Daniel Maisonnette

Didier Bardel

Vincent Robin

Daniel Nélias

Michel Suéry

D Nelias

M Suery

Mechanical behaviour at high

Keywords: 6061 Aluminium Alloy, Thermomechanical properties, Electron Beam Welding, Stress-strain curves, Yield strength, Finite Element modeling

Introduction

The 6061-T6 aluminium alloy (AA) is an age hardenable alloy, it means that its mechanical properties are mainly controlled by the hardening precipitates contained within the material [START_REF] Simar | Par-572 doen, Integrated modeling of friction stir welding of 6xxx series Al alloys: 573 Process, microstructure and properties[END_REF]. The optimum mechanical properties in terms of yield strength are gotten after a solution heat treatment followed by a quenching and a tempering treatment at 175 • C [START_REF] Bardel | Coupled precipitation 577 and yield strength modelling for non-isothermal treatments of a 6061 alu-578 minium alloy[END_REF][START_REF] Bardel | 580 Cyclic behaviour of a 6061 aluminium alloy: coupling precipitation and 581 elastoplastic modelling[END_REF]. The present work focuses on the mechanical properties of this alloy after and during various rapid heat treatments as encountered during electron beam (EB) welding. To achieve this goal the mechanical properties of the AA6061 alloy will be investigated after different temperature histories. This kind of data are necessary for Computational welding Mechanics (CWM) DB in order to predict residual stresses and distortions due to welding operations [START_REF] Pickett | Model 583 characterisation and failure analysis of welded aluminium components 584 including process history[END_REF][START_REF] Smith | Measured and predicted residual stresses in thick section 588 electron beam welded steels[END_REF][START_REF] Robin | From the numerical modeling of manufacturing processes and 591 welding in particular to the mechanical behavior of weldments[END_REF][START_REF] Robin | Fracture analyses of welded 594 components[END_REF]. In this paper, it will be shown that, in a single pass welding process a simple isokinetic model [START_REF] Grong | Microstructural modelling in metals processing[END_REF] provides a good estimation of the mechanical properties after transient heat treatments when coupled to the work-hardening model proposed in this work. Note that the recently performed simulations of laser welding by [START_REF] Zain-Ul Abdein | 600 Finite element analysis of metallurgical phase transformations in aa 6056-601 t4 and their effects upon the residual stress and distortion states of a laser 602 welded t-joint[END_REF], friction stir welding [START_REF] Yuan | 605 Effect of tool design and process parameters on properties of Al alloy 606 6016 friction stir spot welds[END_REF], laser shot peening [START_REF] Ren | Metallographic structure evolution of 6061-t651 aluminum alloy pro-610 cessed by laser shock peening: Effect of tempering at the elevated tem-611 peratures[END_REF] and gas metal arc welding [START_REF] Ahmad | Effect of a post-weld heat treatment on the me-613 chanical and microstructure properties of aa6061 joints welded by the gas 614 metal arc welding cold metal transfer method[END_REF] require this type of database, as well as additive manufacturing modeling.

The behavior of AA6061 at room temperature -after various thermal histories -has been investigated by Bardel et al. * daniel.nelias@insa-lyon.fr [START_REF] Bardel | Gar-617 nier, Integrated modelling of a 6061-t6 weld joint: From microstructure 618 to mechanical properties[END_REF]. It shows that the mechanical properties, such as the yield al. [START_REF] Pickett | Model 583 characterisation and failure analysis of welded aluminium components 584 including process history[END_REF] and Abotula and Chalivendra [START_REF] Abotula | An experimental and numerical investigation 620 of the static and dynamic constitutive behaviour of aluminium alloys[END_REF], but also the work of Robin et al. [START_REF] Robin | 3d modelling of electro-magneto 623 forming of metal sheets[END_REF] who applied their model to Electro-Magnetic Forming (EMF). These data will be then used as input for a two-phase thermo-mechanical model and we will demonstrate the good prediction that this model can provide thanks to a confrontation between welding simulations and instrumented experiments (including neutron diffraction).

Mechanical testing at high temperature

Experimental procedure

A home made heating device using Joule effect with continuous current was adapted to a tensile machine to monitor the temperature history during the tensile test by the mean of thermocouples connected to a temperature controller. Here, the specimens are heated to the test temperature (see Fig. 1) and the tensile test starts immediately at a controlled strain rate ε=8.10 -5 s -1 . Moreover, an extensometer with ceramic tips, a load cell and a video camera are used to record the local uniaxial strain, the strength and the evolution of the diameter, respectively. [START_REF] Chanda | A comparative study on iso-speed extru-633 sion and isothermal extrusion og 6061 Al alloy using 3D FEM simulation[END_REF][START_REF] Massardier | Correlation between the microstruc-636 tural evolution of a 6061 aluminium alloy and the evolution of its thermo-637 electric power[END_REF]) to 81 get the so-called O temper state (all the hardening precipitates 82 are dissolved even without dwell time [START_REF] Bardel | Coupled precipitation 577 and yield strength modelling for non-isothermal treatments of a 6061 alu-578 minium alloy[END_REF] ). Then, thanks to 83 a special design of the machine and the specimens to get fast 84 cooling (above 15 • C/s as suggested in [START_REF] Li | Quench 639 sensitivity of 6351 aluminum alloy[END_REF]), the microstruc-85 tural state is kept before the tensile test at a given temperature 86 started (see Fig. 1b).

