Clement Rochet

Eric Andrieu

Babak Arfaei

Jean-Paul Harouard

Adrien Laurino

Terry C Lowe

Grégory Odemer

Christine Blanc
email: christine.blanc@ensiacet.fr

\ E Andrieu

J.-P Harouard\

Influence of equal-channel angular pressing on the corrosion fatigue behaviour of an Al-Mg-Si aluminium alloy for automotive conductors

Keywords:

des établissements d'enseignement et de recherche français ou étrangers, des laboratoires publics ou privés.

Introduction

Electrical conductors are found in most of technological systems such as high voltage lines, cars and aircrafts. In the automobiles, the electrical distribution system can reach almost 5 km in length, which represents a significant part of the total vehicle weight. Presently, Cu and Cu alloys comprise most wiring because of their high electrical conductivity. However, lighter and cheaper alternatives exist. Indeed, in high voltage lines for example, Al-Mg-Si alloys are increasingly used instead of Cu alloys, because of their low weight, relatively high con ductivity (50--60% International Annealed Cu Standard), good me chanical properties and high corrosion resistance. Therefore, car man ufacturers are currently evaluating the replacement of Cu wires by Al based wires in order to reduce the weight and cost of next-generation vehicles. Automotive wiring must resist degradation due to corrosion and fatigue embrittlement, to which Al alloys are known to be sus ceptible [START_REF] Kairy | Exploring the electrochemistry of 6.xxx series aluminium alloys as a function of Si to Mg ratio Cu content, ageing conditions and microstructure[END_REF][START_REF] Jiang | Influence of microstructure on defonnation behavior and fracture mode of Al-Mg-Si alloys[END_REF]. In the case ofa combination of vibrations and corrosive fluids, Al wires may experience corrosion fatigue embrittlement which is known to be very critical. Furthennore, car manufacturers are looking into new ways to improve the electrical conductivity and the mechanical properties of Al wires. Severe plastic defonnation processes (SPD) such as equal-channel angular pressing (ECAP) have already proven to be effective in this malter [START_REF] Murashkin | Enhanced mechanical properties and electrical conductivity in ultra:fine-grained Al 6101 alloy processed via ECAP-confonn[END_REF]. Therefore, ECAP could be of great interest for Al wires; however, before replacing Cu wires by 6xxx Al wires in automotive vehicles, it is necessary to measure the influence of ECAP on the corrosion fatigue resistance of Al-Mg-Si alloys.

Corrosion fatigue embrittlement includes both corrosion processes and a cyclic mechanical stress. The influence of SPD on the fatigue behaviour of Al alloys has already been studied in the literature and it is generally known that the fine grains obtained with SPD are beneficial in tenns of fatigue life [START_REF] Chung | Improvement of high-cycle fatigue life in a 6061 Al alloy produced by equal channel angular pressing[END_REF]. However, this general conclusion must be qualified taking into account the numerous cyclic plasticity mechan isms that are likely to occur in relation to the alloy metallurgical states, the SPD parameters, and the operating environment. For instance, Murashkin et al. studied the fatigue behaviour of a 6061 aluminium alloy (M6061) processed by high-pressure torsion (HPT) [START_REF] Murashkin | Fatigue behavior of an ultrafine-grained Al-Mg-Si alloy processed by high-pressure torsion[END_REF]. They observed that the endurance limit was improved by a factor of two due to the formation of a very homogeneous ultrafine-grained (UFG) mi crostructure with homogeneous resistance to fatigue crack initiation in the high cycle fatigue (HCF, > 10 5 cycles) regime. In the low cycle fatigue (LCF, < 10 5 cycles) regime, the UFG alloy showed somewhat lower fatigue resistance due to its lower strain hardening ability. The corresponding fracture surfaces showed ductile striations in the stage of stable crack propagation in the coarse-grained (CG) T6-treated alloy, whereas brittle striations seemed to dominate on the fatigue fracture surface of the UFG allo y. The final fracture presented dimples in both https:/ /doi.org/10.1016/j.ijfatigue.2020.105812 material conditions, but with a smaller size for the UFG alloy. Similarly, Chung et al. studied the HCF life of an AA606 l processed by ECAP [START_REF] Chung | Improvement of high-cycle fatigue life in a 6061 Al alloy produced by equal channel angular pressing[END_REF]. They observed a large enhancement in fatigue lifetime by a factor of about 10 after only one single pass of ECAP as compared to con ventionally processed AA6061-T6, and for both low and high cycle regimes contrary to Murashkin et al. [START_REF] Murashkin | Fatigue behavior of an ultrafine-grained Al-Mg-Si alloy processed by high-pressure torsion[END_REF]. Similar conclusions about the improvement of fatigue lifetime in the two regimes after ECAP were obtained by Esmaeili et al. for an AA7075 processed by ECAP [START_REF] Esmaeili | Fatigue behavior of AA7075 aluminium alloy severely defonned by equal channel angular pressing[END_REF].

Concerning the corrosion behaviour of conventional Al-Mg-Si al loys, literature is abundant, and published works have shown that the corrosion resistance in chloride media of these alloys is strongly related to their microstructure and mainly controlled by the precipitation state [START_REF] Blanc | Susceptibility to pitting corrosion of 6056 aluminium alloy[END_REF][START_REF] Andrieu | Corrosion behavior of 6101 aluminum alloy strands for automotive wires[END_REF][START_REF] Bhattamishra | Microstructural srudies on the effect of Si and Cr on the intergranular corrosion in Al-Mg-Si alloys[END_REF][START_REF] Larsen | Effect of excess silicon and small copper content on intergranular corrosion of 6000-series aluminum alloys[END_REF][START_REF] Wang | Effect of artificial aging on intergranular corrosion of 6063 Al alloy[END_REF][START_REF] Zou | The intergranular corrosion behavior of 6000-series alloys with different Mg/Si and Cu content[END_REF]. In this sense, many authors showed a significant influence of ECAP processes on the corrosion behaviour, in particular, pitting cor rosion of these alloys related to the fragmentation of coarse precipitates [START_REF] El | Influence of ECAP as grain refinement technique on microstructure evolution, mechanical properties and corrosion behavior of pure aluminum[END_REF] and to the decrease in grain size [START_REF] Korchef | Corrosion behavior of commercial aluminum alloy processed by equal channel angular pressing[END_REF].

On the contrary, very few works focused on the influence of SPD on the corrosion fatigue behaviour of Al-Mg-Si alloys, where multiple works about conventionally processed alloys have shown a decrease in fatigue lifetime due to premature crack initiation on corrosion defects [START_REF] Laurino | Effect of corrosion on the fatigue life and fracture mechanisms of 6101 aluminum alloy wires for car manufacturing applications[END_REF][START_REF] Weber | The effect of surface corrosion da mage on the fatigue life of 6061-T6 aluminum alloy extrusions[END_REF][START_REF] Chanyathunyaroj | Fatigue characteristics of 6061 aluminum alloy subject to 3.5% NaCl environment[END_REF][START_REF] Almaraz | Fatigue endurance and crack propagation under rotating bending fatigue tests on aluminum alloy AISI 6063-T5 with controlled corrosion attack[END_REF][START_REF] Lee | The effect of pitting corrosion on fatigue life[END_REF][START_REF] Corsetti | The effect of mean stress and environment on corrosion fatigue behavior of 7075-T6 aluminum[END_REF]. The plastic deformation localisation around the corrosion defects and hydrogen embrittlement at defect/crack tip are the common mechanisms mentioned to explain this decrease in fatigue lifetime under corrosion fatigue conditions [START_REF] Magnin | The corrosion-enhanced plastieity model for stress corrosion cracking in ductile fcc alloys[END_REF][START_REF] Delafosse | Dislocation-hydrogen interac tions during stress corrosion cracking in fcc metals: experiments on single crystals and numerical simulations[END_REF]. Sorne authors proposed a synergistic effect of corrosion and cyclic stress. Indeed, cyclic stress might enable an enhancement of oxide film embrittlement [START_REF] Mueller | Dependence of corrosion fatigue crack initiation mechanisms on the corrosion behavior oftwo stainless chromium steels[END_REF], the dissolution of the matrix surrounding metallurgical defects [START_REF] Magnin | Mechanisms of corrosion-fatigue of metallic alloys[END_REF] and hydrogen penetration and trapping at the crack tip [START_REF] Corsetti | The effect of mean stress and environment on corrosion fatigue behavior of 7075-T6 aluminum[END_REF][START_REF] Burns | Driving forces for localized corrosion-to-fatigue crack transition in Al-Zn-Mg-Cu[END_REF][START_REF] Pressouyre | Trap theory of Hydrogen embrittlement[END_REF]. As stated above, the influence of SPD, and in particular ECAP, on such mechanisms is still not well documented at this time. Laurino et al. studied the corrosion fatigue behaviour and fracture mechanisms of an AA6101-T9 obtained by a two-step cold drawing process [START_REF] Laurino | Effect of corrosion on the fatigue life and fracture mechanisms of 6101 aluminum alloy wires for car manufacturing applications[END_REF]. In the T9 metallurgical state, very elongated grains were observed in the wire drawing direction with a 500-nm width in the transverse plane. For this metallurgical state, no difference in fatigue lifetime was observed for corrosion fatigue tests as compared to fatigue tests performed on non-corroded and pre-corroded samples. In addi tion, similar to fatigue tests perforrned for pre-corroded samples, the crack initiation was localised on corrosion defects and a mainly inter granular fracture mode was observed in the propagation zone for samples tested under corrosion fatigue conditions. Sharma et al. studied the corrosion fatigue behaviour of UFG AA5083 obtained by a combi nation of cryomilling and hot isostatic pressing [START_REF] Shanna | A microscopie study on the corrosion fatigue of ultra-fine grained and conventional Al-Mg alloy[END_REF]. The UFG alloy presented a higher fatigue limit than the conventional alloy in air. In a marine environment, the UFG alloy was found to be more susceptible to corrosion fatigue embrittlement. The experimental results suggested that at high stress ranges, slip-enhanced dissolution and hydrogen embrittlement were the dominant mechanisms for fracture in the UFG alloy. Conversely, at lower applied maximum stress values, pitting and crack closure were the prevailing mechanisms that led to crack initia tion and ultimate fracture.

