Y Li

C Robertson

M Shukeir

L Dupuy

Screw dislocation interaction with irradiation defect-loops in α-iron: evaluation of cross-slip effect using dislocation dynamics simulations

Keywords: composite screw dislocation source, radiation-induced defects, cross-slip segment, dislocation dynamics simulation

Plastic strain spreading in post-irradiated ferritic steels usually takes the form of wavy shear bands, where the mobile dislocations randomly interact with the radiation-induced, dispersed defect populations. In these conditions, it is believed that the dislocation/defect interactions are significantly affected by the ubiquitous, stochastic cross-slip mechanism.

The cross-slip effect is assessed by means of specific three-dimensional nodal DD simulation setups, consistent with available experimental evidence of cross-slip activity in post-irradiation straining conditions, at room temperature. A significant, cross-slip induced reduction of the effective dislocation/loop interaction strength is thereby evidenced, for various applied stress conditions and dislocation configurations. The corresponding, local interaction mechanisms are consistent with grain scale DD simulations results, in terms of post-irradiation plastic strain spreading.

Introduction

Ferritic steels are widely used as structural nuclear materials, thereby subjected to neutron irradiation-induced degradation, including hardening and embrittlement [START_REF] Bloom | The challenge of developing structural materials for fusion power systems[END_REF]- [START_REF] Matijasevic | Characterization of irradiated single crystals of Fe and Fe-15Cr[END_REF]. These detrimental evolutions are concurrent with and generally ascribed to the gradual accumulation of dispersed defect cluster populations, in the form of sessile (immobile) dislocation loops. The defect dispersion characteristics, size and number density, depend on many different factors such as the irradiation temperature, cumulated dose and the material chemical composition [START_REF] Nikolaev | Radiation embrittlement of lowalloy steels[END_REF]- [START_REF] Victorian | The microstructure and associated tensile properties of irradiated fcc and bcc metals[END_REF]. Mobile dislocations generated during post-irradiation straining strongly interact with these loop/defect populations [START_REF] Zinkle | Observation and analysis of defect cluster production and interactions with dislocations[END_REF]- [START_REF] Robach | In-situ transmission electron microscopy observations and molecular dynamics simulations of dislocation-defect interactions in ion-irradiated copper[END_REF]. At the grain scale, plastic strain spreading takes the form of wavy shear bands, controlling the subsequent stress-strain and fracture toughness responses [START_REF] Wei | Evolution and microstructure of shear bands in nanostructured Fe[END_REF]- [START_REF] Chaussidon | Dislocation dynamics simulations of plasticity in Fe laths at low temperature[END_REF]. Understanding the shear-band scale plasticity mechanisms is thus a crucial factor, in the management of nuclear structural material lifetime.

Dislocation/loop interactions have been recently investigated using atomistic Molecular Dynamics (MD) [START_REF] Liu | Molecular dynamics simulations of the interactions between screw dislocations and self-interstitial clusters in body-centered cubic Fe[END_REF]-Erreur ! Source du renvoi introuvable. or mesoscopic Dislocation Dynamics (DD) simulations [START_REF] Shi | Interaction of <100> dislocation loops with dislocations studied by dislocation dynamics in α-iron[END_REF]- [START_REF] Drouet | A direct comparison between in-situ transmission electron microscopy observations and Dislocation Dynamics simulations of interaction between dislocation and irradiation induced loop in a zirconium alloy[END_REF]. These studies usually assume periodic boundary conditions, i.e. infinitely long mobile dislocations. Shear band development also involves the ubiquitous cross-slip mechanism, which possibly have a strong influence on the dislocation/loop interaction, with finite length (screw) dislocations [START_REF] Shin | Dislocation-impenetrable precipitate interaction: a three-dimensional discrete dislocation dynamics analysis[END_REF]- [START_REF] Kelly | Strengthening method in crystals[END_REF]. Cross-slip is a time-dependent, stochastic phenomenon [START_REF] Chaussidon | The glide of screw dislocations in bcc Fe: Atomistic static and dynamic simulations[END_REF]- [START_REF] Bulatov | Computer simulations of dislocations[END_REF] and for this reason, its specific contribution to the effective radiation-induced loop strength has not yet been measured or evaluated. Loop interaction strength assessment is attempted hereinafter, by means of DD simulations. This technique allows implementing a specific initial configuration, associated with a (screw) dislocation glide plane change, due to a single and well defined cross-slip event. That initial configuration is called a «composite dislocation source» (please refer to Fig. 1), which has been developed based on TEM observation of cross-slip activity, in poststrained ferritic steel [START_REF] Robertson | Experimental analysis of the plastic behavior of ionirradiated bainitic RPV steel[END_REF]- [START_REF] Robertson | Dislocation structures in 16MND5 pressure vessel steel strained in uniaxial tension at different temperatures from -196°C up to 25°C[END_REF]. Such evaluation is generally not accessible to any other known investigation method, including experiments. For instance, crystal plasticity cannot treat cross-slip at the scale of individual dislocation/defect interaction, unless informed by other, lower scale simulation methods. The typical time step of atomistic (molecular dynamics) simulations is generally too brief however, to generate a composite source configuration within a reasonable simulation timeframe.