96

To distinguish the microstructural and the mechanical effects the use of true stress and logarithmic strain is needed, instead of engineering starin and stress. To distinguish the softening due to the microstructure and the one due to necking, the stresses that becomes no longer uniaxial are reassessed thanks to the Bridgman factor [START_REF] Bridgman | Stress distribution at the neck of tension specimen[END_REF] which takes into account the triaxiality. This factor requires the identification of the specimen radius (noted a) in the necking zone and the radius of the necking contour (noted R). These two radii (see Figs. 2a and2b) are measured thanks to image recording, which in turns allows to correct the strain-stress curve according to necking with the following equations:

F Bridgman = 1 (1 + 2R/a) ln (1 + a/ (2R)) (1)
Then, the Bridgman stress is calculated in such way:

σ Bridgman = F Bridgman × σ Cauchy (2)
Consequently the remaining softening visible on the strain-97 bridgman stress curves of Fig. 3 in the next section is only due 98 to microstructural effects.

99

During the tensile test, the homogeneity of the strain between the ceramic tips allows to get an accurate measurement of the strain. Nevertheless, when necking occurs, the provided strain is affected and the measurements are erroneous. To bypass this problem, the logarithmic strain of Hencky (ε log = ln L L 0 ) is corrected at the beginning of necking and the definition becomes:

ε log = ln A 0 A i = ln D 0 2 D i 2 = 2 ln D 0 D i ( 3
)
where A 0 , A i and D 0 , D i = 2a are: the initial/instantaneous 100 cross-section and the initial/instantaneous diameters.

Results

During heating (from the T6 state)

In Fig. 3 the strain-stress curves obtained at various temperatures ranging from 20 to 560 • C are presented. Note that here the effect of the thermal expansion has not been substracted. In the T6 state at room temerature a slight hardening is observed while a significant softening is observed for tempeatures ranging between 250 and 400 • C. On the light of simulations conducted in [START_REF] Bardel | Coupled precipitation 577 and yield strength modelling for non-isothermal treatments of a 6061 alu-578 minium alloy[END_REF] this effect is attributed to the dissolution/coasening of the β ′′ -β ′ hardening precipitates. Such behaviour is not commonly reported in the literature. Nevertheless, this kind of softening has already been reported in [START_REF] Zain-Ul Abdein | Thermo-Mechanical 644 Characterisation of AA 6056-T4 and Estimation of its Material Properties 645 using Genetic Algorithm[END_REF] for an AA6056 alloy in the T4 temper state (under aged condition). Note that here, the hardening precipitates are already dissolved when a temperature of T = 500 • C is reached (see results in [START_REF] Bardel | Coupled precipitation 577 and yield strength modelling for non-isothermal treatments of a 6061 alu-578 minium alloy[END_REF]). Consequently, no softening is encountered during the tensile test between T = 500 • C and 560 • C because the microstructural state is already stabilized before the mechanical sollicitation). DB

The yield strength σ y can be distinctly identified on these curves DB and its evolution is then reported in Fig. [START_REF] Develay | Propriétés de l'aluminium et des alliages d'aluminium cor-647 royés[END_REF] and [START_REF] Sharma | Effect of albite particles on the coefficient of thermal expan-649 sion behavior of the Al6061 alloy composites[END_REF], respectively.

During cooling (from the O state)

The strain-stress curves at temperatures ranging from R.T. to

Additionnal tests at T = 200 • C and T = 300 • C

In this subsection some additional tests are presented to highlight some of the previous results. cent studies of the authors on AA6061-T6 [START_REF] Bardel | Coupled precipitation 577 and yield strength modelling for non-isothermal treatments of a 6061 alu-578 minium alloy[END_REF] have shown that for this kind of transient heat treatments, the main mechanism is the decrease of the density of precipitates by pure dissolution or coarsening. Thus, a phenomenological microstructural approach composed by a hard and a soft phase is employed where the evolution of the soft one (in the initial T6 state the proportion is 0) is deduced thanks to the isokinetic model proposed by Myhr et al. [START_REF] Myhr | Process modelling applied to 6082-T6 aluminium[END_REF]. In addition to its simplicity, the chemistry of the alloy is not taken into account, it allows a relative general application for the 6xxx series.

