C Rigoudy

K Makasheva

M Belhaj

S Dadouch

G Teyssedre

L Boudou

Rational Engineering of the Dielectric Properties of Thin Silica Layers with a Single Plane of AgNPs

Interface properties of dielectrics are of outmost importance when dealing with phenomena as charge injection or secondary electron emission. When the dielectric thickness scales down to the nm range, a deviation from the bulk material properties occurs; therefore, a possibility to tune the surface properties comes by modifying dielectrics at its sub-surface. This paper focuses on the precision of plasma-based processes in order to elaborate thin silica dielectric layers with a single plane of silver nanoparticles (AgNPs) at its sub-surface and thus to control their behavior under electrical stress. An illustration of the modulated properties of such nanostructured dielectrics is given through current measurements and electron emission.

I. INTRODUCTION

The dielectric charging phenomenon, i.e. the ability of dielectric materials to trap electrical charges, leads to modification of the local electric field. It causes change in the electrical behavior, and on a longer term degradation of the dielectric properties of materials under electrical stress [START_REF] Rezende | Detection of charge distributions in insulator surfaces[END_REF][START_REF] Laurent | Charge dynamics and its energetic features in polymeric materials[END_REF]. The dielectric charging phenomenon is twofold: (i) it lays down the principle of operation of many applications, among which non-volatile memories [START_REF] Rao | Silicon nanocrystal based memory devices for NVM and DRAM applications[END_REF] and electret based microphones [START_REF] Rebeiz | RF MEMS: Theory, Design, and Technology[END_REF], and (ii) it is considered the main cause of electrostatic sticking in MicroElectroMechanical Systems (MEMS) [START_REF] Hillenbrand | Electret microphones with stiff diaphragms[END_REF], charge build-up and breakdown in insulations [START_REF] Laurent | Charge dynamics and its energetic features in polymeric materials[END_REF][START_REF] Makasheva | Dielectric Engineering of Nanostructured Layers to Control the Transport of Injected Charges in Thin Dielectrics[END_REF][START_REF] Milliere | Silver nanoparticles as a key feature of a plasma polymer composite layer in mitigation of charge injection into polyethylene under dc stress[END_REF] or being at the origin of potential gradients that may generate ElectroStatic Discharges (ESD), thus damage electronic devices and limit their lifetime, especially in space environment [START_REF] Balcon | Secondary electron emission on space materials: Evaluation of the total secondary electron yield from surface potential measurements[END_REF]. The later example is closely related to another physical phenomenon, namely electron emission from dielectrics under irradiation. In general, the emission efficiency is extremely sensitive to the surface properties of irradiated materials (composition, surface state, etc.), since the secondary electrons are typically emitted from the first 10 nm beneath the surface of the irradiated material. Main solutions to avoid ESD in those cases are oriented to geometric optimization of the devices and to design of surface coatings.

Even though the required outcome is on macroscopic level the solution is often existing at nanoscale and can be achieved by rational engineering of the dielectric properties of materials [START_REF] Makasheva | Dielectric Engineering of Nanostructured Layers to Control the Transport of Injected Charges in Thin Dielectrics[END_REF]. The lately developed nanometric dielectrics, or so-called nanodielectrics, have shown a considerable progress on many of the global electrical parameters of nanocomposite dielectrics (permittivity, losses, low field conductivity, space charge, etc.) since the pioneer works in the early 90s [START_REF] Lewis | Nanometric dielectrics[END_REF]. The Total Electron Emission Yield (TEEY) from dielectric materials under irradiation is one of the global parameters to control, when the featured application is intended for space environment. Therefore, sustainable solution calls for fine tuning of the material composition and structure, and a thorough analysis of the dielectric charging phenomenon.

This work aims to explore the impact of metal inclusions (silver nanoparticles, AgNPs), embedded in thin silica (SiO 2 ) layers, on the physical mechanisms responsible for dielectric charging and electron emission, in order to modulate them. Our strategy relies on synthesis of thin stratified coatings and is based on the principle of combined effect from metal/ dielectric interfaces. An illustration of the modified properties of such nanostructured dielectrics is given through current measurements and electron emission in space environment.

Elaboration of stratified dielectric thin films requires facilities/methods permitting strict control of alignment of the metal inclusions at a given nanometric distance beneath the surface. Chemical or physical-based methods, or combination of both, can reply to the above requirements, however with different accuracy. Although well adapted for technological transfer the chemical (wet) methods for synthesis present various drawbacks as wasted materials (solvents), lack of precision in the metal NPs plane depth, uncertainties of the NPs distribution in the plane or different types of ligands that may preferentially trap electrical charges. Combined elaboration methods like dry dipping followed by plasma deposition [START_REF] Deng | Antimicrobial nano-silver non-woven polyethylene terephthalate fabric via an atmospheric pressure plasma deposition process[END_REF], or physical methods like gas aggregation source [START_REF] Peter | High rate deposition system for metalcluster/SiOxCyHz-polymer nanocomposite thin films[END_REF], ion implantation [START_REF] Carles | The synthesis of single layers of Ag nanocrystals by ultra-low-energy ion implantation for largescale plasmonic structures[END_REF], or plasma-based processes as the one employed in this work [START_REF] Makasheva | Dielectric Engineering of Nanostructured Layers to Control the Transport of Injected Charges in Thin Dielectrics[END_REF][START_REF] Milliere | Silver nanoparticles as a key feature of a plasma polymer composite layer in mitigation of charge injection into polyethylene under dc stress[END_REF][START_REF] Pugliara | Controlled elaboration of large-area plasmonic substrates by plasma process[END_REF], should be privileged because of the accuracy in the process and the lack of uncontrolled chemical species surrounding the metal NPs.

II. PLASMA PROCESS FOR SAMPLE ELABORATION AND CHARACTERIZATION METHODS

The targeted structures consisted in a single plane of AgNPs embedded at 10 nm beneath the surface of 200 nm thick SiO 2 layer deposited on mirror polished stainless still and intrinsic Si-substrates. The searched total thickness of this stratified structure was of 220 nm. The selection of stainless still substrates (3.3cm × 3.3 cm) was related to the possibility to ensure good electrical contact on the back side for the charge current and electron emission measurements.

To achieve performance of the dielectric structure, the AgNPs should be well aligned in a plane, small enough so that their surface to volume ratio remains large, well distributed in the plane, with density that ensures charge transport regime far from the electrical percolation limit, however with interparticle distance allowing charge transport through tunneling or trap-assisted processes.

A. Plasma process and real-time monitoring

The above defined nanostructured thin dielectric layers have been elaborated in a plasma process successfully combining in the same reactor sputtering of a metallic target (physical vapor deposition -PVD) and plasma polymerization (plasma enhanced chemical vapor deposition -PECVD) [START_REF] Makasheva | Dielectric Engineering of Nanostructured Layers to Control the Transport of Injected Charges in Thin Dielectrics[END_REF][START_REF] Milliere | Silver nanoparticles as a key feature of a plasma polymer composite layer in mitigation of charge injection into polyethylene under dc stress[END_REF][START_REF] Pugliara | Controlled elaboration of large-area plasmonic substrates by plasma process[END_REF].

The plasma was sustained in an argon axially-asymmetric radio-frequency (RF, f = 13.56 MHz) capacitively-coupled discharge at low pressure (p =5.3 Pa). The RF-driven top electrode (smaller, 10 cm in diameter) was placed in front of a bottom electrode (larger, 12 cm in diameter) with a gap distance of 3.5 cm. The bottom electrode and the walls were grounded. The discharge powered electrode was an Ag-made target to bear the silver sputtering. The axially-asymmetric design of this RF discharge induces a self-bias voltage V dc on the powered electrode, controlling in that way the sputtering. The self-bias voltage is at the origin of the acceleration of argon ions towards the silver target to achieve bombardment. The size and density of the deposited AgNPs were controlled by the injected in the plasma power P = 80 W (V dc = -950 V) and the sputtering time (t s = 5 s) for a given pressure.