In this framework, the present paper deals with the influence of ECAP on the corrosion fatigue behaviour of an Al-Mg-Si alloy. Fatigue tests were first perforrned in air for samples before and after ECAP, on both non-corroded and pre-corroded samples; the results obtained with non-corroded samples were considered as reference data, whereas those on pre-corroded samples allowed the effect of corrosion defects on the fatigue behaviour to be evaluated. Both sets of results were useful in analysing the data obtained after corrosion fatigue tests. To complete the results, fracture surfaces were analysed by scanning electron mi croscopy to characterise fatigue and corrosion fatigue fracture modes. Comparison of ail the results was helpful to distinguish between the influence of corrosion and cyclic mechanical loading on the corrosion fatigue behaviour of the alloy, and to deterrnine whether a synergy occurred between these two parameters, i.e. corrosion and mechanical loading. Furthermore, the influence of ECAP on these mechanisms was analysed.

Material and experimental techniques

Material

The study focused on an Al-Mg-Si alloy with the following chemical composition (wt. %, Al balance): Mg 0.57, Si 0.37, Fe 0.10, Mn 0.04, Cu 0.01, others 0.10. The as-received condition, considered as the re ference metallurgical state, was provided as extruded 9.5 mm rods. As received samples were ECAP-processed under two different conditions, referred to as ECAP-RR and ECAP-RH.

ECAP-RR samples were submitted to two passes in an ECAP

Conform machine (120 ° die angle) [START_REF] Xu | Prineiples of ECAP-Conform as a continuons process for achieving grain refinement: Application to an aluminum alloy[END_REF] at room temperature with a rotation of 90 ° between passes [START_REF] Furukawa | The shearing char acteristics assoeiated with equal-channel angular pressing[END_REF]. For both passes, the displace ment rate was 2.04 mm s-1 to impart a total effective strain e = 1.34. 2. For ECAP-RH samples, the same machine was used, but the first pass was perforrned at room temperature and the second one at 135-155 °C. There was also a rotation of 90 ° between passes, and for both passes, the displacement rate was 2.04 mm s-1 to impart a total effective strain e = 1.34.

Experimental techniques 2.2.1. Mechanical tests

To determine the influence of ECAP on the corrosion fatigue be haviour of the alloy, preliminary tensile tests were performed using an MTS testing machine equipped with a 30 kN Joad cell, to evaluate the yield strength (YS) and the ultimate tensile strength (UTS) of both the as-received and ECAP-processed samples. Cylindrical tensile samples, with a gauge length of 24 mm and a 4-mm diameter, were used. Tests were all performed at room temperature under a constant strain rate of 10-3 s-1 using a 20-mm extensometer to record sample deforrnation during testing. In order to assess the reproducibility of the mechanical response, at least three tensile samples were tested for each condition.

Then, fatigue tests were perforrned in laboratory air on both the as received and ECAP-processed samples to obtain reference data. Fatigue tests were perforrned for non-corroded samples and pre-corroded samples, so that it was then possible, by comparison to the results ob tained during corrosion fatigue tests, to evaluate the synergy between the corrosion mechanisms and the mechanical loading during corrosion fatigue tests. Tests were carried out by using an Electroforce 3300 machine from TA instruments equipped with a 3.3 kN Joad cell. They were ail perforrned at a frequency of 50 Hz with a Joad ratio R = 0 """ = 0.1 under a sinusoïdal stress loading. Those testing condi tions°';,�re chosen in agreement with standard vibration tests performed at LEONI Wiring Systems. Cylindrical samples with a gauge length of 28 mm and a minimum diameter of 2.5 mm were used (Fig. 1). Prior to fatigue tests, samples were ground using SiC abrasive paper to P4000 (per the ISO 6344 standard designation of abrasives) and then polished using diamond paste to 0.25 µm. For the fatigue tests on pre-corroded samples, Silicone (CAFl) was used to mask the cylindrical sample, ex cept for 1 cm of the gauge length, which was exposed to the corrosive medium after polishing. Two different pre-corrosion tests were con sidered:

PC72h: immersion during

Characterisation of the microstructure and analysis of the fracture surfaces

The microstructure of the as-received, ECAP-RR and ECAP-RH samples was described in detail in a previous paper [START_REF] Rochet | Influence of equal-channel angular pressing on the microsnucture and corrosion behaviour of a 6:xxx aluminium alloy for automotive conductors[END_REF]. Only major results are reproduced here. First, electron backscatter diffraction (EBSD) analysis showed a significant decrease in area-weighted average grain size from 24 µm for the as-received sample to 16 µm and 4 µm for ECAP-RR and ECAP-RH samples, respectively. This was associated with an increase in high-angle grain boundaries (HAGBs) density and a de crease in the global texture of the samples after ECAP. Additionally, an image analysis based on scanning electron microscope (SEM) observa tions enabled a full characterisation of the precipitation state of the alloy before and after ECAP. SEM observations combined with an en ergy dispersive spectroscopy (EDS) analysis highlighted the presence of numerous Fe-rich coarse intermetallic (!MC) precipitates inside the matrix. The IMCs fraction of surface area was found to decrease from 0.53% for the as-received samples to 0.32% and 0.34% for ECAP-RR and ECAP-RH samples, respectively. This was most likely related to fragmentation of the IMCs during ECAP, which resulted in a higher amount of small IMCs not detectable during image analysis [START_REF] Rochet | Influence of equal-channel angular pressing on the microsnucture and corrosion behaviour of a 6:xxx aluminium alloy for automotive conductors[END_REF]. This particular effect was also observed by other authors [START_REF] Korchef | Corrosion behavior of commercial aluminum alloy processed by equal channel angular pressing[END_REF][START_REF] Xu | Effect of hot extrusion and subsequent T6 treatment on the microsttucture evolution and tensile properties of an Al-6Si-2Cu-0.5Mg alloy[END_REF]. No ]3", j3' and j3-Mg 2 Si precipitates formed during ECAP processes.

A Quanta 450 SEM manufactured by FEI was also used to observe fracture surfaces and characterise crack initiation sites on fatigue samples.

