Thierry Baudin

François Brisset

Anatoliy Zavdoveev

Hiba Azzeddine
email: hiba.azzeddine@univ-msila.dz

Pr Hiba Azzeddine

Microstructure and texture characterization in friction stir lap welded TIMETAL 21S

Keywords: Friction stir welding, Microstructure, Texture, TIMETAL 21S, Titanium

The evolution of microstructure and texture during friction stir lap welding (FSLW) of TIMETAL 21S (β-type Ti-15Mo-3Nb-3Al-0.2Si, wt.%) sheets were investigated through electron backscatter diffraction (EBSD). Excellent grain refinement is obtained through stir zone (SZ) thickness (1.2-1.8 µm). The microstructure of the thermomechanically affected zone (TMAZ) is characterized by elongated deformed grains surrounded by small recrystallized grains indicating the occurrence of discontinuous dynamic recrystallization (DDRX). The microstructure of heat affected zone (HAZ) is quite similar to the base metal (BM). The texture transformed from weak rolling-recrystallization texture in BM and HAZ to a typical shear texture with the domination of D 1 or D 2 components in the SZ and TMAZ area. A net shear texture gradient is formed across the SZ thickness which is connected with the heterogeneity of deformation. It is believed that the concomitant occurrence of grain size, dislocation and texture strengthening is responsible for the mechanical property distribution in different parts of FSLW joint.

Introduction

Friction-stir welding (FSW) is a well-known solid-state, hot-shear joining process using a non-consumable tool invented and patented by the welding institute (TWI) in 1991 [START_REF] Besharati Givi | 1 -General introduction[END_REF].

Nowadays, FSW is used in welding applications for a variety of metals and alloys with low melting temperature such as Aluminum [START_REF] Singh | Recent research progress in solid state friction-stir welding of aluminium-magnesium alloys: a critical review[END_REF][START_REF] Isa | Recent research progress in friction stir welding of aluminium and copper dissimilar joint: a review[END_REF], and Magnesium [START_REF] Singh | Review on friction stir welding of magnesium alloys[END_REF] and also for high melting temperature like Iron [START_REF] Mironov | Microstructural evolution during friction stirprocessing of pure iron[END_REF] and Titanium [START_REF] Gangwar | Friction stir welding of titanium alloys: A review[END_REF]. The principle of FSW is very simple consisting of a tool, typically formed of a pin and shoulder, is rotated and plunged into two workpieces and then translated in a specific direction. The deformation and frictional heat induced by the rotating tool ensure the bonding between the two workpieces by the formation of a solid-state weld [START_REF] Kumar | The role of friction stir welding tool on material flow and weld formation[END_REF].

Different welding joint types can be produced by FSW such as butt joint, lap joint, edge joint, and T-joint depending on the way the pieces of metal are put together or aligned with each other [START_REF] Mishra | Friction stir welding and processing[END_REF]. The butt and lap joints welding are the most common type used in the sheet welding industry. In the butt joint, two pieces of metal are placed together in the same plane, and the side of each metal is joined by welding. While in the lap joint, the welding is produced on two pieces of metal placed on top of each other [START_REF] Mishra | Friction stir welding and processing[END_REF]. The friction stir lap welding (FSLW) is suitable for joining thin sheet structures [START_REF] Mishra | Friction stir welding and processing[END_REF].

Based on thermomechanical history, the cross-section of the welded material is usually separated into four zones, namely base metal (BM), thermomechanically affected zone (TMAZ), heat affected zone (HAZ), and stir zone (SZ) [START_REF] Kumar | The role of friction stir welding tool on material flow and weld formation[END_REF]. The SZ undergoes severe plastic deformation and high temperature after friction stir which assure the solid-state connection.

TMAZ and HAZ are transition zones between the SZ and BM, in which the material is affected by a low amount of deformation and high temperature from the welding. Consequently, each zone exhibits a specific microstructural and textural evolution. Besides, the microstructure and texture are heterogeneously distributed through the SZ because of the effect of the FSW tool, on the strain and the temperature distribution across the SZ [START_REF] Yoon | Effect of rotation rate on microstructure and texture evolution during friction stir welding of Ti-6Al-4V plates[END_REF][START_REF] Yoon | Microstructure and texture distribution of Ti-6Al-4V alloy joints friction stir welded below β-transus temperature[END_REF][START_REF] Han | Impact of friction stir welding on recrystallization of oxide dispersion strengthened ferritic steel[END_REF][START_REF] Ahmed | Effect of Tool Geometry and Heat Input on the Hardness, Grain Structure, and Crystallographic Texture of Thick-Section Friction Stir-Welded Aluminium[END_REF].

It must be noted that the inhomogeneous distribution of microstructure affects directly the mechanical properties of the weld. Nevertheless, FSW is considered as severe plastic deformation (SPD) processing that produces significant ultra-fine grains through dynamic recrystallization (DRX) [START_REF] Ma | Superplastic deformation mechanism of an ultrafinegrained aluminum alloy produced by friction stir processing[END_REF]. The degree of grain refinement during FSW is strongly related to the deformation temperature and strain velocity [START_REF] Reynolds | Texture in friction stir welds of Timetal 21S[END_REF][START_REF] Zhang | Improving performance of friction stir welded AZ31/AM60 dissimilar joint by adjusting texture distribution and microstructure[END_REF][START_REF] Kumar | Effect of FSW process parameters on mechanical properties and microstructure of dissimilar welded joints of AA2024 and AA6082[END_REF]. Small grain size can be obtained through a combination of low deformation temperature and high strain rate [START_REF] Reynolds | Texture in friction stir welds of Timetal 21S[END_REF][START_REF] Zhang | Improving performance of friction stir welded AZ31/AM60 dissimilar joint by adjusting texture distribution and microstructure[END_REF][START_REF] Kumar | Effect of FSW process parameters on mechanical properties and microstructure of dissimilar welded joints of AA2024 and AA6082[END_REF].

Extensive investigations were performed during the last decade on FSW of α-type Ti alloys with hexagonal close-packed (HCP) crystal structure such as Ti-6Al-4V (wt.%) alloy [START_REF] Yoon | Effect of rotation rate on microstructure and texture evolution during friction stir welding of Ti-6Al-4V plates[END_REF][START_REF] Yoon | Microstructure and texture distribution of Ti-6Al-4V alloy joints friction stir welded below β-transus temperature[END_REF][START_REF] Mironov | Development of grain structure in β-phase field during friction stir welding of Ti-6Al-4V alloy[END_REF][START_REF] Edwards | Effect of process conditions on superplastic forming behaviour in Ti-6Al-4V friction stir welds[END_REF][START_REF] Liu | Microstructural characteristics and mechanical properties of friction stir welded joints of Ti-6Al-4V titanium alloy[END_REF][START_REF] Wu | Achieving superior superplasticity from lamellar microstructure of a nugget in a friction-stir-welded Ti-6Al-4V joint[END_REF][START_REF] Yoon | Effect of initial microstructure on Ti-6Al-4V joint by friction stir welding[END_REF][START_REF] Edwards | Material flow during friction stir welding of Ti-6Al-4V[END_REF][START_REF] Kim | α/β phase transformation and dynamic recrystallization induced microstructure development in fine-grained Ti-6Al-4V friction stir weld[END_REF][START_REF] Jia | Effect of static spheroidization on superplasticity of fine lamellae in nugget of a friction stir welded Ti-6Al-4V joint[END_REF]. Some studies were focused on the impact of FSW tool parameters such as the rotational and welding speed values on the microstructural evolution and the mechanical properties of the welded joints [START_REF] Yoon | Effect of rotation rate on microstructure and texture evolution during friction stir welding of Ti-6Al-4V plates[END_REF]. Other investigations focused on improving the superplasticity of the alloy through grain refinement induced by FSW [START_REF] Edwards | Effect of process conditions on superplastic forming behaviour in Ti-6Al-4V friction stir welds[END_REF][START_REF] Wu | Achieving superior superplasticity from lamellar microstructure of a nugget in a friction-stir-welded Ti-6Al-4V joint[END_REF][START_REF] Jia | Effect of static spheroidization on superplasticity of fine lamellae in nugget of a friction stir welded Ti-6Al-4V joint[END_REF].

