Ion-induced electron emission by keV-range energy indium ions: Influence of material and geometry.

V.Hugonnaud, 1, 2, a) S.Mazouffre, 2, b) and D.Krejci 1, c) 1) Enpulsion, Viktor kaplan straße 2700, Wiener Neustadt, Austria 2) CNRS-ICARE, 3 Avenue de la Recherche Scientifique, 45100 Orléans, France (Dated: 13 October 2022)

This work provides measurements of the ion-induced electron emission (SEE) yield caused by keV-range energy indium ions (3 to 8 keV). In our experiment the ion sources, fuelled with indium, are based on the working principle of fieldemission electric propulsion (FEEP) technologies. The measured yields are given for different material properties (plain structure vs foam) and geometries (flat surface vs conical). Study outcomes show that electrons induced by keV-range indium ions can lead to overestimate the measured ion current by 30% to 200% depending on the collector materials and geometry. We spotlighted that molybdenum, tungsten and stainless steel 316L have relatively low yields compared to aluminium 2017A. Moreover, it is possible to lower down a specific material yield by two third when using a foam structure. Finally, it is shown that off-axis ions influence on a collector increase the resulting electron emission yield. In the case of an indium-FEEP thruster ion-induced electron emissions are a major perturbation when measuring ion current. Mainly due to the energy range of ions studied, yields retrieved are orders of magnitudes larger than what it is usually measured when studying the plume of an electrostatic electric propulsion system. It is therefore strongly recommended to use a so-called Faraday cup to study the ion beam from an indium-FEEP thruster as the probe close architecture allows to greatly mitigate these perturbations.

I. INTRODUCTION

Over the last decade the demand in electric propulsion (EP) systems has increased. Therefore, plume characterisation by mean of electrostatic probes is a subject of interest as it provides meaningful information on thruster performances and plume-spacecraft interaction. Electrostatic probes are often used to measure plasma properties within the plume created by an electric thruster [START_REF] Sheridan | How big is a small Langmuir probe[END_REF][START_REF] Demidov | Electric probes for plasmas, the link between theory and instrument[END_REF][START_REF] Mazouffre | Characteristics and Performances of a 100-W Hall Thruster for Microspacecraft[END_REF][START_REF] Keerl | Two-dimensional plasma plume density characterisation of the IFM Nano Thruster[END_REF] . Probe types are numerous [START_REF] Sheridan | How big is a small Langmuir probe[END_REF][START_REF] Demidov | Electric probes for plasmas, the link between theory and instrument[END_REF][START_REF] Marerese | Development of a single-orifice Retarding Potential Analyser for Hall Thruster Plume Characterization[END_REF][START_REF] Chen | Langmuir probe analysis for high density plasmas[END_REF][START_REF] Merlino | Understanding Langmuir probe current-voltage characteristics[END_REF][START_REF] Terhune | Ion and Droplet Mass Measurement of an Electrospray Emitter using ExB filter[END_REF][START_REF] Sheehan | Emissive probes[END_REF][START_REF] Teshigahara | Development of 2D Laser-Induced Fluorescence (LIF) System in High-Density Helicon Plasma[END_REF][START_REF] Rawat | Design and simulation of 10kW Faraday cup for ion beam current[END_REF][START_REF] Mazouffre | Evaluation of various probe designs for measuring the ion current density in a Hall thruster plume[END_REF][START_REF] Brown | Recommended Practice for Use of Faraday Probes in Electric Propulsion Testing[END_REF][START_REF] Krejci | Emission Characteristics of Passively Fed Electrospray Microthruster with Propellant Reservoirs[END_REF][START_REF] Vincent | A compact new incoherent Thompson scattering diagnostic for low-temperature plasma studies[END_REF] and enable acquisition of current density, particles energy and velocity. To enhance probe current collection efficiency, designs have evolved to sophisticated architecture using systems of electrode or close geometry structure [START_REF] Mazouffre | Evaluation of various probe designs for measuring the ion current density in a Hall thruster plume[END_REF][START_REF] Brown | Recommended Practice for Use of Faraday Probes in Electric Propulsion Testing[END_REF] . Despite many effects, current measurement accuracy still depends largely on the collector material, the probe geometry and the plasma properties. Ion-induced electron emission (SEE) is a fundamental surface interaction which can strongly influence the total ion current measured by the device. Lately, strong interest from satellite manufacturers in electric thrusters using molten metal indium as propellant spotlighted scarce data availability of electrons induced by keV-range energy indium ion bombardment. Consequently, we propose in this study to experimentally measure the yield of ion-induced electrons (γ SEE ) emitted by different materials when bombarded by energetic (3 keV to 8 keV) indium ions. γ SEE is given for aluminium 2017A, stainless steel 316L, tungsten and molybdenum for an ion beam incidence normal to the material surface. To this end, we updated the design of an electrostatic probe termed Faraday cup (FC) [START_REF] Rawat | Design and simulation of 10kW Faraday cup for ion beam current[END_REF][START_REF] Hofer | A Comparison of Nude and Collimated Faraday Probes for Use with Hall Thrusters[END_REF][START_REF] Hugonnaud | Faraday cup design for low power electric thrusters[END_REF] to properly measure the yield. A FC is a well known device which measures electron or ion cura) Author to whom correspondence should be addressed: valentin.hugonnaud@enpulsion.com b) stephane.mazouffre@cnrs-orleans.fr c) david.krejci@enpulsion.com rents when the collector is biased positively or negatively, respectively. γ SEE is also measured for different ion collector shapes modifying the ion beam incidence angle.

II. EXPERIMENTAL APPARATUS

A. Test bench, mechanical interface and instrument

Measurements were done in a cylindrical stainless-steel vacuum vessel of 0.91 m in diameter and 1.75 m in length located at the aerospace engineering's department laboratory of Wiener Neustadt University of Applied Sciences (FH Wiener Neustadt). During operation the pressure level typically was 2 × 10 -6 mbar. The chamber frame is connected to ground which is the voltage reference for this study.

An aluminium rotating arm allows automatic alignment of the probe with the thruster equatorial plane. The probe holder is mounted on a URS1000BCC motorized rotation stage from Newport. A Newport's SMC100 enables to control the motor from the atmospheric side. The thruster centreline is referred to as the 0 angular position. The pivot point of the rotating structure is aligned with the thruster exit plane. The system enables a scan from -90 to +90 in the horizontal plane with respect to the thruster. The distance R between the FC aperture and the thruster exit plane is 25.5 cm. The entire mechanical structured is grounded. Further description can be found here [START_REF] Hugonnaud | Faraday cup sizing for electric propulsion ion beam study: Case of a Field-Emission Electric Propulsion thruster[END_REF] .

A calibrated Keithley 2050 sourcemeter is used to measure the ion current. The device is operated in voltage source and measures current with 0.012% basic accuracy. This is the author's peer reviewed, accepted manuscript. However, the online version of record will be different from this version once it has been copyedited and typeset. The ENPULSION NANO is engineered and produced by the Austrian company Enpulsion GmbH [START_REF] Krejci | Recent flight data from IFM Nano Thrusters in a low earth orbit[END_REF][START_REF] Schoenherr | Development, Production, and Testing of the IFM Nano FEEP Thruster[END_REF] . The thruster working principle is based on FEEP technology [START_REF] Marcuccio | Experimental Performance of Field Emission Microthrusters[END_REF][START_REF] Genovese | Indium FEEP microthruster: Experimental characterization in the 1-100µN range[END_REF][START_REF] Tajmar | Experimental Performance of Field Emiision Microthrusters[END_REF] and has been studied over 20 years by FOTEC GmbH 4,24-27 . The thruster uses indium, a molten metal, as propellant. The latter is used to wet 28 a porous emitter crown constituted of 28 sharp needles, also called injectors. To provide enough potential difference to form Taylor cones at the tip of each injectors [START_REF] Taylor | Desintegration of water drops in an electric field[END_REF][START_REF] Mora | The current emitted by highly conducting Taylor cones[END_REF] and start the ionization process a counter electrode termed extractor (V ex ) is placed around the emitter. The resulting local field strength at the apex of the Taylor cone (∼ 10 9 V/m) has twofold purpose i) it leads to particle extraction and ionization and ii) it accelerates the ions.

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Results obtained in this work rely on measurements and analysis of the ion beam produced by two laboratory units (LU) version of this thruster. The lab units are not equipped with neutralisers considering the low current densities at stakes but also to ease current trace understanding. Emission current (I em ) and voltage (V em ) are controlled with a power processing unit (PPU) which allows accurate recordings with 1% reading accuracy. Due to experimental and time constraints the LUs had different injector distribution as showed in figure 1. This difference brings information on the influence of the firing conditions and ion beam velocity vector upon the ion-induced electron yield measured as described in section IV B. LU-A was used to asses the evolution of SEE yield with different foam material pore sizes, see section IV A. LU-B was used for the rest of the study. The thruster LU-A was fired at 1 mA while LU-B was fired at 1, 2, 3 and 4 mA. For each emission current different V em are applied to vary the ion energy. Ions in the beam produced by the ENPULSION NANO are singly-charged. Without energy losses the ion energy reads E i = eV em [START_REF] Mühlich | IFM Nano Thruster performance studied by experiments and numerical simulations[END_REF] .

