Edouard Chauvet

Paraskevas Kontis

Eric A Jägle

Baptiste Gault

Dierk Raabe

Catherine Tassin

Jean-Jacques Blandin

Rémy Dendievel

Benjamin Vayre

Stéphane Abed

Guilhem Martin
email: guilhem.martin@simap.grenoble-inp.fr

Hot cracking mechanism affecting a non-weldable Ni-based superalloy produced by selective electron Beam Melting

Keywords: Electron beam Melting, Ni-based superalloys, Hot cracking, Grain boundaries, Liquid film, micro-segregations

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Introduction

Precipitation strengthened Ni-based superalloys have been a one of the great breakthroughs in materials science enabling the development of advanced and fuel efficient aeroengines. Since the 1950s, this class of superalloys is in continuous development to achieve better creep and oxidation resistance, lower fatigue crack growth rate and high yield stress for service temperatures up to 1000°C. These superalloys are typically strengthened by ordered, coherent intermetallic precipitates such as γ'-Ni3(Al,Ti) or γ''-Ni3Nb. Different processes can be used to produce parts made of this class of Ni-based superalloys: casting [START_REF] Zhang | Effect of grain-boundary characteristics on castability of nickel-base superalloys[END_REF], directional solidification [START_REF] Wang | Solidification cracking of superalloy single-and bicrystals[END_REF][START_REF] Zhang | Hot tearing of nickel-based superalloys during directional solidification[END_REF][START_REF] Grodzki | Effect of B, Zr, and C on Hot Tearing of a Directionally Solidified Nickel-Based Superalloy[END_REF][START_REF] Zhou | Effect of grain boundary characteristics on hot tearing in directional solidification of superalloys[END_REF][START_REF] Zhou | Effect of grain boundary fraction on castability of a directionally solidified nickel alloy[END_REF][START_REF] Zhou | Effect of dendrite arm spacing on castability of a directionally solidified nickel alloy[END_REF], and powder metallurgy [START_REF] Lemarchand | Investigation of grain-boundary structuresegregation relationship in an N18 nickel-based superalloy[END_REF][START_REF] He | Microstructure evolutions and nucleation mechanisms of dynamic recrystallization of a powder metallurgy Ni-based superalloy during hot compression[END_REF][START_REF] Jiang | Microstructurally sensitive crack nucleation around inclusions in powder metallurgy nickel-based superalloys[END_REF]. With the recent development of new processing and advanced manufacturing routes, in particular additive manufacturing, there is an increasing demand for producing high-temperature components made of such Nibase superalloys. Key drivers for this include the ability to produce complex netshape components without the restrictions of traditional machining, thereby enabling optimized cooling channels that allow for higher service temperatures as well as the ability to rapidly produce small batches of complex parts without the prohibitive costs and lead times of traditional casting techniques.

