Angélique Benoit

M Besse

R Louahdi

H Paul

P Paillard
email: pascal.paillard@univ-nantes.fr

T Baudin

Study of 6061 weld heat affected zones made with GMAW

Keywords: 6061 aluminum alloy, precipitation hardening, GMAW, infrared thermography, microhardness mapping, TEM

whether they are published or not. The documents may come

Introduction

6xxx aluminum alloys are widely used in the industry for their properties such as mechanical properties, corrosion resistance, weldability or good formability. Among the 6xxx aluminum alloys, the 6061 is one of the alloys most employed for structural application [START_REF] Parson | Light metals 1989 Campbell PG[END_REF][START_REF] Buha | Microstructural Development and Mechanical Properties of Interrupted Aged Al-Mg-Si-Cu Alloy[END_REF][START_REF] Ozturk | Influence of aging treatment on mechanical properties of 6061 aluminum alloy[END_REF] and is commonly welded. Welding is a joining operation which impairs base metal because of the high heat input used. It creates a zone where mechanical properties are different from the base metal which is called heat affected zone (HAZ). In the 6xxx aluminum alloy the HAZ is softer than base metal. Softening is due to the reversion of hardening precipitates β'' which are heat sensitive [START_REF] Dutta | A calorimetric study of precipitation in commercial aluminium alloy 6061[END_REF]. So, the characterization of this part is critical for the design of structural parts. 6061 welding have been already studied, especially using laser welding [START_REF] Cieslak | On the Weldability, Composition, and Hardness of Pulsed and Continuous Nd:YAG Laser Welds in Aluminum Alloys 6061,5456 and 5086[END_REF][START_REF] Hirose | Quantitative Evaluation of Softened Regions in Weld Heat-Affected Zones of 6061-T6 Aluminum Alloy-Characterizing of the Laser Beam Welding Process[END_REF], arc welding [START_REF] Moreira | Fatigue Behaviour of FS, LB and MIG Welds of AA6061-T6 and AA6082-T6" Multiscale Fatigue Crack Initiation and Propagation of Engineering Materials: Structural Integrity and Microstructural Worthiness[END_REF][START_REF] Ahmad | Effect of a post-weld heat treatment on the mechanical and microstructure properties of AA6061 joints welded by the gas metal arc welding cold metal transfer method[END_REF][START_REF] Malin | Study of Metallurgical Phenomena in the HAZ of 6061-T6 Aluminum Welded Joints[END_REF] or friction welding [START_REF] Murr | A TEM study of precipitation and related microstructures in friction-stirwelded 6061 aluminium[END_REF][START_REF] Woo | Prediction of hardness minimum locations during natural aging in an aluminum alloy 6061-T6 friction stir weld[END_REF]. Some kinds of processes have been compared for example laser, friction stir and metal inert gas (MIG) [START_REF] Moreira | Fatigue Behaviour of FS, LB and MIG Welds of AA6061-T6 and AA6082-T6" Multiscale Fatigue Crack Initiation and Propagation of Engineering Materials: Structural Integrity and Microstructural Worthiness[END_REF] or tungsten inert gas (TIG) and laser [START_REF] Hirose | Quantitative Evaluation of Softened Regions in Weld Heat-Affected Zones of 6061-T6 Aluminum Alloy-Characterizing of the Laser Beam Welding Process[END_REF]. In this study, we focus on arc welding processes. Different HAZ produced with four arc welding processes are examined. MIG and derivates, pulsed MIG and Cold Metal Transfer (CMT®) MIG, are compared to TIG process. TIG is a reference due to its ability to produce quality weld. The processes are compared through the analysis of the HAZ. Making easier the comparison, the linear welding energy of MIG weld was similar. Regarding TIG, because the process is too different from MIG, it was not possible to use same linear welding energy. So parameters were adjusted to obtain same fusion zone (ZF). Microstructures and microhardness results are compared to thermal measurements realized during the four welding. This study presents also a new way to characterize welding using infrared camera. It shows too an extension of the microhardness characterization, the mapping, made possible by using automated measurements.

Experimental

6 mm (0.24 in.) thickness plates in 6061 T6 were used for this study. The MIG, pulsed MIG and CMT® MIG beads were realized with Fronius Transpuls Synergic 3200 CMT®. TIG bead was made with a Esseti Quad 450 AC/DC welding power supply. The filler wire used both for MIG and TIG was 5356 aluminum alloy of 1.2 mm (0.05 in.) of diameter.

Struers Durascan Vickers hardness tester was used to create a microhardness mapping from the weld bead and up to 25 mm (0.98 in.) in length. Thank to fully automatic measuring, about 1800 points, spaced of 300 µm were performed, under a load of 0.1 kg.

The thermal measurements were realized using the infrared thermography. The camera is a FLIR Orion SC7500 equipped by an InSb sensor. A quasi-monochromatic filter of 3.35-3.45 µm was used with two different integration times allowing to be sensitive from 175°C to 860°C (347°F to 1580°F). The acquisition frequency is 50Hz. The camera was placed in front of the plate. Figure 1 describes the experimental setup. Prior to measurements, plates were painted in black in order to work with a constant emissivity of 0.96.

Results

1-Infrared thermography

The temperature distribution under the bead was accurately measured by infrared thermography during the welding process. These measurements were performed at 1 mm (0.04 in.) under the weld toe, for all the processes. This non destructive and non invasive testing of the temperature allows saving thermal cycles of welding processes in real time. This information has great interest in welding processes and can be directly related to the microstructure and mechanical properties of weld joints. Indeed, the initial microstructure can be deeply affected by the welding heat. It is thus possible to predict the different types of precipitates formed as a function of the temperature reached, and this at any point of the welded plate metal. According to thermograph in figure 3a, we found that thermal cycles of MIG and TIG are quite different. Indeed, for the same length of bead, the time of MIG welding are shorter than TIG welding. This is due to the characteristic of the TIG welding process with the feed rate can not exceed a few cm/min. Moreover, it is normal to observe very fast cooling rates for the three MIG processes. And as expected, the cooling rate of TIG process is longer. In addition, after 60 seconds, the arc breaking on the latter process leads to an acceleration of the cooling rate. This is represented by a change in slope on the thermograph curve. Concerning the temperatures reached, for equivalent welding energies, MIG and pulsed MIG allow approaching 530°C (986°F), while the CMT® process achieved 640°C (1184°F) (figure 3b). This significant difference can affect the microstructure with the formation of precipitates or the solution treatment of alloying elements. With the TIG process, the peak temperature (maximal) reached during welding is higher than that achieved by the MIG "classic" process (close to 600°C -1112°F).

2-Microhardness mapping

In order to evaluate the effects of the different thermal cycles on mechanical properties, microhardness mapping were performed on welded plate metal, from bead and on 25 mm (0.98 in.) in length and in cross-section (figure 4). The base metal, the aluminum alloy 6061-T6 has a hardness Vickers of 115, which corresponds to the blue color on the hardness scale. First, the hardness gradients are parallel and normal to the surface. This means that the temperature of the surface of the plate is at point very close to that just below it. Then, we see that the HAZ is very extensive with the TIG process since the hardness of the base metal is not recovered after 25 mm. The hardness is less than 80 HV, while with MIG processes the hardness of metal base is reached more quickly. The HAZ measures about 13 mm (0.51 in.) for MIG processes. This is explained by the low cooling rate of the TIG process, as observed in the thermograph (figure 3a). In the middle of the HAZ, the hardness varies between 50 and 70 HV (zone orange and red). These very low hardness values occur only with the TIG process. Further along in the HAZ, the hardness increases slightly (dark and light yellow areas) but these values are still quite low. This can be explained considering the hardening precipitation of the 6061 aluminum alloy [START_REF] Edwards | The Precipitation Sequence in Al-Mg-Si Alloys[END_REF]. The strengthening precipitates β'' present at T6 state are metastable and evolve spontaneously with temperature raising. So, β'' changes into β' phase and finally into the β phase which is stable [START_REF] Malin | Study of Metallurgical Phenomena in the HAZ of 6061-T6 Aluminum Welded Joints[END_REF]. However the strengthening ability of these precipitates is low. This induces a decreasing of the mechanical properties and therefore a decrease of hardness. It is therefore very likely that in the yellow and orange-red areas, we can observe this transformation sequence of metastable phases, including the presence of precipitates β' in large amount. Then these precipitates probably change partially into β phase, in the red-orange area. The well-defined hardness gradient in several layers is so easily explained by the transformation of β'', by the coexistence of one or more phases and by their density. Finally, the hardness increases again in the beginning of the HAZ, near the molten zone. This can be explained by a partial reversion of the precipitates because the temperature reached exceeds 530°C (986°F) with this process. And, this temperature corresponds to the equilibrium solid solubility limit of the 6061 alloy, according to the pseudo-binary diagram Al-Mg2Si. Thus, the dissolution of precipitates leads to a release of alloying elements (Mg and Si) into the aluminum matrix. So this solutioning of β' and β particles produces a supersursatured solid solution α'. And by aging, the hardness slightly increases until a moderate stage [START_REF] Malin | Study of Metallurgical Phenomena in the HAZ of 6061-T6 Aluminum Welded Joints[END_REF]. The evolution of the hardness of the three MIG processes is very similar between them, with some details yet. The extent of the HAZ is about 13 mm (0.51 in.) for all the three cases and the average hardness levels are similar. Indeed, in all cases, the hardness is greater than the hardness obtained by the TIG process. Arguably, we can state that, for similar welds in form, the TIG process further deteriorates mechanical properties of weld joints, with an extremely extended HAZ and lower hardness. With the pulsed MIG process, it seems that the hardness is less homogeneous in the different layers, especially approaching the base metal (mixed light and dark blue). So the size of the HAZ is maybe larger. Others measurements are needed to be sure.

Regarding the CMT® process, it is interesting to note that in the HAZ near the molten zone, the hardness is generally higher than for the two other MIG processes. This difference can be explained by a higher temperature reached during welding. Indeed, as noted above, the MIG CMT® process allows rising 640°C (1184°F). Thus, as with the TIG process, a partial solutioning is possible, and the dissolved precipitates release alloying elements to produce a supersursatured solid solution. With these hardness values obtained (between 80 and 90 HV -dark green area), this area can be considered as T4 temper state.

-TEM results

To confirm the results of hardness as well as hypothesis made above, electron microscopy in transmission was undertaken on the alloy welded by Cold Metal Transfer (figure 5).

The bright field image figure 5a corresponds to an extracted zone in the base metal. The needles shape of the precipitates and their size suggest that this is the β'' metastable phase, normally observed in the Al-Mg-Si alloys, in the T6 state.