Characterisation of corroswn defects

The corrosion defects morphology was analysed using a Sensofar S Neox-confocal microscope set in interferometry mode (vertical resolu tion of 0.1-1 nm). The depth, diameter and fraction of surface area of corrosion defects after pre-corrosion and corrosion fatigue tests were determined using Gwyddion Analysis software. To obtain re presentative values of the corrosion defect morphology, 25 images per sample were analysed. Only corrosion defects deeper than 0.1 µm were considered, which allowed an average of 100 defects per image (351 µm x 264 µm) to be characterised. Another point of interest of these tests was the decrease in U:: ratio after both ECAP-RR and ECAP-RH processes (Table 1) essentially due to a significant increase in the yield stress. Such a result suggested a de crease in the plastic zone radius al the tip of a propagating crack so that the plastically deformed volume might be smaller than the re presentative elementary volume (REV) of the material and finally very sensitive to size and distribution of heterogeneities in terms of fatigue crack propagation rate. This point is important to consider in analysing the influence of ECAP on the fatigue behaviour of the alloy. As a .. Fig. 3 thus shows S-N curves obtained for the as-received, ECAP-RR and ECAP-RH samples at 93%, 90% and 80% of the UTS. The fatigue limit (fixed arbitrarily at 10 7 cycles in this work) was reached for the as received, ECAP-RR and ECAP-RH samples as soon as the applied max imum stress was lower than 90% of the UTS. This clearly showed that it was necessary to reach a very high stress level ( > 90% of the UTS and > YS 0 _ 2 ) to induce a fracture whatever the samples. At 93% of the UTS, the numbers of cycles to failure were very close for ail samples. However, this stress level (expressed as a percentage of UTS) corre sponded to a higher effective stress level (expressed in MPa) for ECAP RH sample, and particularly for ECAP-RR sample. This was related to the enhancement of mechanical properties already observed on tensile curves, and was explained by the grain refinement induced by the SPD process. Indeed, the increase in HAGBs density could create efficient barriers to crack propagation. This should be responsible for a decrease in the cumulated plastic deformation during the first cycles, and should lead to a delay in the crack initiation for ECAP samples. In addition, during fatigue tests, this should induce a decrease in the dislocation Table 1 Tensile properties of as-received, ECAP-RR and ECAP-RH samples. Therefore, it was concluded that due to a significant grain refine ment during ECAP, both ECAP-RR and ECAP-RH samples showed a better fatigue resistance than the as-received sample. Then, for a better understanding of corrosion fatigue behaviour of ECAP-processed sam ples, the effects of corrosion and cyclic mechanical loading were con sidered separately in a first approach. For that reason, in the following part, the influence of pre-corrosion on fatigue lifetime in air was ex amined. The corresponding fracture surfaces are reported in Fig. 7. Contrary to non-corroded samples, crack initiation sites were systematically re lated to corrosion surface defects (Fig. 7g, h andi). SEM observations of the samples revealed that those defects were homogeneously

Results and discussion

Effect of pre-corrosion on the fatigue lifetime in air

UTS (MPa) 218 ± 2 253 ± 1 238 ± 5

A%(%)

13.9 ± 0.6 11.4 ± 0.7 13.2 ± 0.9 2). Moreover, the density of corrosion defects (the corroded surface ratio) was found to increase from 1.71 10 3 mm-2 (0.91%) for the as-received sample to 5.36 10 3 mm-2 (1.38%) and 3.52 10 3 mm-2 (1.44%) for ECAP-RR and ECAP-RH samples, respectively .

As already stated, the decrease in fatigue lifetime for pre-corroded samples could be related to corrosion defects acting as stress con centrator sites [START_REF] Pidaparti | Analysis of pits induced stresses due to metal corrosion[END_REF][START_REF] Chen | Transition from pitting to fatigue crack growth-modeling of corrosion fatigue crack nucleation in a 2024-T3 alu minum alloy[END_REF][START_REF] Chlistovsky | Corrosion-fatigue behaviour of 7075-T651 aluminum alloy subjected to periodic overloads[END_REF]. The decrease in N p cnh/N Aic ratio for ECAP-RR and ECAP-RH samples, as compared to the as-received sample, sug gested that ECAP induced a higher susceptibility to the pre-corrosion step. Results shown above (Table 2) with less deep and fewer large defects for ECAP-processed samples did not seem consistent with this hypothesis. Indeed, it is generally expected that deeper and larger de fects are responsible for a higher contribution to stress concentration phenomenon [START_REF] Ambat | Effect of iron-containing inter metallic particles on the corrosion behaviour of aluminium[END_REF]. Many studies, focused on the effect of pitting cor rosion on fatigue crack initiation, indicated that pit depth and diameter were critical parameters [START_REF] Guérin | Corrosion-fatigue lifetime of AlAl--Cu-Lithium alloy 2050 in chloride solution[END_REF][START_REF] Li | Effect ofprior corrosion state on the fatigue small cracking behaviour of 6151-T6 aluminum alloy[END_REF][START_REF] Rokhlin | Effect of pitting corrosion on fatigue crack initiation and fatigue life[END_REF][START_REF] Co | Effects of macro-scale corrosion damage feature on fatigue crack initiation and fatigue behavior[END_REF][START_REF] Van Der Walde | Initiation and shape development of corrosion nucleated fatigue cracking[END_REF]. Rokhlin et aL developed an empirical relation predicting fatigue life of M2024-T3 after pre-corrosion, and showed that the deeper the corrosion pits, the lower the fatigue life of the sample [START_REF] Rokhlin | Effect of pitting corrosion on fatigue crack initiation and fatigue life[END_REF]. In the present study, from this consideration, the as received sample should present the highest susceptibility to the pre corrosion step. However, Burns et aL pointed out that pit depth and width were not the only parameters controlling fatigue crack initiation [START_REF] Burns | Driving forces for localized corrosion-to-fatigue crack transition in Al-Zn-Mg-Cu[END_REF]. Defect shape and defect surface density are also critical para meters in this regard. In agreement with this last statement, the slight increase in both density of corrosion defects and corroded surface ratio for ECAP-RR and ECAP-RH samples could explain the decrease in fa tigue lifetime for those samples. This should be associated with the fragmentation of Fe-rich IMCs during ECAP processes [START_REF] Korchef | Corrosion behavior of commercial aluminum alloy processed by equal channel angular pressing[END_REF][START_REF] Rochet | Influence of equal-channel angular pressing on the microsnucture and corrosion behaviour of a 6:xxx aluminium alloy for automotive conductors[END_REF], those particles acting as preferential corrosion sites [START_REF] Shaeri | Effect of ECAP tem perature on microsttucture and mechanical properties of Al-Zn-Mg-Cu alloy[END_REF]. The presence of a higher amount of cathodic sites led to the formation of a higher number of corrosion defects at the sample surface and to premature crack in itiation in the case of ECAP samples. A similar result was already ob served by Korchef et aL [START_REF] Korchef | Corrosion behavior of commercial aluminum alloy processed by equal channel angular pressing[END_REF].

Therefore, the results clearly demonstrated that the corrosion defect density was the most significant parameter to explain the deleterious effect of ECAP on the crack initiation step and finally on fatigue life time, for the pre-corrosion conditions used. However, during in-service use, corrosion mechanisms and cyclic mechanical loading affect the fatigue behaviour of the wire concurrently. Therefore, the influence of ECAP on the corrosion fatigue behaviour of the Al-Mg-Si alloy had to be examined. Fracture surfaces observed after corrosion fatigue tests are reported in Fig. 1 O. The fracture modes were identical to those observed for non corroded and PC72h pre-corroded samples. As for PC72h pre-corroded samples, crack initiation was always localised on corrosion defects (Fig. 10g, h andi), i.e. dissolution areas surrounding Fe-rich IMCs. No pits were observed for corrosion fatigue samples, which was consistent with the fact that these tests were performed at E,orr• Previous results, showing a similar number of cycles to failure for both corrosion fatigue tests and fatigue tests on PC72h pre-corroded samples, but corresponding to shorter immersion times for corrosion fatigue tests, suggested a synergy between cyclic mechanical loading and corrosion processes. This had already been observed in previous studies on Al alloys [START_REF] Laurino | Effect of corrosion on the fatigue life and fracture mechanisms of 6101 aluminum alloy wires for car manufacturing applications[END_REF][START_REF] Corsetti | The effect of mean stress and environment on corrosion fatigue behavior of 7075-T6 aluminum[END_REF][START_REF] Guérin | Corrosion-fatigue lifetime of AlAl--Cu-Lithium alloy 2050 in chloride solution[END_REF]. To confirrn this assertion, fatigue tests were perforrned at 90% of the UTS on samples pre-corroded during 1 h 30 min, i.e. approximatively the corrosion fatigue test duration for as received, ECAP-RR and ECAP-RH samples at 90% of the UTS. Results showed that, for this pre-corrosion duration, ail samples reached a number of cycles to failure significantly higher than for corrosion fa tigue tests (Fig. 11). This demonstrated that, for the same immersion time in 0.5 M NaCI (at E 00 ,,), the interaction between the cyclic me chanical loading and the corrosion processes led to a stronger decrease in fatigue lifetime of the alloy: a synergy effect between corrosion processes and the cyclic mechanical loading was clearly showed. As previously said, this synergy effect seemed to be more efficient for ECAP-processed samples.