Knowing the proportion of these two phases:

• the first represents the maximal hardening (the T6 state, here called hard phase) ;

• the other represents the minimal hardening (the O state, here called soft phase). a law of mixtures is employed in the mechanical part of the model to provide the strength of the alloy for each temperature. However, this methodology requires the knowledge of the behaviour of the hard and the soft phases at each temperature without the microstructural effect which is already provided by the isokinetic simulation [START_REF] Myhr | Process modelling applied to 6082-T6 aluminium[END_REF]. These behaviours are non measurable because the precipitation state evolves continuously during each experiment. Thereby, the purpose of this investigation is to determine, thanks to an inverse method, the stress-strain cuves of these soft and hard phases with temperature. For that, a dedicaced Finite Element (FE) model is used to reproduces the experients and to identify these curves. Note that the identification of the mechanical behaviour at room temperature of each of these phases does not require additional effort since it corresponds to what has been measured from T6 and O temper states, as quickly presented above and also in greater details in [START_REF] Maisonnette | Effects of 660 heat treatments on the metallurgy and mechanical properties of a 6061 661 aluminum alloy[END_REF].

Thermo-electrical, microstructural and mechanical simulations of the tensile tests

To best fit with experiments the simulation of tensile tests is needed. These simulations consists in two uncoupled steps:

• a electro-kinetic or thermo-electrical (TE) analysis, calibrated thanks to the thermocouple measurements, that provides as input of the mechanical analysis the heat treatment and temperature gradient in the specimen. These data are needed in order to accurately reproduce the mechanical behaviour. The resulting thermal history allows to compute, at each time step, the proportion of the soft and hard phases in the specimen.

• Then a microstructural and mechanical (MM) analysis is performed to estimate at each integration point the resulting constitutive law and then simulate the tensile test.

Finally, the resulting stress-strain curves are confronted to the experimental ones to identify the behaviour of the hard and soft phases (that are not known during a temperature treatment, just at room temperature) for each thermal treatment.

Presentation of the model 324

Electro-thermal analysis.

To reproduce the Joule heating used to heat the specimen, an electro-thermal analysis is performed with the commercial software Sysweld R [29] from the heat equation:

div λ t ---→ gradT + P = dc p ∂T ∂t (4)
where T , t, λ t (W.m -1 .K -1 ), P (W.m -3 ), d (kg.m -3 ), and c p 325 (J.kg -1 .K -1 ) are respectively the temperature, time, thermal con-326 ductivity, the energy produced in the material per unit volume, 327 the density and the specific heat.

328

In this thermal problem, the energy P corresponds to the electrical energy defined by [START_REF] Feulvarch | Resistance spot welding simulation: a general finite element formulation of electrothermal contact conditions[END_REF] (Q is a potential volume heat source added to the electrical one):

P = ---→ gradV. J + Q = ---→ gradV. σ(T ) ---→ gradV + Q (5)
where V, σ(T ) and J are the electrical potential, the electrical 329 conductivity tensor, and the current flux computed at each node 330 through the electro-kinetic equation div J = 0.

331

The equation solved by Sysweld R [29] assumes that the thermal conductivity and the specific heat (now written as function of the enthalpy H: c p = ∂H/∂T ) of each phase proportion p i may differ from one phase to another. In this case the total enthalpy and conductivity become: H = phases p i H i and λ t = phases p i λ t i . Finally, by remplacing dc p ∂T ∂t as function of d Ḣ:

d Ḣ = phases d i H i ∂p i ∂t + phases d i p i ∂H i ∂t (6)
we get the solution of this thermo-electrical problem:

332 phases d i (T )H i ∂p i ∂t + phases d i (T )p i ∂H i ∂t . . . -div phases p i λ t i (T ) ---→ gradT . . . (7)
- ---→ gradV. σ(T ) ---→ gradV -Q = 0
Calculation of phase proportion.

The thermal analyis takes into account the microstructural evolution of the alloy through the hard and the soft phase proportions, namely p 1 and p 2 (with p 1 + p 2 = 1 and p 1 = 1 for the T6 state and p 2 = 1 for the O state). The computations presented in this section aim to describe the kinetic of this process thanks to the Myhr et al. [START_REF] Myhr | Process modelling applied to 6082-T6 aluminium[END_REF] isokinetic model. Indeed, it was shown that the dissolution rate of the hard phase is defined by:

ṗ2 = n × p 2 (1-1 n ) t * (8)
with n and t * equal to:

333 n = 0, 5 -ap 2 b (9) t * = t * r 1 exp Q s nR + Q d R 1 T - 1 T r 1 ( 10
)
in these equations a, b are two constants and Q s , Q d , t * r 1 are the enthalpy of dissolution, the activation energy of the diffusion and the time needed to reach a total dissolution at the temperature T r 1 . These parameters have to be fitted on microstructural evolutions. Here, the values provided by Hirose et al. [START_REF] Hirose | Evaluation of properties in laser welds of A6061-T6 aluminium alloy[END_REF] will be chosen (cf. table 1) because they have already provided good hardness predictions in weld joint [START_REF] Hirose | Evaluation of properties in laser welds of A6061-T6 aluminium alloy[END_REF]. in the middle of the specimen where high gradients are involved and because all data will be measured there. 2800 quadrilateral linear elements (called Q4 in Sysweld R ) are employed to mesh the whole test specimen.

Several boundary conditions are needed to solve eq. 8, they are schematically drawn in Fig. 10. First of all, to compute J electrical boundary conditions are needed: σ(T ).