Various publications [START_REF] Andersen | The Crystal Structure of the β'' Phase in Al-Mg-Si Alloys[END_REF][START_REF] Yao | A TEM study of precipitation in Al-Mg-Si alloys[END_REF] of the study of β'' phase describe its particles as needles directed along the three axes <001>Al. These fine needles about 50 nm in length also appear as dark spots pointing in the viewing direction, with an average diameter of 4 nm. Thanks to the presence of such precipitates, it is possible to achieve the maximum hardness values. According to [START_REF] Andersen | The Crystal Structure of the β'' Phase in Al-Mg-Si Alloys[END_REF], the maximum hardness is achieved when the alloy contains a combination of very fine Guinier- Preston zones (GP-1) and larger needles β'' phase, with a number density in the matrix of about 10 4 /µm 3 . The bright field image figure 5b corresponds to an extracted zone in the HAZ near the base metal. The density of precipitates seems to be lower in this area than in the base metal. β'' phase is always present. However, its changes of morphology in some places (green circles). It appears coarser. Therefore, it could be also the β' phase. Indeed, according to Dumolt [START_REF] Dumolt | Formation of a Modified β' Phase in Aluminium Alloy 6061[END_REF] increasing the aging temperature leads to a softening of the alloy due to coarsening of β'' precipitates.

Because of the nanometric size of these particles, it is difficult to distinguish them. High resolution microscopy (TEM/HREM) or atom-probe field-ion microscopy (APFIM) investigations for example, would help to have more specific evidences [START_REF] Hasting | Comparative study of the be-phasae in a 6xxx Al alloy by 3DAP and HRTEM[END_REF]. This has not been done in this study. However, the presence of a new phase (red circle) in laths shape can be observed. According to [START_REF] Chakrabarti | Phase relations and precipitation in Al-Mg-Si alloys with Cu additions[END_REF], it could be the Q' phase, metastable phase in the ternary alloys Al-Mg-Si where silicon content is in excess, the Q' phase equilibrium in the quaternary system Al-Mg-Si-Cu alloys. Both phenomena (β'' transformation and presence of a new phase) contribute to a slightly but significant softening of the alloy. The bright field image figure 5c corresponds to an extracted zone in the HAZ near the molten zone. The density of precipitates decreases again. The β'' no longer seems to be present. The β'' fade-out, combined with the presence of precipitates coarser and therefore fewer, lead to a high decrease of hardness. This microstructure, with a very low density of precipitates is consistent with the hypothesis of partial dissolution during the welding.

Conclusions

In this study we tried to correlate the physical parameters (temperatures, thermal cycles) associated with the four welding processes (MIG, pulsed MIG, MIG CMT®, and TIG) with mechanical properties (hardness) and microstructure (TEM study of precipitates). This study only focused on the HAZ zones of the 6061-T6 alloy. Thermography provides an accurate estimation of the temperature reached during welding, as well as the heating and cooling rate. With the MIG CMT®: "Cold Metal Transfer" the temperature reached is higher than those obtained with MIG and pulsed MIG, for same welding linear energies. This is quite surprising, considering the name given of this process. However, this process provides the best results of hardness. Indeed, the size of the HAZ is identical to those of the two others MIG processes. But the average hardness, in particular near the weld bead, is higher. Comparative hardness tests (conducted from the bead and up to 25 mm in length) carried out by mapping is a new visual way to present the evolution of mechanical properties. Finally, with an initial TEM study, it was possible to establish a direct relationship between the evolution of hardness and the presence of different precipitates. Among the outlook, high resolution TEM studies could allow to have more precise identifications of the different precipitates. Calorimetric measurements are also planned to identify the sequence of different phase transformations, and to know at which temperatures these transformations occur.

Figure 1 :

 1 Figure 1: experimental setup of thermograph measurement Transmission electron microscope used for this study was a FEI TECNAI G² 200kV. Transmission electron microscope

Figure 2 :

 2 Figure 2: extraction zones in the plate

Figure 3 :

 3 Figure 3: thermographs a) comparison of TIG and MIG processes, b) comparison of MIGs

Figure 4 :Figure 5 :

 45 Figure 4: Hardness mapping HV0.1 of the four welding processes 6 mm (0.24 in.)

Acknowledgments

The authors would like to greatly acknowledge Lidia Lityńska-Dobrzyńska for TEM examinations.


Fady El Haber
email: fady.elhaber@cnrs-imn.fr

Xavier Rocquefelte
email: xavier.rocquefelte@cnrs-imn.fr

Christine Andraud

Bouhalouane Amrani

Stéphane Jobic

Olivier Chauvet

Gerard Froyer

Gérard Froyer

Prediction of the transparency in the visible range of x-ray absorbing nanocomposites built upon the assembly of LaF 3 or LaPO 4 nanoparticles with poly(methyl methacrylate)

à la diffusion de documents scientifiques de niveau recherche, publiés ou non, émanant des établissements d'enseignement et de recherche français ou étrangers, des laboratoires publics ou privés.

Prediction of the transparency in the visible range of x-ray absorbing nanocomposites built upon the assembly

of LaF3 or LaPO4 nanoparticles with poly(methyl methacrylate)

INTRODUCTION

During the past few years, lanthanum trifluoride (LaF 3 ) and lanthanum phosphate (LaPO 4 ) nanoparticles (NPs) have been extensively investigated as antiwear materials [START_REF] Zhou | Synthesis and room temperature ionic conductivity of nano-LaF 3 bulk material[END_REF], polymer fillers or lubricating grease additives [START_REF] Dekker | Proceedings Symposium IEEE/LEOS Benelux Chapter[END_REF][START_REF] Stouwdam | Near-infrared emission of redispersible Er 3 ,N d 3 , and Ho 3 doped LaF 3 nanoparticles[END_REF], and host lattices for luminescent properties [START_REF] Grzyb | Photoluminescent properties of LaF 3 :Eu 3 and GdF 3 :Eu 3 nanoparticles prepared by co-precipitation method[END_REF][START_REF] Pi | Luminescence behavior of Eu 3 doped LaF 3 nanoparticles[END_REF][START_REF] Lim | Near infrared luminescence properties of nanohybrid film prepared from LaPO 4 :Er 3 ∕LaPO 4 core/shell nanoparticles and silica-based resin[END_REF][START_REF] Jung | Preparation and luminescence properties of LaPO 4 :Er, Yb nanoparticles[END_REF][8][START_REF] Wei | Polyol-mediated synthesis of water-soluble LaF 3 :Yb, Er upconversion fluorescent nanocrystals[END_REF][START_REF] Xie | Synthesis of highly fluorescent LaF 3 Ln 3 ∕LaF 3 core/shell nanocrystals by a surfactantfree aqueous solution route[END_REF]. Recently, El Haber and Froyer also showed that these materials also exhibit remarkable properties for protection against x-ray radiation [START_REF] Haber | Transparent polymers embedding nanoparticles for x-rays attenuation[END_REF]. Hence, these compounds, once embedded in a highly transparent organic matrix such as the poly(methyl methacrylate) (PMMA), can be regarded as potential substitutes for lead-containing glasses. Such an alternative has great merit both to respect the European directives in terms of toxicology [START_REF]The RoHS Regulation[END_REF] and to lighten the protective devices, making easy their manipulation. However, the challenge remains so far to achieve a good compromise between the x-ray absorbing capacity and the high transparency of the x-ray shield to envision applications in the medical domain, for instance. Very preliminary investigations carried out on ball-milled PMMA∕LaF 3 or PMMA∕LaPO 4 powdered mixtures with 10% volume of 6 nm large NPs exhibit a lead equivalency of 0.1 mm for x-ray attenuation for 3 mm thick white pellets. Based on this result, specific synthetic routes were developed for the elaboration of LaF 3 and lanthanum orthophosphate NPs with modified surfaces to favor their dispersion in PMMA by dispersing surface modified NPs in methyl methacrylate monomer for radical polymerization [START_REF] Haber | Contribution à l'obtention de verres organiques sans plomb atténuateurs des rayons X-Synthèse modification de surface et caractérisations de nanoparticules incorporables dans les polymères acryliques[END_REF]. Nevertheless, if LaPO 4 and LaF 3 in PMMA may be envisioned in the future as efficient x-ray shielding system in the 38-80 keV energy domains [START_REF] Haber | Transparent polymers embedding nanoparticles for x-rays attenuation[END_REF], some prerequisites are expected before commercialization. Hence, in addition to the x-ray absorption capabilities, such screens must exhibit for applications in the radiology department (Fig. 1) a high transparency in the visible range to facilitate the observation of the patient by the xray manipulator throughout the examination. For this reason, a light transmission higher than 70% in the visible range (380-800 nm) of the x-ray shielding panel is requested.

The aim of the present paper is to predict the transparency properties of such nanocomposite material in the visible range using a multiscale theoretical approach, taking into account the refractive index of the inorganic particles but also their size, their concentration in the organic medium, the thickness of the x-ray attenuator, etc. Namely, after analyzing the intrinsic optical properties of LaPO 4 and LaF 3 based on first-principles calculations, the refractive indices and the extinction coefficients of these two compounds were extracted and used as input parameters in four-flux calculations to shed light on the extrinsic optical properties of the nanocomposites (i.e., LaPO 4 or LaF 3 embedded in PMMA) materials.

THEORETICAL APPROACH A. Our Strategy

In general, the light interaction with nanocomposites can be roughly estimated via the Beer-Lambert law for highly diluted systems. When the concentration in the scattering centers increased, multiple scattering phenomena appear and cannot be neglected. Namely, the electromagnetic light scattering is a complex phenomenon because of the existence of secondary radiations dissipating energy in all directions. Therefore, the total scattered light depends strongly on the relative size of the particles compared to the incident wavelength and how these particles are dispersed in the matrix with respect to the incident light and scattered directions [START_REF] Mishchenko | Scattering, Absorption, and Emission of Light by Small Particles[END_REF].

In case of diluted media, for well-separated small particles, single-scattering approximation can be used [START_REF] Maheu | Four-flux models to solve the scattering transfer equation: special cases[END_REF][START_REF] Maheu | Four-flux models to solve the scattering transfer equation in terms of Lorenz-Mie parameters[END_REF]. However, in the present case, the scattering medium contains a large number of particles. As a consequence, the single-scattering approximation is no longer valid, and multiple scattering phenomena have to be taken into account, which implies solving the radiative transfer equation (RTE) [START_REF] Auger | Dependent light scattering in dense heterogeneous media[END_REF]. Considering that there is no interference between the radiation scattered by neighboring particles [START_REF] Niklasson | Comparison between four flux theory and multiple scattering theory[END_REF], the resolution of the RTE leads to four-flux calculations to estimate transparency properties of nanocomposites. To carry out such radiative transfer calculations, few inputs are needed: the optical constants (refractive index and extinction coefficient) of the organic medium and light scatterers, the shape of the scattering centers (here supposed to be spherical), their size and their concentration, and the thickness of the medium to pass through.

