Harish Banda

Sandy Barbara Pé

Daffos

Lionel Dubois

Olivier Crosnier

Patrice Si

Pierre-Louis Tabe Rna

Flore Nce Duclairoir

Investigation of ion transport in chemically tuned pillared graphene materials through electrochemical impedance analysis

Keywords:

ehemically tuned pillared graphene structures show ability to limit restacking of graphene sheets for electrochemical energy storage in ses. A comprehensive electrochemical characterization using various ion sizes allowed identification of ion sieving in the cross linked galleries of reduced pillared graphene materials (RPs ). The access to the cross linked galleries, which provide additional ion sorption sites, offered slightly increased capacitances in RPs compared to completely restacked sheets in reduced graphene oxide (RGO). We performed electrochemical impedance analyses on RPs and RGO to under stand the ion transport inside the cross linked graphene galleries. RGO adsorbs ions in the inter particle micro/meso pores and the ion access to such sites from the bulk electrolyte occurs with relative ease. RPs sieve ions into their inter layer gallery pores based on effective ion sizes and the ion transport process is resistive compared to RGO. A control study using 3D pillared graphene hydrogel with improved macro porosity assigns this resistive behavior and the moderate capacitances to limited ion access to the active sites due to excess number of pillars. The obtained results on the ion transport dynamics between graphene layers provide perspectives towards further optimization of these graphene materials for ses.

Recently, our group proposed pillared graphene structures with

Introduction

Supercapacitors (Ses) or electrochemical double layer capaci tors are electrochemical energy storage devices able to deliver high power densities with excellent cycle life [START_REF] Miller | Electrochemical capacitors for energy management[END_REF][START_REF] Conway | Electrochemical Supercapacitors -Scientific Fundamentals and Technological Applications[END_REF]. These characteristics are made possible by the charge storage process involving a charge separation at the electrode electrolyte interface via fast adsorption of electrolytic ions at the electrode surface. However, these systems are also characterized by limited energy densities and thus are widely studied as complementary to batteries [START_REF] Simon | Materials for electrochemical capacitors[END_REF]. To improve the energy densities, researchers are working on new systems with porous carbons, or the development of pseudocapacitive materials with additional faradaic processes, or hybrid systems that operate E-mail addresses: tabema@cbimie.ups-tlse.fr (P.-L. Taberna), florence.duclairor@ cea.fr (F. Duclairoir). bttps://doiorg/10.1016/j.electacta.2018.11.122 between batteries and ses [START_REF] Presser | Carbons and electrolytes for advanced supercapacitors[END_REF][START_REF] Zhang | Carbon-based materials as supercapacitor electrodes[END_REF][START_REF] Wang | A review of electrode materials for electrochemical supercapacitors[END_REF].

Among various carbons, activated carbons (Aes) have been widely studied for ses applications owing to their high specific surface areas (SSA) and low material costs [START_REF] Presser | Carbons and electrolytes for advanced supercapacitors[END_REF]. Interest in Aes also lies in their tunable porous structures that allowed analysis of the impact of pore sizes on specific capaà tanoes [START_REF] Chmiola | Anomalous increase in carbon capacitance at pore sizes less than 1 nanometer[END_REF][START_REF] Salitra | Carbon electrodes for double-layer capacitors I. Relations between ion and pore dimensions[END_REF]. These studies revealed that in addition to high surface areas, matching pore sizes to electrolytic ion sizes also enable high capaà tances [START_REF] Largeot | Relation between the ion size and pore size for an electric double-layer capacitor[END_REF][START_REF] Raymundo-Pinero | Relationship between the nanoporous texture of activated carbons and their capacitance properties in different electrolytes[END_REF][START_REF] Salanne | Efficient storage mechanisms for building better supercapacitors[END_REF]. Extensive analyses on ion sieving in these pores have discovered that the high capacitances are a result of a partial desolvation and confinement of ions in these pores [START_REF] Chmiola | Desolvation of ions in subnanometer pores and its effect on capacitance and double-layer theory[END_REF][START_REF] Pean | Confinement, desolvation, and electrosorption effects on the diffusion of ions in nanoporous carbon electrodes[END_REF][START_REF] Merlet | Highly confined ions store charge more efficiently in supercapacitors[END_REF]. Among other carbon derivatives tested, graphene has been studied as an electrode ma terial candidate [START_REF] Raccichini | The role of graphene for electrochemical energy storage[END_REF][START_REF] Liu | Graphene-based supercapacitor with an ultrahigh energy density[END_REF][START_REF] Stoller | Graphene-based ultracapacitors[END_REF]. The derived research paths have mainly focused on tuning graphene sample porosity [START_REF] Tao | Towards ultrahigh volumetric capacitance: graphene derived highly dense but porous carbons for supercapacitors[END_REF][START_REF] Qu | Dense 3D graphene macroforms with nanotuned pore sizes for high performance supercapacitor electrodes[END_REF], for example, by creating hales inside the graphene sheets, or by forming gra phene aerogels displaying porosities adapted for efficient electro lyte ion transfer [START_REF] Xu | Holey graphene frameworks for highly efficient capacitive energy storage[END_REF][START_REF] Hantel | Partially reduced graphite oxide for supercapacitor electrodes: effect of graphene layer spacing and huge specific capacitance[END_REF]. In parallel, many research paths involve the use of graphene as a support for a faradaic or non faradaic component to improve the overall capacity of the system [START_REF] Li | Functional gels based on chemically modified graphenes[END_REF][START_REF] Sidhureddy | From graphite to interconnected reduced graphene oxide: one-pot synthesis and supercapacitor application[END_REF]. organic molecules as pillars to prevent the significant restacking of graphene sheets during chemical reduction. We predict that such pillared structures would offer higher specific capacitances with more electrochemical active sites available for ion sorption. Hence, an interesting class of materials with varied inter graphene layer separation (d spacing), named as reduced pillared materials (RPs), were readily synthesized by varying the size of the pillar (alkyldi amine with 5, 6, 8 Cs) used [24e26]. In the course of related elec trochemical studies, ion sieving behavior was observed in these graphene derived materials [START_REF] Banda | Ion sieving effects in chemically tuned pillared graphene materials for electrochemical capacitors[END_REF]. A clear dependency between d spacing of the inter layer galleries and electrolytic ion sizes was noted: ions with sizes smaller than d spacing can access gallery pores whereas larger ions cannot access the surface area inside the galleries [START_REF] Banda | Ion sieving effects in chemically tuned pillared graphene materials for electrochemical capacitors[END_REF]. The cross linked galleries in these materials sieve ions based on their size and hence demonstrate a potential for enhanced capacitances through optimization of d spacing/ion size couple [START_REF] Eliad | Ion sieving effects in the electrical double layer of porous carbon Electrodes: estimating effective ion size in electrolytic solutions[END_REF]. While this phenomenon has been clearly identified and explained using cyclic voltammetry (CV) experiments, additional electrochemical impedance spectroscopy (EIS) studies are pre sented in this paper with further insights on the electrolytic ion transport in between the graphene layers. This analysis of ion dy namics inside the graphene galleries gives new perspectives and is of direct interest towards optimizing graphene materials for SCs. So, herein, pillared graphene materials with a d spacing of 0.78 and 0.80 nm obtained with pillars of 1,5 diaminopentane and 1,6 diaminohexane respectively e were tested in various tetraalky lammonium tetrafluoroborate salt solutions in acetonitrile. Impedance analyses were performed under positive and negative polarizations to obtain the distinct anion and cation responses in these materials. Reduced graphene oxide (RGO), which only dis plays inter particle porosity but does not have any inter layer cross linked galleries, has been studied as a control material. Pillared graphene hydrogel using 1,6 diaminohexane (6 GH) with 3D porous network is also synthesized as a control material to inves tigate the role of porosity in 6 RP. EIS data of the materials have been analyzed using conventional Nyquist and Bode plots and a new set of analyses (evolution of imaginary part of capacitances vs. real impedances) is also proposed to investigate their ion transport characteristics.

Experimental methods

The pillared graphene based materials studied here were pre pared using graphene oxide as a precursor with alkyldiamines as pillars [START_REF] Banda | Ion sieving effects in chemically tuned pillared graphene materials for electrochemical capacitors[END_REF]. The amine groups react with oxygen containing func tions such as epoxide and carboxylic acid on GO and bridge GO sheets together (Scheme 1) [START_REF] Herrera-Alonso | homme, Intercalation and stitching of graphite oxide with diaminoalkanes[END_REF]. Graphene oxide (GO), used as a precursor for pillared graphene materials, was synthesized using a modified Hummers Offeman's method [START_REF] Hummers | Preparation of graphitic oxide[END_REF][START_REF] Kumar | Synthesis of high quality reduced graphene oxide nanosheets free of paramagnetic metallic impurities[END_REF]. GO sheets were first reacted with respective diamines to form cross linked gra phene materials which are then chemically reduced with hydrazine to form reduced pillared materials (RPs) (Scheme 1) [START_REF] Banda | One-step synthesis of highly reduced graphene hydrogels for high power supercapacitor applications[END_REF][START_REF] Li | Processable aqueous dispersions of graphene nanosheets[END_REF]. Depending on the diamine used in the first step, materials were named as 5 RP (1,5 diaminopentane) and 6 RP (1,6 diaminohexane). RGO was synthesized by a direct chemical reduction of GO with hydrazine hydrate. 6 GH was synthesized by treating GO and 1,6 diaminohexane at 180 C in a sealed autoclave. The detailed syntheses protocols are published elsewhere and a brief description can be found in supplementary information [START_REF] Banda | Ion sieving effects in chemically tuned pillared graphene materials for electrochemical capacitors[END_REF]. Electrodes were fabricated by preparing an ink using a N methyl 2 pyrrolidone (NMP) slurry of the synthesized materials (80%) with polyvinylidene fluoride binder (10%) and carbon black (10%) conductive additive. This ink is drop casted on stainless steel disks of 0.4 cm 2 area and dried at 65 C under air for 6 h followed by 80 C under vacuum overnight. The dried electrode specific weights were around 3e4 mg/cm 2 .