SiO 2 -base and cover layers were deposited in the same reactor in the plasma sustained by P = 120 W (V dc = -820 V). Optimal parameters for silica deposits have been identified in a previous work for the very same reactor [START_REF] Makasheva | Dielectric Engineering of Nanostructured Layers to Control the Transport of Injected Charges in Thin Dielectrics[END_REF][START_REF] Pugliara | Controlled elaboration of large-area plasmonic substrates by plasma process[END_REF]. Argon was the vector gas (Ar flow = 2.8 sccm). The precursors used for SiO 2 deposition were hexamethyldisiloxane (HMDSO, C 6 H 18 Si 2 O) and O 2 (O 2flow = 1.25 sccm). The HMDSO was injected by pulses (t on = 3.1 s and period T = 5 s) giving an average HMDSO flow of 0.248 sccm. The total pressure was of p = 7.9 Pa. The thicknesses of the two silica layers were controlled by the deposition time (t d-SiO2-base = 25 min and t d- SiO2-cover = 90 s). The use of oxygen has two purposes: to bring oxygen atoms needed for the SiO 2 -deposition and to ensure plasma combustion of the C-content by creating by-products like CO, CO 2 , COH 2 in order to evacuate the carbon atoms originating from the HMDSO decomposition.

The plasma state was monitored at each step of the deposition by measurements of the self-bias voltage and by Optical Emission Spectroscopy (OES) with a Princeton Instrument Acton Advanced Sp 2500A in the range 500 -600 nm with resolution of 0.2 nm.

B. Structural characterization methods

Structural characterization of the resulting samples has been performed to determine the thickness and composition of the plasma silica matrix as well as to obtain the AgNPs size, shape, density and distribution and the total thickness of the structure.

The thicknesses of SiO 2 layers were determined by spectroscopic ellipsometry using a Semilab ellipsometer SE-2000 platform in the range 250-850 nm. Bruggeman's model was used to analyze the recorded spectra.

The types of molecular bonds within the silica layers were determined by Fourier Transform InfraRed (FTIR) spectroscopy using a VERTEX70 in transmission mode in the range 400 -4000 cm -1 with resolution of 2 cm -1 . The transmission mode was achieved due to the transparency of intrinsic Si-substrates in the infrared domain.

Transmission Electron Microscopy (TEM) analysis of the elaborated structures in plane and cross-section views was performed with a JEM 2100F microscope operating at 200 kV.

C. Characterization of the electrical behavior of the structures

The electrical properties of the nanostructured layers have been studied through measurements of the polarization currents under strong electric fields and of the TEEY under electronic radiation conditions.

The study of polarization currents under strong electric fields was performed through a protocol of successive and increasing voltages [START_REF] Makasheva | Dielectric layers for RF-MEMS switches: design and study of appropriate structures preventing electrostatic charging[END_REF]. Each voltage step lasts for 1000 s and is followed by a depolarization time at 0 V during 1000 s. The voltage was applied between the metallic substrate and an upper Au-electrode (150 nm thick and 1 cm in diameter).

The TEEY measurement were realized under high vacuum (10 -6 Pa) in the energy range 0 -2000 eV, using a 2 keV Kimball electron gun, according to the protocol described in details elsewhere [START_REF] Rigoudy | Study of required conditions to limit the dielectric charging phenomenon when measuring the electron emission yield from thin dielectric layers[END_REF].

III. RESULTS AND DISCUSSION

A. Importance of plasma stability during sample elaboration

The plasma stability during deposition was real-time monitored by OES and measurements of the self-bias voltage. The intensity of the emission lines is proportional to the electron density and the electron temperature in the plasma and the self-bias voltage is related to the power absorbed in the plasma. Thus, these two indicators reflect the stability of plasma deposition process and consequently, the precision and quality of the deposited layers. Typically, stabilization of the plasma process was achieved 30 min after plasma ignition. OES spectra of a stable plasma are shown on Fig. 1. For the silica deposition (Fig. 1a) only Ar-lines can be observed in the spectrum. A stable intensity of the Ar-line at 549.6 nm is a prerequisite to start deposition onto the substrates. Another condition is the absence of Ag-lines. It is achieved when the silver target is completely covered by SiO 2 and no Ag-atoms are sputtered into the plasma. For the AgNPs-deposition however, the Ag-lines are clearly observed (Fig. 1b). To correlate the amount of sputtered silver we use as indicator the ratio between the intensities of Ag-(at 546.5 nm) and Ar-(at 549.6 nm) lines. The same intensity ratio (I Ag(546.5nm) /I Ar(549.6nm) ) is a criterion for reproducibility of the AgNPs-deposition.

Stability of the emission line intensities and of the self-bias voltage are the two key parameters in the deposition process, especially for long deposition times, as they give possibility to associate the deposition time to the composition (in particular the silica layer quality) and the thickness of deposited layers.

B. Quality of the plasma deposited SiO 2 layers and quantitative analysis of the single layer of AgNPs

The homogeneity of the SiO 2 -layer thickness was evaluated by ellipsometry measurements mapping the entire surface of the stainless still substrates (10.9 cm 2 ). The thickness map gives less than 10% difference between the thinnest and the thickest point.

According to the above given plasma conditions and deposition times the SiO 2 base layer is found 187.1 nm thick and the thickness of the SiO 2 cover layer is 9.1 nm. FTIR measurements were performed systematically in order to control the structural properties of the elaborated SiO 2 layers. The three characteristic peaks of Si-O-Si vibrations are detected on the FTIR spectrum of SiO 2 layers (Fig. 2). Indeed, the rocking at 457 cm -1 , the symmetrical stretching at 814 cm -1 and the asymmetrical stretching at 1063 cm -1 , including the shoulder at 1184 cm -1 are visible on the spectrum. The plasma silica FTIR spectrum differs from a thermal SiO 2 one for the slightly larger full width at half maximum (FWHM) of the asymmetrical stretching: 87 cm -1 for the plasma SiO 2 compared to 78 cm -1 for thermal SiO 2 . Although the used HMDSO precursor contains carbon, no carbon-traces are detectable on the FTIR spectrum. Presence of only SiO 2associated bonds means that the plasma process conditions and stability are met to grow a dielectric layer with composition and structure very near to that of thermal silica.

TEM imaging allows a direct observation of the elaborated structure (Fig. 3). The total thickness of the structure is measured from the cross section view (Fig. 3a). It is found 206 nm. The base SiO 2 layer is of 182 nm and the SiO 2 cover layer is of 9.5 nm. These values confirm the ones found by spectroscopic ellipsometry. Moreover, they are very close to the targeted layer thicknesses 190 nm and 10 nm, respectively. These results confirm the importance of control during the plasma elaboration, and validate the chosen model to analyze the ellipsometric data. The insert in Fig. 3a shows well that for very thin SiO 2 cover layers (< 10 nm) the sample surface topography presents slight waviness reproducing the shape of AgNPs, due to the conformal deposition. Better planarity of the surface is achieved when increasing the depth of the AgNPs plane or when decreasing the size of the nanoparticles. The plane view (Fig. 3b) of the sample shows that the AgNPs are well distributed in the plane and are far from percolation. Statistics on the AgNPs size are shown on Fig. 3c. The AgNP-distribution in size is bimodal with major axis centered at 4.1 nm and 17.8 nm for smaller and larger NPs. The AgNPs are of prolate spheroid shape, with eccentricity factor of 0.38 and 0.44 for the small and the large AgNPs respectively, in agreement with previous work [START_REF] Makasheva | Dielectric Engineering of Nanostructured Layers to Control the Transport of Injected Charges in Thin Dielectrics[END_REF][START_REF] Pugliara | Controlled elaboration of large-area plasmonic substrates by plasma process[END_REF]. Dark Field TEM imaging allowed us to conclude that the AgNPs are polycrystalline. No oxidation of the AgNPs was detected. It means that the SiO 2 cover layer protect the AgNPs plane from oxidation, thus their metallic character is preserved.