Despite the suitable properties of β-Ti alloys having a single body-centered cubic (BCC) crystal structure in the aerospace and automobile industries or biomedical engineering [START_REF] Boyer | Aerospace applications of beta titanium alloys[END_REF][START_REF] Banerjee | Perspectives on Titanium Science and Technology[END_REF][START_REF] Cotton | State of the Art in Beta Titanium Alloys for Airframe Applications[END_REF], only limited investigations are available dealing with their microstructural and textural characterization after FSW [START_REF] Reynolds | Texture in friction stir welds of Timetal 21S[END_REF][START_REF] Liu | Microstructural and mechanical properties of a beta-type titanium alloy joint fabricated by friction stir welding[END_REF][START_REF] Liu | Elucidation of microstructural evolution of beta-type titanium alloy joint during friction stir welding using liquid CO2 cooling[END_REF]. In fact, the microstructure, texture and mechanical properties were investigated in detail for β-type Ti-15-3 alloy (Ti-15V-3Cr-3Al-3Sn, wt.%) after normal FSW and liquid CO 2 cooling FSW [START_REF] Liu | Microstructural and mechanical properties of a beta-type titanium alloy joint fabricated by friction stir welding[END_REF][START_REF] Liu | Elucidation of microstructural evolution of beta-type titanium alloy joint during friction stir welding using liquid CO2 cooling[END_REF]. But only a qualitative texture presentation as a function of welding speed was reported for β-type TIMETAL 21S alloy (Ti-15Mo-3Nb-3Al-0.2Si, wt.%) fabricated by FSW butt-joint [START_REF] Reynolds | Texture in friction stir welds of Timetal 21S[END_REF].

Hence, the aim of the present study is the characterization in detail of the microstructure and texture evolution in different parts of a FSW joint (BM, HAZ, TMAZ, and SZ) of β-type TIMETAL 21S obtained by lap welding. It is believed that the present results help to elucidate the response of the present alloy to the FSW and extend our fundamental knowledge on the characteristic of FSW β-type Ti alloys.

Experimental procedure

The TIMETAL β-21S alloy was kindly received from TIMET Savoie, Ugine, France in the form of cold-rolled and subsequently annealed sheets with 1.6 mm of thickness. The chemical composition of the sheet TIMETAL β-21S alloy is given in Table 1. The alloy is produced with a single BCC structure that is supposed stable at room temperature. In addition, the precipitation of α-phase was not detected in similar TIMETAL β-21S alloy [START_REF] Reynolds | Texture in friction stir welds of Timetal 21S[END_REF] and is not studied in the present work after welding. Two sheets were superimposed and then lap welded by FSLW at the French Welding Institute (Institut de Soudure). The surface appearance of the 21S alloy weld joint is shown in Figure 1.

The welding reference frame is welding direction (WD), transverse direction (TD), and normal direction (ND). In the present case, the rolling direction (RD) is parallel to WD. The advancing side (AS) and retreating side (RS) are also indicated in the Figure.

The tool geometry was quite simple but confidential details were not given by the welding center. The welding was realized with welding speed of 50 mm/min with a constant tool rotation rate of 200 rpm. Argon shielding was used to avoid any oxidation of the sample and welding tool. The selection of specific welding speed and tool rotation was selected based on the investigation reported by Reynoled et al. [START_REF] Reynolds | Texture in friction stir welds of Timetal 21S[END_REF] indicating that the combination of 50 mm/min and 200 rpm results in good welding appearance. The conic pin diameter was about 5 mm at the surface of the stacked sheets and the shoulder diameter was 30 mm (see Figure 1). An excellent surface appearance with no defects can be noticed indicating that a defectfree 21S alloy weld joint was successfully fabricated by FSW using the present welding parameters. The microstructure observations and texture analysis of the joints were carried out by EBSD measurements on the transverse cross-section (TD-ND plane) of the weld alloy.

Sample preparation consists of mechanical polishing using abrasive paper up to 2400 and then followed by a final polishing in 50 % colloidal silica (OP-S) + 49 % H 2 O + 1% H 2 O 2 solution in vibratory polisher for 17 h.

The EBSD measurements were carried out using a TSL-EDAX-Hikari system mounted on a scanning electron microscope FEG-SEM ZEISS Supra 55 VP operating at 20 kV. The data acquisition and analysis were processed by the Orientation Imaging Microscopy, OIM TM software.

To manage these EBSD acquisitions, a batch including combo scans (beam control) and single scans (motor control) was programed. So, a scan area of about 16 × 4 mm 2 with a step size of 2 µm was firstly recorded using the combo function to get a general overview of the FSLW microstructure. Then, several single scans and combo scans were analyzed with step size of 0.2 or 0.4 µm to analyze more precisely different parts of the FSLW joint.

For this study, the system was run continuously during 9 days without any user operation,

showing the very good stability of the present SEM-EBSD analytical system. During this long-time acquisition, about 100 million Kikuchi patterns were acquired (exactly 97 922 014 points). The majority of these points was used to describe the microstructure and the texture of the studied sample.

A grain dilation and grain CI standardization clean-up with a grain tolerance angle of 5° and a minimum grain size of 5 pixels were applied to improve the quality of EBSD raw data. Due to the orientation noise, spurious boundaries with misorientation angles θ < 2° were excluded from EBSD maps [START_REF] Liu | Elucidation of microstructural evolution of beta-type titanium alloy joint during friction stir welding using liquid CO2 cooling[END_REF]. In addition, all datum points with a confidence index (CI) lower than 0.05 were excluded from the quantitative analyses, where CI quantifies the reliability of the indexed pattern [START_REF] Mikami | Effect of reference point selection on microscopic stress measurement using EBSD[END_REF]. The mean grain size was measured using the equivalent diameter approach.

The grain orientation spread (GOS) and Kernel Average Misorientation (KAM) approaches in OIM™ software were used for the identification of recrystallized grains and strain distribution (dislocation density), respectively [START_REF] Cho | Determination of a mean orientation in electron backscatter diffraction measurements[END_REF][START_REF] Betanda | Measurement of stored energy in Fe-48%Ni alloys strongly cold-rolled using three approaches: Neutron diffraction, Dillamore and KAM approaches[END_REF]. GOS is defined as the average deviation between the orientation of each point in the grain and the average orientation of the grain, where grains with GOS < 1° are considered fully recrystallized [START_REF] Cho | Determination of a mean orientation in electron backscatter diffraction measurements[END_REF].

KAM is known as the average misorientation angle of a given point with all its neighbors belong to the same grain [START_REF] Badji | Texture, misorientation and mechanical anisotropy in a deformed dual phase stainless steel weld joint[END_REF]. In the present case, the θ KAM value was calculated from the mean misorientation angle between a given point and its 3 th neighbors excluding misorientations greater than 5°. Hence the dislocation density can be obtained following the equation [START_REF] Azzeddine | An investigation of the stored energy and thermal stability in a Cu-Ni-Si alloy processed by high-pressure torsion[END_REF]:

(

where α = 3 for mixture of tilt and twist boundary types [START_REF] Azzeddine | An investigation of the stored energy and thermal stability in a Cu-Ni-Si alloy processed by high-pressure torsion[END_REF], b = 0.28 nm is the Burgers vector [START_REF] Castany | In situ TEM study of dislocation slip in a metastable β titanium alloy[END_REF], n is the nearest neighbor and d is the scan step size.

MTEX software was used to analyze the evolution of texture by calculating the orientation distribution function (ODF) using the harmonic method (L = 22), and a gaussian function with a half-width of 5° to model each orientation [START_REF] Bachmann | Texture Analysis with MTEX -Free and Open Source Software Toolbox[END_REF].

The texture strength can be calculated by the texture index, I, using the following equation [START_REF] Kocks | Texture and Anisotropy: Preferred Orientations in Polycrystals and Their Effect on Materials Properties[END_REF]:

(

where f(g) is the ODF value, G is the Euler space.

Results and discussion

3.1.Microstructure and texture of TIMETAL 21S before FSW (BM)

Figure 2 shows the orientation imaging micrography (OIM) in inverse pole figure The initial texture of 21S alloy is shown in Figures 3a andc in term of the {110} pole figure

and ODF section at φ 2 = 45°, respectively. The ideal deformation and recrystallization texture of rolled BCC alloys are shown for comparison in Figures 3b andd and their descriptions are given in Table 2. The texture of the as-received alloy is weak and characterized by the presence of deformed rolling texture components (A 1 , A 2 , B 2 and C 2 ) and recrystallized texture components (Cu and RCu). 4a shows the wide field orientation map (about 16 x 4 mm 2 ) of the transversal crosssection of the weld 21S alloy. The AS and RS of the weld sample are identified in the Figure .   Basically, the AS and RS regions are created due to the nature of the stirring process [START_REF] Nandan | Recent advances in friction-stir welding -Process, weldment structure and properties[END_REF]. In AS region, the rotation and translation have the same direction which leads to an increase of plastic deformation and heat in this region. In contrast, the rotation and translation directions are opposite in RS region which results in lower deformation [START_REF] Nandan | Recent advances in friction-stir welding -Process, weldment structure and properties[END_REF].

The microstructure and texture of zones 1 to 7 indicated by black rectangles will be inspected in the next sections. The typical FSW zones SZ and BM can be easily distinguished in Figure 4a. The transition zones (TMAZ and HAZ) between the SZ and BM can be identified with smaller EBSD areas (see next sections). As can be seen, the two sheets are well bonded in the SZ region. The SZ is free of defects and has a basin shape with a notably spreading near the top surface due to the effect of the tool shoulder on material flow during FSW. In addition, the SZ is not symmetric and it is shifted toward RS which is a characteristic of the tool movement on both sides (AS and RS) during the welding.