Until recently, the community strongly focused on the well-known alloy Inconel 718, see e.g. [START_REF] Amato | Microstructures and mechanical behavior of Inconel 718 fabricated by selective laser melting[END_REF][START_REF] Helmer | Additive manufacturing of nickel-based superalloy Inconel 718 by selective electron beam melting: Processing window and microstructure[END_REF][START_REF] Jia | Selective laser melting additive manufacturing of Inconel 718 superalloy parts: Densification, microstructure and properties[END_REF][START_REF] Strondl | Microstructure and mechanical properties of nickel based superalloy IN718 produced by rapid prototyping with electron beam melting (EBM)[END_REF][START_REF] Strossner | Mechanical and Microstructural Investigation of Nickel-Based Superalloy IN718 Manufactured by Selective Laser Melting (SLM)[END_REF][START_REF] Wang | The microstructure and mechanical properties of deposited-IN718 by selective laser melting[END_REF], as it can be quite efficiently produced by additive manufacturing without producing critical defects. Wide processing windows have been identified for the two powder bed-based techniques selective laser melting process (SLM) [START_REF] Jia | Selective laser melting additive manufacturing of Inconel 718 superalloy parts: Densification, microstructure and properties[END_REF][START_REF] Wang | The microstructure and mechanical properties of deposited-IN718 by selective laser melting[END_REF] and S-EBM [START_REF] Helmer | Additive manufacturing of nickel-based superalloy Inconel 718 by selective electron beam melting: Processing window and microstructure[END_REF][START_REF] Dehoff | Crystallographic texture engineering through novel melt strategies via electron beam melting: Inconel 718[END_REF]. This is not surprising as Inconel 718 is usually considered to be a weldable superalloy due to its low content of γ' forming elements (Al and Ti) [START_REF] Lippold | Welding Metallurgy and Weldability of Nickel-Base Alloys[END_REF], instead relying on γ''-Ni3Nb as strengthening precipitate phase. However, to achieve higher creep resistance under high load, superalloys with between 40 and 80% volume fraction of γ' are required. Such superalloys include Inconel 738 [START_REF] Cloots | Investigations on the microstructure and crack formation of IN738LC samples processed by selective laser melting using Gaussian and doughnut profiles[END_REF][START_REF] Engeli | Processability of different IN738LC powder batches by selective laser melting[END_REF], CMSX-4 [START_REF] Ramsperger | Solution Heat Treatment of the Single Crystal Nickel-Base Superalloy CMSX-4 Fabricated by Selective Electron Beam Melting[END_REF][START_REF] Ramsperger | Microstructure of the Nickel-Base Superalloy CMSX-4 Fabricated by Selective Electron Beam Melting[END_REF], René 142 [START_REF] Murr | Microstructures of Rene 142 nickel-based superalloy fabricated by electron beam melting[END_REF] or CM247LC [START_REF] Carter | The influence of the laser scan strategy on grain structure and cracking behaviour in SLM powder-bed fabricated nickel superalloy[END_REF], which have hence recently attracted attention for usage in additive manufacturing processes. Those alloys contain comparatively high contents of Al and Ti, are classified as non-weldable, and, as a result, are prone to cracking and hence turn out to be very challenging to produce defect free by additive manufacturing [START_REF] Engeli | Processability of different IN738LC powder batches by selective laser melting[END_REF][START_REF] Carter | The influence of the laser scan strategy on grain structure and cracking behaviour in SLM powder-bed fabricated nickel superalloy[END_REF][START_REF] Ruttert | Impact of hot isostatic pressing on microstructures of CMSX-4 Ni-base superalloy fabricated by selective electron beam melting[END_REF][START_REF] Attallah | Additive manufacturing of Ni-based superalloys: The outstanding issues[END_REF][START_REF] Carter | Process optimisation of selective laser melting using energy density model for nickel based superalloys[END_REF].

Different mechanisms have been suggested in the welding literature to explain the observed cracking tendency [START_REF] Lippold | Welding Metallurgy and Weldability of Nickel-Base Alloys[END_REF]: solidification cracking [START_REF] Wang | Solidification cracking of superalloy single-and bicrystals[END_REF] and liquation cracking [START_REF] Chen | Dendritic microstructure and hot cracking of laser additive manufactured Inconel 718 under improved base cooling[END_REF] that require the presence of liquid, or strain-age cracking [START_REF] Lippold | Welding Metallurgy and Weldability of Nickel-Base Alloys[END_REF] and ductility-dip cracking [START_REF] Lippold | Welding Metallurgy and Weldability of Nickel-Base Alloys[END_REF] that occur at the solid state.

The objective of the present work is to elucidate the mechanism and the origin of cracking in a non-weldable Ni-based superalloy fabricated by S-EBM. Using detailed microstructural characterization, combining techniques such as electron microscopy, electron backscattered diffraction and atom probe tomography, we identify a number of common features that provide insights into the origin of the cracks and discuss those in light of the models generally accepted for hot cracking in conventional processing routes.

