L Stermann

G Simon

L Vanel

D Tanguy
email: dome.tanguy@univ-lyon1.fr

In situ measurement of plasticity accompanying Hydrogen induced cracking

Keywords: Micro-mechanics, Hydrogen, Embrittlement, Digital Image Correlation PACS 62.20.Mk

Single crack propagation is studied in a Hydrogen embrittled aluminum alloy. Hydrogen is introduced in the system by electrochemical reactions in an acid aqueous medium. After Hydrogen charging, tensile tests are performed in air, on notched samples, with a micro-tensile machine. An optical microscope is used to follow single crack initiation and propagation at a high magnification. Digital Image Correlation gives the displacement field on the surface with a spatial resolution of about 1 micron. It enables the determination of the position of the crack tip and the local velocity at a sub-grain scale. The equivalent Von Mises strain is calculated and gives a precise measure of the local plastic field which accompanies crack propagation. In addition to the primary plasticity which is emitted from the crack tip or its immediate neighborhood in the form of two intense slip bands, it is systematically found a secondary plastic zone which spreads over several microns ahead of the tip. The characteristics of the plastic zone are measured, together with the velocity and the applied stress intensity factor. In addition, different fracture mechanisms are found on the fracture surface. In particular there are transitions in the fracture mode from intergranular smooth to transgranular parallel to the grain boundary plane. The local fracture mechanisms, in the vicinity of the surface, are linked to the local velocities and plastic deformations. Surprisingly no strong velocity/plasticity correlations are found while the velocities are scattered over a wide range.

Introduction

Environmentally assisted cracking, including hydrogen embrittlement, is a major concern for structural metallic alloys because it causes premature failures at mechanical loads well below those required for purely mechanical fracture. In addition, there is a recent increase of interest for this problem in the context of the reduction of CO 2 gas emission either by the deployment of fuel cell technology or by the development of hydrogen engines for airplanes. The mechanisms at the origin of this phenomenon are intricate and diverse, depending on the material, the environment and the way the mechanical load is applied [START_REF]Pydic is a free suite of Python tools for local digital image correlation analysis[END_REF]. Corrosion induced anodic dissolution can be mechanically localized, for example at the surface steps where slip bands emerge [START_REF]Pyramidal Implementation of the Lucas Kanade Feature Tracker description of the algorithm[END_REF], or chemically localized, for example at grain boundaries hosting segregations or fine precipitation with one element being preferentially dissolved [START_REF] Ai | Hydrogen diffusivity during corrosion of high-purity aluminum[END_REF]. This might cause crack initiation. The possibility that localized anodic dissolution is also responsible for fast, crystallographic, transgranular crack propagation, i.e. beyond the initiation stage, in conjunction with localized plasticity or cleavage [START_REF] Ai | Effective hydrogen diffusion in aluminum alloy 5083-H131 as a function of orientation and degree of sensitization[END_REF], is still debated (see the references in [START_REF] Badwe | Decoupling the role of stress and corrosion in the intergranular cracking of noble-metal alloys[END_REF]). Recently, atomic scale observations demonstrated that a strained nanoporous layer, obtained by preferential dissolution of a concentrated binary alloy, can inject a cleavage intergranular crack in the non-dealloyed material, microns ahead of the layer [START_REF] Ali | Caractérisation et modélisation micromécanique de la propagation de fissures fragiles par effet de l'hydroène dans les alliages AA7XXX[END_REF]. This mechanism could be active, locally, in alloys with intense intergranular segregation. In addition to this complexity, anodic dissolution often implies the reduction of hydrogen coming from the dissociation of water, in aqueous media. A fraction of the hydrogen produced is absorbed and diffuses in the material where it causes embrittlement [7]. Here also, the mechanism of fracture is not completely established [8]. Recent high-resolution experimental observations concern: hydrogen-induced dislocation cell refinement in the plastic wake [START_REF] Bendo | Atomic scale haadf-stem study of η' and η 1 phases in peak-aged AlZnMg alloys[END_REF], quantified plasticity localization in deformation bands [START_REF] Bertsch | Hydrogen-induced compatibility constraints across grain boundaries drive intergranular failure of Ni[END_REF], the formation of vacancy-hydrogen clusters revealed by the pinning/depinning dynamics of dislocations in nano pillars [START_REF] Bond | On the determination of the hydrogen fugacity in an environmental cell tem facility[END_REF]. In particular, the fracture of micro-beams in the presence of water vapor inside an environmental SEM, coupled with transmission electron backscatter diffraction, revealed the intensity of plasticity accompanying brittle transgranular fracture with nanometer scale resolution [START_REF] Christodoulou | Hydrogen embritlement and trapping in Al-6%-Zn-3%-Mg[END_REF]. On the modeling side, atomistic simulations have contributed by evaluating several elementary mechanisms such as: the modification of dislocation emission at crack tips by adsorbed hydrogen [START_REF] Deng | Hydrogen embrittlement revealed via novel in situ fracture experiments using notched micro-cantilever specimens[END_REF][START_REF] Du | Hydrogen-enhanced vacancy diffusion in metals[END_REF], the enhancement of equilibrium vacancy concentrations [START_REF] Du | Diffusion of hydrogen within idealized grains of bcc Fe: A kinetic Monte Carlo study[END_REF][START_REF] Dumont | Characterisation of the composition and volume fraction of η' and η precipitates in an AlZnMg alloy by a combination of atom probe, small-angle Xray scattering and transmission electron microscopy[END_REF], the formation of bubbles by aggregation of vacancies produced by dislocations and stabilized by hydrogen [START_REF] Flanagan | What do we understand about stress corrosion cracking in aqueous environments?[END_REF], the decrease of cohesion related to the segregation of hydrogen [START_REF] Flanagan | Authorsŕeply[END_REF][START_REF] Gangloff | Hydrogen assisted cracking of high strength alloys[END_REF], vacancy-hydrogen clusters [START_REF]Gaseous Hydrogen Embrittlement of Materials in Energy Technologies[END_REF], voids [START_REF] Girardin | Correlation between dislocation organization and slip bands: TEM and AFM investigations in hydrogen-containing nickel and nickelchromium[END_REF] or dynamic segregation [START_REF] Gjønnes | An electron microscope investigation of the microstructure in an aluminium-zincmagnesium alloy[END_REF] at grain boundaries. Diffusion of vacancies at low [START_REF] Gupta | Relating the early evolution of microstructure with the electrochemical response and mechanical performance of a Cu-rich and Cu-lean 7XXX aluminum alloy[END_REF] and high hydrogen concentration [START_REF] Holroyd | Environment-induced cracking of high strength aluminium alloys[END_REF] and trapping of hydrogen are also very important issues that have been intensively studied by atomistic simulations [START_REF] Holroyd | Improved understanding of environment-induced cracking (EIC) of sensitized 5XXX series aluminium alloys[END_REF][START_REF] Holroyd | Crack propagation during sustained-load cracking of Al-Zn-Mg-Cu aluminum alloys exposed to moist air or distilled water[END_REF], including pipe diffusion at dislocations [START_REF] Ilin | Simulation of hydrogen diffusion affected by stressestrain heterogeneity in polycrystalline stainless steel[END_REF].

In this context, we develop new experiments with high magnification in situ observations of the fracture processes, that could be useful for developing mesoscale models capable of bridging the gap between atomistic simulations and macroscopic experiments. In this paper, an aluminum alloy is charged with hydrogen and strained with a micro-tensile stage to obtain single cracks. Digital image correlation is used to measure the crack advance and the related plastic deformation at the micron scale.

Aluminum alloys are interesting model systems because they can be prepared in metallurgical states where they are sensitive enough to hydrogen to promote single cracks facilitating in situ monitoring. More specifically, there is a wealth of studies [START_REF] Ismer | Interactions between hydrogen impurities and vacancies in Mg and Al: A comparative analysis based on density functional theory[END_REF] in the AlZnMg and AlMg [START_REF] Knight | Some effects of alloy composition on stress corrosion cracking in AlZn-MgCu alloys[END_REF] aluminum alloys. They are immune to dry hydrogen gas because of the high dissociation barrier of the molecule on the oxide [START_REF] Lervik | Atomic structure of solute clusters in AlZnMg alloys[END_REF]. Exposure to a plasma [START_REF] Li | The interaction of dislocations and hydrogen-vacancy complexes and its importance for deformation-induced proto nano-voids formation in α-fe[END_REF][START_REF] Lozano-Perez | SCC in PWRs: Learning from a Bottom-Up Approach[END_REF], which skips this dissociation step, shows hydrogen penetration is effective with the formation of bubbles at the oxide-metal interface [START_REF] Lynch | Discussion of some recent literature on hydrogen-embrittlement mechanisms: adressing common misunderstandings[END_REF] or the formation of vacancy-hydrogen clusters further inside the material in nano pillars [START_REF] Bond | On the determination of the hydrogen fugacity in an environmental cell tem facility[END_REF]. In aqueous media, hydrogen comes from electrochemical reactions at the surface in-volving the dissociation of the water molecule [START_REF] Lervik | Atomic structure of solute clusters in AlZnMg alloys[END_REF] and the subsequent absorption of a fraction of the produced hydrogen. The result is a degradation of the mechanical properties: on macroscopic specimen, a mechanical threshold appears at a stress intensity factor (K) about 5 to 10 MPa √ m, well below the critical K in the absence of hydrogen [START_REF]Corrosion-Deformation Interactions[END_REF]. Beyond this threshold, the crack velocity shows a weak dependence on K, until the critical value for catastrophic failure is reached. Water vapor (including lab air), in particular hot water vapor, and chloride containing solutions, are hydrogen producing environments. At room temperature and low relative humidity, crack propagation occurs at velocities of the order of 10 -10 ms -1 that continuously grow with relative humidity until condensation occurs reaching 10 -8 ms -1 (alloy 7075, T651 state) in distilled water (see figure 19.12 and references in [START_REF] Lervik | Atomic structure of solute clusters in AlZnMg alloys[END_REF]). Metallurgical state influences this velocity by one order of magnitude, underaged (resp. overaged) being more sensitive (resp. less sensitive) [START_REF]Corrosion-Deformation Interactions[END_REF]. Exposure to hot water vapor followed by subsequent tensile test in air results in hydrogen embrittlement [START_REF] Magnin | The corrosionenhanced plasticity model for stress corrosion cracking in ductile fcc alloys[END_REF]. Fracture morphology can be either transgranular or intergranular [START_REF] Magnin | The corrosionenhanced plasticity model for stress corrosion cracking in ductile fcc alloys[END_REF][START_REF] Malitckii | Hydrogen uptake from plasma and its effect on EUROFER 97 and ODS-EUROFER steels at elevated temperatures[END_REF]. Hydrogen bubbles have been observed at intergranular precipitate/matrix interface in TEM but are assumed to decrease embrittlement by trapping hydrogen otherwise present along the grain boundaries [START_REF] Magnin | The corrosionenhanced plasticity model for stress corrosion cracking in ductile fcc alloys[END_REF][START_REF] Mcmurtrey | Mechanism of dislocation channel-induced irradiation assisted stress corrosion crack initiation in austenitic stainless steel[END_REF]. Coarse intergranular precipitates can indeed have a beneficial effect on intergranular hydrogen embrittlement [START_REF] Malitckii | Hydrogen uptake from plasma and its effect on EUROFER 97 and ODS-EUROFER steels at elevated temperatures[END_REF].