B. Intrinsic Optical Properties-Density Functional Theory Calculations

An accurate description of the optical properties of LaF 3 and LaPO 4 bulk compounds may be achieved via a quantum mechanical treatment. Namely, both scattering and absorption properties of an insulating compound are related to the dielectric function εωε 1 ωiε 2 ω, which describes the linear response of the electronic structure to the electric field of the impinging light carrying an ℏω energy [START_REF] Lynch | Interband Absorption-Mechanisms and Interpretations-Handbook of Optical Constants of Solids[END_REF][START_REF] Bassani | Electronic States and Optical Transitions in Solids[END_REF]. The spectral dependence of the optical properties can be theoretically determined from the electronic band structure by calculating the ε 2 ω spectrum, which is related to the capacity of the material to absorb light [START_REF] Bassani | Electronic States and Optical Transitions in Solids[END_REF]. The ε 1 ω function is then obtained using the Kramers-Kronig relations [START_REF] Ambrosch-Draxl | Linear optical properties of solids within the full-potential linearized augmented plane wave method[END_REF]. It should be noted that, in contrast to the treatment of experimental optical data, no extrapolations are necessary to perform the integrals in the Kramers-Kronig analysis in the present DFT calculations due to an estimation of the ε 2 values for a high energy range. Finally, the complex refractive index, Nωnωikω, could be deduced from ε 1 ω and ε 2 ω using the following formulas:

nω ε 1 ωω 2 1 ω 0.5 0.5 2 p ; (1a) kω -ε 1 ωω 2 1 ωω 2 2 ω 0.5 0.5 2 p ; (1b)
where nω and kω are the refractive index and the extinction coefficient at a given optical angular frequency of the incoming radiation, respectively. These two quantities, used as input in the four-flux calculations, will be determined here on the basis of DFT calculations using the WIEN2k program package [START_REF] Blaha | WIEN2k: An Augmented Plane Wave+LO Program for Calculating Crystal Properties[END_REF].

In order to correctly describe the position of the very localized 4f states of La 3 ions (4f °electronic configuration) and the related optical properties, we have carried out DFT calculations within the generalized gradient approximation (GGA) [START_REF] Perdew | Generalized gradient approximation made simple[END_REF] plus on-site Coulomb repulsion (U). The validity of the GGA U approach has been tested by changing the value of the effective Hubbard parameter, U eff U -J, where J is the Hund's exchange interaction, the U value being determined by trial and error. As will be discussed in Subsection 3.A,aU eff value of about 10.3 eV has been shown to provide a coherent picture both in terms of x-ray photoelectron spectroscopy (XPS) combined with bremsstrahlung isochromat spectroscopy (BIS) and reflectivity spectroscopy of LaF 3 .

The density of states and optical properties are deduced from a self-consistent calculation, using the full-potential linearized augmented plane-wave method, as embodied in the WIEN2k code. The maximum l value in the expansion of the basis set inside the atomic sphere was 12. The convergence of the basis set is controlled by a cutoff parameter RMT × K max 7, where RMT is the smallest atomic sphere radius in the unit cell and K max is the magnitude of the largest k vector. The self-consistency was carried out on a 32 k-point mesh in the irreducible Brillouin zone, with the following radii: RMTLa2.3 a.u., RMTF2.1 a.u., RMTP1.67 a.u., RMTO1.2 a.u., and GMAX 12 bohr -1 , where GMAX is the magnitude of the largest vector in the charge density Fourier expansion.

C. Extrinsic Optical Properties-Four-Flux Calculations

It is important to note that materials under consideration consist of NPs dispersed in PMMA (Fig. 2). To ensure a good transparency to the nanocomposite, light scattering caused by the NPs must be minimized as described in the Rayleigh scattering equation [Eq. ( 2)]. Namely, scattering phenomena is enhanced when the diameter (D p ) and the volumetric concentration of the inorganic particles in the matrix increase (short interparticles r distances) but can be minimized when n p ∕n m tends to 1:

I R 8π 4 n m D 6 p λ 4 r 2 n p n m 2 -1 n p n m 2 2
1 cos θI 0 ; (2)
the Rayleigh scattering equation where I R and I 0 are the scattered and the incident intensities, respectively, D p the particle diameter, λ the wavelength, r the distance between the particles, and θ the scattering angle. n p ∕n m is the refractive indices ratio of the inorganic particles (n p ) and the organic medium (n m ). Namely, the RTE [Eq. ( 3)] accounts for both the simple scattering and the multiple one. Emission phenomena are not considered at all in this type of analysis: This four-flux model was phenomenologically established by Maheu and Gouesbet [START_REF] Maheu | Four-flux models to solve the scattering transfer equation: special cases[END_REF]. The solution of Eq. 3 gives access to the values of the four fluxes, that is, the specular and the scattered reflected beams (hereafter labeled Rcc and Rcd, respectively) and the specular and the scattered transmitted beam (hereafter labeled Tcc and Tcd, respectively) [START_REF] Vargas | Generalized four-flux radiative transfer model[END_REF][START_REF] Rozé | Average crossing parameter and forward scattering ratio values in four-flux model for multiple scattering media[END_REF] (Fig. 2).

The multiple reflections inside the layer are taken into account with the Fresnel coefficient for the collimated beam, and for scattered beam, the diffuse reflection coefficient is calculated using an angular average over all angles of incidence of the Fresnel reflection coefficient rθ (where θ is the angle of incidence of the light), considering that the scattered beam is isotropic. To gain simplicity, it should be noticed that, in our calculations, the following restrictions have been considered: (i) incident light is supposed to be monochromatic light with a quasi-constant electric field; (ii) the scattering light has the same frequency as the incident light; and (iii) the medium surrounding the NPs is homogeneous, linear, isotropic and nonabsorbing (i.e., PMMA).

RESULTS AND DISCUSSION

A. Intrinsic Optical Properties-LaF 3 and LaPO 4 Bulk Compounds In order to extract the refractive indices and extinction coefficients of LaF 3 and LaPO 4 , we have estimated the optical properties of these two compounds in their bulk configuration based on DFT calculations. Before discussing these results, we first consider the effect of the U eff and scissors-operator (SO) parameters on the electronic structure of LaF 3 . The onsite U eff parameter will be taken the same for LaPO 4 (see below).

The total densities of states of LaF 3 obtained using GGA, GGA U and GGA U SO are represented in Fig. 3.I n the present case, a Gaussian broadening has been applied to the theoretical density of states (DOS) in order to reproduce the experimental broadening related to XPS and BIS experiments of LaF 3 [START_REF] Park | Electron-spectroscopy study of rareearth trihalides[END_REF]. To have a proper description of the optical properties of these two compounds, it was necessary to apply two corrections. The first one consists in applying a Hubbard term (U eff 10.3 eV) to the 4f La states in order to have a more localized description of these levels and then an adequate position of the energy levels. The second one consists in having a proper description of the bandgap of these two compounds. There are several ways of correcting this DFT deficiency [START_REF] Park | Electron-spectroscopy study of rareearth trihalides[END_REF], the simplest one being to use an SO, i.e., shifting the conduction bands rigidly with respect to the valence bands to have a bandgap similar to the experimental one (SO 2 and 1.8 eV for LaPO 4 and LaF 3 , respectively). It should be noticed that, in contrast to the correction of the bandgap, which is system dependent, the U eff value is expected to be the same for LaF 3 and LaPO 4 due to the atomiclike nature of the 4f La states, in a first approximation.

The general features of these three calculated DOSs given in Fig. 3 are similar except a significant change of the La4f band positions as a function of the U eff value and of the bandgap as a function of the SO value. Figures 3a and3b show the effect of applying a Hubbard term (U eff 10.3 eV). However, while the correction of the 4f La position using GGA U allows us to have a proper description of the shape of the BIS (related to the conduction band) and the XPS (related to the valence band) spectra, it does not provide a good description of the bandgap value, as expected. The final result is then obtained using GGA U SO and shown in Fig. 3c, leading to good agreements with the reported XPS/BIS experimental data [START_REF] Held | LDA DMFT investigations of transition metal oxides and f -electron materials[END_REF].

In summary, the electronic structure of LaF 3 could be described as follows. Below the Fermi level, the occupied bands are mainly composed of 2pF states interacting with 5dLa states. Above the Fermi level, the first empty states are mainly composed of 5dLa states, and the 4f La states are inside the 5dLa block and not in the bandgap, as found in GGA calculations. This observation has a strong impact in terms of the understanding of the optical properties of LaF 3 : at low energy, the optical properties will be mainly controlled by the 2pF → 5dLa optical transitions.

Figures 4a and4b compare the measured reflectivity at normal incidence on a basal plane of LaF 3 [START_REF] Olson | Optical properties of single crystals of some rare-earth trifluorides, 5-34 eV[END_REF] to the simulation deduced from GGA and GGA U calculations. In the two cases, the simulations before and after applying the SO are provided. The arrow in Fig. 4a shows the optical transitions involving the 4f La states and that are not expected at such low energy. In contrast, using GGA U SO approach allows having a proper description of the LaF 3 reflectivity curve up to energy of 20 eV, at least. Such an observation confirms the validity of the present approach to properly describe the optical properties of LaF 3 . It should be noticed that the SO for LaPO 4 (SO 2 eV) has been chosen in such a way to reproduce its experimental optical gap of 7.8 eV [START_REF] Nakzawa | Vacuum ultraviolet luminescenceexcitation spectra of RPO 4 :Eu 3 (R Y , La, Gd and Lu)[END_REF]. Calculated band structure of LaPO 4 is not given here but compares well with the one determined by Mishra et al. [START_REF] Mishra | First principles investigation of host excitation of LaPO 4 ,L a 2 O 3 and AlPO 4[END_REF].

Extinction coefficient (k) and refractive index (n) values for these two compounds could then be extracted from the present GGA U SO calculations. Figures 5a and5b represent k and n, respectively, for LaF 3 and LaPO 4 . The present calculations evidenced that LaPO 4 and LaF 3 start absorbing light at about 7.8 and 10 eV, respectively. As expected, these Fig. 4. Experimental [START_REF] Olson | Optical properties of single crystals of some rare-earth trifluorides, 5-34 eV[END_REF] (open circles) and theoretical reflectance spectra of LaF 3 : (a) simulations deduced from GGA (dashed line) and GGA SO (solid line), (b) simulations deduced from GGA U (dashed line) and GGA U SO (solid line). The arrow indicates the region that is incorrectly simulated by the GGA SO approach due to an underestimation of the 4f La energy. two compounds are intrinsically transparent in the visible light range (1.55-3.1 eV), i.e., no absorption occurs in the 380-800 nm domain. In addition, it should be noticed that the high bandgap value implies a resulting small refractive index value in the visible range, due to the causality rule that links k and n, expressed by the Kramers-Kronig relation. In fact, the refractive indices deduced from our DFT calculations in the energy window (1.55-3.1 eV) ranged from 1.80 to 1.84 for LaPO 4 and 1.62 to 1.64 for LaF 3 , in accordance with the experimental measurements [START_REF] Chindaudom | Determination of the optical constants of an inhomogeneous transparent LaF 3 thin film on a transparent substrate by spectroscopic ellipsometry[END_REF][START_REF] Gramaccioli | A uranium-and thoriumrich monazite from a south-alpine pegmatite at Piona, Italy[END_REF].

To gain simplicity, the refractive indices of LaF 3 and LaPO 4 would be taken constant on the whole visible range, i.e., 1.62 and 1.82, respectively, while k would be settled to zero.