The following investigation approach is therefore adopted, based on the simulation method and setups as described in Section 2. Simulations using periodic boundary conditions (without pining points) are carried out first, as a benchmarking case (Section 3.1.1: loop case; Section 3.2.1: loop case). These results allow validating the adopted simulation parameters and setup, by comparison with well-established MD simulation results. The role of cross-slip is evaluated next, using DD simulations with "composite dislocation" sources. The results are compared with the preliminary case results: Section

Simulation method and setups

Dislocation Dynamics simulation setup and model: benchmarking case

All DD simulations results presented in this work are performed using a 3D nodal code called NUMODIS (e.g. [START_REF] Shi | Interaction of <100> dislocation loops with dislocations studied by dislocation dynamics in α-iron[END_REF], [START_REF] Drouet | Dislocation dynamics simulations of interactions between gliding dislocations and radiation induced prismatic loops in zirconium[END_REF]- [START_REF] Drouet | A direct comparison between in-situ transmission electron microscopy observations and Dislocation Dynamics simulations of interaction between dislocation and irradiation induced loop in a zirconium alloy[END_REF]), developed in CEA. The dislocation lines are described by a series of inter-connected nodes. Computation of the internal elastic stress and corresponding nodal force is carried out within the frame of the non-singular continuum elastic theory [START_REF] Cai | A non-singular continuum theory of dislocations[END_REF]. Nodal velocity is proportional to the effective resolved shear stress, as per:

eff s s b v B   (1)
where τ eff is the effective resolved shear stress, b s is the dislocation Burgers vector and B s the viscous drag coefficient, characterizing the phonon scattering effect in Fe at room temperature [START_REF] Alshits | Mechanisms of dislocation drag[END_REF]- [START_REF] Urabe | Diloscation mobility in potassium and iron single crystals[END_REF], [START_REF] Po | A phenomenological dislocation mobility law for bcc metals[END_REF].

Typical ferritic steels microstructures include 1-2 µm thick sub-laths (or platelets) crystallites, formed during the first steps of the material elaboration [START_REF] Robertson | Dislocation structures in 16MND5 pressure vessel steel strained in uniaxial tension at different temperatures from -196°C up to 25°C[END_REF], [START_REF] Renevey | Ductile -brittle transition of ferritic steels modelled by the local approach to fracture[END_REF]. The resulting material is then subjected to tempering at intermediate temperature, which aims at relieving the internal stress and removing the impurities remaining the lath matrix, as these tend to diffuse towards the lath interfaces. The remaining impurity contain in lath crystal is then minimal and the latter can be regarded as pure Fe, as far as dislocation motion is concerned.