Influence of ECAP on the corrosion fatigue behaviour

To explain this result, the corrosion defects formed during the cor rosion fatigue tests and during the PC1h30 pre-corrosion tests were characterised by using interferometric microscopy. The results are shown in Fig. 12. In order to have a statistical approach, results were presented using the boxplot representation. Each box encloses 50% of the data set with the median value displayed as a straight line. The top and the bottom of each box mark the limits of ± 25% of the data set (upper and lower quartiles respectively). The whiskers for each box refer to the minimum and maximum values from the data set that fa!! within an acceptable range. Any value outside of this range is re presented by a circle. First, a decrease in both depth and diameter of corrosion defects when the immersion time decreased was observed for ail samples when results obtained after a pre-corrosion of 72 h were compared to those obtained after a pre-corrosion of 1 h 30 min (Fig. 12a andb). Moreover, no significant difference in depth and diameter was observed between corrosion defects formed during corrosion fatigue tests and those forrned during PClh30 tests. This observation showed that a cyclic mechanical loading influenced neither the type nor the morphology of corrosion defects since previous observations of the fracture surfaces showed that, for corrosion fatigue tests, the corrosion defects corresponded to the dissolution of the matrix around IMCs as for PC72h and PC1h30 samples. In addition, no significant effect of ECAP on the depth of the corrosion defects forrned during corrosion fatigue tests was observed; a slight increase in the diameter of the corrosion defects for ECAP-processed samples could be noticed by comparison to the as-received samples. This was also true for samples pre-corroded during 1 h 30 min. On the contrary, as previously noted, significantly shallower corrosion defects were observed on ECAP samples pre-cor roded during 72 h as compared to the as-received samples. Ali these results were in perfect agreement with the conclusions drawn in a previous paper [START_REF] Rochet | Influence of equal-channel angular pressing on the microsnucture and corrosion behaviour of a 6:xxx aluminium alloy for automotive conductors[END_REF] stating that the influence of ECAP on corrosion defects morphology was only observed when defects reached critical dimensions. However, the main result here was that an increase in both the density and the corroded surface ratio was observed for corrosion fatigue samples by comparison to PClh30 samples (Fig. 12c andd). Comparison of Fig. 12b, c andd promoted by the electrochemical coupling effect existing between the tip and the edges of the crack that act as cathodic sites [START_REF] Jones | A unified mechanism of stress corrosion and corrosion fatigue cracking[END_REF]. li could be assumed that the differences in microstructure shown in our previous paper [START_REF] Rochet | Influence of equal-channel angular pressing on the microsnucture and corrosion behaviour of a 6:xxx aluminium alloy for automotive conductors[END_REF] between the as-received sample and ECAP-processed sam ples, but also between ECAP-RR and ECAP-RH samples, should con tribute to explain some differences related to the crack propagation rate. Those differences of microstructure should also explain differences in the susceptibility to hydrogen embrittlement. Indeed, another me chanism which may be cited here is based on hydrogen-assisted crack propagation processes during corrosion fatigue tests or during corrosion defect propagation un der cyclic thermal exposures [ 46, 4 7]. Preliminary tests on hydrogen embrittlement (HE) susceptibility of the Al alloy considered in this study suggested that HE was not a predominant da mage mechanism. However, it is necessary to keep in mind that the process zone al a fatigue crack tip is very small and that an effect of hydrogen in this area is still possible but very difficult to detect.

Conclusions

The influence of ECAP processes on the resistance to corrosion fa tigue of an Al-Mg-Si alloy was investigated. The main conclusions, in relation to previous studies performed on the same alloy [START_REF] Rochet | Influence of equal-channel angular pressing on the microsnucture and corrosion behaviour of a 6:xxx aluminium alloy for automotive conductors[END_REF]

Declaration of Competing Interest

The authors declare that they have no known competing financial interests or persona! relationships that could have appeared to influ ence the work reported in this paper.

72 h in 0. 5 MFig. 1 .

 51 Fig. 1. Fatigue sample geometry (NF EN ISO 11782-1). the same device as used for the corrosion fatigue tests. Corrosion fatigue tests were performed on both as-received and ECAP-processed samples by using a corrosion cell directly mounted on the Electroforce fatigue machine. The electrolyte was a 0.5 M NaCI solution, at a controlled temperature of 25 •c, as for the pre-corrosion conditions described before. It was maintained at a constant flow rate by means of a peristaltic pump. Corrosion fatigue tests were ail per formed at E co ,,. For reproducibility, fatigue tests and corrosion fatigue tests were repeated at least three times.

3. 1 .

 1 Influence of ECAP on the fatigue behaviour of non-corroded samples To determine the influence of ECAP processes on the corrosion fa tigue resistance of the Al-Mg-Si alloy, fatigue tests were first carried out in laboratory air on non-corroded samples, to be used as reference tests. It is of major importance to note that laboratory air is not a well-con trolled or inert environment; it can be responsible for embrittlement of materials during fatigue due to the adsorption and penetration of hy drogen present in water vapor. Sorne models were proposed to describe the influence of a gaseous environment such as moist air on crack propagation. Lynch's mode! of Adsorption Induced Dislocation Emission (AIDE) shows that the adsorption of aggressive species such as hydrogen at a crack tip could promote dislocations emission by a de crease in surface energy [32]. Therefore, such species mainly influence crack propagation rather than crack initiation during fatigue tests. Since crack initiation is the main mechanism influencing the fatigue lifetime in the present work, the effect of hydrogen from air moisture during endurance tests carried out in laboratory air was neglected. However, for fatigue tests performed on pre-corroded samples and corrosion fatigue tests, hydrogen could be produced by corrosion me chanisms around corrosion defects and then al the crack tip [15]; in that sense, an H effect was likely to be present. Fatigue tests were performed for different stress levels. Considering the differences in mechanical properties between ail samples, the stress levels for a given sample were chosen to correspond to varions per centages of the UTS measured for this sample. Therefore, in order to determine YS and UTS values of the as-received, ECAP-RR and ECAP RH samples, preliminary !ensile tests were performed al 25 •c and 10-3 s-1 in air. Fig. 2 shows the !ensile curves obtained for the as-received, ECAP RR and ECAP-RH sampi es at 10 -3 s-1 • An increase in both UTS and YS 0 .2 was observed as well as a decrease in the elongation to failure, for both ECAP samples as compared to the as-received sample. This evo lution was mostly related to the significant grain refinement induced by ECAP. Results also showed that the mechanical properties of ECAP-RH sample, with the smallest grain size, i.e. 4 µm, were lower than those of ECAP-RR sample, with a grain size of 16 µm. This was explained by the observation of an incomplete recrystallisation process for the ECAP-RR sample, whereas a complete recrystallization occurred for the ECAP-RH sample [30,33]. Indeed, the increase in processing temperature for ECAP-RH during the second pass led to a complete recrystallisation process, and then to a decrease in the deformation stored in grains, which induced a slight softening of the alloy.