---→ gradV. n = j(V) (11)

V = V 0 (12)
here, V 0 is the imposed electrical potential and j(V) the imposed current density which is located at the beginning and at the end of the cylindrical area of the specimen (red arrows in Fig. 10).

The current intensity is close to that measured during the heating of the specimens during tensile tests but some adjustments are required to obtain the desired thermal loading.

Concerning the thermal boundary conditions, a temperature of 12 • C during the whole simulation is chosen in the specimen heads as during the experiments (the clamping system is water-cooled). Then, an air-specimen heat transfer coefficient is chosen equal to 50×10 -6 W.mm -2 .K -1 .

The material properties, the electrical conductivity, the specific heat, the density and the thermal conductivity (both are temperature dependant) are taken from the litterature [START_REF] Asm | Properties of Wrought Aluminium and Aluminium Alloys[END_REF][START_REF]Handbook, Metallic Materials and Elements for Aerospace Vehicle Structures[END_REF]:

ρ = 3.7×10 -8 Ω.m ; C p =

Results of the thermo-microstructural analysis

Note that heat exchange by radiation is assumed to be negligible compared to heat transfer by conduction and convection, and then is not included here for sake of simplicity. Then, in this section, two examples of microstructural results are presented: precipitates (β ′′ -β ′ ) should be dissolved [START_REF] Bardel | Coupled precipitation 577 and yield strength modelling for non-isothermal treatments of a 6061 alu-578 minium alloy[END_REF].

373

The evolution of temperature at the middle of the specimen In this analysis, the constitutive law used at each integration is determined thanks to a mixture law given by:

Y(ǫ, T ) = p 2 × Y O (ǫ, T ) + (1 -p 2 ) × Y T6 (ǫ, T ) (13)

FE simulation of electron beam girth weld 452

In this section, to evaluate the capabilities of our metallurgi-453 cal model we will confront residual stresses data measured in a 454 girth weld to the corresponding finite element simulation. Note 455 that the calibration of thermal finite element model is not de-456 scribed here because it is similar to previous work [START_REF] Zain-Ul Abdein | 600 Finite element analysis of metallurgical phase transformations in aa 6056-601 t4 and their effects upon the residual stress and distortion states of a laser 602 welded t-joint[END_REF]: only the 457 calculation steps will be recalled in the first subsection. The heat transfer analysis rests upon the solution of the classical heat equation with appropriate boundary conditions [START_REF] Bergheau | Finite Element Simulation of Heat Transfer[END_REF].

It is computed on a three-dimensional model in order to properly reproduce the thermal cycles in the heat-affected region. The mechanical analysis is based on the momentum balance equation where inertial effects are neglected. The internal heating due to plastic dissipation is neglected considering the small strain rates generated by a welding operation. As this effect as well as the influence of stresses on metallurgical transformations is neglected, the mechanical analysis can be uncoupled from the thermo-metallurgical simulation. The computation is thus achieved in a second stage using the temperatures and phase proportions calculated previously. Heat transfer and microstructural evolution are considered in the mechanical analysis through four effects:

• the thermal strains.

• The influence of temperature and of the precipitate state evolution on the constitutive behavior (hard and soft proportions are state variables).

• The stress-strain relation must be temperature-dependent, and representative of the hard and soft material mixture.

• The nonlinear transient mechanical calculation is achieved using the temperature fields computed by a 3D Eulerian steady state thermal analysis as in [START_REF] Bergheau | Finite element simulation of processes involving moving heat sources application to welding and surface treatment[END_REF]. From a numerical point of view, the 3D thermal eulerian sim-524 ulation is transformed into a temporal description on one sec-525 tion which is then used for the mechanical analysis. Indeed, 526 as the structure is similar to a tube, a section perpendicular to 527 the welding direction (Fig. 17a) is used with a relevant rota-528 tional symmetry conditions. The stresses are computed on this 529 section after the welding operation (Fig. 17b), the inner and ACCEPTED MANUSCRIPT

FE results and comparison with experiments

26 strength,

 26 are strongly dependent on the microstructure which 27 is varying with time due to the growth or dissolution of hard-28 ening precipitates. It was for example shown that: (i) the me-29 chanical properties diminishe when the maximum temperature 30 reached increases or when the heating rate decreases, (ii) all 31 β ′′ -β ′ hardening precipitates in AA6061-T6 are dissolved at 32 560 • C whatever the heating rate. However, the knowledge of 33 the mechanical properties at room temperature is not sufficient 34 to model accurately a process involving high temperatures and 35 thermal gradients encountered both in space and time during 36 welding or additive manufacturing. The knowledge of mechan-37 ical properties at temperature ranging from 20 • C and 560 • C 38 are required to simulate the welding process. This is the pur-39 pose of the present investigation. Note that the solidification or 40 the properties of the melted alloy are not studied here. 41 In this study, to be as representative as possible of the precip-42 itation state evolution encountered in the HAZ under welding 43 condition and to model the mechanical properties during weld-44 ing, the thermal loadings applied to specimens will not include 45 dwell time before the beginning of the tensile tests. These ex-46 periments include tensile tests at high temperature during either 47 heating or cooling. Results are used to propose non measurable 48 stress-strain curves at various temperatures for the T6 and O 49 states (the latter meaning that all β ′′ -β ′ hardening precipitates 50 are dissolved) for transient heat treatments. These data com-51 plete the high strain rate investigation performed by Pickett et 52

Figure 1 :

 1 Figure 1: Tensile test procedure from T6 state: (a) after heating and (b) after overheating and partial cooling.

87 2 . 2 .