The paper is organized as follows: the choice of the material and its preparation are presented, then the digital image correlation (DIC) methodology is illustrated by an example showing the plastic blunting of an intergranular crack and finally the main results are given, including: the measured deformation state at initiation at the notch tip, the different fracture morphologies and the statistics of the local velocities, plastic zone sizes and plastic deformation. Some correlations that were attempted are discussed.

Material and methods

The alloy chosen is the aluminum alloy 7108 which chemical composition is given in table 1. It was received as a rolled plate of thickness 8 mm with the microstructure shown on figure 1a. AlZnMg alloys harden by precipitation. The present precipitate phases are the solute clusters, called "Guinier-Preston zones", on the lattice sites of aluminum [START_REF] Mütschele | Segregation and diffusion of hydrogen in grain boundaries of palladium[END_REF], the semi-coherent metastable η' (MgZn 2 ) and the stable incoherent η phases [START_REF] Nagao | Hydrogen-enhanced-plasticity mediated decohesion for hydrogen-induced intergranular and "quasi-cleavage" fracture of lath martensitic steels[END_REF]. Figure 1a shows intragranular precipitates in the shape of plates and rods which are likely different forms of η [START_REF] Nazarov | Ab initio study of H-vacancy interactions in fcc metals: Implications for the formation of superabundant vacancies[END_REF] precipitates in an overaged state, owing to their large size in comparison to the peak aged microstructure [START_REF] Robertson | Hydrogen embrittlement understood[END_REF]. The figure also shows larger, globular, precipitates along the grain boundaries and a precipitate free zone as reported in [START_REF] Sauzay | Prediction of grain boundary stress fields and microcrack initiation induced by slip band impingement[END_REF].

The copper content (Tab. 1) is significantly lower than in other commercial alloys where it is around 2.5%. Cu is reported to enter the composition of the intergranular MgZn 2 precipitates, increase their density, reduce the localized anodic dissolution rate and accordingly the stress corrosion cracking plateau velocity [START_REF] Scamans | Pre-exposure embrittlement and stress corrosion failure in AlZnMg alloys[END_REF], in a saturated NaCl solution. In other environments, this trend is not confirmed [START_REF] Scully | Gaseous Hydrogen Embrittlement of Materials in Energy Technologies[END_REF]. Moreover different quench rates can also have a drastic effect on the plateau velocity and the link with microstructure is not established [START_REF] Scully | Gaseous Hydrogen Embrittlement of Materials in Energy Technologies[END_REF]. Even with these limitations, we consider that having a reduced Cu content reduces the protective effect Cu segregation could have on GBs with respect to anodic dissolution in wet air conditions, which are the conditions of our tensile tests. In a previous study [START_REF] Shen | Segregation of hydrogen and vacancies at the σ5(210)[001] symmetric tilt grain boundary in Ni and influence on cohesion[END_REF], it was also shown numerically that the presence of generic brittle precipitates (representing H embrittled η precipitates [START_REF] Magnin | The corrosionenhanced plasticity model for stress corrosion cracking in ductile fcc alloys[END_REF][START_REF] Mcmurtrey | Mechanism of dislocation channel-induced irradiation assisted stress corrosion crack initiation in austenitic stainless steel[END_REF]) has some minor effect on intergranular cleavage, as modeled by cohesive zones. As a consequence, the choice was made to simplify the microstructure by heat treating the samples, after machining, to eliminate the intergranular precipitates. Samples were held at 400 • C for 1h, followed by a water quench and natural aging. The result is the dissolution of all the precipitates of the as received microstructure and a recrystallisation (Fig. 1b). Some grains are equiaxe, others kept an elongated aspect ratio (Fig. 2a). To take advantage of the latter, the micro tensile samples are machined (by electrical discharging machining) along the short direction (S) (Fig. 2b). The gauge length is 4 mm, the width either 1 mm or 1.5 mm and the total length including the heads 8 mm. A U notch, of initial thickness U 0 =70 µm, is cut in the plane normal to S with a front along the long direction (L). As a consequence, the number of grains along the crack front is minimized. The dimension of the samples in this direction (the sample's thickness B) is a parameter. Ideally, the samples could be thinned to have only one grain in the thickness so that the slip orientation observed on the side surface are representative of the crystallography all along the crack front. The real grain structure is far from this idealness. The grain size is heterogeneous (Fig. 2a) as observed on a EBSD map in the S-TL plane and on the fracture surfaces on the L-TL plane, with some equiaxed grains of diameter of the order of 20 µm and others keeping an elongated aspect with dimensions up to 200 µm in the L direction. The average grain thickness (the dimension in the S direction) is 17 µm (+/-9 µm). Finally, there are indeed portions of the samples with long, flat intergranular facets oriented perpendicular to the S direction, but usually, the thickness contains many grains, some of them equiaxe. The thickness B is varied between 0.2 mm and 1 mm.

The samples are polished mechanically with diamond pastes containing particles of diameter down to 1 µm. Hydrogen is introduced in the material by an electrochemical process. The samples are immersed in a sulfuric acid solution (pH 2.0) in a 3-electrodes cell. A cathodic potential E c is imposed during a charging time t c . All the electrochemical potentials are measured with respect to a reference which is a saturated calomel electrode (SCE). The mechanical tests are done in lab air after the sample is taken out of the acid and dried. The relative humidity is in between 40 and 60% and the temperature about 22 • C in the room.

H induced cracking is strain rate dependent [START_REF] Shimizu | Local deformation and fracture behavior of high-strength aluminum alloys under hydrogen influence[END_REF][START_REF] Sun | Quantum mechanical modeling of hydrogen assisted cracking in aluminum[END_REF]. Therefore, it is important to define how the sample is stressed and what is the criterion to decide if a crack propagates. The force on the sample is raised in steps at the minimum jaws' velocity allowed by the machine (0.005 mm/min i.e. a strain rate of ǫ0 ≈ 2 × 10 -5 s -1 ). When the target force F is reached, the displacement of the jaws is fixed. A mechanical relaxation occurs which 2b) is measured on the side of the sample by DIC. The mechanical load is characterized by the stress intensity factor K defined from F and a as (see [START_REF] Tada | The stress analysis of cracks handbook[END_REF] for a "Singled Edge Notch Test" specimen):

K = f (a/W )σ ∞ √ πa (1)
f (a/W ) = 1.12 -0.23 a W + 10.56( a W ) 2 -21.74( a W ) 3 +30.42( a W ) 4
where W is the width of the sample and σ ∞ the remote stress. The applicability of the formula in our case was checked by finite element calculations. Finally, the parameters of the experiment are: those of the hydrogen charging (E c and t c ) and the thickness B of the sample. The load is controlled by the application of the force. The micro tensile stage is placed under an optical microscope and images are periodically taken enabling the determination of the crack length on the side of the sample by DIC. K is therefore known all along the test and is reported at initiation. The deformation of the notch ǫ notch = U U0 at initiation is also measured. U is the opening of the notch at the edge of the circular tip, as defined on figure 2b. Before the results are given, the DIC method is detailed.

Digital image correlation

The exposure of the polished surface to the H 2 SO 4 solution generates an attack of the surface which reveals the grains and the grain boundaries. When observed in optical microscopy, the rough surface state creates a speckle (Fig. 3) that is used to measure surface displacements by digital image correlation [START_REF] Tanguy | Cohesive stress heterogeneities and the transition from intrinsic ductility to brittleness[END_REF]. The method used is the "Lucas Kanade Feature Tracker" as implemented in the OpenCV library following the algorithm by Bouguet [START_REF] Tanguy | Plasticity localization ahead of a crack tip modeled by 2D discrete dislocation dynamics[END_REF] and called with the Python interface "pydic" [START_REF] Tehranchi | Atomistic study of hydrogen embrittlement of grain boundaries in nickel: I. fracture[END_REF]. The correlation between two images gives a displacement field that is called "incremental" to differentiate it from the displacement field which is the sum over all the incremental fields of the sequence of images, starting from an initial configuration with no plastic strain, crack and stress.

At the microscope's maximum magnification of ×2000, the pixel size is 110 nm. The grid size is 10 × 10 px 2 , i.e. the maximum spatial resolution is about one micron. The noise level is estimated by taking consecutive pictures without strain and correlating them. With a correlation window of 40 × 40 px 2 , the noise level is about 0.1 px, meaning a background noise of amplitude 1% in the measure of the deformation. Therefore, only large plastic deformations are measured which is already enough to meet our goals: determination of the crack tip position and measurement of the plastic deformation in the vicinity of the crack tip. An example is shown on figure 3. It represents the blunting of an intergranular crack followed by a fracture event. Two slip bands are present, in the grains above and below the intergranular crack plane, originating from two distinct sources slightly shifted along the crack propagation direction. They are originally both active during the blunting stage while, during fracture, only the upper one, associated to the crack tip, is active. At a later stage, the crack continues to propagate intergranularly in a semi-brittle manner but with a reduced plasticity production. The displacement field associated to the blunting in the near tip region could not be obtained because the speckle changes too much at this location. This two stages process is visible in the asymmetry of the "incremental" crack opening profile (the opening component of the displacement field u y , y being the orientation of the tensile axis) (Fig. 4a). Two plastic lobes are also visible in the "incremental" strain map of figure 4b. On the upper left corner, the focus is lost and this part of the image could not be used for DIC. The measured "strain" is the in plane equivalent strain ǫ eq :

ǫ eq = 2 3 (ǫ xx -ǫ yy ) 2 + 3ǫ 2 xy (2)
For comparison, an example of intergranular crack propagation is also given (Fig. 4c andd). Note the reduced plastic deformation (in intensity and spreading) and the difference in angle between the crack plane and the plastic lobes which is the signature of the crystallography. The opening angle is defined as α = 2 tan where ∆u yc is the opening of the crack measured on the "incremental" crack opening profile and ∆a is the increment of crack length. It was found that the distance in between the end of the flat part of the "incremental" crack opening profile and the crack tip position (f on Fig. 4c) is close to the value of ∆a measured by taking the difference in crack tip position between two images, within the limitation of the precision, about 1 µm. This means that the plasticity which emerges at the crack lips is essentially produced at or in the immediate vicinity of the moving crack tip and not from fixed remote sources. The position of the crack tip can now be defined. A particle sitting exactly on the crack path, ahead of the crack tip, would have zero displacement, by symmetry, while particles slightly off plane would be displaced further away either elastically or by the processes active ahead of the crack tip in the non linear "process zone". Behind the crack tip, such particles do not exist, they are all either above or below the crack plane. Therefore, the displacement field changes sign, with a discontinuity behind the crack tip while this occurs continuously, going through zero, ahead of the crack tip. Depending on the position of the grid with respect to the crack plane, this effect might be directly visible on the grid points. Otherwise, if the crack runs in between two rows of grid points, the displacements have to be extrapolated, and this introduces some subjectivity. In addition, the fact that intense slip bands intercept the crack plane at or behind the crack tip [START_REF] Tehranchi | A decohesion pathway for hydrogen embrittlement in nickel: Mechanism and quantitative prediction[END_REF] is also used as a guide to decide were the crack tip is. And finally, it is also considered that the opening deformation at the crack tip, estimated by the difference in the displacement of the grid point one row above and one row below the crack path, divided by one grid spacing, cannot be too large (smaller than 100%). In the two examples of figure 4 a and c, the crack tip positions are respectively 650 and 760 px. The incremental deformation at the tip ǫ tip yy is defined as:

