Gautier Doumenc
email: gautier.doumenc@hotmail.com

Laurent Couturier

Bruno Courant

Pascal Paillard

Alexandre Benoit

Eric Gautron

B Girault

Thilo Pirling

Sandra Cabeza

David Gloaguen

Investigation of microstructure, hardness and residual stresses of wire and arc additive manufactured 6061 aluminium alloy

Keywords: Wire and Arc Additive Manufacturing, 6XXX, Microstructure, Residual stress

Wire and Arc Additive Manufacturing Cold Metal Transfer® enable the defect-free soldering of poorly weldable alloys, such as AA6061. Thin-walls have been successfully built with two deposition strategies that is either a back and forth movement of the torch (alternate deposition vectors) or a repeating movement along one direction (unidirectional deposition vectors). The obtained macro-and micro-structures are highly dependent on the WAAM process parameters and the deposition strategy. Both deposition strategies produce stacked structures consisting of successive alternating layers of columnar and equiaxial grains. The AA6061 alloy contains iron-rich intermetallic compounds and β-phase. The precipitates of the latter have been melted during the WAAM deposition process leading to an even higher hardness as compared to the commercial T6 material. At the nanoscale, Transmission Electron Microscopy analyses reveal β' phase, β'' phase and small dispersoids at T6 state, as well as a multitude of phases (L, Q') in the studied material. It shows the coexistence of β''-and β'phases. The conventional T6 treatment applied to the WAAM material resulted in a slight over-ageing of this material. Neutron diffraction has been used for residual stress measurements showing tensile stresses (up to 130 MPa) in the built parts and compressive stresses (up to -80 MPa) in the substrate. Despite a difference between the generated microstructure (grain shape orientation and subsequent crystallographic texture) by the two building strategies, no significant variation has been observed on their mechanical properties.

INTRODUCTION

Additive Manufacturing (AM) shows a great expansion due to the technological stake it represents. It enables repairing worn-out parts (building up), lightening structures, manufacturing parts with complex geometries and displays a large industrial adaptability [START_REF] Caffrey | 3D Printing and additive manufacturing state of the industry annual worldwide progress report[END_REF][START_REF] Benoit | Soudage à l'arc semi-automatique ( MIG / MAG )[END_REF][START_REF] Jacono | Quality of gas metal arc welds for rapid prototyping and wear replacement[END_REF][START_REF] Debroy | Additive manufacturing of metallic components -Process, structure and properties[END_REF]. Wire and Arc Additive Manufacturing (WAAM) is a DED (Direct Energy Deposition) process that represents an interesting solution for large scale structures requiring low cost fast manufacturing. It combines an electric arc as a heat source and a wire filler as a feedstock. Compared to other AM processes, WAAM techniques provide a high material deposition rate, a relatively low cost and the ability to manufacture large metre-scale components of medium geometric complexity [START_REF] Debroy | Additive manufacturing of metallic components -Process, structure and properties[END_REF][START_REF] Cunningham | Cost Modelling and Sensitivity Analysis of Wire and Arc Additive Manufacturing[END_REF]. However, a final operation of machining is usually required in order to obtain precise geometrical features from the produced near-net-shape part. Most of the time, a material surplus is deposited to foresee this machining operation [START_REF] Silva | Formability of a wire arc deposited aluminium alloy[END_REF][START_REF] Kazanas | Fabrication of geometrical features using wire and arc additive manufacture[END_REF]. As for the other AM processes, a preliminary study focused on microstructure and mechanical properties in relation with process parameters is also mandatory to overcome uncontrolled microstructural evolution, large residual stresses, cracks and geometrical distortions [START_REF] Debroy | Additive manufacturing of metallic components -Process, structure and properties[END_REF].

Thanks to their properties such as high specific strength combined with natural corrosion resistance, age hardening aluminium alloys such as Aluminium Alloy 6061 (AA6061) are widely used in many structural applications. Hence, AA6061 stands as a good candidate when using a high deposition rate process such as WAAM for structural applications. Its mechanical properties mainly stem from the nanometric β''-phase precipitation, an intermetallic compound rich in Mg and Si (Mg5Si6) [START_REF] Hasting | Composition of β″ precipitates in Al-Mg-Si alloys by atom probe tomography and first principles calculations[END_REF]. β''-phase is metastable and generally precipitates in a wellcontrolled manner during dedicated heat treatment but is only one of the many intermetallic phases that can be encountered in this type of alloy. However, the actual microstructure of the WAAM-produced part remains unknown due to the complex out of equilibrium thermal history.

6XXX alloy series are also well-known for their poor weldability [10,[START_REF] Benoit | Développement du soudage MIG CMT pour la réparation de pièces aéronautiques. Application aux pièces en alliage base aluminium 6061[END_REF]. The Al/Mg ratio set around 0.8 makes it highly susceptible to hot cracking and, as all aluminium alloys, it has a natural trend to generate porosities because of the high solubility of hydrogen in the liquid aluminium that drops down to almost zero in the solid aluminium [START_REF] Talbot | Solubility of hydrogen in liquid aluminium[END_REF][START_REF] Opie | Hydrogen solubility in aluminum and some aluminum alloys[END_REF]. It is common to use filler wire metals with other compositions (4XXX or 5XXX alloys) to avoid harmful hot cracking [START_REF]com%2Fus %2Fen%2Fsupport%2Fupload%2FAluminum_Filler_Alloy_Selection_Chart.pdf&usg=AOvVaw3 SDCY2[END_REF][START_REF] Ahmad Ibrahim | Effect of porosity on tensile behaviour of WELDED AA6061-T6 aluminium alloy[END_REF]. The same drawbacks naturally appear in WAAM. In the literature, a substantial work on AM of aluminium alloys with PBF-Laser technique shows that all types of aluminium alloys are printable using this particular process [START_REF] Galy | Main defects observed in aluminum alloy parts produced by SLM: From causes to consequences[END_REF]. In contrast, recent studies on WAAM show that only specific 2XXX [START_REF] Fang | Microstructure evolution and mechanical behavior of 2219 aluminum alloys additively fabricated by the cold metal transfer process[END_REF][START_REF] Gu | Wire+Arc Additive Manufacturing of Aluminium[END_REF][START_REF] Cong | A comparative study of additively manufactured thin wall and block structure with Al-6.3%Cu alloy using cold metal transfer process[END_REF][START_REF] Gu | The effect of inter-layer cold working and post-deposition heat treatment on porosity in additively manufactured aluminum alloys[END_REF][START_REF] Pickin | Characterisation of the cold metal transfer (CMT) process and its application for low dilution cladding[END_REF][START_REF] Icardi | Wire + arc additive manufacturing (WAAM) for large scale components: aluminium alloys 2319 and 5554[END_REF][START_REF] Hönnige | Control of residual stress and distortion in aluminium wire + arc additive manufacture with rolling[END_REF][START_REF] Gu | The strengthening effect of inter-layer cold working and post-deposition heat treatment on the additively manufactured Al-6.3Cu alloy[END_REF][START_REF] Montevecchi | Idle time selection for wire-arc additive manufacturing: A finite element-based technique[END_REF], 4XXX [START_REF] Ortega | Effect of process parameters on the quality of aluminium alloy Al5Si deposits in wire and arc additive manufacturing using a cold metal transfer process[END_REF][START_REF] Qi | Microstructure and mechanical properties of doublewire + arc additively manufactured Al-Cu-Mg alloys[END_REF][START_REF] Köhler | Wire and arc additive manufacturing of aluminum components[END_REF][START_REF] Gu | Deformation microstructures and strengthening mechanisms for the wire+arc additively manufactured Al-Mg4.5Mn alloy with interlayer rolling[END_REF][START_REF] Oyama | Heat source management in wire-arc additive manufacturing process for Al-Mg and Al-Si alloys[END_REF] and 5XXX [START_REF] Gu | The effect of inter-layer cold working and post-deposition heat treatment on porosity in additively manufactured aluminum alloys[END_REF][START_REF] Icardi | Wire + arc additive manufacturing (WAAM) for large scale components: aluminium alloys 2319 and 5554[END_REF][START_REF] Qi | Microstructure and mechanical properties of doublewire + arc additively manufactured Al-Cu-Mg alloys[END_REF][START_REF] Köhler | Wire and arc additive manufacturing of aluminum components[END_REF][START_REF] Oyama | Heat source management in wire-arc additive manufacturing process for Al-Mg and Al-Si alloys[END_REF][START_REF] Yehorov | Exploring solutions to mitigate unidirectional grain growth and to set operational envelop for wire +arc additive manufacturing of aluminium alloys[END_REF][START_REF] Horgar | Additive manufacturing using WAAM with AA5183 wire[END_REF][START_REF] Gomes | Wire-arc additive manufacturing of Al-Mg alloy using CMT and PMC technologies[END_REF][START_REF] Derekar | Influence of Interpass Temperature on Wire Arc Additive Manufacturing (WAAM) of Aluminium Alloy Components[END_REF][START_REF] Elrefaey | Effectiveness of cold metal transfer process for welding 7075 aluminium alloys[END_REF][START_REF] Ayarkwa | Assessing the effect of TIG alternating current time cycle on aluminium wire + arc additive manufacture[END_REF] series can be successfully deposited. Thus, an upstream development is absolutely required to obtain WAAM-built parts using 6XXX and 7XXX series of aluminium alloys [START_REF] Derekar | A review of wire arc additive manufacturing and advances in wire arc additive manufacturing of aluminium[END_REF][START_REF] Wu | A review of the wire arc additive manufacturing of metals: properties, defects and quality improvement[END_REF]. Recent investigations on process parameters enabling the welding of the 6061 aluminium alloy were carried out by Benoit et al. [START_REF] Benoit | Soudage homogène MIG de l'alliage d'aluminium 6061[END_REF] and Kumar et al. [START_REF] Kumar | Investigations on the parametric effects of cold metal transfer process on the microstructural aspects in AA6061[END_REF]. But to date, even though Al-Mg-Si-Cu alloys are widely used for structural applications, there is no detailed studies regarding the ability to use AA6061 for Wire and Arc Additive Manufacturing [START_REF] Debroy | Additive manufacturing of metallic components -Process, structure and properties[END_REF]. The recent development of a Metal Inert Gas (MIG) transfer mode based on a low-energy shortcircuit, called Cold Metal Transfer (CMT®), made possible the use of 6XXX age hardening aluminium alloys for welding [START_REF] Elrefaey | Effectiveness of cold metal transfer process for welding 7075 aluminium alloys[END_REF][START_REF] Benoit | Soudage homogène MIG de l'alliage d'aluminium 6061[END_REF][START_REF] Selvi | Cold metal transfer (CMT) technology -An overview[END_REF][START_REF] Serindag | Investigation of the Influence of Weld Defects on the Mechanical Behaviour of Cold Metal Transfer ( CMT ) and Cold Metal Transfer Pulsed ( CMT-Pulsed ) Arc Welded AA[END_REF][START_REF] Çam | Recent developments in joining of aluminum alloys[END_REF]. However, those studies have been limited to simple welding. To our knowledge, no study on AM with AA6061 has been achieved until now.