III. METHOD A. Ion current and collection efficiency

The ion current density angular distribution is obtained from the ratio between the ion current, averaged over 10 consecutive measurements at each angular position, and the collection area. The ion current I int can be retrieved from the integration of the current measured on a plane which contains the thruster axis, e.g. following the angle θ from -π/2 to π/2 (sections IV D). We use the hypothesis of cylindrical symmetry around the thruster axis [START_REF] Brown | Recommended Practice for Use of Faraday Probes in Electric Propulsion Testing[END_REF][START_REF] Hugonnaud | Faraday cup design for low power electric thrusters[END_REF][START_REF] Hugonnaud | Optimization of a Faraday cup collimator for electric propulsion device beam study: Case of a Hall thruster[END_REF] . However, this assumption holds only if the amount of firing needles is sufficient and well distributed (e.g. LU-B). Therefore, I i int reads:

I i int = π R 2 π 2 -π 2 j i (θ ) | sin(θ ) | dθ , (1)
with R being the distance from the probe inlet to the ion source emission plane, j i the current density (A/m 2 ) and θ the angle between the probe centre axis and the ion source firing axis. Consequently, the ion collection efficiency reads [START_REF] Hugonnaud | Faraday cup design for low power electric thrusters[END_REF][START_REF] Hugonnaud | Faraday cup sizing for electric propulsion ion beam study: Case of a Field-Emission Electric Propulsion thruster[END_REF] :

η p = I i int I em . ( 2
)
Where I em is the ion current emitted by the ion source and controlled by a power processing unit. Uncertainties are computed following equation 3 where σ is the standard deviation and u the corresponding uncertainty.

u = σ √ n , (3)
with n the number of samples. The sum of each uncertainties is computed following a linear error propagation with 95% confidence (k=2) level which reads:

u = n ∑ i=1 u 2 i , u 95% = 2u. (4)
Previous studies [START_REF] Hugonnaud | Faraday cup design for low power electric thrusters[END_REF][START_REF] Hugonnaud | Faraday cup sizing for electric propulsion ion beam study: Case of a Field-Emission Electric Propulsion thruster[END_REF] showed that with the right architecture a FC can accurately measure the ion current giving a value close to the one emitted by the thruster with ±3% with 95% confidence. During ion current measurement I em and V em are controlled and recorded by the thruster laboratory unit onboard telemetry. In the case of ion current density angular distribution set, thruster firing conditions are constantly monitored. Then, an average of 800 data samples acquired over 25 minutes with accuracy of 0.01%, negligible compared to the telemetry reading accuracy (∼1%), gives the mean firing parameters.

B. Ion-induced electron emission yield (γ SEE )

Ion-induced electron emission (SEE) is known to be the cause of an artificial current rise when measuring ion current. This is the author's peer reviewed, accepted manuscript. However, the online version of record will be different from this version once it has been copyedited and typeset. Therefore, it is fundamental to either properly recollect ioninduced electrons (SE) or to determine the rate of emitted SE to correct output consequently. The yield of SEE, γ SEE , depends on projectile ion energy, incidence angle and collector material properties. Commonly, γ SEE is determined by applying a voltage sweep on an ion collector V coll placed downstream an electrode V ele with a fixed negative voltage [START_REF] Corbella | Revising secondary electron yields of ion-sputerred metal oxides[END_REF][START_REF] Habl | Secondary electron emission due to multi-species iodine ion bombardment of different target materials[END_REF] . When V coll < V ele , the current measured on the collector is I c = I SEE , with I SEE the ion current including the contribution from SEE. Then, when V coll > V ele , the current measured on the collector corresponds to the real ion current, I c = I i . Therefore, γ SEE reads [START_REF] Eder | Electron emission from clean gold bombarded by slow Au q+ (q=1-3) ions[END_REF] :

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γ SEE = I SEE -I i I i . (5)
However, if one uses the above mentioned probe architecture, and make the hypothesis that the whole ion flux is only collected by the probe rear part, two problems arise when V coll > V ele : 1) the collector will start to collect SE emitted by the electrode placed upstream, 2) a part of the ion current will be collected by the electrode and not the collector. Both case will induce a wrong estimation of SE yield. In our study we modified the architecture of a FC to get around this artificial yield rise, see figure 2. There, V ele either refers to V rep or V cup depending on the probe configuration, A vs B. When configuration A is used the ion collector is the "cup + rear disk". For configuration B the ion collector is only the "rear disk". In both cases the electrode placed upstream right behind the probe aperture, termed repeller (V rep ), [START_REF] Hugonnaud | Faraday cup design for low power electric thrusters[END_REF][START_REF] Hugonnaud | Optimization of a Faraday cup collimator for electric propulsion device beam study: Case of a Hall thruster[END_REF] is biased negatively to screen thermal electrons and ion-induced electrons due to plasma-probe interaction. Two IV curves, for an ion collector made of aluminium, are displayed in figure 3. They illustrate the problem stated above. Ion current is plotted as positive values. The ion current without SE contribution should correspond to the current measured in zone 2 [START_REF] Hugonnaud | Faraday cup sizing for electric propulsion ion beam study: Case of a Field-Emission Electric Propulsion thruster[END_REF] . However, in this zone the two configurations A and B give different values with I B < I A . Figure 4 provides explanations for current measured in zone 1 and 2 for configuration A and B. In zone 1, SE are not recollected with both configurations, since V coll is lower than the other one, hence an overestimated current. The cur-FIG. 3: I-V curves measured with different probe configurations. Probe A uses the collector and the cup as one electrode while with configuration B these two elements are electrically uncoupled. The two coloured dashed lines corresponds to V rep (configuration A) or V cup (configuration B) at -60 V. The black dashed line ease the visualisation for the transition from negative to positive voltages.

rent measured by B is increased due to SE from the ion collector located at the rear of the probe. The current measured by A is solely perturbs by ion-induced electrons that escape the cup which is characterized by the yield γ L SEE [START_REF] Hugonnaud | Faraday cup sizing for electric propulsion ion beam study: Case of a Field-Emission Electric Propulsion thruster[END_REF] . Therefore, in both cases the current is artificially increased but not with the same factor. Configuration A manages to recollect an important part of ion-induced electrons in zone 1. In zone 2, all ion-induced electrons are recollected as V coll >V ele , V ele being V rep and V cup with configuration A and B, respectively. However, in B an additional parameter contributes to the current measured. The source cannot be properly characterized with the current experiment but is expected to be either due to 1) ions being now collected by the cup walls instead if the collector due to the electrode potential change or 2) ion induced electrons from the cup walls being now attracted by the collector due to the potential change. Figure 5 and6 show the current change with the voltage inside probes A and B with a 60 conical collector made of aluminium 2017A. In configuration A (figure 5), the repeller is biased to -60 V and SE are recollected when V coll overtakes V rep . The shift observed is caused by a potential dip at the repeller aperture centre. While a current drop is observed on the collector (top), no specific change is seen on the repeller (bottom) at the same voltage. Therefore, it indicates that in zone 1 SE from the cup are not collected by the repeller, electrons leave the probe instead. Moreover the ion current measured in zone 2 by the repeller is 100 times lower than the collector current and therefore, can be neglected. A detailed description of ion current behaviour with A can be found elsewhere [START_REF] Hugonnaud | Faraday cup design for low power electric thrusters[END_REF][START_REF] Hugonnaud | Faraday cup sizing for electric propulsion ion beam study: Case of a Field-Emission Electric Propulsion thruster[END_REF] , molybdenum (c) and stainless steel 316L (d) collector. These materials were chosen as they are widely used among the EP community as collector for electrostatic probes. Data is displayed for different ion current emission ranging from 1 to 4 mA. The extractor voltage is also displayed to assure similar conditions between measurements. One should note that yields are not given for every emission current at a given ion energy. This is due to the limit of the on-board power processing unit used for this experiment (i.e., total discharge, input power). It has been proved [START_REF] Eder | Electron emission from clean gold bombarded by slow Au q+ (q=1-3) ions[END_REF][START_REF] Baragiola | Critical potentials in secondary electorn emission from Iron, Nickel and Molybdenum[END_REF][START_REF] Alonso | Kinetic electron emission from solids induced by slow heavy ions[END_REF] that SEE yield can result from a combination of potential (PE) and kinetic (KE) electron emission. The two mechanisms rely on different energy transfer principles. PE emission results from electron excitation due to conversion of internal energy while KE emission occurs in the presence of kinetic energy transfer during collisions. The main difference is that PE is independent of the incoming ion energy and is driven by the collector material work function (φ ) while KE strongly depends on the ion energy. In this study E i φ , as φ is of the order of a few eV [START_REF] Eder | Electron emission from clean gold bombarded by slow Au q+ (q=1-3) ions[END_REF] for studied materials. Therefore, PE mechanism can be neglected and we consider KE to be the only phenomenon responsible for electron emission. Figure 7 shows that SEE yield increases monotonically for all material. However, while the yield is of the same order for molybdenum, tungsten and steel 316L, γ SEE obtained with aluminium 2017A is 3.5 to 4.5 larger. Moreover, between the lowest and highest mean ion energy the growth rate differs for the four materials. Indeed, the rise is around 57% for molybdenum, 67% for tungsten, 72% for aluminium 2017A and approaches 95% for steel 316L. We note a deviation of the yield at 2mA for the aluminium collector even after several additional measurements. This behaviour is not yet explained. However, section IV B gives information on how the firing conditions (needles distribution, incidence angle of the ion influence) can easily modify the yield measured. Besides, some collector materials are more influenced than others.