Even if the deformation in SZ is complicated during the FSW, it is recognized as a simple shear deformation mode [START_REF] Reynolds | Texture in friction stir welds of Timetal 21S[END_REF][START_REF] Liu | Microstructural and mechanical properties of a beta-type titanium alloy joint fabricated by friction stir welding[END_REF][START_REF] Park | Basal plane texture and flow pattern in friction stir weld of a magnesium alloy[END_REF]. Accordingly, the texture in this zone is a simple shear texture and it is usually presented in the SPN-SD plane where SPN is the shear plane normal and SD is the shear direction [START_REF] Li | Texture formation during equal channel angular extrusion of fcc and bcc materials: comparison with simple shear[END_REF]. Before going further, it is important to identify the relation between (ND, TD) and (SPN, SD) references frames. In FSW, the shear plane is considered parallel to the pin column surface (SZ/TMAZ border as indicated by dash line in Figures 4b andc) [START_REF] Reynolds | Texture in friction stir welds of Timetal 21S[END_REF][START_REF] Liu | Microstructural and mechanical properties of a beta-type titanium alloy joint fabricated by friction stir welding[END_REF][START_REF] Park | Basal plane texture and flow pattern in friction stir weld of a magnesium alloy[END_REF]. In this case, the shear plane is not constant within the SZ and follows a curved surface across the SZ due to the rotation of the pin [START_REF] Liu | Microstructural and mechanical properties of a beta-type titanium alloy joint fabricated by friction stir welding[END_REF][START_REF] Mironov | Development of grain structure during friction stir welding of pure titanium[END_REF]. Based on that, in the center of the SZ (zones 1 to 4), the TD and ND should be parallel to SD and SPN as shown in Figure 4b, respectively. Nevertheless, the plastic deformation in zones 2 and 3 is driven by the conical-shaped tool probe [START_REF] Liu | Microstructural and mechanical properties of a beta-type titanium alloy joint fabricated by friction stir welding[END_REF] which results in the deviation of SPN from ND. Hence, suitable rotations along TD and WD must be performed to present the texture of zones 2 and 3 in the simple shear (SPN, SD) reference frame. For zone 7, Figure 4c shows that the angle between SPN and ND is around 45° and SD coincide with WD on the AS region since it is well accepted that the SD is tangential to the tool rotation. Thus, to present the texture in the simple shear (SPN, SD) reference frame the texture data of SZ of the zone 7 must be rotated by 45° around WD and then by 90° around ND.

Microstructure and texture of SZ through thickness

The variation of microstructure through the thickness of SZ is shown in Figure 5a. The black zones in the maps correspond to the areas with CI values lower than 0.05 and may be attributed to distorted nature of the sample due to the high deformation. Excellent grain refinement is occurred in the SZ zone, in which the grain size decreases drastically from 43 µm (BM) to a range of 1.2-1.8 µm. It was reported that the grain size of SZ increases from 2 µm (5.08 mm/s) to 3 µm (0.85 mm/s) with the decrease of welding speed of FSWed S21 alloy [START_REF] Reynolds | Texture in friction stir welds of Timetal 21S[END_REF]. Basically, the microstructure of SZ is similar through thickness except for a higher fraction of recrystallized grains in the top of SZ (zone 1) which can be associated with the stirring effect of the tool shoulder and the possible wear of the welding tool rather than with the temperature gradient. Indeed, several investigations reported that the presence of grain size gradient through the thickness of SZ is caused by the temperature gradient generated from the FSW tool and the low thermal conductivity of the materials [START_REF] Yoon | Effect of rotation rate on microstructure and texture evolution during friction stir welding of Ti-6Al-4V plates[END_REF][START_REF] Yoon | Microstructure and texture distribution of Ti-6Al-4V alloy joints friction stir welded below β-transus temperature[END_REF][START_REF] Han | Impact of friction stir welding on recrystallization of oxide dispersion strengthened ferritic steel[END_REF][START_REF] Ahmed | Effect of Tool Geometry and Heat Input on the Hardness, Grain Structure, and Crystallographic Texture of Thick-Section Friction Stir-Welded Aluminium[END_REF]. However, the similarity of grain size in the three zones demonstrated that the FSW parameters used for the present study produce low heat input and a small temperature gradient across the thickness. In contrast, Figures 5b andc indicate that this small temperature gradient affects the evolution of grain boundaries type: very low grain boundaries (VLAGB) with misorientation 2°< θ < 5°, LAGBs with 5°< θ < 15° and (High) HAGBs with θ > 15°. As can be observed, the three zones contain a similar fraction of LAGBs (~ 12 %). While the microstructure of the top of SZ (zone 1) contains the highest fraction of HAGBs (about 60 %). Zones 2 and 3 are characterized by a high fraction (about 39 %) of VLAGBs indicating a high dislocation density which still not transformed to LAGBs or HAGBs due to low temperature in these zones.

The GOS map indicates the presence of layer structure in zone 3 as shown by the dashed line in Figure 5a. The formation of layer structure in the SZ results from the plastic flow of the transferred metal during welding, which is arising from the tool shoulder and stir pin [START_REF] Kumar | The role of friction stir welding tool on material flow and weld formation[END_REF][START_REF] Han | Impact of friction stir welding on recrystallization of oxide dispersion strengthened ferritic steel[END_REF].

The interaction between different layers leads to an inhomogeneous deformation and plastic strains which affect directly the microstructure and texture. As already shown in Figure 5a, the grain refinement is very similar through the SZ thickness. In contrast, a net difference in texture evolution can be noticed through thickness as shown in Figure 6. Zone 1 exhibits a typical shear texture with the presence of all shear texture components except the D 1 component. The ideal shear texture components of BCC materials are shown for comparison in Figure 6 and their description is given in Table 3. The original and rotated {110} pole figures and ODF sections at φ 2 = 45° of zone 2 and 3 are shown in Figure 6. The texture of zone 2 seems quite similar to the zone 1. While, the texture of zone 3 shows the domination of D 1 , , and F texture components. The degradation of the typical shear texture of the SZ zone through the thickness is attributed to the complexity of deformation and strain heterogeneity.

The development of texture gradient through SZ thickness could explain the difference in the grain boundary distribution in these zones as shown in Figures 5b andc.

Microstructure and texture of zone 4 (SZ, TMAZ and HAZ)

The WD-IPF map of zone 4 shown in Figure 7 enclosed three regions: region 1: SZ, region 2:

TMAZ and region 3: HAZ. The WD-IPF and GOS maps, grain size distribution and grain boundaries distribution of the three regions are given at higher magnifications in the same 4. The evolution of the texture in the three regions is presented in Figure 8. Obviously, the different regions of the joint, SZ, TMAZ, and HAZ, exhibit different grain size, grain morphology and texture. As already seen in section 3.2, region 1 which presents the bottom of SZ exhibits an excellent grain refinement (1.6 µm) and typical shear texture with a domination of D 2 component. Several studies reported the domination of D 2 or D 1 components in FSW-BCC materials [START_REF] Mironov | Microstructural evolution during friction stirprocessing of pure iron[END_REF][START_REF] Reynolds | Texture in friction stir welds of Timetal 21S[END_REF][START_REF] Liu | Microstructural and mechanical properties of a beta-type titanium alloy joint fabricated by friction stir welding[END_REF][START_REF] Liu | Elucidation of microstructural evolution of beta-type titanium alloy joint during friction stir welding using liquid CO2 cooling[END_REF][START_REF] Fonda | Thermal Stir Welds in Titanium[END_REF][START_REF] Rahimi | Development of Microstructure and Crystallographic Texture in a Double-Sided Friction Stir Welded Microalloyed Steel[END_REF]. It is believed that the domination of grains with D 2 or D 1 texture component is attributed to the activation of {112}<111> slip system [START_REF] Liu | Microstructural and mechanical properties of a beta-type titanium alloy joint fabricated by friction stir welding[END_REF]. The similarity of microstructure in region 1 with zones 1 to 3 (Figure 5) confirms the uniform grain refinement throughout the SZ thickness. It is known that the grain refinement in FSW material is caused by the occurrence of dynamic recrystallization. This is in good agreement with the present results since the SZ zone shows a higher fraction of recrystallized grains and HAGBs compared to the TMAZ and HAZ zones (see Table 4).