The material parameters corresponding to pure Fe, hereby used a model material, are listed in Tab. 1 below. Tab. 1. Pure Fe materials parameters at 300K [START_REF] Shi | Interaction of <100> dislocation loops with dislocations studied by dislocation dynamics in α-iron[END_REF], [START_REF] Alshits | Mechanisms of dislocation drag[END_REF]- [START_REF] Urabe | Diloscation mobility in potassium and iron single crystals[END_REF].

Viscous

As a first investigation step, preliminary DD simulation results are benchmarked by comparison with MD simulations using a similar configuration as shown in Fig. 1. The simulated crystal orientations as X, Y and Z axis are parallel to the , and directions, respectively. The DD simulated volume dimensions are: L X = 400 nm, L Y = 300 nm and L Z = 400 nm, which is consistent with the shear bands thickness observed in postirradiated materials [START_REF] Odette | Modeling the multiscale mechanics of flow localization-ductility loss in irradiation damaged bcc alloys[END_REF], [START_REF] Gururaj | Channel formation and multiplication in irradiated FCC metals: a 3D dislocation dynamics investigation[END_REF]. One screw dislocation source, with its Burgers vector b parallel to the Z direction, is placed at the center of the simulation volume. The total length of the source is L, which is comparable to the dimensions of the simulation volume. In all the cases, one loop/obstacle is placed at a short distance from the mobile dislocation source, as discussed in the next section and shown in Fig. 1 below. The center of the loop is coplanar with the dislocation line and L 2 = L/4 from point B (at the center of segment BC). The chosen defect position (typically 15 nm or less) ensures early contact between the dislocation and the immobile defect/loop. Small variations of L 1 has a minor effect on the results since L 1 << L, while the influence of L 2 will be discussed in Section 3. The right-hand part of Fig. 2 shows the two segments AC and BC forming the segment BC is L p long and glides in the primary slip system. The loop information is same as former case with L 2 = L p /2.

Results and discussion

Interaction with loop

Interaction mechanism and obstacle strength: planar dislocation source

DD simulation results obtained using periodic boundary conditions are compared with MD simulation results. In this case, controlled strain rate loading conditions (along Z direction) are used, where = 10 6 s -1 . Interaction with the screw dislocation changes the loop Burgers vector from to b = , after the interaction is completed and the mobile dislocation breaks away from the loop (not shown). The interaction strength corresponding to this mechanism is τ c = 0.4µb/(L-D), in good agreement with corresponding MD results (i.e. τ c = 0.38 µb/(L-D) from [START_REF] Liu | Molecular dynamics simulations of the interactions between screw dislocations and self-interstitial clusters in body-centered cubic Fe[END_REF]). The next simulation case is carried out using exactly the same strain rate as before, this time using a finite-length dislocation source (L = 300 nm),

where nodes A and B are pinned (c.f. Fig. 2) and both segments AB and BC glide in the same primary slip plane ( ). The resulting interaction mechanism is presented in Fig. 3.

The screw dislocation is initially attracted by the loop and reacts with segment «2» (of the loop) to form a junction (Fig. ). This particular reaction is embedded in the NUMODIS DD code, based on the interaction mechanism predicted by atomistic MD simulations, in the selected case studies [START_REF] Terentyev | Reactions between a ½ <111> screw dislocation and <100> interstitial dislocation loops in alpha-iron modelled at atomic scale[END_REF]- [START_REF] Terentyev | Simulation of the interaction between an edge dislocation and a <100> interstitial dislocation loop in α-iron[END_REF]. Thereafter, the new segments «3», «4» and the initial screw dislocation segments rearrange as shown in Figs. This configurational change also affects the effective loop interaction strength as shown in Fig. 4 for different source cases. The critical loop strength evolution obtained from these simulations can be described using the following expression:

1 () extra c eff LT eff LT Gb LD           ( 2
)
where G is the shear modulus (see Tab.

Interaction mechanism and effective obstacle strength: composite dislocation source

The composite dislocation source case is systematically investigated with four different simulation sets, using different constant applied stress (τ p , τ cs ) conditions and finite L p long dislocation segments (see Tab. 3 and caption), gliding in the primary slip plane. contacts the obstacle at an early stage of each simulated case.