ǫ tip yy = u(y + ) -u(y -) y + -y - (3)
where y + and y -are the positions above and below the crack plane, at the crack tip. ǫ tip yy is incremental, however, the consecutive images are seldom taken close enough in time to resolve the dynamics of the formation of the plastic zone in the first microns ahead of the crack tip (∆a is always larger than the plastic process zone, as will be shown below), except when the velocity is very low (4 µm/h and below), for example in figure 4 a (blunting). Therefore, the measurements represent the complete deformation of the tip. If the grid runs in between two rows of grid points, y + -y -= 10 px, otherwise, if a row of points have small displacements, it means that the crack runs close to this row and the deformation is estimated from the rows above and below this row (y + -y -= 20 px). When the correlation window is bigger that the grid step, the displacements are smoothed. This introduces a bias in the deformation. Window sizes of 20, 30 and 40 px were tested. Qualitatively, the displacement field remains the same, except in the vicinity of the tip. Indeed, ǫ tip yy takes the values, respectively, of 2.3%, 8.4% and 26%, while if the measure is made by taking the values one grid step further away from the crack plane, i.e. y + -y -= 30 px, ǫ tip yy is respectively: 14%, 15% and 18%. On the one hand, this window size effect comes from the proximity of the primary slip bands with the crack tip and the smoothing inherent to the DIC method which forces the continuity of the displacement field. On the other hand, if not broken the matter is continuous and the displacement should be a continuous interpolation between the displacement of the two primary slip bands. The method used is qualitatively correct, but is not accurate enough and spatially resolved enough to precisely measure ǫ tip yy . A more performant method, based on particle tracking, is discussed at the end of the paper. Together with ǫ tip yy , the length Λ of the deformed region ahead of the crack tip is also measured (Fig. 4c). It starts at the crack tip and ends when u(y + ) -u(y -) falls within the noise, i.e. below 0.1 px.

Internal hydrogen induced cracking results

A reference test, without hydrogen charging, was done in laboratory air. The relative humidity was of the order of 65%. The force was raised above the yield point (at 5% strain, σ max = 175 MPa, ǫ notch ≈ 65%, Tab. 2) after which the displacement of the heads of the sample was fixed. The force dropped at a constant rate. After 18 days, a 3 µm long crack appeared on the surface at the tip of the notch, together with a sudden increase in the force drop rate. This is considered as the initiation time. The corresponding stress intensity factor is

K=11 MPa √ m.
During the steady propagation stage, the crack was stopped and restarted by placing the experiment under dry nitrogen gas followed by a return in wet air after 2 hours (restarting the crack took 12 h) which confirmed the role of the wet environment in agreement with literature [START_REF] Lervik | Atomic structure of solute clusters in AlZnMg alloys[END_REF]. The crack velocity, estimated by the evolution of the trace on the surface, was of the order of 3 µm/h (10 -9 m/s also in good agreement with [START_REF] Lervik | Atomic structure of solute clusters in AlZnMg alloys[END_REF], figure 19.12). The fracture surface (Fig. 5a) is the result of a single planar transgranular crack which propagated through the entire section of the sample except for a reduced final shear lip.

Eleven other tests were also conducted in lab air but after hydrogen charging in different conditions presented in table 2. A macroscopic side view of the sample (sample G in table 2) after the complete propagation of a crack is shown on figure 6. During hydrogen charging, the electrolyte attacked the surface and revealed some of the grains. Multiple short cracks are visible from the notch tip. The main crack emerged from one of them and propagated as a single crack perpendicular to the tensile axis for about 300 µm. The data are always collected from this phase of the tests. In the final stage, the crack deviates, plastic deformation is intense and multiple cracks form.

The macroscopic results of table 2 show that the major parameter for embrittlement (low K at initiation and high average propagation velocity) is the applied cathodic potential E c at the expense of the charging time, within the range studied, i.e. all samples were exposed to the electrolyte for more than 24 h. The trend is that the deformation required for initiation decreases with applied potential. In the absence of hydrogen charging, it is as high as 65% and decreases down to an average of 20% at -1600 mV/ECS. The average crack velocity is also clearly dependent on E c with a slow velocity of 7 µm/h at -1300 mV/ECS, an average of 24 µm/h for the four samples at intermediate potentials -1400 and -1500 mV/ECS and a significantly higher velocity at -1600 mV/ECS of 200 µm/h. In the appendix, figure 10 presents a more detailed view of the deformation before the formation of the first cracks at the bottom of the notch. DIC shows the extent of the deformation and its repartition ahead of the notch. Initiation always occurs at the point of maximum deformation, which can be at one corner of the notch or at its center, depending on the geometry and the way the notch changes shape when the mechanical load is raised. It is interesting to see that the scale of the area over which the concentration of deformation, responsible for initiation, occurs is at least 20 µm (the area deformed beyond 10% is 2 × 2 px 2 in figure 10f), which is of the order of the grain size. A physical criterion for initiation should involve the local deformation within this zone and the stress concentrations that it creates [START_REF] Valle | Digital image correlation development for the study of materials including multiple crossing cracks[END_REF][START_REF] Van Der Ven | The thermodynamics of decohesion[END_REF][START_REF] Wan | Hydrogen embrittlement effect observed by in-situ hydrogen plasma charging on a ferritic alloy[END_REF][START_REF] Wang | Hydrogen influence on diffusion in nickel from first-principles calculations[END_REF] but this would require observations at a spatial resolution higher than the one reached in this study. The deformation maps (Fig. 10) also show that, in the most brittle cases, the prestraining due to the notch is very limited passed the first 100 µm ahead of the notch tip. For example, it is below 0.01 on figure 10f. Therefore, the single cracks propagate in samples essentially free from pre-deformation (even if this might enhance propagation).

Local crack velocities

Once the main crack has detached from the assembly of short cracks formed at the notch tip, local propagation velocities can be measured. Out of the 11 samples fractured (Tab. 2), only 90 crack tip positions could be determined. The limitations to such measurements are: the quality of the pattern which is not homogeneous among the samples, and the out of plane plastic deformation, which can be too large close to the crack tip to allow the measurement of the incremental crack opening profile. 80 local velocities were determined, 60 were from DIC, the others only from the gray level of the images. Among the 60, only 50 were considered representative of "propagation". For example, the very large velocities, some as high as 1800 µm/h, result from the sudden emergence of cracks running just below the surface and their apparent merging with an arrested surface crack. Such cases were discarded and the maximum local velocity v max is about 500 µm/h. As already mentioned, there is a marked difference in velocity depending on the electrochemical potential applied during the hydrogen charging. The propagation velocity is markedly enhanced when the potential is decreased to -1600 mV/ECS. Even if the hydrogen content was not estimated, it is reasonable to consider that a lower potential involves a higher hydrogen content. Therefore, the samples are split in two groups: B, D and E constitute the "low hydrogen content" group (LH group) (A, C and F did not allow local measurements) and G, H, I, J and K constitute the "high hydrogen content" group (HH group). The corresponding velocity distributions are shown on figure 7a for the LH group and figure 7b for the HH group with an average of 24 µm/h and a maximum of 111 µm/h, and an average of 200 µm/h and a maximum of 480 µm/h, for LH and HH respectively. In order to obtain the best possible images, the focus is adjusted by hand before every image which limits the time resolution to one minute maximum. Even then, the quality of the image does not alway allows to correlate properly two consecutive images and the time (and therefore spatial) resolution in crack velocity is usually larger. It is illustrated by the distribution of crack length increments ∆a in figure 7c that can be compared to the distribution of grain size along the crack propagation direction, the "transverse long direction" (figure 11 in the appendix). The distribution of crack increments goes from 1 to 36 µm, but is peaked around 6 µm (average 11 µm). The grain size distribution is bimodal (Fig. 11b) with small grains of size about 15 µm and larger grains of size about 45 µm (up to 100 µm). Apart from the small grains, the amplitude of the crack length increments ∆a is such that several measurements can be obtained along a GB facet, in case of intergranular cracking, or inside a grain in case of transgranular cracking. The local velocities were also systematically correlated with the fracture mechanism, determined by post-mortem observation of the fracture surface. No correlation appeared between sudden changes in v and changes in fracture mechanisms (for example a transition between intergranular and transgranular cracking). It should be specified that LH samples showed only intergranular crack propagation with different (not quantified) roughness, while HH experienced intergranular cracking with a low roughness (not quantified) and transgranular cracking (quasicleavage). This difference in roughness, between LH and HH samples, could be a sign of less embrittlement and more plasticity. The distribution of ∆a also gives an estimate of the spreading of the error made on the velocity. Knowing that the accuracy on the crack tip position is at best 1 DIC grid step (1 µm at maximum magnification), the relative error is ∆v/v ≈ 1/∆a which is in between 100% and 2.5% depending on the value of ∆a (16% for ∆a =6 µm). It can be mentioned that whatever the H content and the fracture mechanism, all cracks arrest or nearly arrest at obstacles (they are restarted by applying a small force increment). This is frequent when a triple point at grain boundaries is met. An intergranular crack was observed to blunt when approaching a triple point, arrest, and re-initiate slightly ahead on the next GB facet (off plane). Finally, the region in between the two fractured grain boundaries was fractured by plastic glide. This plastic process modifies the appearance of the triple points on the fracture surface (bottom of Fig. 5b andc). They appear blunted while they are usually sharp when the grains are equiaxed and the fracture can change GB facet by pure decohesion.

Plastic deformation

The displacement field obtained by DIC also enables to measure the plastic deformation tens of microns around the crack tip. The strain field, although noisy and limited to deformations higher than 1%, highlights the plastic lobes which accompany crack advance (Fig. 4b andd). This is called "primary plasticity". It originates from the shear stress concentration at the crack tip. In the two examples on figure 4b andd, the plastic lobes are not straight and spread over several grains, therefore they are not composed of a single glide system but more likely by several, furthermore in several grains. However, there is a clear difference in the angle between the two lobes between figure 4b andd, which might be of crystallographic origin. As a consequence, it is likely that several sources are activated and not only those at the crack tip or at its immediate vicinity. The incremental crack opening profiles often show that most of the plasticity that emerges at the crack surfaces originates from "moving sources" related to the crack tip. Therefore the crack opening angle α is a measure of only part of the plasticity created during crack ad- vance. Nevertheless, it is taken as a first measure of plasticity because it is a robust quantity, easy to measure on the profiles. The distributions of α for the LH and HH groups are given on figure 8a and d respectively. Plasticity is clearly reduced in the HH case in comparison to the LH case.