Additively built parts using WAAM undergo a complex thermal history: high heating and cooling rates leading to strong out of equilibrium state, subjected to successive solidification, thermal expansion, high-temperature gradients, phase transformations, ageing of precipitated minor phases. This affects both the microstructure and the residual stress distribution. The optimization of MIG-CMT to widen the use of WAAM in miscellaneous industries necessarily involves the in-depth knowledge of the mechanical and metallurgical states of the produced parts to ensure the best in-service performances. Therefore, the present work aims to link microstructural features, hardness and residual stresses in additively manufactured AA6061. The purpose is to investigate the process parameters enabling the built of a defect-free material using WAAM as well as the understanding of the relationship between the particular WAAM built microstructure and the resulting mechanical properties. Two deposition strategies (namely alternate (zigzag) or unidirectional vectors) have also been investigated in order to control the microstructures as shown by Dinda with DED-Laser process [START_REF] Çam | Recent developments in joining of aluminum alloys[END_REF] and by Yehorov et al. with WAAM process [START_REF] Gu | The strengthening effect of inter-layer cold working and post-deposition heat treatment on the additively manufactured Al-6.3Cu alloy[END_REF]. In addition, after a heat treatment, a higher hardness is obtained for WAAM parts as compared to conventional forged products. The microstructure of AA6061 parts produced by WAAM has been evaluated in order to assess and understand these hardening phenomena.

EXPERIMENTAL DETAILS

Filler wire

As mentioned above, AA6061 is considered to be non-weldable, hence there is no commercially available filler wire. AA6061-filler wire has been drawn on demand for the study from a single 6061 aluminium alloy billet. The composition, measured using (Electron Dispersive Spectroscopy) EDS, is presented and discussed in section 3.1.

Deposition process

AA6061 thin-wall samples have been deposited on AA1050 plates with a size of 2×30×250 mm 3 by MIG-CMT process developed by Fronius® [START_REF]CMT -Cold metal transfer: the cold welding process for premium quality[END_REF]. CMT is a specific MIG transfer mode based on dip transfer deposition, characterised by a well-controlled material drop deposition during the short circuit of the wire electrode to the substrate, decreasing the energy required for the material deposition. Hence, it is called cold metal transfer due to the low energy used compared with the other transfer modes available in MIG welding [START_REF] Pickin | Evaluation of cold metal transfer (CMT) process for welding aluminium alloy[END_REF]. Compared with other welding techniques, MIG-CMT seems a suitable candidate for WAAM of age hardening aluminium alloys due to the low applied heat input. Temperature has been measured with three thermocouples positioned within the baseplate to follow the temperature evolution during WAAM.

A Cartesian robot with step-by-step motors equipped with a welding device Fronius® TransPuls Synergic 3200 CMT compose the experimental set-up (Fig. 1(a)). A water-cooling system regulates the bottom of the clamped AA1050 baseplate at 20°C, as described in Figure 1 WAAM is a complex multi-parameter process. Deposition parameters have been chosen by trial and error starting with the standard parameters pre-recorded within the welding stationwelding synergiesand appropriate adjustments. Only parts obtained with suitable parameters to ensure a good material health are considered in this study. Torch speed has been kept constant at 40 cm.s -1 . A parameter optimization protocol based on Mezrag et al. [START_REF] Mezrag | Control of mass and heat transfer for steel / aluminium joining using Cold Metal Transfer process To cite this version : HAL Id : hal-01289341 Control of mass and heat transfer for steel / aluminium joining using Cold Metal Transfer process[END_REF] and Ortega et al. [START_REF] Ortega | Effect of process parameters on the quality of aluminium alloy Al5Si deposits in wire and arc additive manufacturing using a cold metal transfer process[END_REF] works ensuring the optimal stability of the WAAM process (based on the respect of the CMT cycle monitored by a high-speed camera and electrical measurements) allowed us to find the optimal Wire Feed Speeds (WFS) at 4.3 m.min -1 . A commercial mixture composed of 80% argon and 20% helium has been used as a shielding gas [START_REF] Benoit | Soudage à l'arc semi-automatique ( MIG / MAG )[END_REF]10]. A higher wire feed speed (5 m.min -1 ) has been used for the preliminary adhesion layer deposition to prevent necking and delamination from the base plate.

Additionally, two different deposition strategies (namely alternate or unidirectional vectors) have been studied using the same set of process parameters (set-up, heat input level, travel speed, WFS) to understand how these strategies influence the microstructural features and the development of residual stresses within the built part. An adaptive cooling time has been taken in order to reach 20°C in the baseplate before each layer deposition. This time never exceeded 40 s

Heat treatments

6XXX aluminium alloys series are precipitation hardened. The temperature and the duration of ageing enabling an optimal strengthening are approximately 170°C and 8 hours, respectively. This particular heat treatment is known as T6. The microstructure formed during the additive manufacturing process within the produced part is unknown and should present marked heterogeneities. The as-deposited state is therefore studied. This as-deposited microstructure is not stable over time, so we are also interested in characterising the stable microstructure obtained after natural ageing.

Different metallurgical states have been characterised:  As-Deposited (denoted AD): material has been kept at -80°C after printing to prevent natural ageing,  Naturally Aged (NA/T4): material has been kept at room temperature during at least two weeks after deposition (no solutionizing in-between),  Forced Maturation (M): material has undergone a maturation at 90°C during 2 h to evaluate its hardness stability,  Peak aged (T6): material has been solutionized at 530°C during 1 h, then water quenched and finally aged at 175°C during 8 h.

Specimen preparation and equipment

X-ray radiography has been conducted with a GE TITAN E 160 at 160 kV and 40 mA with a pause time of 15 s. The X-ray shot has been carried out with a single-film using the plane to plane technique. Specimen surfaces for porosity investigation and hardness tests were prepared under a standard procedure of surface dressing and prepolishing with series of 180-grit, 500-grit, 800-grit, 1200-grit, and 2000-grit silicon carbide (SiC) papers.

The Vickers microhardness tests were carried out with a STRUERS DuraScan under a load of 0.1 kgf and a holding time of 10 s. Hardness maps were measured using a square grid with a 2 mm step.

Samples for metallographic observations and EBSD specimens have undergone the same preparation, completed by a polishing step with a 3 μm and 1 µm diamond solution. Finally, they were polished using a colloidal neutral silica suspension (0.04 μm) for 8h to reach a strain-free mirror finish. Metallographic samples have been subsequently etched with Barker's electrolytic reagent (5 mL HBF4, 200 mL H2O) during 2 minutes at 30 V. Microstructures have been examined under polarised light using a ZEISS AXIO Imager M2 optical microscope.

Fine microstructure features and intermetallic phases have been characterised employing a ZEISS Merlin Field Emission Gun Scanning Electron Microscope (FEG-SEM) equipped with an energy dispersive spectrometer Oxford X-Max. The latter instrument has been used to determine the chemical compositions of the wire and the WAAM deposited material. Electron Back Scatter Diffraction (EBSD) analyses were performed using the ZEISS Merlin FEG-SEM together with an EBSD Nordlys 2 HKL (Oxford Instruments) probe. An accelerating voltage of 20 kV and a beam current of 1 nA were used for the acquisitions. The results have been collected and processed with the Channel 5 software.

Nanometric scale phases have been studied using Transmission Electron Microscopy (TEM). Samples have been grinded down to 75 µm thick foils. Then, 3 mm-diameter disks were punched out of these foils and prepared for TEM observations by electropolishing with a twin-jet polisher TENUPOL-5 at -25 °C and under 15 V with an etchant composed of 30% commercially pure (around 70%) nitric acid and 70% methanol in volume. A Thermo Fisher Scientific S-TEM Themis Z G3 has been used for fine metallurgical studies. TEM micrographs and electron diffraction patterns were acquired with a CMOS One View camera (Gatan).

Residual stress measurements

To understand the stress development, Neutron Diffraction (ND) experiments have been carried out on SALSA (Strain Analyser for Large and Small scale engineering Applications) beamline at Laue Langevin Institute (ILL). More details of this instrument are available in the work of Pirling et al. [START_REF] Pirling | SALSA-A new instrument for strain imaging in engineering materials and components[END_REF]. The gauge has a volume of 2×2×2 mm 3 and the {311} diffraction peak of aluminium (2θ = 88.5° with a wavelength of 1.62 Å) is considered for the measurements. As presented in Figure 2 (a), the stress field investigation has been achieved along a column in the middle of the thin-wall sample. Measurement points ensure a reliable determination of the stress gradient regarding to the microstructure length scale (Figure 2 (d)). Asymmetrically distributed gauge positions over the height of the wall have been probed. An additional point has been taken at the centre of the baseplate to understand its role in the residual stress equilibrium. Lattice strain have been measured along the three orthogonal directions, matching the sample coordinates system at each measurement point, in order to provide an assessment of the three normal strain components εxx εyy, and εzz. The experimental set-up for the three measurement directions is described in Figure 2.

RESULTS & DISCUSSION

Part geometries

Defect-free AA6061 thin-walls have been successfully built with two different deposition strategies (alternate or unidirectional deposition vectors). The welding parameters have been chosen in order to ensure the CMT cycle stability, so that no spatter was produced during material deposition. Geometries of the deposited walls have been measured using a Mitutoyo 500 digital calliper. The aim was to produce walls of 150×5×50 mm 3 dimensions once machined. To achieve this, as reported in the introduction, a 1 mm additional thickness is added to the targeted 5 mm to allow for the finishing machining step [START_REF] Silva | Formability of a wire arc deposited aluminium alloy[END_REF][START_REF] Kazanas | Fabrication of geometrical features using wire and arc additive manufacture[END_REF]. Figure 3 illustrates the as-deposited wall geometries achieved for both strategies prior to the final machining operation. Figure 3 presents two of the built samples: one for the alternate strategy (a and b) and one for the unidirectional strategy (c and d). One can observe from the pictures of the walls (i.e. Figure 3.a and Figure 3.c) that the shape of the built walls is different. Indeed, the slight excess of material at the start of the welding bead and the slight lack of it at the end, added over the successive bead depositions for the unidirectional strategy, lead to an asymmetric shape with a large excess of material on the starting side and a large lack of it on the ending side of the welding beads. This effect has already been reported in the literature by Zhang et al. [START_REF] Zhang | Weld deposition-based rapid prototyping: A preliminary study[END_REF]. As proposed by the authors, we can observe on Figure 3.a that the use of an alternate strategy enables to compensate these excesses and lacks of material over two welding beads in opposite directions, leading to a final wall with a correct symmetric shape, as desired. It may be noted here that the choice of a building strategy strongly modifies the grain structure within the building. This aspect, shown in section 3.5, especially using EBSD, will be discussed later in section 4 since it may have a strong influence on the mechanical strength of the resulting part and has to be taken into consideration. The effective wall thickness is 5 mm and the surface waviness is 0.8 mm [START_REF] Kazanas | Fabrication of geometrical features using wire and arc additive manufacture[END_REF]. There is no preferable strategy regarding the surface waviness and the effective wall thickness. Therefore, we aimed for a 6 mm as-built mean thickness.