 22 Post treatment 88 Material softening can be observed during tensile tests at 89 some specific temperatures. Its occurence can be explained by 90 two phenomena: 91 • the precipitation state is evolving during the test and the 92 degraded microstructure leads to less hardening. 93 • in the middle of the specimen a necking can occur and 94 produce a significant decrease of the section compared to 95 the initial one.

  101

Figure 2 :

 2 Figure 2: (a) Necking zone geometry and definition of the radii used for the calculation of the Bridgman factor, (b) example of R measurement.

Figure 3 :

 3 Figure 3: Bridgman stress versus logarithmic strain for various temperatures from T6 state (thermal expansion is not subtracted)

Figure 4 :

 4 Figure4: Evolution of (a) the yield strength σ y and (b) the Young modulus versus test temperature from T6 state. Experimental data from the literature are taken from[START_REF] Develay | Propriétés de l'aluminium et des alliages d'aluminium cor-647 royés[END_REF] and[START_REF] Sharma | Effect of albite particles on the coefficient of thermal expan-649 sion behavior of the Al6061 alloy composites[END_REF], respectively.

  560 • C are presented In Fig. 5, after the O treatment presented in Fig. 1b. These results are of interest since they represent what happens during cooling for hardening precipitate alloys. One may notice a significant hardening of the material that occurs during the tensile tests at R.T. and at 200 • C. To the authors' knowledge such data has not yet been published for AA6xxx series. The explanation is provided next.

Figure 5 :

 5 Figure 5: Bridgman stress versus logarithmic strain for various temperatures from O state (thermal expansion is not subtracted) In the O state (i.e. at T = 560 • C: all hardening precipitates are dissolved) the results are naturally similar to those reported in the previous section. However, conversely to the previous results, no softening is reported for all the other curves in the plastic domain and a large work hardening appear from T = 200 • C. Indeed, this temperature is close to the tempering one and in these conditions the 0 state rapidly evolves [2] because many hardening precipitates nucleate during the tensile test and harden the alloy throughout the test. Note that, in these tests, the fracture comes much faster: the ductility is here around 4 % compared to a range of 20-80 % for the previous tests. For these experiments, the evolution of the yield strength is plotted in Fig. 6. Between T = 560 • C and 400 • C the temperature is still high and no significant nucleation is produced: the yield strength slightly increases when the test temperature decreases. Below 400 • C, a sharp increase of the yield strength is produced thanks to the progressive formation of β ′′ -β ′ hardening precipitates during this tensile test at low strain rate. When the temperature is still decreasing from 300 • C to room temperature the nucleation rate (see [2] for details) increases as well as the yield strength. It should be noted that the yield strength, which ranges from σ y = 9.7 MPa at T = 560 • C to σ y = 62 MPa at T = 200 • C when the alloy is tested from the O temper state, is significantly lower than that measured during heating e.g. from the T6 temper state (278 MPa at 200 • C, see Fig. 3).

Figure 6 :Figure 7 :Figure 8 :

 678 Figure 6: Yield strength versus test temperature from O state (during cooling)

  209

Figure 9 :

 9 Figure 9: Strain-stress curve at T = 300 • C for various maximum pre-heating temperatures (Thermal expansion is not subtracted)The specimen which has undergone pre-heating at T Max =

217 3 . 4 .

 34 Discussion: comparison with published data 218 Usually, for this kind of age hardening alloys, tensile tests 219 are performed after a long dwell time at the test temperature 220 in order to stabilize the material microstructure. However, the 221 final goal of this work is to provide experimental data for weld-222 ing and additive manufacturing applications where the kinetics 223 are very fast. So it requires new database where no dwell time 224 are involved during the thermo-mechanical treatment. It means 225 that the microstructure can evolve significantly during the ten-226 sile test and the data are more consistent with the loadings un-227 dergone during welding or additive manufacturing. To the best 228 of the authors' knowledge, little data in the literature are repre-229 sentative of this type of processes for AA6xxx alloys. 230 To check the results presented above the evolution of yield 231 strength given in [23] after a dwell time of 30 min was con-232 fronted with our results in Fig. 4a. The differences can be 233 explained in light of recent Small Angle Neutron Scattering 234 (SANS) and thermodynamical simulations reported in [2]: 235 • when T < 250 • C the results are similar because the tem-236 perature and the dwell time are too low to significantly 237 impact the precipitation state.