In addition to the crack opening angle, the distribution of ǫ tip yy and Λ, which characterize the deformation ahead of the crack tip are given on figure 8b, e andc, f respectively, for the HH and LH groups. When excluding the highest values of ǫ tip yy which correspond to crack arrests on triple points, the average values are respectively 5.8% and 17% for the HH and LH groups. The difference in Λ is much less pronounced with averages of 4.4 µm and 4.7 µm for the HH and LH groups respectively. There have been post mortem characterization of this deformation in TEM [START_REF] Bendo | Atomic scale haadf-stem study of η' and η 1 phases in peak-aged AlZnMg alloys[END_REF] and in high resolution EBSD [START_REF] Christodoulou | Hydrogen embritlement and trapping in Al-6%-Zn-3%-Mg[END_REF]. It was suggested that [START_REF] Bendo | Atomic scale haadf-stem study of η' and η 1 phases in peak-aged AlZnMg alloys[END_REF] hydrogen accelerates the evolution of the dislocation structure and stabilizes it in configurations where the dislocation density is higher than at comparable deformation without hydrogen, and that this microstructure favors the development of cracks. In our system the samples also contain hydrogen, initially distributed in the bulk, but they are not pre-strained prior to the expansion of the cracks. The "secondary plastic zone", which is situated just ahead of the crack tip and moves along with it, can reach high tensile deformations, in particular in the case were the H content is supposed to be low (LH), but where the intergranular fracture is conserved. This observation seems compatible with the idea promoted in [START_REF] Bendo | Atomic scale haadf-stem study of η' and η 1 phases in peak-aged AlZnMg alloys[END_REF] that some plastic deformation acts together with hydrogen to embrittle the grain boundaries. At the same time "primary plasticity" shields the external load and acts against crack propagation.

Discussion

The cracks observed in this work stem from several intricate phenomena and related parameters. The initial distribution of hydrogen after cathodic charging c 0 is supposed to be homogeneous. On the contrary, the strain field produced by the notch and the step by step rise of the force is highly inhomogeneous (Fig. 10). The time taken by initiation is long enough (typically 2 hours) to allow hydrogen redistribution, even at the scale of the whole sample by a combination of desorption, trapping in the strain field, with possibly additional very localized absorption from wet air in relation with localized strain at the notch tip. Later on, dur-ing the single crack propagation stage, the primary and secondary plasticity (Fig. 4 and8) might have several roles: to elastically shield the crack tip from the applied mechanical load K, to redistribute hydrogen in more elastically deformed regions or trap it in the plastic wake, or maybe even, in the case of secondary plasticity, to directly participate in the cracking process [START_REF] Bendo | Atomic scale haadf-stem study of η' and η 1 phases in peak-aged AlZnMg alloys[END_REF].

A first point to discuss is the delicate problem of effective diffusion in aluminum alloys. At the scale of the whole sample, the distribution of H after cathodic charging, and also during the fracture test, was evaluated by macroscopic diffusion calculations using effective diffusion coefficients and Fick's law (assuming ideal desorption under stress with zero concentration at the surfaces and ignoring the trapping related to plasticity [START_REF] Wang | Site stability and pipe diffusion of hydrogen under localised shear in aluminium[END_REF]). The hydrogen effective diffusion coefficient D ef f describes the diffusion of H as an interstitial atom in the lattice taking into account the trapping/detrapping from crystalline defects, precipitates... Therefore, it depends on the alloy and on its metallurgical state. Young and Scully [START_REF] Wunderlich | TEM studies of grain boundaries in cyclically deformed alznmg bicrystals[END_REF] have measured diffusion coefficients by isothermal desorption at moderate temperature and extrapolated the data to 25 • C. They obtained D ef f = 10 -11 m 2 s -1 for the alloy 7050 in the underaged state, a material which is similar to ours. However, it was recently demonstrated that the apparent diffusion coefficient is quite strongly influenced by the initial state of the system. For pure Al, D ef f = 10 -14 m 2 s -1 if the sample is free of hydrogen initially and increases up to 10 -12 m 2 s -1 when the sample already contains H [START_REF] Xie | Hydrogenated vacancies lock dislocations in aluminium[END_REF][START_REF] Young | Hydrogen production, absorption and transport during environment assisted cracking of an Al-Zn-Mg-(Cu) alloy in humid air[END_REF]. This is interpreted as a concentration dependence [START_REF] Zamora | Ab initio prediction of environmental embrittlement at a crack tip in aluminum[END_REF], with a faster apparent diffusion when the traps are filled (i.e. when the average concentration is high). Note that in a previous study [START_REF] Malitckii | Hydrogen uptake from plasma and its effect on EUROFER 97 and ODS-EUROFER steels at elevated temperatures[END_REF], in the same material as ours, H embrittlement was found to be coherent with D ef f = 2 × 10 -12 m 2 s -1 and not with the slower effective diffusion coefficients. We consider that the diffusion coefficient from Ai et al. [START_REF] Xie | Hydrogenated vacancies lock dislocations in aluminium[END_REF] is the most reliable because it was determined at room temperature and not extrapolated from high temperatures. Similar values have also been found independently by Hebert [START_REF] Xie | Hydrogenated vacancies lock dislocations in aluminium[END_REF]. Therefore, D ef f = 10 -12 m 2 s -1 is taken as the diffusion coefficient after the H uptake from the cathodic charging (assuming the traps are filled). This is a strong statement. If the effective diffusion coefficient is 10 -14 m 2 s -1 , then the uptake by diffusion would be limited to the first 100 µm below the surface. It would not be coherent with the hydrogen affected depth which extends in our experiments over the whole sample thickness, even when the thickness is 1 mm long. An argument, sometimes found in the literature, is that "dislocation transport" is the reason

v = v 0 (K, C H )(1- α α 0
) with v 0 = 500 µm/h and α 0 = 15 • . why H affects fracture deep in the material. In our case, the primary plasticity is oriented in such a way that it would transport hydrogen towards the mid thickness of the sample and therefore these arguments are not relevant to our case. Finite elements calculations, with this effective diffusion coefficient, show that a sample 500 µm thick retains 20% of its original hydrogen content after 10 hours, which is the maximum duration of the test. Crack propagation is also not diffusion rate limited at the maximum velocity of 500 µm/h [START_REF] Zhou | Is hydrogen diffusion along grain boundaries fast or slow? atomistic origin and mechanistic modeling[END_REF] with this diffusion coefficient. Embracing all the different aspects is not the goal of this discussion, but we already have a series of parameters: K, E c , and a number of quantities measured, v, α, Λ... from which we would like to attempt extracting meaningful correlations. Atomistic simulations of intergranular fracture by nano-ductility [START_REF] Girardin | Correlation between dislocation organization and slip bands: TEM and AFM investigations in hydrogen-containing nickel and nickelchromium[END_REF] have shown how a nano scale process zone develops in the case of a strong embrittlement by nano voids. It was shown that the nanoscale non linear process zone expands when the applied load increases. During this stage a dislocation can be emitted if the local cohesive stress is large enough so that the applied elastic stress creates a bending of the crack opening profile at the origin of an incipient dislocation core. Therefore, the ease of producing dislocations at the crack tip (or in the nanoscale fracture process zone) is intimately related to the competition between an increasing elastic stress applied to the process zone and the relaxation of this stress by the fracture mechanism itself.

The opening angle α is expected to be proportional to the ratio between the dislocation emission rate ν and the crack velocity, both result of the competition mentioned above. An extra complexity comes from the emission of the dislocations which creates the primary plastic zone which in turn shields the applied stress and therefore has a retroaction on the dynamics of the process zone formation, i.e. on ν and v. It is expected that, if the system is more embrittled, the dislocation emission is reduced. As a consequence, we take α as a parameter which expresses the degree of brittleness of the fracture process zone and plot v (the other local data representative of fracture that we could measure) as a function of α (Fig. 9). HH and LH are represented on different graphs, but the same curve is used as a guide to the eye to compare the two data sets. The trend is as expected: if α is large, v is small, eventually zero (complete blunting, like in the case of a mechanical overload). In the LH case, no crack propagation was observed if α is below 4 • , while at higher H content, large velocities are observed with a low opening angle. This suggests that either the shielding by primary plastic lobes is not completely efficient so that raising K gives a higher local stress, necessary for GB fracture at low hydrogen concentration, or, the secondary plasticity is part of the embrittlement process, or both. It can be noticed that transgranular fracture (in full symbols on figure 9a) is slower that intergranular fracture, if the maximum velocity at constant α is considered. No transgranular cracks ran faster than 250 µm/h.

One puzzling aspect of the data is that whatever the α value, there are cracks with almost zero velocity. This is conceavable if crack advance requires a threshold local stress and if the stress felt at the crack tip falls below this threshold because of the turtuousity of the crack (many grains exist in the thickness, even in the thinnest cases, and secondary cracks, parallel to the main crack, are observed on the fracture surface).

Conclusion

Embrittlement of an AlZnMg alloy by internal hydrogen was studied experimentally by means of micro tensile tests. Single cracking was favoured by notching the samples, and the mechanical load was imposed by fixing the displacement of the grips. Crack propagation was followed on the surface by optical microscopy. The displacement field was measured by digital image correlation with a resolution of 1 µm. The crack positions, and therefore the local propagation velocities, were determined, but also the plasticity produced by the crack advance. A correlation was attempted between the velocity and the opening angle, measured from the crack opening profile, representing the rate of dislocation emission from the fracture process zone. The expected tendency where the crack velocity decreases when more dislocations is injected in the system is retrieved but, whatever the crack opening angle, low propagation velocities are always observed. Indeed, the velocities are scattered between 4 and 500 µm/h in the most brittle cases. Furthermore, fracture occurs by a rich variety of mechanisms mixing transgranular and intergranular cracking. No strong correlation was found between the local crack velocity and the nature of the crack, even if the largest velocities (beyond 250 µm/h) always correspond to smooth intergranular cracks.

The experiment suffers from a number of limitations, in particular: the number of exploitable images per tensile test is low, because the quality of the surface pattern and the focus of the image are not always ideal for DIC; the resolution (1 µm) is sufficient to image primary plasticity but is too limited to have a clear picture of the secondary plasticity. Future work will be dedicated to the development of cracking experiments within an SEM to better characterize secondary plasticity, in interaction with simulations.

Fig. 1 Fig. 2

 12 Fig. 1 Microstructure viewed by high resolution scanning electron microscopy (a) as received, (b) after heat treatment and one week natural aging, the precipiates are dissolved.

Fig. 3 A

 3 Fig. 3 A zoom on the crack tip of sample D viewed by optical microscopy (10 px = 1.1 µm). (a) The DIC grid superimposed to image A (step 10 px = 1.1 µm). (b) Image B, after a crack advance, with the displacement field (arrows) obtained by correlating image A with image B (correlation window 40×40 px 2 ).

Fig. 4

 4 Fig. 4 Incremental crack opening profiles (a and c) and equivalent strain maps (b and d) obtained by DIC of two pairs of images. (a) and (b) illustrate crack blunting followed by a transgranular crack advance (the two consecutive crack configurations are the same as those used for figure 3). (c and d) are representative of brittle intergranular cracking. The symbols in black on the profiles represent the displacement in y, ahead of the crack tip. For one position in x, the different symbols represent different planes (y=-20, -10, 0, +10 and +20 px from the crack plane). The open symbols represent the opening at the back of the crack tip. The crack tip is at x=650 px, y=585 px, in between two rows (resp. x=760 px, y=565 px) in (a) and (b) (resp. in (c) and (d)). The insert in (c) defines the crack opening angle α, the length Λ of the "plastically deformed zone" along the crack path, ahead of the tip and the increment of crack opening ∆u yc . f is the distance between the end of the flat region of the profile and the crack tip (see text).