TEM imaging approves the quality of the plasma elaboration process. This verification is mandatory in order to accurately evaluate the impact of the nanostructuration, i.e. the nanometer scale parameters on the global electrical properties.

C. Current density under high electric field

As discussed before the dielectric charging leads to numerous complications over time, over use or under irradiation. Inclusion of AgNPs is a way to modify the electrical properties of a material, however preserving its dielectric behavior. The metallic nature of AgNPs increases the apparent relative permittivity of the nanocomposite. However, the AgNPs can also behave as artificial charge traps, which once occupied generate an internal electric field that ) FWHM opposes to the charges injection [START_REF] Milliere | Silver nanoparticles as a key feature of a plasma polymer composite layer in mitigation of charge injection into polyethylene under dc stress[END_REF]. Comparison of the current density as a function of the electric field for a SiO 2 layer and the presented above nanostructured layer is given in Fig. 4. The precision in the elaboration process allowed us to compare the global response of the SiO 2 dielectric layers that differ at nanoscale level. One of the main differences found in conductivity data is that the current density at low field is 4 orders of magnitude higher with AgNPs than without. Present results show that a single plane of AgNPs within a plasma silica matrix deeply impacts the transport, trapping and release of electrical charges, which depends among others on the distance between the AgNPs plane and the sample surface. The influence on the electrical properties of the AgNPs size, and the AgNPs plane depth will be detailed in coming works.

D. Modifications in the Total Electron Emission Yield

Under irradiation, the TEEY is globally higher for dielectrics, being around 3 for a bulk SiO 2 , compared to metals for which it is around 2. Fig. 5 shows the TEEY of the thin SiO 2 layer alone and the nanostructured one with embedded at 9.5 nm AgNPs. The atypical shape of the TEEY is due to the small thickness of the dielectric layer [START_REF] Rigoudy | Study of required conditions to limit the dielectric charging phenomenon when measuring the electron emission yield from thin dielectric layers[END_REF]. However, one can see that addition of a single plane of AgNPs lowers the TEEY curve toward metal-like sample at low energy (Fig. 5). The global shape remains the same, the specific behavior of the thin dielectric is preserved. With this structure, the lowering can be modulated by varying the size as well as the depth of the AgNPs plane.

IV. CONCLUSION

In this study we explored the capacity of plasma-based deposition method to elaborate thin nanocomposite dielectric layers consisting of silica matrix with embedded single plane of AgNPs at a controlled distance from the surface. The precise control down to nm on the dielectric layer thickness, as well as on the dimensions of the AgNPs allows evaluation of the impact of nanoscale structuration on the global electrical properties of thin dielectrics and opens the avenue for rational engineering of the dielectric properties of thin silica layers.

Fig 1 .

 1 Fig 1. OES spectra recorded during: (a) plasma silica deposition and (b) deposition of the single plane of AgNPs. The spectra are normalized to the intensity of Ar-line at 549.6 nm.

Fig 2 .

 2 Fig 2. FTIR spectrum for a sample of plasma SiO2. Insert: Deconvolution of the asymmetrical stretching peak and shoulder.

Fig 3 .

 3 Fig 3. (a) Cross-section view TEM image of the SiO2/AgNPs sample. Insert: enlargement of the near surface part. (b) Plane view TEM image of the SiO2/AgNPs sample, (c) Statistical distribution of the AgNPs' major axis.

Fig 4 .

 4 Fig 4. Current density vs. electric field of thin SiO2 layer (black triangle) and the SiO2/AgNPs nanostructure (red points). The open red points represent instabilities in the current density measurements.

Fig 5 .

 5 Fig 5. Total Electron Emission Yield of thin SiO2 layers with AgNPs (red points) and without (black triangle).

ACKNOWLEDGMENT Research supported by the program IDEX Actions Thématiques Stratégiques -ATS 2015 of the Université de Toulouse under project SEPHIR (2016-066-CIF-D-DRVD). The authors acknowledge support from the UMS Raymond Castaing of Université de Toulouse and thank Dr. A. Pugliara and Mr. L. Weingarten for the TEM observations.


L Boudou

L Berquez

G Teyssedre

Combined-PSD and LIMM Measurements for Traps Characterization in BoPP Thin Films

Understanding charge trapping phenomena in dielectrics requires a good description of the distribution in density, energetic levels, and if possible nature of the centers stabilizing charges. This topic is of great interest for many application fields as electronics or energy storage. Along this objective, fractionated-PSD (Photo-stimulated discharge) measurements, combined with space charge measurements using LIMM (Light Intensity Modulation Method) are developed in order to discriminate the origin of the different peaks observed in the PSD spectrum of Bi-oriented Polypropylene thin films.

I. INTRODUCTION

Charge trapping and transport phenomena in solid dielectrics are extensively studied since several decades. In particular, the study of charge accumulation is of great interest when considering its effects on electric field distortion and electrical properties modification. The space charge properties are closely linked to the characteristics of the traps (origin, density and depth in energy) present in the dielectric. Traps can be originated from physical or chemical defects [START_REF] Teyssedre | Charge transport modeling in insulating polymers: From molecular to macroscopic scale[END_REF], the corresponding depth in energy varying from low energies 0.1-0.4 eV (shallow traps) to energies higher than 1 eV (deep traps). Different complementary methods, essentially with thermal or optical excitation, can be used to probe the traps. Thermally stimulated current and luminescence are well adapted to probe the shallow traps distribution [START_REF] Van Turnhout | Thermally stimulated discharge of electrets[END_REF], whereas Photo-Stimulated Discharge (PSD) is mostly used for traps in the range 1 eV to more than 4 eV, depending on the light source characteristics [START_REF] Mellinger | Photostimulated discharge in electret polymers: an alternative approach for investigating deep traps[END_REF][START_REF] Ma | Analysis and interpretation of photo-stimulated discharge spectrum for polypropylene films under different electric and geometrical conditions[END_REF]. Recombination induced luminescence can also give information on deep traps acting as recombination centres [START_REF] Teyssedre | Spectral analysis of optical emission due to isothermal charge recombination in polyolefins[END_REF]. Finally, space charge techniques can give the in-depth charge density profile and allow to follow the kinetics of the charge release [START_REF] Dissado | Space charge Injection and extraction in high divergent fields[END_REF].

The main advantage of PSD method is that it can be applied without any thermal effect that can affect the material microstructure and the charge stability. Nonetheless the interpretation of the PSD spectra is not straightforward. Actually, the discharge current can originate from different mechanisms: space charge detrapping, photoconduction effect, photo-emission from the electrode or from the dielectric material surface under test. It can also be perturbed by dipolar relaxations during measurement. In order to discriminate the origin of the different PSD peaks recorded during current measurements, the influence of different parameters can be tested: the geometry of the electrode facing the scanning light (full electrode, semi-transparent electrode, finger-like shape electrode), the nature of the metallic electrodes or the application of a small voltage during PSD measurements [START_REF] Ma | Analysis and interpretation of photo-stimulated discharge spectrum for polypropylene films under different electric and geometrical conditions[END_REF].