Chemical composition

The global composition of the deposit has been measured using EDS and compared to the composition measured in the filler wire. The results presented in Table 1 show that both compositions are close. Preferential evaporation from the weld pool of some elements, presenting a high vapour pressure such as Zn, Mn and Mg, can be significant regardless of the temperature [START_REF] Yuan | Loss of elemental Mg during wire + arc additive manufacturing of Al-Mg alloy and its effect on mechanical properties[END_REF][START_REF] Block-Bolten | Metal vaporization from weld pools[END_REF]. Even if several elements could be vaporized, sample composition prepared by WAMM remains within the nominal composition range given by the standard for AA6061, except for Si (excess) and Cu (depletion) contents [START_REF] Brandes | Smithells Metals Reference Book: Seventh Edition, Smithells Met[END_REF]. However, the Si content was already out of range in the filler material and only slightly varies. Cu content, although remaining in the composition range for the wire filler, was in the low range and its slight variation may be attributed to evaporation that causes a slight depletion with respect to the standard.

Material health

X-Ray inspection

The visual inspections on machined samples showing neither cracks nor obvious porosities, X-Ray inspection has been carried out to detect any defects that would not be visible to the naked eye (within the material or on the surface but too small to be detected). Walls have been machined on both sides to avoid artefacts due to the raw surface condition. As we can see in Figure 4, walls built with both strategies are almost free from defects. Only a limited number of porosities can be found in the very first layer (close to the baseplate), which is the adhesion layer. This adhesion layer has been built with parameters reducing the CMT stability in order to enhance the wettability and to counteract the very high thermal extraction from the baseplate. Porosities are also present near the top surface in the last layer, being the only welding bead not reheated and partially melted by a coming up welding pass. Hence, the largest porosities are found in that last deposited layer which is more likely to be machined for most industrial applications to get the net shape of the additively manufactured part.

Optical microscope investigations

Porosity measurement has been performed by image analysis with the Trainable Weka Segmentation of Fiji software [START_REF] Arganda-Carreras | Trainable Weka Segmentation: A machine learning tool for microscopy pixel classification[END_REF]. Evaluation has been carried out on macrographs along the longitudinal direction (XZ plane) and extracted in the centre region of the specimen to evaluate the efficiency of the chosen parameters. Based on OM observations, the porosity fraction appears very low (< 0.2%) and the maximum measured pore size is 170 µm. One reason stands in the careful cleaning and degreasing of the baseplate prior to the first layer deposition. This step is very important in order to avoid pollutants like grease to be cracked within the electric arc producing hydrogen that can be solutionized in the liquid aluminium and then leading to the apparition of porosities upon solidification when the hydrogen solubility drops down to almost zero. An additional cause for the low fraction of porosities is the use of a He-rich shielding gas making possible to raise the arc plasma temperature, maintaining a less viscous melt pool and thus a better hydrogen degassing from the liquid pool. Porosities can however still be found in the first layer which is an adhesion layer made with a higher wire feed speed (5 m.s -1 ) and in the last layer due to incomplete hydrogen degassing.

Hardness

As can be seen in Figure 6, the hardness is relatively constant throughout the thickness (i.e. along the Y direction) of the samples regardless of the building strategy. The deposition strategies seem to lead to the same hardness values. The different temperature gradients, solidification rates as well as the complex thermal effects which appear during the process seem to have no influence on the hardness values of each layer. Due to the material high thermal conductivity and the time between layers, it is possible that the last layer leads to homogenise the metallurgical state giving rise to a homogeneous hardness. Table 2 gives a summary of the hardness of the built material in the different metallurgical states investigated in this work. As a result, for a given heat treatment the two strategies give the same hardness. As-deposited material has an average hardness of 64 HV0.1 approaching the hardness found for the quenched material. The deposited material is probably quenched due to the high cooling rates. Indeed, it goes from liquid temperature to ambient temperature in few seconds. The other successive layers are rapidly deposited and the time spent in the high to moderate temperature range is brief. Thus, the alloying element mobility is not sufficient to precipitate hardening phases. The as-deposited part is in the solutionized state. Therefore, hardness is low because the only hardening mechanism present in this case is the solid solution strengthening. Naturally aged material, obtained after two weeks at room temperature, has an average hardness of 75 HV0.1 corresponding to the value found for the T4 state. A forced maturation heat treatment of 2 hours at 90°C did not allow to modify the hardness further, highlighting that metallurgical stability is reached in the material after the natural ageing. Lastly, the T6 heat treatment carried out on the walls leads to higher hardness (124 HV0.1 for the alternate strategy and 118 HV0.1 for the unidirectional strategy) than that obtained on forged material (equal to 107 HV0.1 according to ASM standard [START_REF] Asm | Metals Handbook: Vol 2 -Properties and selection : Nonferrous alloys and special-purpose materials[END_REF]). The unidirectional strategy gives a hardness close to the alternate one. It lies within the dispersion of measurements (5 to 10 HV0.1).

Microstructural characterization

Microstructure and macrostructure observations by optical microscopy

As illustrated in Figure 6 (a), two types of microstructures are revealed: columnar and equiaxed grains with OM. Both deposition strategies induce the same stacking structure of alternate layers of equiaxed and columnar grains. The substrate (i.e. the baseplate) can be considered as equiaxed. Each layer in the built walls is constituted of columnar grains growing from the bottom of the layer with, on top, a finer layer of equiaxed grains as schematised on Figure 6 (a). Each new built layer partially melts the previous one. The characteristic depth of the remelted layer has been measured around 200 µm. The measurements of the heat-affected zones and molten zones were determined during a test campaign in which different settings were studied. A first measurement was carried out on a single layer and then on 2, 3, 4, 5, 7 and finally 10 layers. Analysis of heat-affected zones and molten zones have been achieved thanks to optical microscope on chemically attacked polished cross-cut samples. Due to the low heat input produced by the CMT technology, only a very narrow heat affected zone of about 10 µm in depth is present where a very slight grain coarsening is observed. The same observations have been made on longitudinal-sections (plane X-Z) on samples cut from the wall thickness. Such an observation of a sample built with the alternate strategy is presented in Figure 6 (b) while Figure 6 (c) displays it for a sample produced with the unidirectional. Although the two strategies give similar microstructures in cross-section, they influence the direction of the columnar-grain growth in the longitudinal plane. Indeed, the growth direction of the columnar grains is linked to the welding direction (via the heat gradient developing during the welding operation). Thus, the alternate strategy (Figure 6 (b)) -alternating between two opposite welding directions -produces columnar grains which grow along two directions of solidification symmetrical to the transverse plane (Y-Z). The unidirectional strategy (Figure 6(c)) retains the same solidification direction layer after layer. For both strategies, the equiaxed grain proportion seems to be equivalent in terms of size distribution. These microstructural observations in the X, Y and Z directions allow to evaluate the morphology of the grains. In the first instance, the grain size is measured according to the intercept method presented in the ASTM E112 standard [START_REF] Astm | Standard test methods for determining average grain size[END_REF]. The use of this method is the most appropriate to evaluate the size of the columnar grains as well as their aspect ratio. As detailed in this standard, the three main measurement directions have been identified. The mean lineal intercept length l ̅ , for columnar grains, is obtained from the equation [START_REF] Caffrey | 3D Printing and additive manufacturing state of the industry annual worldwide progress report[END_REF].

l ̅ = √l 𝐿 ̅ . l 𝑇 ̅ . l 𝑃 ̅ 3 (1)
Where l 𝐿 ̅ , l 𝑇 ̅ and l 𝑃 ̅ are the mean lineal intercept length on the longitudinally, transversely and planar oriented surface, respectively. For equiaxed grain structure, the grain size is obtained more directly through the mean lineal intercept length. Measurements in the longitudinal plane are reported in Table 3. Concerning the equiaxed areas, whatever the strategy being used, the grain size remains very similar around 100 µm. However, considering the columnar grains, these measurements show that the unidirectional strategy produces slightly longer grains than the alternate one, the mean intercept in the two perpendicular directions being almost equal for both strategies, leading to an aspect ratio of 4.7 for the unidirectional strategy and 3.2 for the alternate one.

EBSD Analysis

Grain size, shape and crystallographic distributions have been evaluated through EBSD and compared with the OM measurements. EBSD inverse Pole Figure (PF) maps and PF measured on samples built using each strategy are presented in Figure 7. They corroborate the previous analysis and provide additional information. Whatever the strategy, deposited parts are composed of two alternate microstructure zones: equiaxed and columnar dendritic grains. This hierarchical stacking of grains morphologies can explain the narrow distribution of microhardeness values along the Z axis as seen on Figure 5. It may be noted that the columnar grains are elongated along the solidification direction and have a significant aspect ratio (see Table 4). These grains grow in epitaxy on the previous layer so they did not go through a nucleation stage. This phenomenon occurs because each layer has the same composition and the same crystallographic structure. Growth can then occur spontaneously without requiring activation energy as soon as the temperature drops below the liquidus [START_REF] Kou | Welding Metallurgy[END_REF][START_REF] Basak | Epitaxy and Microstructure Evolution in Metal Additive Manufacturing[END_REF]. The phenomenon of competitive growth explains the preferential grain growth direction leading to the columnar shape of the grains. This appears during solidification in polycrystalline materials with dendrites with various crystallographic orientations. Some dendrites have preferred growth orientations at the liquid / solid interface along the maximum heat flow direction.

The favoured growth direction depends on the crystal structure, this direction for face-centred cubic (FCC) materials, as aluminium based alloys, is <100> [START_REF] Kou | Welding Metallurgy[END_REF][START_REF] Dinda | Evolution of microstructure in laser deposited Al-11.28%Si alloy[END_REF]. The unidirectional strategy gives a single grain growth direction causing a sharp crystallographic <100> texture in the direction X+65° as illustrated in the PF in Figure 7(a). The alternate strategy gives a more balanced texture with two preferential solidification directions (Figure 7(b)). As shown in Figure 8(a), the n layer displays a marked <100> texture component in the X+65° direction while the n+1 layer shows a <100> texture component with the same intensity in the X-65° direction. Therefore, while the overall texture obtained with the alternate strategy seems more homogeneous, we can observe that each layer exhibits a strong <100> crystallographic orientation enhancement. The grain morphologies determined using EBSD correspond to the results obtained with OM. The observed grain size distributions determined for both deposition strategies are presented in Figure 9 as a histogram. One can see that the alternate strategy produces a higher proportion of fine grains than the unidirectional one. A closer look at grains with an equivalent diameter larger than 15 µm, shows that the unidirectional strategy produces more of those large grains. Although coherent, the average grain sizes observed by the EBSD technique are much smaller than those observed by OM. This may be explained by the fact that very small grains of size less than 2 µm, which are not measurable with OM technique, were considered for the determination of size distributions by EBSD. In addition, EBSD results may also be affected by preparation artefacts and indexation errors. The deposition conditions promoted the formation of an equiaxed grain layer at the top of each deposited layer. The heterogeneous nucleation leading to this equiaxed layer may stem from constitutive phases that are able to remain unmelted and form the non-dendritic equiaxed zone. It may also stem from torn out arms of dendrites. Indeed, due to high welding speed, important liquid flow in the melting pool is emerging which can sweep dendrite branches away [START_REF] Kou | Welding Metallurgy, Second edi[END_REF]. These dendrite arms, torn out from the dendrites by important liquid flow within the melting pool, can serve as heterogeneous nucleation sites for the growth of new grains with random orientations in the rapid cooling top layer of liquid. Therefore, equiaxed grains can be found at the top of each layer. In the present study, dispersoids or constitutive phases have been rarely found at the centre of the equiaxed dendritic grains on the crosssections. The assumption of the dendrite arm detachment should therefore not be ruled out. It is also possible that the two phenomena are both participating to the formation of these equiaxed grains. Equiaxed grain layer ensures a large amount of grain boundary surfaces, hence reducing the hot cracking formation. Thus, their presence is most likely responsible for the absence of hot cracking in the WAAM 6061 aluminium alloy. An alternative assumption, explaining the hot cracking absence, is that the CMT process enables the rapid filling of the inter-dendritic spaces, fast enough to accommodate the stresses applied during the cooling of the weld bead.