238•

  above 250 • C, a microstructural evolution (coarsen-239 ing/dissolution) starts to take place at the beginning of our 240 test. Furthermore, the dwell used in [23] amplifies this 241 microstructural evolution and the lower values obtained at 242 the beginning of the tensile test are the consequence of a 243 larger degradation of the initial T6 state. 244 Concerning the Young modulus, the results given by Sharma 245 et al. [24] and our are close but not similar. These difference 246 are probably due to the chemical composition which is slightly 247 different (they studied a Al6061 alloy composites). In addition 248 the test conditions, that are not provided in their paper, differ 249 also probably.

252

  In order to get a constitutive law for thermo-mechanical load-253 ings, as for welding applications, a phenomenological model is 254 proposed in this section. The objective of this model is to re-255 produce the behaviour of this age hardening alloy, initially in 256 the T6 state, from room to solidus temperatures, especially the 257 hardening or the softening that was highlighted in the previ-258 ous section (e.g. from 200 to 400 • C). The metallurgical model 259 proposed in this section is a coupling between a microstructural 260 one and a work hardening model, it aims to predict the mechan-261 ical behaviour during a single pass welding process where the 262 hardening precipitates are dissolved. 263 To simulate the evolution of the precipitation state for the 264 AA6xxx series, multi-class precipitation models are good can-265 didates [8, 25, 26]. Recently, the authors successfully cou-266 pled this physical approach to strengthening models to pre-267 dict, at room temperature, the monotonic and cyclic behaviour 268 of this Al 6061 alloy previously subjected to complex heat 269 treatments [2, 3]. Nevertheless, the physical understand-270 ing of the precipitation-hardening coupling during thermo-271 mechanical loadings is not yet well understood. However, re-272

  896 J/(Kg.K) ; d = 2700 and d = 2617 Kg/m 3 (0 and 400 • C) ; K = 150, K = 184 and K = 173 W(m.K) (0, 350 and 540 • C).

Figure 10 :

 10 Figure 10: Schematic illustration of the electro-kinetic and thermal boundary conditions

374

  and during the heating up to 300 • C is plotted in Fig.11aalong 375 with the phase proportions of the soft and hard phases. As 376 demonstrated in the experimental section, even before the be-377 ginning of the tensile test, a very fast dissolution of the hard 378 phase occurs and keeps going during the tensile test. This re-379 sult again justifies why the heating phase is of interest and why 380 no dwell time must be used to characterize the mechanical be-381 haviour of these alloys. 382 Similar results are presented in Fig. 11b for the heating up to 383 500 • C. Here, conversely to the previous example, a complete 384 dissolution is reached before the beginning of the tensile test. 385 More precisely, the hard phase is completely dissolved when 386 492 • C is reached, it confirms that a stable state is obtained at 387 500 • C prior to the mechanical loading. It explains why there is 388 no softening in this test.

389 4 . 3 .

 43 Mechanical modeling of tensile tests 390In this section a mechanical model, which describes the 391 elasto-plastic behaviour of the hard and the soft phases at each 392 temperature, is coupled to the thermo-microstructural analysis 393 presented in the previous section.

394 4 . 3 . 1 .

 431 Description of the mechanical model395

Figure 11 :

 11 Figure 11: Evolution of temperature and phase proportion (a) at T = 300 • C (b) at T = 500 • C

400 4 . 3 . 2 .Figure 12 :

 43212 Figure 12: Comparison of tensile curves obtained by experimental and numerical investigations (a) at T = 300 • C (b) at T = 500 • C

Figure 13 :

 13 Figure 13: Constitutive laws obtained by inverse method at various temperatures (a) for the hard phase (b) and for the soft phase

442A

  thermal loading composed by a linear heating (rate r = 443 15 C/s DB ) up to a maximum T = 400 • C and followed by an 444 air cooling stage was simulated. For transient heat treatment, 445 the thermo-microstructural simulation provides a proportion of 446 hard phase with p 1 = 0.37. The resulting stress-strain curve 447 is compared in Fig. 14 with an experimental one given by a 448 previous work of the authors [28]. We can see that the general 449 trends are well reproduced by this comprehensive two phase 450 model, except in the beginning of the plastic domain.

Figure 14 :

 14 Figure 14: Confrontation between experimental and numerical curves during a tensile test performed at room temperature after a linear heating (rate r = 15 C/s DB ) up to T = 400 • C.

458 5 . 1 .

 51 Computational Welding Mechanics 459 Welding issues such as weldability, distortions and residual 460 stresses can be solved by mathematical modeling using Finite 461 Element Analysis (FEA) to compute coupled transient non-462 linear problems. This activity refers to Computational Weld-463 ing Mechanics (CWM) which has become a decision-making 464 tool for welding as well as mechanical engineers [7, 34]. The 465 interactions between heat transfer, metallurgy and mechanics 466 have to be taken into account. Strong coupling is performed 467 to solve thermo-metallurgical part and weak coupling is suffi-468 cient to predict mechanical states. Strong coupling means that 469 temperatures and phase proportions are solved in the same set 470 of equations unlike weak coupling which only needs thermometallurgical results as input data.