Fig. 5

 5 Fig. 5 (a) Fracture surface of the reference test in humid lab air. (b) Typical fracture surface in air after electrochemical hydrogen charging. It is composed of a mixture of transgranular and intergranular fracture (sample J, Tab. 2). (c) Blunted/sharp broken tripple points (sample G, Tab. 2).

Fig. 6

 6 Fig. 6 Macroscopic optical view of the side surface of sample G (E c = -1600 mV/ECS, t c = 24 h, K = 5 MPa √ m, Tab. 2) close to final fracture.

Fig. 7

 7 Fig. 7 Local crack propagation velocity distribution v for samples B, D and E (a) which are charged with hydrogen at moderate potential, and H,I,J,K which were charged at -1600 mV/ECS (b). (c) The distribution of the crack increments ∆a obtained from the crack tip positions measured on the incremental crack opening profiles from DIC.

Fig. 8

 8 Fig. 8 Graphs a, b and c correspond to the group with a low hydrogen content (LH) while d, e, f correspond to the one with a high hydrogen content (HH). (a, d) Crack opening angle α, (b, e) deformation in the first micron ahead of the tip ǫ tip yy and (c, f) length of the process zone (Λ) ahead of the tip.

Fig. 9

 9 Fig. 9 Local velocity as a function of the crack opening angle: (a) high hydrogen content, (b) low hydrogen content. The lines are a guide to the eye corresponding to v = v 0 (K, C H )(1-

Fig. 10

 10 Fig. 10 Illustration of the local strain necessary for crack initiation from a notch. Images (a,b), (c,d) and (e,f) correspond respectively to samples D, E and J (Tab. 2), i.e. to H charging conditions (E c =-1500 mv/SCE, t c =72 h), (E c =-1500 mv/SCE, t c =24 h), (E c =-1600 mv/SCE, t c =24 h), respectively. Pictures on the left are optical images of the surface used for DIC. The grid is superimposed. On the right are the corresponding Von Mises deformation maps calculated from the displacement field. The axis units are in pixels (px) with 20 px = 12 µm, which is also the grid step length. For example, (b) is 700 px x 1000 px = 414 µm x 593 µm. The inserts show the cracks emerging from the notch tip.

Fig. 11 (

 11 Fig.11(a) Crack length increment distribution and (b) grain size distribution in the "long transverse" direction in the long transverse/short plane (Fig.2).

Table 1

 1 Measured composition of the AA7018 alloy (weight%). Beyond this time, the onset of the force drop due to crack propagation requires a certain initiation time. It was observed that if no drop is detected after 20 min, propagation never starts within the next 2 hours (a few tests were made with a waiting period of 12 hours, which compares with the total duration of the experiment) and therefore, a waiting time of about 20 min in between two force raises is appropriate to detect crack initiation and propagation. The crack length a (Fig.

	Si	Fe	Cu	Mn	Mg	Cr	Ni	Zn	Pb	Ti	Zr	Al
	0.0308 0.0804 0.0116 0.0099 1.23 0.0016 0.0104 4.995 0.0004 0.0352 0.1417 base
	lasts about 5 min.

Table 2

 2 Characteristics of the tests done: sample name, cathodic electrochemical potential E c , hydrogen charging time t c , maximum stress σ max (force divided by the initial section, not taking into account the notch), notch deformation ǫ notch , applied stress intensity factor K (at initiation and at the end of the single crack propagation), average crack velocity v, maximum local velocity v max , sample thickness B. The velocities marked with a star were not obtained by DIC.

	Sample	E c mV/SCE	t c h	σ max ǫ notch MPa %	K MPa √	v m µm/h µm/h mm v max B
	air	-	-	175	65	11		3	-	1.0
	A	-1300	48	190	64	22		7*	7*	0.4
	B	-1400	72	158	39	10 -11	15	96	0.5
	C	-1400	24	164	53	13 -15	25*	25*	0.4
	D	-1500	72	126	34	23 -26	33	101	0.6
	E	-1500	24	121	39	18 -22	22	79	0.7
	F	-1600	48	145	11	5 -9		120*	130*	0.2
	G	-1600	24	156	24	5 -13	262	435	0.2
	H	-1600	48	145	19	7 -12	176	1800	0.6
	I	-1600	28	162	25	8 -14	254	528	0.6
	J	-1600	24	155	18	8 -12	65	144	0.4
	K	-1600	65	155	22	7 -12	338	480	1.0

nm 200 nm

Acknowledgements This work used the equipment from the Optolyse platform (ILMTech). We gratefully acknowledge V. Motto-Ros from Institut Lumière Matière, P. Pittet from Institut des Nanotechnologies de Lyon and O. Dezellus and B. Gardiola from Laboratoire des Multimatériaux et Interfaces for their help in preparing the experiments.


S Scharnholz
email: sigo.scharnholz@isl.eu

R Hassdorf

D Bauersfeld

B Vergne

L V Phung

D Planson

Investigation of SiC thyristors with varying amplifying gate design

Due to their superior performance in terms of current density and blocking voltage, silicon carbide (SiC) thyristors are of particular interest for pulsed power applications [1,[START_REF] Agamy | Proceedings of the IEEE Energy Conversion Congress and Exposition[END_REF]. Furthermore, in view of the tremendous progress made in SiC crystal quality and wafer size, large-area SiC thyristors appear to be feasible today. In this context, both electrically and optically triggered SiC thyristors require adequate amplifying gate structures. With the aim to investigate this aspect, a mask set featuring numerous SiC thyristor devices with varying amplifying gate structures has been conceived [START_REF] Scharnholz | Proceedings of the Symposium de Génie Electrique[END_REF]. And based thereon, a 100 mm epitaxial SiC wafer has been processed recently. Resulting from this process run, the present paper reports on the electrical characteristics measured on-wafer as well as on singularized and encapsulated SiC thyristor devices with and without amplifying gate.

The fabricated devices are asymmetrically blocking vertical SiC thyristors, of which the 10 µm thick p-base (drift) layer (5×10 15 cm -3 ) is separated by a 4 µm p-type buffer (5×10 17 cm -3 ) from the wafer substrate (> 1×10 19 cm -3 ). The n-base (gate) and the anode layers have a nominal thickness (and doping concentration) of 2 µm (1×10 17 cm -3 ) and 1 µm (1×10 19 cm -3 ), respectively. The gate and the single-step junction termination extension (JTE) are formed simultaneously by reactive ion etching (RIE). A second RIE step was used to define the JTE length and thus the device mesa. Following the etching, the primary metallization of gate/cathode (Ti/Ni) and anode (Ni/Ti/Al) was created separately by e-beam evaporation, lift-off and subsequent annealing. The surface was passivated with 1 µm thick, deposited SiO2. The over-metallization consists of 1 µm Al and 0.7 µm Ti/Ni/Au on the top-side (anode/gate) and the backside (cathode) contacts, respectively Fig. 1 shows a scheme of circular shaped SiC amplifying gate thyristors with the characteristic dimensions (radii). All fabricated amplifying gate thyristors have a central, circular gate (G) with a radius of rEA = 160 µm and a main anode (A), expanding from an inner radius of rEM = 500 µm to a distance of rSM = 1200 µm, where an optional anode short and the beginning of the JTE is situated. With a JTE length of 150 µm and a security margin to the cutting edge, the die size measures 3.2 mm × 3.2 mm (0.1 cm 2 ). In-between gate and main anode lies the amplifying gate, consisting of the anode A' of the pilot thyristor and the auxiliary gate G'. The etch-step separating A' and G' forms a circle with varying radius of rSA. This way, the resulting thyristors meet the amplifying gate design criterion (RP > RM Þ rSA/rEA > rSM/rEM) [START_REF] Scharnholz | Proceedings of the Symposium de Génie Electrique[END_REF][START_REF] Baliga | Fundamentals of Power Semiconductor Devices[END_REF] to different degrees: non-adequate (rSA = 300 µm), adequate (rSA = 400 µm), and safe (rSA = 460 µm).

On-wafer (in a vacuum probing chamber), the best devices demonstrated stable blocking up to 2 kV. After dicing, bonding and encapsulation they were successfully tested up to 1.2 kV. Fig. 2 shows the micrograph of one of the fabricated and bonded devices, a square shaped amplifying gate thyristor with dotted anode shorts. Fig. 3 shows on-state characteristics (at varying gate currents) of one of the amplifying gate thyristors (lower graph) compared to a device without amplifying gate. As expected, the amplifying gate thyristor shows a characteristic snap-back and requires a higher gate current to trigger. Fig. 4 shows the gate-anode characteristics of an amplifying gate thyristor, allowing to determine the characteristic resistances RP and RM of pilot and main thyristor, which confirm the application of the above mentioned design rule. Fig. 5 shows turn-on characteristics of SiC thyristors with circular and square shaped anode (cf. Fig. 2), each with varying amplifying gate design. The missing peak-shaped inversion in the gate current transient confirms that the pilot thyristor of non-adequately designed devices is not triggered.

Fig. 1 .Fig. 2 .

 12 Fig. 1. Scheme of an SiC amplifying gate thyristor.

Fig. 3 .

 3 Fig. 3. On-state characteristics of SiC thyristors with and without amplifying gate (AG).

Fig. 4 .

 4 Fig. 4. Gate-anode characteristics of SiC thyristors with differently dimensioned amplifying gate.

Fig. 5 .

 5 Fig. 5. Turn-on characteristics of SiC thyristors with circular (solid lines) and square shaped (dotted lines) anode, each with differently dimensioned amplifying gate (cf. legend).


H Krawiec
email: krawiec@agh.edu.pl

V Vignal

J Lelito

A Krystianiak

P Ozga

In-situ monitoring of the corrosion behaviour of austempered ductile iron (ADI) under cyclic salt spray exposure

Keywords:

The corrosion behavior of austempered ductile iron (ADI) under cyclic salt spray exposure was studied using a specific three-electrode cell located in the salt fog chamber. Three steps were identified, namely pitting corrosion of ferrite (very first cycles), followed by inward oxidation of the ausferrite matrix and then (iii) mixed oxidation of the matrix associated with the formation of a rust layer at the specimen surface. Evolution of the corrosion rate calculated from in-situ electrochemical measurements was correlated with surface changes quantified from optical and field emission scanning electron (FE-SEM/EDS) microscopies. Corrosion products were identified by means of micro-Raman spectroscopy.

Introduction

In austempered ductile iron (ADI), graphite spheres are randomly distributed in a dual-phase matrix (ferrite (α-phase) and high carbon austenite (noted γ HC -phase)). This particular microstructure leads to good mechanical properties compared to ductile iron (high strength and good wear resistance, for example) [START_REF] Kovacs | On the terminology and structure of ADI[END_REF][START_REF] Shih | Mechanical properties and microstructures of copper alloyed ADI[END_REF][START_REF] Rao | Influence of microstructure on fracture toughness of austempered ductile iron[END_REF][START_REF] Dorazil | Mechanical properties of austempered ductile iron[END_REF]. ADI is then used in many industrial sectors (automotive and railroad industries, truck and bus components, construction and mining components, agricultural applications, defense sector…).