Micrometric phases (intermetallic and β-phase)

Micrometric phases have been studied through SEM observations. Back-Scattered Electrons (BSE) contrast is strongly related to the mean atomic number of the irradiated zone. Therefore, it enables to highlight the presence of composition variations at a scale down to a few hundred of nanometres. These composition variations in metallic alloys are usually due to the presence of secondary phases distributed within the material matrix. As we can see in Figure 10, two minor phases are present in the additive manufacturing AA6061-wire, a globular dark-grey phase and a more polygonal white phase. We have measured the composition of one precipitate of each of these minor phases using local EDS analysis (Figure 10(a)), whose results are displayed on the diagram presented in Figure 10(b). Their compositions enable to identify the two types of observed minor phases: the globular dark-grey phase is the β-phase (Mg2Si) and the more polygonal white phase rich in Fe and Si is an iron InterMetallic phase (IM). The SEM images in BSE contrast in Figure 11, shows the AA6061 alloy at different metallurgical states being: (a) before the WAAM (wire), (b) as-deposited and (c) after an additional T6 heat treatment. In the wire filler, the microstructure is similar to the one found for conventional forged AA6061. The globular dark-grey β-phase and the polygonal white iron-rich IM phases are both present. In the as-deposited material, a silicon-rich phase network is observed in light-grey. This network is formed in interdendritic spaces and has already been deeply studied in the foundry field [START_REF] Samaras | Modelling of microsegregation and homogenization of 6061 extrudable Al-alloy[END_REF]. β-phase is also observed in proportion and in size equivalent to that observed in the wire. The T6 heat treated material no longer presents the previously observed network of Si-rich phase. In addition, all the intermetallic and β-phase are fewer and smaller than those found in the wire and the as-deposited material. Two main precipitated phases are commonly found in AA6061 [START_REF] Flament | Etude des évolutions microstructurales sous irradiation de l'alliage d'aluminium 6061-T6[END_REF]:

 Coarse β-Mg2Si precipitates are formed during the cooling of the alloy and dissolved thanks to the stages of homogenisation/solutionizing heat treatment. However, too slow quenching after solutionizing results in heterogeneous precipitation of this phase at the grain boundaries. Due to their large size and low number density, these precipitates do not take part in the structural hardening of the alloy. Precipitation of coarse β particles even contributes to the limitation of Al-Si-Mg alloys hardening since it decreases the amount of Mg and Si available for nanoscale precipitation of the hardening nanophases [START_REF] Flament | Etude des évolutions microstructurales sous irradiation de l'alliage d'aluminium 6061-T6[END_REF].  Iron-rich intermetallic (IM) phases are characterised by the presence of iron. The αc-Al12(FexMn1-x)3Si phases are formed during the homogenisation heat treatment. In a similar way, as the β precipitates, iron-IM are coarse and then do not participate directly to the material hardening [START_REF] Flament | Etude des évolutions microstructurales sous irradiation de l'alliage d'aluminium 6061-T6[END_REF].

After a T6 heat treatment, the silicon-rich phase network is no longer visible and micrometric phases (iron-IM and β-phase) are smaller and fewer than in the wire. The iron-IM and β-phase have probably been melted within the electric arc during deposition and, due to the extremely high cooling rate and the induced rapid solidification, new particles could not nucleate and grow significantly. Due to very low solubility in solid aluminium, iron segregates within the interdendritic spaces (solidified last drops of liquid). After the T6 heat treatmentand especially after its first step of solutionizing -, Fe becomes more homogeneously distributed within the material, and all the β-phase precipitates have been dissolved in the -Al matrix. -Al matrix, enriched with alloying elements, could significantly change the precipitation kinetics leading to different precipitation states after T6 heat treatment [START_REF] Myhr | Modelling of the microstructure and strength evolution in al-mg-si alloys during multistage thermal processing[END_REF].

Hardened precipitates characterization

AA6061 is an age hardening aluminium alloy, its high mechanical properties stem thus from its precipitation state. Therefore, minor phase nature, spatial and size distributions play an important role. These phases appear during a controlled heat treatment -to master these three aforementioned factors -and are widely influenced by the alloy composition and its thermomechanical history. However, additively manufactured parts undergo a complex thermal history possibly altering the alloy composition (i.e. vaporisation as seen in 3.2) and causing the evolution of the micrometric and nanometric precipitated phases. Hence, it is important to investigate the nature, the localization and the size of these minor phases present in AA6061 additively manufactured parts and their evolution during subsequent heat treatments.

The nanometric precipitates have been observed by TEM. Specimens have been cut out from the middle region of the parts built by WAAM. The same precipitation state has been observed in the parts produced using both building strategies. Thus, we report in this paper results obtained from samples built using only the alternate strategy. Even if the nanostructures in AA6061 produced by traditional route of forging have been the subject of numerous works [START_REF] Hasting | Composition of β″ precipitates in Al-Mg-Si alloys by atom probe tomography and first principles calculations[END_REF][START_REF] Marioara | The influence of alloy composition on precipitates of the Al-Mg-Si system[END_REF][START_REF] Flament | Etude des évolutions microstructurales sous irradiation de l'alliage d'aluminium 6061-T6[END_REF], an AA6061 sheet in the T6 state stemming from this route has also been studied as a reference to compare with. Two main types of precipitates have been observed: dispersoids and nanometric phases (β'', Q', L…).

Dispersoids

Dispersoids are incoherent with the matrix. They have a characteristic length in the range of 10 to 500 nm and a spherical or parallelepipedal morphology. Figure 12 shows bright-field TEM images of dispersoids for (a) the T6 forged material (reference), (b) the as-deposited WAAM material and (c) the T6 WAAM material. The observation of the T6 forged material (Figure 12(a)) highlights the presence of dispersoids with various sizes and shapes. These dispersoids are coarse (60 -600 nm), incoherent with the matrix and homogeneously dispersed within the grains. Figure 12 (b) shows a bright-field TEM image of the as-deposited WAAM material. One can observe the segregation of a Si-rich phase along the grain boundaries, part of the Si-rich phase network presented in section 3.5. It could be noted the presence of very large dispersoids with a rather low number density, way lower than the one observed in the reference material. As explained earlier, this phenomenon is encountered when long soaking durations at rather high temperatures are undergone by the material. Additively manufactured material underwent temperatures up to 400°C, which could explain the presence of large dispersoids with low number density. In Figure 12 (c), we observe smaller dispersoids in the T6 WAAM material. They are in lower quantity and their shape are still widely varying. Other small phases appear: large β'-phase and coarse Q'-phase are observed. The characteristics of those phases are presented in section 3.6.2.

Dispersoids are usually formed during the homogenisation heat treatment, their size varies depending on its duration and temperature [START_REF] Kenyon | Effect of Dispersoids on the Microstructure Evolution in Al-Mg-Si Alloys[END_REF]. Indeed, higher temperatures and longer holding times promote their size to increase and their number density to decrease [START_REF] Strobel | Dispersoid phases in 6xxx series aluminium alloys[END_REF]. Although the primary function of dispersoids is to improve fracture toughness by limiting grain growth, this type of precipitates has a rather limited direct effect on the alloy strengthening [START_REF] Strobel | Dispersoid phases in 6xxx series aluminium alloys[END_REF][START_REF] Lodgaard | Precipitation of dispersoids containing Mn and/or Cr in Al-Mg-Si alloys[END_REF]. Their effect on mechanical behaviour comes from the following phenomena: they may act as nucleation sites for β-particles (Mg2Si) after homogenisation. Therefore, a high number density and a uniform size distribution of small β-particles lead to a better dissolution during further heat treatment prior to the final age hardening step [START_REF] Remøe | A methodology to assess the spatial distribution of dispersoids in Al-Mg-Si alloys[END_REF]. High temperature (up to 400°C) are repetitively applied during WAAM process, which could explain the presence of large dispersoids. This configuration is not favourable to obtain suitable mechanical properties in terms of toughness. Prior to the homogenisation step of the T6 heat treatment, a uniform distribution of small dispersoids with high number density is preferred because this can lead to the precipitation of a higher volume fraction of strengthening particles after ageing step of the T6 heat treatment.

Nanometric phases

Nanometric β''-phases are the main hardening phases of the alloy. These precipitates, also called in the literature "GPII" [START_REF] Vissers | The crystal structure of the β′ phase in Al-Mg-Si alloys[END_REF], heterogeneously nucleate on GPI [START_REF] Murayama | Pre-precipitate clusters and precipitation processes in Al-Mg-Si alloys[END_REF]. They are needle-shape of around 4 nm in diameter and up to 10 nm long [START_REF] Vissers | The crystal structure of the β′ phase in Al-Mg-Si alloys[END_REF]. The β" needles lie in the direction <001> of the aluminium matrix, and their number density, in the AA6061 T6, is around 10 22 /m 3 according to Flament [START_REF] Flament | Etude des évolutions microstructurales sous irradiation de l'alliage d'aluminium 6061-T6[END_REF]. The stoichiometry of the β"-precipitates, according to Zandbergen et al. [START_REF] Zandbergen | Structure determination of Mg5Si6:particles in Al by dynamic electron diffraction studies[END_REF], is Mg5Si6. This phase, in the form of small size precipitates coherent with the surrounding matrix and finely dispersed within the latter, is usually associated with the mechanical resistance peak of the AA6XXX. The temperature and the duration of the ageing heat treatment enabling these phases to be obtained is approximately 170°C for less than 10 hours (175 °C for 8 h is commonly used as heat treatment) [START_REF] Myhr | Modelling of the age hardening behaviour of Al-Mg-Si alloys[END_REF]. Over-ageing of the material would lead to the further transformation of the metastable β"-phase, mainly into the less metastable β'-phase, which is well documented [START_REF] Marioara | The influence of alloy composition on precipitates of the Al-Mg-Si system[END_REF]. The transformation of the β"-phase towards the β'-phase takes place at the very beginning of the transformation; these two phases coexist until the complete dissolution of the β''-phase.