  The shape of the heat sources and the input energy are fitted to experimental data (i.e. recording of thermocouples and the dimensions of the weld pool and the heat affected zone obtained from macrograph transversal cuts). The heat input Q is represented by a heat flux distribution. Electron beam welding vaporizes the material and requires a specific shape for heat input distribution. A gaussian volume heat source is used to model such processes. The larger molten zones observed at the top and bottom of the plate is obtained by the mean of two gaussian surface heat sources and allow to get the shape provided by the macrographs (see fig.15).

Figure 15 :

 15 Figure 15: The calibrated molten zone in red.

508Y+Figure 16 :

 16 Figure 16: Representation of the girth weld, the line where the neutron diffraction experiments was performed is represented by an orange arrow, it is colinear to the transverse direction.

530Figure 17 :

 17 Figs. 18 and 19: a very good correlation is observed. Further-536

Figure 18 :

 18 Figure 18: Comparison between iHD experiments and numerical simulation at inner surface (X=10 mm).

Figure 19 :

 19 Figure 19: Comparison between iHD experiments and numerical simulation at outer surface (X=-10 mm).

Figure 20 :

 20 Figure 20: Comparison between neutrons diffraction (ND) experiments and numerical simulation (2 mm below the inner surface).

Table 1 :

 1 Numerical values of parameters employed in the microstructural model

	Parameters Numerical values
	a	0.394
	b	0.107
	Q s	30 kJ/mol
	Q d	130 kJ/mol
	t * r 1 T r 1	600 s 375 • C
	Mesh, properties and boundary conditions. The FE model
	used here is axisymmetrical with a fine mesh (0.16 × 0.3 mm 2 )

  The mechanical behaviour of a 6061 age hardening alloy 541 was investigated at medium and high temperature from T6 to 542 O states. To be as representative as possible of the thermo-543 mechanical loading encountered during welding, no dwell time 544 is used before the start of the tensile test. The O state can be obtained, even thanks to a transient 551 treatment, if the maximum temperature reaches at least

	540	6. Conclusions
	545
	546	The main experimental results are:
	547	• from a T6 state, a softening is obtained at intermediate
	548	temperature (from 200 to 400 • C) due to the dissolution
	549	and/or coarsening of the hardening precipitates during the
		tensile tests.

550 • 552 T = 500 • C. 553 In a second part, a thermo-electrical model is used to repro-554 duce the experimental tests from a T6 state. These FE simu-555 lations was coupled with a phenomenological microstructural-556 mechanical model proposed in this study. For this purpose, 557 the stress-strain curves corresponding to the T6 and the O state 558 (without microstructural evolution) have been identified by in-559 verse method at all temperatures. Then, thanks to a mixture law 560 and an isokinetic dissolution model, the resulting behaviour of 561 the alloy can be reproduced for any heat treatments and then 562 used for a welding thermal cycle simulation. It was shown that 563 this kind of model can provide a good estimation of residual 564 stresses for complex welding simulations.

565 7. Acknowledgements 566 The financial support of the CNRS through FEDERAMS, 567 FR2145, is gratefully acknowledged as well as AREVA for the 568 financial support. The authors wish to thank the TG7 (NeT) net-569 work and especially Robert Wimpory for ND measurements.


Exploring Green Chemistry's Contribution in Circular Economy: Transforming Territorial Challenges into Innovative Solutions

Dr. Gregory CHATEL

The significant role of green chemistry in circular economy projects will be discussed, with a particular emphasis on territorial approaches for waste and biomass valorization and the utilization of unconventional activation methods such as sonochemistry, microwave chemistry, and pressurized fluids. The presentation will showcase several examples of projects that utilize these innovative techniques for the production of valuable molecules within the context of the bioeconomy and circular economy. Emphasis will be placed on the importance of multidisciplinary collaboration and territorial strategies in achieving sustainable solutions. Several case studies will be presented, including a local-scale project that focuses on the valorization of viticultural waste through a cascade of valorization processes and various applications. Additionally, the establishment of a start-up dedicated to the valorization of invasive plants for favouring the biodiversity preservation, utilizing an innovative laboratory-developed process, will be highlighted. Furthermore, an initiative that utilizes supercritical CO2-based processes to valorize coffee grounds collected from local restaurants for the production of molecules of interest in cosmetic and nutraceutical applications will be discussed. The presentation will also delve into recent research on the valorization of apple pomace for antibacterial applications. More globally, the integration of green chemistry strategies through multidisciplinary and multi-actor approaches is essential for effectively transforming territorial challenges into innovative solutions.