Experimental procedures

Powder characteristics

The prealloyed powders were produced using vacuum induction melting (VIM) gasatomization with Ar and provided by ERASTEEL (Spain). The material under investigation is a non weldable Ni-Co-Cr-Mo-Al-Ti-B nickel-based superalloy containing significant amounts of Cr, Co and Mo and with Ti+Al wt.% = 8.6. The exact chemical composition of the atomized powder cannot be indicated for industrial confidentiality reasons. The asreceived powder morphology and microstructure were characterized using a Zeiss Ultra field emission gun scanning electron microscope (FEG-SEM). The as-received powders mostly exhibit a spherical morphology with a few irregular particles and a relatively high density of satellites, see Figure 1a. Every powder particle is polycrystalline as revealed by the inverse pole figure electron backscattered diffraction map (IPF-EBSD) shown in Figure 1b. Some of the particles contain spherical pores corresponding to entrapped gas during the atomization process (Figure 1c). The powder size distribution was determined by laser granulometry (MALVERN Mastersize 3000), giving an average powder particle size of approximately ~75 μm. The flowability and relative powder density were measured based respectively on the Metal Powders Industry Federation standards 3 and 4. The powder bed relative density was measured to be 53.6 % and a flow time of 16.0 ± 0.1 seconds was found for 50 g of powders going through a 2.54 mm diameter orifice (Hall flowmeter).

Sample manufacturing

The as-received prealloyed powders were loaded into an ARCAM A1 Electron Beam Melting (S-EBM) machine operating at 60 kV accelerating voltage under a controlled pressure of He set to 2.10 -3 mbar. The powder was deposited by layers of 50 μm on a stainless steel plate (XY-plane), then slightly consolidated with a widely defocused beam and finally selectively melted using the automatic mode according to the input CAD geometry. The powder bed was slightly consolidated during the preheating stage in order to improve the mechanical behavior and electrical conductivity. This preheating stage turns out to be the key parameter to achieve a stable process: a smoke-free 1 processing windows as well as a degree of consolidation that further enables easy removal of the powder at the end of the fabrication.

A suitable temperature for preheating is therefore a temperature that roughly corresponds to the first stage of sintering, i.e. formation of necks between particles without densification.

Here, the preheating temperature was selected so as to have an operating temperature of approximately 1050°C. Specimens with dimensions 23 × 23 × 30 mm (see Figure 2a) were produced following a standard melting strategy, in which the beam scans the area defined by the CAD file in a snake-like way (meander) with a line offset of 0.1 mm. The temperature profile measured below the start plate is displayed in Figure 2b where the different processing stages are also indicated.

In the S-EBM, the processing parameters vary spatially according to the geometry of the part (here cubes) when using the automatic mode. This is a typical characteristic of the S-EBM process in comparison to other AM technologies. Here, every layer was melted with the exact same melting parameters: there was no change between layer n th and layer n+1 th except for a 90° rotation of the scanning direction. Please note that we only focused on the center of the cube samples, i.e. in the regions unaffected by the contours. 1 Smoke is defined as the phenomenon occurring when electrostatic repulsive forces between negatively charged powder particles becomes significant leading to powder particle splashing.

Microstructural Characterization

The S-EBM-samples were polished down to 0.04 μm colloidal silica finish and were observed by optical microscopy along the vertical cross sections which corresponds to the XZ-plane, i.e. along the build direction (BD = Z). Electron Backscattered Diffraction (EBSD), Energy Dispersive X-ray Spectroscopy (EDX) and SEM imaging were performed in a ZEISS Ultra FEG-SEM and data were collected using the OIM-TSL software. EBSD maps were collected using an acceleration voltage of 20 kV on polished samples so as to analyze the microstructure and microtexture of the as-built specimens. 1 mm²-EBSD maps were acquired in the region affected by cracks (upper part of the samples). Ten EBSD maps of 1 mm² were analyzed so as to ensure the statistical reliability of the results. EDX analyses were performed using an acceleration voltage of 15 kV and an aperture of 90 μm on flat polished samples in order to investigate qualitatively element segregations in the as-built samples. In addition, the microstructure was revealed using a fresh glyceregia etchant (15 mL HCl + 5 mL HNO3 + 10 mL Glycerol). The γ/γ' microstructure was observed using an Ultra Zeiss FEG-SEM in secondary electron mode (SE) for etched samples and backscattered electron (BSE) contrast for polished samples. The size of the precipitates was estimated by image analysis based on SE-SEM micrographs acquired on etched specimens. At each position, at least 5 images were used to have an estimation of the precipitate size.