Electrochemical analyses of the materials were performed in three electrode cell configuration using a silver wire as the quasi reference electrode, an overcapacitive (10 times by weight) YP50 porous carbon pellet as counter electrode and 20 mm cellulose piece as the separator in Swagelok ® cells. 1 M tetraalkylammonium tet rafluoroborate (TAABF 4 ) salt solutions in acetonitrile with constant anion size and varying cation sizes; ethyl (TEABF 4 ), propyl (TPABF 4 ), butyl (TBABF 4 ) and hexyl (THABF 4 ) were used as electrolytes. The unsolvated sizes of the electrolyte ions tested are TEA þ e 0.68 nm, TPA þ e 0.76 nm, TBA þ e 0.82 nm, THA þ e 0.95 nm and BF 4 À e 0.48 nm. Anhydrous acetonitrile and electrochemical grade elec trolyte salts for the study were purchased from Sigma Aldrich. Electrochemical cells were assembled inside an Ar filled glovebox and once sealed have been electrochemically tested outside using a VMP 3 potentiostat device from Bio Logic. Prior to EIS measure ments, the cells have been cycled at 50 mV/s scan rates in CVs until the electrolyte wets the whole porous material volume and the electrochemical signature reached gets a steady state e the area under the curve stabilizes (~100 cycles). For EIS measurements, cells were held at þ1 V or 1 V vs. Ag in this 3 electrode configu ration until the current is almost null and then studied with sinu soidal voltage amplitude of 5 mV within a frequency range spanning from 100 kHz to 10 mHz. Obtained impedance values were normalized with electrode weights for plotting.

Results and discussion

Reduced pillared materials: characteristics

In our previous work, pillared graphene structures in RPs were analyzed using a combination of spectroscopic and electrochemical characterization methods (Fig. S1) [24e26]. X ray diffraction (XRD) patterns of RGO, synthesized as a control, showed a broad peak around 28 corresponding to the graphite like arrangement (002) with an inter graphene sheet distance of 0.37 nm (Fig. S1a). This graphite like arrangement suggested an efficient reduction of GO leading to a partial restacking of the reduced graphene sheets [START_REF] Moon | Reduced graphene oxide by chemical graphitization[END_REF]. Whereas, the XRD patterns of RPs displayed two peaks corre sponding to partial graphitic restacking (~25 ) and diamine induced pillaring of graphene layers (~15 ). Thus, RPs demonstrate pillared structures characterized with d spacing of 0.78 nm in 5 RP and 0.80 nm in 6 RP. The N 2 gas sorption analyses have assigned a high SSA of 330 m 2 /g for RGO and lower values of ~330 m 2 /g for RPs. Lower surface areas in RPs compared to RGO could arise from excessive inactive weight of alkyldiamines in RPs, structural changes during degassing and the diamine pillars may also steri cally hinder the gas molecules and render the inter layer gallery pores inaccessible [START_REF] Jagiello | 2D-NLDFT adsorption models for carbon slit-shaped pores with surface energetical heterogeneity and geometrical corrugation[END_REF][START_REF] Srinivas | Porous graphene oxide frameworks: synthesis and gas sorption properties[END_REF]. The pore size distributions (PSDs), calculated using 2D non local density functional theory (2D NLDFT), attribute high volume of micropores and a significant contribution of mesopores for RGO and a high mesopore content with some micropores in RPs (Fig. S1c) [START_REF] Burress | Graphene oxide framework materials: theoretical predictions and experimental results[END_REF].

The electrochemical responses of RGO and RPs in various elec trolytes with increasing cation sizes and constant anion size investigated the ion sieving effect in pillared graphene materials (Fig. 1). CVs of RGO and RPs showed similar rectangular quasi ideal capacitive behavior in TEABF 4 but, an increase in cation size from TPABF 4 to TBABF 4 and THABF 4 resulted in current decay under negative polarizations for RPs (Fig. 1 & Fig. S2a). Thus, the access to inter layer gallery pores in RPs is regulated by the d spacing/ion size couple whereas RGO doesn't have any pillared structures and hence, exhibits similar curves in different electrolytes. The calcu lated capacitances values were lower in RPs when the inter layer graphene galleries have a smaller d spacing than the ion sizes, implying that the galleries restrict larger ions (Fig. S2b). For instance, 5 RP shows a specific capacitance of 130 F/g in TEABF 4 but when larger cations were used, a steep decline from 115 F/g (TPABF 4 ) to 42 F/g (THABF 4 ) is noted. Whereas, RGO, which adsorbs ions only at the surface of graphene layers (SSA 330 m 2 /g), ex hibits similar capacitances of ~110 F/g in all the electrolytes studied. In short, this electrochemical study revealed ion sieving based on ion sizes inside the pillared graphene materials [START_REF] Banda | Ion sieving effects in chemically tuned pillared graphene materials for electrochemical capacitors[END_REF].

Impedance studies: Nyquist plots

To further characterize this ion sieving and also to gain insights into the ion transport inside the graphene galleries, EIS measure ments were performed on RGO and RPs under negative and posi tive polarizations. The normalized complex plane Nyquist plots are displayed in Fig. 2. These plots can be divided into three regions [36e38]. At high frequencies, only the electrical conductivity of the electrode (ionic and electronic) can be seen; thus, SC electrodes behave like a resistor (Z" 0). The sharp increase of the imaginary part of the impedance ( Z 00 ) in the low frequency region where the plot is nearly parallel to the y axis is related to a capacitive behavior. The intermediate 45 region is associated to the ion transport in the porous network of the electrode that corresponds to the electrolyte resistance in the pores.

At 1 V vs. Ag, the electrodes are negatively polarized with respect to open circuit voltages (OCV) and cations are predomi nantly adsorbed at the electrode surfaces [START_REF] Segalini | Qualitative electrochemical impedance spectroscopy study of ion transport into sub-nanometer carbon pores in electrochemical double layer capacitor electrodes[END_REF]. The Nyquist plots show differences between RGO and RPs in all the electrolytes tested (Fig. 2aec). Firstly, in each electrolyte, as compared to RGO, RPs show a larger electrolyte resistance in the porous electrode network that extend into lower frequencies. Due to this resistance, the low frequency capacitive behavior in RPs shifts into higher re sistances. Secondly, with increasing cation size, the vertical line indicating capacitance in each RP gradually diverges away from the ideal 90 . The extent of this deviation from the ideal capacitance can be estimated by replacing the capacitance C with a Constant Phase Element (CPE, see Equation ( 1)) using a parameter [START_REF] Bard | Evidence for electronic and ionic limitations at the origin of the second voltage plateau in nickel electrodes, as deduced from impedance spectroscopy measurements[END_REF]. a values are obtained by measuring the slopes in equation ( 1) at low frequencies. log z 00 ðwÞ a,logðCuÞ 0 < a < 1

A value of 1 for a indicates a fully capacitive behavior of the material. 6 RP shows decline of a values from 0.87 in TEABF 4 to 0.77 in TBABF 4 while, 5 RP shows low value of ~0.80 in all the electro lytes indicating a non ideal capacitive behavior (Table S1). Thus, there is an important influence of the cation size on the frequency responses of RPs under negative polarization. RGO exhibits a values close to 1 from TEABF 4 (0.95) to TBABF 4 (0.93) electrolytes, evidencing a faster ion transport in the pores.

Ion sorption in RGO occurs in the micro/meso inter particle pores and thus the TEA þ , TPA þ and TBA þ cations access them with relative ease. Therefore, the impedance measurements of RGO show ideal capacitive like behavior and do not vary a lot in these electrolytes. In RPs, ion sorption occurs significantly in the inter layer porous network in addition to the inter particle pores. Ac cess to the inter layer pores in RPs is regulated by d spacing/ion size couple. Therefore, the frequency responses of RPs are .;

E ê ê 0 0 0 z z z .1 -1 -1 -1 0 -1 0 -1 0 Potentlal (V) vs. Ag Potentlal M vs. Ag Potential (V) vs. Ag
dependent on their d spacing and the effective size of the elec trolyte ions. 6 RP, with a d spacing of 0.80 nm, has a performance similar to that of RGO in TEABF 4 but extended diffusion regions are seen with larger cations. Impedance data in THABF 4 electrolyte further supports the different ion sorption mechanisms for RGO and RPs (Fig. S3a). In RPs, THA þ cation (0.95 nm) is larger than the d spacing values, like with TBA þ , and hence even larger imped ances were recorded. Notably, RGO possesses a major contribution from inter particle pores with a diameter of ~0.9 nm and thus only when THA þ cations are used, a sudden increase in the Warburg region and decrease in a parameter (0.83) are observed (Table S1).