β'-precipitates have a rod shape with a characteristic dimension of its long axis of about 100 nm and contrary to the β''-precipitates they are only semi-coherent with the matrix. Vissers et al. [START_REF] Vissers | The crystal structure of the β′ phase in Al-Mg-Si alloys[END_REF] determined their structure to be hexagonal and their stoichiometry to be Mg9Si5. This phase finally evolves during further over-ageing to form the stable FCC β-phase, of Mg2Si stoichiometry. It results in rather large precipitatesplatelets of micrometric characteristic sizewith lower number density and incoherent with the surrounding matrix leading to the softening of the alloy. The latter can only be observed by orienting the crystal along the axis zone <001> [START_REF] Hasting | Comparative study of the β″-phase in a 6xxx Al alloy by 3DAP and HRTEM[END_REF], but since TEM enables only local characterisation, it is difficult to establish a statistic description of the size and volume fraction of these precipitates The typical diameter of the precipitates is around 4 nm and their characteristic length is around 50 nm which is the usual size according to Vissers et al. [START_REF] Vissers | The crystal structure of the β′ phase in Al-Mg-Si alloys[END_REF]. β'-phase precipitates are more rarely found in this material condition. They are distinguished from the β''-precipitates by their size: their diameter is above 10 nm and their length ranges between 100 and 500 nm according to Myhr et al. [START_REF] Myhr | Modelling of the microstructure and strength evolution in al-mg-si alloys during multistage thermal processing[END_REF]. One can also found large dispersoids in the matrix, but no other phase is observed. Figure 13(b) shows the TEM image of the as-deposited material. β''-and β'precipitates are found nearby the large dispersoids. At a distance of 1 µm far from these large dispersoids, precipitates are no longer observed. β''-precipitates have also been observed in the T6 WAAM material as shown in Figure 13(c). Isolated precipitates were measured with High-Resolution TEM (HR-TEM). The obtained β''-precipitates have a length set between 5 and 200 nm and a diameter between 2 and 10 nm. Simultaneous presence of large β''-precipitates and β'precipitates is the sign of a slight over-ageing undergone by the material. Even if the material is homogenised during the first step of the T6 heat treatment, the β''-phase does not have the dimensions usually reported for the T6 material [START_REF] Vissers | The crystal structure of the β′ phase in Al-Mg-Si alloys[END_REF]. β''-phase is not the only nanoscale phase present in the material. A broad range of metastable phases can be found in Al-Mg-Si-Cu alloys. These phases are formed from particular compositions in the T6 state and can coexist with β''-phase [START_REF] Saito | The effects of low Cu additions and predeformation on the precipitation in a 6060 Al-Mg-Si alloy[END_REF][START_REF] Marioara | Improving thermal stability in Cu-containing Al-Mg-Si alloys by precipitate optimization[END_REF]. L-and Q'phases are observed in WAAM deposited material as illustrated in Figure 14. These phases are not found in the forged material (Figure 13) for the same heat treatment. Hardening β-type precipitates have been studied, showing the coexistence of β''and β'phases. This may be due to compositional fluctuations (matrix enrichment of the alloying elements) as well as the complex thermal cycles applied during the process. In addition, the dispersoids being observed in the previous section are quite different as compared to the reference. Their size and volume fraction may influence the precipitation kinetics. These observations lead us to the conclusion that the conventional T6 treatment applied to the WAAM material resulted in a slight over-ageing of this material. L-phase is considered as a hardening phase and Q'-phase as a softening precipitate. It is therefore possible that the complex coexistence of these different phases (GPI, β'', β', L, Q', dispersoids, etc.) could induce significant hardening of the material. It is however difficult to clearly explain the formation mechanism as well as the role of each of these phases taken independently on the global mechanical behaviour.

Residual stress analysis

Residual stresses, developed during the building process, clearly have a significant influence on the mechanical performance and can lead not only to delamination from the supporting structures or unwanted shape distortion but also to early crack formation [START_REF] Martina | Residual stress of as-deposited and rolled wire+arc additive manufacturing Ti-6Al-4V components[END_REF]. WAAM process develops large residual stresses depending on the material, the desired geometry and the clamping conditions. The origin of residual stresses is the non-uniform thermal contraction and expansion of the material during the different thermal cycles related to this process. In the literature, Coules et al. [START_REF] Coules | Experimental measurement of biaxial thermal stress fields caused by arc welding[END_REF] describe the tensile residual stress formation in welded joints through material shrinkage phenomena that occur during the cooling. Several studies, focused on residual stresses within parts made by WAAM, are available [START_REF] Martina | Residual stress of as-deposited and rolled wire+arc additive manufacturing Ti-6Al-4V components[END_REF][START_REF] Szost | A comparative study of additive manufacturing techniques: Residual stress and microstructural analysis of CLAD and WAAM printed Ti-6Al-4V components[END_REF][START_REF] Colegrove | High pressure interpass rolling of Wire + Arc additively manufactured titanium components[END_REF]. As explained by Szost et al. [START_REF] Szost | A comparative study of additive manufacturing techniques: Residual stress and microstructural analysis of CLAD and WAAM printed Ti-6Al-4V components[END_REF], when molten metal is solidifying on a substrate, the thermal contraction generates tensile stresses in the deposited layer and compression stresses in the substrate. Additively manufactured structures consist of a layer-by-layer building, accumulating residual stresses within each layer. Therefore, high stress level can be obtained during the building process and trapped within the material as residual stresses. To understand the stress development, ND experiments have been carried out on SALSA beamline at ILL (see section 2.5). The measured average lattice strains within the diffracting volume are calculated from the Bragg's angle shift (Eq. ( 2)) based on the true (rational) strain definition [START_REF] He | Two-Dimensional X-Ray Diffraction[END_REF]. For these experiments, the strain-free Bragg's angle, 𝜃 0 , is estimated through an averaging of 𝜃 values over the sample height in the three orthogonal directions X, Y and Z [START_REF] Colegrove | High pressure interpass rolling of Wire + Arc additively manufactured titanium components[END_REF].

𝜀 𝑖𝑖 = ln ( 𝑠𝑖𝑛𝜃 0 𝑠𝑖𝑛𝜃 ) i =x, y, z (2)
This assumption, which supposes that the mean stress in the whole part is equal to zero, due to the residual stress equilibrium, does not consider a possible microstructure evolution. ND data was processed using the LAMP software [START_REF] Richard | Analysis and visualisation of neutron-scattering data[END_REF] to obtain the peak positions. The correlation between residual normal stresses 𝜎 𝑥𝑥 , 𝜎 𝑦𝑦 , 𝜎 𝑧𝑧 and the 3 orthogonal measured lattice strains along the macroscopic coordinate axes is given by Eq. ( 3).

𝜎 𝑖𝑖 =

(2𝑆 1 (ℎ𝑘.𝑙)+

𝑗 = 𝑥, 𝑦, 𝑧𝑗 ≠ 𝑖 1 2 𝑆 2 (ℎ𝑘. 𝑙) and 𝑆 1 (ℎ𝑘. 𝑙) are the X-ray Elastic Constants (XEC) [START_REF] Hauk | Structural and Residual Stress Analysis by Nondestructive Methods[END_REF]. They depend on the measured {hkl} reflection and consider the elastic anisotropy of the diffracting grain group. The XEC have been calculated from single-crystal elastic constants (c11 = 108.2 GPa, c12= 60 GPa and c44 = 28.5 GPa [START_REF] Gu | Deformation microstructures and strengthening mechanisms for the wire+arc additively manufactured Al-Mg4.5Mn alloy with interlayer rolling[END_REF]) using an elastic-self consistent model [START_REF] Gloaguen | Measurement and prediction of residual stresses and crystallographic texture development in rolled Zircaloy-4 plates: X-ray diffraction and the selfconsistent model[END_REF].

The calculated XEC values of Alternate or unidirectional deposition strategies give similar stress levels and the same type of variations in all the three directions. Slightly higher stress levels are obtained in the longitudinal direction. This result has already been observed in the literature [START_REF] Hönnige | Control of residual stress and distortion in aluminium wire + arc additive manufacture with rolling[END_REF][START_REF] Martina | Residual stress of as-deposited and rolled wire+arc additive manufacturing Ti-6Al-4V components[END_REF][START_REF] Szost | A comparative study of additive manufacturing techniques: Residual stress and microstructural analysis of CLAD and WAAM printed Ti-6Al-4V components[END_REF]. As shown in Figure 15,a Longitudinal direction follows the global trend. For alternate strategy, large compressive stresses (-129 MPa) are found in the middle of the baseplate (z = -4 mm) followed by maximum tensile stress (+80 MPa) in the first layers of the wall (z = 5 mm). The second maximum tensile stress is observed in the last layers (z = 37 mm ; +87 MPa). Unidirectional strategy also shows large compressive stresses (-116 MPa) in the middle of the baseplate (z = -4 mm) followed by maximum tensile stress (+78 MPa) in the first layers of the wall (z = 3 mm). The second maximum tensile stress is determined in the last layers (z = 36 mm ; +66 MPa). For both strategies, the transition zone between the baseplate and the wall is free from stress. Concerning the transversal direction, after a maximum reached at the beginning, tensile stress remains still globally low. To complete the description, one should precise that one single measurement point deviates significantly towards a high tensile stress in one case and toward a high compressive stress in the other. The mechanical properties of the specimen at ambient temperature have been characterised by monotonic tensile tests: RE = 189 ± 38 MPa (yield strength) and 𝜎UTS = 344 ± 43 MPa (ultimate tensile strength). In our study, stress level is always smaller than the yield strength of the samples: the stress component 𝜎 𝑥𝑥 = +/-90 MPa (z = 4 mm) reaches 47% of the yield strength.

In order to reduce the residual stresses and depending on the part design and the overall desired aim, it is sometimes possible to remove the baseplate [START_REF] Szost | A comparative study of additive manufacturing techniques: Residual stress and microstructural analysis of CLAD and WAAM printed Ti-6Al-4V components[END_REF]. Hence, one sample was machined and removed from the baseplate to understand how it impacts the residual stress equilibrium. Figure 15 (c) shows the residual stresses in the sample produced with the unidirectional strategy after baseplate removal. A stress redistribution is then observed. The global trend found before baseplate removal shifts and a magnitude reduction is observed (i.e. z = 4 mm ; 𝜎 𝑥𝑥 ≈ +90 MPa before removal ; z = 4 mm ; 𝜎 𝑥𝑥 ≈ +49 MPa after removal). As found before substrate removal, similar stress levels and are found in all directions. From the first maximum compressive stress at the starting point (i.e. z = 0 mm ; 𝜎 𝑖𝑖 ≈ -87 MPa) , the residual stresses increase to reach a maximum tensile stress in the fourth layer (z = 8 mm ; 𝜎 𝑖𝑖 ≈ +45 MPa). The highest tensile stress is observed in the last layer (z = 36 mm ; 𝜎 𝑖𝑖 ≈ +57 MPa). After baseplate removal, a stress reduction up to 30% is observed.