To investigate small scale segregations, site specific lift-outs were prepared for atom probe tomography (APT) using an FEI Dual Beam FIB Helios 600. Specimens were extracted from regions at the tip of cracks, following the procedures described in [START_REF] Kontis | On the effect of boron on grain boundary character in a new polycrystalline superalloy[END_REF]. Samples were analyzed using a Cameca LEAP3000X HR instrument operating in laser mode with pulse rate at 200 kHz, pulse energy 0.4 nJ and temperature 50 K. Data reconstruction and analyses were performed in the commercial package IVAS 3.6.14. and compositions were extracted from composition profiles in the form of proximity histograms, see [START_REF] Hellman | Analysis of threedimensional atom-probe data by the proximity histogram[END_REF].

Mechanical Characterization

Microhardness measurements were performed using a Wilson Tukon 1102 Vickers Tester. A 1 kg load was used and 20 measurements were carried out at each position along the Z-direction (=BD). The error was estimated by computing the standard deviation over the 20 measurements made at each position.

To identify the reasons for cracks formation, tensile micro-specimens (gauge length = 6 mm, width = 2 mm, thickness = 1 mm) were machined from the as-built S-EBM cubic samples at different positions along the build direction with the tensile axis aligned along the X direction as shown in Figure 2c. Tensile tests were performed at room temperature using a 3300 Instron universal testing machine with an initial displacement rate of 5mm/min. Note that the objective of the tensile tests was less to provide quantitative data regarding the mechanical properties of the fabricated materials but rather to break the samples under wellcontrolled and reproducible conditions and subsequently observe the fracture surface under secondary electron contrast. In total, ten micro specimens were deformed under tension up to fracture.

Experimental Results

Microstructural gradients

Careful observations of the XZ cross section along the entire sample height reveal typical features of parts manufactured by S-EBM. In particular, columnar grains that grow epitaxially along the build direction across several layers were found, which also corresponds to the main direction of the thermal gradient, as shown in Figure 3a. EBSD maps acquired at different positions along the Z-direction are displayed in Figure 3b and confirm these observations. Also, the EBSD maps reveal that the samples are strongly textured and exhibit a typical fiber texture of directionally solidified Ni-based superalloy, i.e. a preferential growth of the <001> direction along Z and random orientation in the XY plane [START_REF] Dehoff | Crystallographic texture engineering through novel melt strategies via electron beam melting: Inconel 718[END_REF][START_REF] Dehoff | Site specific control of crystallographic grain orientation through electron beam additive manufacturing[END_REF][START_REF] Helmer | Grain structure evolution in Inconel 718 during selective electron beam melting[END_REF]. Significant microstructural heterogeneities were observed at different scales. At the grain scale, a gradient of columnar grain width along the build direction is readily visible in Figure 3b. The width of the columnar grain increases slightly from ~ 30 μm at Z = 1 mm to ~ 150 μm at Z = 10 mm, as illustrated in Figure 4a. As the height of the sample further increases, the columnar grain width increases substantially to achieve very large columnar grains of approximately 1 mm width at Z = 25 mm. In contrast, the grain boundary density decreases from bottom to top as readily visible in Figure 4a. Within the grains themselves, a significant variation in the size of the γ' precipitates was observed as illustrated in Figure 3c and Figure 4b. Note that in Figure 3 the EBSD maps and BSE-SEM images are not necessarily shown at the same positions because the grain and precipitate size evolution do not occur at the same position along Z. The general observation was that the size of the γ' precipitates was increasing from the top to the bottom of the specimen. In particular, close to the top surface (Z = 29 mm) the size of the γ' precipitates is 100 ± 20 nm. The size of the γ' precipitates dramatically increased to approximately 500 nm ± 100 nm at Z = 25 mm, to almost plateau at approximately 600 nm ± 100 nm at Z = 5 mm. This gradient of precipitate size is also consistent with the hardness profile performed along the build direction, see Figure 4b. As expected, the maximum hardness (490 HV) was found at Z = 29 mm (i.e. 1 mm below the top surface) where the precipitates exhibit the smallest observed size whereas a rather constant hardness around 400 HV was measured from Z = 0 to 27 mm. The hardness map displayed in Figure 4c is also consistent with the previous reported results and show that the hardness is relatively uniform at each Z-position. Low hardness values are sometimes observed and has to be linked to the presence of defects such as cracks or pores.