At þ1 V vs. Ag, the materials are positively polarized with respect to open circuit voltages and BF 4

À ions are predominantly adsorbed on the electrode surfaces. The Nyquist plots of RGO and RPs show similar shapes among themselves in the three different electrolytes tested (Fig. 2def). The BF 4 À ions are smaller than the gallery pores in RPs and hence display capacitive like behavior similar to RGO. However, the vertical part in the low frequency region slightly shifts away from the ideal behavior from TEABF 4 to TBABF 4 . a values for RGO range from 0.97 to 0.95 whereas RPs range from 0.93 to 0.86 in TEABF 4 to TBABF 4 (Table S1). This interesting shift from ideal behavior despite analyzing the same BF 4 À ions in different electrolytes hints at a possible role of cations even under positive polarization, such as previously observed [START_REF] Griffin | In situ NMR and electrochemical quartz crystal microbalance techniques reveal the structure of the electrical double layer in supercapacitors[END_REF]. In summary, RGO and RPs show similar responses in all electrolytes under positive potentials and an increasing influence of larger cations, in addition to anions, is also noted. This observation was described as an ion exchange mechanism in earlier reports [41e43]. With the largest cation in THABF 4 , a values of RGO and RPs decrease to 0.90 and ~0.76 respectively, indicating a greater cation exchange effect.

Impedance studies: complex capacitance analysis

In order to further explore the frequency responses of the syn thesized materials, we performed a complex capacitance analysis that was proposed by Taberna et al. and derived from Cole and Cole [44e46]. In the proposed analysis, complex capacitance is studied in order to focus on the capacitive behavior and ionic resistance within the active material. The closest equivalent circuit repre sentation of the capacitor in this analysis has a series resistance, Rs, together with an angular frequency dependant R p C p bridge: a combination of resistance, R p (u), in parallel with a capacitance, C p (u) (Fig. S5).

This complex analysis allowed determination of the real capacitance, C 0 (w), which represents the deliverable capacitance of the cell under DC or low frequency AC conditions. C 0 (w) is (2)

The evolution of C 0 with respect to frequency for each sample at 1 V and 1 V vs. Ag was plotted accordingly (Fig. 3). At higher frequencies, the system acts as a pure resistor with capacitance values close to zero and at lower frequencies, a pure capacitive behavior is seen as a plateau if the maximal C 0 occurs in the tested frequency range. The intermediate zone gives information about electrolyte access to the electrode porosity (gallery and inter particle pores) [START_REF] Taberna | Electrochemical characteristics and impedance spectroscopy studies of carbon-carbon supercapacitors[END_REF].

At 1 V vs. Ag, C 0 vs w plots for RGO approach a maximal C 0 plateau in TEABF 4 , TPABF [START_REF] Presser | Carbons and electrolytes for advanced supercapacitors[END_REF] and TBABF 4 at the lowest frequency tested (10 mHz), revealing a fast transport of electrolytic ions in the electrode material (Fig. 3aec). However, in THABF 4 , RGO displays lower C 0 values and a non ideal behavior as the cation size ap proaches the inter particle pore sizes in RGO (Fig. S4a). For 6 RP, C 0 is closely equivalent to that of RGO in TEABF 4 and increasing cation sizes shifts the pure capacitive behavior to lower frequencies. 5 RP, even in TEABF 4 , does not approach a plateau and the ion transport is gradually limited in other electrolytes. Additionally, with larger cations in RPs, the maximal C 0 values and the slopes in the inter mediate region gradually decrease.

This EIS data under negative polarization really shows a slower ion transport in RPs when larger cations are used. In RPs, steric effects when ion size is in the same range as the inter layer gallery size lead to limited ion transport and accessibility to adsorption sites. Thus, with a complex dynamic process, RPs do not reach a plateau with full capacitance in this frequency range and confirm resistive behavior. At þ1 V vs. Ag, RGO and RPs all show curves with similar maximal C 0 values and slopes in the intermediary region (Fig. 3def). However, with increasing cation size, it can be noted that the plateaus of C 0 slightly drift towards lower frequencies. This observation, as noted in the Nyquist plots, could be seen as an in fluence of cations via the existence of an ion exchange mechanism [START_REF] Griffin | In situ NMR and electrochemical quartz crystal microbalance techniques reveal the structure of the electrical double layer in supercapacitors[END_REF]. Results obtained with THABF 4 under positive potential show a steep fall of C 0 values in all the materials, supporting this effect (Fig. S4b). Moreover, an increase in the alkyl chain length from TEA þ to THA þ ions could give rise to stronger interactions between graphene sheets and alkyl carbons of the electrolytic ions, resulting in enhanced influence of cations.

The complex capacitance model also allows the determination of imaginary capacitance (C 00 ) which corresponds to the energy loss arising from irreversible processes at the electrodes [START_REF] Taberna | Electrochemical characteristics and impedance spectroscopy studies of carbon-carbon supercapacitors[END_REF]. This C 00 (w) parameter was calculated from EIS data using equation [START_REF] Simon | Materials for electrochemical capacitors[END_REF], where Z 0 (w) is the real part of impedance Z:

C 00 ðwÞ z 0 ðwÞ
w zðwÞj 2 (3)
The imaginary C 00 goes through a maximum at a frequency f 0 , which defines the transition between purely resistive and capaci tive behaviors of a SC. This frequency, known as knee frequency, defines the time constant (t 0 ) as minimum time required to extract half of the capacitance from a SC [START_REF] Cole | Dispersion and absorption in dielectrics I. Alternating current characteristics[END_REF]. C 00 vs w plots of RGO and RPs at þ1 and 1 V vs. Ag show similar and dissimilar curves respec tively, supporting the role of ion discussed above. (Figs. S6 andS7).

The t 0 values obtained from these curves assign better power capability for RGO than RPs (Table 1). Interestingly, the t 0 values for

þ1 V -1 V þ1 V -1 V þ1 V -1 V þ1 V -1 V R G

TEABF4 TPABF4 TBABF4

(a)

. RGO (b)

• RGO (c) positive polarization [START_REF] Cole | Dispersion and absorption in dielectrics I. Alternating current characteristics[END_REF]. For example, t 0 of 5 and 6 RPs in TBABF 4 are 11 and 20 s respectively. This difference between RPs is counter intuitive as one would expect 6 RP, with a higher d spacing, to facilitate better transport and lower t 0 value for BF 4 À anion. As t 0 values are affected by a combination of ionic, electrical and cell resistances, it seems important to determine the distinct impact of the material properties on ionic resistances.

To do so, herein, we studied C 00 with respect to Z 0 . We believed that such an analysis could enable facile quantitative determination of the effective ionic resistances (R i ) involved in ion transport. The R i values directly obtained here could be seen as equivalent to the one third of the in pore ionic resistance calculated by employing the Transmission Line Model [START_REF] Qu | Studies of activated carbons used in double-layer capacitors[END_REF]. Similar to other plots, three different regions could be identified in such a curve (Fig. 4). At high frequencies the SC performs like a resistor and C 00 would be close to zero (from Eq. ( 3)). The mid frequency range is related to the ion transport inside the electrode porous structure hence, a steep in crease in C 00 would be seen since there is no dielectric loss in an ideal capacitance when u/0. At low frequencies, an ideal SC would display a steep fall in C 00 (from Eq. ( 3)). The R i values were obtained from the real Z values at the start and end of the mid range frequencies.

The C 00 vs Z 0 plots at 1 V vs. Ag show a clear difference between RGO and RPs in TEABF 4 , TPABF 4 and TBABF 4 (Fig. 4aec). The re sponses of RGO and RPs in various frequency regions and their evolution in electrolytes could be distinctly identified here compared to Nyquist and Bode plots. RGO displays curves with a clear maxima of C 00 occurring at a low resistance and then rapidly decreasing along the x axis. The calculated R i show close values of around 10 mOhm.g from TEABF 4 to TBABF 4 (Fig. 4g). 6RP Ri is closely equivalent to that of RGO in TEABF 4 (Fig. 4a) but, with larger cations, the C 00 extends to even higher resistances and eventually shows a continuous plateau with TBABF 4 (Fig. 4c). The plateau symbolizes an extended limitation of ion transport inside the porous electrode while the cell is not reaching a capacitor like behavior. This increasingly resistive behavior of 6 RP with larger cations is quantified with rise in R i values from 20 mOhm.g in TEABF 4 to 70 mOhm.g in TBABF 4 . 5 RP plot shows continuous plateau with no transition peak even with the smallest cation tested in TEABF 4 and already shows a high R i of 55 mOhm.g (Fig. 4a). Further increase in cation sizes lead to even higher R i up to 172 mOhm.g in TBABF 4 (Fig. 4c). This continuous loss of energy even at the low frequencies led us to examine the leakage currents for all cells. The measured leakage currents show similar values for RGO and RPs in all electrolytes and suggest that the observed continuous energy losses are characteristics of the materials (Table S2). At þ1 V vs. Ag, RGO and RPs show similar curves, however, occurring at different values of real Z (Fig. 4def). The C 00 values for RPs occur at greater resistances than RGO indicating greater irreversible energy losses. Among the two RPs, 6 RP has slightly larger ionic resistances than 5 RP in all the electrolytes (Fig. 4 h). This resistive behavior of 6 RP supports the earlier observed greater -to values than 5 RP. We speculate that the conformation (trans & gauche) of the diamines inside the galleiy could be the reason behind this behavior. lt has been reported that with increasing alkyl chains lengths in the diamine the ratio of trans gauche conformations shifts towards higher trans content [START_REF] Herrera-Alonso | homme, Intercalation and stitching of graphite oxide with diaminoalkanes[END_REF]. Hexanediamine (in 6 RP) may locally occupy higher volume in the galleiy than pentanediamine (in 5 RP) and thus have higher ionic resistanoes. Moreover, the odd and even number of carbons in the diamines used here could also affect the overall polarity and ion sorption in the gallery. The cation sorption at 1 V, masks these fine details among RPs as the steric effects result in large Ri values.