The TMAZ area, which is the transition area between SZ and HAZ, is characterized by a mixture of elongated and equiaxed grains owing to an upward flow of material near the SZ during the operation of plunging and stirring. The dislocation density is higher in this zone which means this zone is highly deformed. It can clearly see from the GOS map of the TMAZ that the elongated grains are surrounded by the small recrystallized grains. The magnification of the rectangle zone in the GOS map (Figure 7) is shown in Figure 9. The serration or the bulging of the original grain boundaries (seen in grain G1) indicates the nucleation of new grains through strain-induced boundary migration (SIBM) mechanism which is ranked as discontinuous dynamic recrystallization (DDRX) [START_REF] Sakai | Dynamic and post-dynamic recrystallization under hot, cold and severe plastic deformation conditions[END_REF]. The bulging of the grain boundaries is caused by the difference in dislocation density between the two sides of the grain boundaries [START_REF] Sakai | Dynamic and post-dynamic recrystallization under hot, cold and severe plastic deformation conditions[END_REF]. With further strain, the misorientation between the bulges and deformed grains increases, leading to the development of sub-grain boundaries and LAGBs as shown in the grain boundary map (Figure 9b) and then transformed into HAGBs and the bulged regions become DRX grains as can be seen in grain G2 (Figure 9a). It must be noted that grain G3

shows an orientation gradient and contains sub-grain boundaries and LAGBs inside it. This is a characteristic of continuous DRX (CDRX) by sub-grain development (SD) mechanism [START_REF] Sakai | Dynamic and post-dynamic recrystallization under hot, cold and severe plastic deformation conditions[END_REF].

In this mechanism, the LAGBs are developed and gradually transformed to HAGBs by the accumulation and continuous absorbing of dislocations inside the deformed grain [START_REF] Sakai | Dynamic and post-dynamic recrystallization under hot, cold and severe plastic deformation conditions[END_REF]. However, the HAGBs are still not developed at this stage of processing. Similar observation was reported for β-type Ti-15V-3Cr-3Al-3Sn (wt.%) alloy processed by FSW [START_REF] Liu | Microstructural and mechanical properties of a beta-type titanium alloy joint fabricated by friction stir welding[END_REF]. It is believed that the grain refinement in the SZ is dominated by the CDRX mechanism due to the severe plastic deformation and high temperature in this zone [START_REF] Liu | Microstructural and mechanical properties of a beta-type titanium alloy joint fabricated by friction stir welding[END_REF].

The microstructure and texture of the HAZ zone are quite similar to the BM except the grain size is smaller and high fraction of deformed grains. It is recognized that the HAZ is only affected by the heat generated during welding which caused the increase of grain size compared to the BM. However, it is not the case in the present investigation. The decrease of grain size and increase of deformation features (Table 4) in HAZ indicated that this zone is also affected by deformation. In fact, this zone is very close to the TMAZ which can be considered as a transition zone between TMAZ and the real HAZ. This could explain the presence of deformation features in this zone.

Microstructural, textural evolution and related strengthening mechanism

The WD-IPF map of zone 7 (indicated in Figure 4) is shown in Figure 10a. The wide field orientation map of this zone allows to visualize the SZ, TMAZ and HAZ in the AS region.

The grain size gradient from the SZ to the HAZ is well seen from the WD-IPF map. The The evolution of microstructural parameters like grain size, dislocation density and grain boundary type (VLAGBs, LAGBs, and HAGBs) and the texture strength within the 6 areas are shown in Figure 12. The grain size (Figure 12a) gradually decreases from HAZ to SZ due to the increase of dislocation density (Figure 12b). However, the dislocation density decreases in SZ due to the occurrence of DRX which results in grain refinement in this zone. The increase of HAGBs and the decrease of VLAGBs (Figure 12c) within these zones are in good agreement with the occurrence of grain refinement mechanism. In contrast, the texture index shown in Figure 12d indicates that the texture strength decreases between HAZ and TMAZ

where M T is the Taylor factor predicted for the deformed material, M R is the Taylor factor for a random texture, σ 0 is the friction stress, α 1 is a constant, G is the shear modulus and k HP is the Hall-Petch constant.

According to the results shown in Figure 13, it can be suggested that the mechanical properties in HAZ are govern by grain size and dislocation strengthening. The three strengthening mechanisms are responsible for mechanical properties in the TMAZ. While, grain size and texture strengthening mechanisms are dominated in SZ. Obviously, measurements of mechanical properties such as microhardness or tensile test are needed to confirm these speculations.

Conclusion

The microstructure and texture evolutions of TIMETAL 21S alloy processed by friction stir lap welding were investigated using EBSD, and the following conclusions were obtained:

 The two TIMETAL 21S sheets were successfully welded by friction stir lap welding.

 According to the microstructural and textural evolution, the weld region of 21S alloy can be separated into four zones: BM, HAZ, TMAZ, and SZ.

 The microstructure of SZ was similar through the thickness and an excellent grain refinement (1.2-1.8 µm) was obtained during the FSLW. In contrast, a shear texture gradient is developed across the thickness due to the complexity of deformation.

 The characteristic of the microstructure of TMAZ demonstrates the occurrence of DDRX. A typical shear texture is developed in this zone.

 The microstructure and texture of HAZ are quite similar to those of BM.

 The concomitant occurrence of grain size, dislocation and texture strengthening is responsible for the mechanical property distribution in different parts of FSLW joint.

Figure 1 :

 1 Figure 1: Photograph of the weld surface appearance. The welding reference frame are: welding direction (WD), transverse direction (TD), and normal direction (ND). Advancing side (AS), and retreating side (RS) are indicated.

(

  RD-IPF) map and GOS map of the as-received (BM) S21 alloy. The black line indicates high angle boundaries (HAGBs) with misorientation θ > 15°. The average grain size and the grain recrystallized fraction are shown in the upper side of the RD-IPF and GOS maps, respectively.

Figure 2 :

 2 Figure 2: (a) RD-IPF and (b) GOS maps of the as-received (BM) 21S alloy.

Figure 3 :

 3 Figure 3: (a, c) {110} pole figure and ODF section at φ 2 = 45° of as-received (BM) 21S alloy and (b, d) main ideal texture component positions in {110} pole figure and ODF section at φ 2 = 45° of rolled BCC alloys.

Figure 4 :

 4 Figure 4: (a) WD-IPF map of the transversal cross-section of the weld 21S alloy, (b, c) relation between (TD, ND) and (SPN, SD) reference frames in the SZ center and SZ near the AS region, respectively.

Figure 5 :

 5 Figure 5: (a) WD-IPF, KAM, and GOS maps, (b) histogram of grain boundary misorientation distribution and (c) evolution of VLAGBs, LAGBs and HAGBs of the zone 1, 2 and 3 in SZ of 21S alloy.

Figure 6 :

 6 Figure 6: Original and rotated {110} pole figures and ODF sections at φ 2 = 45° of (a) zone 1, (b) zone 2, and (c) zone 3 of SZ of welded 21S alloy. Main ideal shear texture component positions in {110} pole figure and ODF section at φ 2 = 45° of BCC alloys are presented for comparison.

Figure .

 . Figure. The mean grain size, HAGBs fraction, mean KAM value and dislocation density of SZ, TMAZ and HAZ regions are summarized in Table4. The evolution of the texture in the

Figure 7 :

 7 Figure 7: WD-IPF map, GOS map, grain size distribution, misorientation angle distribution of SZ, TMAZ and HAZ regions in zone 4 of FSW 21S alloy, respectively.

Table 4 .

 4 Mean grain size, HAGBs fraction, mean KAM value and dislocation density of SZ, TMAZ and HAZ region in zone 4 of FSW 21S alloy. Region d (µm) HAGBs (%) KAM (°) ρ × 10

Figure 8 :

 8 Figure 8: {110} pole figure and ODF section at φ 2 = 45° of region 1 (SZ), region 2 (HAZ), region 3 (BM) of zone 4 of welded 21S alloy. Main ideal BCC texture component positions in {110} pole figure and ODF section at φ 2 = 45° are presented for comparison.

Figure 9 :

 9 Figure 9: (a) WD-IPF, and (b) grain boundary map: sub-boundaries (1°-5°) highlighted by red lines, LAGB (5-15°) indicated by green lines, and HAGB>15° presented by blue lines.

texture of area 1

 1 to area 6 is shown in Figure 10b via {110} pole figure.

Figure 10 :

 10 Figure 10: (a) WD-IPF map of zone 7, as indicated in Figure 4 and (b) {110} pole figure corresponding to the 1-6 areas given on the WD-IPF map, respectively.

Figure 11 :

 11 Figure 11: Example between (a) original pole figure and ODF section at φ 2 = 45° and (b) rotated pole figure and ODF to coincide with the SPN-SD plane of area 1 (SZ) and blue box in area 2 shown in Figure 10.

Figure 12 :

 12 Figure 12: Evolution of (a) grain size, (b) dislocation density, (c) grain boundaries type and (d) texture index within HAZ, TMZ and SZ (1 to 6 areas).

Table 1 .

 1 Chemical composition of TIMETAL β-21S alloy

	Element	Mo	Nb	Al	Si	Fe	O	C	N	H	Ti
	Weight %	14-16	2.4-3.2	2.5-3.5	0.15-0.25	0.4	0.11-0.17	0.05 0.05	0.015 Bal

Table 2 .