Residual stresses developed during the building process have clearly an influence on the mechanical performance [START_REF] Martina | Residual stress of as-deposited and rolled wire+arc additive manufacturing Ti-6Al-4V components[END_REF]. As explain by Szost et al. [START_REF] Szost | A comparative study of additive manufacturing techniques: Residual stress and microstructural analysis of CLAD and WAAM printed Ti-6Al-4V components[END_REF], a general trend, for which stresses are rather tensile in the wall and compressive in the baseplate, is usually found for the stress state reached after deposition. Once unclamped, the component tends logically to contract at the top and expand at the bottom. For this reason, bending takes place together with a stress redistribution. According to Martina et al. [START_REF] Martina | Residual stress of as-deposited and rolled wire+arc additive manufacturing Ti-6Al-4V components[END_REF], during stress redistribution, stress magnitude decreases along the height of the building and can even become compressive at the top. The stress in the baseplate also evolves in the opposite way until being tensile at the bottom of the baseplate. Residual stress measurements show that alternate or unidirectional deposition strategies give similar stress levels and the same kind of variations in all measurement directions. Slightly higher stress levels are obtained in the longitudinal direction, which have already been observed in the literature [START_REF] Hönnige | Control of residual stress and distortion in aluminium wire + arc additive manufacture with rolling[END_REF][START_REF] Martina | Residual stress of as-deposited and rolled wire+arc additive manufacturing Ti-6Al-4V components[END_REF][START_REF] Szost | A comparative study of additive manufacturing techniques: Residual stress and microstructural analysis of CLAD and WAAM printed Ti-6Al-4V components[END_REF][START_REF] Colegrove | High pressure interpass rolling of Wire + Arc additively manufactured titanium components[END_REF]. The stress is generally maximal in the longitudinal direction and minimal in the transversal one where less matter restrains the geometry to deform [START_REF] Martina | Residual stress of as-deposited and rolled wire+arc additive manufacturing Ti-6Al-4V components[END_REF][START_REF] Szost | A comparative study of additive manufacturing techniques: Residual stress and microstructural analysis of CLAD and WAAM printed Ti-6Al-4V components[END_REF][START_REF] Colegrove | High pressure interpass rolling of Wire + Arc additively manufactured titanium components[END_REF]. A global trend along the three directions is observed. The residual stress passes through a maximum compressive stress at the centre (≈ -100 MPa) of the baseplate, then a maximum tensile stress is encountered in the first layer ( ≈ +90 MPa). A second maximum tensile stress is observed in the last layers (≈ +80 MPa), remaining low in between these two zones. The low stress magnitudes found are consistent with those reported by Hönnige et al. [START_REF] Hönnige | Control of residual stress and distortion in aluminium wire + arc additive manufacture with rolling[END_REF] on a AA2319 with CMT pulsed advanced deposition process. However, the global trend described by these authors seems to be different and the stresses in the transverse and normal directions are much lower than the values obtained in the present study. No element from literature and performed analyses enables us to explain this variation. Several studies, focused on residual stresses within parts built by WAAM, relate significant stress levels and distortions strongly depending on the material, the process parameters, the geometry and the clamping conditions [START_REF] Hönnige | Control of residual stress and distortion in aluminium wire + arc additive manufacture with rolling[END_REF][START_REF] Martina | Residual stress of as-deposited and rolled wire+arc additive manufacturing Ti-6Al-4V components[END_REF][START_REF] Szost | A comparative study of additive manufacturing techniques: Residual stress and microstructural analysis of CLAD and WAAM printed Ti-6Al-4V components[END_REF][START_REF] Williams | Wire + Arc additive manufacturing[END_REF][START_REF] Hönnige | Residual Stress Characterization and Control in the Additive Manufacture of Large Scale Metal Structures[END_REF][START_REF] Colegrove | Microstructure and residual stress improvement in wire and arc additively manufactured parts through high-pressure rolling[END_REF][START_REF] Colegrove | Application of bulk deformation methods for microstructural and material property improvement and residual stress and distortion control in additively manufactured components[END_REF][START_REF] Martina | Investigation of the benefits of plasma deposition for the additive layer manufacture of Ti-6Al-4V[END_REF][START_REF] Ding | Thermomechanical analysis of Wire and Arc Additive Layer Manufacturing process on large multi-layer parts[END_REF][START_REF] Colegrove | High pressure interpass rolling of Wire + Arc additively manufactured titanium components[END_REF]. It could be noticed that our results are not in agreement with these studies. But it should be also noted that the comparison with these works is very difficult since the process parameters and materials are not equivalent. The stress magnitude reduction of approximately 30% when the substrate is removed is however consistent with the literature [START_REF] Szost | A comparative study of additive manufacturing techniques: Residual stress and microstructural analysis of CLAD and WAAM printed Ti-6Al-4V components[END_REF].

CONCLUSIONS

From this experimental study, the following observations and conclusions can be drawn:  It is possible to build parts free from significant defects with the AA6061 using the WAAM process. A very limited quantity of porosities is found thanks to a careful cleaning of the wire and substrate. Hot cracking can be avoided by the formation of equiaxed grains on the top portion of each deposited layer. This phenomenon is promoted by the presence of constitutive phases within the wire filler material and by the use of a high welding speed.  The as-deposited material is quenched due to the rapid cooling of the deposit. During the holding step at room temperature, it undergoes natural ageing which leads to a relative metallurgical stability of the microstructure after two weeks, its hardness being then equivalent with the value obtained on the commercial T4 material. After a T6 heat treatment, the hardness value of the WAAM material is higher than the one measured on the commercial T6 material. This phenomenon can be explained by a state of precipitation which is different from the one observed in the forged material, and more particularly, by the absence of large dispersoids and the coexistence of different phases (GPI, Q', β', β'', L) at the nanoscale.  Both strategies produce stacked structures consisting of successive alternating layers of columnar and equiaxed grains. The as-deposited material shows a strong preferential crystallographic orientation, but less pronounced with the alternate deposition strategy. The micrometric phases found in the filler wire (iron-IM and β-phase) may be melted during the deposition process and are then no longer present in the deposited material. This variation changes the material concentration, containing more solutes that affect the precipitation kinetics obtained during the hardening T6 heat treatment, leading to a WAAM material harder than the usual material in the T6 state.  TEM analyses have highlighted significant differences between the forged material and the WAAM material.

At T6 state, there are no large dispersoids. However, it reveals a multitude of β'-, L-, Q'-and β''phases in the material. This coexistence could explain the improvement of hardness properties.  Finally, the stress analysis shows that both manufacturing strategies lead to similar stress profiles.

Compressive stresses are found in the baseplate and residual tensile stresses along three orthogonal directions in the walls. Substrate removal reduces by 30% the maximum compressive stress.

  (b). The welding torch working distance is set at 15 mm.

Figure 1 :

 1 Figure 1 : (a) Schematic experimental set up for deposition. (b) details of base plate clamping and cooling devices.

Figure 2 :

 2 Figure 2 : Schematic illustration of the set-up for neutron measurements (a) for transverse direction (Y) ; (b) for normal direction (Z) ; (c) for longitudinal direction (X) ; (d) gauge positions.

Figure 3 :

 3 Figure 3: (a) geometry of the as deposited thin-wall for the alternate strategy ; (b) final part geometry for the alternate strategy ; (c) geometry of the as deposited thin-wall for the unidirectional strategy ; (d) final part geometry for the unidirectional strategy.

Figure 4 :

 4 Figure 4 : X-Ray radiographs of the deposited walls (a) alternate strategy (b) unidirectional strategy.

Figure 5 :

 5 Figure 5: Hardness mapping for the two strategies and different metallurgical states

Figure 6 :

 6 Figure 6: Macrographs illustration of the microstructure for the as-deposited WAAM AA6061 with magnification of enclosed regions by the white square for : (a) longitudinal section similar for both strategy, and in the X-Z plane section for (b) alternate strategy, (c) unidirectional strategy.

Figure 7 :

 7 Figure 7: Grain morphology and crystallographic texture for the (a) unidirectional strategy and (b) alternate strategy.

Figure 8 :

 8 Figure 8: Texture evolution (a) layer n, (b) layer n+1.

Figure 9 :

 9 Figure 9: Grain size histogram obtained by EBSD

Figure 10 :

 10 Figure 10: (a) SEM image with BSE contrast. (b) EDS of β-phase (red) and iron-IM (black).

Figure 11 :

 11 Figure 11: AA6061 at different metallurgical states (a) wire filler, (b) WAAM as-deposited material, (c) T6-treated WAAM material.

Figure 12 :

 12 Figure 12: Bright-field TEM images of dispersoids for (a) the forged material at T6 state taken as the reference, (b) the WAAM material as-deposited and (c) the WAAM material at T6 state.

Figure 13 :

 13 Figure 13: Bright-field TEM images of dispersoids, nanometric β'' and β' phases and associated diffraction pattern for (a) the forged T6 material taken as the reference, (b) the WAAM as-deposited material and (c) the WAAM T6 material.

Figure 14 :

 14 Figure 14: Bright-field TEM images of nanometric phases observed in the WAAM T6 material highlighting the presence of Q', L and β'' phases.

1 2

 1 𝑆 2 (ℎ𝑘𝑙) and 𝑆 1 (ℎ𝑘𝑙), necessary for the stress calculation, are:

1 2

 1 𝑆 2 (31.1) = 19.18710 -6 MPa -1 , 𝑆 1 (31.1) = -4.94510 -6 MPa -1 .Figure15shows the residual stress profiles along the height of the wall in the longitudinal X, transverse Y and normal Z directions for each strategy.

Figure 15 :

 15 Figure 15: Evolution of the residual stresses along the height of the wall for each direction (longitudinal X, transverse Y and normal Z) within the samples built with (a) the alternate strategy ; (b) the unidirectional strategy and (c) the unidirectional strategy after substrate removal.

  global trend along the three directions is observed. Residual stresses are compressive with a maximum in the centre of the baseplate (z = -4 mm ; 𝜎 𝑖𝑖 ≈ -100 MPa). The residual stress increases rapidly toward tensile values with increasing z values. The highest normal stress values are obtained in the first layers (z = 4 mm ; 𝜎 𝑥𝑥 ≈ +90 MPa). A second maximum tensile stress is observed in the last layers (z = 40 mm ; 𝜎 𝑥𝑥 ≈ +80 MPa), remaining low in between.

Table 1 :

 1 Chemical composition of the AA6061wire filler and for the AA6061-part built by WAAM

	Elements	Mg	Si	Cr	Fe	Ti	Zn	Mn	Cu	Al	Total
	AA6061 Wire Filler wt%	0.91	0.99	0.13	0.29	0.01	0.16	0.07	0.20	97.24 100.00
	AA6061 as-deposited wt%	0.87	1.05	0.09	0.29	0.01	0.10	0.05	0.13	97.41 100.00
	AA6061 Nominal Composition wt% [51]	0.81 -1.20	0.40 -0.80	0.04 -0.35	max 0.70	max 0.15	max 0.25	max 0.15	0.15 -0.40	bal. 100.00

Table 2 :

 2 Hardness properties for the different heat treatments and deposition strategies

	Metallurgical State		As deposited	Naturally aged (2 weeks /RT)		Maturated (2h/90°C)	T6 (530°C/1h + 175°C/8h)
	Strategy	Alternate Unidirectional Alternate Unidirectional Alternate Unidirectional Alternate Unidirectional
	Maximum hardness (HV0.1)	75	73	100	89	88	100	137	140
	Mean hardness (HV0.1)	64	64	80	75	78	76	124	118
	Minimum hardness (HV0.1)	48	48	49	53	62	49	112	100
	Standard deviation (HV0.1)	5	6	10	6	5	9	6	6

Table 3 :

 3 Measurements obtained by the intercept method.Results on average area, diameter and size of grains for the two sets of grain population (equiaxed and columnar), determined according to ASTM E112 standard test methods for the different deposition strategies, are summarized in Table4.