As the Ri values at + 1 V are smaller, this subtle influence of the alkyl chain conformation on 5 and 6 RPs is seen distinctly. Overall, the proposed C' vs. Z' plots show an amplified differenoe between the materials compared to Nyquist, Bode plots or other representation proposed by Pickup et al. [START_REF] Moghaddam | Electrochemical impedance study of the polymerization of pyrrole on high surface area carbon electrodes[END_REF]. lt is also noteworthy that we obtain effective ionic resistances directly from the experimental data without using TLM or PSO TLM models [START_REF] Qu | Studies of activated carbons used in double-layer capacitors[END_REF][START_REF] Song | Electrochemical impedance spectroscopy of porous electrodes: the effect of pore size distribution[END_REF].

lnvestigating role of porosity

Earlier analysis of capacitance values from CVs have shown marginally higher values for RPs than for RGO with appropriately smaller ions (TEABF 4 ). However, even in the case of TEABF 4 , the impedance analysis has indicated that RPs exhibit sluggish ion transport relative to RGO. This could be explained by hindered mobility of the ions in the galleries due to unadapted porosity or the presence of alkyl diamine pillars. To investigate if the impeded ion transport is arising from unsuited porous network, we syn thesized and analyzed a pillared graphene hydrogel (GH) that has a 30 percolating network using 1,6 hexanediamine as the pillar (6 GH) (details in SI). The 30 percolating network offers m ultiple ion pathways and is known to show enhanoed ion transport char acteristics in the hydrogels [START_REF] Li | Functional gels based on chemically modified graphenes[END_REF]. The N 2 gas sorption isotherms and corresponding PSOs characterize 6 GH with enhanoed micro and meso pores compared to 6 RP (Fig. S9 a&b). The calculated SSAs indicate increase in the values from 130 m 2 /g in 6 RP to 182 m 2 /g in 6 GH (Table S3). Scanning electron microscopie images also reveal more open macro porosity in 6 GH than 6 RP (Fig. S9 c&d ). Addi tionally, 6GH shows greater d spadng of 0.86 nm compared to 0.80 nm of 6 RP probably due to the stronger reaction conditions and diamine consumption during 6 GH synthesis. The TGA profiles of the two materials support this expia nation with a greater weight Joss due to decomposition of covalently attached diamines.

Electrochemical analysis of 6 GH in 1 M TEABF 4 , which has sufficiently smaller ions than the d spacing, shows slightly higher capacitances of 142 F/g than 6 RP (130 F/g) and RGO (117 F/g). This shows that the 30 assembly with improved SSAs, along with the cross linking at nanoscale, yields enhanced capacitances in these pillared materials. However, the impedanoe analysis under nega tive polarizations shows that the ion motion remains veiy much impeded in 6 GH despite its 30 structure and the larger d spacing (0.86 nm) (Fig. Sa andb ). ln fact, the obtained Ri and '<o values show greater resistances in 6 GH with higher values of 57 mOhm.g and 23 s compared to 20 mOhm.g and 17 s for 6 RP (Table S3 ). These observations suggest that the alkyl diamine pillars inside the gal le ries play an important role in ion transport between the graphene layers. Moreover, the obtained capadtances of both 6 RP and 6 GH are only slightly better than RGO des pite enhanced ion access to the cross linked galleries. This is possibly due to excessive filling of the galleries with diamine pillars leading to impeded ion transport to active sites. This understanding of the impact of pillars and the porosity provides new strategies to improve energy and power 0 . 4 ------------ densities of pillared materials. Oecreasing the number of pillars between the graphene layers could be a straightforward strategy for enhanced power and energy values. lntegrating additional inter particle pores could be another strategy for enhancement of energy densities. Collectively, pillared materials with low pillar density and additional inter particle pores could offer both high energy and power densities.

Conclusions

lmpedance measurements have been performed to investigate the ion transport dynamics in pillared graphene based materials. A range of electrolytes with increasing cation sizes and a constant anion size were studied under positive and negative electrode polarizations. Oepending on the cation size, Nyquist and Bode plots ofRPs exhibited large Warburg like regions and delayed capadtive behaviors whereas RGO responses remained similar. Ionie re sistances obtained through original C' vs. 'l! plots enabled a quan titative analysis of these observations and show that the ion transport in RPs is mainly controlled by graphene inter layer galleiy sizes. As RPs sieve ions into the pillared graphene layers, steric ef fects impede the ion transport when ion size is in the same range as the inter layer galleiy size. RGO adsorbs ions in the inter particle micro/mesa pores and the ion access to the active sites is rela tively easier. The observed high ionic resistanoes for RPs compared to RGO are thus in accordance with the earlier CV experiments that displayed poor power capability for RPs at high scan rates. This knowledge from cyclic voltammetry and the dynamic ion transport between graphene layers materials opens up perspectives towards enhanced charge storage in graphene materials.

Scheme 1 .

 1 Scheme 1. Schematic representation of reduced graphene oxide (RGO) and reduced pillared graphene materials (RPs) showing the specific surface areas and d-spacing values determined from XRD. XRD patterns adapted with permission from Chem. Mater. 30, 2018, 3040-3047. Copyright (2018) American Chemical Society [24].

Fig. 1 .

 1 Fig. 1. Normalized CVs of RGO and RPs obtained at a scan rate of 10 mV/s in the voltage range of À1 V to 1 V vs. Ag in 1 M solutions of (a) TEABF 4 , (b) TPABF 4 and (c) TBABF 4 in acetonitrile. Normalization was performed with respective currents observed at open circuit voltages (~0.1 V).

Fig. 2 .

 2 Fig. 2. Nyquist plots of RGO and RPs in 1 M solutions of (a,d) TEABF 4 , (b,e) TPABF 4 and (c,f) TBABF 4 in acetonitrile at À1 V vs. Ag (a,b,c) and 1 V vs. Ag (d,e,f). A zoom of the high and mid-frequency regions is shown in the insets. Unsolvated ion sizes of TEA þ , TPA þ and TBA þ are shown on the right.

- 1 V

 1 vs. Ag 0.6 , -"'1n .1i::::== ::::= = =::::._ __ ___. ___ _._ __ __. o ----,,.,... ------------=--=--=--=--=--=.--, . .. 011L.---------='o.1 0.0 ---........ _ _ _ ......._ _ _ __.

  ---........ _ _ _ ......._ _ _ __.calculated using equation[START_REF] Conway | Electrochemical Supercapacitors -Scientific Fundamentals and Technological Applications[END_REF], where Z 00 (w) is the imaginary part of impedance Z and w is the frequency.

Fig. 3 .

 3 Fig. 3. Real capacitance (C') vs frequency plots for RGO and RPs in 1 M solutions of (a,d) TEABF 4 , (b,e) TPABF 4 and (c,f) TBABF 4 in acetonitrile at (a,b,c) À1 V vs. Ag and (d,e,f) 1 V vs. Ag.

  10"' 10• 1 10" 10 1 10 2 10' 10 4 10 1 0 10 .. 10"' 10•' 10' 10' 10 2 10' 10 4 10' 0 10 .. 10• 2 10•' 10' 10' 10 2 10' 10 4 10• Frequency (Hz) Frequency (Hz) Frequency (Hz)

Fig. 4 .

 4 Fig. 4. Imaginary part of capacitances (C'') vs. real impedances (Z') for RGO and RPs in 1 M solutions of (a,c) TEABF 4 , (b,d) TPABF 4 and (c,f) TBABF 4 in acetonitrile at (a,b,c) 1 V vs. Ag and (d,e,f) À1 V vs. Ag. The arrows in (a), (b) and (c) show increasingly larger ionic resistances (R i ) in 6 RP. The obtained R i values for RGO and RPs at À1 V vs. Ag and 1 V vs. Ag.

  .. L-_.__ _ _ .__ _ _ .__ _ __, o_ ..___.__ __ .__ __ .__ _ __,

-Fig

  Fig. S. ( a) Nyquist plots and (b) lmaginary part of capacitances vs. real impedance plots of RGO, 6 RP and 6 GH in 1 M solution of TEABF 4 at -1 V vs. Ar, A zoom of the high frequency regions in Nyquist plot is shown in the inset

Table 1

 1 Time constant (t 0 ) values (seconds) of RGO and RPs were calculated from the frequencies at which maximum values of C 00 occurred.

	Material	TEABF 4	TPABF 4	TBABF 4	THABF 4

Acknowledgments

This work was supported by CEA (H. Banda) and by French ANR Program (SUGGEST ANR 15 CE05 0022). The access to mutualized equipment and PFNC Nano characterization Minatec platforms was appreciated by the authors. G. Bidan and D. Aradilla are also thanked for fruitful discussions and advices.

Appendix A. Supplementary data

Supplementary data to this article can be found online at https://doi.org/10.1016/j.electacta.2018.11.122.


Marie-Laetitia De Bonfils-Lahovary

Claudie Josse

Lydia Laffont

Christine Blanc
email: christine.blanc@ensiacet.fr

Influence of hydrogen on the propagation of intergranular corrosion defects in 2024 aluminium alloy

Keywords: A, Aluminium A, Hydrogen B, EBSD C, Intergranular corrosion C, Interfaces

To study the influence of hydrogen on the intergranular corrosion mechanism of a 2024 aluminium alloy, samples were hydrogen precharged by cathodic polarisation and then exposed to a NaCl solution. EBSD analyses and SEM observations showed that hydrogen increased the number of corroded interfaces and led to the embrittlement of low-angle grain boundaries which were not susceptible to corrosion without hydrogen precharging. The increase of the reactivity of the 2024 aluminium alloy in the presence of hydrogen gave a new insight into the intergranular corrosion mechanism: corrosion-induced hydrogen promoted the intergranular corrosion propagation and partially controlled the corrosion defect morphology.