B. Extrinsic Optical Properties-Nanocomposite Materials

To anticipate the transparency properties of LaF 3 and LaPO 4 NPs in PMMA and to direct the preparation of x-ray absorber/ polymer composites with the requisites for medical applications, for instance, simulations of the optical characteristics in the visible range were carried out by the four-flux method with a special calculation code developed by Institute of Nanoscience in Paris, France [START_REF] Silva | Multiple light scattering in multilayered media: theory, experiments[END_REF]. The refractive index of PMMA (n PMMA ) was taken equal to 1.49 [START_REF] Zieman | POF Handbook: Optical Short Range Transmission Systems[END_REF]. The thickness (e) of the x-ray shield was taken at 3 mm or more, and the volume percentage of NP in polymer (vol. %) was equal or higher than 10% to ensure a sufficient absorption of x ray. By reference with preliminary synthetic investigations [START_REF] Haber | Contribution à l'obtention de verres organiques sans plomb atténuateurs des rayons X-Synthèse modification de surface et caractérisations de nanoparticules incorporables dans les polymères acryliques[END_REF], particle size was taken around 6 nm.

Figure 6 displays the evolution of light transmission [Tλ] versus the nanocomposite thickness (e) along the whole visible range for spherical absorbing particle with 6 nm diameters. We can clearly notice that the thickness in the 3-5 mm domain (step of 0.5 mm) does not impact significantly the transparency of the materials even if a more pronounced decrease of Tλ is observed for wavelengths lower than 600 nm. Namely, for LaF 3 ∕PMMA nanocomposite, light transmission shifts from 88.9% to 88.4% at 800 nm and from 79.3% to 73.8% at 380 nm going from e 3 mm to e 5 mm. For LaPO 4 ∕PMMA composites, the transparency is less than with LaF 3 due to a higher refractive index of the scattering centers. The transmission shifts then from 84.1% to 83.6% at 800 nm and from 71.2% to 63.3% at 380 nm for e 3 mm and e 5 mm, respectively. Nevertheless, transmission remains higher than 70% for any compositions at 555 nm, i.e., the wavelength corresponding to the maximum photopic sensitivity of the eyes. This means that LaF 3 ∕PMMA and LaPO 4 ∕ PMMA composites may be convenient to replace 0.1 mm lead equivalency foils to protect against x ray while insuring a good transparency in the visible range.

To evidence the role of the size particle of the scattering particles, Fig. 7a depicts the evolution of transparency at 555 nm for layers whose thickness changes from 0.5 to 5 mm. Calculations were carried out with LaF 3 and LaPO 4 NPs with diameter of 3 and 15 nm for a constant volume percentage of 10%. While it turns out that transmission at 555 nm is almost not impacted by panel thickness for particle size of 3 nm, T555 nm falls down very fast with (e) for 15 nm diameter particles. The higher the particle size, the lower the transparency of the absorbing panel for a given thickness, at least for D p ≪ λ. This is well illustrated in Figs. 7a and7b, while Fig. 7c represents the evolution of the transmission versus D p ∕λ in the UV-visible-IR domain for LaF 3 and LaPO 4 NPs at vol. % 10% in PMMA, for e 3 mm.

Let consider now the evolution of T555 nm for LaF 3 and LaPO 4 6 nm diameter NPs in PMMA for different volume percentages of the inorganic phase for a panel thickness of 3 mm. Figure 8 shows the T555 versus n p ∕n PMMA for volumetric concentration of 10%, 30%, 50%, and 70%. It appears clearly that the higher the n p ∕n PMMA ratio, the higher the negative impact of an increase of the particle/polymer volumetric ratio. Namely, T555 nm shifts from 90% to 85% and from 83% to 63% when n p ∕n PMMA increases from 1.09 to 1.21 going from LaF 3 to LaPO 4 . For a reminder, T555 nm remains constant and equal to 93% for a particle with the same refractive index as the medium. The attenuation of the intensity of the passing beam is observed because particles hit by the impinging beam act as scattering centers emitting radiation in any direction. In particular, it clearly appears that a way to overcome the transmission light decrease when the refractive index of the inorganic particle is increased is to reduce significantly the diameter of the inorganic particle. In the present case, with particles diameters of 5 nm and less, the use of inorganic compounds with a refractive index up to 1.9 will do not induce too strong an opacity of the nanocomposite material. In contrast with particles having a diameter higher than 5 nm, the opacity will be strongly dependent of the refractive index value of the particle.

CONCLUSION

To conclude, the present theoretical investigation has provided a way to estimate the impact of intrinsic and extrinsic parameters on the optical properties of a nanocomposite material composed of inorganic inclusions (LaF 3 and LaPO 4 ) embedded in an organic medium (PMMA). More specifically, the role of the extrinsic parameters n p ∕n m and D p ∕λ have been discussed based on a multiscale approach combining DFT and four-flux calculations. After validating our DFT results based on comparison with XPS/BIS and reflectivity experimental data, the intrinsic optical characteristics of LaF 3 and LaPO 4 have been extracted and used as input parameters in the four-flux calculations. Our results provide the magnitude of the light transmission reduction when n p ∕n m deviates from 1 and when D p ∕λ is raised (impact of the multiple scattering). In particular to avoid too strong an opacity of a panel composed of such nanocomposite material with PMMA as the organic medium, the use of LaF 3 has to be more privileged than LaPO 4 because it provides an n p ∕n m ratio closer to 1. Whatever, if we limit our analysis to a 3 mm thick shield with a 10 vol. % inorganic absorber concentration to ensure 0.1 mm lead equivalency of protection against x rays, NPs of a size lower than 6 nm should be sufficient to lead transmission higher than 70%. Let us notice that another strategy allowing the insertion of particles with a larger size would consist in the use of an organic medium with a higher refractive index, that is polystyrene (n m 1.55) in place of PMMA (n m 1.49), for instance.

Fig. 1 .

 1 Fig. 1. X-ray shielding structure used in radiology departments where the x-ray manipulator has to be protected against x rays in everyday life. The transparent x-ray shield (upper part) should transmit at least 70% of visible light with lead equivalency of 0.1 mm

  dI ds -α abs α sca I α sca R 4π p • I 0 dω 0 ↑↑ single scattering multiple scattering ; 3 the RTE where I is the intensity of radiation, p the phase function depending on the angular distribution of scattered light, α sca NC sca the scattering coefficient (N number of inorganic inclusions in the organic medium per volume unit), and α abs NC abs the absorption coefficient (C sca and C abs are the scattering and absorption cross sections, respectively, calculated by Mie theory).

Fig. 2 .

 2 Fig. 2. Schematic representation of the interaction of light with an xray absorbing panel built upon inorganic NPs embedded in PMMA (Rcc, specular reflectance; Rcd, scattered reflectance; Tcc, specular transmission; Tcd, scattered transmission).

Fig. 3 .

 3 Fig. 3. Comparison of the total densities of states of LaF 3 deduced from (a) GGA, (b) GGA U, and (c) GGA U SO approaches. The open circles correspond to the XPS/BIS experimental data [26].

Fig. 5 .

 5 Fig. 5. (a) Theoretical extinction coefficients and (b) refractive indices of LaF 3 and LaPO 4 deduced from GGA U SO calculations.

Fig. 6 .

 6 Fig. 6. Evolution of the total transmission (Tcc Tcd) versus the thickness e of the x-ray absorbing layer for e ranging from 3 to 5m m( D p 6 mm, vol. % 10%). (a) LaF 3 NPs n p 1.62; (b) LaPO 4 NPs np 1.8.

Fig. 7 .

 7 Fig. 7. (a) Transmission at 555 nm of x-ray attenuators whose thickness ranges from 0.5 to 5 mm for LaF 3 and LaPO 4 NPs embedded in PMMA (volume fraction of 10%). (Left) D p 3 nm; (Right) D p 15 nm. (b) Transmission at 555 nm versus inorganic NP diameters a 3 mm thick x-ray shield (volume fraction of NPs 10%). (c) Light transmission of LaF 3 ∕PMMA x-ray attenuator versus D p ∕λ (where λ 380-800 nm) ratio for particle diameters of 3, 5, 10 and 15 nm (e 3 mm, vol. % 10).


Bernède

Y Lare

M Banéto

L Cattin

M Morsli

K Jondo

K Napo

J C Bernède

Á M Bane

Á K Jondo

Á K Napo

Á J C Berne

Effect of a zinc oxide, at the cathode interface, on the efficiency of inverted organic photovoltaic cells based on the CuPc/C 60 couple

Effect of a zinc oxide, at the cathode interface, on the efficiency of inverted organic photovoltaic cells based on the CuPc/C(60) couple

Y.

Introduction

The active heart of organic photovoltaic cells (OPV cells) consists either of two organic layers or an homogeneous mixture of two organic materials [START_REF] Berne `de | Organic Solar Cells Performances Improvement Induced by Interface Buffer Layers[END_REF][START_REF] Cai | [END_REF]. One of them, either an organic dye or a semi conducting polymer, donates the electrons. The other serves as electron acceptor. OPV cells with power conversion efficiency (PCE) of 5-6% under simulated AM1.5 illumination have been obtained [3][4][5]. Usually, in these devices, an indium tin oxide, (ITO), coated glass substrate is used as transparent anode and an aluminium thin film as cathode. The classical multi-heterojunction structures are based on ITO/anode buffer layer/ CuPc/C 60 /cathode buffer layer/Al devices where CuPc is the copper phthalocyanine and C 60 the fullerene. The buffer layers (BF) are necessary in view of the difficulties in organic optoelectronic devices of the charge carrier transport between the organic materials and the electrodes. In the case of the anode/electron donor contact, a common solution is to introduce a thin anode buffer layer (ABL), which adjusts the electronic behaviour of the adjacent materials. We have shown that an ultra-thin metal film or a thin oxide film deposited onto the conductive substrate, whatever the transparent conductive oxide (TCO), can be used to improve the interface TCO/donor and therefore the devices performances [6][7][8].

It has been shown that, when a large band gap material, such as bathocuproine (BCP) [9] or aluminium tris(8hydroxyquinoline) (Alq 3 ) [ 10,11], is introduced between the cathode and the electron acceptor, there is a strong improvement of the OPV cells performances. This layer is called exciton blocking layer (EBL). It should be transparent to the solar spectrum to act as a spacer between the photoactive region and the metallic cathode. It must, also, transport electrons to avoid high series resistance.

If the classical devices described above have given the best performances, development of inverted solar cells could allow more flexibility on designing tandem structures, which can be formed using a semitransparent cathode. In an inverted device the TCO is the cathode and the metal electrode is the anode. Up to now, the inverted organic solar cells studied are usually based on bulk heterojunctions and promising results have been obtained [12][13][14], while no many works have been dedicated to inverted multi-heterojunction PV cells [15]. The challenge to reversing the layer sequence of multi-heterojunction PV cells is achieving cathode and anode ohmic contacts. We have already shown that the presence of a thin BCP layer at the interface ITO cathode/C 60 allows improving OPV cells performances in the same way that in classical ones [16]. In the present work we study the effect of a zinc oxide buffer layer, at the cathode interface, on the efficiency of inverted OPV cells based on the BCP/CuPc/C 60 multi-heterojunction structure.