Space charge measurement can also give important information on the in-depth trapped charge distribution of previously charged sample, and allows to follow the space charge dynamics before and after light irradiation. For example, Mellinger et al [START_REF] Mellinger | Photostimulated discharge in electret polymers: an alternative approach for investigating deep traps[END_REF] studied the space charge profiles, using piezoelectrically generated pressure step (PPS) method, of a gold-coated PTFE film corona-charged at 200 °C before and after irradiation at 240 nm. After irradiation at 240 nm, part of the charge was removed from the bulk of the polymer. The authors also point out the possibility of detrapping of near-surface charges, without experimental validation, because the PPS method cannot distinguish between near-surface charges and influence charges on the metal electrodes.

In this work, combined LIMM/PSD measurements are performed using the same apparatus offering the advantage to switch from one measurement to another without any sample manipulation. The first results obtained on Bi-oriented Polypropylene (BoPP) thin film are presented and discussed.

II. METHODOLOGY

A. Combined PSD/LIMM apparatus

Fig. 1 shows the diagram of the circuit for combined PSD/LIMM apparatus developed for this study. The K component allows to switch from PSD to LIMM and reciprocally, without any sample manipulation.

We used a fingered electrode as front electrode (facing the light during PSD/LIMM measurements) and full electrode as back electrode, connected either to an ammeter for PSD measurements or to a lock-in amplifier for LIMM measurements. For PSD measurement, the monochromatic light (from 600 nm to 200 nm) is produced by a 150 W Xenon arc lamp associated with a monochromator equipped with a 300 nmblaze grating (1200 g/mm) as detailed in a previous communication [START_REF] Boudou | Photostimulated discharge current measurements on biaxially oriented polypropylene thin films (BOPP)[END_REF]. The incident irradiation angle is normal to the sample surface and the lighted area is about 8 cm 2 .

For LIMM measurement, an intensity-modulated laser diode is used to heat locally the sample surface. The laser hits the sample surface on a "finger" of the fingered electrode (spot size about 2 mm in diameter) with an incident angle of approximatively 45°. The interaction between the produced thermal wave and the trapped charge allows to extract space charge density profiles from measured pyroelectric current, temperature profile modeling and mathematical treatment. The basic principle of the method, the theory and details on the setup can be found in [START_REF] Velazquez-Salazar | Towards space charge measurements by (F)LIMM under DC electric field[END_REF].

The front electrode is grounded during the PSD and LIMM measurement and all the experiments were performed in air atmosphere and at room temperature.

B. Material

The sample under study is a 18 μm-thick BoPP, with rough surfaces to promote impregnation of capacitors. It was provided by Kopafilm, Germany. Test samples are discs of 8 cm in diameter, with gold electrodes sputtered on each face. For the electrode facing to the scanning light or to the laser diode (termed as the front electrode), we used a fingered electrode having a diameter of 4 cm with a thickness of 50 nm, and fingers width of 4 mm [START_REF] Boudou | Photostimulated discharge current measurements on biaxially oriented polypropylene thin films (BOPP)[END_REF]. For the electrode connected to the ammeter or to the LIMM detection system (termed back electrode), we used a circular electrode of 4 cm diameter and 50 nm thickness.

C. Experimental protocol

The experimental protocol used is represented in Fig. 2. Before the polarization step, blank PSD and LIMM spectra are recorded (PSD0 and LIMM0). Then the sample is polarized at a constant voltage of 3 kV, corresponding to an electric field of 167 kV/mm, during 30 min. After the polarization step, a first partial PSD measurement is performed by light scanning from long wavelength (600 nm) to an intermediate wavelength (350 nm), i.e. in-between the two PSD peaks previously reported for BoPP (400 nm and 220 nm) [START_REF] Boudou | Photostimulated discharge current measurements on biaxially oriented polypropylene thin films (BOPP)[END_REF]. The second PSD measurement is performed by scanning over the entire wavelength range from 600 to 200 nm. LIMM recordings are performed at different periods during the protocol to follow the space charge evolution.

III. RESULTS AND DISCUSSIONS

A. PSD/LIMM spectra

The PSD spectrum recorded on the fresh sample is not totally flat. Two small current peaks are detected at 395 nm and 225 nm, with amplitude lower than 1 pA (Fig. 3). The LIMM measurement performed before the voltage application (LIMM0 in Fig. 2) shows very low currents in the tested frequency range, corresponding to a negligible density of trapped space charge.

On Fig. 4 are represented the PSD spectra obtained from partial (PSD1) and full (PSD2) light scanning. A first peak with a maximum located at 397 nm is recorded in both cases. This peak is relatively broad (spreading from 480 nm to 360 nm with a FWHM of 50 nm) and does not correspond to known absorption band in BoPP investigated herein [START_REF] Qiao | Fieldand electron beaminduced luminescence phenomena in polypropylene thin films[END_REF]. Either it can correspond to additives that could efficiently trap charges, or possibly to localized states formed under the effect of space charge [START_REF] Duke | Optical absorption due to space-charge-induced localized states[END_REF].

A second peak is recorded in the wavelength range 260-200 nm. The mathematical deconvolution (insert of Fig. 4) of this peak shows two contributions, located at 229 and 204 nm, in agreement with earlier work [START_REF] Mellinger | Photostimulated discharge in electret polymers: an alternative approach for investigating deep traps[END_REF]. These two contributions have to be correlated with the peaks observed in the absorption spectrum [START_REF] Qiao | Fieldand electron beaminduced luminescence phenomena in polypropylene thin films[END_REF] of BoPP, that shows an absorption maximum at 200 nm and a first shoulder at 230 nm, typical of π→π* transitions in C=C double bonds. A second shoulder at 280 nm is rather due to the n→π* transition in carbonyl compounds. These two wavelengths (230 and 280 nm) are those providing maximum photoluminescence in polyolefins [START_REF] Qiao | Fieldand electron beaminduced luminescence phenomena in polypropylene thin films[END_REF]. Fig. 5 shows the charge density profiles obtained from LIMM measurements performed before PSD1, after PSD1 and after PSD2. After the polarization step and before PSD1 measurement, positive charges are accumulated near the Au/BoPP interface (anode), corresponding to homocharges. Furthermore, a negative charge peak appears within the bulk (between 2 and 7 m) but we don't know at this stage if it corresponds to real negative trapped charges or to an artifact due to the LIMM signal processing. Finally, probing the trapped charges near the back electrode is difficult using LIMM. One can assume that homocharges, that is to say negative charges, can be accumulated at the cathode during the polarization step. If we consider the first 2 m, corresponding to a depth range where the LIMM sensibility is highest, charge density is maximum after voltage application, then it decreases after the PSD1 and PSD2 scans are performed. Even though the difference between LIMM1 and LMM2 profiles shown on Fig 5 is weak, the trend is a decrease of the charge density when sample is lighted from 600 to 350 nm (PSD1) In addition, the LIMM3 profile performed after complete scanning (PSD2) also shows a decrease of charge density compared to LIMM2 profile.

According to these first combined PSD/LIMM results, charge detrapping can be one of the mechanisms responsible for the 395 nm-peak recorded during PSD measurement.

Concerning the double PSD peak located in the 260-200 nm range, optical detrapping is not excluded, but other phenomena like photo-emission from the metallic electrode can be at play. Actually, by changing the metal nature of the front electrode, Ma et al [START_REF] Ma | Analysis and interpretation of photo-stimulated discharge spectrum for polypropylene films under different electric and geometrical conditions[END_REF] have shown a direct correlation between the work function of the metal and the cut-off wavelength of the peak observed in the 200-300 nm range for BoPP. The authors suggest that this peak results from electron emission via photoelectric effect from the front finger electrode when a local strong electric field exists.