Introduction

2024 Al alloy (AA 2024) is one of the most widely used Al alloys in aerospace industries. It is a structural-hardening alloy with Cu and Mg as major alloying elements. The formation of hardening precipitates (Guinier-Preston (GP) zones or θ'-Al 2 Cu/S'-Al 2 CuMg phases depending on the metallurgical state) leads to high mechanical properties (yield strength and ultimate tensile strength) [1,[START_REF]Success stories: air force; material substitution and new sealing, 4, s[END_REF], but it is combined with a more largely heterogeneous microstructure, characterised also by the presence of intergranular precipitates, dispersoids and intermetallic coarse particles (IMC), that increases the susceptibility of the alloy to localised corrosion, such as intergranular corrosion (IGC) [START_REF] Wang | The intergranular corrosion susceptibility of 2024 Al alloy during re-ageing after solution treating and cold-rolling[END_REF][START_REF] Bonzom | Intergranular corrosion propagation rate of 2024 alloy investigated via the "one-dimensional artificial pit" technique[END_REF][START_REF] Knight | In situ X-ray tomography of intergranular corrosion of 2024 and 7050 aluminium alloys[END_REF]. Numerous papers in the literature show that pits initiated preferentially on IMCs in AA 2024 [START_REF] Derose | Localised corrosion initiation and microstructural characterisation of an Al 2024 alloy with a higher Cu to Mg ratio[END_REF][START_REF] Hu | A localized approach to study corrosion inhibition of intermetallic phases of AA 2024-T3 by cerium malate[END_REF][START_REF] Prakashaiah | Corrosion inhibition of 2024-T3 aluminum alloy in 3.5% NaCl by thiosemicarbazone derivatives[END_REF]; concerning IGC, the mechanism is usually described as galvanic coupling between intergranular Cu-rich particles (θ-Al 2 Cu; S-Al 2 CuMg) and the adjacent matrix [START_REF] Zhang | Transitions between pitting and intergranular corrosion in AA2024[END_REF][START_REF] Wloka | Detection of nanoscale η-MgZn 2 phase dissolution from an Al-Zn-Mg-Cu alloy by electrochemical mocrotransients[END_REF][START_REF] Guillaumin | Localised corrosion of 2024 T351 aluminum alloy in chloride media[END_REF][START_REF] Galvele | Mechanism of intergranular corrosion of Al-Cu alloys[END_REF].

Nevertheless, even if the precipitation state is a first order parameter for IGC, the AA 2024 susceptibility towards IGC depends on other parameters as, for example the level of misorientation (θ) of the grain boundaries. The influence of this parameter on corrosion phenomena has been studied on various materials, such as stainless steels, Ni base alloys and Al alloys [START_REF] Gertsmann | Study of grain boundary character along intergranular stress corrosion crack paths in austenitic alloys[END_REF][START_REF] Randle | The coincidence site lattice and the 'sigma enigma[END_REF][START_REF] Fang | Corrosion behaviors of random and special grain boundaries in a sensitized 304 stainless steel[END_REF][START_REF] Kim | Grain boundary character distribution and intergranular corrosion behavior in high purity aluminum[END_REF]. For AA 5182, Davenport et al. [START_REF] Davenport | Intergranular corrosion and stress corrosion cracking of sensitised AA5182[END_REF] have reported that high-angle grain boundaries (HAGB, θ > 20°) are susceptible to corrosion mechanisms, such as IGC due to the presence of β-Mg 2 Al 3 -phase precipitates. In another study performed on AA 2050, Guérin et al. have shown that, even if most of the interfaces in the alloy are low-angle grain boundaries (LAGB, θ < 10°), only HAGB (θ > 10°) are corroded after exposure to an aggressive NaCl solution [START_REF] Guérin | Identification of the metallurgical parameters explaining the corrosion susceptibility in a 2050 aluminium alloy[END_REF]. Other authors, such as Luo et al. have explained the susceptibility to intergranular corrosion of AA 2024 T351 taking the grain-stored energy into account [START_REF] Luo | Observations of intergranular corrosion in AA2024-T351: the influence of grain stored energy[END_REF]. As grain-stored energy reflects the dislocation density, they proposed that a high level of dislocation density within individual grains made them more susceptible to preferential attack at the grain boundary.

However, the corrosion processes associated with the galvanic coupling between intergranular Cu-rich particles (θ-Al 2 Cu, S-Al 2 CuMg) and the adjacent matrix lead to the production of Al cations; the hydrolysis of Al cations generates protons. This induces an acidification of the electrolyte trapped inside the IGC defects. Further, previous studies showed the formation of a thin metallic Cu-nanoparticles layer at the interface alloy-IGC defect [START_REF] De Bonfils-Lahovary | Characterization of intergranular corrosion defects in a 2024 T351aluminium alloy[END_REF], that could be due to dissolution of the Cu-rich intergranular precipitates [START_REF] Luo | Observations of intergranular corrosion in AA2024-T351: the influence of grain stored energy[END_REF][START_REF] Ma | Discontinuities in the porous anodic film formed on AA2099-T8 aluminium alloy[END_REF], and/or to the matrix dissolution [START_REF] Hashimoto | Structure of the copper-enriched layer introduced by anodic oxidation of copper-containing aluminium alloy[END_REF]. This Cu-rich layer could protect the interior of the attacked grains from further corrosive environment in AA 2024 [START_REF] Luo | Observations of intergranular corrosion in AA2024-T351: the influence of grain stored energy[END_REF][START_REF] Zhou | Grain stored energy and the propagation of intergranular corrosion[END_REF][START_REF] Hashimoto | Investigation of dealloying of S phase (Al2CuMg) in AA 2024-T3 aluminium alloy using high resolution 2D and 3D electron imaging[END_REF]. However, Cu could also provide effective cathodic support for oxygen and/or protons reduction on the walls of the IGC defects [START_REF] Hashimoto | Investigation of dealloying of S phase (Al2CuMg) in AA 2024-T3 aluminium alloy using high resolution 2D and 3D electron imaging[END_REF]. Bonfils et al. showed hydrogen (H) penetration inside the material during the corrosion processes with H amount reaching 20 ppm and even 120 ppm after 30 and 500 h of exposure to a 1 M NaCl solution, respectively [START_REF] De Bonfils-Lahovary | Characterization and control of the intergranular corrosion defects in a 2024 T351 aluminium alloy[END_REF]. Similarly, Larignon et al. have shown that H can be detected at HAGB and LAGB in corroded AA 2024 samples by scanning kelvin probe force microscopy (SKPFM) [START_REF] Larignon | Investigation of Kelvin probe force microscopy efficiency for the detection of hydrogen ingress by cathodic charging in an aluminium alloy[END_REF]. Lafouresse et al. have also

H precharging

Cathodic charging in 10 mM H 2 SO 4 solution was selected to introduce H inside the material since a large H amount can be inserted into pure Al with this method [START_REF] Birnbaum | Hydrogen in aluminium[END_REF]. A Pt counter electrode and a saturated calomel reference electrode (SCE) were used. The samples were 1 cm 3 cubes, the precharging duration was 40 h and the precharging potential was -0.8 V vs SCE; these parameters allowed a sufficient amount of H to be introduced inside the sample in a reasonable amount of time. The precharging side was polished with diamond paste down to 3 μm with distilled water as a lubricant before precharging; it was chosen to be perpendicular to the elongated grain direction (L) [START_REF] Lafouresse | A Kelvin probe force microscopy study of hydrogen insertion and desorption into 2024 aluminum alloy[END_REF][START_REF] Lafouresse | Hydrogen mapping in an aluminium alloy using an alternating current scanning electrochemical microscope (AC-SECM)[END_REF]. Therefore, one of the long transverse (LT) -short transverse (ST) sides of the samples was exposed to the electrolyte; all the other sides were protected by a varnish. After cathodic precharging, the H amount was determined by gas fusion analysis using a Bruker G8 GALILEO ON/H Instrumental Gas Analyser (IGA); the H content was 19 ppm in agreement with previous results [START_REF] Lafouresse | A Kelvin probe force microscopy study of hydrogen insertion and desorption into 2024 aluminum alloy[END_REF].

Corrosion tests

Corrosion tests were performed for both non-charged and H-precharged samples to make IGC defects grow along the L direction. The tests corresponded to continuous immersion of the samples (1 cm 3 cubes) in a 5 mM NaCl aerated solution at room temperature (25 °C). A low chloride concentration was chosen because, as previously said, corrosion phenomena lead to H production. To be able to study the influence of H on the corrosion behaviour of AA 2024, it was essential that the H amount generated by the corrosion processes remained lower than the H introduced by cathodic precharging for H-precharged samples. The duration of the corrosion tests was also adapted to allow the influence of H introduced by cathodic precharging to be revealed. Some corrosion tests were performed in a more concentrated NaCl solution (1 M) and for long duration (24 h) for two non-charged samples to study the evolution of the IGC defect morphology. Only the LT-ST faces of the samples were exposed to the NaCl solution, the other faces being protected by a varnish [START_REF] Zhang | Localized corrosion growth kinetics in AA2024 alloys[END_REF][START_REF] Zhang | Anisotropy of localized corrosion in AA2024-T3[END_REF]. Before exposure to the NaCl solution, the non-charged samples were abraded using SiC papers down to 5 μm; then, they were polished down to 3 μm using diamond paste with distilled water as a lubricant. For the H-precharged samples, only a very short polishing step with 3 μm diamond paste was done since they have been abraded then polished before the H-precharging. Unless it was short, this polishing step could lead to some H desorption. However, it corresponded to a perfectly defined protocol used for all samples so that it was assumed that this did not lead to misinterpretation of the results.