20 Fig. 2 .

 202 Fig. 2. Tensile curves plotted for the as-received, ECAP-RR and ECAP-RH samples at 10-3 s-1 •reminder, it was shown previously, for the as-received sample, that the microstructure was characterised by a heterogeneous grain structure (with an average grain size of 24 µm) and the presence of IMCs[START_REF] Rochet | Influence of equal-channel angular pressing on the microsnucture and corrosion behaviour of a 6:xxx aluminium alloy for automotive conductors[END_REF]. Both ECAP-RR and ECAP-RH processes were found to induce a sig nificant grain refinement (associated with an increase in HAGB density) and a slight breakdown of IMCs. As discussed in the introduction, these microstructural heterogeneities influence the fatigue behaviour of Al alloys. Considering that ECAP induced an increase in the density of such heterogeneities, i.e. an increase in grain boundaries and !MC density, it was assumed that the plastically deformed volume will remain re presentative of the microstructure during fatigue tests, even for ECAP RR and ECAP-RH samples. This assumption is based on consideration of the decrease in the plastically deformed volume through a decrease in the microstructural scale. Fig.3thus shows S-N curves obtained for the as-received, ECAP-RR and ECAP-RH samples at 93%, 90% and 80% of the UTS. The fatigue limit (fixed arbitrarily at 10 7 cycles in this work) was reached for the as received, ECAP-RR and ECAP-RH samples as soon as the applied max imum stress was lower than 90% of the UTS. This clearly showed that it was necessary to reach a very high stress level ( > 90% of the UTS and > YS 0 _ 2 ) to induce a fracture whatever the samples. At 93% of the UTS, the numbers of cycles to failure were very close for ail samples. However, this stress level (expressed as a percentage of UTS) corre sponded to a higher effective stress level (expressed in MPa) for ECAP RH sample, and particularly for ECAP-RR sample. This was related to the enhancement of mechanical properties already observed on tensile curves, and was explained by the grain refinement induced by the SPD process. Indeed, the increase in HAGBs density could create efficient barriers to crack propagation. This should be responsible for a decrease in the cumulated plastic deformation during the first cycles, and should lead to a delay in the crack initiation for ECAP samples. In addition, during fatigue tests, this should induce a decrease in the dislocation

  slip bands leading to a Jess pronounced extrusion/ intrusion phenomenon at the sample surface resulting in a further delay in the crack initiation for ECAP-RR and ECAP-RH samples[START_REF] Chung | Improvement of high-cycle fatigue life in a 6061 Al alloy produced by equal channel angular pressing[END_REF]. Fur thermore, as explained previously, the increase in ECAP processing temperature for ECAP-RH sample, leading to a complete recrystallisa tion process, induced a slight softening of the alloy. This could explain the differences observed between the two ECAP-processes samples. Fig. 4 shows SEM observations of the fracture surfaces of as-re ceived, ECAP-RR and ECAP-RH samples after fatigue tests performed at 93% of the UTS. No difference in fracture modes was observed between the as-received, ECAP-RR and ECAP-RH samples. For ail samples, the fracture surfaces showed a single initiation site localised on the sample surface (Fig. Sa). Crack initiation site was most likely associated with the presence of IMCs (Fig. Sc) or extrusion/intrusion zones. The crack propagation zone was characterised by a granular aspect (Fig. 5d), which was consistent with the grain size of as-received, ECAP-RR and ECAP-RH samples measured in a previous study [30]. Therefore, it was considered that the crack propagated by intergranular fracture. No striations were observed for ail samples along this intergranular crack propagation zone; this could suggest an unsteady crack propagation, which did not progress at each cycle. This type of fracture had already been observed in previous work [15] and by Esmaeili et aL for AA7075 [6]. The numerous grain boundaries observed in ECAP-processed samples, in relation to the decrease in grain size due to the deformation process, acted as microstructural barriers, which finally promoted an intergranular fracture. Finally, it was observed that the final fracture zone was classically ductile and showed a high density of dimples (Fig. Sb) for ail samples.

Fig. 6

 6 Fig.6shows S-N curves obtained in laboratory air at 90% and 80% of the UTS, for the as-received, ECAP-RR and ECAP-RH samples after a pre-corrosion step of 72 h in 0.5 M NaCI at E,orr• Results from tests performed without pre-corrosion were reported for comparison (grey areas). At 90% of the UTS, a decrease in fatigue lifetime was observed for pre-corroded samples for both the as-received and ECAP-processed samples. The ratio Npc72llNAfr (where Npcnh and NAfr were the average numbers of cycles to failure for the tests performed on pre-corroded and non-corroded samples, respectively) decreased from 0.07 ± 0.03 for the as-received sample to 0.05 ± 0.02 and 0.02 ± 0.01 for the ECAP RR and ECAP-RH samples, respectively. At 80% of the UTS, ail pre corroded as-received samples reached the fatigue limit, whereas pre corroded ECAP-RR and ECAP-RH samples both exhibited premature fractures ( < 10 6 cycles).The corresponding fracture surfaces are reported in Fig.7. Contrary to non-corroded samples, crack initiation sites were systematically re lated to corrosion surface defects (Fig.7g, h and i). SEM observations of the samples revealed that those defects were homogeneously

Fig. 3 .

 3 Fig. 3. S-N curves (maximum stress (in % UTS or in MPa) vs. nurnber of cycles to failure) plotted for the as-received, ECAP-RR and ECAP-RH samples. The fatigue tests were perfonned in laboratory air for non-corroded sam.ples at 25 °C (f = 50 Hz, R = 0.1, sinusoidal stress loading). Black arrows refer to non broken sam.ples.

Fig. 9 Fig. 4 .Fig. 5 . 10 7Fig. 6 .

 945106 Fig. 9 shows S-N curves obtained for the as-received, ECAP-RR and ECAP-RH samples tested at 90% and 80% of the UTS in 0.5 M NaCl at E corr • Previous results obtained for non-corroded and PC72h pre-

Fig. 7 .Fig. 8 . 2

 782 Fig. 7. SEM observations of the fracture surface of (a), (d) and (g) as-received sample; (b), (e) and (h) ECAP-RR sample; (c), (f) and (i) ECAP-RH sample tested at 90% of the UTS after pre-corrosion during 72 h in 0.5 M NaCl at Ecorr• White arrows G), (k) and (1) indicate corrosion defects acting as crack initiation sites. The fatigue tests were performed in laboratory air at 25 °C (f = 50 Hz, R = 0.1, sinusoidal stress loading).(a)

  showed that the increase in the cor roded surface ratio for corrosion fatigue tests as compared to fatigue tests on PCl h30 pre-corroded sampi es was mainly due to an increase in the corrosion defect density, with no significant effect of the cyclic mechanical loading on the diameter of the corrosion defects, as ob served above. Similar analyses were performed after fatigue tests at 80% of the UTS on samples pre-corroded during 16 h in 0.5 M NaCI by comparison with corrosion fatigue tests performed at 80% of the UTS (the mean duration of these tests was about 16 h). For brevity, these results are not reported in this paper but the conclusions were very similar, even if the differences observed between corrosion fatigue tests and fatigue test.s on pre-corroded samples were less pronounced th.an for tests perforrned at 90% of the UTS. Therefore, these results con firrned that an interaction occurred between the cyclic mechanical loading and the corrosion processes leading to an increase in the density of corrosion defects. Severa! mechanisms had been proposed in the literature to explain this synergy effect. The ones of interest in our case are those describing the influence of cyclic mechanical loading on oxide

Fig. 9 .

 9 Fig. 9. S-N curves of as-received, ECAP-RR and ECAP-RH samples plotted fol lowing corrosion fatigue tests performed at 25 °C (f = 50 Hz, R = 0.1, sinu soidal stress loading) in 0.5 M NaCI at Ecorr• Plain and outlined grey areas refer to previous results obtained without pre-corrosion and with pre-corrosion during 72 h at Ecom respectively. Arrows refer to non-broken samples.

Fig. 10 .

 10 Fig. 10. SEM observations of the fractnre surfaces of (a), (d) and (g) as-received sample; (b), (e) and (h) ECAP-RR sample; (c), (f) and (i) ECAP-RH sample tested at 90% of the UTS in 0.5 M NaCl at Ecorr• The corrosion fatigue tests were performed at 25 °C (f = 50 Hz, R = 0.1, sinusoidal stress loading).

Fig. 11 . 1 .

 111 Fig. 11. Comparison of the number of cycles to failure determined for the as received, ECAP-RR and ECAP-RH samples after corrosion fatigue test.s and PClh30 tests. Results obtained for PC72h tests are reported for comparison. Both corrosion fatigue tests and fatigue tests on pre-corroded samples were performed at 25 °C (f = 50 Hz, R = 0.1, sinusoidal stress loading). following: 1. Preliminary fatigue tests perforrned in laboratory air showed an increase in fatigue lifetime at high stress levels for ECAP-RR and

Fig. 12 .

 12 Fig. 12. (a) Depth; (b) equivalent diameter; (c) corroded surface ratio and (d) density of corrosion defects measured for the as-received (AR), ECAP-RR and ECAP-RH samples after pre-corrosion in the PC72h and PClh30 conditions and after corrosion fatigue tests.

1.54 ± 0.08 Diameter (µm) 2.72 ± 0.25 2.05 ± 0.09 2.38 ± 0.11 Density (10 3 mm -2 ) 1.71 ± 0.05 5.36 ± 0.05 3.52 ± 0.05 Corroded Surface Ratio(%) 0.91 ± 0.21 1.38 ± 0.22 1.44 ± 0.21

Acknowledgements

The authors thank the ANRT for their financial support (Clement Rochet's PhD thesis). They also thank LEONI Wiring Systems for their financial support.