Experimental

The device structure used in this study is:

ITO=CBL=BCPð9nmÞ=C 60 ð40nmÞ=CuPcð35 nmÞ=ABL=Al
Here CBL is a buffer layer which allows achieving a good matching between the ITO cathode and the organic material. The multi-layer heterojunction structure used in the present work has ITO/CBL as bottom cathode and ABL/Al as anode. The buffer layer ABL is used to achieve a good matching between the work function of the anode and the highest occupied molecular orbital (HOMO) of CuPc. We have already shown that Au, MoO 3 [6,7] or, better, MoO 3 /Au, allows achieving this goal [8], by simple vacuum deposition of a thin (3.5 ± 1 nm) MoO 3 film, of an ultrathin (0.5 nm) gold film or of MoO 3 /Au superposed films. The best results were achieved with this last combination. Therefore, we used MoO 3 /Au as ABL buffer layer.

The effect of the CBL on the properties of the inverted organic solar cell is the subject of this work. Here, CBL is a conductive or insulating ZnO film. Therefore, the structures probed were:

ITO=ZnO=BCPð9nmÞ=C 60 ð40 nmÞ=CuPcð35 nmÞ= ABL=Al;
the ZnO being conductive (ZnO:Al) or insulating (ZnOsc).

The ITO coated glass substrates used in this study were commercially obtained from the SOLEMS. Since ITO covered the whole glass substrates, some ITO must be removed to obtain the bottom electrode. After masking a broad band of 30 mm by 20 mm, the ITO was etched by using Zn ? HCl. After carefully scrubbing, with alcohol and then soap, these substrates were rinsed in running deionised water. The substrates were dried and immediately loaded into a vacuum chamber (10 -4 Pa) when the ZnO films were obtained by rf magnetron sputtering (ZnO:Al) or introduced into a hood, when they were deposited by spin coating (ZnOsc).

The chemical products (CuPc, C 60 , BCP…Au, MoO 3 ) have been provided by Aldrich. All the thin films of the structures have been deposited in the same run in a vacuum of 10 -4 Pa. The thin film deposition rates and thickness were estimated in situ with a quartz monitor. The deposition rate and final thickness were 0.1 nm/s and 9 nm for BCP, 0.05 nm/s and 35 nm in the case of CuPc and 0.05 nm/s and 40 nm in the case of C 60 . These thicknesses have been chosen after optimisation.

In the case of the buffer layer ABL, the known optimum thickness have been used, 3 nm for MoO 3 [8] and 0.5 nm for Au [6,7]. For the CBL, the ZnO thickness probed were 0, 3 and 5 nm.

After ABL (MoO 3 /Au) deposition, the aluminium top electrode was thermally evaporated, without breaking the vacuum, through a mask with 2 9 8 mm active areas, and following by 50 nm thick encapsulating layer of amorphous selenium (Se-a). The selenium protective coating layer has been proved to be efficient to protect the under layers from oxygen to water vapour contamination [17], at least during the first hours of room air exposure [11].

The conductive ZnO films were deposited by ac magnetron sputtering on glass substrates at room temperature. A cylindrical target of diameter 75 mm and thick of 5 mm was prepared by sintering the mixture of ZnO powder (purity of 99.99%) with 2wt% Al 2 O 3 (purity 99.99%). Ar (99.99%), source is used as sputtering gas for deposition of the ZnO:Al films. The distance between target and substrate was about 60 mm. The flow rate of Argon was controlled by using flow meter. The sputtering pressure was maintained at 0.13 Pa. The pressure of the sputtering system before the deposition of the sample films was under 2 9 10 -3 Pa. The films were deposited onto substrates under 150 W. These deposition conditions were fixed in order to achieve ZnO:Al conductive transparent films. The conductivity of the obtained films has been checked, using a ZnO:Al film thick of 100 nm, it is around r = 2.32 9 10 3 (X cm) -1 .

The insulating films were deposited by spin coating. Appropriate amount of zinc nitrate (0.05 mol/l) and ammoniac (0.1 mol/l) were mixed. The pH of the bath was 8.0. The solution was maintained at a temperature 60 °C \ T \ 65 °C, while the bath was continuously stirred. 0.5 ml of this solution was deposited by spin coating onto ITO glass coated substrate, the spin speed was 2,000 rpm and the spin time was 60 s. The coated films were then annealed at 300 K for 1 h in order to change the Zn(OH) 2 into ZnO polycrystalline films. The thickness of the ZnO films was estimated using an atomic force microscope apparatus (AFM).

The ZnO films have been characterized by different techniques. The morphology was observed through scanning electron microscopy (SEM) with a JEOL 6400F at the ''Centre de Microcaracte ´risation de l'Universite ´de Nantes''. The crystalline structure of the films was analyzed by X-ray diffraction (XRD) by a Siemens D 5000 diffractometer using Ka radiation from Cu (k = 0.15406 nm). AFM images on different sites of the film were taken ex-situ at atmospheric pressure and room temperature. All measurements have been performed in tapping mode (Nanoscope III, (Digital Instruments, Inc.). Classical SiN4 cantilevers were used (Ultrasharp, AXESSTECH, France). The average force constant and resonance were approximately 14 N/m and 320 kHz, respectively. The cantilever was excited at its resonance frequency.

Electrical characterizations were performed with an automated I-V tester, in the dark and under one sun global AM 1.5 simulated solar illumination. Performances of photovoltaic cells were measured using a calibrated solar simulator (Oriel 300 W) at 100 mW/cm 2 light intensity adjusted with a PV reference cell (0.5 cm 2 CIGS solar cell, calibrated at NREL, USA). Measurements were performed at an ambient atmosphere. All devices were illuminated through TCO electrodes.

Experimental results

XRD and microprobe analysis have been measured on 100 nm ZnO thick films, while the OPV cells have been achieved with 3-7 nm thick films.

XRD patterns (not shown here) of ZnO:Al films show that they crystallize in the hexagonal system with a wurtzite structure. Their crystallites are preferentially oriented along the c-axis (002) perpendicular to the substrate surface [18]. The spin coated ZnO films (ZnO SC ) are also crystallized in the hexagonal wurtzite structure; however, their crystallites are randomly oriented.

The measurements of films composition have been done by EPMA analysis. This technique allows estimating oxygen atomic concentration, however, the precision of the measure decreases with the atomic weight of the element, therefore, the results show that the films, whatever the deposition technique used, are nearly stoichiometrics with a slight oxygen deficiency (1-3 at.%). For ZnO:Al films 2 at.% of Al is measured.

The surface morphology of the ZnO thin films has been studied by SEM and AFM. The SEM study shows that the surface of the sputtered ZnO:Al is smooth (Fig. 1a), while that of ZnO SC films is inhomogeneous (Fig. 1b). A high density of clusters randomly distributed on the film surface is clearly visible. However, the layer itself appears quite smooth. These films have been submitted to the scotch tape test. Such test is usually used to measure the adhesion of a deposited film to the substrate. Here it is used to clean the deposited film since the ZnO clusters at the surface of the films are not good adhesion to the homogeneous layer of ZnO. It is known that precipitation can take place either in the bulk of the solution with the formation of colloids by an homogeneous reaction or at the surface of the substrates leading to the formation of an homogeneous film by an heterogeneous reaction. A predominantly homogeneous reaction often terminates the heterogeneous, and the film grows by absorption of colloids, leading to superficial heterogeneities. Therefore, the compact layer, which corresponds to the heterogeneous reaction, is covered during the growth process by colloids formed by the homogeneous reaction in the solution bulk (Fig. 1b).

The visualization, in Fig. 2 of the surface of a ZnO SC film, in which half surface (right side) has been submitted to this test, shows that the clusters have been removed by the scotch tape. It allows seeing that the surface of the film itself is smooth and homogeneous. These surfaces, before and after the scotch tape cleaning, have been checked by AFM study. Before cleaning, large ''maximum roughness peaks'', out of scale, are present. After cleaning, one can see in Fig. 3 that the ZnO SC buffer layers are quite homogeneous with a Root Means Square (RMS) surface roughness of 1 nm. The averaged thickness estimated from AFM is around 4 nm.

The conductivity of the ZnO:Al films is around r = 2.32 9 10 3 (X cm) -1 . As expected the ZnOsc films are insulating, with r less than 10 -5 (X cm) -1 . These values have been measured on 100 nm thick film.

After characterization the different ZnO thin films have been used as buffer layer in the inverted organic solar cells ITO/ZnO/BCP/C 60 /CuPc/MoO 3 /Au/Al.

In a first time, we have optimized the ZnO thickness using ZnO:Al films. The ZnO:Al film thickness was 0, 3 and 5 nm. The results are summarized in Table 1. It can be seen that the optimum ZnO:Al thickness is 3 nm. For thicker or thinner ZnO films all the characteristics are degraded. The short circuit current and open circuit voltage decreases significantly.

In the case of ZnO SC , the results depend strongly on the morphology of the anode surface. Following the AFM results, the mean thickness of the films used, was around 4 nm. When the ZnO SC cleaned film is used as buffer layer, the J-V characteristics exhibit diode rectifying effect, while they do not, when it is not cleaned. For such inhomogeneous films, the rectifying effect is destroyed. The clusters, as visible in Fig. 1b, have mean diameter far higher than the whole thickness of the organic active films and therefore they induce strong inhomogeneity in the growth of the organic film with short path, pinholes…, which results in high density of defaults in the organic thin films and small shunt resistance. As a consequence, the J-V characteristics are more or less ohmics.

After cleaning, the film surface is homogeneous and a photovoltaic effect is obtained. Table 2 and Fig. 4 summarize the effect of the nature of the ZnO film on the solar cells performances, it can be seen that the best results are achieved using the conductive ZnO:Al buffer layer. It should be noted that the highest Voc value is obtained with ZnOsc. The ZnOsc being nearly insulating it increases the shunt resistance and therefore the Voc value. However, the high resitivity of the ZnOsc induces also smaller short circuit current and finally smaller efficiency that that obtained with conductive ZnO:Al. From these results, if we exclude the rough ZnO SC , it can be said that the presence of a smooth ZnO film, conductor or not, improves the cells performances even if the best results are achieved with the conducting ZnO:Al.