	Strategy	Zone	Mean Intercept (µm) Longitudinal	Mean Intercept (µm) Transverse	Mean Intercept (µm) Planar	Mean Intercept (µm)	Aspect Ratio (L/T)
	Unidirectional	Equiaxed Columnar	/ 464	/ 99	/ 93	95 163	1.0 4.7
	Alternate	Equiaxed Columnar	/ 319	/ 99	/ 93	86 143	1.0 3.2

Table 4 :

 4 Results calculated from measurements in Table3according to the ASTM E112 standard test's methods.

	Strategy	Zone	Average Grain Area (µm²)	Average Diameter (µm)	Grain size n° (G)
	Unidirectional	Equiaxed Columnar	11405 32258	107 180	3.5 2.0
	Alternate	Equiaxed Columnar	9735 19056	98 142	3.7 2.7

Acknowledgments

This work is based upon experiments performed at the SALSA instrument (ILL, Grenoble, France). The authors thank the ILL neutron Facilities scientific comities for the allocated experimental days (experiment 81346).

Funded by the French Contrat Plan État-Région and the European Regional Development Fund of Pays de la Loire, the CIMEN Electron Microscopy Centre in Nantes is greatly acknowledged.

Declaration of Competing Interest

The authors declare that they have no competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.


Thomas Clausse
email: thomas.clausse@univ-lorraine.fr

Léopold V B Diop

Nicolas P Martin

Romain Sibille

Thomas Mazet

Zr 3 Mn 3 Sn 4 Ga : a new hexagonal Ti 6 Sn 5 -type quaternary intermetallic

Keywords: Intermetallic, X-Ray diffraction, Neutron diffraction, Kagome

We report the discovery of a new quaternary intermetallic compound of composition Zr 3 Mn 3 Sn 4 Ga during a survey of the Zr-Mn-Sn-Ga phase diagram. Single crystal x-ray diffraction reveals that the new compound crystallizes in a hexagonal lattice (space group P 6 3 /mmc # 194). The analysis of single-crystal x-ray diffraction data indicates that the atomic arrangement is an ordered variant of the high temperature hexagonal Ti 6 Sn 5 crystal structure. From a simple geometry point of view, the unit cell consists in GaZr 3 octahedra chains and Mn chains growing along the sixfold symmetry axis c. The Mn chains are stacked within the ab basal plane in a kagome geometry, with large interatomic distances. We discuss the possible slight levels of Sn/Ga and Zr/Mn mixed occupancy based on additional results of electron microprobe analysis and single-crystal neutron diffraction. It is noteworthy that this compound is the first

quaternary intermetallic discovered in this structure, and the first Ti 6 Sn 5 derivative bearing a magnetic metal, namely Mn, in the 6g position, which may give rise to interesting physical properties.

Introduction

Exploratory works have been undertaken in the Zr-Mn-Sn-Ga quaternary system in order to synthesize ZrMn 6 (Sn, Ga) 6 single crystals by means of the self-flux method [START_REF] Canfield | [END_REF]. Indeed, this compound is one of the representatives of the HfFe 6 Ge 6 structure, which is known to hold a 3d transition metal kagome sublattice. This geometry has received increasing attention as a platform to investigate intriguing electronic transport and magnetic phenomena (Liu et al., 2020;Meier et al., 2020;Mielke III et al., 2022;Yin et al., 2020;Li et al., 2021;Ghimire et al., 2020).

ZrMn 6 (Sn, Ga) 6 compounds exhibit incommensurate magnetic structures influenced by the tin/gallium substitution (Mazet et al., 2010), which motivates their synthesis as single crystals.

To the best of our knowledge, no studies of both the Zr-Mn-Sn-Ga quaternary system and the M-Sn-Ga (M=Zr,Mn) ternary systems are reported in the literature.

Isothermal sections (T = 773 K) have been published for the Zr-Mn-Sn (Romaka et al., 2014) and the Zr-Mn-Ga (Markiv, 1981) ternary systems. The former shows ZrMn 6 Sn 6 as the only ternary compound, while the latter reveals five ternary compounds (Zr 3 Mn 2 Ga 4 , Zr 3 Mn 2 Ga 6 , ZrMn 0.5 Ga 2.5 , Zr 6 Mn 10.5 Ga 12.5 and ZrMn 3 Ga 6 ).

Unexpectedly, our experiments yielded single crystals of a new quaternary phase, derivating from the Ti 6 Sn 5 high temperature hexagonal structure, which was solved using single-crystal x-ray diffraction. Their stoichiometry is close to Zr 3 Mn 3 Sn 4 Ga, and is discussed in light of electron probe microanalysis and neutron diffraction experiments. These results are disclosed here.

Experimental details

Synthesis

The crystals were synthesized using a tin self-flux method [START_REF] Canfield | [END_REF], starting with pure metals used as received : Zr (100 mesh powder, 99.5%, Strem chemicals), Mn (100-200 mesh powder, 99.6%, Cerac speciality inorganics), Sn (3 mm shots, 99.8%, Sigma Aldrich) and Ga (pieces, 99.99%, Alfa Aesar). A 15 g batch of molar ratio Zr/Mn/Sn/Ga 1/2/6.3/0.7 was prepared and loaded in a fused silica tube, which was thoroughly rinsed and flame-sealed under a pressure of 200 mbar of pure argon. The ampoule was subsequently placed into a resistive vertical tube furnace for melting. The mixture of the constituent elements was heated up to 1100 • C and kept at this temperature for 36 h. Then the ampoule was slowly cooled down to 600 • C at a rate of 5 • C/h, and finally, centrifuged to remove the excess of flux.

The quaternary nature of the crystals was assessed using energy dispersive x-ray spectroscopy on a Jeol JSM-6010LA scanning electron microscope. No traces of oxidation were detected on the crystals.

The main product is single crystals of the title compound, in the form of dark gray metallic luster prismatic crystals that were mechanically extracted from the bulk.

The flux remains sticking to the crystals were mechanically removed. A typical raw specimen is shown in figure 1.

Electron probe microanalysis

Quantitative elemental analysis was carried out using wavelength dispersive xray spectroscopy on a Jeol JXA-8530F Castaing Microprobe. Several crystals were removed from the bulk, coated in epoxy resin and polished. Quantitative analysis was performed on multiple locations of various single crystals, using pure metals as standards for zirconium, manganese and tin, while GaP was used for gallium.

X-ray diffraction

Single crystal x-ray diffraction data were collected at room temperature and at 100 K using a Bruker Kappa APEX-II diffractometer equipped with a molybdenum x-ray source (Mo K α = 0.71073 Å). Acquisition and further treatment of raw data were performed using the APEX2 program package [START_REF] Bruker | Bruker advanced X-ray solutions[END_REF]. The structure was solved and refined using SHELXS (Sheldrick, 2008) and SHELXL 2018/3 (Sheldrick, 2015), respectively. Unit cell drawings were produced with Vesta (Momma & Izumi, 2011).

Details on the data collection and refinement are available in the supporting information.

Neutron diffraction

Single-crystal neutron diffraction data were collected at 10 K using the Zebra instrument at the Swiss Spallation Neutron Source SINQ in Villigen PSI, Switzerland. The sample was cooled down using a helium closed-cycle refrigerator mounted inside a Eulerian cradle. The beam of thermal neutrons was monochromatized using the (220) reflection of germanium crystals, yielding a neutron wavelength of 1.383 Å with less than 1% of λ/2 contamination. A small dataset of about 300 reflections (50 independent), covering reciprocal space up to sin(θ)/λ ≃ 0.4 Å-1 , was measured using a single 3 He-tube detector in front of which slits were appropriately adjusted. Reflections were collected using ω-scans (step size 0.1 • ) and then numerically integrated above the variable linear flat background of fits to a Gaussian peak shape. The Lorentz correction was applied and the data reduction using the Fullprof Suite software (Rodríguez-Carvajal, 2001) indicated a satisfactory R internal for equivalent reflections of 0.0749.

Results and discussion

The crystal structure has been solved using only x-ray diffraction data, and is described in section 3.2. Neutron diffraction data have been refined on the basis of this strctural model. The corresponding results are consistent with the x-ray data, and are available in the supplementary information.

Phase stoichiometry

The averaged stoichiometry of the crystals deduced of the electron probe microanalysis is Zr 2.84(4) Mn 3.12(3) Sn 3.81(4) Ga 1.23( 5) . These results are close to Zr 3 Mn 3 Sn 4 Ga, with possible Zr/Mn and Sn/Ga substitutions.

Given the big difference of atomic numbers of Zr and Mn on one hand, and Sn and Ga on the other, we hoped the chemical substitution would be confirmed by the refinement of the x-ray diffraction data using the structural model described in section 3.2. Unfortunately, attempts to account for this eventual disorder led to unsatisfactory R factors, or occupancies of shared sites within the error bars. These data were therefore refined with constrained full occupancy of all crystallographic sites, and bring no additional information about the phase stoichiometry.

Manganese and zirconium also have very distinct neutron scattering lengths, respectively -3.75 fm and 7.16 fm, which led us to further test an eventual substitution of Zr for Mn in the corresponding data refinement. The refinement of the single-crystal neutron diffraction data using the fully ordered model deduced from the x-ray diffraction data refinement is shown in figure 2. A total of 7 intensity-dependent parameters was used : the 2 free atomic coordinates for Zr and Sn on their respective 6h Wyckoff positions, and 1 isotropic displacement factors for each atom. This model leads to an excellent agreement with a R F of 0.0642. Adding Zr/Mn disorder leads to only a slight improvement, with R F = 0.0589, corresponding to a substitution of Zr for Mn of 0.065 per Zr atom and estimated standard deviation of 0.012. Sn and Ga sites being kept constrained to full occupancy, the stoichiometry of the crystal deduced from the refinement of the neutron data is Zr 2.80(4) Mn 3.20(4) Sn 4 Ga, which agrees with the Zr and Mn contents derived from the electron microprobe analysis.

A Zr 3 Mn 3 Sn 4 Ga formula is therefore suggested for the quaternary phase, which will be used in the following discussion for simplicity. Electron probe microanalysis leans toward some levels of substitution of Ga for Sn and Mn for Zr, the latter being supported by neutron diffraction data refinement. This substitution suggests the existence of an homogeneity domain, in accordance with the well-known chemical affinity of Zr and Mn, as well as Sn and Ga. However, the inability to achieve a compatible result with the x-ray diffraction data, as well as the limited amount of data used in the neutron refinement and the only slight improvement of the corresponding fit quality make it difficult to fully argue against the ordered model. Further synthesis and experiments will therefore be required to explore this matter.