Defects: pores

XZ as well as XY cross sections showed no "lack of fusion-type defects", which correspond to unmelted zones leading to incomplete bonding between successive layers [START_REF] Tammas-Williams | XCT analysis of the influence of melt strategies on defect population in Ti-6Al-4V components manufactured by Selective Electron Beam Melting[END_REF].

The samples were almost fully dense (>99.5% measured by image analysis). However, other types of defects were identified such as pores and cracks, as illustrated in Figure 5.

Regarding the pores, two different types were identified based on their morphology and location in the sample. A first type corresponds to spherical pores randomly distributed with a diameter in the range of 20-50 μm (Figure 5a). Their morphology of near-perfect spheres agrees with the internal pores observed in the as-received powders (see Figure 1c). This suggests that those pores were partly inherited from the powder atomization process of the initial powder particles, and they will be referred to as gas porosity. This is a relatively classical observation, often reported for parts made by powder based additive manufacturing technologies, see e.g. [START_REF] Tammas-Williams | XCT analysis of the influence of melt strategies on defect population in Ti-6Al-4V components manufactured by Selective Electron Beam Melting[END_REF]. The porosity within the as-received powder batch as well as in the as-built samples were quantified. The initial powder batch contained 0.5% porosity while the as-built samples exhibited only 0.4% porosity. This shows that some of the initial spherical pores have disappeared during melting (Δ = -0.1% porosity). This can be accounted by the fact that in S-EBM, the melting occurs up to 2-3 layers down and even more depending on the processing conditions, therefore there is extensive time for some gas porosity to escape to the ambient low vacuum environment. However, one has to keep in mind that other type of porosity like the ones described below was observed, see Figure 5b. This shows that the gas porosity has likely been reduced even more significantly.

The second type consists of well-aligned chains of rounded-shape pores along the columnar grain growth direction with a typical diameter < 5 μm (Figure 5b). We assume that their origin is a consequence of the intragranular dendrites secondary arms coalescence during the last stage of solidification [START_REF] Zhang | Hot tearing of nickel-based superalloys during directional solidification[END_REF][START_REF] Zhou | Effect of grain boundary fraction on castability of a directionally solidified nickel alloy[END_REF][START_REF] Rappaz | Last-stage solidification of alloys: Theoretical model of dendrite-arm and grain coalescence[END_REF]. The distance between two chains of pores can be defined as the interdendritic spacing (distance between dendrite trunks). Those pores hence form as bridging occurs between dendritic secondary arms: the liquid can no more reach the isolated channels to compensate for the volume contraction associated with the liquid-to-solid phase transformation (i.e. solidification). Those pores are referred to as microshrinkage.

Defects: cracks and crack path propagation

The most critical defects were undoubtedly cracks running along the build direction over several millimeters, as shown in Figure 5c. Interestingly, cracks were only observed in the range Z=12-30mm when the columnar grain width was larger than approximately 150 μm. In the range Z=0-12mm, no cracks were detected. The EBSD analysis demonstrates that crack propagation always occurs along grain boundaries, as clearly shown in Figure 6a. The cracks can therefore be defined as intergranular. ~100 boundaries over the region of interest were analyzed and classified based on their misorientation angle and their cracking susceptibility, as reported in Figure 6b, which suggests that high angle grain boundaries (HAGB, misorientation > 15°) are sensitive to cracking whereas the low angle grain boundaries (LAGB, misorientation < 15°) remain uncracked. This result is consistent with the literature dealing with welding and casting of non-weldable nickel-base superalloys [START_REF] Zhang | Effect of grain-boundary characteristics on castability of nickel-base superalloys[END_REF][START_REF] Wang | Solidification cracking of superalloy single-and bicrystals[END_REF][START_REF] Rong | The influence of grain boundary angle on the hot cracking of single crystal superalloy DD6[END_REF].

Presence of a liquid film

The fracture surface of the S-EBM samples machined in the upper part (20 < Z < 30 mm) shows a dendritic morphology with a limited development of secondary arms, giving evidence of the presence of a liquid film [START_REF] Zhang | Hot tearing of nickel-based superalloys during directional solidification[END_REF], visible in Figure 7a. In other words, cracking most likely must have occurred when a liquid film was wetting the dendrites during the last stage of solidification. Samples machined at lower Z positions (5 < Z < 15 mm), where no cracks were detected, exhibit a fracture surface with some regions showing dimples typical of a ductile failure and other smaller regions indicating the presence of liquid, see Figure 7b.