Several papers [START_REF] Krawiec | Relationships between microstructure and pitting corrosion of ADI in sodium chloride solution[END_REF][START_REF] Krawiec | Corrosion behaviour and structure of the surface layer formed on austempered ductile iron in concentrated sulphuric acid[END_REF][START_REF] Krawiec | Influence of the heat treatment on the microstructure and corrosion resistance of austempered ductile iron (ADI)[END_REF][START_REF] Krawiec | Macroscopic and local electrochemical studies of austempered ductile iron in perchlorate solutions[END_REF][START_REF] Krawiec | Influence of corrosion inhibitors on the electrochemical behaviour of ADI in aqueous environments: local and macroscopic approaches[END_REF] have described the corrosion behaviour of ADI in bulk electrolyte solutions. In these papers, the influence of the elaboration parameters (tempering temperature, time of isothermal treatment,…), microstructure and physical-chemical parameters (nature of ions, temperature, pH…) has been studied and discussed. By contrast, the corrosion behavior of ADI under salt spray conditions has only been investigated according to ASTM B117 standard (continuous salt spray exposure) by means of weight loss measurements [START_REF] Hsu | Effects of copper and austempering on corrosion behavior of ductile Iron in 3.5 pct sodium chloride[END_REF]. No electrochemical investigations have been performed after (continuous / cyclic) salt spray test exposure.

Specific cells have been designed to perform in-situ electrochemical measurements under salt spray conditions, in droplets and in thin electrolyte films. They have been applied to some metals (Fe [START_REF] Tang | Atmospheric corrosion local electrochemical response to a dynamic saline droplet on pure Iron[END_REF], Zn [START_REF] Simillion | Atmospheric corrosion modeling[END_REF][START_REF] Song | Initial corrosion of pure zinc under NaCl electrolyte droplet using a Zn-Pt-Pt three-electrode system[END_REF], Cu [START_REF] Zhang | Electrochemical behavior of Cu underneath thin electrolyte film in atmosphere[END_REF], Mg [START_REF] Chen | Corrosion of pure magnesium under thin electrolyte films II -influence of thin electrolyte films on anodic process of pure magnesium corrosion[END_REF]) and alloys (Fe-Ni [START_REF] Boughrara | Thin electrolyte layer thickness effect on corrosion behaviour of invar in sulphate solutions[END_REF], 316 L stainless steel [START_REF] Ledvina | The thin electrolyte layer approach to corrosion testing of dental materials-characterization of the technique[END_REF][START_REF] Remita | A thin layer cell adapted for corrosion studies in confined aqueous environments[END_REF], 907 steel [START_REF] Wang | Corrosion behavior of 907 steel under thin electrolyte layers of artificial seawater[END_REF] and aluminum alloy 2024-T3 [START_REF] Cheng | A study of the corrosion of aluminum alloy 2024-T3 under thin electrolyte layers[END_REF]), but not to ADI. In-situ electrochemical measurements offer several advantages over ex situ approaches (salt spray exposure followed by electrochemical measurements in the bulk electrolyte). The influence of the development / evaporation of the electrolyte film or droplets, the formation of corrosion products (dense or compact layers), the electrolyte flow… on corrosion mechanisms and kinetics can be investigated. In ex-situ approaches, corrosion products are modified by contact with air, there is no electrolyte flow…

The corrosion behavior of pure iron and various steels under atmospheric conditions has been addressed in numerous papers. These data may be useful to understand the behaviour of ADI. The reaction mechanisms of formation of corrosion products have been studied and discussed in review articles [START_REF] Melchers | Long-term corrosion of cast irons and steel in marine and atmospheric environments[END_REF][START_REF] Tewary | Microstructural characterisation and corrosion performance of old rail way girder bridge steel and modern weathering structural steel[END_REF][START_REF] Guo | A critical review of corrosion development and rust removal techniques on the structural/environmental performance of corroded steel bridges[END_REF][START_REF] Oh | Characterization of iron oxides commonly formed as corrosion products on steel[END_REF][START_REF] Cornell | The Iron Oxides[END_REF] and references therein]. Fe first dissolves to Fe 2+ ions. These ions undergo hydrolysis (formation of Fe (OH) + ) and oxidation (formation of Fe 3+ ions). Fe(OH) + can be quickly oxidized to Fe(OH) 2

+

, which then transforms into γ-FeO(OH) (lepidocrocite). Fe(OH) + can also be slowly oxidized to Fe a

2+ Fe b 3+ O x (OH) y .
This compound then transforms into green rust (slow oxidation) or into magnetite (rapid oxidation) [START_REF] Oh | Characterization of iron oxides commonly formed as corrosion products on steel[END_REF]. The formation of green rust stage is followed by the appearance of lepidocrocite, which then transforms into end-products such as goethite (α-FeOOH), magnetite (Fe 3 O 4 ), maghemite (γ-Fe 2 O 3 ) and hematite (α-Fe 2 O 3 ). The nature of end-products depends upon the environmental conditions (pH, temperature, solution composition and oxidation rate). In the presence of halogen ions (Cl -or F -), akaganeite (β-FeOOH) can be formed from green rust. Corrosion products play an important role in atmospheric corrosion processes. Fe 2+ based compound may undergo oxidation whereas Fe 3+ based compounds may undergo reduction (making the oxygen reduction reaction unnecessary and promotes oxidation processes).

In the present paper, the corrosion behaviour of ADI is studied under salt spray conditions for 42 wet / dry cycles (0.05 M NaCl at 30 • C, 0.4 mL/h). Electrochemical measurements (OCP vs time measurements and polarisation curves) were performed in the salt spray chamber, at the beginning of each wet cycle. The morphology and composition of corrosion products was determined by means of field-emission scanning electron microscopy (FE-SEM/EDS) observations and micro-Raman experiments. The role of the ADI microstructure (studied by FE-SEM/EDS) and corrosion products in the corrosion processes is discussed.

Experimental

Material, specimen and austempering heat treatment

The chemical composition of the ductile iron (DI, pearlitic structure) which was used as initial alloy is C: 3.48 wt.%, Mn: 0.63 wt.%, Si: 2.58 wt.%, P: 0.03 wt.%, S: 0.017 wt.%, Cr: 0.06 wt.%, Ni: 0.04 wt.%, Cu: 0.32 wt.%, Mg: 0.054 wt.%, Mo: 0.07 wt.%, Ti: 0.015 wt.%, Sn: 0.006 wt.%, Pb: 0.001 wt.%, V: 0.009 wt.%, W: 0.038 wt.%, Zn: 0.003 wt.%. Liquid DI was poured into Y block molds. After solidification, the blocks were cut into cylinders (diameter =10 mm, length =70 mm). To obtain ADI samples, cylinders were placed in a salt bath furnace. A temperature of 900 • C was first applied for 2 h (austenitizing), followed by 280 • C for 10 min (austempering). Water cooling to room temperature was then carried out.

Prior to salt spray tests, ADI specimens were mechanically ground with silicon carbide (SiC) emery papers down to 4000 grit and polished with diamond pastes down to 1 μm. They were ultrasonically rinsed in ethanol between each step.

Salt spray test

Fig. 1 shows the three-electrode arrangement used in the electrochemical experiments. The counter electrode is a graphite carbon disk of very high purity (chemical composition: 99.997 wt.% C, 0.2 wt.%ppm Al, 1 wt.%ppm B, 0.05 wt.%ppm Ca, 0.01 wt.%ppm Cu, 0.2 wt.%ppm Fe, 0.05 wt.%ppm Mg, 0.02 wt.%ppm Na and 1 wt.%ppm Si). Its surface area is 0.8 cm 2 . The pseudo reference electrode is a platinum wire (length =2 mm, diameter =1 mm) and the working electrode is ADI (surface area of 0.8 cm 2 ). This three-electrode arrangement was embedded in an epoxy resin and placed in the salt spray chamber (Ascott S120iP) at a tilt angle of 5 • .

The corrosion test consists of successive wet (0.05 M NaCl, 0.4 mL/h) and dry periods. A cycle corresponds to a wet period (12 h at 30 • C) followed by a dry one (12 h at 30 • C). The total number of cycles is 42. Electrochemical measurements were performed for each cycle, 4 h after the onset of the wet period. The OCP measurement was first carried out for 900 s (acquisition frequency of 1 Hz). Linear sweep voltammetry (LSV) measurements were then performed from -1400 mV up to +1000 mV. at a potential scan rate of 3 mV/s. The potentiostat/Galvanostat/EIS Analyzer Parstat 4000 (AMTEK Princeton Applied Research, USA) was used. To validate the experimental procedure used, a preliminary test (not shown in the paper) was performed. Two identical samples were introduced into the salt spray chamber. During cycles 4 and 30, a polarization curve was plotted on only one of the two samples, 4 h after the start of the wet phase. As soon as this curve was stopped, two other polarization curves were plotted simultaneously on the two samples (the sample on which a curve has already been plotted and the so-called 'fresh' sample). These three curves were quantitatively compared (shape, values of current density for given applied potential and corrosion rates). It was found that for the experimental conditions selected in this paper, all the polarization curves are very close, indicating that repeatability is very good and changes induced by plotting these curves are very small. The same experiments were performed for the OCP measurements and the same conclusions were drawn. In addition, XRD experiments on the two samples did not revealed any changes in the corrosion products.

FE-SEM/EDS observations

The morphology and the chemical composition of the rust layer were studied at the microscale by using a field-emission type scanning electron microscope (JEOL 7600 F) with an integrated electron dispersion spectrometer (FE-SEM/EDS).

Micro-Raman spectroscopy measurements

Micro-Raman spectra were recorded using a Renishaw inVia Raman microscope (backscattering geometry, 100×objective lens with numerical aperture 0.85 (Leica objective), 633 nm excitation laser (HeNe, spatial resolution 1 μm), 1 s exposure time, 1000 accumulations and laser power reduced to 700 μW by filters to avoid sample damage in the effect of local heating). When the laser power is too high ferrihydrite, lepidocrocite, magnetite, maghemite, siderite and wustite can form hematite by oxidation [START_REF] Hanesch | Raman spectroscopy of iron oxides and (oxy)hydroxides at low laser power and possible applications in environmental magnetic studies[END_REF]. A Peltier-cooled CCD matrix detector was used for detection (diffraction gratings of 1800 grooves mm -1 , resolution below 1.34 cm -1 ). The spectra were recorded in fast mode (fixed position of grating) according to ranges: i) the range of wavenumber 50-1300 cm -1 (main oxide bands); ii) the range of wavenumber 900-2000 cm -1 (the remaining oxide bands, graphite bands (D,G) as also band of amorphous carbon); iii) the range of wavenumber 2500-3250 cm -1 (graphite 2D band and the remaining second-order graphite bands). The spectrum model profile was determined by fitting each fragment of the spectrum after background correction with a pseudo-Voigt profile (linear combination of the Gaussian curve and Lorentzian curve) using Renishaw Wire ver. 3.4 software.