It is interesting to compare how the evolution of γ SEE for indium ions to yields induced by primary electron influences available in the literature. Both depend on the material surface quality (roughness, chemical composition) and the angle of incident [START_REF] Shih | Secondary electron emission studies[END_REF] (see section IV C). However, we note that the secondary electron yield induced by ions increases with the ion energy whereas for primary electrons the yield increases until it reaches a maximum before decreasing [START_REF] Shih | Secondary electron emission studies[END_REF][START_REF] Baglin | The secondary electron yield of technical materials and its variation with surface treatments[END_REF] . Shih et al. [START_REF] Shih | Secondary electron emission studies[END_REF] showed that secondary electron emitted by a surface greatly depends on the penetration depth (in microns) of the incoming primary particle. If P (the penetration-depth) is less than E (the escape-depth) the number of secondary electrons generated and then emitted increase with the primary energy. However, when P becomes larger than E, the escape rate of internal secondary electrons decreases quicker than the generation rate of the latter. Therefore, the yield decreases for higher primary energies. The maximum is reached when P is near E, which gives a bell-shape profile to the secondary primary electron curve. It is understood [START_REF] Young | Penetration of Electrons and Ions in Aluminum[END_REF][START_REF] Li | Recent advances in focused ion beam nanofabrication for nanostructures and devices: fundamentals and applications[END_REF] that electrons have a penetration depth around 10 times larger than ions for a given energy. Consequently, the threshold value to see the decline of the yield with an increase of the energy for ions is at higher level than for primary electrons, hence not visible at the current energy studied.

B. Solid vs foam metal

Figure 8 shows γ SEE obtained with different ion collector aluminium properties and geometries. Figure 8a presents the evolution of the yield for aluminium foams AlSi7Mg characterized by different pore diameters. F1 refers to pore size ranging from 0.2 to 0.35 mm and F6 has pore sizes from 0.63 to 3 mm. Note that the porosity remains 55%-65% for all foams. Here, the thruster is LU-A and fires at 1 mA. The yield increases monotonically for all foams. However, the yield is smaller for F6. In average γ SEE drops by 26%±8.4% between F1 and F6. Nevertheless, the yield variation between low and high mean ion energy approaches 40%±3% for both collectors. The observed yield reduction can be caused by different SE emission sites. When increasing the pore size, SE originates from deeper locations inside the collector. Hence the probability to recollect SE before they escape the foam structure increases. Figure 8b shows the evolution of the SEE yield for a flat bulk aluminium 2017A collector (Al) and the same foam aluminium collector used in figure 8a. Here, ions are ejected by LU-B. First, γ SEE is near 60%±8.5% lower for foam. Second, the difference between both yields slightly increases with the ion energy indicating that a foam structure is less influenced by the increase of energy.

It is observed that γ SEE for F6 measured with LU-A and LU-B is not the same. Figure 9 shows the variation of γ SEE for molybdenum (bottom) and F6 (top) obtained with different ion sources. Note that it was not possible to reproduce the exact same operation points as LU-A needed more extraction voltage to provide similar emission parameters. The PPU internal limits prevented to reach emission voltage below 5 kV for LU-A and above 6 kV for LU-B at 1 mA. γ SEE obtained for F6 with LU-A increased by 84% and 98% for 5 kV and 6 kV, respectively. On the contrary, there are no major change for the molybdenum collector as the variations are within the uncertainties of the measurement method. Note that for both materials the rate of change between the minimum and maximum yield is identical +40% and +58% for LU-A and LU-B, respectively. The increase of the measured yield for F6 could be caused by higher dependency to the ion incidence angle than molybdenum. Indeed, due to the needle distribution and the large extractor voltage used with LU-A the ions going through the probe might not hit the collector with a velocity vector identical to those from LU-B. Owing to its foam structure, the probability to have ions hitting its surface with an incidence angle different than 0 increases. The next section brings insights into the effect of the incidence angle upon This is the author's peer reviewed, accepted manuscript. However, the online version of record will be different from this version once it has been copyedited and typeset. Once more the yield increases monotonically with the bombarding mean ion energy. Here, we assume that ions entering the probe have a velocity vector purely perpendicular to the probe inlet plane. Therefore, when the angle which defines the conical shape of the collector decreases, the incidence angle (θ ) increases. In this case, ions hit Alc with θ = 37.5 and Alc2 with θ = 60. We observe that increasing the incidence angle increases the yield of SEE. For higher θ it is easier for an ion to rip off electrons from the surface. In average, between Al (flat) and Alc the yield increases by 19.5%±3.4%. From Alc to Alc2 the yield increases by 14.3%±3% at the highest mean ion energy.

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D. Ion-induced electron mitigation.

We have shown in the aforementioned sections that the ion current from an indium-based ion source acquired with an electrostatic probe depends significantly on the ion collector material properties and geometry. Nevertheless, all these perturbations can be counteract with the right probe architecture and voltage configuration, as showed in figure 11. Figure 11a shows the ion current density distribution acquired with a flat aluminium collector disturbed or not by ion-induced electron emission. The right plot of figure 11a displays the corresponding probe ion collection efficiency using equation 1 and 2. When the ion-induced electron effects are completely suppressed, the probe efficiency is the same no matter the thruster firing condition. Similarly, figure 11b shows that without SE This is the author's peer reviewed, accepted manuscript. However, the online version of record will be different from this version once it has been copyedited and typeset. effects the ion current density distribution is independent of the collector material. The plots show clearly that the ion current density distribution acquired from a FEEP thruster strongly depends on ion-induced electron effects. Consequently, to retrieve ions information from the thruster plume a Faraday cup is warmly recommended as its architecture makes ion-induced electron emission perturbations simpler to mitigate, hence improving measurement reliability.

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V. CONCLUSION

In this study, we have measured the ion-induced electron yield γ SEE of different collector materials bombarded by keVrange energy indium ions. To enable such measurements a new probe design, based on a Faraday cup architecture, was developed. The collected ion flux is defined by the aperture of the probe casing front part. A negative voltage is applied to a repeller placed behind the casing aperture to prevent perturbations induced by plasma-probe interaction to reach the ion This is the author's peer reviewed, accepted manuscript. However, the online version of record will be different from this version once it has been copyedited and typeset. collector. The latter has a cup-like shape (FC A). It can be divided into a simple cylinder and a disk (FC B) without modifying the experiment ambient conditions. From the currents acquired with FC A and FC B, with specific voltage variations, it was possible to estimate the SEE yield. It was shown that the yield varies linearly with the beam energy no matter the material. In the case of solid and flat materials such as tungsten and molybdenum, the SEE yield contributes to a variation from 10 to 30% of the measured current for low and high energy respectively. For solid and flat steel 316L collector, the yield does not exceed 10% for ion energy up to 5 keV and becomes negligible below 3 keV. For more energetic ions the SEE yield can contribute to rise the current by 30%. On the contrary, γ SEE for flat aluminium 2017A collectors is significantly larger and can increase the measured current by 40 to 150%. In a second step, the structure of the material was studied. Aluminium foams with different pore sizes and constant porosity (55-60%) were used. We observed a diminution of the yield between the solid and the foam version. We also observed that changing the incidence angle between the collector and the ion trajectory affects the SEE yields. Nonetheless, the variation is less with low SEE yield materials like molybdenum. The lowest yield recorded for an aluminium 2017A collector was at 0 and the most important was for the largest incidence angle. Overall, we demonstrated that ioninduced electron is a major perturbation for ion current measurements with an electrostatic probe in the plume of a FEEP indium-based thruster. The yield depends largely on material and shape disturbances can be properly mitigated by the use a probe architecture similar to a Faraday cup. The probe design allows to greatly mitigate perturbations induced by SEE.

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FIG. 1 :

 1 FIG. 1: Emitter distribution for both laboratory units, LU-A (left) and LU-B (right), during measurements.

FIG. 2 :

 2 FIG. 2: Schematic of probe configuration used to measureγ SEE

FIG. 4 :FIG. 5 :FIG. 6 :

 456 FIG. 4: Explanation of the measurement method to obtainγ SEE

Figure 7 displays

 7 Figure7displays SEE yield measured following the method explained in section III B. The yields are plotted as a function of the mean indium ion energy for flat aluminium 2017A (a), tungsten (b), molybdenum (c) and stainless steel 316L (d) collector. These materials were chosen as they are widely used among the EP community as collector for electrostatic probes. Data is displayed for different ion current emission ranging from 1 to 4 mA. The extractor voltage is also displayed to assure similar conditions between measurements. One should note that yields are not given for every emission current at a given ion energy. This is due to the limit of the on-board power processing unit used for this experiment (i.e., total discharge, input power). It has been proved[START_REF] Eder | Electron emission from clean gold bombarded by slow Au q+ (q=1-3) ions[END_REF][START_REF] Baragiola | Critical potentials in secondary electorn emission from Iron, Nickel and Molybdenum[END_REF][START_REF] Alonso | Kinetic electron emission from solids induced by slow heavy ions[END_REF] that SEE yield can result from a combination of potential (PE) and kinetic (KE) electron emission. The two mechanisms rely on different energy transfer principles. PE emission results from electron excitation due to conversion of internal energy while KE emission occurs in the presence of kinetic energy transfer during collisions. The main difference is that PE is independent of the incoming ion energy and is driven by the collector material work function (φ ) while KE strongly depends on the ion energy. In this study E i φ , as φ is of the order of a few eV[START_REF] Eder | Electron emission from clean gold bombarded by slow Au q+ (q=1-3) ions[END_REF] for studied materials. Therefore, PE mechanism can be neglected and we consider KE to be the only phenomenon responsible for electron emission. Figure7shows that SEE yield increases monotonically for all material. However, while the yield is of the same order for molybdenum, tungsten and steel 316L, γ SEE obtained with aluminium 2017A is 3.5 to 4.5 larger. Moreover, between the lowest and highest mean ion energy the growth rate differs for the four materials. Indeed, the rise is around 57% for molybdenum, 67% for tungsten, 72% for aluminium 2017A and approaches 95% for steel 316L. We note a deviation of the yield at 2mA for the aluminium collector even after several additional measurements. This behaviour is not yet explained. However, section IV B gives information on how the firing conditions (needles distribution, incidence angle of the ion influence) can easily modify the yield measured. Besides, some collector materials are more influenced than others.It is interesting to compare how the evolution of γ SEE for indium ions to yields induced by primary electron influences available in the literature. Both depend on the material surface quality (roughness, chemical composition) and the angle of incident[START_REF] Shih | Secondary electron emission studies[END_REF] (see section IV C). However, we note that the secondary electron yield induced by ions increases with the ion energy whereas for primary electrons the yield increases until it reaches a maximum before decreasing[START_REF] Shih | Secondary electron emission studies[END_REF][START_REF] Baglin | The secondary electron yield of technical materials and its variation with surface treatments[END_REF] . Shih et al.[START_REF] Shih | Secondary electron emission studies[END_REF] showed that secondary electron emitted by a surface greatly depends on the penetration depth (in microns) of the incoming primary particle. If P (the penetration-depth) is less than E (the escape-depth) the number of secondary electrons generated and then emitted increase with the primary energy. However, when P becomes larger than E, the escape rate of internal secondary electrons decreases quicker than the generation