Discussion

First of all, for classical multilayer heterojunction OPV cells, it is known that one way for circumventing the diffusion length limitation of the excitons is to use cells with multiple interfaces. Peumans et al. [9] have shown that the introduction of a thin large band gap organic material allows improving significantly the device performances. He called electron blocking layer (EBL) this thin film, because its band gap is substantially larger than those of the organic donor and acceptor, which blocks excitons in the organic semiconducting layer far from the cathode avoiding any quenching effect at the interface cathode/organic The EBL confines the photogenerated excitons to the domain near the interface where the dissociation takes place and prevents parasitic exciton quenching at the photosensitive organic/electrode interface. Also it limits the volume over which excitons may diffuse. For these reasons we have introduced a BCP EBL at the interface ITO/electrode acceptor in our inverted OPV cells [16]. This EBL should be transparent to the solar spectrum to act as a spacer between the photoactive region and the metallic cathode. It must, also, transport electrons to avoid high series resistance. If the EBL block the excitons, it should not block all charge carriers. Therefore, the EBL should be chosen so that it allows electrons collection at the cathode. However, the optimum EBL thickness is 8-9 nm, which is too thick to allow high tunnelling current. So it has been suggested that the charge transport in the EBL is due to damage induced during deposition of the upper layers, which introduces conducting levels below its LUMO and explains the reason why the transport of electron is not weakened. [19] (Fig. 5). Therefore, since it is admitted that the EBL is transparent to electron, in a first approach, we will discuss the behaviour of the contact with the cathode as a simple contact ITO/ZnO/organic electron acceptor.

We have shown in a preceding paper [20] that, even if some dipole effect is often present at the interface due to interface states, in a first approach, the behaviour of the contact electrode/organic material can be roughly anticipated using the Schottky-Mott model. Therefore, the different band scheme configurations of the interface anode/ CBL/C 60 , are presented in Fig. 6. The excitons are generated from irradiated CuPc, then they are dissociated by Coulomb attraction to electrons and holes at the interface CuPc/C 60 . The holes may be transported directly to the anode through the ABL, which improves the hole collection efficiency [6][7][8]. At the same time, a good band matching between the cathode work function and the electron acceptor for achieving ohmic contact and good electron collection is necessary.

The treatment of the ITO surface usually has strong influence on the performances of the ITO devices. For example, O 2 plasma treatment increases the work function of ITO, thus enhancing the hole collection in classical organic solar cells. In the present work we have only removed its surface contaminants. However, the work 6a). The collection barrier for electrons from C 60 to ITO corresponds roughly [17] to the energy difference between the work function of ITO and the LUMO of C 60 . The work function of ZnO:Al being 4.3 eV, the electron collection barrier decreased from 0.3 eV for cleaned ITO to nearly 0 for ZnO:Al capped cathode (Fig. 6b). Therefore, the interfacial layer plays an important role in lowering the potential barrier between ITO and C 60 and consequently it should increase the collected electron current and the solar cells efficiency. The significant decrease of the potential barrier between ITO and C 60 is confirmed by the large decrease of the series resistance (Table 2). This ZnO positive effect is obtained only if ZnO is degenerated, which is the case of ZnO:Al, in the case of the insulating ZnOsc the behaviour of the contact is different. Effectively, if we introduce an insulating film, (I), between the degenerated semiconductor (ITO), and the semiconductor, (S), the contact behaves like a MIS structure. It means that there is not band alignment between the TCO, the insulating layer and the organic semiconductor. However, an electric field appears on the insulating layer, which contributes to some decrease of the potential barrier [8]. This results from the fact that the Fermi level must be constant throughout the metal/insulating layer/semi-conductor. Here also it can be seen Table 2 that there is some decrease of the series resistance. However, the most significant information in Table 2 for ZnOsc is the strong increase in the shunt resistance which justifies the high Voc value.

Conclusion

As a conclusion, the improved inverted OPV cells performances can be ascribed to the inserted ZnO interfacial layer. When a conductive ZnO:Al buffer layer is introduced between the ITO cathode and the organic layer it allows a good band matching between the cathode and the LUMO of the organic material, which decreases the potential barrier at the interface cathode/organic material. Such effect is corroborated by the large decrease of the series resistance of these OPV cells. When an insulating ZnOsc is used as buffer layer, its high resistivity increases significantly the shunt resistance of the OPV cells, which increases significantly the open circuit voltage of the OPV cells.

Fig. 1

 1 Fig. 1 Surface visualization of a ZnO:Al (a) and a ZnOsc film (b)

Fig. 2

 2 Fig. 2 Surface visualization of a ZnO film deposited by spin coating: left side before scotch test and right side, after the scotch test

Fig. 4 JFig. 6 a

 46 Fig.4J-V characteristics of organic solar cells using different cathode buffer layers between the ITO cathode and the C 60 electron acceptor

Table 1

 1 Effect of the ZnO:Al thickness on the solar cells performances

	OPV structure	Jsc (mA/cm 2 )	Voc (V)	FF (%)	g (%)
	ITO/BCP/C 60 /CuPc/MoO 3 /Au/Al	3.50	0.355	36.1	0.44
	ITO/ZnO:Al(3 nm)/BCP/C 60 /CuPc/MoO 3 /Au/Al	6.73	0.415	39.7	1.11
	ITO/ZnO:Al(5 nm)/BCP/C 60 /CuPc/MoO 3 /Au/Al	3.11	0.335	38.3	0.40

Table 2

 2 Performances of solar cells using different ZnO CBL layers

	OPV structure	Jsc (mA/cm 2 )	Voc (V)	FF (%)	g (%)	Rs(X)	Rsh(X)
	ITO/BCP/C 60 /CuPc/MoO 3 /Au/Al	3.55	0.350	35	0.40	787	4,096
	ITO/ZnO SC /BCP/C 60 /CuPc/MoO 3 /Au/Al	2.67	0.485	38.2	0.49	375	7,325
	ITO/ZnO:Al/BCP/C 60 /CuPc/MoO 3 /Au/Al	7.46	0.415	49.7	1.54	156	3,125

Acknowledgments This work has been financially supported by the AUF Convention PCSI 6313PS839.


Emmanuel Cadel

Nicolas Barreau
email: nicolas.barreau@univ-nantes.fr

John Kessler

Philippe Pareige

Atom probe study of sodium distribution in polycrystalline Cu(In,Ga)Se 2 thin film

Keywords: Thin film solar cell, Atom probe, Sodium, Grain boundary

come

Atom probe study of sodium distribution in polycrystalline Cu(In,Ga)Se-2 thin film

Introduction

Laboratory-scale solar cells based on CuIn 1Àx Ga x Se 2 (CIGSe) chalcopyrite absorber thin films can reach conversion efficiency close to 20% [1]. However, due to the polycrystalline nature of the CIGSe layers, such high performances could appear unexpected. Although the influence of the grain boundaries on the device operation is still under debate [2][3][4][5][6][7][8][9], several studies suggest that this issue cannot be considered regardless the presence of sodium within the layers [4,6,7]. Indeed, independently of the CIGSe deposition process, the presence of sodium within the CIGSe films is found to improve the solar cells [7,[START_REF] Keyes | In: 26th PVSC[END_REF]. The gain in efficiency (mainly V oc ) is imputed to the increased acceptor density in Na-containing CIGSe [START_REF] Keyes | In: 26th PVSC[END_REF][START_REF] Rockett | [END_REF][12]. Further electrical characterizations also reveal that rather than acting as a dopant, sodium passivates compensating defects [12]. However, the chemical mechanisms leading to the defect passivation are still the seeds for many assumptions. Basically, in bulk CIGSe two crystal point defects are recognized to generate donor states [13]: (i) indium in copper antisite, In Cu and (ii) selenium vacancies V Se . The presence of sodium could hinder the formation of these defects by (i) occupying the copper antisites instead of indium, Na In Cu and (ii) enhancing the chemisorption of oxygen atoms in the selenium vacancies, O Se . In order to evaluate these hypotheses, the first requirement is the location of the sodium within the CIGSe layers. This issue has been addressed by several studies [START_REF] Rockett | [END_REF]14,15] which all suggest that there is sodium at the grain boundaries. Nevertheless, no evidence has been shown on whether the grains contain sodium or not. A reason for this is the lack of available analysis techniques combining both accurate spatial resolution and high elemental detection limits.

Amongst the nano-scale-sensitive analytical techniques (electron holography, electron energy loss spectroscopy), atom probe tomography (APT) is one of the most promising tools thanks to its unique capability to combine atomic scale resolution in three dimensions with quantitative chemical analysis [16,[START_REF] Miller | Atom probe field ion microscopy[END_REF]. For a long time limited to the investigation of conductive materials, mainly metallic alloys, the laser pulsed version of the APT technique is nowadays open to semiconductor materials as shown in the fields of nanoelectronic or photovoltaic devices for instance [START_REF] Kelly | [END_REF][19][20][21][22][23]. The successful use of APT for the analysis of thin film reactions and segregation to linear and planar defects, the aim of the present work in the case of coevaporated CIGSe thin films, has already been demonstrated [24][25][26]. The possibility of sampling all kinds of material configurations (layers, multilayers, thin films, bulks, etc.) using a dual-beam SEM allows new perspectives of selected area atomic scale characterizations. In the present paper, we report for the first time a direct observation of intraand inter-granular distribution of sodium within CIGSe thin films.

Experimental

CuIn 0.75 Ga 0.25 Se 2 polycrystalline thin films have been grown with Na precursor (NaF) on p-doped Si(1 0 0) substrates by the three-stage process at constant substrate temperature (575 °C); note that the investigated CIGSe grown on standard SLG/Mo yielded a 14.2% cell efficiency. One should note that a reference sample has also been prepared without Na precursor. Because of the low density of grain boundaries in CIGSe thin films and the small volume investigated with APT (typically 100 3 nm 3 ), the films have been analyzed perpendicular to the grain growth direction in order to increase the probability of probing grain boundaries. With this aim, as illustrated in Fig. 1, the films have been removed from their silicon substrates using a lift off method, then mounted and glued on a truncated tungsten support-tip using Pt-ion-beam-induced deposition. The whole sample is finally aligned along a Ga-30 kV focused ion beam and successive annular pattern millings transformed it to a very sharp atom probe tip (see Fig. 1). One should note that in order to minimize beam damage, lower energy Ga-2 kV beam was used with low incidence angles at the end of the milling process. Further details on this procedure can be found in Ref. [27]. Atom probe investigations were carried out at 80 K using a CAMECA laser assisted tomography atom probe (LaWaTAP) [28]. The laser system is an amplified ytterbium-doped laser (AMPLITUDE SYSTEM s-pulse, k = 515 nm) used in the following conditions: $100 nJ/350 fs pulse focused onto a $0.1 mm 2 spot. The base pressure within the analysis chamber is 3 Â 10 -10 mbar.

Results and discussion

A typical mass spectrum related to a Na-rich region shown in Fig. 2 illustrates that Na is clearly identified as a single charged ion. For this particular atom probe data set, the noise floor in the region close to the 23 Na peak is estimated at 20 atomic ppm per atom mass unit (a.m.u.). One should note that no signal relative to 23 Na has been detected on the mass spectrum corresponding to the Nafree CIGSe reference sample. The presence of Al + (<0.1 at.%) observable on the mass spectrum could originate from the use of alumina crucibles as metal evaporation sources.