 2 Position

	Notation		TD-ND plane
	Rolling component	Miller indice {hkl}<uvw> *	Euler angles (°) ( 1 , Φ,  2 )
		{011} < 1>	(144.7, 45, 0)
		{011} <2 >	(35.2, 45, 0)
		{211} < 1>	(50.7, 65.9, 63.3)
		{121} < 01> (129.2, 65.9, 26.5)
		{111} <0 1>	(60, 54.7, 45)
		{111} < 1>	(120, 54.7, 45)
	Cu	{001} <	>	(0, 0, 0)
	RCu	{011} < 1>	(90, 45, 0)
	* {hkl}//(TD, ND) plane and <uvw>//RD
	3.2

of ideal rolling texture components for BCC materials projected in TD-ND plane

[START_REF] Hölscher | Rolling and recrystallization textures of bcc steels[END_REF]

.

.Microstructure and texture of TIMETAL 21S after FSW Figure

Table 3 .

 3 Position of ideal shear texture components for BCC materials projected in SPN-SD plane[START_REF] Azzeddine | Texture evolution in high-pressure torsion processing[END_REF].

	Notation	SPN-SD plane
	Shear component	Miller indice	Euler angles (°)
		{hkl}<uvw>	( 1 , Φ,  2 )
		{011} <1 1>	(54.7, 45, 0)
		{101} < 11>	(125.2, 45, 90)
	J	{111} < 2>	(90, 54.7, 45)
		{111} <1 1>	(30, 54.7, 45)
	E	{112} < 1>	(90, 35.2, 45)
		{121} <1 1>	(39.2, 65.9, 26.5)
	F	{011} <100>	(0, 45, 0)

14 (m -2 ) SZ

	TMAZ 13.7	34.6	1.86	5.79
	HAZ	32.7	45.0	0.88	2.74
		1.6	62.9	1.76	5.48

Acknowledgements

The authors thank Dr. Y. Millet (TIMET Savoie, Ugine, France) for supplying the TIMETAL 21S sheets. A. Robineau and J.C. Goussain (Institut de Soudure, Goin, France) are gratefully acknowledged for the FSW.

Data availability

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


Salaheddine Sadi

Abdelkader Hanna

Thierry Baudin

François Brisset

Jose Maria Cabrera

Hiba Azzeddine
email: hiba.azzeddine@univ-msila.dz

Pr Hiba Azzeddine

Microstructure and texture evolution of ECAP-processed Mg-Ce alloy during isothermal annealing

Keywords: Annealing, Equal channel angular pressing, Magnesium, Microstructure, Rare earth, Recrystallization, texture

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

Introduction

Severe plastic deformation, especially equal-channel angular pressing (ECAP) processing has been shown to be effective in improving simultaneously strength, ductility, and corrosion resistance of Mg-based alloys [START_REF] Ma | Improving both strength and ductility of a Mg alloy through a large number of ECAP passes[END_REF][START_REF] Gopi | Electrochemical Behaviour of ECAP-Processed AM Series Magnesium Alloy[END_REF][START_REF] Bryła | Microstructure and mechanical characterisation of ECAP-ed ZE41A alloy[END_REF][START_REF] Prithivirajan | Analysing the combined effect of crystallographic orientation and grain refinement on mechanical properties and corrosion behaviour of ECAPed ZE41 Mg alloy[END_REF][START_REF] Sankuru | Effect of processing route on microstructure, mechanical and dry sliding wear behavior of commercially pure magnesium processed by ECAP with back pressure[END_REF]. Such improvement is attributed principally to grain refinement and texture modification [START_REF] Sankuru | Effect of processing route on microstructure, mechanical and dry sliding wear behavior of commercially pure magnesium processed by ECAP with back pressure[END_REF][START_REF] Li | Microstructure evolution and mechanical properties of Mg-Gd-Y-Zn-Zr alloy during equal channel angular pressing[END_REF][START_REF] Minárik | Effect of different c/a ratio on the microstructure and mechanical properties in magnesium alloys processed by ECAP[END_REF][START_REF] Sabat | Ductility enhancement in Mg-0.2%Ce alloys[END_REF][START_REF] Beyerlein | Texture evolution in equal-channel angular extrusion[END_REF][START_REF] Azzeddine | Texture evolution in high-pressure torsion processing[END_REF]. Due to the limited active slip systems in Mg-based alloys, grain refinement is obtained by dynamic recrystallization (DRX). Several DRX mechanisms have been evidenced in Mg-based alloys depending on the deformation processing conditions and alloying elements [START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF][START_REF] Minárik | The effect of Zr on dynamic recrystallization during ECAP processing of Mg-Y-RE alloys[END_REF][START_REF] Liu | Effect of Multi-Pass Equal Channel Angular Pressing on the Microstructure and Mechanical Properties of a Heterogeneous Mg88Y8Zn4 Alloy[END_REF][START_REF] Guerza-Soualah | The deformation and recrystallization behaviour of an Mg-Dy alloy processed by plane strain compression[END_REF][START_REF] Guerza-Soualah | Microstructural and textural investigation of an Mg-Dy alloy after hot plane strain compression[END_REF]. For example, twin-induced dynamic recrystallization (TDRX) and particle stimulated nucleation (PSN) are the mechanisms responsible for the occurrence of DRX at low strain and then transited to sub-grain development (SD) and grain boundary bulging dynamic recrystallization (GBBDRX) with increasing strain in ECAP processed Mg-0.3Ce (wt.%) alloy [START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF]. Subsequently, each mechanism has a direct consequence on texture development [START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF]. In addition, increasing deformation temperatures completely change the microstructure and texture due to the occurrence of grain growth [START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF].

Besides, several investigations demonstrated that the homogeneity of the microstructure of Mgbased alloys during ECAP processing is imposed by the initial grain size [START_REF] Figueiredo | Grain refinement and mechanical behavior of a magnesium alloy processed by ECAP[END_REF][START_REF] Fatemi | Review on ultrafined/nanostructured magnesium alloys produced through severe plastic deformation: microstructures[END_REF]. Based on the model proposed by Figueiredo and Langdon, there is an initial grain size criterion dc, in which the deformation microstructure is homogenously distributed [START_REF] Figueiredo | Grain refinement and mechanical behavior of a magnesium alloy processed by ECAP[END_REF]. However, up to now, this initial grain size criterion is not quantified since many parameters must be taken into account such as ECAP processing temperature and second particles distribution.

Recent researches demonstrated that DRX is usually restricted in rare-earth-containing Mg (Mg-RE) alloys which results in texture weakening due to solute segregations and/or particle pinning mechanisms that directly affect the grain boundary mobility and dislocation motion [START_REF] Guerza-Soualah | The deformation and recrystallization behaviour of an Mg-Dy alloy processed by plane strain compression[END_REF][START_REF] Hadorn | Effects of Solute and Second-Phase Particles on the Texture of Nd-Containing Mg Alloys[END_REF][START_REF] Imandoust | A review on the effect of rare-earth elements on texture evolution during processing of magnesium alloys[END_REF]. Subsequently, the static recrystallization and grain growth behaviours in Mg-RE alloys during annealing treatments will be different than in conventional Mg-based alloys [START_REF] Hantzsche | Effect of rare earth additions on microstructure and texture development of magnesium alloy sheets[END_REF][START_REF] Imandoust | Unraveling Recrystallization Mechanisms Governing Texture Development from Rare-Earth Element Additions to Magnesium[END_REF].

Several investigations can be found on the thermal stability of ECAP processed Mg-based alloys (like pure Mg and AZ31 alloy) mostly after isochronal annealing treatments [START_REF] Fan | Effect of heat treatment on internal friction in ECAP processed commercial pure Mg[END_REF][START_REF] Biswas | Thermal Response on the Microstructure and Texture of ECAP and Cold-Rolled Pure Magnesium[END_REF][START_REF] Victoria-Hernández | Strain-induced selective grain growth in AZ31 Mg alloy sheet deformed by equal channel angular pressing[END_REF][START_REF] Lei | Effect of annealing on the texture and mechanical properties of pure Mg by ECAP at room temperature[END_REF][START_REF] Naik | Effect of Annealing and Aging Treatment on Pitting Corrosion Resistance of Fine-Grained Mg-8%Al-0.5%Zn Alloy[END_REF][START_REF] Lei | Microstructural Stability and Mechanical Behavior of Pure Mg After One-Pass Equal Channel Angular Pressing Plus Intersection Annealing[END_REF].

Nevertheless, only limited studies have explored the static recrystallization and grain growth behaviours of ECAP processed Mg-RE alloys [START_REF] Wei | Effects of ECAP and Annealing Treatment on the Microstructure and Mechanical Properties of Mg-1Y (wt. %)[END_REF][START_REF] Yu | The Effect of Annealing on the Microstructure and Properties of Ultralow-Temperature Rolled Mg-2Y-0.6Nd-0[END_REF].

