2+ -concentration is Structure and luminescence of Eu-doped (Ca,Sr) 2 SiS 4 phosphors increased. The S1-peak is much weaker, and shifts from 480 nm to about 490 nm. For both peaks, the shift towards longer wavelengths is explained by the shrinking lattice, as this causes a stronger crystal field and thus longer wavelength emission. The strength of the emission shift is strongly different for both emission centers (S1 hardly shifts while S2 shows a 40nm shift for 50% substitution). This again points at a possibly preferential lattice site for the Ca 2+ ions upon substitution in the S-phase.

For the two-phase compositions (50-90% Ca 2+ ), one can try to fit the emission spectra to a sum of four Gaussian distributions. For the samples with little C-phase (50-60% Ca 2+ ), this is not feasible because the intensity of C-phase peaks is too low. The powders with more C-phase (70-90% Ca 2+ ) can be fitted very well, as is illustrated in figure 8 for Sr 0.4 Ca 1.6 SiS 4 .

Thermal Quenching

The use of phosphor materials in white-light emitting diodes is seriously hampered if thermal quenching occurs at relatively low temperatures. For high-power LEDs, the LED chip can reach temperatures of 450K [19]. Unless a remote phosphor approach is used, both the emission colour and the quantum efficiency should remain stable at this elevated temperature. favorable thermal quenching behaviour, with T 0.5 = 467K (i.e. the temperature at which the emission intensity has dropped to half its value at low temperature), which is in line with earlier reports.

Sr 2 SiS 4 :Eu 2+ has a much lower T 0.5 value of 380K, pointing at the excited Eu 2+ 5d state being closer to the bottom of the conduction band [20]. Partial substitution of Sr 2+ by Ca 2+ leads to an increase in T 0.5 value.

Additional detailed investigations are required to explain the thermal quenching behaviour, based on an accurate energy level scheme incorporating the (Ca,Sr) 2 SiS 4 band structure and the Eu 2+ 4f and 5d energy levels.

Conclusions

The (Ca,Sr)-thiosilicate phosphors have interesting features. First, as the excitation spectrum is very broad, they can readily be excited by commercially available LEDs, both near-UV and blue. Secondly, the emission peaks can be shifted by changing the Sr 2+ /Ca 2+ concentration. Thirdly, the sintering temperatures used in the preparation are relatively low as compared to e.g. nitride phosphors. The europium-doped (Ca,Sr)-thiosilicate phosphors have a relatively high emission efficiency, which could be further improved by tuning the synthesis conditions, such as the sintering temperature and duration, the dopant concentration and the (CaS,SrS):Si ratio. A drawback of the (Ca,Sr)-thiosilicates is their degradation upon contact with moisture. Preparing and processing in inert conditions and encapsulating the particles with an Al 2 O 3 -coating as in [21]

Confidential: not for

  distribution. Submitted to IOP Publishing for peer review 25 November 2009

Figure 1 :

 1 Figure 1: XRD pattern of monoclinic Sr 2 SiS 4 without Eu 2+ -doping.

Figure 2 :

 2 Figure 2: PL spectra of Sr 2 SiS 4: Eu 2+ (a) Emission spectra at room temperature, for different excitation wavelengths: low wavelength reveals extra peak. (b) Emission spectrum spectra for 4% Eu 2+ substitution at low temperature (70 K), excitation wavelength 375 nm.

Figure 3 :

 3 Figure 3: Pattern of (Ca,Sr) 2 SiS 4 as a function of Ca 2+ -concentration. The large box emphasizes the monoclinic signature of the S-phase. The three smaller boxes designate the orthorhombic C-phase.

Figure 4 :

 4 Figure 4: PL emission of (Ca,Sr) 2 SiS 4 :Eu 2+ as a function of Ca 2+ -concentration, upon excitation at 400 nm.

Figure 5 :

 5 Figure 5: Normalized emission spectrum of SrCaSiS 4 as a function of Eu 2+ -concentration (measured at 75 K).

Figure 6 :

 6 Figure 6: Excitation spectra for different compositions of (Ca,Sr) 2 SiS 4 :Eu 2+ monitored at the emission maxima. The curves are shifted vertically for clarity. (a) Ca 2 SiS 4 (b) Ca 1.6 Sr 0.4 SiS 4 (c) Ca 1.2 Sr 0.8 SiS 4 (d) Ca 0.8 Sr 1.2 SiS 4 (e) Ca 0.4 Sr 1.6 SiS 4 (f) Sr 2 SiS 4 .

) 2 SiS 4 :

 4 Eu 2+ powders as function of the composition and to relate this to the XRD results, SEM-CL measurements on Ca 1.5 Sr 0.5 SiS 4 :Eu 2+ powder were performed. CL maps from an area of about 22 µm by 35 µm were measured. The measurement temperature was lowered to 10K, which reduced the width of the emission bands and enabled to distinguish different emission features more easily. The CL spectrum averaged over the measured area clearly shows the contribution of at least 4 emission bands. Evaluation of the individual CL emission spectra clearly reveals strongly different emission spectra over the measured area. By carefully choosing the spectral windows for monochromatic maps (figure 7), A (460-525nm) and C (610-625nm) on the one hand, and B (530-555nm) and D (660-750nm) on the other hand show a striking similarity in intensity distribution, while both sets of maps are each other's inverse.