The Na distribution within the reconstructed volume resulting from the APT analysis of the Na-containing CIG-Se thin film is presented in Fig. 3; each black dot corresponds to one detected Na atom. The Na atoms are randomly distributed (typically 85 ± 20 atomic ppm) within the investigated volume at the exception of 2D-like areas where its density is much higher. This observation confirms that, as already suggested in the literature, Na atoms are present within the CIGSe grains. In the present article the attention is focused on the upper Na-enriched zone, referred to as interface in Fig. 3, which extends throughout the whole reconstructed cylinder. The width of this interface has been evaluated to be close to 1.3 nm; however, the actual value is likely to be lower taking into account that a solute enriched zone may be artificially enlarged on APT images because of a local magnification effect. Indeed, the atom probe is a projection microscope Fig. 1. SEM images of the CIGSe thin film and an atom probe sample (tip) extracted by the "lift off" method from the thin film. Fig. 2. Mass spectrum related to ions detected in the region named (A) from Fig. 3. and the presence of a structural defect (such as an interface) emerging at the surface of the sample may produce a slight deviation of atom trajectories [25]. Fig. 4a shows the Na concentration profile measured by moving a 1 nm thick sampling slice by steps of 0.1 nm within the box (A) shown in Fig. 3a. Sampling errors in Na-APT profile (see error bars in Fig. 4a) are given by the standard deviation 2r with r =(C(1 -C)/N) 1/2 with N the number of detected ions contained in the sampling slice and C the solute concentration. This profile illustrates the abrupt increase of the Na density at the interface. Within this interface, the Na concentration reaches 0.89 ± 0.12 at.% and the Gibbsian interfacial excess C Na , defined as the number of segregated Na per unit area, is close to 8 Â 10 13 at. cm -2 . Although no CIGSe atomic planes can be observed in the 3D-APT map (i.e. the existence of a misorientation between the two grains cannot be demonstrated), this interface is assumed to be a grain boundary separating two grains, namely grain 1 and grain 2.

In Fig. 4a are also shown the composition profiles of Cu, In, Ga and Se. This plot suggests that the Cu, Ga and Se contents are different in grain 1 and grain 2, whereas that of In appears similar in both grains. No significant enrichment or depletion of Cu, In, Ga and Se is detected through the grain boundary. One should nevertheless be cautious before drawing conclusions from these latter observations. Indeed, this film has also been investigated by energy dispersive X-ray spectroscopy (EDS) and its composition is 24.3 at.% Cu, 21.5% In, 4.2% Ga, 50.0% Se, which is, as expected, near the pseudo binary Cu 2 Se-In 2 Se 3 tie line. The compositions deduced from APT and EDS are compared on a ternary phase diagram (Cu, group III, Se) in Fig. 4b. This graph clearly shows that selenium is under-estimated in our APT measurement, suggesting that some Se ions are not detected during the field evaporation of the specimen by laser pulses. This effect is known as preferential field evaporation: ions can sometimes be evaporated between two laser pulses and therefore not be detected, which results in differences between the expected and the measured compositions [START_REF] Miller | Atom probe field ion microscopy[END_REF]. Preferential field evaporation is mainly observed for alloys containing elements with high evaporation field differences, which is actually the case in CIGSe; it is reported for singly charged ions the following field values: F Cu =30Vnm -1 , F In =12Vnm -1 , F Ga =15Vnm -1 whereas Se is not tabu- lated [START_REF] Miller | Atom probe field ion microscopy[END_REF]. Instrumentation parameters such as pulse laser energy (i.e. pulse fraction) and specimen base temperature must therefore be further optimized in order to ensure accurate quantitative measurements. However, such an artifact only slightly affects the qualitative evolution of the profiles throughout the investigated volume. Therefore, the difference observed in the adjacent grain composition is most probably effective. Moreover, the conclusions concerning the presence of Na within the grains and its accumulation at the grain boundary remain valid. Unfortunately, the present investigations do not allow the determination of the crystal sites occupied by the Na neither within the CIGSe grain structure nor at the grain boundary. Nevertheless, within the grain, the Na concentration being very low, the calculations reported by Wei et al. [13] predict that these atoms most probably occupy the Cu antisites. Such a location could minimize the density of In Cu donor defects, thus effectively improving the acceptor density as experimentally observed.

Conclusions

In the present article the first results of APT investigations on CIGSe polycrystalline layers are presented. Both the presence of Na within the grains and its accumulation at the grain boundaries are demonstrated, which confirms many of the conclusions suggested in the literature. This study shows that the use of APT can be a powerful tool for the analysis of the CIGSe.

Fig. 3 .

 3 Fig. 3. (a) Spatial distribution of Na atoms (black dots) within CIGSe thin film. (b) View of the same three-dimensional atom probe map of Na atoms with a 60°rotation around the z-axis.

Fig. 4 .

 4 Fig. 4. (a) Cu, In, Ga, Se and Na concentration profiles through the region (A) of Fig. 3. (b) APT and EDS measurements of CIGSe grains composition. The dashed lines represent the Cu 2 Se-In 2 Se 3 pseudobinary and the Cu-rich/III-rich dividing line (vertical).


J B Leriche

S Hamelet

J Shu

M Morcrette

C Masquelier

G Ouvrard

M Zerrouki

P Soudan

S Belin

E Elkaïm

F Baudelet

An Electrochemical Cell for Operando Study of Lithium Batteries Using Synchrotron Radiation

A new electrochemical cell has been specially designed for operando experiments at synchrotron facilities both for X-ray diffraction and X-ray absorption. It allows the investigation of insertion materials under high current densities ͑up to 5C rate͒ and hence to study complex phenomena of structural and electronic changes out of equilibrium. The LiFePO 4 -FePO 4 system has been chosen as a case study to validate this cell, and tricky phenomena, with apparent delays in phase formation compared with the number of electrons exchanged, have been spotted.

The study of the structural and electronic evolution of lithium battery electrodes during charge/discharge processes is crucial to understand the Li storage/release mechanisms and to optimize these materials so as to achieve high performance and cyclability. To this end, in the past 20 years, several in situ and ex situ techniques, such as X-ray diffraction ͑XRD͒, [1][2][3][4][5][6][7][8][9][10][11] X-ray absorption spectroscopy ͑XAS͒, [12][13][14][15] and Mössbauer, 16 Raman, IR, and NMR 17,18 spectroscopies, have been developed.

In situ evaluation of battery materials, i.e., observation inside a closed electrochemical cell, brings online information and eliminates the risk of manipulating highly reactive powders with ambient atmosphere. It allows the investigation of complex reaction mechanisms and proves to account very satisfactorily for voltagecomposition profiles in various chemical systems due to structural and electronic transitions within the electrode͑s͒.

In situ XRD investigations can be performed both in standard laboratory diffractometers and synchrotron source facilities, the latter providing X-ray beams with a photon flux higher by several orders of magnitude than the ones delivered by conventional X-ray tubes. To this end, several kinds of electrochemical cells have been designed both for reflexion or transmission geometries. The recent development of high quality position sensitive detectors in standard X-ray diffractometers allows easier use of such techniques in laboratories. Quite universal in terms of materials to be investigated ͑crystalline and amorphous powders͒ and of phenomena to be described ͑structural changes and electronic transfer͒, in situ XAS experiments have been widely performed recently both in extended X-ray absorption fine structure ͑EXAFS͒ and X-ray absorption near edge structure ͑XANES͒ modes. For instance, although the EXAFS part of the signal provides direct structural information on the close environment of a given atom ͑selected by its own absorption en-ergy͒, the XANES part of the spectrum can be roughly seen as a picture of the empty electronic states of a given atom and allows the monitoring of the filling of these levels ͑charge-transfer processes͒ during lithiation ͑and reverse͒. Recent approaches using band structure calculations and data simulation were very successful in allowing a precise analysis of chemical bonds during electrochemical lithium insertion/extraction. 19 Additionally, the development and use of bended single crystals in synchrotron facilities ͑dispersive X-ray absorption fine structure͒ and the possibility of fast rotation of the monochromator ͑QuickXAS͒ paved the way for new approaches for the study of battery materials at synchrotron facilities. The possibility of using very short acquisition times, typically on the order of a few seconds for both XRD and XAS, indeed permits us to investi-gate structural and electronic properties of electrodes out of equilibrium ͑contrary to standard in situ studies͒ at various current loadings, which we further refer to as the operando mode.

In this paper, we report on the development and use at the Synchrotron facility SOLEIL ͑France͒ of a new, reliable, and easy-touse electrochemical cell designed for photon transmission or reflection geometry ͑both XRD and XAS͒ that allows fast ͑up to 5C, 850 mA g -1 of active material͒ and highly reversible electrochemical cycling. To perform experiments in different energy ranges, with various time and spatial resolutions, our goal, in the long term, is to use four different beamlines on the third-generation synchrotron source SOLEIL. These beamlines are CRISTAL for XRD, ligne utilisée pour la caractérisation par imagerie et absorption ͑LUCIA͒ for high spatial resolution and medium energy range XAS, spectroscopies applied to materials based on absorption ͑SAMBA͒ for classical and Quick XAS experiments at higher energy, and optique dispersive EXAFS for dispersive XAS ͑high time and spatial resolutions͒. 20 High rate LiFePO 4 powders are used as an illustrative example of positive electrode materials for Li-ion batteries to validate this cell working operando with synchrotron radiation.

Description of the Electrochemical Cell

One of the objectives of this study is to develop, so as to make it available to the scientific community, a user-friendly electrochemical cell to be used in synchrotron radiation facilities for both XRD and XAS experiments in transmission or reflection geometry. By user-friendly, we mean that the cell͑s͒ should be mounted easily by the user in a standard Ar-filled glove box by stacking the positive electrode, the separator ͑soaked with electrolyte͒, and the negative electrode layer by layer in-between two XRD-transparent Be windows ͑the manipulation of which necessitates extreme care and knowledge of risks associated with its oxide, BeO͒. Additionally, it should be highly hermetic, stable in time, dismountable, reusable, easy to clean, easy to mount within the beamline, and easy to connect to the cycling potentiostat.

To this end, we extensively built and tested a stainless steelbased cell derived from the Swagelok configuration ͑Fig. 1͒. mostly transparent to X-rays ͑only 14.6% of the beam is absorbed at 5 keV and 0.71% at 20 keV͒. When positive electrode materials with operating voltages greater than 3.5 V vs Li are to be used, a thin ͑3-4 m͒ aluminum foil is placed onto the Be window so as to protect it from oxidation. This part of the cell is then assembled to the main body with a rubber o-ring in contact with the Be window and through six stainless steel screws.

2. The second important piece of the cell ͑specially dedicated to transmission͒ is the cylindrical plunger ͑described in Fig. 1 as parts A, B, and C͒ that constitutes in most cases the "negative" part of the cell. As for the positive part, the plunger was designed so as to allow easy replacement of the Be window ͑200 m thick, 2 cm in diameter in this case͒. Essential for the proper adjustment of the mechanical pressure within the cell, the two parts ͑A and B͒ of this plunger can translate along each other from the use of a spring. In transmission geometry, the incoming X-rays enter part A of the plunger ͑2 cm wide͒ through a cylindrical hole ͑1 cm wide, much larger than the beam size͒ and pass through the small Be window held with an o-ring between parts B and C. Part C is a capped cylinder with a rather small hole of 4 mm through which the beam goes before hitting the active part of the cell ͑Li/separator/positive electrode͒.