Furthermore, during PSD measurement, charges can be detrapped and can move over short distances, depending on the internal electric field. The charge detrapping and transport can lead to a recordable current during PSD measurements, but the changes induced on the charge density profile are not necessary visible on LIMM results, because of the sensitivity and resolution of the method limited to 1 m. . Charge density profiles extracted from LIMM measurements performed at different time of the experimental protocol described in Fig. 2 One way to increase the transport of detrapped charge is to apply a small driving voltage during PSD measurement. Positive voltage (on top electrode) should drive positive charges in the bulk direction. Negative voltage should drive positive charges from bulk to the metal/polymer interface.

Another point to consider concerns the fingered electrode used as front electrode. It can first induce a field enhancement du to edge effect and as a consequence, charge injection and trapping at the edge of the fingers of the top electrode can be increased [START_REF] Ma | Analysis and interpretation of photo-stimulated discharge spectrum for polypropylene films under different electric and geometrical conditions[END_REF]. Besides, the fingered electrode plays a great role when considering the area (or volume) of the polymer material probed when LIMM and PSD are performed. A gold layer of 50 nm of thickness cannot be considered as transparent to the light: the transmission spectrum in the visible-UV range has to be considered for power generation and for qualitative analysis of the PSD spectrum. Another difficulty is that LIMM measurement probe the space charge trapped under the gold metallization and in a small area (≈3 mm 2 ), corresponding to the spot size of the laser diode. Therefore, PSD and LIMM measurements do not probe exactly the same part of the polymer surface of the BoPP sample.

In conclusion, the interpretation of the origin of the different PSD peaks of BoPP is not trivial. The feasibility of fractionated PSD combined with LIMM has been demonstrated, which will help in the interpretation of the peak wavelength positions, but further investigations are needed.

B. PSD transient current

Measuring the current decay as a function of time under fixed wavelength is another way of analyzing the discharge phenomenon. Fig. 6 shows the transient photo-discharge current recorded at a fixed wavelength of 225 nm during 5 h, by step of 1 h, on a BoPP sample previously polarized at 3 kV for 30 min. LIMM measurements have been performed at different times during the decay and the obtained space charge density profiles are plotted in Fig. 7. During LIMM measurements, which last 10 min each for scanning frequencies from 10 Hz to 10 kHz, the 225 nm-light irradiation is cut. The first space charge profile (LIMM0) shows positive charge accumulation near the anode, in agreement with results presented in section III.A. The PSD current quickly decreases from 60 pA to around 12 pA during the first minute, before a much slower PSD current decrease is observed (cf. insert of Fig. 6). LIMM1 performed after 1 h of discharge shows a small decrease of the trapped charge at the anode (insert of Fig. 7) compared to LIMM0 (before PSD). After LIMM1 and when 225 nm-light irradiation is turned on again, the recorded PSD current is slightly higher (12.3 pA) when compared to the current recorded after 1 h of discharge (10.5 pA) (Fig. 6). The next space charge profiles recorded and numbered from 2 to 5 in Fig. 7 show a slight decrease of the charge density at the anode/BoPP interface. The corresponding PSD current recorded during 1 h between each LIMM measurement shows a slow decrease and a non-monotonous shape of the decay.

One can assume that discharge current under long term illumination in the UV-range is due to several mechanisms in competition: photoemission from the electrode, detrapping of charge, etc. However, LIMM profiles are very similar and the difference of charge density at the anode/BoPP interface is weak. The relation between the light power reaching the material, the light absorbed locally and the efficiency of the light to detrap charges is far from being at reach.

The obtained results need to be confirmed by further investigation on samples polarized at higher field or during longer time, in order to have higher density of trapped charge, or charged by other means as electron gun. The transient discharge behavior has also to be investigated at a fixed wavelength of 395 nm, 229 nm and 204 nm, corresponding to the maximum peak position shown in Fig. 4. Work has also to be achieved on LIMM improvement in order to get information on trapped charge at the rear electrode.

IV. CONCLUSION

This work is a first attempt to perform combined LIMM/PSD measurements using the same apparatus offering the advantage to switch from one measurement to another without any sample manipulation. The results obtained on a 18 m-thick BoPP film show three PSD peaks (395, 229 and 204 nm) and positive charge accumulation at the anode/BoPP interface is revealed by LIMM measurements. These results show also the difficulties and the necessary improvement of the combined LIMM/PSD method in order to correctly interpret the PSD spectra and space charge dynamics in BoPP thin films.

Fig. 1 .

 1 Fig. 1. Combined PSD/LIMM circuit diagram

Fig. 2 .

 2 Fig. 2. Experimental protocol

Fig. 3 .

 3 Fig. 3. Blank PSD spectrum recorded before polarization

Fig. 4 .

 4 Fig. 4. PSD1 (red line) and PSD2 (black line) spectra after charging under 3kV for 30 min. Insert: PSD2 spectrum in the range 280-200 nm showing the mathematical deconvolution

Fig. 5

 5 Fig.5. Charge density profiles extracted from LIMM measurements performed at different time of the experimental protocol described in Fig.2

Fig. 6 .

 6 Fig. 6. Transient PSD current under continuous illumination at 225 nm.Sample previously polarized under 3 kV for 30 min Insert: zoom on the first 10 min of the transient current.

Fig. 7 .

 7 Fig. 7. Charge density profiles extracted from LIMM measurements. Insert: Evolution of charge density at the cathode/BoPP interface versus time

ACKNOWLEDGMENT

The research was carried with support from the IDEX Actions Thématiques Stratégiques -ATS 2015 of University of Toulouse, projet SEPHIR (2016-066-CIF-D-DRVD).


Space Charge Measurement by PEA Method on an Aeronautical Cable Amin Benyoucef [START_REF] Cella | Partial discharges measurements at the constituent's level of aerospace power electronics converters[END_REF][START_REF] Yasuda | Measurement of charge distributions in coaxial cable using the pulsed electroacoustic method[END_REF], Laurent Berquez (1), Gilbert Teyssedre (1), Eddy Aubert [START_REF] Yasuda | Measurement of charge distributions in coaxial cable using the pulsed electroacoustic method[END_REF][START_REF] Hozumi | Space charge distribution measurement in a long size XLPE cable using the pulsed electroacoustic method[END_REF] (1) Laplace, Université Paul Sabatier Bat 3R3, 118 Route de Narbonne 31062 Toulouse, France (2) IRT Saint Exupery, 3 Rue Tarfayan 31405 Toulouse, France (3) Safran Electrical and Power, 1 Rue Louis Blériot, 31700 Blagnac, France Abstract-Technological evolutions in aircraft networks let us foresee an increase of the DC voltage up to several kilovolts while keeping minimal insulation thickness. This trend results in a strong increase of the electric field in the insulations and space charge can become an issue for the reliability of the systems. Indeed, space charge accumulation modifies the electric field distribution and may lead to a premature aging of the dielectric. Besides, the field distribution in complex structures such as multilayers cannot be anticipated under DC as easily as under AC because conductivity is more difficult to estimate and varies much more than permittivity. In this paper, we present the first results obtained on aeronautical cables. For this, it was necessary to develop a deconvolution strategy adapted to aeronautical cables to obtain electric field and charge density profiles in a part of the insulation thickness.

I. INTRODUCTION

The increase in the power of electrical systems of future aircraft goes with an increase of the AC and DC voltages of power networks. For DC systems, the phenomenon of partial discharges, well-known in case of AC voltage, is actively studied [START_REF] Cella | Partial discharges measurements at the constituent's level of aerospace power electronics converters[END_REF] with much less consequences as under AC as the frequency of occurrence is strongly reduced. However, other effects, such as space charge accumulation, are to be considered. For the insulation of terrestrial or submarine HVDC energy cable systems, with voltages of several hundreds of kV, the problem of space charge is known and deeply investigated for long.