Data availability

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


Egor V Verbitskiy

Sylvain Achelle

Filip Bureš

Pascal Le Poul

Alberto Barsella

Yuriy A Kvashnin

Françoise Gennady L Rusinov

Robin-Le Guen

Oleg N Chupakhin

Valery N Charushin

Françoise Robin-Le Guen

3,4-b]pyrazine-based linear push-pull systems

Keywords: Synthesis, photophysical and nonlinear optical nitrogen heterocycles, donor-acceptor systems, nonlinear optic, fluorescence, intramolecular charge transfer

come

Introduction

Organic nonlinear optical (NLO) materials have attracted considerable attention in recent years because of their potential applications in various fields of optoelectronics and photonics, including optical communications, optoelectronic transfer, optical signal process and optical data storage [1]. The main advantages of these materials are high NLO coefficients, ultrafast optical responses, flexible molecular design and synthesis and higher optical damage threshold, as compared to traditional inorganic solids. Generally, the NLO properties of organic materials are dependent on the molecular structures of chromophores. Typically, in order to show a significant second-order NLO response, a molecule must be non-centrosymmetric, with intramolecular charge transfer (ICT) transitions at relatively low energy and characterized by a large transition dipole moment and a large difference between the excited state and the ground state molecular dipole moment [1e]. This can be achieved in linear organic molecules by connecting an electron-donor and an electron-acceptor group through a π-conjugated polarizable spacer, as it occurs in classical organic dipolar push-pull systems [START_REF] Bureš | Fundamental aspects of property tuning in push-pull molecules[END_REF].

Pyrazine is a nitrogen based heterocyclic compound with a rigid planar conjugated structure. Due to the electron-deficient character of the pyrazine, the pyrazinyl fragment can be used as electron-withdrawing group in push-pull structures. The incorporation of pyrazine units in luminescent materials has been recently reviewed [START_REF]1:1, 50 mL), and the organic layer was separated. The aqueous layer was extracted with EtOAc (2 × 25 mL). The combined organic extracts were dried with MgSO 4 and the solvents evaporated. Purification by silica gel column chromatography with CH 2 Cl 2 /hexane (1:2, v/v) as an eluent to afford the title compounds (7 and 8). 4[END_REF]. Moreover, it has been shown that pushpull pyrazine derivatives exhibit interesting second-order (frequency doubling) [4] and thirdorder (two photon absorption) [5] NLO properties.

It has been demonstrated that [1,[START_REF] Bureš | Fundamental aspects of property tuning in push-pull molecules[END_REF]5]oxadiazolo [3,4-b]pyrazine (furazanopyrazine) exhibits higher electron-withdrawing character than similar non-annulated pyrazine analogues [START_REF] Sheremetev | Advances in the chemistry of furazano[3,4-b]pyrazines and their analogues[END_REF]. Recently, the [1,[START_REF] Bureš | Fundamental aspects of property tuning in push-pull molecules[END_REF]5]oxadiazolo [3,4-b]pyrazine core has been used as an acceptor part of push-pull molecules exhibiting pronounced fluorescence in the solid state and sensitivity towards nitroaromatic compounds [7].

This work represents further extension of our previous research focusing on the design of novel fluorophores that incorporate diazine scaffold. In this contribution we describe the synthesis of a new series of D-π-A [where A equals to the 1,2,5-oxadiazolo [3,4-b]pyrazine acceptor, D represents triphenylamine or carbazole donors, and π involves thienylene or phenylene groups] push-pull chromophores 7a,b and 8a,b and systematic investigation of their photophysical and second order NLO properties. The influences of the type of the πlinker as well as the nature of the electron-donating group on the photophysical properties were thoroughly studied and structure-property relationships were elucidated.

Results and Discussion

Synthesis

The synthesis of the linear push-pull systems required 5-(4-bromophenyl)- [1,[START_REF] Bureš | Fundamental aspects of property tuning in push-pull molecules[END_REF]5]oxadiazolo [3,4-b]pyrazine [START_REF]1:1, 50 mL), and the organic layer was separated. The aqueous layer was extracted with EtOAc (2 × 25 mL). The combined organic extracts were dried with MgSO 4 and the solvents evaporated. Purification by silica gel column chromatography with CH 2 Cl 2 /hexane (1:2, v/v) as an eluent to afford the title compounds (7 and 8). 4[END_REF] and 5-(5-bromothiophen-2-yl)- [1,[START_REF] Bureš | Fundamental aspects of property tuning in push-pull molecules[END_REF]5]oxadiazolo [3,4-b]pyrazine (5) as starting materials. Compound 3 was prepared via a two-step procedure from commercially available 4'-bromoacetophenone (1) in 73% overall yield (Scheme 1).

Scheme 1. Synthesis of 5-(4-bromophenyl)- [1,[START_REF] Bureš | Fundamental aspects of property tuning in push-pull molecules[END_REF]5]oxadiazolo [3,4-b]pyrazine [START_REF]1:1, 50 mL), and the organic layer was separated. The aqueous layer was extracted with EtOAc (2 × 25 mL). The combined organic extracts were dried with MgSO 4 and the solvents evaporated. Purification by silica gel column chromatography with CH 2 Cl 2 /hexane (1:2, v/v) as an eluent to afford the title compounds (7 and 8). 4[END_REF].

The second starting material (5) was obtained by bromination of previously described 5-(1benzothiophen-3-yl)pyrimidine (4) [7] with an excess of N-bromosuccinimide in 85% yield. Scheme 4. Synthesis of 5-[5-(heteroaryl)thiophen-2-yl]- [1,[START_REF] Bureš | Fundamental aspects of property tuning in push-pull molecules[END_REF]5]oxadiazolo [3,4-b]pyrazines (8a,b).

Electrochemical properties

Electrochemical behavior of compounds 7 and 8 was studied by cyclic voltammetry (CV) in CH 2 Cl 2 containing Bu 4 NPF 6 electrolyte at a scan rate of 0.1 V/s. The working electrode was a glassy carbon disk, Pt wire was used as the counter electrode and an Ag wire as a reference electrode. Ferrocene was used as an internal reference for potential measurements. The first oxidation/reduction peak potentials and their differences are listed in Table 1 and CV diagrams are shown in Figure 1.

Table 1: Electrochemical data of compounds

Comp. Compounds 7b and 8b bearing the carbazole moiety as an electron-donating group exhibit a first irreversible oxidation process at 0.95 and 0.96 V vs. ferrocene. When the donor is the diphenylamino moiety, the first oxidation is reversible and measured E 1/2 are much lower (0.55 V for 7a and 0.57 V for 8a; see Figure 1) in accordance with the stronger electron-donating effect of the diphenylamino group. For these two compounds, a second irreversible oxidation peak can be observed at the higher potential that modifies the reversibility of the first system, on the reverse scan.

E 1/2 ox1 [V] a E 1/2 red1 [V] a ∆E [V] b E HOMO [eV] c E LUMO [eV] c λ max [nm] d 7a
All compounds bear the same oxadiazolopyrazine electron-withdrawing group. For this reason, the same behavior can be observed on reduction when scanning from 0 to -2 V. All compounds show a first reversible reduction around -1.1 V vs. Fc and a second quasi-reversible system around -2 V (scan rate = 0.1 V/s). The latter becomes reversible when the scanning rate is increased to 1 V/s (Figure S25). The obtained half-wave potentials for both systems are similar regardless of the nature of the electron-donating moieties and of the π-conjugated linker. This trend indicates a weak interaction between the donor and the acceptor.

The HOMO/LUMO energies were calculated from potential values (the first oxidation and reduction processes) to evaluate the electrochemical band gap. The E LUMO remains unaltered as all chromophores showed the same reduction behavior. On the other hand, the E HOMO of diphenylamino derivatives is higher leading to a decreased band gap for compounds 7a and 8a.

These results correspond to the evolution of λ max values obtained from the electronic spectra.