FIG. 7 :

 7 Figure8shows γ SEE obtained with different ion collector aluminium properties and geometries. Figure8apresents the evolution of the yield for aluminium foams AlSi7Mg characterized by different pore diameters. F1 refers to pore size ranging from 0.2 to 0.35 mm and F6 has pore sizes from 0.63 to 3 mm. Note that the porosity remains 55%-65% for all foams. Here, the thruster is LU-A and fires at 1 mA. The yield increases monotonically for all foams. However, the yield is smaller for F6. In average γ SEE drops by 26%±8.4% between F1 and F6. Nevertheless, the yield variation between low and high mean ion energy approaches 40%±3% for both collectors. The observed yield reduction can be caused by different SE emission sites. When increasing the pore size, SE originates from deeper locations inside the collector. Hence the probability to recollect SE before they escape the foam structure increases. Figure8bshows the evolution of the SEE yield for a flat bulk aluminium 2017A collector (Al) and the same foam aluminium collector used in figure8a. Here, ions are ejected by LU-B. First, γ SEE is near 60%±8.5% lower for foam. Second, the difference between both yields slightly increases with the ion energy indicating that a foam structure is less influenced by the increase of energy.It is observed that γ SEE for F6 measured with LU-A and LU-B is not the same. Figure9shows the variation of γ SEE for molybdenum (bottom) and F6 (top) obtained with different ion sources. Note that it was not possible to reproduce the exact same operation points as LU-A needed more extraction voltage to provide similar emission parameters. The PPU internal limits prevented to reach emission voltage below 5 kV for LU-A and above 6 kV for LU-B at 1 mA. γ SEE obtained for F6 with LU-A increased by 84% and 98% for 5 kV and 6 kV, respectively. On the contrary, there are no major change for the molybdenum collector as the variations are within the uncertainties of the measurement method. Note that for both materials the rate of change between the minimum and maximum yield is identical +40% and +58% for LU-A and LU-B, respectively. The increase of the measured yield for F6 could be caused by higher dependency to the ion incidence angle than molybdenum. Indeed, due to the needle distribution and the large extractor voltage used with LU-A the ions going through the probe might not hit the collector with a velocity vector identical to those from LU-B. Owing to its foam structure, the probability to have ions hitting its surface with an incidence angle different than 0 increases. The next section brings insights into the effect of the incidence angle upon This is the author's peer reviewed, accepted manuscript. However, the online version of record will be different from this version once it has been copyedited and typeset.PLEASE CITE THIS ARTICLE AS DOI: 10.1063/5.0102206

  FIG. 8: Influence on computed γ SEE of the ion collector material property. (a) γ SEE for different aluminium AlSi6Mg foams. (b) γ SEE for different aluminium collector types.

FIG. 10 :

 10 FIG. 10: Ion-induced electron yield obtained with the ion source LU-B for different collector geometries. LU-B emission current is set to 2 mA.

  FIG. 11: Effects of ion-induced electron emission during ion current angular density distribution measurements for different ion source operation points (a) and different collector material (b).

TABLE I :

 I Details summary of each materials assessed as ion collector.This is the author's peer reviewed, accepted manuscript. However, the online version of record will be different from this version once it has been copyedited and typeset.

	Material	Geometry	Properties	Name
	Aluminum 2017A conical Plain, α 1 = 105 Alc
	Aluminum 2017A conical Plain, α 2 = 60 Alc2
	Aluminum 2017A	flat	Plain	Al
	Aluminum 2017A	flat	Foam #1 36	F1
	Aluminum 2017A	flat	Foam #3 36	F3
	Aluminum 2017A	flat	Foam #6 36	F6
	Molybdenum	flat	Plain	Mo
	Tungsten	flat	Plain	W
	Stainless steel 316L	flat	Plain	S

PLEASE CITE THIS ARTICLE AS DOI: 10.1063/5.0102206 IV. RESULTS AND DISCUSSION A. γ SEE for different materials

ACKNOWLEDGMENTS

V.Hugonnaud benefits from an ENPULSION PhD grant. Authors would like to thank Dr. Pr. Aumayr for his support on theoretical subject regarding plasma diagnostics design and behaviour. Additionally, we would like to thank C.Scharlemann, head of the FH Wiener Neustadt Aerospace department, for sharing their facilities.


U-Chan Chung

Catherine Elissalde
email: elissald@icmcb-bordeaux.cnrs.fr

Fre´de´ric Mompiou

Je´roˆme Majimel

Sonia Gomez

Claude Estourne

Sylvain Marinel

ÃÃ Andreas Klein

Fran @ Ois Weill

Dominique Michau

Ste´phane Mornet

Mario Maglione

Interface Investigation in Nanostructured BaTiO 3 /Silica Composite Ceramics

Silica-coated ferroelectric particles are promising building blocks for functional bulk composites such as dielectric resonators, supercapacitors, or multiferroics. The ferroelectric/silica interface was fully investigated by means of high-temperature in situ X-ray diffraction, high-resolution scanning electron microscopy, and X-ray photoelectron spectroscopy. Mechanisms occurring at the interface were visualized using high-resolution postmortem and in situ transmission electron microscopy performed at different temperatures. On the light of this interface investigation, we have used advanced sintering processes such as spark plasma sintering and microwave sintering to obtain nanostructured composite ceramics and to evaluate their dielectric properties.

I. Introduction

T HE coating of oxide or metallic particles by a continuous silica shell (core-shell architecture) has led to significant advances in functional materials, thus increasing drastically the range of applications (photonic crystals, luminescent semiconductors, biomedical magnetic systems, etc.). [START_REF] Caruso | Nanoengineering of Particle Surfaces[END_REF][START_REF] Davies | Engineered Particle Surfaces[END_REF] As high-temperature processes are not required to obtain the macroscopic composite, the initial nanoscale design is preserved. Property control is based on the periodicity and tunability of interparticle spacing (particle size, coating thickness) that drives the coupling between neighboring particles. [START_REF] De | Chemical Strategies to Design Textured Materials: From Microporous and Mesoporous Oxides to Nanonetworks and Hierarchical Structures[END_REF][START_REF] Garcia Santamaria | Synthetic Opals Based on Silica-Coated Gold Nanoparticles[END_REF][START_REF] Grasset | Synthesis and Magnetic Characterization of Zinc Ferrite Nanoparticles with Different Environments: Powder, Colloidal Solution, and Zinc Ferrite-Silica Core-Shell Nanoparticles[END_REF][START_REF] Vestal | Synthesis and Magnetic Characterization of Mn and Co Spinel Ferrite-Silica Nanoparticles with Tunable Magnetic Core[END_REF][START_REF] Liz-Marza´n | The Assembly of Coated Nanocrystals[END_REF] However, in high-density materials such as nanostructured electroceramics, preserving the nanoscale architecture is yet to be achieved, mainly because of the high-temperature sintering step. Core-shell building blocks were already used to favor solid-state reaction at the nanoscale, in order to (i) tune dielectric properties, (ii) delay or enhance densification, (iii) avoid grain growth during sintering, and (iv) improve the mechanical properties of sintered nanoparticles. [START_REF] Buscaglia | Coating of BaCO 3 Crystals with TiO 2 : Versatile Approach to the Synthesis of BaTiO 3 Tetragonal Nanoparticles[END_REF][START_REF] Mornet | Controlled Growth of Silica Shell on Ba 0.6 Sr 0.4 TiO 3 Nanoparticles Used as Precursors of Ferroelectric Composites[END_REF][START_REF] Lee | Coating BaTiO 3 Nanolayers on Spherical Ni Powders for Multilayer Ceramic Capacitors[END_REF][START_REF] Sacks | Fabrication of Mulitte-Matrix Composites by Transient Viscous Sintering of Composite Powders[END_REF][START_REF] Giesche | Preparation, Characterization, and Sinterability of Well-Defined Silica/Yttria Powders[END_REF][START_REF] Tian | Core-Shell Structure of Nanoscaled Ba 0.5 Sr 0.5 TiO 3 Self-Wrapped by MgO Derived from a Direct Solution Synthesis at Room Temperature[END_REF][START_REF] Hakim | Nanoparticle Coating for Advanced Optical, Mechanical and Rheological Properties[END_REF] In these dense 3D bulk materials, macroscopic properties are intimately resulting from the thermal treatment. Studies devoted to BaTiO 3 (BT)silica (SiO 2 )-based composites provide an illustration of the drastic role of temperature and of the initial composite architecture. [START_REF] Chen | Structure, Sintering Behavior and Dielectric Properties of Silica-Coated BaTiO 3[END_REF][START_REF] Park | Nano Size BaTiO 3 Powder Coated with Silica[END_REF][START_REF] Abicht | Preparation, Characterization and Sintering Behavior of Barium Titanate Powders Coated with Ba-, Ca-, Si-and Ti-Containing Components[END_REF][START_REF] Vo¨lzke | The Influence of Different Additives and The mode of their Addition on the Sintering Behaviour and the Properties of Semiconducting Barium Titanate Ceramics[END_REF] Other critical issues are the specific role and composition of the core-shell interface.