In this frame, the aims of the present study are the characterization of static recrystallization and grain growth microstructure and texture of Mg-0.3Ce (wt.%) alloy after ECAP for N = 1, 2, and 4 passes at 300 °C using route Bc and followed by isothermal annealing at 450 °C for 5, 30, 60, and 180 min. The results are discussed based on the effect of strain level, DRX mechanisms, second particles, and initial state.

Experimental procedure

Cylindrical samples of Mg-0.3Ce alloy were processed by ECAP at 300 °C via route Bc for N = 1 (1 ECAP), 2 (2 ECAP), and 4 (4 ECAP) passes. ECAP processing was conducted using a die having a channel angle of Φ = 90° and an outer arc of curvature of Ψ = 37°. Details on the evolution of deformation microstructure and texture can be found in a previous study [START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF]. The deformed samples underwent an annealing treatment in a radiation furnace at 450 °C for 5, 30, 60, and 180 min.

Electron backscatter diffraction (EBSD) areas of 400 × 400 μm 2 with a step size of 0.4 μm near the centre of the TD-ND plane (normal to the extrusion direction (ED)) of the samples were measured using a FEG-SEM SUPRA 55 VP field emission gun scanning electron microscope operating at 20 kV, where TD and ND mean the transversal and normal directions, respectively.

The EBSD data were analyzed using the TSL Orientation Imaging Microscopy (OIM™) software. A clean-up procedure consisting of a grain tolerance angle of 5°, a minimum grain size set at 5 pixels, and excluding all datum points with a confidence index CI > 0.05 was used to analyze the EBSD data.

The evolution of recrystallization fraction was calculated by using the grain orientation spread (GOS) approach implemented in OIM software [START_REF] Cho | Determination of a mean orientation in electron backscatter diffraction measurements[END_REF]. In Mg-based alloys, the recrystallized grains are considered having GOS values ≤ 2° [START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF][START_REF] Guerza-Soualah | Microstructural and textural investigation of an Mg-Dy alloy after hot plane strain compression[END_REF].

The analysis of the crystallographic texture from EBSD measurements was performed using MTEX software by calculating the orientation distribution function (ODF) using the harmonic method (L = 22), and a gaussian function with a half-width of 5° [START_REF] Bachmann | Texture Analysis with MTEX -Free and Open Source Software Toolbox[END_REF].

The microhardness of the samples was measured by a Vickers microhardness tester (SHIMADZU type HMV-2 tester) with a load of 100 g (Hv0.1) and a dwell time of 10 s.

Results

Figure 1 shows the evolution of the microstructure as a function of the annealing time of ECAPprocessed Mg-0.3Ce alloy. The black lines indicate the grain boundaries with high angle misorientation (HAGBs >15°). The evolution of the fraction of recrystallization and recrystallized grain size as a function of the annealing time of all samples are shown in Figure 2. The fraction of DRX during ECAP processing is also plotted in Figure 2a. The deformation microstructures of the samples were reported earlier in which the 1 and 2 ECAP samples exhibit a heterogeneous microstructure formed by large deformed grains and a very small amount of dynamically recrystallized grains [START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF]. On the contrary, the microstructure of the sample after 4 ECAP passes was more homogeneous and formed of equiaxed grains due to the occurrence of a significant amount of DRX [START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF]. The microstructures of 1 and 2 ECAP samples are still heterogeneous after annealing for 5 min as shown in Fig. 1. The microstructure of 1 ECAP sample changes with increasing annealing time but is still very heterogeneous and formed by a mixture of large and small recrystallized grains with a mean grain size of 49.0 µm. A high fraction of twins mainly extension twins 86° <112 ̅ 0> can be noticed. The microstructure of 1 ECAP sample is still not completely recrystallized (85%) even after a long annealing time (180 min) in comparison with the 91% recrystallized volume fraction for 2 and 4 ECAP samples. In contrast, a homogeneous microstructure is gradually formed in 2 ECAP sample with increasing annealing time. The mean recrystallized grain size increases from 2. 3a-c). Apparently, the different grain boundary fractions (Figures 3d-f) do not change with annealing time for all samples, except, the decrease of LAGBs after annealing for 30 min in 1 and 2 ECAP samples. However, the fraction of LAGBs and extension twin were higher in 1 ECAP sample compared with 2 and Figure 4 shows the recalculated pole figures {0002} of ECAP-processed and annealed Mg-0.3Ce alloy. The main ideal ECAP shear texture components for HCP metals are also shown in the pole figures and their descriptions in Euler angles are summarized in Table 1 [START_REF] Beyerlein | Texture evolution in equal-channel angular extrusion[END_REF]. The texture of 1 ECAP sample behaves similarly to the microstructure evolution. The texture of 1 ECAP sample oscillates between C1-fiber, B-fiber, and P-fiber with increasing annealing time. A new orientation named S (80°, 90°, 0°) is formed starting from 60 min of annealing. This S texture component was formed in the same alloy during ECAP processing for 4 passes at 350 °C [START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF].

Figure 4 shows that a B-fiber is formed in 2 ECAP sample during annealing between 5 and 30 min and gradually changes to P-fiber after 60 min of annealing and then to C1-fiber after 180 min. The texture of 4 ECAP sample was characterized by retained C1-fiber and this texture is stable through the entire annealing duration. Y-fiber (345°, 90°, 0-60°) C1-fiber (45°, 30°, 0-60°) C2-fiber (225°, 30°, 0-60°)

Figure 5 presents the Vickers microhardness evolution as a function of the annealing time of ECAP-processed and annealed Mg-0.3Ce alloy. The Hv values of as-cast and the deformed states of 1, 2, and 4 ECAP samples were added to the plot [START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF].

The microhardness of the as-cast Mg-0.3Ce alloy increases considerably after ECAP processing mainly due to the grain refinement as known by the Hall-Petch relationship [START_REF] Figueiredo | Deformation mechanisms in ultrafine-grained metals with an emphasis on the Hall-Petch relationship and strain rate sensitivity[END_REF]. However, besides the grain size, it was reported that the residual stress could be also an important factor in increasing the microhardness of severely deformed alloys [START_REF] Zhang | The tribological behavior of a surfacenanocrystallized magnesium alloy AZ31 sheet after ultrasonic shot peening treatment[END_REF][START_REF] Moon | Effects of residual stress on the mechanical properties of copper processed using ultrasonicnanocrystalline surface modification[END_REF].

The

Discussion

Effect of ECAP processing on static recrystallization and grain growth

The results of the microstructure and microtexture evolution shown in the present study demonstrated that the homogeneity of the ECAP-processed Mg-0.3Ce alloy during annealing process strongly depends on the number of ECAP passes, i.e., accumulated strain. As consequence, several parameters can influence the resulted microstructure and texture such as the initial grain size, the distribution of second particles and the DRX fraction. The equivalent strain, εeq is expressed as a function of number of ECAP passes, N, as following [START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF][START_REF] Iwahashi | Principle of equal-channel angular pressing for the processing of ultra-fine grained materials[END_REF]:

𝜀 𝑒𝑞 = 0.98𝑁 (1)
The homogeneity of the microstructure and the stabilization of the texture throughout the annealing time increase with increasing number of ECAP passes (equivalent strain). The distribution of second particles affects also the primary recrystallization process of the ECAP samples. Figure 7 displays the scanning electron microscope (SEM) micrography with energy-dispersive X-ray spectroscopy (EDS) analyses (points 1 and 2) of 1, 2 and, 4 ECAP samples after ECAP processing and annealing for 30 min, respectively. It must be noted that the as-cast sample contains two types of second particles: small particles in the form of line present inside the grains and along the grain boundaries identified as Mg12Ce phase and individual particles having a globular morphology and containing 9 % (% wt.) of Ce element [START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF].

On one hand, the SEM images (Figure 7) show that the grain boundaries in the deformed and annealed 1 ECAP sample still contain Mg12Ce second particles which restricted their migrations and retarded the formation of dynamic and static recrystallized grains as well. Obviously, the new grain boundaries of static recrystallized grains of the annealed 1 ECAP sample, as indicated by arrow, are free from these second particles. In contrast, the microstructures of deformed and annealed 2 and 4 ECAP samples indicate that their grain boundaries were free from these Mg12Ce second particles due to their dissolution during ECAP processing resulting in less pinning effect and allowing the occurrence of DRX. The zoom of the yellow box shown in the SEM image of annealed 1 ECAP sample evidences the existence of a precipitate free zone (PFZ) which was already developed during ECAP processing [START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF]. In addition, it can be observed the presence of ultra-fine precipitates homogenously distributed throughout the grains. Unfortunately, the identification of the composition of these precipitates was not possible. However, based on the precipitation sequence reported for Mg-Ce solid solution known as [START_REF] Saito | TEM study of real precipitation behavior of an Mg-0.5at%Ce agehardened alloy[END_REF]: Mg supersaturated solution → GP-zone → β1 (Mg3Ce) → β (Mg12Ce), then, these precipitates could be Mg3Ce and/or Mg12Ce phases. It was reported earlier that the solute depletion and the dissolution of the Guinier-Preston (GP) zone are the main mechanisms responsable for the formation of PFZ in the present alloy [START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF].