  Some recent publications on the subject: M. Zwingelstein, M. Draye, C. Piot, J.-L. Besombes, G. Chatel,* ACS Sustain. Chem. Eng. 2019, 7, 8310 G. Chatel,* Chem. Eur. J. 2020, 26, 9665. A. Vandeponsee, M. Draye, C. Piot, G. Chatel,* Green Chem. 2020, 22, 8544 G. Chatel,* R. Duwald, M. Draye, P. Fanget, C. Piot, Patent WO2021250251A1, 2021. V. Quinty, C. Colas, R. Nasreddine, R. Nehmé, C. Piot, M. Draye, E. Destandau, D. Da Silva,* G. Chatel,* Plants 2023, 12, 83 Dr. Gregory CHATEL is an Associate Professor at the EDYTEM laboratory (Savoie Mont Blanc University/CNRS). He obtained his PhD degree from the University of Grenoble in 2012 under the supervision of Prof. Micheline DRAYE and Prof. Bruno ANDRIOLETTI. Afterward, he joined the research group of Prof. Robin D. ROGERS at The University of Alabama (USA) as a postdoctoral fellow to work on ionic liquids for applications in the fields of space propulsion, biomass valorization, and separation. Towards the end of 2013, he was recruited as an Assistant Professor at the Institut de Chimie des Milieux et Matériaux de Poitiers (IC2MP, Université de Poitiers/CNRS, France). Additionally, at the national level, he founded and served as the first president of the Young Chemists' network of the Société Chimique de France. In 2016, Dr. CHATEL joined the Savoie Mont Blanc University to focus on research in green chemistry for the valorization of biomass and waste, as well as the reevaluation of fine chemical processes. He successfully completed his habilitation in 2018 and currently holds the position of Associate Professor at the EDYTEM Institute (Université Savoie Mont Blanc/CNRS). In recognition of his contributions, he was awarded the silver medal of the European Young Chemists Award by the European Chemical Society (EuChemS) in 2020. In addition, Dr. CHATEL established a new Master's Degree program in Green Chemistry and Eco-Innovations at Savoie Mont Blanc University from 2021 and was honoured with the teaching-training prize from the French Chemical Society in 2023. Presently, he actively co-supervises the Circular Economy axis of the CLEE Chair sponsored by USMB Foundation.


Fabien Caroir

Frank Chanier

Dimitris Sakellariou

Fabien Paquet

Julien Bailleul

Louise Watremez

Virginie Gaullier

Agnes Maillard

Strike-slip faulting in the western prolongation of the North Anatolian Fault: the Lichades -Oreoi Channel -Skiathos Basin lineament

The North Anatolian Fault (NAF) is one of the major active structures in the Eastern Mediterranean. Its right-lateral strike-slip fault initiated in eastern Turkey 13 Ma ago. The NAF westward propagation during Neogene and Quaternary times delineates the plate boundary between Eurasia and Anatolia-Aegean. The western termination of NAF is currently located in the North Aegean Trough (NAT) where NAF displays a NE-SW direction. In the NAT, the NAF termination is located near to the Sporades Islands. In the western prolongation of this termination, there is a wide domain characterised by distributed deformation. This major extensional area is mainly constituted by the Corinth rift and the North Evia domain, our study area. The whole zone experience a relative high seismicity with strike-slip focal mechanisms, especially right-lateral displacements along NE-SW-striking faults, which are mainly located between the North Evia domain and the Southern Thessaly.

Our study is mostly based on new very-high-resolution seismic reflection profiles (Sparker) acquired during the WATER surveys (Western Aegean Tectonic Evolution and Reactivations) in July-August 2017 and 2021, onboard the R/V "Téthys II". We also analysed several seismicity catalogues in order to connect the recent structures from seismic lines to active tectonics over the region. The interpretations from these datasets emphasize the evolution of the deformation of the North Evia domain, in particular, along the NE-SW striking lineament "Lichades Area -Oreoi Channel -Skiathos Basin" (L-O-S).

The deformation in the Lichades Area is dominated by numerous active normal faults striking W-E or WNW-ESE and showing metric-scale offsets (up to 5 m.) within the Holocene sequence. One of the largest sub-active to active fault is striking NE-SW, parallel to the Oreoi Channel, and thus strongly oblique to the main rift deformation. The Oreoi Channel is a marine straight linking the Lichades Area and the Skiathos Basin. The seismic profiles highlight normal faults of different ages with a NE-SW direction. In the south-east, the Oreoi Channel is delineated by the Oreoi Fault, a mainly onshore normal fault which is dipping towards north-west. The Skiathos Basin is a newly discovered structure from our seismic dataset that is separated from the Skopelos Basin by a NE-SW striking acoustic ridge. The Skiathos Basin presents two main depocenters individualized by areas of rising acoustic basement. Some normal faults, oriented NE-SW and W-E, have been identified in the basin. Finally, many earthquakes focal mechanisms located in the Skiathos Basin and the Oreoi Channel indicate strike-slip faulting, with a right-lateral motion along the NE-SW direction.

This detailed structural analysis together with the synthesis of seismic activity allow to propose a tectonic map with new insights on the recent deformation of the key-area "L-O-S" in the south-western prolongation of NAF. The Skiathos basin development shows indications of transtensive deformation. The Oreoi Channel is controlled by NE-SW-striking faults with a rightlateral component and the Lichades Area displays several fault with oblique direction and pure extension. We propose that the L-O-S tectonic system prolongs the NAF system and may progressively evolve as the future plate boundary.

Powered by TCPDF (www.tcpdf.org)