One typical dislocation-loop interaction case is presented in Fig. 5, for loading conditions τ p = τ cs = 120 MPa of set 3. Segment BC interacts with the obstacle in Fig. 5 As node C contacts with the loop/obstacle, a mutual attraction occurs between segments AC and BC then segment BC gets past the loop. Soon after, the whole dislocation source is transferred into the cross-slip plane (Fig. 5(d)). The loop is released at this time, while its Burgers vector returns to . Simulation results associated with in Tab. 3 cases are presented in Fig. 7 in the form of triplet number series (τ p , τ cs , S). The S = 0 (deep green) case indicates that the source is blocked by the obstacle; the S = 1 (purple) case indicates the source overcomes the defect while gliding in the primary slip pane; the S = 2 (white) case indicates the source overcomes the defect while gliding in the cross-slip plane (as in Fig. 5). Fig. 7 shows that a screw dislocation can directly cut through the obstacle provided τ p > 150 MPa which is consistent with Fig. strain rate jump at t = 0.12 ns. The second peak appears as the source is entirely transferred into the primary slip plane. In the L p = 300 nm case (Fig. 8(b) curve A), the dislocation velocity is faster due to the lack of the competition between segments AC and BC.

Similarly, in τ cs > τ p case (Fig. 8(b), curve D), loop by-passing mechanism occurs with a slower strain evolution. The presence of the cross-slipped segment AC (for the case of L cs >1/3L) systematically helps the primary segment BC to get past the obstacle, including for τ p and τ cs levels below the critical obstacle strength  (from Fig. 4:   150 MPa). This effect reduces with the decrease of the cross-slip segment length L cs .

In order to evaluate the separate contribution of cross-slip on the dislocation-loop interaction strength, we finally replaced the loop with a hard, impenetrable platelet (or facet). The facet position, size and orientation are exactly the same as those of the loop. The studied case corresponds to L p = 150 nm with τ p = τ cs = 120 MPa. The reader should note that the implemented facet does not generate any long-range stress. The total strain and the corresponding strain rate evolutions versus the simulation time are shown in Fig. 9. In the loop case, the entire dislocation-loop interaction time is ~ 1 ns. The total reaction time is nearly the same in the facet case (~10 ps shift), with a relative total strain error of about 1.78%, at the reaction completion time. This demonstrates that the interaction with a loop is mainly controlled by the cross-slip mechanism. The reaction map (not shown) corresponding to the hard facet case is exactly the same as that of the loop case (c.f. Fig. 6).

Interaction with loop 3.2.1 Interaction mechanism and obstacle strength: coplanar dislocation source

The screw dislocation without pinning points and periodic boundary condition case is first examined, using fixed strain rate conditions as in section 3.1.1. In this case, the sessile loop is absorbed in the form of a helical turn [START_REF] Terentyev | Reactions between a ½ <111> screw dislocation and <100> interstitial dislocation loops in alpha-iron modelled at atomic scale[END_REF] which then closes itself and leaves a loop is behind, as the screw dislocation breaks away. This mechanism is associated with a critical interaction stress τ c = 0.72µb/(L-D) in agreement with [START_REF] Terentyev | Reactions between a ½ <111> screw dislocation and <100> interstitial dislocation loops in alpha-iron modelled at atomic scale[END_REF].

The coplanar finite-length (pinned) dislocation source case is examined next (see Fig. 10).

A helical turn is formed during the first stages of the interaction (Fig. 10(b)), then after a significant bow-out of the dislocation, the helical turn reconnects and re-emits the initial loop in a process similar to Hirsch's mechanism [START_REF] Hirsch | Proceedings of a conference on 'point defect behavior and diffusional processes[END_REF] (Fig. 10

Interaction mechanism and obstacle strength: composite dislocation source

The results associated with L cs = 150 nm case are presented in Fig. 11. In Fig. 11 total strain rate strongly depends on the τ cs level.