Electrochemical Tests

To validate this cell as a reliable tool for in situ/operando experiments at synchrotron facilities, we undertook, before any XRD or XAS experiment, a systematic electrochemical characterization of several positive electrode materials such as LiMn 2 O 4 , LiFePO 4 , and LiVPO 4 F, which operate between ϳ2.5 and 4.5 V vs Li.

The positive electrode materials, mixed with 10 wt % of carbon SP and poly͑tetrafluoroethylene͒ binder, were deposited on 4 m aluminum foil as plastic films for typically ϳ6 mg of active material/cm 2 . Li metal was used as the counter ͑negative͒ electrode, and 1 M LiPF 6 in ethylene carbonate/diethyl carbonate ͑1:1, v/v͒ as the electrolyte impregnated in one layer of Whatman glass fiber or Celgard polymer was used as separators.

Electrochemical responses were recorded in a systematic way under several charge/discharge rates up to 5C ͑850 mA/g of active material, i.e., charge/discharge in 12 min͒ and demonstrated very good stability of the cell for at least 10 cycles, as shown for three representative voltage/composition data plotted in Fig. 2. Each experiment was repeated three to five times and showed very good reliability and reproducibility, essential for taking full advantage of allocated beam time at synchrotron facilities without worries for the electrochemistry part of the operando experiments.

Operando XRD on LiFePO 4 -Based Powders

Triphylite LiFePO 4 , a material of choice to be used as a positive electrode in the next generation of Li-ion batteries [21][22][23][24][25][26] particularly suited for applications in power tools and hybrid electric vehicles, was chosen as the first demonstrative example of the feasibility of obtaining high quality XRD data in very short periods of time. We investigated several pristine powders of LiFePO 4 with different particle sizes, morphologies, and stoichiometry. The electrochemical cell was mounted for transmission geometry and connected to a galvanostat/potentiostat ͑VMP3 of Biologic͒. The incident X-ray wavelength chosen was 0.727 Å, allowing easy transmission through the many components of the cell, and the data were collected on an image plate detector, allowing very small recording times of 10 s per pattern. Altogether, seven batteries were cycled very reliably for an overall amount of ϳ2000 patterns recorded, from which the first series of interesting phenomena was spotted, as described below.

Be

Figure 3 shows, for instance, the electrochemical and XRD data recorded operando during a charge lasting nearly 2 h during which 38 patterns were recorded ͑1 pattern of 10 s every 3 min͒. Classically described in the literature as a two-phase reaction between LiFePO 4 and FePO 4 , the process of Li deintercalation appears to occur here with a significant delay; i.e., the diffractogram recorded at midcharge ͑20͒ is very far from being characteristic of a theoretical equal amount of both LiFePO 4 and FePO 4 phases. This apparent "crystallization delay" for the appearance of FePO 4 at the expense of the disappearance of LiFePO 4 was very recently observed by two other groups. [27][28][START_REF] Chiang | ICMAT Conference[END_REF] Chang et al. 27 interpreted their observations as the result of slow nucleation kinetics of the resulting phase, while process disturbances tend to stimulate the formation of nuclei. Chiang et al. 28,[START_REF] Chiang | ICMAT Conference[END_REF] envisaged that the delay in the appearance of crystallized FePO 4 was due to the existence of transient intermediate amorphous FePO 4 components, the amount of which is dependent on ͑i͒ the particle size, ͑ii͒ the applied electrical overpotential, and ͑iii͒ the amplitude of the misfit strain between the lithiated and delithiated crystalline phases.

We carefully analyzed by full pattern matching refinements the variation in the intensities of several sets of diffraction peaks ͑several crystallographic directions͒ under two charge/discharge rates, i.e., within2hofcharge ͑C/2; 85 mA/g of LiFePO 4 ͒ or 10 h ͑C/10; 17 mA/g of LiFePO 4 ͒. Much care was taken in selecting an appro-priate background profile, in introducing anisotropic crystallite size parameters, and in normalizing the diffracted intensities vs the diffraction peak of Be used as a standard here. Figure 4a illustrates the quality of both the diffraction data and the refinements from which we were able to extract the variation in the normalized intensities of several sets of diffraction peaks of the two observed crystalline phases, i.e., Li 1-FePO 4 and Li Ј FePO 4 . Figure 4b illustrates, for instance, the decrease in the intensity of the ͑200͒ Pnma peak of LiFePO 4 as a function of x Li + overall content for both C/10 and C/2 rates. It shows a significant experimental deviation from the theoretical 50% decrease in intensity at midcharge ͑dotted line͒ more obvious at C/2 than at C/10. In an apparent contradiction with the recent interpretation of Chiang et al., 28,[START_REF] Chiang | ICMAT Conference[END_REF] the appearance of ͑200͒ FePO 4 is also delayed but "grows" experimentally at the same rate as ͑200͒ LiFePO 4 decreases ͑Fig. 4c͒. Indeed, the variations in the normalized intensities of both Li 1-FePO 4 and Li Ј FePO 4 are somehow symmetrical and, as a consequence, cross each other at 50% of relative intensity for both phases. This observation suggests that, all over the charge process, only two crystalline phases are present, without the formation of a third amorphous component in the system. Figure 4d shows that this observed phenomenon is less pronounced at a C/10 rate. This prompted us to address porous electrode effects in more detail, which are under experimental investigation as of this date.

To get a deeper insight into this apparently new mechanism, we tried to precisely analyze the variation in the full width at halfmaximum of the diffraction peaks ͑hence, the variation in crystallite sizes in several crystallographic directions͒, without success unfortunately, as the resolution of the diffracted pattern was not high enough ͑price to pay for fast acquisition with the image plate͒. This very important issue will be addressed very soon through a series of experiments designed for recording diffraction patterns with a higher instrumental resolution and a higher crystallinity of the LiFePO 4 powder.

Operando XAS on LiFePO 4 -Based Powders

In addition to XRD, the technical development of our cell allowed us to investigate the local electronic and atomic structure of the LiFePO 4 electrodes at the thermodynamic equilibrium and during the charge and discharge process ͑operando͒ using in situ XAS at the Fe K edge. XAS data were collected in the transmission detection mode. The data reduction of the XAS spectra was performed using the software program ATHENA for Windows. [START_REF] Ravel | [END_REF] Stan- dard procedures given in these programs were used to obtain the normalized and calibrated spectra. The thermodynamic equilibrium study was carried out on the SAMBA beamline using a Si͑111͒ double-crystal sagittal focusing monochromator. The size of the beam was 300 ϫ 500 m. The XAS measurements were recorded using three ionization chambers in a series to measure the intensities of the incident beam, the beam transmitted by the sample, and the beam transmitted by the Fe reference foil. For each ⌬x = 0.2 corresponding to the lithium ions extracted/inserted during charge/discharge, the electrode was stored at an open-circuit voltage for 30 min after which the XAS spectra were recorded ͑Fig. 5͒. For this experiment, a rate of 1C ͑170 mA/g of active material͒ was used ͑charging curve plotted in the insert of Fig. 5͒. The superposition of XAS spectra recorded at various states of charge displays clear isosbestic points showing a two-phase process at the spatial scale used. These results agree well with the literature. 12 Within the XANES spectra, the edge positions shift progressively to higher energies during the charge process as a result of the oxidation of Fe 2+ to Fe 3+ . 31 As a consequence of the biphasic nature of the electrochemical process, each XANES spectrum is a linear combination of the LiFePO 4 and FePO 4 spectra with approximately the same percentage as in the electrochemical data. For example, the best linear combination fit for the Li 0.4 FePO 4 composition ͑0.6 Li extracted͒ yields 37.4% of LiFePO 4 and 62.6% of FePO 4 .

For dynamic studies, XANES data have been collected using a charge-coupled device camera on the ODE beamline dedicated to energy-dispersive XAS. In this setup, each spectrum is recorded during 10 s. Some spectra collected at different Li contents during charge at the 2C rate ͑340 mA/g of LiFePO 4 ͒ are gathered in Fig. 6a. As for the data recorded at equilibrium ͑Fig. 5͒, XANES spectra show isosbestic points confirming the biphasic character of the process with, again, an apparent delay during charge and discharge of the battery compared to the number of Li + extracted electrochemically. Indeed, during discharge/charge, each spectrum is a linear combination between LiFePO 4 and FePO 4 ͑taken at the equilibrium on the SAMBA beamline͒. Figure 6b shows the simulation done for a global composition of Li 0.36 FePO 4 ͑0.64 Li + extracted͒ for which the best fit gives 61% of LiFePO 4 and 39% of FePO 4 ͑0.39 Li + extracted͒. The extraction process seems to occur with a delay with reference to the electrochemistry. This delay confirms our results obtained by XRD.

Conclusion

High quality XRD and XAS data can be obtained by using this in situ cell to reveal the complex nature of the transformation dynamics during the charge/discharge of electrodes for lithium-ion batteries. Such a cell allows the operando observation of these phenomena and will be used in the near future as a powerful tool to investigate possible inhomogeneities in porous electrodes.

Figure 1 .Figure 2 .

 12 Figure 1. ͑Color online͒ Electrochemical cell for operando measurements. ͑a͒ Photo and ͑b͒ detailed view of the cell with incoming and outcoming beam paths in transmission geometry, and ͑c͒ incoming and ͑d͒ outcoming beam paths in reflexion geometry.

Figure 3 .Figure 4 .

 34 Figure 3. ͑a͒ Electrochemical data and ͑b͒ XRD patterns recorded operando during a charge at C/2 of LiFePO 4 .

Figure 5 .Figure 6 .

 56 Figure 5. ͑Color online͒ Normalized absorption at the iron K edge ͑recorded at the equilibrium͒ for various states of charge within the LiFePO 4 /FePO 4 system. The top insert is the corresponding charge curve, and the bottom insert is an expanded view of the pre-edge part.

The top part of the cell ͑used in most cases as the "positive" side of the cell vs metallic Li͒ is composed of a rigid, 5 cm large, outer ring with a hole ͑2 cm in diameter͒ behind which a metallic Be window ͑200 m thick, 4 cm in diameter͒ is placed. This part is very similar to the cell already described in our recent papers6,9 dealing with "standard" in situ XRD investigations within laboratory diffractometers using either Cu K␣ or Co K␣ radiations in reflection geometry. The electrode material to be investigated ͑either in powder, dried slurry, or plastic film form͒ may be placed directly behind the Be window, which acts as a metallic current collector z E-mail: christian.masquelier@u-picardie.fr

Acknowledgment

We thank Professor Jean-Marie Tarascon and Dr. Charles Delacourt for fruitful discussions. The present work was undertaken within the ANR program Stock-E under the contract ANR-07-Stock-E-09-01. We also acknowledge the financial support of UMICORE ͑Belgium͒ for SH.

Centre National de la Recherche Scientifique assisted in meeting the publication costs of this article.