The different features highlighted on the fracture surface could be associated to two different types of defects described previously. Indeed, the fracture surface with a full dendritic morphology was associated to cracks that propagate along HAGB as shown in Figure 6. The fracture surface showing a ductile failure, with only few traces of dendritic morphology, had to be linked to regions where the coalescence of secondary arms was significant. These observations finally led us to exclude solid-state cracking mechanisms, such as strain-age cracking or ductility dip cracking, and point to a mechanism requiring the presence of liquid films. Therefore, the origin of the presence of a liquid film has to be clarified.

Local enrichment at grain boundaries

Characterizing the vicinity of grain boundaries is of major interest as compositional local enrichment to grain boundaries can be considered as a signature of the original Scheil type enrichment in the liquid during the very last stages of solidification.

Regions of interest near the tip of cracked HAGB in the upper part of the build (Z = 20mm) are marked in Figure 8a-b. Intergranular semi-continuous particles were observed in terms of a bright contrast in the backscattered image in Figure 8b. Similar particles can also be seen decorating some of the grain boundaries in the lower part of the build, at Z = 5mm (Figure 8c). EDX mapping of the main alloying elements was performed at one of the cracks tips, significant partitioning of Cr and Mo was found within these particles, whereas other elements were not detected as illustrated in Figure 9.

In additively manufactured materials, the solidification velocity and the thermal gradient are relatively high (V ~ 10 1 -10 2 °C.mm -1 and G ~ 10 2 -10 4 °C.s -1 ) compared to those measured in conventional casting (V ~ 10 -1 -10 -2 °C.mm -1 and G ~ 1-10 °C. s -1 ). Thus, the degree of segregations is expected to be much lower in S-EBM samples than in conventionally cast ones, see e.g. [START_REF] Ramsperger | Solution Heat Treatment of the Single Crystal Nickel-Base Superalloy CMSX-4 Fabricated by Selective Electron Beam Melting[END_REF][START_REF] Ramsperger | Microstructure of the Nickel-Base Superalloy CMSX-4 Fabricated by Selective Electron Beam Melting[END_REF][START_REF] Ruttert | Impact of hot isostatic pressing on microstructures of CMSX-4 Ni-base superalloy fabricated by selective electron beam melting[END_REF].

However, the presence of particles containing significant amounts of Cr and Mo at the grain boundaries suggests a local compositional enrichment of these elements. Thus, in order to accurately measure the composition of these particles, site specific lift-outs were prepared for atom probe tomography (APT). Specimens were extracted from regions near the tip of cracks, typically regions such as the one depicted in Figure 8b. The second phase particles showing a Cr-and Mo-enrichment revealed by EDX were clearly identified by APT as intergranular borides. Three different types of intergranular borides were identified with stoichiometry ratios corresponding to MB, M2B and M5B3, respectively, as summarized in Table 1. All three types of borides were observed to be Mo-and Cr-rich whereas other elements contribute in smaller quantities. An APT 3D reconstruction containing both MB and M2B borides alongside an intergranular particle is shown in Figure 10a-b. Note that the larger the columnar grain width (i.e. the lower the grain boundary density), the more the grain boundaries were decorated by borides, see comparison between Figure 8a and Figure 8b. In other words, it means that the boride density per grain boundary unit length is higher in presence of wide grains. Note that the borides preferentially sit at HAGB and that LAGB were found to be boride-free.