Figure 7 (Figure 7 :

 77 Figure 7: SEM-CL measurement at 10K of Ca 1.5 Sr 0.5 SiS 4 :Eu 2+ . (bottom) secondary electron image and CL emission spectrum integrated over the entire measurement area, (center) CL intensity maps in selected wavelength ranges relative to the total CL intensity. A: 460-525nm, B: 530-550nm, C: 610-625nm, D: 660-750nm, (top) local CL emission spectra from the areas indicated in the CL maps, along with the wavelength ranges used for the CL maps. (colour online)

Figure 8 : 2 (

 82 Figure 8: The emission spectrum (measured at 75 K) of Sr 0.4 Ca 1.6 SiS 4 :Eu 2+ consists of 4 separate peaks: 2 (solid lines) originating in the S-phase, and 2 (dashed lines) originating in the C-phase.

Figure 9 :CaSrSiS 4 :

 94 Figure 9: Normalized PL intensity as a function of temperature, with indication of T 0.5 for (a) Sr 2 SiS 4 :Eu 2+ (b) CaSrSiS 4 :Eu 2+ and (c) Ca 2 SiS 4 :Eu 2+ . The powders were excited by a 400nm-LED.

Table 1 : Lattice parameters for Sr 2 SiS 4 . The numbers between brackets represent errors on the last digit.

 1

		a (pm)	b (pm)	c (pm)	(°)
	This work (no Eu 2+ ) 655.0(1) 664.1(1) 824.8(1) 108.4(1)
	This work (4% Eu 2+ ) 654.8(1) 663.9(1) 824.6(1) 108.4(1)
	Dumail et al. [12]	655.6	664.0	824.8	108.6

  reported for Sr 2 SiS 4 an emission wavelength of 545 nm, both measured on samples with low Eu 2+ -concentration (<0.5%). The PL measurements carried out at room temperature on our Sr 2 SiS 4 samples show a similar value of 548 nm, irrespective of the excitation

Table 2 : Lattice parameters of (Ca,Sr) 2 SiS 4 in function of Ca 2+ -concentration. If a sample contains 2 phases, the refined lattice parameters of both phases are shown in the table.

 2

	% Ca 2+ Space Group a (pm) b (pm) c (pm)	(°)
	0	P2 1 /m (11)	655	664	825	108.4
	10	P2 1 /m (11)	654	664	819	108.3
	20	P2 1 /m (11)	652	663	814	108.2
	30	P2 1 /m (11)	652	663	809	108.1
	40	P2 1 /m (11)	650	662	803	108.0
	50	P2 1 /m (11)	650	663	802	107.8
		Pnma (62)	1357	824	622
	60	P2 1 /m (11)	649	663	799	107.9
		Pnma (62)	1356	824	621
	70	P2 1 /m (11)	649	663	798	107.8
		Pnma (62)	1356	823	621
	80	P2 1 /m (11)	649	662	798	107.9
		Pnma (62)	1356	823	621
	90	P2 1 /m (11)	648	661	796	107.8
		Pnma (62)	1352	821	620
	100	Pnma (62)	1352	819	621

Table 3 : Emission maxima as a function of Ca 2+ -concentration in (Ca,Sr) 2 SiS 4 :Eu 2+ . Main values are emission peaks of the samples at room temperature. Values between brackets are the emission peaks of the samples measured at low temperature (75 K). Values with an asterisk (*) have a larger error, because the associated peak in the emission is small and/or broad, and thus cannot be easily resolved.

 3

	% Ca 2+	S1	2+ 2 SiS 4 -like S2	C1	Ca 2 SiS 4 -like C2
	0	478 (486) 548 (556) -	-
	10	479	559	-	-
	20	482 (489) 570 (574) -	-
	30	484	581	-	-
	40	487	584	-	-
	50	489 (493) 584 (592) -	-
	60	489	585	-	-
	70	496* (496) 590 (590) 556 (557) 660* (658)
	80	489* (496) 591 (591) 560 (559) 664* (665)
	90	-	-	573	659*
	100	-		569 (565) 668 (665)

  The thermal quenching behaviour was determined from 70 to 470K, for a selected number of phosphor materials (i.e. Ca 2 SiS 4 :Eu 2+ , CaSrSiS 4 :Eu 2+ and Sr 2 SiS 4 :Eu 2+ ). Powder samples were illuminated from outside the cryostat by a single Structure and luminescence of Eu-doped (Ca,Sr) 2 SiS 4 phosphors 400nm-LED (operated at 20mA).Figure 9 shows the total emission intensity as function of temperature. Ca 2 SiS 4 :Eu 2+ clearly has the most

  could eliminate this drawback. Nevertheless, the thermal quenching temperature Structure and luminescence of Eu-doped (Ca,Sr) 2 SiS 4 phosphors T 0.5 decreases upon increasing substitution of Ca 2+ by Sr 2+ . Although Sr 2 SiS 4 :Eu 2+ appears not very favourable for application as colour conversion material in LEDs (due to the onset of the thermal quenching near room temperature), the intermediate compositions have a sufficiently high thermal quenching temperature, certainly if a remote phosphor approach is used.

Acknowledgements

PFS is a post-doctoral researcher for FWO-Vlaanderen.