Fresnoite, Ba 2 TiSi 2 O 8 , clearly identified as a secondary phase growing inbetween the two components, is described in the literature as a ferroelectric material with dielectric, pyroelectric, and piezoelectric properties. [START_REF] Foster | Fresnoite: A New Ferroelectric Mineral[END_REF][START_REF] Halliyal | Unusual Pyroelectric and Piezoelectric Properties of Fresnoite (Ba 2 TiSi 2 O 8 ) Single Crystal and Polar Glass-Ceramics[END_REF] None of the previous reports referring to BT/SiO 2 ceramics were focused on a specific study of the interface; silica is mainly used as an additive for lowering the sintering temperature. [START_REF] Felgner | Investigations on the Liquid Phase in Barium Titanate Ceramics with Silica Additives[END_REF][START_REF] Saldan˜a | Preparation of BaTiO 3 Single Crystals Using the Modified SiO 2 -Exaggerated Grain Growth Method[END_REF] The full investigation of interfaces in such composites strongly depends on the complex control of interfaces, including the design of the starting composite. There is an obvious interest in understanding, and thus monitoring, such an intrinsically functional interphase formation, as the final properties and the possibility of additional functionalities will strongly depend on its nature (i.e., amorphous, partially, or completely transformed into crystalline fresnoite). Attaining such a level of interface control is of great interest for property tunability in the large area of functional electroceramics (dielectric resonators, tunable capacitors, antennas, supercapacitors, etc.).

Here, we propose a deep investigation of the BT/SiO 2 interface, down to the atomic resolution, using techniques such as in situ transmission electron microscopy (TEM), postmortem high-resolution transmission electron microscopy (HRTEM), and X-ray photoelectron spectroscopy (XPS). Core-shell particle design, BT@SiO 2 , was chosen as initial building block of the ceramic composites because the full coverage of each particle is definitely a prerequisite for the control of the chemical and structural mismatch between the core and shell, at the grain scale. Such a coating quality is also mandatory to uniformly transfer the property modification from the grain scale to the 3D dense ceramic. As previously reported, [START_REF] Mornet | Controlled Growth of Silica Shell on Ba 0.6 Sr 0.4 TiO 3 Nanoparticles Used as Precursors of Ferroelectric Composites[END_REF] the soft chemistry route we used for BT@SiO 2 enables a uniform, continuous, and homogeneous nanoscale coating of each individual ferroelectric particle. An additional advantage of ferroelectric@silica architecture is that silica enables additional functionality of the ferroelectric component, by ways of surface modification: this way, either the coexistence of ferroelectricity and ferromagnetism or low loss supercapacitor were obtained. [START_REF] Mornet | Ferroelectric-Based Nanocomposites: Towards Multifunctional Materials[END_REF][START_REF] Chung | Controlling Internal Barrier in Low Loss BaTiO 3 Supercapacitors[END_REF] II. Experimental Procedure (1) Synthesis and Standard Sintering of BT@SiO 2 Particles Barium titanate particles of mean diameter 300 and 500 nm (BT500) were purchased from Sakai Chemical Co. (Fukushima, Japan). All the particles had been individually coated with a homogeneous amorphous silica shell of 5-nm-thickness (Si5) using a method derived from the so-called Sto¨ber process. The surface of the BT particles had to be activated by acidic treatment with nitric (1M) and citric (0.01M) acids before silica coating. The reaction occurred in a water/alcohol/ammonia solution using tetraethoxysilane as the silica source. The procedure is described in details elsewhere. [START_REF] Mornet | Controlled Growth of Silica Shell on Ba 0.6 Sr 0.4 TiO 3 Nanoparticles Used as Precursors of Ferroelectric Composites[END_REF] Standard sintering was performed on BT500-Si5 powder at 8001, 10001, and 12101C during 3 h with two additional steps at 2501 and 6001C of 1 h each, under oxygen atmosphere and with a heating rate of 1001C/h.

(2) Characterization of the BT/SiO 2 Interface Temperature-dependent X-ray diffraction (XRD) has been performed using a Philips X'Pert automatic diffractometer (Philips, Almelo, the Netherlands) (CuKa radiation) equipped with a variable-temperature Anton Paar HTK16 Pt stage. The temperature of the sample has been increased from room temperature up to 12051C at a heating rate of 51C/min and the patterns have been recorded at 151C intervals in the range 151r2yr601 with a step size of 0.021, counting two seconds per step.

The microstructural characterization of the sintered compacts was performed using a scanning electron microscope (SEM) JEOL JSM 6360A (JEOL, Cruissy-sur-Seine, France) and a high-resolution scanning electron microscope (HRSEM) JEOL 6700F. A thin Au/Pd coating was deposited on the ceramics fracture surface before observation.

TEM in situ heating experiments were carried out in a JEOL 2010 operated at 200 kV using a Gatan high-temperature TEM holder (Gatan, Evry, France) and SPI silicon nitride membrane window grids. Temperature was controlled by a thermocouple with an accuracy of a few degrees. The dynamical response under thermal solicitation was monitored by means of DVD/HD recording using a MEGAVIEW II CCD camera.

STEM-EDX and HRTEM observations were performed using a JEOL 2200 FS equipped with a field emission gun, oper- ating at 200 kV and with a point resolution of 0.19 nm. HRTEM images were recorded using a Gatan Ultrascan CCD 2 k  2 k and Digital Diffractograms were calculated using the Gatan Digital Micrograph s software. XPS experiments were performed using an ESCALAB 220i XL spectrometer with monochromatic AlKa excitation. Sample charging was partially compensated by an electron flood gun. BaCO 3 , TiO 2 , and SiO 2 were used as precursors for the preparation of fresnoite powder via a solid-state reaction. Fresnoite ceramics were obtained by conventional sintering at 13501C during 3 h. The BT thin films were processed using standard radio frequency sputtering deposition on amorphous SiO 2 substrates.

(3) Advanced Sintering Processes and Dielectric Measurements The BT(500 nm)-SiO 2 (5 nm) powders were sintered using spark plasma sintering (SPS) equipment (Dr Sinter SPS-2080 SPS Syntex INC Japan of the Plateforme Nationale de Frittage Flash (PNF2) of CNRS located at Toulouse, France). Coreshell powders (without any sintering aids) were loaded into a cylindrical die of 8 mm inner diameter. The pulse sequence was 12-2 (pulses-dead time or zero current). The temperature was raised to 6001C over a period of 3 min, and from this point it was monitored and regulated by an optical pyrometer focused on a small hole located at the surface of the die. A heating rate of 1001C/min was used to attain the final temperature of 10501C under an argon (Ar) atmosphere. SPS experiments were performed under a uniaxial pressure of 50 MPa. To observe the microstructure of the SPS ceramics, thin slices were obtained using an ''ultracut S'' ultramicrotome. The ceramic has been clamped directly in the microtome sample holder and a diamond knife has been used to obtain slices of 65-nm-thickness; the slices were then deposited on a copper grid.

Microwave sintering experiments were performed in a TE102 rectangular microwave cavity. The details of the microwave device have been described in a previous paper. [START_REF] Marinel | A New Inductive Furnace Based on Microwave Irradiation for Growing Long YBa 2 Cu 3 O 7Àd Single-Domain Bars[END_REF] A cylindrical LaCrO 3 susceptor was used and placed perpendicular to the electrical field component. The samples used for microwave sintering were thermally pretreated at 2501 and 6001C for 1 h with a heating rate of 1001C/h. As in standard sintering conditions, these two additional annealing steps were performed to ensure removal impurities resulting from the soft chemistry route used.

Then, the preheated disk was placed in an alumina crucible and positioned in the middle of the cylindrical susceptor. The incident power was fixed at 600 W and the thermal cycle duration was 5 min. The sample is, hence, quenched in order to freeze the high-temperature microstructure.

The permittivity and dielectric losses of the ceramics were recorded as a function of temperature (300-500 K) and frequency (100 Hz-100 kHz) using a Wayne-Kerr component analyzer 6425. Dielectric measurements were performed on the ceramic discs (diameter 6 mm, thickness 1 mm) with gold electrodes sputtered on the circular faces.