Characterisation of the IGC defects: preparation of the samples and description of the techniques

The evolution of the IGC defect morphology was first studied by using a MA200 NIKON OM. After the corrosion tests (6 h in a 5 mM NaCl solution), cross-sections (L-ST plane) of the corroded samples were prepared and observed to characterise the morphology of the corrosion defects for both non-charged and H-precharged samples. The results were obtained from the observations of 5 samples for each case (non-charged and H-precharged samples), which corresponded to 5 cm analysed in the ST direction for each case. In the following, mean values were given. For a more accurate description of the corrosion defects, transmission electron microscopic (TEM) observations were performed using a JEOL JEM2100 F equipped with a field-emission gun and an energy-dispersive X-ray spectrometer (EDX, SDD Bruker) operating at 200 kV at UMS Castaing. In order to obtain a thin sample in a localised region, i.e. one containing an IGC defect, the preparation was done using conventional lift out procedure using a FEI HELIOS Nanolab 600i dual beam FIB / Scanning Electron Microscope (SEM) [START_REF] De Bonfils-Lahovary | Characterization of intergranular corrosion defects in a 2024 T351aluminium alloy[END_REF]. Three corroded H-precharged samples were observed and representative used SKPFM analyses on corroded AA 2024 in order to study H desorption kinetics from corrosion defects and have distinguished between H trapped at different l ocations [START_REF] Lafouresse | A Kelvin probe force microscopy study of hydrogen insertion and desorption into 2024 aluminum alloy[END_REF]. T herefore, d ue t o t he large amount of H trapped inside the corroded materials, questions are being raised about its subsequent influence on the corrosion processes. Larignon et al. have clearly shown, for cathodically H-precharged AA 2024 samples, that H induced a shift towards more negative values of the corrosion potential [START_REF] Larignon | Investigation of Kelvin probe force microscopy efficiency for the detection of hydrogen ingress by cathodic charging in an aluminium alloy[END_REF]. Further, Lafouresse et al. have performed alternative current scanning electrochemical microscope (AC-SECM) measurements on H-precharged AA 2024 samples and have shown an increase of the local reactivity, i.e. higher local currents associated with a lower effective resistance, due to the H diffusion inside the samples [START_REF] Lafouresse | Hydrogen mapping in an aluminium alloy using an alternating current scanning electrochemical microscope (AC-SECM)[END_REF].

Similarly, the influence of H on the oxygen reduction reaction in a martensitic stainless steel has been shown by Schaller et al. [START_REF] Schaller | Spatially resolved mapping of the relative concentration of dissolved hydrogen using the scanning electrochemical microscope[END_REF][START_REF] Thomas | The effect of absorbed hydrogen on the dissolution of steel[END_REF]. For other materials and especially for Ni base alloys, the influence o f H localised at the grain boundaries on the propagation of intergranular cracks has been investigated. B arrows et al. have used molecular dynamics simulations for Ni base alloys to highlight the influence o f H localised at the grain boundaries on the intergranular de-cohesion and also showed that the increase of the H coverage of the Ʃ3 grain boundaries influenced t he c rack-tip p ropagation k inetics [START_REF] Barrows | Traction-separation relationships for hydrogen induced grain boundary embrittlement in nickel via molecular dynamics simulations[END_REF]. I n another study concerning Ni 3 Al in a H 2 -containing environment, Takahashi et al. have observed brittle fracture nucleation for randomly oriented grain boundaries, except for LAGB [START_REF] Takahashi | Direct evaluation of grain boundary hydrogen embrittlement: a micromechanical approach[END_REF]. On the contrary, for coherent Ʃ3 boundaries, no grain-boundary fracture was observed. The H diffusion along the grain boundaries has also been studied by Feaugas et al. in Ni base alloys [START_REF] Oudriss | On the implication of hydrogen on inter-granular fracture[END_REF][START_REF] Oudriss | Grain size and grain boundary effects on diffusion and trapping of hydrogen in pure nickel[END_REF]; they showed a fast diffusion and a high solubility of H along high-angle random boundaries. These authors also showed that H was trapped at special Ʃ3 grain boundaries, which might reduce the susceptibility of the metals to H embrittlement due to the low H diffusion [START_REF] Oudriss | Meso-scale anisotropic hydrogen segregation near grainboundaries in polycrystalline nickel characterized by EBSD/SIMS[END_REF]. The results have been confirmed by Bechtle et al. [START_REF] Bechtle | Grain-boundary engineering markedly reduces susceptibility to intergranular hydrogen embrittlement in metallic materials[END_REF].

However, to the author knowledge, except for Lafouresse et al. [START_REF] Lafouresse | Hydrogen mapping in an aluminium alloy using an alternating current scanning electrochemical microscope (AC-SECM)[END_REF] and Schaller et al. [START_REF] Schaller | Spatially resolved mapping of the relative concentration of dissolved hydrogen using the scanning electrochemical microscope[END_REF][START_REF] Thomas | The effect of absorbed hydrogen on the dissolution of steel[END_REF], there are very few publications that provide data concerning the influence of H on the electrochemical reactivity of the materials. However, if H is introduced around a corrosion defect and if it can modify the electrochemical reactivity of the material, then it could have an influence on the corrosion defect growth and then on its morphology. Therefore, to provide a better understanding of the IGC mechanism and the propagation kinetics of the IGC defects in AA 2024, the present work aims to study the influence of H on the propagation of the IGC defects. H-precharged samples were studied after exposure to a corrosive environment (5 mM NaCl solution) and the results were compared with those obtained for non-charged samples. Firstly, optical and transmission electron microscopic (OM and TEM) observations were used to characterise both the morphology and chemical composition of the IGC defects accurately. Then, electron backscatter diffraction (EB SD) analyses were performed to identify the corroded interfaces accurately for both non-charged and H-precharged samples.

Experimental procedure

. Material

The material used for this study was a AA 2024 T351 provided by Airbus Group (France). The material was received as a 50 mm thick plate obtained by hot rolling; this step was followed by solution heat treatment at 495 °C ( ± 5 °C), water quenching, straining and tempering in ambient conditions for 4 days to achieve the final T351 metallurgical state. The chemical composition of the material was the following: Al base, 4.46% Cu, 1.44% Mg, 0.60% Mn and 0.13% Fe, wt. %. The microstructure of this alloy corresponded to grains highly elongated in the rolling direction (longitudinal, L) with the average sizes of 700, 300 and 100 μm in the longitudinal (L), long transverse (LT) and short transverse (ST) directions, respectively. observations are shown in the following. The results were compared to those described in previous work [START_REF] De Bonfils-Lahovary | Characterization of intergranular corrosion defects in a 2024 T351aluminium alloy[END_REF] for corroded non-charged samples. EDX analyses were helpful in identif}-ing the precipitates at the grain boundaries.

EBSD analysis

EB SD analyses were performed on three samples before H precharging and corrosion exposure using a JEOL JSM 7100F TTLS LV field-emission SEM and an OXFORD INIBUMENTS HKL EBSDsystem (Nordlys Nano EB SD detector -version 3.1 of the Aztec software). Samples were analysed at 20kV with a 12.5nA current. EBSD maps were performed on a 1. 5 mm X 3 mm area delimited by micro-hardness indents with a step size of 1.5 µmin the LT-ST plane. This selected area was large enough to obtain representative results. Overall, approximately 1000 grains were analysed for each sample. The commercial orientation imaging software package Oxford Channel 5 was then used for the post-processing of the EBSD data. Prior to data analysis, in order to minimise measurement errors, ail grains comprising less than 9 pixels were automatically removed from the maps. Further, a lower limit boundary misorientation eut-off o f 3 • w as u sed t o eliminate spurious boundaries caused by orientation noise. The analyses allowed the determination of the nature of the interfaces (Fig. l a) in relation with their level of misorientation in agreement with Priester who defines a grain boundary as the interface between two crystals, with the same crystallographic structure, with a greater or lesser degree of misorientation [START_REF] Priester | Les joints de grains: De la théorie à l'ingénierie[END_REF]. The word 'interfaces' refers, in this work, to the grain boundaries independent of their level of misorientation. Usually, a rotation angle of 15' was defined as the transition angle from LAGB, i.e. subgrain boundaries, and HAGB, i.e. grain boundaries. For pure aluminium, the critical angle for the transition LAGB /HAGB was between 8.6' and 14.4' [START_REF] Winning | Transition between low and high angle grain boundaries[END_REF]. In this study, the critical misorientation angle to discriminate the subgrain boundaries and the grain boundaries was chosen to be equal to 10'. For grain boundaries, a distinction was made between coïncidence site lattice grain boundaries (CSL), where the degree of fit ( E) b etween the t wo structures of t be t wo grains is described by the reciprocal of the ratio of coïncidence sites to the total number of sites oriented, and random grain boundaries where the structure was observed as being disorganised. After the corrosion tests (24 h of immersion in 5 mM NaCl aerated solution), the selected area was analysed using SEM with the backscattered electron (BSE) detector to easily identify any IGC defects (Fig. l b).The corrosion map obtained was superimposed on the previous EB SD map of the non-corroded material (Fig. ., t -

• 1 ' q + ' ' / 1
..

~~ t

non-<:harged samples and two corroded H-precharged samples; it permitted the determination of the misorientation level for each corroded interface (Fig. 1 c).