Even if the voltage levels in aeronautics are much lower, an increase of the DC buses beyond 3 kV is foreseen while the insulation thickness is minimized to optimize the mass of the aeronautical network. This means that a significant increase in the electric field in the insulations could occur, possibly to levels where space charges could be accumulated and the field threshold for non-linear conduction be reached. Therefore, space charge effects must be investigated, mastered and taken into account in future technologies.

To meet these challenges, our objective is to develop a test bench to measure space charge distributions on aeronautical cables and to estimate the evolution of the electric field distribution in particular in multilayer arrangements of cable insulation and to identify the stress regimes where these phenomena occur.

II. PEA METHOD APPLIED TO AN AERONAUTICAL CABLE

A. PEA method

For the first time in the 90's, a PEA system was introduced for the measurement of the space charge on cable for high voltage energy transmission [START_REF] Yasuda | Measurement of charge distributions in coaxial cable using the pulsed electroacoustic method[END_REF][START_REF] Hozumi | Space charge distribution measurement in a long size XLPE cable using the pulsed electroacoustic method[END_REF]. The method consists in applying a voltage pulse to the sample, producing oscillation of the charges around their equilibrium positions, thus generating acoustic waves. The acoustic waves are then converted into an electrical signal through a piezoelectric sensor and this signal is amplified and digitally processed by a deconvolution algorithm (Fig. 1).

Resorting to PEA device for cables, various variants were produced, depending on the way the voltage pulse is applied: through a capacitor connected to cable conductor, or on a part of the cable sample having its outer sheath separated from that at the measurement point, or even with pulse application directly on the measurement cell [START_REF] Rizzo | Review of the PEA Method for Space Charge Measurements on HVDC Cables and Mini-Cables[END_REF]. As often in aeronautic domain, the samples had no outer sheath. Therefore, the most convenient way to apply the voltage pulses was with using a different cable as decoupling capacitor.

B. Aeronautical cable

The aeronautical cable studied here is a typical example of an aircraft cable section in which the insulation is wound on the multi-stranded copper conductor of 3.47 mm radius, designed to avoid skin effect in the conductor. A polyimide (PI) layer is usually wound directly onto the conductor and is used because of its high mechanical strength and its dielectric insulation and thermal properties. The outer tape of the cable is a polytetrafluoroethylene (PTFE) which limits fire and arc tracking problems. The insulation system has a total thickness of approximately 360 µm (Fig. 2.a). In Fig. 2.b, we can see that on each face of the PI ribbons are layers of fluorinated ethylene propylene (FEP) with a thickness of about 16 µm. This FEP layer is introduced to improve the adhesion between PI layers and between PI and PTFE layers.

C. Protocol and calibration

The protocol used for present PEA measurements is shown in Fig. 3. It consists of two steps, a first step called calibration where a voltage of 1 kV is applied which is to be used to evaluate the transfer function of the system [START_REF] Vissouvanadin | Deconvolution techniques for space charge recovery using pulsed electroacoustic method in coaxial geometry[END_REF]. Then in a second step, so-called test phase, a voltage of 9 kV is applied for 120 min with a relaxation phase lasting 120 min. The step is repeated with a reverse possibility (-9 kV for 120 min followed by a relaxation phase of 120 min). Fig. 4 shows a PEA signal during the calibration phase. When a voltage of 1 kV is applied, a first peak appears at around 75 ns which corresponds to the capacitive charge at the ground electrode. From 215 ns on, a multitude of peaks are observed very close together. This is presumably due to the difference in permittivity between the PI and FEP layers (ε r =2.4 and 2.1 for PI and FEP, respectively), which generates a signal at the interface of the two materials [START_REF] Bodega | Space charge measurements on multidielectrics by means of the Pulsed Electro Acoustic method[END_REF]. Furthermore, due to the low thicknesses of PI and FEP, multiple reflections are picked up by the piezoelectric sensor [START_REF] Holé | Influence of divergent electric fields on space-charge distribution measurements by elastic methods[END_REF]. Further, with considering an obvious dispersion of the signal, then the different signals merge and overlap as the signal corresponding to the inner electrode (cable conductor) is approached. Ultimately, it is not possible to distinguish the different signals necessary to evaluate the attenuation and dispersion factors during the deconvolution step. Indeed, in previous PEA measurements on samples with cylindrical geometry, the attenuation and dispersion spectra were determined considering that the acoustic signal generated at the inner interface (at cable conductor) due to capacitive charges is modified during its propagation. As in the present case the signal due to internal charges is spoiled, the method for space charge recovery had to be modified as depicted below.

D. Deconvolution method

In the calibration signal, an acoustic response is produced due to the difference in permittivity between PI and FEP. It leads to complexity with the response due to the capacitive charge at the conductor as well as the multiple reflections generated. Consequently, the information on attenuation and dispersion can not be retrieved from this calibration signal. The strategy addressed to be able to proceed to deconvolution was to characterize the outer layer of PTFE insulation for acoustic properties. As shown in Fig. 5, a calibration measurement is first performed on a flat PTFE sample without charges to recover attenuation and dispersion information. Next, we perform a measurement on an aeronautical cable in the calibration stage to allow us to have the transfer function of the PEA cell. Thus, by multiplying the attenuation and dispersion information with the transfer function of the cell, we can have the transfer function of the system. By carrying out the convolution product between a PEA signal carried out on an aeronautical cable and the transfer function of the system, the evolution of the charge density and the evolution of the electric field on the outer layer of the PTFE insulation can be followed.

III. RESULTS

Preliminary space charge measurements have been achieved following the protocol shown in Fig. 3. Our deconvolution method allows us to evaluate the electric field as well as the charge density in the PTFE insulation layer, which represents the major part of the insulation thickness (see Fig. 2). In Fig. 6.a we have represented a mapping of the electric field over the depth of the outer layer of cable insulation (232 µm in thickness) along the established protocol duration. Fig. 6.b shows the variation of the electric field at the lower electrode as a function of the protocol duration. The geometrical electric field calculated at the outer radius, taking into account the multilayer dielectric, for a voltage of 9 kV is 24 kV/mm. During the entire application of the 9 kV voltage, the measured electric field decreases from about 30 kV/mm to 23 kV/mm. Moreover, during the relaxation phase a residual electric field of opposite sign appears and decreases over time.

When a voltage of -9 kV is applied, the field distribution is symmetrical: it varies from about -30 kV/mm to -23 kV/mm as a function of time. During the relaxation phase, an electric field of the opposite sign also appears and decreases in the course of time. The decrease of the electric field, and the appearance of an electric field of opposite sign during the relaxation phase reflect the formation of image charges [START_REF] Roulet | Image charges revisited: Beyond classical electrostatics[END_REF].

The appearance of image charge at the outer electrode of the cable is due to the presence of space charge in the cable insulation volume. In Fig. 7.a the charge density profile at 1 kV during the calibration step is represented. It gives an illustration on the way a capacitive charge is pictured by the system used. The image charge peak width at half-maximum is about 20 µm.

In Fig. 7.b, we have represented the charge density profiles under an applied voltage of 9 kV at different application times. At the lower electrode, we notice a decrease of the image charge density in the course of time. We believe that the origin of this space charge is related to the Maxell-Wagner effect [START_REF] Vu | Correlating conductivity and space charge measurements in multidielectrics under various electrical and thermal stresses[END_REF]: because of the presence of multilayers of different materials, a field redistribution may occur due to the difference in electrical conductivity in the different materials. The formation of these space charges corresponds to the redistribution of the electric field over the entire thickness of the insulation. In another respect, the oscillations found near the outer electrode are probably related to the fact that the insulation is wrapped: an imperfect propagation of acoustic waves when crossing the different PTFE layers may induce spurious effects on the signal. Preliminary space charge measurements results obtained by PEA method on aeronautical cable were presented. These measurements revealed that the acoustic signals generated by the difference in permittivity between the PI and FEP multilayers, that issued from multiple reflections at the internal interfaces, and the signal generated by the influence charges on the conductor, are merging. As a result, the origins of the different signals near the cable conductor can no longer be distinguished and the information on the dispersion and attenuation cannot be recovered. This is why we concentrated on the outer layer of insulation.