Lastly, a simulation case where the loop is replaced by a oriented hard facet (not shown) is carried out. It is recalled that unlike the loop, the facet has no associated stress field; whereas interaction with a facet involves none of the dislocation recombination mechanisms associated with loop interaction. In this way, we found that interaction with the facet is similar to interaction with a loop, as presented in Section 3.1.2. This comparison further confirms that, in presence of cross-slip, the effective loop/obstacle interaction strength weakly depends on the loop-induced elastic stress field or the particular dislocation-loop interaction mechanism.

General discussion

The interaction mechanisms reported in Sections 3.1.1 and 3.2.1 involve un-dissociated dislocations and strictly no cross-slip contribution. In these conditions, it can generally be assumed that these DD and MD evolutions are strain-rate independent and therefore, applicable to much slower strain rate conditions [START_REF] Terentyev | Reactions between a ½ <111> screw dislocation and <100> interstitial dislocation loops in alpha-iron modelled at atomic scale[END_REF]- [START_REF] Shi | Interaction of <100> dislocation loops with dislocations studied by dislocation dynamics in α-iron[END_REF], [START_REF] Nogaret | Atomic-scale plasticity in the presence of Frank loops[END_REF]. The results reported in Sections 3.1.2 and 3.2.2 are obtained under constant applied stress conditions, however. In the composite-source case studies, it is important to note that the simulation time originates immediately after the glide plane change of one of the incoming, mobile dislocation arms.

In usual straining conditions, this glide plane change would take place after a definite dwell time (associated with the cross-slip mechanism) which is not counted in the total reaction time reported in Figs. 8 and 11(b). Sections 3.1.2 and 3.2.2 results thus focus on the strainrate independent steps of the interaction, which may be applicable to much slower strain rate conditions as well.

For instance, disperse defects were implemented in massive, grain scale calculations, based on high strain rate MD calculation results [START_REF] Terentyev | Reactions between a ½ <111> screw dislocation and <100> interstitial dislocation loops in alpha-iron modelled at atomic scale[END_REF]- [START_REF] Terentyev | Simulation of the interaction between an edge dislocation and a <100> interstitial dislocation loop in α-iron[END_REF]. Grain-scale simulations were then carried out at strain rates up to 5 orders of magnitude slower. The resulting dislocation microstructures and corresponding stress-strain response are, nonetheless, fully consistent with available experimental evidence [START_REF] Gururaj | Channel formation and multiplication in irradiated FCC metals: a 3D dislocation dynamics investigation[END_REF], [START_REF] Gururaj | Post-irradiation plastic deformation in bcc Fe grains investigated by means of 3D dislocation dynamics simulations[END_REF]- [START_REF] Li | Irradiation defect dispersions and effective dislocation mobility in strained ferritic grains: A statistical analysis based on 3D dislocation dynamics simulations[END_REF]. This situation can be rationalized by noting that most of the simulation time (or the actual, experimental time) is then spent waiting between the different cross-slip events. It is also important to note that strain reported strain rates strongly depends on the mobile dislocation density, according to the simulated space sizes. The situation is significantly different in massive DD calculations, where the mobile dislocation density is limited by interaction with various obstacles, thus limiting the mobile dislocation densities [START_REF] Arsenlis | A dislocation dynamics study of the transition from homogeneous to heterogeneous deformation in irradiated body-centered cubic iron[END_REF].