Structure description

Zr 3 Mn 3 Sn 4 Ga crystallizes in a quaternary ordered variant of the high tempreature hexagonal Ti 6 Sn 5 -type crystal structure (space group P 6 3 /mmc # 194) (Schubert et al., 1963). The room temperature cell parameters, deduced from the x-ray diffraction experiment, are a = 9.2073(4) Å and c = 5.6287(2) Å. The fractional atomic coordinates and atomic displacement parameters are given in the supporting information. No signature of a temperature-induced structural phase transition was perceived down to the lowest temperature of the x-ray diffraction experiment (100 K). The crystal symmetry is preserved down to 100 K, i.e., the hexagonal symmetry of the atomic arrangement remains unaltered. It is worth noting that an orthorhombic allotrope has been reported for Ti 6 Sn 5 at temperatures below 943 K (Okamoto, 2010;Wang et al., 2014).

The crystal structure of Zr 3 Mn 3 Sn 4 Ga derives from the hexagonal Ti 6 Sn 5 unit cell by splitting the two titanium positions between Zr (Wyckoff site 6h) and Mn (Wyckoff site 6g), while Ga occupies the Wyckoff 2a tin position, and tin remains on their original 6h and 2c crystallographic sites.

The unit cell is depicted down the [001] direction and in perspective in figure 3 and4 respectively. The coordination polyhedra are listed in table 1, and depicted in figure 5.

Most of them are distorted due to the variability of elements on their vertexes. The Sn atoms (2c Wyckoff position) have an axial symmetry ( 6m2) and a local environment which is made of 15 near neighbours, 3 of which are Zr ones, 6 Sn(6h) type and 6 are Mn ones. The Zr(6h) position is also surrounded by 15 atoms, here one finds only one Sn(2c) neighbour, 4 Sn located at 6h, 2 Ga neighbours, 4 Mn atoms and 4 Zr neighbours. The local environment of Mn(6g) atoms consists of 4 Sn(6h) neighbours, 2 Sn(2c) neighbours, 4 Zr neighbours and 2 Mn neighbours. The Ga(2a) atoms have an axial symmetry ( 3m) and 8 near neighbours: 6 Zr and 2 Ga atoms. The Sn(6h) position is surrounded by 12 near neighbours: 2 Sn located at 2c, 2 Sn(6h) type, 4 Zr neighbours and 4 Mn atoms.

A simplified description of the structure can be made from the Ga and Mn polyhedra. Indeed, the Mn centered icosahedra are packed in a kagome geometry in the ab basal plane with shared faces. This packing generates gaps around the 2a Wyckoff position, which host the Ga centered dodecahedra. This layout is depicted in figure 6. This polyhedra association generates slabs that are stacked along the sixfold symmetry axis c, with shared Ga (respectively Mn) vertexes between the Ga (respectively Mn) polyhedra. This stacking is depicted in figure 7.

Purely geometric considerations lead to the following description. From that point of view, gallium atoms (Wyckoff site 2a) lie at the center of zirconium octahedra (Wyckoff site 6h). These octahedra are stacked with shared faces along the highsymmetry direction c, building GaZr 3 octahedra chains. Three of these chains are bridged by a tin atom (Wyckoff site 2d). Manganese (Wyckoff site 6g) chains grow along the c-axis between the octahedra chains. They generate two families of trigonal prisms. Every other prism is centered by a tin atom (Wyckoff site 2c), while the empty trigonal prisms are capped by tin atoms (Wyckoff site 6h) on their rectangular faces.

Figures 8 and9 illustrate the unit cell on that description basis.

Structure discussion

The Ti 6 Sn 5 hexagonal structure consists of five sites, which will be referred as titanium and tin sites in the upcoming discussion, eventhough they may be occupied by other elements.

Only a few intermetallics share isotypism with Ti 6 Sn 5 . Ordered binary compounds are V 6 Ga 5 (Schubert et al., 1963), Ta 6 Ga 5 (Popova & Putro, 1978), Ti 6 Pb 4.8 (Kleinke, 2001). Some ternary variants have been reported. They correspond to solid solutions derivating from the previously mentionned binaries, with mixed occupancies on the titanium sites.

Ordered ternary compounds have also been identified, with a RE 3 RuMg 7 formula, where RE = Sm, Gd, Tb (Linsinger & Pöttgen, 2011). Mixed occupancy has been reported in this structure, with the 2c tin site shared between samarium and magnesium. The RE 3 RuMg 7 compounds are also special since one element -the magnesium -occupies both a titanium and a tin site. Zr 3 Mn 3 Sn 4 Ga is the first quaternary ordered isostructural compound. It is remarkable that this structure occurs in that quaternary system, while none of the associated binaries or ternaries include a Ti 6 Sn 5 isotype. In this compound, zirconium and manganese occupy the titanium positions, while tin and gallium atoms reside on the tin positions. It is noteworthy that this is the first example of a splitting of the tin site between two elements in this structure.

The magnetic properties of the compound might be of significant interest thanks to the manganese in the structure. Indeed, Zr 3 Mn 3 Sn 4 Ga is the only Ti 6 Sn 5 compound bearing a 3d magnetic metal on the 6g position, i.e., the site forming a kagome lattice.

The room temperature intra and interplanar Mn-Mn distances in Zr 3 Mn 3 Sn 4 Ga are summarized in table 2. For comparison purposes, the interatomic distances for ZrMn 6 Sn 6 , which is a kagome compound with the same elements (Mazet et al., 1999) are also given in table 2. In Zr 3 Mn 3 Sn 4 Ga, the in-plane Mn-Mn distances are far greater than along the [001] direction. The Mn sublattice should therefore be described as Mn chains stacked in a kagome geometry, rather than kagome planes stacked in the [001] direction as in ZrMn 6 Sn 6 . This contrast might however not be an obstacle to kagome driven physics. Indeed, hints of an interesting correlated electron behavior in Ti 6 Sn 5 were reported, based on first-principle calculations (Jeong, 2007) and experimental results (Drymiotis et al., 2003).

The low dimensionnality of the Mn sublattice, the possible strong electronic correlations, or the kagome driven physics might make of Zr 3 Mn 3 Sn 4 Ga an interesting subject for solid state physics, and demand further examination. Magnetic and electronic properties are currently undergoing studies by means of bulk magnetization and specific heat measurements, neutron scattering, and 119 Sn Mössbauer spectroscopy.

Conclusion

Attempts to synthesize ZrMn 6 Sn 6 -x Ga x single crystals in a tin self-flux unexpectedly led to a new quaternary phase with a stoichiometry close to Zr 3 Mn 3 Sn 4 Ga.

X-ray diffraction experiments show that it is a quaternary variant of high tempera-

IUCr macros version 2.1.15: 2021/03/05 ture hexagonal Ti 6 Sn 5 . A simple geometric description of the cell consists in GaZr 3 octahedra chains and manganese chains growing along the c-axis. The latter chains are stacked in the ab-planes in a kagome geometry. A homogeneity range could exist for the quaternary phase, as suggested by electronic microprobe measurements as well as by structural refinement of the neutron diffraction data.

The low dimensionnality of the manganese sublattice, as well as the existence of a kagome geometry, motivate the study of the magnetic properties of this compound.

Synopsis

The crystal structure of Zr 3 Mn 3 Sn 4 Ga, a new quaternary intermetallic crystallizing in the high temperature hexagonal Ti 6 Sn 5 type structure, is reported.

Fig. 1 .Fig. 2 .

 12 Fig. 1. Photograph of a Zr 3 Mn 3 Sn 4 Ga single crystal, as obtained after centrifugation of the synthesis flux.

Fig. 3 .

 3 Fig. 3. Zr 3 Mn 3 Sn 4 Ga unit cell viewed down the [001] direction. Zirconium, manganese and gallium are drawn as blue, pink and orange spheres, respectively. Tin atoms are drawn in black on the 6h position, and in gray in the 2c.

Fig. 4 .

 4 Fig. 4. Perspective view of the Zr 3 Mn 3 Sn 4 Ga unit cell. Zirconium, manganese and gallium are drawn as blue, pink and orange spheres, respectively. Tin atoms are drawn in black on the 2c position, and in gray in the 6h.

Fig. 6 .

 6 Fig. 6. Zr 3 Mn 3 Sn 4 Ga unit cell viewed down the [001] direction, with apparent Ga and Mn coordination polyhedra. Zirconium, manganese and gallium are drawn as blue, pink and orange spheres respectively. Tin atoms are drawn in black on the 2c position, and in gray in the 6h. Some atoms outside the cell are drawn to complete the polyhedra.

Fig. 7 .

 7 Fig. 7. Zr 3 Mn 3 Sn 4 Ga unit cell perspective view with apparent Ga and Mn coordination polyhedra. Zirconium, manganese and gallium are drawn as blue, pink and orange spheres respectively. Tin atoms are drawn in black on the 2c position, and in gray in the 6h. Some atoms outside the cell are drawn to complete the polyhedra.

Fig. 8 .

 8 Fig. 8. Zr 3 Mn 3 Sn 4 Ga unit cell viewed down the [001] direction showing the GaZ 3 octahedra, and layout of the manganese chains in a kagome geometry. The manganese trigonal prisms, capping tin, and bridging tin atoms are emphasized. Some atoms outside the cell are drawn to complete the octahedra.

Fig. 9 .

 9 Fig. 9. Perspective view of the Zr 3 Mn 3 Sn 4 Ga unit cell showing the GaZ 3 octahedra chains, and layout of the manganese chains in a kagome geometry. The manganese trigonal prisms, capping tin, and bridging tin atoms are emphasized. Some atoms outside the cell are drawn to complete the octahedra.

Table 1 .

 1 Zr 3 Mn 3 Sn 4 Ga coordination polyhedra and nearest neighbours of each atoms in the unit cell. Atoms are considered nearest neighbours when they are closer than the sum of

		atomic radii plus 0.5 Å. Z15 : 15-vertex Frank-Kasper polyhedron.
	Central		Nearest neighbours		Polyhedron
	atom	Zr (6h) Mn (6g) Sn (6h) Sn (2c) Ga (2a)	(fig. 5)
	Zr (6h)	4	4	4	1	2	Z15
	Mn (6g)	4	2	4	2	-	Icosahedron
	Sn (2c)	3	6	6	-	-	Z15
	Sn (6h)	4	4	2	2	-	Icosahedron
	Ga (2a)	6	-	-	-	2	Dodecahedron

Table 2 .

 2 Mn -Mn room temperature interatomic lengths comparison between Zr 3 Mn 3 Sn 4 Ga and ZrMn 6 Sn 6 . d / Å Zr 3 Mn 3 Sn 4 Ga ZrMn 6 Sn 6(Mazet et al., 1999)

	c	2.8144(1)	4.4890(5)
	⊥ c	4.6037(2)	2.7310(5)

IUCr macros version 2.1.15: 2021/03/05

The authors are indebted to Christine Gendarme of the CC3M (Centre de Compétences Microscopies, Microsonde et Métallographie, Institut Jean Lamour) for the Castaing Microprobe measurements and analysis of the single crystals, and to Thierry Schweitzer (IJL, team 103), who prepared the microprobe samples. This work is based on experiments performed at the Swiss spallation neutron source SINQ, Paul Scherrer Institute, Villigen, Switzerland.