For this purpose, we carried out measurements on PTFE plane samples to recover the attenuation and dispersion information. We then carried out measurements on aeronautical cables and then traced back to the electric field information. During our tests at 9 kV and -9 kV, we noted a decrease in time of the electric field at the outer part of the cable presumably due to the Maxwell-Wagner effect in the multilayer insulation of the cable. This results in a redistribution of the electric field.

Fig. 1 .

 1 Fig. 1. Schematic diagram PEA method

Fig. 3 .

 3 Fig. 3. Calibration and test protocol

Fig. 5 .

 5 Fig. 5. Deconvolution synoptic for a multilayer insulation cable

Fig. 7 .

 7 Fig. 7. Charge density profile during (a) the calibration step, and (b) when applying a voltage of 9 kV for different times


Y Kemari

G Teyssèdre

A Mekhaldi

M Teguar

Dielectric Properties and -Relaxation in Crosslinked Polyethylene: Effect of Thermal Aging

Keywords: cross-linked polyethylene, dielectric relaxation, thermal aging, differential scanning calorimetry, permittivity

The present paper aims at investigating the influence of thermal aging on the crystallinity, the permittivity and relaxation of cross-linked polyethylene (XLPE) used in power cables. For this, accelerated thermal aging is performed on circular XLPE samples for 5040 h (210 days) at 80, 100 and 120 °C and for 1512 h (63 days) at 140 °C. Changes in the crystallinity of XLPE are assessed by Differential Scanning Calorimetry (DSC). Frequency domain dielectric spectroscopy is used to analyze the variations of real and imaginary parts of the permittivity, before and after aging. Moreover, -relaxation is investigated through loss factor ( r ") analysis in the temperature range from 0 to 100 °C at several different frequencies 10, 10 2 , 10 3 and 10 4 Hz. We show that dielectric spectroscopy is not the best suite method to reveal the earlier steps of aging.

I. INTRODUCTION

With the advent of synthetic insulating materials in the middle of the 20 th century, cross-linked polyethylene (XLPE) continues to be extensively used in underground cables manufacturing due to its excellent dielectric properties [START_REF] Metwally | The evolution of medium voltage power cables[END_REF][START_REF] Teyssedre | Advances in high-field insulating polymeric materials over the past 50 years[END_REF]. Nevertheless, cables could be exposed to severe thermal constraints during their operating conditions. Consequently, insulation may suffer degradation and structural changes due to thermal oxidation [START_REF] Teyssedre | Semi-quantitative analysis of photoluminescence in thermoelectrically aged cables: I-identification of optical signatures[END_REF]. In a previous work [START_REF] Kemari | Correlations between structural changes and dielectric behavior of thermally aged XLPE[END_REF], it has been found that at the beginning of aging, the mechanisms of degradation are inhibited by the antioxidant introduced into the material. Also, post-cross-linking reactions could occur leading to a decrease of orientation polarization in the material due to the recrystallization of some parts in XLPE. On the other hand, the chemical degradation of XLPE during prolonged thermal stress consists mainly in the consumption of the antioxidant followed by the generation of polar groups such as carbonyl, peroxide and hydroxyl groups.

In polyethylene material, three principal relaxation modes can be observed. The α relaxation is generally observed in crystalline polymers and it is usually related to the motion of chain units within the crystalline portion [START_REF] Suljovrujic | Structural changes and dielectric relaxation behavior of uniaxially oriented high-density polyethylene[END_REF]. The α-process was reported as dielectrically active due to the reorientation of carbonyl groups in the chains in the crystalline phase [START_REF] Ashcraft | A dielectric study of molecular relaxation in oxidized and chlorinated polyethylenes[END_REF]. However, space charge processes seem to be at play as well. Therefore, it cannot be easily used as an aging marker. In contrast, β and γ processes have their origins in the amorphous phase [START_REF] Suljovrujic | Structural changes and dielectric relaxation behavior of uniaxially oriented high-density polyethylene[END_REF].

Several studies were performed to investigate dielectric relaxation modes of polyethylene in the perspective of aging diagnosis development. D. Kostoski et al [START_REF] Kostoski | Effects of charge trapping in gamma irradiated and accelerated aged low-density polyethylene[END_REF] studied the effects of gamma irradiation in air and accelerated aging on low-density polyethylene (LDPE) using thermally stimulated discharge current (TSDC) measurements. Results show the increase of the intensities of the β and γ dielectric relaxations and the shift of their positions due to the influence of radiation dose and accelerated aging. Moreover, E. Suljovrujić et al [START_REF] Suljovrujić | Dielectric relaxation study of gamma irradiated oriented low-density polyethylene[END_REF] showed that the decrease of the β relaxation intensity can be related to a decrease of the amorphous fraction in LDPE. Besides, irradiation and accelerated aging induced a decrease of the activation energy of this relaxation mode.

In the current work, we focus on β-relaxation of XLPE and its changes produced due to a different aging method, which is accelerated thermal aging. For this, XLPE samples have been subjected to aging for 5040 h (210 days) at 80, 100 and 120 °C and for 1512 h (63 days) at 140 °C. The evolution of the loss factor,  r ", associated with the β-relaxation is proposed as a parameter to reflect aging degree. Therefore, frequency domain dielectric spectroscopy measurements have been performed. In fact, the variations of phase transitions of XLPE due to thermal aging have been monitored in our previous work [START_REF] Kemari | Correlations between structural changes and dielectric behavior of thermally aged XLPE[END_REF] using differential scanning calorimetry (DSC). Some results will be discussed in this study in order to show the correlation between the crystalline structure and the dielectric relaxation. This will also lead to a better understanding of aging and degradation mechanisms and to assess the relevance of using the changes in the β relaxation process as an aging marker.

II. EXPERIMENTAL SECTION

A. Material and Samples Preparation

The studied material is used for extruded ac power cables of nominal voltages U 0 /U = 18/30 kV. First, the pellets of crosslinkable resin (LE4201R from Borealis) are provided with 2 wt. % of di-cumyl peroxide (DCP) as cross-linking agent. The material contains also di-tert-butyl peroxide and Irganox antioxidant. For manufacturing the XLPE samples, the pellets were preheated for 5 minutes at 125 °C then molded and crosslinked at 185 °C under a pressure of 300 MPa using a heating molding press. These conditions (300 MPa and 185 °C) are maintained for about 10 minutes to ensure complete crosslinking of the material. At the end, the press is cooled with water with maintaining always the same pressure.

Finally, XLPE plates were cut into circular shape samples with a diameter of about 46 mm and a thickness of (2 ± 0.2) mm. Samples were cut using a cutting press with a circular cutting die tool.

B. Thermal Aging Experiments

Accelerated thermal aging tests were carried out by arranging the XLPE circular samples vertically inside aircirculating ovens for 5040 hours (210 days) at 80, 100 and 120 °C. Moreover, since the degradation process is very fast at 140 °C, samples have been aged for a maximum duration of 1512 hours (63 days).

C. Frequency Domain Dielectric Spectroscopy

Real and imaginary parts of the permittivity were measured using Novocontrol Alfa-A High Performance Analyzer dielectric spectrometer with a ZGS active sample cell. The system can control the temperature under N 2 atmosphere by a Quatro cryosystem within ± 0.1 °C. Prior to the dielectric measurements, gold electrodes were deposited onto both surfaces of XLPE circular samples by sputtering. The diameter and the thickness of the sputtered electrodes are respectively 35 mm and 30 nm. Real and imaginary permittivity,  r ' and  r " were analyzed before and after aging in the temperature range from 0 to 100 °C at 10, 10 2 , 10 3 and 10 4 Hz.