Results and discussion

I. Influence of Hon the IGC defect mo,phology

The morphology of the IGC defects was first investigated using OM after a 6 h exposure in a 5 mM NaCl aerated solution at room temperature (25 'C). Fig. 2a) and b) show OM micrographs of IGC defects formed on non-charged and H-precharged samples, respectively. For the non-charged sample (Fig. 2a), only a few IGC defects were detected (mean value = 6 defects/cm for a cross-section in the L-ST plane). Defects were thin and 86% of them were unbranched. For the H-precharged sample (Fig. 2b), IGC defects were more numerous (mean value = 28 defects/cm) and more branched: 26% of the IGC defects showed the morphology shown in Fig. 2b. The observations therefore highlighted an increase of the local reactivity, i.e. the anodic dissolution of the alloy, which was assumed to be related to the presence of H inside the sample. Further, IGC defects formed for a non-charged sample, after exposure to a more concentrated NaCl solution (1 M) and for longer duration (24 h) (insert in Fig. 2a) showed that IGC defects formed could become branched, but the level of ramifications observed was lower than for H-precharged samples. The evolution of the morphology of the IGC defects for the non-charged samples when the exposure to the NaCl solution was longer was assumed to be related to H diffusion all around the IGC defect, H being produced by the corrosion processes [START_REF] De Bonfils-Lahovary | Characterization and control of the intergranular corrosion defects in a 2024 T351 aluminium alloy[END_REF]. Therefore, the comparison of the IGC defect morphology for non-charged and H-precharged samples led us to assume a significant influence of H on the IGC defect propagation. Furthermore, the evolution of the IGC defect morphology with exposure time to NaCl solution seemed to confirm this hypothesis.

To go further in the characterisation of the IGC defects, conventional lift out procedure using a FEI HELIOS Nanolab 600i dual beam FIB / scanning electron microscope (SEM) was then used to obtain a thin sample in a specific region, i.e. near an IGC defect. Fig. 3 shows representative bright field (BF) TEM images showing a ramification of an IGC defect formed for an H-precharged sample of M 2024 T351 after a 24 h immersion in a 5 mM NaCl solution. A continuous precipitation was observed along the two walls of the IGC defects in the upper ramification. The precipitates were identified as Al-Cu-Mn (71 % Al -18% Cu -11 % Mn, at. %) and Al-Cu-Mg (83% Al -15% Cu -2% Mg, at. %) by EDX analyses. The morphology of Al-Cu-Mg precipitates suggested that they were partially dissolved. This was in agreement [START_REF] Garcia-Vergara | Morphology of enriched alloy layers in an anodized Al-Cu alloy[END_REF]. Overall, the morphology at the TEM scale of the IGC defects formed for the H-precharged sample was similar to that observed for a non-charged sample [START_REF] De Bonfils-Lahovary | Characterization of intergranular corrosion defects in a 2024 T351aluminium alloy[END_REF]. Therefore, the results suggested that H did not impact the IGC elementary mechanisms. However, it probably modified the corrosion kinetics, by increasing the corrosion susceptibility of grain boundaries and subgrain boundaries. This was consistent with OM observations (Fig. 2) that showed the increase of the IGC defect density and their ramification for H-precharged samples. To provide experimental proof about this hypothesis, EBSD analyses were therefore performed on H-precharged samples by comparison to non-charged samples.

H influence on the susceptibility to corrosion of the interfaces depending on their level of misorientation

The distribution of the levels of misorientation of the interfaces for non-corroded AA 2024 alloy was studied for three samples: similar results were obtained with the three samples. Therefore, for more clarity, Fig. 4a represents the results for one sample only. In the selected zone (1.5 × 3 mm²), more than 1000 grains were present, which was a good indication of the number of analysed interfaces. Subgrain boundaries constituted the majority of interfaces, i.e. 59% of the studied interfaces, and only 41% were grain boundaries (Fig. 4a). The dark line in Fig. 4a represents the distribution of the boundary misorientations of a set of grains spread randomly and having a cubic symmetry [START_REF] Humphreys | Recrystallization and Related Annealing Phenomena[END_REF]. Some similarities between the experimental grain boundaries distribution and the theoretical one were observed for θ > 10°, but the theoretical distribution did not completely fit with the experimental results. This could be explained considering the shaping treatment during the plate manufacturing. The presence of subgrains could be linked to the recovery process during the rolling. Moreover, in order to characterise the nature of the interfaces accurately, the distribution of the coincidence site lattice (CSL) grain boundaries was also studied (Fig. 4b). Results showed that 93% of the grain boundaries were random grain boundaries and only 7% were CSL grain boundaries, with Ʃ3 interfaces representing the majority of the CSL (14%) (Fig. 4b). Therefore, in the following, for the CSL grain boundaries, only the Ʃ3 grain boundaries were considered.

X""cu-nanopartlcles

Then, as explained in the experimental section (Fig. 1), the levels of misorientation of the corroded interfaces were recorded to establish a relationship between the level of misorientation of an interface and its susceptibility to corrosion. For this analysis, the duration of the corrosion tests was 24 h in order to make enough IGC defects grow without introducing a too high H amount. For a non-charged sample, after the corrosion test, 22 corroded interfaces were identified on the corroded area (4.5 mm²); all of them were analysed. Clearly, the number of interfaces analysed for the corroded non-charged sample was significantly lower than for the non-corroded sample, but the results were assumed to be representative considering that the same analysis was performed for three samples and similar results were obtained (Table 1). Fig. 5 thus shows, as an example, the distribution of the levels of misorientation of these 22 corroded interfaces; in this figure, the distribution plotted for the non-corroded sample, with all the interfaces analysed in the selected zone of 4.5 mm 2 , is reported for comparison. Results showed that grain boundaries corresponded to 95% of the corroded interfaces whereas they represented only 41% of all the interfaces in the non-corroded material; further, their susceptibility to corrosion increased with their level of misorientation with a larger number of corroded interfaces identified for the highest levels of misorientation. In particular, it was calculated that 67% of the corroded interfaces had a misorientation angle higher than 40°whereas those interfaces constituted only half of the total interface population. Finally, none of the 22 corroded interfaces observed were Ʃ3 CSL grain boundaries. The results were in perfect agreement with Guerin's work performed for an AA 2050 alloy [START_REF] Guérin | Identification of the metallurgical parameters explaining the corrosion susceptibility in a 2050 aluminium alloy[END_REF] and confirmed the strong reactivity of the grain boundaries with the highest levels of misorientation, which may be linked to their disarranged structure and correlated with their high energy [START_REF] Humphreys | Recrystallization and Related Annealing Phenomena[END_REF]. For the Ʃ3 CSL grain boundaries, Kim et al. also showed that these interfaces had a good resistance to IGC [START_REF] Kim | Grain boundary character distribution and intergranular corrosion behavior in high purity aluminum[END_REF]. The

Table 1

Summary of the results obtained after the corrosion tests in a 5 mM NaCl solution. *For the corrosion tests in 3.1., results concerned IGC defects: when several corroded interfaces were connected together, they were considered as only one IGC defect. **For the corrosion tests in 3.2., each corroded interface was considered individually so that the concept of "ramification" (branched and unbranched) was not relevant in this part of the study. GB = grain boundaries and sGB = subgrain boundaries. boundaries are more susceptible to IGC than subgrain boundaries. However, in the presence of a large H amount, the subgrain boundaries susceptibility was increased and Ʃ3 CSL grain boundaries also became susceptible to corrosion.

New insight into the IGC mechanism for AA 2024

The results presented in the previous two paragraphs permitted to provide a new insight into the IGC mechanism for the AA 2024, highlighting the influence of H produced during the corrosion processes. Fig. 6a is a schematic representation of the AA 2024 T351 microstructure for a non-corroded sample. The most numerous interfaces are subgrain boundaries (59%). Fig. 6b shows the IGC initiation and propagation of short IGC defects for a corroded non-charged sample. The IGC defects initiate following the common mechanism described in the literature [START_REF] Galvele | Mechanism of intergranular corrosion of Al-Cu alloys[END_REF], where intergranular Cu-rich particles promote the corrosion processes due to a galvanic coupling with the adjacent matrix. The intergranular precipitates are mostly present at the grain boundaries, rather than at the subgrain boundaries, because the high energy of the grain boundaries promotes the precipitation phenomenon. As a consequence, the corrosion processes mostly concern the grain boundaries during the first time of exposure to the chloride media; therefore, the main part of the corroded interfaces are grain boundaries (Fig. 2a). However, the corrosion processes lead to the formation of a Cu-rich layer at the interface IGC defect/matrix, that promote the reduction of protons formed due to the cations hydrolysis. H diffused into the sample from the IGC defects where protons are produced as illustrated by SKPFM study [START_REF] Larignon | Investigation of Kelvin probe force microscopy efficiency for the detection of hydrogen ingress by cathodic charging in an aluminium alloy[END_REF], with the interfaces (both grain boundaries and subgrain boundaries) adjacent to the corroded interfaces acting as preferential diffusion paths. The H enrichment of those interfaces enhances their electrochemical reactivity (Fig. 6d), promoting their dissolution, as shown by AC-SECM measurements [START_REF] Lafouresse | Hydrogen mapping in an aluminium alloy using an alternating current scanning electrochemical microscope (AC-SECM)[END_REF] so that IGC defects can progress laterally and become more and more branched when the duration of exposure to the chloride solution increased (insert in Fig. 2a). Further, Larignon et al. [START_REF] Larignon | Investigation of Kelvin probe force microscopy efficiency for the detection of hydrogen ingress by cathodic charging in an aluminium alloy[END_REF] showed a decrease of the corrosion potential of AA 2024 in the presence of H; therefore, a galvanic coupling could occur between the H-enriched zone and the adjacent material where the H-enriched zone is the sacrificial anode. This could also explain the propagation of the IGC defects along the subgrain boundaries connected to the corroded grain boundaries. Overall, the H enrichment of the material during the corrosion processes provides an additional explanation to the evolution of the morphology of the IGC defects with increasing duration of exposure to the chloride solution. Clearly, when H was artificially inserted inside the material before the corrosion processes by cathodic charging, the corrosion processes were enhanced more rapidly as shown in Fig. 6c with a stronger ramification of the IGC defects (Fig. 2b).