As a concluding remark, all the results presented in Section 3 were obtained using the specific, materials-dependent parameters of Section 2. It can be shown, however, that crossslip mechanism affects dislocation/loop interactions in a wide range of BCC alloys (see for example: [START_REF] Robertson | Experimental analysis of the plastic behavior of ionirradiated bainitic RPV steel[END_REF]- [START_REF] Robertson | Dislocation structures in 16MND5 pressure vessel steel strained in uniaxial tension at different temperatures from -196°C up to 25°C[END_REF], [START_REF] Farrell | Mapping flow localization process in deformation of irradiated reactor structural alloys-final report[END_REF]), corresponding to different sets of physical, material parameters (see Tab. 2). The reported cross-slip effect is in any case active and mainly depends on the sub-grain stress landscape, acting on the incoming (screw) dislocation arms [START_REF] Gururaj | Channel formation and multiplication in irradiated FCC metals: a 3D dislocation dynamics investigation[END_REF], [START_REF] Gururaj | Post-irradiation plastic deformation in bcc Fe grains investigated by means of 3D dislocation dynamics simulations[END_REF]- [START_REF] Li | Irradiation defect dispersions and effective dislocation mobility in strained ferritic grains: A statistical analysis based on 3D dislocation dynamics simulations[END_REF].

Conclusions

Interaction between screw-type dislocation sources with and loops is investigated using 3D nodal dislocation dynamics simulations. The comparative interaction strength levels associated with and loops are evaluated using co-planar source cases first, where all the initial source segments glide in the same primary slip plane.

Coplanar sources are used: I) infinitely long dislocation segments, due to the periodic boundary conditions, II) in the form of a finite-length, pinned dislocation segments. Case-I is adopted as a benchmarking case, for validating our DD simulation model and setup by comparison with well-established MD simulations results.

1. Interaction strength associated with or loop is significantly larger in finite-length source case (case-II) than in periodic boundary conditions case (case-I). Pinned source nodes induce a local dislocation curvature and associated extra line tension contribution, adding up to the total effective interaction strength.

The case of composite dislocation sources is further investigated. This configuration includes two distinct (L p , L cs ) long segments, gliding in the primary and cross-slip planes, respectively. The effect of various loading conditions (τ p , τ cs ) on the effective interaction strength is examined, in terms of interaction mechanisms and time evolution of the strain rate.

2. The presence of a cross-slipped segment L cs could systematically reduce the resolved shear stress needed to unpin the screw dislocation if L cs >1/3L. The interaction strength level directly depends on the cross-slip segment length L cs .

3. The cross-slip effect is dominant regardless of the particular loop type involved ( or ), i.e. regardless of the particular loop-induced interaction mechanism and loop-induced stress field.

In conclusion, the present results indicate that cross-slip is possibly the dominant strain rate limiting mechanism, in presence of disperse loop populations. The corresponding local interaction mechanisms are consistent with grain scale DD simulations results, in terms of post-irradiation plastic strain spreading.

Fig. 1 . 2 . 2

 122 Fig. 1. DD simulation volume adapted to dislocation-loop interaction investigation. (a) This

Fig. 2 .

 2 Fig. 2. Composite dislocation source configuration. Left-hand sketch: total source length L

  3(b)), according to Frank's rule (

Fig. 3 .

 3 Fig. 3(c) highlights the effect of using periodic boundary condition (Case I) on the

  1), b is the Burgers vector modulus,  = 0.4 and  = 0.42. The correction term  represents the extra line tension contribution associated with the difference in local dislocation curvature (see Fig.3(c)), due to the pinning points of the finite length source[START_REF] Foreman | The bowing of a dislocation segment[END_REF]-[START_REF] Hull | Introduction to dislocations[END_REF]. It is important to note that the screw dislocation is systematically released before adopting a semi-circular bowed-out configuration, while the loop is mostly immobile (with respect to the dislocation). As a result, the line tension correction and the loop strength  (150 MPa) are both nearly constant, regardless of L 2 (or L p ), within the [L/6, L/2] range (see also Fig.7data, for different L p and therefore, L 2 values).

Fig. 4 .

 4 Fig. 4. loop strength evolution with reciprocal of dislocation source length 1/(L-D).

  (b) and a junction is formed at the point of contact. Segment BC is then blocked since the resolved shear stress τ p is lower than the critical obstacle strength (from Fig. 4:      150 MPa). Meanwhile, segment AC propagates in the cross-slip plane and its length L cs increases as node C moves along the line direction towards node B (Fig. 5(c)).

Fig. 5 .Fig. 6 .Fig. 7 .