On another hand, the microstructures of all deformed samples show the presence of the same globular particles (point 2) and it is evidenced from the EDS analysis that the concentration of Ce element slightly increases compared to those found in the as-cast state (13 vs. 9 %). These particles are stable through the annealing for 30 min and seems not affecting the static recrystallization since these particles are present inside the SR grains. It has been reported that Mg12Ce second particles were the main reason for the deformation retained texture during isochronal annealing at 250 and 450 °C for 60 min of severe deformed (εeq = 1.6) Mg-1Ce (wt.%) alloy [START_REF] Basu | Shear band-related recrystallization and grain growth in two rolled magnesium-rare earth alloys[END_REF]. However, the present Mg-0.3Ce alloy exhibits completely different texture evolution during the isothermal annealing treatment. The mechanisms responsible of microstructure and texture evolution in Mg-RE alloys are not completely elucidated but it is established that they are strongly affected by the concentration of RE elements and their limited solubility in Mg matrix [START_REF] Imandoust | A review on the effect of rare-earth elements on texture evolution during processing of magnesium alloys[END_REF][START_REF] Stanford | Micro-alloying Mg with Y, Ce, Gd and La for texture modification-A comparative study[END_REF].

Interesting to note that 2 and 4 ECAP samples have almost similar grain size and microhardness value especially at the grain growth stage (annealing for 180 min). It can be speculated that there is a strain threshold (a critical shear strain) where beyond this value the evolution of recrystallized microstructure is similar. In the present case, this critical strain is εeq = 1.96 (2 ECAP). Nevertheless, this suggestion is not relevant for the texture evolution. As can be seen in Figure 4, the textures of 2 and 4 ECAP samples have completely different behaviour during isothermal annealing. The texture of 4 ECAP sample was retained all over the recrystallization and grain growth stages. While the texture of 2 ECAP sample completely changes from recrystallization to grain growth stage. This difference is related to the fraction of DRX during the ECAP processing. As shown in Figure 2a the DRX fraction for 2 and 4 ECAP samples was ~ 15 and ~55%, respectively. It is well known that the texture is retained during annealing if the DRX occurred during deformation processing [START_REF] Bhattacharyya | Texture enhancement during grain growth of magnesium alloy AZ31B[END_REF][START_REF] Al-Samman | Dynamic recrystallization during high temperature deformation of magnesium[END_REF][START_REF] Azzeddine | On the texture and grain growth in hot-deformed and annealed WE54 alloy[END_REF]. In this case and based on the oriented nucleation theory, the formation and the growth of new grains during subsequent annealing will have the same orientations of the parent grains [START_REF] Doherty | Nucleation and growth kinetics of different recrystallization texture components[END_REF]. If no DRX exists, the recrystallization texture is usually different from the deformation one [START_REF] Tighiouaret | On the evolution of microstructure, texture and corrosion behavior of a hot-rolled and annealed AZ31 alloy[END_REF][START_REF] Su | Texture Weakening and Grain Refinement by High Speed Rolling and Annealing of an AZ31 Magnesium Alloy[END_REF][START_REF] Tighiouaret | Thermal Stability of an Mg-Nd Alloy Processed by High-Pressure Torsion[END_REF] and this is explained by the oriented growth theory [START_REF] Doherty | Nucleation and growth kinetics of different recrystallization texture components[END_REF]. In this theory, the stability of an orientation will depend on the grain boundary mobility, thus some grains with specific orientations will preferably growth at the expense of other orientations [START_REF] Doherty | Nucleation and growth kinetics of different recrystallization texture components[END_REF].

The Kernel Average Misorientation (KAM) approach was used to estimate the stored energy resulting from the Geometrically Necessary Dislocations (GND) dislocations [START_REF] Betanda | Measurement of stored energy in Fe-48%Ni alloys strongly cold-rolled using three approaches: Neutron diffraction, Dillamore and KAM approaches[END_REF]. Figure 8 presents the KAM maps of the 1, 2, and 4 ECAP samples annealed for 5 min. The stored energy of the 4 ECAP sample is lower due to the occurrence of significant DRX during ECAP processing. The stored energy is practically similar for 1 and 2 ECAP samples, hence it is expected that both samples present the same static recrystallization behaviour. However, the recrystallization behaviour is completely different between both samples. In order to track the nature of the stored energy in both samples, Figure 9 presents the IPF, KAM, IQ highlighted by grain boundaries distribution and extension twin maps of selected zones of 1 and 2 ECAP samples annealed for 5 min. The fraction of very low grain boundaries VLAGBs (1°< θ <5°), LAGBs (5°< θ <15°), HAGBs (15°< θ) and extension twin are summarized in Table 2. The stored energy in the 1 ECAP sample seems to be generated from the high fraction of extension twin (Table 2). Indeed, it was demonstrated previously that the extension twins were not a preferred site of DRX during ECAP processing [START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF] nor during static recrystallization since the evolution of the microstructure as a function of annealing time indicates that the twins only grow. Several reports confirm this behaviour of extension twins during annealing in deformed Mg-based alloys [START_REF] Levinson | Influence of deformation twinning on static annealing of AZ31 Mg alloy[END_REF][START_REF] Zhao | Quasi-in-situ observations of low-angle grain boundaries, twins and texture evolution during continuous annealing in a cold-rolled Mg-Zn-Gd alloy[END_REF]. An example of different twin forms can be seen in the selected zone of 1 ECAP sample annealed for 180 min as shown in Figure 10. Several twins can be found inside the same grain in the form of steps as indicated by the CD distance (Figure 10d). Another type of form is the growth and coalescence of extension twins as indicated by arrows in Figure 10a. These types of twins seem to start at grain boundaries and continuous to grow until consuming the parent grain as observed in grains G1 and G2 (Figures 10a and10b) due to the high stored energy resulted from the dislocation/twin interaction. The migration and growth of extension twins is the main reason for the texture modification of 1 ECAP sample during annealing. In fact, it has been demonstrated by quasi-in-situ EBSD observations that the effect of extension twin on texture modification is more important than the orientation of nuclei [START_REF] Zhao | Quasi-in-situ observations of low-angle grain boundaries, twins and texture evolution during continuous annealing in a cold-rolled Mg-Zn-Gd alloy[END_REF].

Interesting observation can be noticed in Figures 10b andc (black box) where grain with

LAGBs is inside another grain. Both grains have the same orientation as can be recognized in the IPF map (Figure 10a). The development of grains in form of isolated grains is a characteristic of abnormal grain growth [START_REF] Victoria-Hernández | Strain-induced selective grain growth in AZ31 Mg alloy sheet deformed by equal channel angular pressing[END_REF]49]. Obviously, the stored energy in the 2 ECAP sample is not originated from the extension twins since a very low fraction was detected (0.5%) (Table 2). In contrast, the 2 ECAP sample contain a significant fraction of LAGBs (15.1%) compared to 1 ECAP sample (5.0%). The presence of LAGBs indicates the dominance of the recovery process [START_REF] Fatemi | Characterization of continuous dynamic recrystallization in WE43 magnesium alloy[END_REF]. The results demonstrated that the presence of LAGBs ensures a homogeneous evolution of the microstructure with further annealing time.

The present results demonstrated the existence of a direct connection between stored energy and the stability of texture. The texture was unstable for 1 and 2 ECAP samples (having high stored energy) and retained for 4 ECAP sample (having low stored energy). This is in good agreement with the assumption that the possibility of recrystallization nuclei with a variety of orientations may increase with increasing the stored energy [START_REF] Jung | Role of RE in the deformation and recrystallization of Mg alloy and a new alloy design concept for Mg-RE alloys[END_REF]. Usually, the development of B-fiber during ECAP processing of Mg-based alloys causes the deterioration of the yield strength and poor ductility [START_REF] Ying Ma | Microstructure, Texture Evolution, and Mechanical Properties of ECAP-Processed ZAT522 Magnesium Alloy[END_REF][START_REF] Janeček | Texture and microstructure evolution in ultrafine-grained AZ31 processed by EX-ECAP[END_REF]. Therefore, the modification of texture during annealing processing could be beneficial for improving the mechanical properties of Mg-0.3Ce alloy.