III. Results and Discussion

(1) Microstructural Investigations of Sintered BT@SiO 2 Particles Calcinations were firstly performed at different temperatures on pellets made of core-shell BT (500 nm)-SiO 2 (5 nm) particles (denoted as [BT500-Si5]) in order to gradually probe the effect of temperature on the interface modification between barium titanate and the initially coated amorphous silica nanolayer. Such a step-by-step approach was performed on both coated and uncoated BT particles. HRSEM pictures of BT500-Si5 calcinated at 8001C showed the formation of interparticle necks growth by a viscous flow (zones pointed by arrows in Fig. 1(a 0 )). [START_REF] Sacks | Preparation of SiO 2 Glass from Model Powder Compacts: II, Sintering[END_REF][START_REF] Jagota | Simulation of the Viscous Sintering of Coated Particles[END_REF] The absence of shrinkage on the dilatometric curve and the density value similar to that of the green body (56%) support a nondensifying mass transport mechanism. The surface of the grains is rough, and an ''orange skin'' aspect on some grains is observed resulting from the dehydration (removal of physically adsorbed water) and deshydroxylation processes involving condensation of the surface silanol groups. [START_REF] Zhuravlev | The Surface Chemistry of Amorphous Silica. Zhuravlev Model[END_REF][START_REF] Sacks | Preparation of SiO 2 Glass from Model Powder Compacts: I, Formation and Characterization of Powders, Suspensions, and Green Compacts[END_REF] When the sintering temperature is raised up to 10001C, a drastic change occurs in the microstructure. The formation of blisters on the surface of the coated grains is clearly observed whereas uniform and glossy grain surfaces appear on the uncoated sample calcinated at 10001C in the same conditions (Figs. 1(b) and (b 0 )). The density remains lower than 90% when the coated powder is sintered at 12101C and grain growth is much more pronounced in the case of the ceramic made of uncoated grains. In situ XRD as a function of temperature was performed to accurately probe the on-set of the fresnoite formation at the BT/ SiO 2 interface and its thermal evolution. The fresnoite, Ba 2 Ti-Si 2 O 8 , occurs on the XRD diagram recorded at 8601C and no titanium-rich phases such as Ba 6 Ti 17 O 40 were detected (Fig. 2(a)). The peculiar microstructure obtained at 10001C can, thus, be correlated to fresnoite crystallization.

(2) Discussion on the Mechanism of Fresnoite Formation The surface of the BT grains is considered here to provide a site for heterogeneous nucleation of fresnoite crystals. This surface, cleaned by an acidic treatment performed before the coating, is rich in barium vacancies, which can enhance diffusion fluxes. It is known for a long time that the Ba of BT surface dissolves at low pH and acts as a potential-determining ion (PDI) in the solution. This explains the difficulty to disperse BT particles in water. [START_REF] Chen | Stabilization and Processing of Aqueous BaTiO 3 Suspension with Polyacrylic Acid[END_REF][START_REF] Shih | Silica Coating on Barium Titanate Particles[END_REF][START_REF] Anderson | Surface Chemistry Effects on Ceramic Processing of BaTiO 3 Powders[END_REF] In previous zeta potential studies, we observed that cleaning the particles with acidic treatment allowed removing these PDI from the surface of Ba 0.6 Sr 0.4 TiO 3 particles. [START_REF] Mornet | Controlled Growth of Silica Shell on Ba 0.6 Sr 0.4 TiO 3 Nanoparticles Used as Precursors of Ferroelectric Composites[END_REF] After HNO 3 treatment, the isoelectric point of these particles was shifted from 8.5 to 5.1, closer to the value of TiO 2 (5.5) and SrTiO 3 (5.1). [START_REF] Gao | Microstructure-Controlled Deposition of SrTiO 3 Thin Film on Self-Assembled Monolayers in an Aqueous Solution of (NH 4 ) 2 TiF 6 Sr(NO 3 ) 2 H 3 BO 3[END_REF] The onset of fresnoite nucleation was clearly determined at 8601C by XRD and thermal analysis. The evolution of the fresnoite formation as a function of temperature, was estimated considering the relative intensities between BT and the fresnoite main peaks (I fresn. /I BT ) from in situ XRD experiments. Close correspondence with the dilatometric investigation carried on BT500-Si5 (shrinkage rate shown in insert) can be evidenced, allowing thus to propose a mechanism for the fresnoite forma-tion and interface evolution (Fig. 2(b)). The increase of I fresn. /I BT ratio in the range 8001-9001C obviously coincides to the first nucleation of fresnoite crystals. This is consistent with the first shrinkage observed in the same temperature range and which does not occur when BT and SiO 2 are sintered separately (not shown). Considering the initial continuous silica coating on each individual particle, we assume here that the fresnoite is formed uniformly and at the interface at the grain scale. A stabilization of the ratio I fresn. /I BT is observed in exactly the same temperature range (9001-11501C) as the plateau of the shrinkage rate. The formation of fresnoite crystallites hinders densification. Note that in the case of pure SiO 2 particles, densification starts at a temperature as low as 10501C. The increase of the volume fraction of fresnoite inclusions, or the fresnoite particle growth, delays sintering by a viscous flow. However, at least up to 10001C, these inclusions are not considered as rigid inclusions in view of the presence of amorphous silica, which contribute to lubricate the contact between them. The further increase in the (

3) In Situ TEM Experiments

To support the above-mentioned sintering model, in situ heating TEM experiments were performed up to 10001C in order to visualize the gradual nanoscale transformation of fresnoite at the BT/SiO 2 interface. The aim was to obtain a direct imaging of the dynamics of the core-shell grains under thermal stress, i.e. the crystallization process at the interface. Initial 300-nm BT cores coated with 5-nm silica were used for imaging scale convenience. It was previously demonstrated that the interface evolution between silica and BT is not dependent on the ferroelectric core size.

Figures 3 and4 show pictures extracted from a video sequence taken at different temperatures. Every picture was recorded after holding the set temperature for a few minutes. Every temperature step was achieved in less than a minute. In Fig. 3(a), a single core-shell particle of 300 nm in diameter is shown at 6001C. Because the particle is seen in projection, the silica shell (labeled S) appears as a thin 5-nm layer in the outer part of the particle (note the dashed line delimiting the shell in Figs. 3(e) and (f)). The absence of thickness fringes in the grain core but close to the boundary indicates that the particle is faceted (Fig. 3(a)). This implies the presence of sharp edges, which show an increase of brightness from the center, opaque to the electrons, to the thin outer shell. The same particle morphology is also seen in Fig. 6. Other particles with the same morphology are visible on the left part of the image (Fig. 3). In the areas labeled N, where grains are in contact, different grain surfaces can be easily distinguished. Electron diffraction indicates that the particles are single crystalline, but some small unidentified surface defects (marked D in Figs. 3(a)-(d)) are usually observed. Below 8501C, the silica shell does not exhibit any visible structural modifications (Figs. 3(b) and (e)). The observation of necks in the contact area seems to indicate that a viscous flow of silica has led to the formation of bridges between adjacent grains. This has also led to a global redistribution of matter around the core; one can see by comparing the particle shape in Figs. 3(a) and (d) that the shell surface is less faceted at higher temperature. At 8501C, some crystallites, already as large as the silica-shell thickness, are visible everywhere in the outer shell (Figs. 3(c) and (f)). They appear as black dots in the silica amorphous shell. This strongly suggests that they should correspond to the fresnoite, as observed during XRD experiments. Such a good agreement with XRD data shows that both e-beam irradiation and vacuum have negligible influence on sintering and crystallization processes. Only a possible electron beam heating and charging effects have to be considered. Silica bridges in the necked regions keep growing due to the combination of capillarity forces and a faster diffusion at higher temperature. This confirms that fresnoite crystallites do not act as pinning centers preventing the viscous flow at least up to 10001C. Despite the existence of fresnoite crystallites that should consume the silica, the amorphous shell is still preserved up to 10001C (Figs. 3(c) and(d)). From 8501 up to 10001C, the surface roughness increases as depicted in Fig. 4, which shows pictures taken in another area. In some other particles observed, this roughening leads to the formation of blisters such as seen on HRSEM pictures (see Fig. 1(b 0 )). Experiments carried out at different temperatures seem to indicate that this roughening is thermally activated.

(4) STEM-EDX and HRTEM Experiments

To supplement this investigation with chemical probing at the nanoscale, STEM-EDX analysis has been carried out postmortem on the sample investigated by in situ TEM at 10001C as well as on the starting BT@SiO 2 material.

One can observe in Figs. 5(a) and (b) that as-prepared BT cores are homogeneously coated with an approximately 20 nm- thick silica shell. Please note that the silica-shell thickness is slightly overvalued compared with the 5 nm mentioned previously. This can be due to the long-time exposure under the fo-cused electron beam, which favors the growth of the outer silica shell because of its collisionnal and local heating effects. The slope of the [Ba]/[Si] profile (Fig. 5(c)) plotted from STEM-EDX map is rather flat for the inner core zone and decreases abruptly for the silica shell, confirming that BT particles are homogeneously coated and pointing out that the interface zone between the core and the shell is well defined.

The same experiment performed on the in situ sample shows a slightly thicker silica shell, highlighting its viscous character once the 10001C range of temperature is attained. But the most important feature is the observation of dark-contrasted zones inside the silica shell (Fig. 6(a)). These zones pointed by arrows on the three color image of Fig. 6(b) and appearing pinker than the purple outer shell (larger amount of green color component), are revealed to be Ba enriched. Moreover, whereas the extreme leftand right hand-side slopes of the [Ba]/[Si] profile plotted in Fig. 6(c) are rather similar to those plotted in Fig. 5(c) for the starting material, it softens for the dark-contrasted zone.