 567 Fig. 5. Interaction between a composite dislocation source and a loop. The source

Fig. 8 .

 8 Fig. 8. Total strain rate evolutions versus time associated with simulation setup 3 (L cs = L p =

Fig. 9 .

 9 Fig. 9. Time evolutions of the total strain (upper frame) and corresponding strain rate

Fig. 10 .

 10 Fig. 10. Interaction between a coplanar pinned dislocation source and a loop. The

Fig. 11 .

 11 Fig. 11(a) (S=3, black color area), where the dislocation bow-out keeps gliding without the

  , τ cs ). Once the dislocation reaches the simulation volume boundary or t = t max a new combination (τ p , τ cs ) is generated according to the selected stress increment. The initial loading stress is set to 100 MPa (close to the obstacle strength reported in Section 3.1.1), so segment BC (gliding in primary slip plane)

				τ p,min	τ p,max	τ p,inc	τ cs,min	τ cs,max	τ cs,inc
		L (nm) L p (nm)
				(MPa)	(MPa)	(MPa)	(MPa)	(MPa)	(MPa)
	Set 1	300	250	100	200	20	100	200	20
	Set 2	300	200	100	200	20	100	200	20
	Set 3	300	150	100	200	20	100	200	20
	Set 4	300	100	100	200	20	100	200	20
	Tab. 3. Simulation sets including a composite dislocation source. Each set corresponds to
	several different simulations, where (τ

p , τ cs ) vary by τ p,inc and τ cs,inc steps of 20 MPa. The selected τ min and τ max values are not arbitrary: τ min exceeds the critical source activation stress, which ensures that the mobile dislocation moves towards the immobile loop; whereas τ max exceeds the dislocation/loop breakaway stress, which ensures finding the critical interaction stress, for the different simulation setups.

Each set is tested for different τ cs levels acting on segment AC and varying from (τ cs,min :τ cs,max ) and likewise, τ p levels acting on segment BC varies from (τ p,min : τ p,max ). Each simulation case is carried out up to a specific simulation time, t max (typically, several nanoseconds) under a specific loading combination (τ p

Acknowledgements

This work has been carried out within the framework of the EUROfusion Consortium and has received funding from the Euratom research and training programme 2014-2018 under grant agreement No 633053. The views and opinions expressed herein do not necessarily reflect those of the European Commission. The complementary support of the Materials Research Program RMATE of the Nuclear Energy Division of the French Atomic Energy Commission (CEA/DEN) is also acknowledged.


Comparison of classic trace element geochemistry, soil partial leaches, on soils and biogeochemistry in the Western Europe Environment. An example from Sb deposits

In Europe, exploration geochemistry is mainly restricted to classical soils geochemical analyses based on total digestion of samples, despites the good results obtained in other areas with approaches related to selective and partial leaches, in particular to reveal anomalies linked to deeply buried deposits. Similarly, examples of biogeochemistry usages in Europe remain anecdotic, but it is now very efficient to help in focusing exploration efforts with the actual level of element detection reached in a large spectrum of plants.

Within the framework of the UpDeep project (EIT Raw Materials), which aims at developing the geochemical expertise on deep buried exploration into Europe, we performed a comparison of some of these new geochemical methods on the Vendée antimony district (France).

One of the challenging questions of exploration geochemistry in Western Europe is the long-term anthropogenic contamination related to past mining activities but also agriculture and forestry, which can strongly affect primary geochemical anomalies.

Within the scope of testing different well-documented approaches, we collected around 160 samples of soils on which we performed separately the following treatment: deionized water leach, sodium pyrophosphate leach, and Aqua Regia digestion on Ah horizon; and hot hydroxylamine leach, Ionic LeachTM (ALS) and Aqua Regia digestion on B horizon. Each soil samples have been analyzed by pXRF.

We also collected 184 samples of oak and 94 samples of bramble twigs in order to compare their geochemical signal. In this Western European context, identification of a homogenously distributed vegetation media represents a complex issue.