4.2.Effect of DRX fraction on the static recrystallization mechanisms

The amount of DRX during ECAP processing affects largely the evolution of microstructure and texture mechanisms responsible for the static recrystallization and grain growth. For 1 ECAP sample where the DRX was absent two main static recrystallization mechanisms were detected after annealing for 30 min as shown in Figure 11. The first mechanism was the subgrain development (SD) as shown in zones 1 and 2, respectively. The recrystallized grain (G1) is formed due to the accumulation of dislocations inside the grain and gradually transformed to LAGBs and then HAGBs with increasing annealing time [START_REF] Xie | Transition of dynamic recrystallization mechanisms of as-cast AZ31 Mg alloys during hot compression[END_REF]. The second mechanism was the particle simulated nucleation (PSN) as depicted in zones 2 and 3. The presence of large particles, as shown in Figure 7, prevents the dislocation motion which led to their accumulation around these particles, gradually transformed to LAGBs and eventually to new grains with HAGBs [START_REF] Humphreys | Chapter 9 -Recrystallization of Two-Phase Alloys[END_REF]. The PSN mechanism was an important mechanism besides the twin-induced dynamic recrystallization (TDRX) for the development of grains during ECAP processing at low strain (1 ECAP at 300 °C) [START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF]. The SD and PSN mechanisms are considered as continuous recrystallization process [START_REF] Xie | Transition of dynamic recrystallization mechanisms of as-cast AZ31 Mg alloys during hot compression[END_REF]. The corresponding {0002} pole figures show that the orientations of recrystallized grains of both mechanisms were different from the deformed parent grains.

This explains the instability and the modification of the texture of the 1 ECAP sample as a function of annealing time (Figure 4). In the case of 2 ECAP sample where a limited DRX (15%) was developed during ECAP, the static recrystallization seems to start in shear bands during annealing for 5 min as evidenced in Figure 12. The recrystallized grains in the shear band developed a C1-fiber in addition to the Bfiber compared to strong B-fiber in the deformed grains. It is well known that the nucleation in the shear bands led to the formation of grains with a various orientation [START_REF] Basu | Shear band-related recrystallization and grain growth in two rolled magnesium-rare earth alloys[END_REF][START_REF] Singh | Influence of simultaneous alloying with Ca and Sc on the high temperature deformation mechanism, texture, and recrystallization behavior of Mg-Ca-Sc alloys[END_REF]. A close inspection of the microstructure of 2 ECAP sample after annealing for 60 min (Figure 1) shows that some shearing traces can still be observed. The evolution of texture of 2 ECAP sample after This is confirmed by the decrease of LAGBs in 1 and 2 ECAP samples between annealing for 5 and 30 min and the formation of new grains as shown in Figure 1. Despite the high static recrystallization fraction, the LAGBs fraction was stable during the remaining annealing treatment. Similar fractions have been reported in rolled Mg-2Y-0.6Nd-0.6Zr alloy [START_REF] Yu | The Effect of Annealing on the Microstructure and Properties of Ultralow-Temperature Rolled Mg-2Y-0.6Nd-0[END_REF]. In this case, the evolution of LAGBs does not follow the traditional mechanism by dislocation recovery, rearrangement and finally transformation to HAGBs [START_REF] Humphreys | Chapter 7 -Recrystallization of Single-Phase Alloys[END_REF]. It was demonstrated by quasi-in-situ EBSD measurements that the pre-existing LAGBs can hardly migrate during annealing and even new LAGBs can formed between two neighbour grains with similar orientation [START_REF] Zhao | Quasi-in-situ observations of low-angle grain boundaries, twins and texture evolution during continuous annealing in a cold-rolled Mg-Zn-Gd alloy[END_REF]. In addition, the presence of very fine second phase particles can prevent the transformation of dislocations to LAGBs and then to HAGBs.

Conclusion

• The recrystallization and grain growth behaviour of Mg-0.3Ce alloy were investigated after ECAP processing at 300 °C for N = 1, 2 and 4 passes (via route Bc) followed by isothermal annealing treatment at 450 °C for 5, 30, 60 and 180 min.

• The recrystallization and grain growth microstructure and texture homogeneity strongly related to the strain level, second phase particle distribution and fraction of DRX.

• Continuous static recrystallization mechanisms were evidenced in low strained sample.

• Beyond a strain threshold εeq = 1.96 (2 ECAP) the evolution of microstructural parameters such as grain size and microhardness in the grain growth stage were the same.

  8 µm after 5 min to 19.4 µm after 180 min. Fine grains with an average value of 3.2 µm were formed in 4 ECAP sample after annealing for 5 min. The microstructure of 4 ECAP sample remains homogeneous with increasing annealing time and the recrystallized grain size increases to 18.6 µm up annealing for 180 min. The formation of extension twins is restricted in 2 and 4 ECAP samples due to the grain size effect.

Figure 1 :

 1 Figure 1: Inverse pole figure maps in extrusion direction (ED-IPF) of ECAP-processed and annealed Mg-0.3Ce alloy.

Figure 2 :

 2 Figure 2: Evolution of (a) fraction of recrystallization and (b) recrystallized grain size as a function of annealing time of 1, 2, and 4 ECAP samples.

Figure 3 :

 3 Figure 3: Distribution of misorientation angle and LAGBs, extension twins and HAGBs fraction evolution as a function of annealing time of ECAP-processed Mg-0.3Ce alloy: (a, d) 1 ECAP, (b, e) 2 ECAP and (c, f) 4 ECAP samples.

Figure 4 :Table 1 .

 41 Figure 4: Texture evolution of ECAP-processed and annealed Mg-0.3Ce alloy.

  Hv value of 1 ECAP sample was stable around 41 Hv when annealing for 5 to 60 min and then decreases slightly to 38.6±1 Hv after 180 min. While the microhardness of 2 and 4 ECAP samples decreases rapidly with increasing annealing time to reach both a similar value of 41.5±1.2 Hv at 180 min of annealing. The decrease of the microhardness is a consequence of the occurrence of recrystallization and grain growth phenomena. However, the Hv values of all samples are still higher than the as-cast value (37 Hv).

Figure 5 :

 5 Figure 5: Vickers microhardness evolution as a function of annealing time of ECAP-processed and annealed Mg-0.3Ce alloy.

Figure 6

 6 Figure6presents a bigger zone (1120 x 880 µm 2 ) with 0.5 µm of step size of 1 ECAP sample annealed for 180 min in which the microstructure is still heterogeneous even after a long annealing time. The low strain results in a low fraction of DRX and a significant volume fraction of unrefined coarse grains. As known, the grain boundaries are favoured nucleation sites, hence the number of available nucleation sites is very low for 1 ECAP sample. In addition, a very low fraction of grain boundaries was already present in the as-cast alloy (before ECAP processing) since the microstructure of the as-cast alloy was characterized by very coarse grains (~ 1 mm)[START_REF] Sadi | Characterization of microstructure and texture of binary Mg-Ce alloy processed by equal channel angular pressing[END_REF]. These conditions result in a decrease in the nucleation rate and consequently delay the static recrystallization and cause the heterogeneity of the microstructure as shown in Figures1 and 2. The 2 and 4 ECAP samples were completely recrystallized after annealing for 30 min while the recrystallization of 1 ECAP sample was retarded to 60 min with only 85 % of recrystallized grains.

Figure 6 :

 6 Figure 6: (a) ED-IPF map, (b) {0002} recalculated pole figure and (c) grain size distribution of 1 ECAP sample annealed for 180 min.

Figure 7 :

 7 Figure 7: SEM micrography with EDS analyses (points 1 and 2) of 1, 2 and 4 ECAP samples after ECAP processing and annealing for 30 min, respectively.

Figure 8 :

 8 Figure 8: KAM maps and their corresponding misorientation of: (a) 1 ECAP, (b) 2 ECAP and (c) 4 ECAP samples annealed for 5 min.

Figure 9 :

 9 Figure 9: (a, b) ED-IPF, (c, d) KAM, (e, f) IQ highlighted by grain boundaries distribution and (g, h) extension twin maps of selected zones of 1 and 2 ECAP samples annealed for 5 min, respectively.

Figure 10 :

 10 Figure 10: Selected zone of 1 ECAP sample annealed for 180 min showing different forms of twin. (a) ED-IPF, (b) IQ highlighted with extension twin, (c) IQ highlighted with VLAGBs, LAGBs and HAGBs maps and (d) Misorientation profiles of AB and BC distances indicated in ED-IPF map.

Figure 11 :

 11 Figure 11: (a) IPF-map and (b) IQ map highlighted with VLAGBs, LAGBs, and HAGBs of specific zone on 1 ECAP sample annealed for 30 min. Different static recrystallization mechanism observed in this sample: Zone 1: SD mechanism, Zone 2: SD and PSN mechanisms and Zone 3: PSN mechanism.

Figure 12 :

 12 Figure 12: (a) Deformed and (b) recrystallized grains and their corresponding {0002} recalculated pole figure of the 2 ECAP sample annealed for 5 min.

Table 2 .

 2 Fraction of VLAGBs, LAGBs, HAGBs and extension twin in annealed 1 and 2 ECAP samples for 5 min.

	Sample	1°-5°	5°-15°	15°-180°	Extension twin
		(%)	(%)	(%)	(%)
	1 ECAP	70.5	5.0	24.5	18.8
	2 ECAP	59.6	15.1	25.3	0.5

Acknowledgments

The authors thank Dr.-Ing Talal Al-Samman (RWTH Aachen University) for supplying the Mg-RE alloy.

Data availability

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

Conflict of Interest:

The authors declare that they have no conflict of interest.