Conclusions

The present work was focused on the influence of H on the intergranular corrosion (IGC) mechanisms of AA 2024 T351. The main results can be summarised as follows:

1 At the optical and scanning electron microscopic scale, the IGC damage was found to be more extended with a significant ramification of the IGC defects for corroded H-precharged samples. At the transmission electron microscopic scale, no obvious difference was observed between corroded H-precharged and corroded noncharged samples; in both cases, the formation of a Cu-rich layer was observed at the matrix-oxide interface. It was therefore assumed that H did not influence the elementary processes of IGC, but it had a significant impact on the local reactivity. 2 EBSD analyses showed, for a corroded non-charged sample, that the most corroded interfaces were those with the highest levels of misorientation (grain boundaries, θ > 40°). Furthermore, the low energy of the subgrain boundaries provided also an explanation to their higher resistance to corrosion [START_REF] Kim | Grain boundary character distribution and intergranular corrosion behavior in high purity aluminum[END_REF][START_REF] Guérin | Identification of the metallurgical parameters explaining the corrosion susceptibility in a 2050 aluminium alloy[END_REF][START_REF] Luo | Observations of intergranular corrosion in AA2024-T351: the influence of grain stored energy[END_REF][START_REF] Bennett | Effect of grain boundary structure on sensitization and corrosion of stainless steel[END_REF][START_REF] Ortner | A study of the relation between grain boundary type and sensitisation in a partially-sensitised AISI 304 stainless steel using electron backscattering patterns[END_REF]. Comparison of the results with those obtained for the H-precharged and corroded samples permitted to study the influence o f H o n the susceptibility to IGC of the AA 2024. In this case, results showed that the corrosion damage was significantly m ore e xtended i n agreement with OM observations (Fig. 2b). Therefore, the EBSD analyses was restricted to an area of 2.3 mm², i.e half-reduced area by comparison to the one studied for the corroded non-charged sample, in order to facilitate the identification of the corroded interfaces. However, even for this reduced area, the number of IGC defects was equal to 336, all being analysed in this work (Fig. 5). As for the corroded non-charged samples, the experiments were repeated, in this case twice, and similar results were obtained (Table 1). Therefore, for more clarity, only the results obtained for the 336 corroded interfaces from the same H-precharged corroded sample are presented here. For the H-precharged corroded sample, a total of 93% of the corroded interfaces were grain boundaries and 7% of them were subgrain boundaries. Therefore, the results showed that H enhanced the susceptibility to IGC of the AA 2024 with an increase by a factor of fifteen of the number of corroded interfaces. Further, in the presence of H, the most susceptible interfaces to IGC remained the grain boundaries. However, the percentage of subgrain boundaries among the corroded interfaces increased from 5 to 7 for corroded non-charged and corroded H-precharged samples, respectively. Clearly, the increase could be considered as low; however, considering the significant number of corroded interfaces for the corroded H-precharged sample, the results clearly showed the significant enhancement of the corrosion susceptibility for subgrain boundaries: 24 corroded subgrain boundaries for a 2.3 mm² analysed area for the corroded H-precharged sample to be compared to only 1 corroded subgrain boundary for a 4.5 mm² analysed zone for the corroded noncharged sample. Furthermore, as indicated before and as shown in Table 1, the results were reproducible. They are also consistent with results obtained after 6 h of immersion in the NaCl solution (Fig. 2 and Table 1). H was thus found to enhance the susceptibility to corrosion of all the interfaces in such an extent that interfaces, which did not corrode when H amount was low, became susceptible to corrosion for a larger H amount. The corrosion behaviour of Ʃ3 CSL grain boundaries confirmed the results since, for the corroded H-precharged sample, 3% of the 336 corroded interfaces were Ʃ3 CSL grain boundaries whereas, for the corroded non-charged sample, no Ʃ3 CSL grain boundaries were corroded. To further quantify the results, the mean level of misorientation of each distribution was calculated. For the non-corroded sample, the mean level of misorientation was 39°, considering here all the interfaces in the selected zone. As explained before, for the corroded samples, the distributions were restricted to the corroded interfaces. For the corroded non-charged sample, a shift of the distribution was observed towards the high degrees of misorientation compared to the noncorroded sample. The mean level of misorientation was equal to 41.5°, which led us to assume that interfaces with a high level of misorientation were more susceptible to IGC. As indicated previously, the result was obtained for a low number of corroded interfaces, but it was confirmed by the analysis of two other samples and was in agreement with literature data [START_REF] Davenport | Intergranular corrosion and stress corrosion cracking of sensitised AA5182[END_REF][START_REF] Guérin | Identification of the metallurgical parameters explaining the corrosion susceptibility in a 2050 aluminium alloy[END_REF]. B y comparison, the distribution of the corroded grain boundaries for the corroded H-precharged sample was more similar to that of the non-corroded alloy, with subgrain boundaries, but also grain boundaries with a low level of misorientation (10° < θ < 40°) that became more reactive leading to a decrease of the mean value of the level of misorientation to 37.5°. In that case, the result was obtained for a large number of corroded interfaces. As done by Guerin et al. [START_REF] Guérin | Identification of the metallurgical parameters explaining the corrosion susceptibility in a 2050 aluminium alloy[END_REF], an IGC susceptibility index (I SC ) was also calculated for the different types of interfaces (subgrain boundaries, random and Ʃ3 CSL grain boundaries). It was determined by the ratio of the proportion of corroded interfaces (% c ) divided by the proportion of the interfaces of the same type identified f or t he n on-corroded sample (% nc ).The results are summarised in Table 2 and confirmed that grain presence of H strongly increased the number of corroded interfaces and made subgrain boundaries (θ < 10°), grain boundaries with low misorientation angle (10°< θ < 40°) and Ʃ3 CSL grain boundaries susceptible to IGC. [START_REF] Wang | The intergranular corrosion susceptibility of 2024 Al alloy during re-ageing after solution treating and cold-rolling[END_REF] The results were analysed considering previous work that showed that H could both lower the corrosion potential of AA 2024 and increase its reactivity probably due to the disturbance of the crystallographic lattice. They allowed to propose a new insight into the IGC mechanism for AA 2024 where the corrosion-induced H that was trapped inside the IGC defects could then modify the local reactivity of the material. The embrittlement of new interfaces, in which H had diffused, was assumed to explain the ramification of the IGC defects when the duration of exposure to the chloride solution was increased.

In this study, corrosion tests were performed under isothermal conditions and results are thus immediately applicable for structural components used in these conditions. However, when dealing with aerospace industry and real flying conditions, some structural parts can experience thermal cycles with exposure to temperatures < 0 °C. In those conditions, the solidification of the electrolyte trapped inside the IGC defect can generate stress concentration at the IGC defect tip, which should lead to larger H enrichment of the corroded alloy. Therefore, the H-enhanced corrosion mechanism proposed in this study might be helpful in predicting the lifetime of those structural components.

Table 2

Calculation of the IGC susceptibility index (I SC ) for the different types of interfaces for AA 2024 samples exposed to a 5 mM NaCl solution for 24 h. Non-charged and H-precharged samples are corroded samples here.

  l a) in order to determine the nature of the corroded interfaces (Fig.le). This procedure was performed for three corroded b)

Fig. 1 .

 1 Fig. 1. Methodology used to study the influence of H on the corrosion behaviour of AA 2024 alloy: a) EBSD map of the surface exposed (1.5 X 3 mm 2 ) to the electrolyte before H precharging and corrosion tests; b) SEM image on BSE mode of the corroded surface (same zone) after immersion in a 5 mM NaCI solution during

Fig. 2 .Fig. 3 .

 23 Fig. 2. Optical micrograph of IGC defects formed after a 6 h exposure at E corr in a 5 mM NaCl solution for a) non-charged and b) H-precharged samples of AA 2024 T351. Insert in Fig. 2a shows an IGC defect after 24 h of immersion in a 1 M NaCl solution for a non-charged sample.

Fig. 4 .

 4 Fig. 4. AA 2024 T351: a) Distribution of the levels of misorientation of the interfaces before H precharging and corrosion tests (non-corroded sample): the black curve represents the distribution of the boundaries misorientation of a set of grains spread randomly with a cubic symmetry [44] b) Distribution of the CSL grain boundaries before H precharging and corrosion tests.

Fig. 5 .

 5 Fig. 5. Distribution of the levels of misorientation of the corroded interfaces for non-charged and H-precharged samples after a 24 h exposure at E corr in a 5 mM NaCl solution. The results obtained for the non-corroded alloy are reported for comparison (all interfaces present in the samples were analysed).

Fig. 6 .

 6 Fig. 6. IGC Mechanism for AA 2024 T351. a) Microstructure of the non-corroded material with identification of the interfaces. Identification of the corroded interfaces for b) non-charged and c) H-precharged samples after immersion in a NaCl solution. d) Scheme of the H influence on the local reactivity.

Acknowledgement

This work was supported by ANR-14-CE07-0027-01 (France) -M-SCOT: Multi Scale COrrosion Testing.

Data availability

The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study.