Photophysical properties

The photophysical properties of compounds 7 and 8 have been studied in various aprotic solvents of increasing polarity and the results are summarized in Table 2. All compounds are emissive in n-heptane. Carbazole derivatives 7b and 8b are also emissive in 1,4-dioxane but no emission is observed in solvents of higher polarity. For push-pull derivatives, such observation is rationalized by decreased emission intensity observed in polar solvents due to lowered energies of the ICT excited state [8]. Figure 2 shows the normalized absorption and emission spectra in nheptane. As expected, the replacement of the 1,4-phenylene linker (compounds 7) by a 2,5thienylene one (compounds 8) induces a red shift of both absorption and emission bands as well as increases the molar extinction coefficient due to a better charge transfer through the heterocycle [9]. Similarly, to previous observation [10], the diphenylamino derivatives exhibited red-shifted absorption and emission. This effect is more pronounced than changing the π-conjugated linker. The absorption and emission maxima increase in the following order 7b < 8b < 7a < 8a. Compound 8b exhibited the highest emission quantum yield in n-heptane (Φ F = 0.96). The lowest value was observed for compound 7b (Φ F = 0.06), but this compound exhibited much higher quantum yield in 1,4-dioxane (Φ F = 0.29). Figure 3 shows the color of nheptane solutions of compounds 7-8 under UV-irradiation. Whereas only moderate absorption solvatochromism was observed for 7 and 8, which does not seem to be correlated to the solvent polarity (the most red-shifted absorption bands are observed in CH 2 Cl 2 ), a strong positive emission solvatochromism was observed for carbazole derivatives 7b and 8b as illustrated in Figures 4 and5. It should be noted that emission solvatochromism was not observed for diphenylamino derivatives since the emission is fully quenched in solvents of higher polarity than n-heptane. This phenomenon, which is due to stabilization of the highly-polar emitting state by polar solvent, has been extensively described for push-pull derivatives [11]. It should be noted however that the quantum yield of compound 7b is higher in 1,4-dioxane than in n-heptane, probably due to aggregation in this latter solvent.

The slope of the regression line of the emission maxima plotted versus the Dimroth-Reichardt polarity parameter (E T (30)) is a way to evaluate the ICT [12]. The slope observed for compounds 7b and 8b (see Figure S1 in ESI) are similar or higher as compared to pyrazine push-pull derivatives with diphenylamino substituents and more extended π-conjugated linker [START_REF] Hoffert | Synthesis and photophysical properties of a series of pyrazine-based push-pull chromophores[END_REF],

indicating that the oxadiazolo part significantly enhances the electron-withdrawing character of the pyrazine. Photographs were taken in the dark upon irradiation with a hand-held lamp (λ em = 366 nm)

Second order nonlinear optical properties

The second order NLO properties have been studied in chloroform solution by the electric-field induced second harmonic generation (EFISH) method at a non-resonant incident wavelength of 1907 nm. The second harmonic at λ = 953 nm is therefore well clear of the absorption bands of the chromophores. This method provides the NLO response as the scalar product between the permanent dipolar moment of the molecule in fundamental state and the vector component of

β

DFT calculations

Spatial and electronic properties of compounds 7a-b and 8a-b were investigated using Gaussian ® 16W software [17] package at the DFT level. Optimized geometries, energies of the frontier molecular orbitals, and ground state dipole moments µ and were calculated at DFT B3LYP/6-311+G(2df,p) level in CHCl 3 as shown in Table 4. The electronic absorption spectra and the first-order hyperpolarizabilities β were calculated by TD-DFT method using CAM-B3LYP/6-311+G(2df,p) in CHCl 3 . The calculated energies of the HOMO and the LUMO range from -5.83 to -5.45 and -3.61 and -3.46 eV (Figure 6), respectively. Both calculated and electrochemical HOMO-LUMO gaps obey the same trends and correlates tightly (Figure S2) and, therefore we can consider the used DFT method as reliable. In general, the lowest HOMO-LUMO gap has been calculated for chromophores in series a bearing N,N-diphenylamino donor. The carbazole-terminated chromophores b showed deepened HOMO. The gaps are slightly lower for chromophores 7 with 1,4-phenylene π-linker. Figure 7 shows optimized geometries, HOMO/LUMO localizations and Mulliken charges in target chromophores; see Figure S3 for complete listing of frontier molecular orbitals. All molecules are unsymmetrical and do not belong to any group of symmetry. Whereas the HOMO is centrally localized either on the diphenylamino or carbazole donors, the LUMO is spread over the adjacent 1,4-phenylene bridge. The Mulliken charges are localized accordingly, involving the donor moieties in 8 and mostly the π-conjugated bridge between the particular donor and acceptor.

The calculated dipole moments reflect the geometry with the out-of-plane orientation of both donor moieties (the torsion angle is about 50°). However, molecules in series a bearing more flexible diphenylamino group possess significantly higher ground state dipole moments of around 11 and 13 D as compared to carbazole derivatives b (Table 4).

The fundamental electronic absorption properties were investigated by TD-DFT method;

the calculated positions of the longest-wavelength absorption maxima λ max DFT are gathered in First-order hyperpolarizabilities β were also calculated as shown in Table 4. The calculated β coefficients range from 85 to 373 × 10 -30 esu. Correlations of the calculated β and experimental µβ or µβ 0 products (EFISH experiment) are provided in Figures S5 andS6. Both correlations are very tight. In general, larger nonlinearities were calculated for chromophores a with N,N-diphenylamino donor, the largest one belongs to chromophore 8a with polarizable 2,5-thienylene π-linker.

Experimental section

General Information. All reagents and solvents were obtained from commercial sources and dried by using standard procedures before use. 1 H and 13 C NMR spectra were recorded on a Bruker AVANCE-500 and AVANCE-600 instruments using Me 4 Si as an internal standard. IR spectra of samples (solid powders) were recorded on a Spectrum One Fourier transform IR spectrometer (Perkin Elmer) equipped with a diffuse reflectance attachment (DRA) in the frequency range 4000÷400 cm -1 . Spectrum processing and band intensity determination were carried out using the special software supplied with the spectrometer. Elemental analysis was carried on a Eurovector EA 3000 automated analyzer. Melting points were determined on Boetius combined heating stages and were not corrected. The chromatographic purification of compounds was achieved with silica gel Alfa Aesar 0.040-0.063 mm (230-400 mesh), eluting with CH 2 Cl 2 /hexane (1:2, v/v). The progress of reactions and the purity of compounds were checked by TLC on Sorbfil plates (Russia), in which the spots were visualized with UV light (λ 254 or 365 nm). The electrochemical studies of the compounds were performed with a homedesigned 3-electrodes cell (WE: glassy carbon disk, RE: Ag wire, Ce: Pt). Ferrocene was added at the end of each experiment to determine redox potential values. UV Vis and fluorescence spectra were recorded with a Fluoromax-3 Jobin-Yvon Horiba spectrophotometer using standard 1 cm quartz cells. Compounds were excited at their absorption maxima (band of lowest energy)

to record the emission spectra. The Φ F values were calculated using a well-known procedure [START_REF] Eaton | References materials for fluorescence measurement[END_REF] with 9,10-bis(phenylethenyl)anthracene in cyclohexane (Φ F = 1.00) as standard [START_REF] Taniguchi | Database of absorption and fluorescence spectra of >300 common compounds for use in PhotochemCAD[END_REF].

Experimental details on EFISH measurements are described elsewhere [START_REF] Ulrich | BODIPY-bridged push-pull chromophores for nonlinear applications[END_REF]. washed with 1,4-dioxane (5 mL). The solvent was evaporated at reduced pressure. The residue was dissolved in a mixture of ethanol (5 mL) and acetic acid (5 mL), 152.0, 149.2, 146.8, 139.0, 137.0, 130.3, 127.9, 125.9, 125.6, 125.5, 124.7, 121.9 4, 154.7, 152.7, 152.2, 152.1, 140.5, 140.3, 138.4, 136.8, 131.8, 128.4, 127.8, 127.3, 126.9, 123.5, 121.1, 120.9, 110.3. ν (DRA, cm

Synthesis

Conclusion

In summary, four new push-pull chromophores bearing [1,[START_REF] Bureš | Fundamental aspects of property tuning in push-pull molecules[END_REF]5]oxadiazolo [3,4-b]pyrazine as electron-withdrawing part and amino group as donor were designed. There electrochemical and photophysical properties indicate that intense ICT occurs in these structures with electrochemical gap below 2.1 eV and strong emission solvatochromism. NLO responses were also measured by EFISH method. All experimental and theoretical results indicate a significant increase of ICT when the diphenylamino electro-donating group is used and when a 2,5thienylene bridge is replacing the 1,4-phenylene linker. Compound 8a, that combines these two characteristics, exhibits a particularly high figure of merit and appears as an interesting candidate for incorporation in a polymeric matrix to obtain a material with the high electro-optic coefficient.

Scheme 2 .Scheme 3 .