Results and discussion

Fig. 2(a-d) shows FE-SEM/EDS and optical images of ADI after etching. This microstructure was already analysed at the microscale in a previous work [START_REF] Krawiec | Influence of the heat treatment on the microstructure and corrosion resistance of austempered ductile iron (ADI)[END_REF]. This was done by means of FE-SEM/EDS, X-ray diffraction (XRD) and microhardness tests. Graphite spheres are embedded in an 'ausferrite matrix' (Fig. 2(a,d)). This matrix consists of a mixture of ferrite plates (generally oriented perpendicularly to the surface) and high carbon austenite (γ HC -phase) between them, Fig. 2(a-d). FE-SEM/EDS analyses were performed in the ferrite plates oriented parallel to the surface (Fig. 2(c)): 6.7 ± 0.6 at.% C, 6.9 ± 0.8 at.% Si, 0.2 ± 0.1 at.% Mn, 1.9 ± 0.6 at.% O, 84.3 ± 0.8 at.% Fe. During the ausferritic transformation, nucleation of ferrite plates occurs at the interface between graphite and austenite, Fig. 2(d). This site is privileged to nucleate ferrite probably due to the highest concentration of elements favoring ausferritic transformation, i.e. silicon and nickel. These alloying elements (Ni, Si) reduce the solubility of carbon in austenite, and silicon additionally accelerates the diffusion of carbon in austenite [START_REF] Chang | Study on isothermal transformation of austempered ductile Iron[END_REF]. These two effects promote nucleation and growth of ferrite plates. In addition, it should be added that the size of the precipitates is largely a function of the isothermal transformation temperature.

Martensite was also detected by means of optical observations and microhardness measurements [START_REF] Krawiec | Influence of the heat treatment on the microstructure and corrosion resistance of austempered ductile iron (ADI)[END_REF], Fig. 2(a-b). Martensite results from the decomposition of a part of the γ-phase (that was not completely dissolved during tempering) during water cooling. EDS analyses in the martensite yields 11.2 ± 0.9at.% C, 3.4 ± 0.1 % Si, 1 ± 0.1 % Mn, 2.2 ± 0.1 % O, 81.6 ± 0.9 % Fe. XRD revealed that the alloy is composed of 20 % of austenite and 80 % of ferrite + martensite.

In-situ monitoring of the corrosion behaviour of ADI and surface observations

Fig. 3(a-b) shows the evolution of the OCP vs time and polarisation curves for different cycles. During the measurement period, the OCP was found to be stable and no transients were observed. Polarisation curves reveal the active behaviour of ADI under salt spray conditions (no passive range). Several electrochemical parameters were derived from these curves, namely the current density at -1250 mV (cathodic domain) and +500 mV (anodic domain), Fig. 4(a), the corrosion current density (Fig. 4(b)) and the average value of the OCP (Fig. 4(c)). The corrosion current density was determined from the polarisation resistance method [START_REF] Scully | Polarization resistance method for determination of instantaneous corrosion rates[END_REF]. The corrosion rate, v corr , was then determined using Eq. ( 1).

v corr = j corr . M nFρ ( 1
)
where M is the molar mass (M = 55.8 g/mol), j corr the corrosion current density (A/m 2 ), n the number of electrons (n = 2), F the Faraday's constant (F = 96485 C/mol) and ρ the density (ρ = 7.8 × 10 6 g/m 3 ).

Fig. 4(a-c) shows that the electrochemical parameters vary significantly vs the number of cycles, indicating that the corrosion behaviour of ADI changes with time. To interpret these changes, optical and FE-SEM/EDS observations at different spatial resolution were carried out after 3, 9, 20, 30 and 42 cycles (Figs. 5 and6). This tendency suggests that the surface is active. Surface observations after cycle #3 confirm this assumption. The microstructure of ADI is still visible in numerous sites, Fig. 5(a1-a2) and 6(a-b). EDS analyses indicate that these sites are slightly oxidized (low oxygen content): 14.9 ± 0.2at.% C, 18.1 ± 4.3 % O, 2.6 ± 1% Na, 2.9 ± 0.5 % Si, 0.05 ± 0.05 % Cl and 60.9 ± 6.1 % Fe. There is no difference level between the graphite spheres (not oxidized) and the "ausferrite" matrix, due to the inward diffusion of oxygen in the matrix. A high density of pits was observed in these sites, Fig. 6(a-b). As it was already observed in bulk electrolyte solutions [START_REF] Krawiec | Influence of the heat treatment on the microstructure and corrosion resistance of austempered ductile iron (ADI)[END_REF], pits initiate in the ferrite. Therefore, pitting corrosion and inward oxidation of ADI mainly occur during stage I. However, in some sites, the surface is covered with a thick rust layer (Fig. 5(a2)), because of the outward diffusion of metallic ions. Numerous cracks are visible in this layer.

Stage II (cycles #3 -#9). By contrast to stage I, the "ausferrite" matrix of ADI is not visible on the whole surface. Only graphite spheres can be observed (Figs. 5(b2) and 6 (c)). They are not dissolved and can still act as cathodic sites. At the end of stage II, the entire surface (except graphite spheres) is strongly oxidized and a dense rust layer is present (Fig. 5(b2)). EDS analyses gives (performed on 4 different sites): 22.4 ± 14.7at.% C, 48.6 ± 10.1 % O, 0.3 ± 0.1 % Na, 0.9 ± 0.5 % Si, 0.3 ± 0.2 % Cl and 27.5 ± 4.9 % Fe. There is again no significant difference level between the rust layer and graphite spheres (Fig. 6(c)), because of inward diffusion of oxygen in the matrix. Small cracks are found around some graphite spheres, Fig. 6(c). The dense rust layer can help to prohibit further corrosion of ADI. This was confirmed from the evolution of the electrochemical parameters. Indeed, a slight decrease of the anodic current density (from 200 down to 150 μA/cm 2 , Fig. 4(a)), the corrosion rate (from 70 down to 40 μm/year, Fig. 4(b)) and the corrosion potential (from -800 down to -700 mV, Fig. 4(c)) is observed when compared to the end of stage I.

Stage III (cycles #10 -#42). The OCP reaches a plateau and its value fluctuates around -680 mV (Fig. 4(c)). By contrast, the corrosion rate increases continuously and sharply (up to 140 μm/year at the end of the test), Fig. 4(b). This increase of v corr is associated with the growth of Pores and cracks are often visible in this layer (Fig. 6(f)). It was already demonstrated on steels that the porous rust can act as a water retention phase [START_REF] Zhang | Fatigue life estimation of existing bridges under vehicle and non-stationary hurricane wind[END_REF][START_REF] Zhang | Influence of outer rust layers on corrosion of carbon steel and weathering steel during wet-dry cycles[END_REF] and absorb the Cl -ions [START_REF] Burger | Use of the gold markers method to predict the mechanisms of iron atmospheric corrosion[END_REF], which can actually accelerate the corrosion rate [START_REF] Hoerlé | Advances in understanding atmospheric corrosion of iron. II. Mechanistic modelling of wet-dry cycles[END_REF].

All graphite spheres are progressively covered with the porous rust layer (outward oxidation). between cycles #20 and #30, nearly all graphite spheres are covered (Fig. 5(c2) and 6(d-e)). This may explain why the current density in the cathodic domain starts to decrease slightly after roughly 20 cycles (Fig. 4(a)). Spheres are preferential sites for the cathodic reactions.

Cross-section observations of the rust layer

After corrosion tests, samples were cut using the Presi Mecatome T210 machine and embedded in an epoxy resin. Cross-section surfaces were ground and then smoothed (down to 1 μm). During this preparation step, most of the outer porous layer was removed. This layer is very brittle and has very low adhesion to the substrate. Fig. 7(a) shows an optical cross-section image of the rust layer. It is composed of an inner dense layer (light grey in Fig. 7(a-c)) and an outer porous layer (brown color in Fig. 7(c)). A large crack is observed in the inner dense layer. Observations at high spatial resolution of this crack indicate that it is filled with corrosion products. The presence of cracks promotes the propagation of the electrolyte (with chloride ions and dissolved oxygen) inside the rust layer, promoting corrosion processes.

Numerous graphite spheres are found in the inner dense layer (sites S4 in Fig. 7(b-c)) confirming that this layer is associated to an inward growth mechanism (diffusion of oxygen to the metallic substrate). Micro-Raman measurements in spheres indicate that are not affected by corrosion processes. Indeed, excatly the same micro-Raman spectra were obtained on spheres located both in the metallic matrix and in the inner dense layer. On the other hand, the outer porous layer (most of which was removed during the preparation stage) is associated to an outward growth mechanism, Fig. 7(c).

The substrate / rust layer interface is very irregular, Fig. 7(a). High spatial resolution observations (Fig. 8) shows that the rust layer is preferentially formed around graphite spheres. As it was previously explained, these sites (around graphite spheres) are enriched in ferrite grains (Fig. 2(a,d)). This observation indicates that the ferrite is preferentially oxidized.

Micro-Raman spectroscopy measurements in the rust layer

Phase identification was performed in the rust layer using micro-Raman spectroscopy. Measurements were carried out on the crosssection surface, in the inner dense layer (sites S3 in Fig. 7(b) and S5 in Fig. 7(c)) and in the outer porous layer (site S6 in Fig. 7(c)). goethite α-FeO(OH) [START_REF] Oh | Characterization of iron oxides commonly formed as corrosion products on steel[END_REF][START_REF]RRUFF database: Goethite R120086[END_REF], hematite α-Fe 2 O 3 [START_REF] Bellot-Gurlet | Raman studies of corrosion layers formed on archaeological irons in various media[END_REF][START_REF] Hanesch | Raman spectroscopy of iron oxides and (oxy)hydroxides at low laser power and possible applications in environmental magnetic studies[END_REF][START_REF] Oh | Characterization of iron oxides commonly formed as corrosion products on steel[END_REF], lepidocrocite γ-FeO(OH) [START_REF] Oh | Characterization of iron oxides commonly formed as corrosion products on steel[END_REF], maghemite γ-Fe 2 O 3 [START_REF] Hanesch | Raman spectroscopy of iron oxides and (oxy)hydroxides at low laser power and possible applications in environmental magnetic studies[END_REF][START_REF] Jacintho | Surface-enhanced Raman spectra of magnetic nanoparticles adsorbed on a silver electrode[END_REF], magnetite Fe 3 O 4 [START_REF] Hanesch | Raman spectroscopy of iron oxides and (oxy)hydroxides at low laser power and possible applications in environmental magnetic studies[END_REF][START_REF] De Faria | Raman microspectroscopy of some iron oxides and oxyhydroxides[END_REF][START_REF] Shebanova | Raman spectroscopic study of magnetite (FeFe 2 O 4 ): a new assignment for the vibrational spectrum[END_REF] and siderite FeCO 3 [START_REF] Hanesch | Raman spectroscopy of iron oxides and (oxy)hydroxides at low laser power and possible applications in environmental magnetic studies[END_REF].