One can see on the HRTEM micrograph in Figs. 7(a) and (b) that the dark-contrasted clusters are crystallized as reticular planes can be observed. The indexation of digital diffractograms and the measurement of reticular distances on HRTEM images unambiguously points toward fresnoite (space group P4bm, a 5 8.518(2) A ˚, c 5 5.211(1) A ˚)[ICDD-PDF: 22-0513]. This strongly suggests that the diffusion process occurring from the BT core (Ba diffusion into silica shell) led to the nucleation of fresnoite clusters into the silica shell. It is worth noting that edges or corners of the crystalline BT grains seem to be energetically favorable for fresnoite nucleation.

Furthermore, the crystallographic investigation of the interface between BT and fresnoite showed that fresnoite clusters grow at the surface of the BT cores according to the following orientation relationship (Fig. 7(c)): ð0 10Þ½100 BT ==ð040Þ½101 fresn: These ð0 10Þ BT and (040) fresn. interface planes are Ba rich. Moreover, similar tilted rectangular patterns (in pink in Fig. 8(d) can be found in the ð0 10Þ BT and (040) fresn. planes and belonging to the (001) BT and ð20 2Þ fresn: facing each other from both sides of the interface plane. The growth of fresnoite clusters on BT cores is then chemically and structurally favored.

As for other nanocomposite systems, one can also observe in Fig. 7(b), the presence of a dislocation network, perpendicular to the interface plane between the BT core and the fresnoite clusters. [START_REF] Majimel | Synthesis of Hybrid Gold-Gold Sulfide Colloidal Particles[END_REF] It seems that the lattice mismatch between the two phases is relaxed by dislocation creation rather than elastic deformation. The Da/a lattice misfit value of about 10.2%, is approximately in good agreement with the presence of three dislocations for 30 lattice plane, as we can see in Fig. 7(b) (Fig. 7(c)). Finally, note that the 71 angle between the (001) BT and ð20 2Þ fresn: corresponds to an angular deviation from the epitaxial relationships mentioned above. This deviation increases as the interface plane length decreases.

(5) XPS Experiment In contrast to TEM, XPS probes a large area and thus provides information over a large amount of particles. XPS spectra of BT powder coated with a 5-nm-thick SiO 2 shell are shown in Fig. 8 in comparison with a sputter-deposited BT thin film and two fresnoite reference samples (ceramic and BT/SiO 2 thin film after heating to 10001C). The as-prepared core-shell powder shows a strong SiO 2 emission at B103.5 eV binding energy and a weak Ba 4d emission with a binding energy of the 4d 5/2 component of B89 eV, which can be attributed to BT. Heating the powder to 10001C in air strongly decreases the SiO 2 intensity and increases the Ba 4d intensity. Two distinct Ba 4d components can be identified, which correspond to the BT and the fresnoite references. The fresnoite phase is identified by the binding energy difference and the intensity ratio of the Si 2p and Ba 4d emissions. Considering the different intensities, the XPS spectrum of the air-heated core-shell powder (10001C) is consistent with a partial coverage of BT by fresnoite particles. This is consistent with postmortem and in situ TEM investigations. The close correspondence between the XPS spectra of the two fresnoite reference samples and the calcinated core-shell particles (c, d, e in Fig. 8) also indicates that XPS studies of thin-film samples can be useful for a detailed study of the interface chemistry between SiO 2 and BT. Thin-film samples might provide a considerable advantage compared with ceramic samples in such studies as they can be deposited onto conducting substrates. Thereby, charging effects, which often lead to considerable broadening and asymmetries of the core level lines, can be avoided.

To summarize all the previous microscopic investigations, interdiffusion between SiO 2 and BT leads to the formation of fresnoite which nucleates in the early stage of sintering, i.e. 8601C (in situ XRD). Even if cross diffusion at the BT/SiO 2 interface occurs, HRTEM confirmed the higher mobility and diffusion of Ba ions into the silica shell. The fresnoite crystallites occur preferentially at energetically and crystallographically favorable sites such as defects, vacancies, edges, corners, etc. As a result, the silica shell roughens (in situ TEM) and as confirmed by XPS, blisters are observed instead of homogeneous fresnoite coating. The dielectric measurements performed on composite ceramics are now described as intrinsic and extrinsic macroscopic probes.

(6) Advanced Sintering Processes and Dielectric Properties

The fresnoite Ba 2 TiSi 2 O 8 is described in literature as a new ferroelectric mineral exhibiting nearly constant permittivity values close to 10 between 300 and 600 K. [START_REF] Foster | Fresnoite: A New Ferroelectric Mineral[END_REF][START_REF] Halliyal | Unusual Pyroelectric and Piezoelectric Properties of Fresnoite (Ba 2 TiSi 2 O 8 ) Single Crystal and Polar Glass-Ceramics[END_REF] This second phase as well as the porosity decreases the permittivity in the composite compared with the uncoated BT but does not affect the Curie temperature of the BT core. As shown previously, a sintering temperature higher than 12001C in standard conditions does not allow attaining 90% of density while the full transformation of the initial amorphous silica into fresnoite is already obtained. The crystallization into fresnoite disrupts the continuity of the silica dielectric phase separating the BT grains. The structural matching between fresnoite and BT, confirmed by HRTEM, suggests that the composite is no longer formed by two distinct phases as in classical ferroelectric/dielectric composites.

Crystallization is governed by two kinetic steps, nucleation and growth. The increase of the heating rate from low temperature, where the nucleation rate of the crystals is fast, can delay the onset of crystallization. [START_REF] Panda | Effect of the Heating Rate on the Relative rates of Sintering and Crystallization in Glass[END_REF]36 In the same way, reducing the sintering time would help limiting fresnoite growth. For this purpose, advanced sintering techniques such as SPS and m-wave sintering were used to densify BT500-Si5 particles. Both techniques provide extremely high heating rates enhancing the consolidation rate of powders. In the case of SPS, the combination of pressure and electric current activates and enhances densification. In particular, pressure can help densification by overcoming the restraining stress at the interface. The extremely accelerated kinetics should prevent excessive nucleation at low temperature and lead to more complete densification before extensive crystallization. In addition, short sintering times (3 min for SPS and 5 for m-wave) and quenching conditions in both SPS and m-wave techniques favor the freezing of the high-temperature microstructure.

The sample, sintered 3 min at 10501C by SPS under Argon atmosphere, exhibits a peculiar microstructure made of rounded grains with a melted looking aspect (Fig. 9(a)). Uncoated BT grains sintered in exactly similar SPS sintering conditions are shown for sake of comparison. The density of the ceramic reaches 82% instead of 70% when using standard sintering at a similar temperature during 2 h. Slightly increasing the sintering temperature significantly raises the density up to more than 90% but with unavoidable reduction phenomena leading to a strong conductivity contribution. Dielectric experiments show a reduction and a flattening of the maximum of losses at T c compared with the ceramic made of uncoated grains sintered in the same SPS conditions (Fig. 9(b)). This proves that the silica shell is preserved in the final composite, as confirmed by HRTEM observations. The existence of an amorphous phase in the vicinity of the grain boundaries has been pointed out in thin slices, obtained using ultramicrotomy, of SPS ceramics (Fig. 10(a)). The micrograph is representative of several zones on the sample. Small fresnoite crystals were also observed (Fig. 10(b)). More generally, an enlarged view of the slice confirms that BT grains are not interpenetrating.

Microwave sintering performed at a temperature estimated to be close to 12001C leads to higher density of 92%. The reproducibility of these experiments was checked and confirmed. In addition, because we used a susceptor, microwave sintering is no longer considered as a volumic sintering and the microstructure remains homogeneous, as confirmed by HRSEM imaging. No blisters occur and grain growth is still avoided whereas the final sintering stage here is attained (Fig. 9(a)). Dielectric measurements can be once more used as an intrinsic probe to check the efficiency of the silica loss barrier. The loss values still remain stable and well below 1%, while the permittivity reaches 3000 at room temperature (Fig. 9(b)).

IV. Conclusion

Silica-coated ferroelectric particles were investigated with, as a front edge challenge, the control of the interfaces in such bulk nanostructured composites. The first step was to perform an accurate and deep investigation of the complex silica/ferroelectric interface. Postmortem and in situ TEM experiments showed the presence of amorphous silica up to 10001C. Neck formation with an associated redistribution of silica around the ferroelectric core was observed below 8501C. Fresnoite crystallites appeared in the temperature range 8501-10001C in agreement with both XRD and dilatometric investigations. Postmortem chemical probing at the nanoscale strongly supports a diffusion process from the BT core, i.e. Ba diffusion into the silica shell. Electron diffraction investigation of the interface showed that the fresnoite cluster growth at the BT surface is chemically and structurally favored. Creation of a dislocation network is believed to help in relaxing the lattice mismatch. In addition, a partial coverage of the ferroelectric core by fresnoite particles was confirmed by XPS probing. Such a high level of interface probing was never reported in such silica-nanostructured ceramics. The second step was to process nanostructured ceramic composites using advanced sintering processes such as SPS and microwave sintering. In both cases, the permittivity attained at least 2000 at room temperature and a stabilization of the dielectric losses, well below 1%, was obtained with, in particular, a significant flattening of the losses at the Curie temperature, thus signing the preservation of the core-shell architecture. This was supported by HRTEM observation on a thin SPS ceramic slice obtained by ultramicrotomy.

The possible structural and chemical matching between BT and fresnoite, and the close correspondence evidenced between XPS spectra of BT/SiO 2 thin films and ceramics made of BT@SiO 2 particles, open routes to process multifunctional composites with multilayer design.

Fig. 1 .