 23 Scheme 2. Synthesis of 5-(5-bromothiophen-2-yl)-[1,[START_REF] Bureš | Fundamental aspects of property tuning in push-pull molecules[END_REF]5]oxadiazolo[3,4-b]pyrazine (5).

Figure 1 .

 1 Figure 1. Cyclic voltammograms of 7a (green), 8a (orange) 7b (blue) and 8b (red) in CH 2 Cl 2

aFigure 2 :

 2 Figure 2: Normalized UV/Vis (solid lines) and emission spectra (dashed lines) of compounds 7

Figure 3 :

 3 Figure 3: n-heptane solution under UV lamp irradiation (from left to right: 7b, 8b, 7a, 8a) c = 1-3×10 -5 M. Photographs were taken in the dark upon irradiation with a hand-held lamp (λ em = 366

Figure 4 :

 4 Figure 4: Normalized emission spectra of 7b in different aprotic solvents

Figure 6 :

 6 Figure 6: Energy level diagram showing the electrochemical (black) and calculated (red)

Figure 7 :

 7 Figure 7: HOMO (red) and LUMO (blue) localizations (above) and Mulliken charges (below) in

  . The spectra consist of two peaks; the low-energy peak corresponds to the HOMO→LUMO transition whereas the high-energy peak involves mostly the HOMO-2→LUMO and HOMO→LUMO+1 transitions. When comparing the corresponding pairs of chromophores, chromophores 8a-b bearing 2,5-thienylene linker possess bathochromically shifted longest-wavelength absorption maxima. The same trend can be seen for chromophores a with N,N-diphenylamino donor, when compared to carbazole derivatives b.

1

 1 (1.994 g, 10 mmol) and selenium dioxide (1.1 g, 10 mmol) in a solution of 1,4-dioxane (15 mL) and water (1 mL) was refluxed for 12 h. Selenium was filtered off,

  6a,b(1.2 mmol), Pd(PPh 3 ) 4 (115 mg, 10 mol %) and K 3 PO 3 (530 mg, 2.5 mmol) was dissolved in 1,4-dioxane 15 mL. The reaction mixture was degassed and refluxed for 15 h under an argon atmosphere. After completion of the reaction (monitored by TLC), the reaction mixture was cooled, filtered, and dissolved in a mixture of EtOAc and water

  -1 ) 3093 (br. w, C-H Ar ), 3080 (br. w, C-H Ar ), 3060 (br. w, C-H Ar ), 3028 (br. w, C-H Ar ), 1599 (s, C-C Ar /C-N Ar ), 1567 (s, C-C Ar / C-N Ar ), 1533 (s, C-C Ar / C-N Ar ), 1449 (s, C-C Ar / C-N Ar ), 1406 (s, C-C Ar / C-N Ar ), 837 (s, C-H Ar ), 749 (s, C-H Ar ), 724 (s, C-H Ar ). Calcd. for C 26 H 15 N 5 OS (445.50): C, 70.10; H, 3.39; N, 15.72. Found: C, 70.00;[START_REF] Castet | Linear and nonlinear optical properties of arylvinyldiazine dyes: a theoretical investigation[END_REF] 3.54; N, 15.56.

Table 2 :

 2 UV/Vis and PL data in various solvents for compounds 7 and 8.

	Compds	Solvent	λ abs (ε)	λ em	Φ F a	Stokes shift
			[nm] [(mM -1 cm -1 ])	[nm]		[cm -1 ]
		n-heptane	473 (18.4)	570	0.54	3598
		1,4-dioxane 466 (12.4)	-b	-b	-b
	7a
		CH 2 Cl 2	485 (11.8)	-b	-b	-b
		MeCN	457 (16.4)	-b	-b	-b
		n-heptane	414 (12.3)	521	0.06	4961
		1,4-dioxane 401 (8.7)	619	0.29	8783
	7b
		CH 2 Cl 2	403 (8.9)	-b	-b	-b
		MeCN	389 (15.1)	-b	-b	-b
		n-heptane	514 (33.2)	607	0.39	2980
		1,4-dioxane 511 (22.7)	-b	-b	-b
	8a
		CH 2 Cl 2	544 (22.7)	-b	-b	-b
		MeCN	518 (20.7)	-b	-b	-b
		n-heptane	452 (25.4)	548	0.96	3875
		1,4-dioxane 450 (22.5)	627	0.35	6273
	8b
		CH 2 Cl 2	463 (20.

  described as β //[14]. The two level corrected µβ values (µβ 0 ) have been also calculated[15]. The results are presented in Table3. It should be noted that positive µβ values are obtained indicating that both ground and excited states are polarized in the same direction and that the excited state is more polarized than ground state, in accordance with the

emission solvatochromism observation (for 7b and 8b). Compounds 8 exhibit higher NLO response than analogues 7 and the diphenylamino derivatives a show better µβ values than carbazole analogues b. These results are coherent with the electrochemical and photophysical observations indicating better ICT for compounds possessing 2,5-thienylene linker and diphenylamino electron-donating group. The figure of merit of compound 8a appears particularly interesting.

Table 3 :

 3 Results for EFISH measurements for compounds 7 and 8.

		7a	7b	8a	8b
	µβ (10 -48 esu) a	700	70	1440	330
	µβ 0 (10 -48 esu) b	485	53	892	237
	µβ/MW c	1.91	0.19	3.88	0.89
	a µβ(2ω) at 1907 nm in CHCl 3 . Molecular concentrations used for the measurements were in the
	range of 10 -3 to 10 -2 M, µβ ± 10% b Two level corrected µβ values (µβ 0 ) [14]. c Figure of merit
	(molecular nonlinearity) 10 -48 esu g mol -1 [16]

Table 4 :

 4 DFT-calculated electronic parameters of 7 and 8.

		7a	7b	8a	8b
	E HOMO (eV) a	-5.45	-5.83	-5.48	-5.83
	E LUMO (eV) a	-3.53	-3.61	-3.46	-3.58
	∆E (eV) a	1.92	2.22	2.02	2.25
	µ (D) a	10.7	6.5	13.1	8.8
	λ max DFT (nm/eV) b	422(2.94)	367(3.38)	479(2.59)	439(2.82)
	β(-2ω,ω,ω) (10 -30 esu) b,c	215	85	373	149

a Calculated at DFT B3LYP/6-311++G(2df,p) level in CHCl 3 . b Calculated at TD-DFT (nstates = 8) CAM-B3LYP/6-311+G(2df,p) level in CHCl 3 . c At 1907 nm in CHCl 3 .

Table 4 ,

 4 the spectra are provided in FigureS4. The calculated λ max

DFT values correlate tightly with the experimental longest-wavelength absorption maxima measured in dichloromethane (Table

2

) -see Figure

S5

  . ν (DRA, cm -1 ) 3085 (w, C-H Ar ), 3062 (w, C-H Ar ), 3035 (w, C-H Ar ), 1589 (s, C-C Ar / C-N Ar ), 1568 (s, C-C Ar / C-N Ar ), 1488 (s, C-C Ar /C-N Ar ), 1438 (s, C-C Ar /C-N Ar ), 1405 (s, C-C Ar / C-N Ar ), 833 (s, C-H Ar ), 751 (s, C-H Ar ), 694 (s, C-H Ar ). Calcd. for C 26 H 17 N 5 OS (447.52): C, 69.78; H, 3.83; N, 15.65. Yield 321 mg (72%), red solid, mp 246-248 °С. 1 H NMR (600 MHz, DMSO-d 6 ) δ 9.82 (d, J = 2.1 Hz, 1H), 8.66 (dd, J = 4.0, 2.2 Hz, 1H), 8.28 (dt, J = 7.8, 1.1 Hz, 2H), 8.21 -8.15 (m, 2H),8.00 (dd, J = 4.0, 2.2 Hz, 1H), 7.83 -7.78 (m, 2H), 7.54 -7.44 (m, 4H), 7.33 (ddd, J = 7.9, 6.9, 1.2 Hz, 2H).13 C NMR (151 MHz, DMSO-d 6 ) δ 155.

	Found: C, 69.83; H, 3.77; N, 15.67.
	5-(5-(4-(9H-Carbazol-9-yl)phenyl)thiophen-2-yl)-[1,2,5]oxadiazolo[3,4-b]pyrazine	(8b).

Acknowledgments

VEV is grateful to the financial support for the synthetic part from the Russian Foundation for Basic Research (Research Project No. 18-29-23045 mk).