III. RESULTS

A. Dielectric constant  r '

Dielectric constant values are presented in Fig. 1 versus temperature at different frequencies. From the obtained results, before and after aging at temperatures of 80 and 100 °C,  r ' remains practically constant as the temperature increases from 0 to 85 °C. Then,  r ' starts to decrease slowly from 85 to 100 °C. In fact, at lower temperature, the segmental motion of the chain is practically null and the dielectric constant is not affected. However, at sufficiently higher temperature, the dielectric constant is again reduced due to strong thermal motion which disturbs the orientational polarization [START_REF] Zulkifli | Polymer dielectric materials[END_REF]. Also, a slight decrease of  r ' was observed in the case of aging at 100 °C at all frequencies from 0 to 85 °C. As already discussed in our pervious report [START_REF] Kemari | Correlations between structural changes and dielectric behavior of thermally aged XLPE[END_REF], the post-cross-linking could optimize the crystallinity of the aged samples at 100 °C, leading to make the orientation of dipoles more difficult. Furthermore, it has been found also that the crystallinity of XLPE slightly increased from 39.8 % before aging to 42.4 % after aging at 100 °C for 5040 h, as given in Table 1.

In contrast, at aging temperatures of 120 and 140 °C,  r ' rises significantly compared to other aging temperatures. Moreover, thermal agitation could be enhanced from 0 to about 40 °C since the intermolecular forces between polymer chains are broken [START_REF] Zulkifli | Polymer dielectric materials[END_REF]. Consequently, the dielectric constant increases with the temperature and then it decreases. The observed threshold temperature (40 °C) is lower than that at other aging conditions. This may be related to the shift of the XLPE melting peak toward lower temperatures after aging at 120 and 140 °C [START_REF] Kemari | Correlations between structural changes and dielectric behavior of thermally aged XLPE[END_REF].

TABLE I RESULTS OBTAINED FROM DSC MEASUREMENTS (SECOND SCAN)

Tm is the melting point, ∆Hm is the enthalpy of fusion and c is the crystallinity rate

B. Imaginary part of permittivity  r " and β-relaxation

In this section we have chosen to study the evolution of  r " in order to investigate the influence of aging on the βrelaxation of XLPE. Fig. 2 presents  r " curves versus temperature at different frequencies before and after aging. The results before and after aging at 80 and 100 °C show that the imaginary part of the permittivity remains nearly temperature-independent (from 0 to 100 °C) at 1 and 10 kHz and the relaxation peak could not be noticed. As the frequency decreases, the curves are shifted to lower temperatures and a minimum appears at 80 °C for the frequency 10 Hz. Also,  r " decreases with the decrease of the frequency in the case of aging at 100 °C which is most likely a consequence of reduced mobility of the chain segments due to the enhancement of the crystallinity [START_REF] Suljovrujić | Dielectric relaxation study of gamma irradiated oriented low-density polyethylene[END_REF].

When the degradation becomes more advanced, at 120 and 140 °C,  r " drastically rises and the peak attributed to the relaxation can be easily observed at lower temperatures. At these aging conditions, chain scission by thermal oxidation is the dominant process and the crystallinity decreases as shown in Table 1. As a result, the amorphous phase becomes more important which makes the macromolecule chains more flexible to interact under the polarization effect. This result remains very important indicating how dielectric losses in XLPE cables could be very significant, even below 40 °C, depending on the frequency. In accordance with previous works [START_REF] Kostoski | Effects of charge trapping in gamma irradiated and accelerated aged low-density polyethylene[END_REF], [START_REF] Suljovrujić | Dielectric relaxation study of gamma irradiated oriented low-density polyethylene[END_REF] and [START_REF] Frubing | Complete relaxation map of polyethylene: filler-induced chemical modifications as dielectric probes[END_REF],  shifts toward lower temperature with decreasing frequency, however its intensity is not significantly affected. As shown in Fig. 2c and2d, a dielectric loss peak occurs at a higher temperature than the β process. The beginning of the a-relaxation peak can be well observed at 10 and 100 Hz. Furthermore, the frequency-temperature shift of the β-relaxation could be fitted by the Arrhenius equation (1) in the cases of aging at 120 and 140 °C:

) exp( 0 max kT E f f   (1)
where E is the activation energy, k is Boltzmann's constant, T is the absolute temperature and f 0 is a constant. The results are plotted in Fig. 3. On the basis of the imaginary part of permittivity, the activation energy for sample aged at 120 °C is 2.22 eV compared to 1.93 eV for aging at 140 °C. A large change in the dielectric behavior of XLPE was observed when changing the aging temperature from 100 to 120 °C. One of the reasons could be the fact that the temperatures are at both sides of the melting point, hence the oxidation kinetics would be slowed down in the partially crystalline state. To bring some light on this point, we have plotted in Fig. 4 the dielectric loss factor measured at 20 °C as a function of aging time at 120 °C. Very clearly, the losses increase abruptly between the last two aging times and so does the dielectric permittivity [START_REF] Kemari | Correlations between structural changes and dielectric behavior of thermally aged XLPE[END_REF]. Similar trends were observed at 140 °C indicating that a critical degradation state exists in the material. Comparatively, the carbonyl index (Fig. 4) follows a more progressive rise before a final step [START_REF] Kemari | Correlations between structural changes and dielectric behavior of thermally aged XLPE[END_REF]. The abrupt changes in properties can reasonably be ascribed to the end step of antioxidant consumption and the subsequent attack of the polymer chains by oxidation.

Regarding activation energies, Geng et al. [START_REF] Geng | Influence of thermal aging on AC leakage current in XLPE insulation[END_REF] reported on the change in high temperature dielectric losses in XLPE aged in the same temperature range but with much shorter times (up to 240 h) than herein. Activation energies of the same order as here were obtained, without a clear trend vs. aging degree. Conduction processes were advocated related to these losses. According to the variation of permittivity and losses, the aging degree after 1512 h at 140 °C would be higher than after 5040 h at 120 °C. With the increase in degradation, the activation energy has significantly decreased, and the peak position has changed. However, monitoring the aging degree through this parameter does not seem feasible as the relaxation is very weak at low degradation degree.

V. CONCLUSION

This paper highlights the evolution of  r ' and  r " in XLPE samples with thermal aging. Also, β-relaxation process is studied and correlated to the physical structure changes of the material. The main conclusions of this study can be summarized as follows: i/ XLPE properties remain practically unchanged in the case of aging at 80 °C whereas some improvement of the crystallinity and the dielectric behavior of XLPE can be noticed at 100 °C. Physico-chemical degradation is effective when aging at 120 and 140 °C, leading to substantial increase of  r ' and  r ". ii/ Dielectric losses in degraded XLPE cables could be large near operating temperatures (10 to 40 °C). iii/ -relaxation peak intensity and position are strongly dependent upon the changes of the amorphous phase. This could be a good indicator of the severe deterioration of XLPE structure. For revealing the earlier stages of degradation, the carbonyl index seems to represent a better parameter.

Fig. 1 .

 1 r' variations with temperature and for different aging temperatures at: (a) 10 kHz, (b) 1 kHz, (c) 100 Hz and (d) 10 Hz

Fig. 2 .

 2 Fig. 2. r" variations with temperature and for different aging temperatures at: (a) 10 kHz, (b) 1 kHz, (c) 100 Hz and (d) 10 Hz

Fig. 3 .Fig. 4 .

 34 Fig. 3. Arrhenius plot of β-relaxation peaks for aging at 120 and 140 °C