 1 Fig. 1. Top: Transmission electron microscopy image of silica (5 nm)-coated BT (500 nm) (BT500-Si5) particles. High-resolution scanning electron microscope fracture images of uncoated BT (500 nm) ceramics and BT500-Si5 sintered at 8001C (a, a 0 ), 10001C (b, b 0 ). (a 0 ) Arrows indicate neck formation.

Fig. 2 .

 2 Fig. 2. (a) Temperature-dependent X-ray diffraction for temperatures ranging from 8301C up to 12051C for BT500-Si5 powders. Fresnoite (Ba 2 TiSi 2 O 8 ) peaks clearly appear at 8601C. (b) Relative intensities between the fresnoite and the BT peaks, I fresn. /I BT from in situ XRD as a function of temperature. Inset: Shrinkage rates as a function of temperature for BT500-Si5. Three regimes can be evidenced: 1, fast nucleation of fresnoite up to 9001C; 2, stabilization of I fresn. /BT corresponding to hindered densification (plateau in the shrinkage rate) between 900 and about 11501C; and 3, abrupt increase in the I fresn. /I BT ratio signing to the onset of densification at temperatures higher than 11501C.

Fig. 3 .

 3 Fig.3. Transmission electron microscopic micrographs of a core-shell particle taken during in situ heating experiment. Note the presence of surface defects (D) before the heating run. Below 8501C (a-b), viscous flow of silica softens necks between neighbouring particles (N). At 8501C (c), crystallites nucleate in the outer shell. The core-shell structure is, however, preserved up to 10001C (d). Note that the necks tend to be smoother. (e) and (f) show the particle at a higher magnification, highlighting the crystallization formation (F) at 8501C in the outer shell (S). The silica shell (S) is indicated by arrows and delimited by dashed lines. The brightness rapidly decreasing from the outer shell to the particle center, is due to the sharp edges.

I

  fresn. /I BT ratio corresponds to the densification on-set at 11501C. At this stage, the amorphous silica layer is practically fully consumed and transformed into fresnoite at the interface. It is difficult to conclude on the sintering mechanism between densifying viscous flow regime (a thin wetting silica layer with appropriate viscosity can be sufficient) and change into volume diffusion or grain boundary diffusion. Isothermal sintering performed at 12101C did not provide a better insight into the sintering mechanism, suggesting that a combination of the two mechanisms occurs.

Fig. 4 .

 4 Fig. 4. Transmission electron microscopic micrographs taken during in situ heating experiment. Between 8501C (a) and 10001C (b), a strong roughening of the surface occurs, leading to the formation of blisters.

Fig. 5 .

 5 Fig. 5. (a) STEM micrograph of a BaTiO 3 -silica starting particle. (b) Three colors image drawn from the EDX map of the particle presented on (a). The red, green, and blue colors correspond to the oxygen, the barium, and the silicon elements, respectively. (c)[Ba]/[Si] profile plotted from (b) the map along the black line of figure (a). Dashed lines indicate the frontiers between the inner BT core and the outer silica shell.

Fig. 6 .

 6 Fig. 6. (a) STEM micrograph of a particle of the BaTiO 3 -silica sample investigated by in situ TEM at 10001C. (b) Three colors image drawn from the EDX map of the particle presented on (a). The red, green, and blue colors correspond to the oxygen, the barium, and the silicium elements, respectively. Fresnoite clusters correspond to the zones marked by arrows. (c)[Ba]/[Si] profile plotted from (b) the map along the black line of figure (a). Dashed lines indicate the frontiers between the inner BT core, the fresnoite clusters, and the outer silica shell.

Fig. 7 .

 7 Fig. 7. (a) High-resolution transmission electron microscopy micrograph of a particle of BaTiO 3 -silica investigated by in situ TEM at 10001C. This image corresponds to the zone delimited by the white square on (a). On the upper left and lower right, insets shown are the indexed digital diffractograms of the BT core and of the fresnoite cluster, respectively. (b) Zoom on the interface zone between the BT core and the fresnoite cluster. White and red dashed lines represented on the inset correspond to BT and fresnoite reticular planes, respectively. The additional half-plane of the BT structure is a misfit dislocation represented by the > symbol. (c) Schematic view of the BT/fresn. interface zone according to the ð0 10Þ[100] BT // (040)[101] fresn . orientation relationship and calculated Da/a lattice misfit (m). (d) Projection of Ba atoms of BT and fresn. interface planes. Atoms highlighted in pink can face each other from both sides of the interface, the rectangles in dashed lines being distorted according to the misfit value.

Fig. 8 .

 8 Fig. 8. X-ray photoelectron spectra of the Si 2p-Ba 4d region. Measured samples are (a) BaTiO 3 thin film on Pt, (b) BaTiO 3 powder covered with 5 nm of SiO 2 , (c) sample (b) after heating to 10001C in air, (d) BaTiO 3 thin film on SiO 2 after heating to 10001C in air, which caused the transformation of the film to fresnoite, and (e) fresnoite ceramic. Component I corresponds to Si 2p in fresnoite, II to Ba 4d 3/2 in fresnoite, and III to Ba 4d 5/2 in BaTiO 3 . Spectra are manually shifted in binding energy for a better identification of chemical species. The asymmetry observed for samples (b), (d), and (e) is attributed to inhomogeneous sample charging.

Fig. 9 .

 9 Fig. 9. (a) High-resolution scanning electron microscope microstructures and (b) the dielectric properties of uncoated BT500 ceramic sintered by SPS, BT500-Si5 sintered by spark plasma sintering (SPS) and microwave.

Fig. 10 .

 10 Fig. 10. (a) High-resolution transmission electron microscopy images of a 60-nm-thick slice obtained by ultramicrotomy of BT500-Si5 sintered at 10501C by spark plasma sintering. (a) Arrow indicates an amorphous phase in the vicinity of the grain boundaries in between two crystallized grains. (b) Small crystal with a reticular distance d 5 2.64 A ˚corresponding to fresnoite. This small crystal is supported by a second crystal, which is expected to be a BT grain in view of both the size and the reticular distance (d 5 2.09 A ˚).

J. Am. Ceram. Soc., 74, 2669-71(1991). &

Acknowledgments

We thank Dominique Denux for TMA and Eric Lebraud for XRD measurements.

M. Menon-contributing editor

This work was financially supported by the French National Center for Scientific Research (CNRS), the National Research Agency (ANR), and the European Community in the frame of the ''FAME'' network of excellence.


Omomene Iwelomen

Jean Toniolo^

Pascale Beloni

Une revue systématique des programmes d'éducation thérapeutique du patient (ETP) sur le diabète en Afrique subsaharienne

Contact : omenewelo@gmail.com -jean.toniolo@unilim.fr

  une maladie chronique qui, si elle n'est pas traitée, peut entraîner la cécité, une insuffisance rénale, un infarctus du myocarde, un accident vasculaire cérébral ou l'amputation d'un membre inférieur [IDF, 2021]. Le fardeau mondial du diabète augmente dans le monde entier et à un rythme plus rapide dans la région de l'Afrique subsaharienne (ASS), où l'accès aux soins et aux médicaments antidiabétiques est complexe, entrainant des difficultés à stabiliser la pathologie [IDF, 2021]. Une éducation thérapeutique du patient (ETP) intentionnelle et structurée favorise les résultats pour la santé des patients diabétiques lignes directrices Cochrane et PRISMA ( PROSPERO ID : CRD42023440701). Bases de données : PubMed, Web of Science, CINAHL et Google Scholar Sélection des études : 14 mars -30 Avril 2023. Sélection des titres et des résumés par deux chercheurs en aveugle à l'aide du logiciel Rayyan QCRI. Une synthèse qualitative à partir des données pertinentes extraites des articles éligibles. Évaluation de qualité : outil « Risk of Bias 2.0 » de Cochrane. Critères d'inclusion : Études rapportant une intervention d'ETP concernant le diabète chez l'adulte comparée aux soins standards, réalisées en ASS. études incluses : 11 Essais Comparatifs Randomisés (ECR) et 2 non-ECR. Répartition géographique : Afrique du Sud [n=5], Nigéria [n=2], Cameroun, Ghana, Éthiopie, Rwanda, Kenya et Mali. 3017 participants dont 932 hommes et 2085 femmes. Diabète de type 2 [n=12] Vs Diabète de type 1 [n=1]. Modes de diffusion : séances de groupe, individuelles, combinées, envoie de SMS et forum WhatsApp 4 études étaient dirigées par une infirmière, 2 par un diététicien, 1 par des pairs, 1 par des éducateurs sanitaires et 5 par une équipe médicale multidisciplinaire La durée moyenne des interventions : 9 mois La fréquence moyenne des séances de groupe : 8 séances. Petits échantillons, pas de mesure standardisée de la glycémie, période de suivi courte, attrition élevée. Educators (AADE) : alimentation saine, être actif, autosurveillance, prise de médicaments, résolution de problèmes, réduction des complications et gestion du stress. Adaptation du matériel et des contenus pédagogiques au contexte local rapportée dans la plupart des études. élément important dans la gestion du diabète pour aider à améliorer la santé des patients. Les programmes actuels ne sont pas toujours efficaces, notamment concernant l'observance médicamenteuse et sur les modifications de la glycémie. Les interventions dirigées par des infirmières et par des pairs sont particulièrement importantes dans cette région confrontée à une désertification médicale. Des études qualitatives doivent être réalisées pour s'enquérir de l'attrition élevée observée et des faibles taux de fréquentation des centres de santé, afin de mieux adapter les programmes d'ETP. Des études médico-économiques sont nécessaires pour déterminer les implications financières d'interventions similaires à des échelles nationales afin d'élaborer des politiques de santé globale.