Fig. 13 shows the relative intensities of Raman bands in relation to intensities of hematite band 610-611 cm -1 from different places in the rust layer after corrosion test. The Raman bands with minimal band overlap were chosen for each iron compound. The two sites located on both sides of the crack in the inner dense layer (S3 and S5 in Fig. 12) contain the same chemical compounds: goethite, akaganeite and hematite. Lepidocrocite was not present in the inner dense layer. Although the content of goethite (which is considered to be the most stable and protective compound [START_REF] Singh | Corrosion behavior of IF steel in various media and its comparison with mild steel[END_REF][START_REF] Lair | Electrochemical reduction of ferric corrosion products and evaluation of galvanic coupling with iron[END_REF][START_REF] Okada | Mossbauer studies on particle volume distribution of α-FeOOH in rust formed on weathering steel[END_REF]) is high in the inner dense layer, the content of other compounds (akaganeite + hematite) is also high. The ratio α/γ* of the inner layer is lower than unity, indicating that the rust layer is not protective [START_REF] Kamimura | Composition and protective ability of rust layer formed on weathering steel exposed to various environments[END_REF][START_REF] Yamashita | Recent Progress in the Study of Protective Rust-layer Formation on Weathering Steel[END_REF][START_REF] Alcántara | Marine atmospheric corrosion of carbon steel: a review[END_REF]. According to [START_REF] Yamashita | Recent Progress in the Study of Protective Rust-layer Formation on Weathering Steel[END_REF][START_REF] Alcántara | Marine atmospheric corrosion of carbon steel: a review[END_REF], the corrosion rate of weathering steels drops for values of α/γ* greater than 1.5. These results are in good agreement with the high values of the corrosion rate found. The lowest content of goethite was found in the outer porous layer (S6 in Fig. 12). A large amount of lepidocrocite was detected, with a small amount of hematite and akaganeite. The largest amount of lepidocrocite was found in the crack (site S2 in Fig. 12), with a small amount of hematite and goethite (no akaganeite). Lepidocrocite is semi-conducting and electrochemically active [START_REF] Singh | Corrosion behavior of IF steel in various media and its comparison with mild steel[END_REF][START_REF] Monnier | XAS and XRD in situ characterisation of reduction and reoxidation processes of iron corrosion products involved in atmospheric corrosion[END_REF][START_REF] Kamimura | Mössbauer spectroscopic study of rust formed on a weathering steel exposed for 15 years in an industrial environment[END_REF]. It can be electrochemically reduced into Fe II soluble species and/or soild Fe II phase. Therefore, the chemical composition was found to depend markedly on the sites analysed: inner dense layer, outer porous layer and cracks.

Fig. 14 shows the schematic sketching for the corrosion mechanism of ADI under salt spray exposure. During the very first cycles (Fig. 14 (a)), preferential dissolution of ferrite occurs. Slight oxidation of the ausferrite matrix is also observed. Iron dissolves into Fe 2+ ions which then undergoes hydrolysis (Fe(OH) + ) and oxidation (Fe 3+ ). Fe(OH) + can transform through rapid oxidation into lepidocrocite and slow oxidation into Fe a 2+ Fe b 3+ O x (OH) y . Within this period, the corrosion rate increases significanty. With increasing number of cycles (Fig. 14(b)), strong oxidation of the ausferrite matrix takes place. Graphite spheres which are not oxidized were used as markers to show that oxidation is due to the inward diffusion of oxygen in the matrix. This compact layer protect ADI against corrosion and then the corrosion rate slightly decreases. From a certain number of cycles (depending of physical-chemical conditions at the specimen surface), a mixed oxidation process appears. The surface is covered with a thick rust layer (Fig. 14(c)), because of the outward diffusion of metallic ions, and the inner layer continues to grow (inward diffusion of oxygen). The outer layer contains numerous microcracks and pores. It acts as a water retention phase and absorbs the Cl - ions. Cracks can also be formed in the inner layer. Then, the corrosion rate sharply increases. Lepidocrocite is found in the outer layer whereas goethite (end-product) is present in the inner layer. Hematite and akaganeite are found throughout the two layers.

Conclusions

The corrosion behavior of ADI was studied under cyclic salt spray exposure. The following conclusions can be drawn: 1) A specific electrochemical cell was used to perform in-situ electrochemical measurements (OCP vs time, polarisation curves) during the wet periods. From these mesurements, the corrosion rate was calculated using the resistance polarisation method. 2) A schematic sketching for the corrosion mechanism of ADI under salt spray exposure was then proposed. It was decomposed into three steps: (i) pitting corrosion of ferrite (very first cycles), (ii) inward oxidation of the ausferrite matrix associated with the formation of a dense inner layer and (iii) mixed oxidation of the matrix associated with the formation of a rust layer at the specimen surface. 3) Evolution of the corrosion rate was correlated with surface changes induced by corrosion in the different steps. In the next work, additional in-situ electrochemical measurements will be carried out to correlate the electrochemical data with the corrosion state of ADI. The purpose of the correlation is to identify the corrosion state simply through the observation/analysis of the electrochemical data so as to establish the guidelines for monitoring the corrosion state from in-situ electrochemical measurements.

4) Goethite (end-product) and lepidocrocite were found in the inner and outer layers, respectively (hematite and akaganeite in both layers). 5) The corrosion rate decreases when inward oxidation occurs (protective layer) and then sharply increases in the presence of the outer porous layer (mixed oxidation).

Fig. 1 .

 1 Fig. 1. Images of the specific cell (WE = working electrode, RE = reference electrode, CE = counter electrode) used to perform electrochemical measurements in the salt fog chamber: (a) after surface preparation and (b) after testing for 42 cycles.

Fig. 2 .

 2 Fig. 2. (a-c) FE-SEM micrographs of ADI after etching in Nital solution for 5 s. (d) Optical image of ADI after etching in the B-M reagent for 30 s.

Fig. 4 .

 4 Fig. 4. Evolution vs the number of cycles of (a) current density at various applied potentials, (b) the corrosion current density and the corrosion rate and (c) the OCP.

Fig. 3 .

 3 Fig. 3. (a) OCP of ADI vs time and (b) polarisation curves (3 mV/s) of ADI for different cycles.

Fig. 5 .

 5 Fig. 5. (a-d) Surface observations of the ADI sample after various numbers of cycles : (1) = optical images and (2) = FE-SEM micrographs at low spatial resolution.

Fig. 6 .

 6 Fig. 6. (a-f) FE-SEM micrograph at high spatial resolution of the ADI sample after various numbers of cycles.

Stage I (cycles # 1 -# 3 ).

 13 Fig.4(a-c) shows that the electrochemical parameters vary significantly vs the number of cycles, indicating that the corrosion behaviour of ADI changes with time. To interpret these changes, optical and FE-SEM/EDS observations at different spatial resolution were carried out after 3, 9, 20, 30 and 42 cycles (Figs.5 and 6). Stage I (cycles #1 -#3). The current density in the anodic domain increases by one order of magnitude (from roughly 25 μA/cm 2 in cycle #1 up to 200 μA/cm 2 in cycle #3) whereas the current density in the cathodic domain increases up to a value of about -50 μA/cm 2 in cycle #3, Fig. 4(a). In addition, the corrosion rate increases, from nearly 0 μm/year (cycle #1) up to 70 μm/year (cycle #3), Fig. 4(b). During these three cycles, the OCP is very low (of about -800 mV, Fig. 4(c)).This tendency suggests that the surface is active. Surface observations after cycle #3 confirm this assumption. The microstructure of ADI is still visible in numerous sites, Fig.5(a1-a2) and 6(a-b). EDS analyses indicate that these sites are slightly oxidized (low oxygen content): 14.9 ± 0.2at.% C, 18.1 ± 4.3 % O, 2.6 ± 1% Na, 2.9 ± 0.5 % Si, 0.05 ± 0.05 % Cl and 60.9 ± 6.1 % Fe. There is no difference level between the graphite spheres (not oxidized) and the "ausferrite" matrix, due to the inward diffusion of oxygen in the matrix. A high density of pits was observed in these sites, Fig.6(a-b). As it was already observed in bulk electrolyte solutions[START_REF] Krawiec | Influence of the heat treatment on the microstructure and corrosion resistance of austempered ductile iron (ADI)[END_REF], pits initiate in the ferrite. Therefore, pitting corrosion and inward oxidation of ADI mainly occur during stage I. However, in some sites, the surface is covered with a thick rust layer (Fig.5(a2)), because of the outward diffusion of metallic ions. Numerous cracks are visible in this layer.Stage II (cycles #3 -#9). By contrast to stage I, the "ausferrite" matrix of ADI is not visible on the whole surface. Only graphite spheres can be observed (Figs. 5(b2) and 6 (c)). They are not dissolved and can still act as cathodic sites. At the end of stage II, the entire surface (except graphite spheres) is strongly oxidized and a dense rust layer is present (Fig.5(b2)). EDS analyses gives (performed on 4 different sites): 22.4 ± 14.7at.% C, 48.6 ± 10.1 % O, 0.3 ± 0.1 % Na, 0.9 ± 0.5 % Si, 0.3 ± 0.2 % Cl and 27.5 ± 4.9 % Fe. There is again no significant difference level between the rust layer and graphite spheres (Fig.6(c)), because of inward diffusion of oxygen in the matrix. Small cracks are found around some graphite spheres, Fig.6(c). The dense rust layer can help to prohibit further corrosion of ADI. This was confirmed from the evolution of the electrochemical parameters. Indeed, a slight decrease of the anodic

Fig. 7 .

 7 Fig. 7. (a) Optical images at low magnification of corrosion products. (b-c) Optical images at high magnification of sites located in (a) where micro-Raman measurements were performed (sites S2, S3, S5 and S6). Sites S1 and S4 correspond to graphite sphere in the substrate and in the corrosion products layer, respectively.

Fig. 8 .

 8 Fig. 8. Optical image of the interface between the substrate and the inner dense layer.

Fig. 9 .

 9 Fig. 9. Raman spectrum from site S6 in Fig. 7(c) (the spectrum range 85-900 cm -1 of Raman shift). The main phases are lepidocrocite (L), hematite (H) and akaganeite (A), with a small quantity of goethite (G).

Fig. 10 .

 10 Fig. 10. Raman spectrum from site S5 in Fig. 7(c) (the spectrum range 85-900 cm -1 of Raman shift). The main phases are akaganeite (A), goethite (G) and hematite (H).

Fig. 11 .

 11 Fig. 11. Raman spectrum from site S2 in Fig. 7(a) (the spectrum range 85-900 cm -1 of Raman shift). The main phases are lepidocrocite (L), hematite (H) and goethite (G).

Fig. 12 .

 12 Fig. 12. Raman spectrum from site S3 in Fig. 7(b) (the spectrum range 85-900 cm -1 of Raman shift). The main phases are akaganeite (A), hematite (H) and goethite (G).

Fig. 13 .

 13 Fig. 13. Relative intensities of lepidocrocite (Raman shift: 250 cm -1 ), akaganeite (Raman shift: 718-720 cm -1 ) and goethite (Raman shift: 718-720 cm -1 ) bands in relation to intensities of hematite band (Raman shift: 610-611 cm -1 ) from different sites (Raman bands with minimal band overlap were chosen).

Fig. 14 .

 14 Fig. 14. (a-c) The schematic sketching for the corrosion mechanism of ADI under salt spray exposure.

Author statement H. Krawiec: project administration, realization of in-situ measurements, analysis of corrosion tests, writing, review, editing, conceptualization, methodology.

V. Vignal: realization of in-situ measurements, SEM observations, analysis of corrosion tests, writing, conceptualization, methodology.

J. Lelito: elaboration of ADI, heat treatments of ADI, microstructural analysis.

A. Krystianiak: realization of salt fog tests, optical observations, preparation of specific cells.

P. Ozga: Raman measurements and analysis.

Declaration of Competing Interest

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