Proximity histograms, which are elemental composition profiles measured as a function of the distance from the γ/M2B interface are presented in Figure 10c. The interface is defined as an iso-concentration surface encompassing the region of the data containing more that 34 at.% Ni. The composition profiles across the γ/M2B interface differ from the typical S-curve observed for such interfaces, see typically [START_REF] Kontis | On the effect of boron on grain boundary character in a new polycrystalline superalloy[END_REF]. Three distinct regions can be seen in Figure 10c: a first region on the left corresponding to the boride, a second region highlighted in red, see the change of slope observed at a distance of about 2.5 nm in Figure 10c, and a third region (on the right) corresponding to the γ precipitate. This 2-4nm wide film at the interface between the boride and the γ precipitate was not previously observed or reported in samples cast by traditional routes. Its composition does not correspond to any of the typical expected phases in this alloy. It is likely that this thin film corresponds to the residual liquid present in the last stage of solidification. Our measurements indicate that this region is highly enriched in B with a concentration measured in this case ~2-3% at.%, i.e. a quantity substantially higher than the B content in the bulk chemistry (0.083 at.%).

Presence of thermal stresses

The presence of an intergranular liquid film itself is not sufficient to result in cracking.

Indeed, solidification shrinkage or thermal stresses should be involved to trigger the hot cracking mechanism.

Cracks are an indirect observation of the presence of thermal stresses pulling on the liquid film during the last-stage of solidification. The development of significant thermal stresses during the solidification and subsequent cooling can lead to residual stresses in the final parts. An original method was employed to demonstrate the presence of significant residual stresses inherited from the S-EBM process. As-built S-EBM samples were annealed to 1190°C/4h and subsequently air quenched. Such an annealing corresponds to the typical solutionizing heat treatment of the alloy investigated. Recrystallized grains can be observed, see in regions marked by black arrows in Figure 11. Note that before annealing, all the grains exhibited a columnar morphology. The observation of recrystallized grains suggests the presence of plastic strain during the additive manufacturing operations, leading to a high density of dislocations, the driving force for recrystallization. The latter observation indicates clearly that substantial strain energy was stored in the microstructure of the as-built S-EBM samples. This supports the idea that significant residual stresses are likely present in the asbuilt parts. Note that even though thermal stresses are thought to play a crucial role in the hot cracking mechanism, it should be highlighted that the level of residual stresses is expected to be significantly lower in S-EBM parts than in SLM parts because of the preheating stage that helps to limit temperature gradients during local melting [START_REF] Sochalski-Kolbus | Comparison of Residual Stresses in Inconel 718 Simple Parts Made by Electron Beam Melting and Direct Laser Metal Sintering[END_REF].

Discussion

Origin of the microstructural gradients

The gradient in the width of the columnar grain can be seen analogous to single crystal growth by the Bridgman method. The Bridgman technique is based on directional solidification by moving a melt from the hot zone to the cold zone of the furnace [START_REF] Reed | THE SUPERALLOYS: Fundamentals and Applications[END_REF]. In the S-EBM samples, the situation is very similar as the hot zone corresponds to the layers under construction (top layers), the underlying layers, typically thermally unaffected by the addition of new layers, act as the cold zone. Over the first millimeter, the microstructure consists of relatively small randomly textured columnar grains resulting from melting of the powder particles. Then columnar grains oriented along the <001> direction grow selectively by epitaxy across several layers along the build direction corresponding to the main direction of the thermal gradient. The re-melting of several layers promotes this epitaxial growth. At 20 mm, the microstructure is strongly textured with a <001>-fiber texture and consists of relatively wide columnar grains (~1 mm). Finally, when the subsequent deposited layers are melted, up to 3-4 layers underneath are also re-melted and therefore the wide columnar grains oriented along the 001 direction act as seed for the next layers. In other words, it means that by selecting appropriate melting strategies, it should be possible to grow single crystal samples within the S-EBM machine.

The gradient in the γ' size can be explained by the high temperature process (preheating ~1050°C). Depending on the Z-position probed, the different layers were annealed for different times and consequently significant growth and coarsening occurred at the bottom of the sample. Just below the top surface, fine precipitates were observed because this region was not re-melted neither re-heated (last layer). The fine precipitates observed in the top region might have formed at low temperature during the cooling stage [START_REF] Sames | Feasibility of in situ controlled heat treatment (ISHT) of Inconel 718 during electron beam melting additive manufacturing[END_REF]. The cooling rate in this particular region was insufficiently slow to achieve significant coarsening of the precipitates.

Cracking mechanism

Solidification cracking vs. liquation cracking

The observation of the fracture surface of the tensile specimens shows that cracks develop in presence of liquid. Both mechanisms, namely solidification cracking and liquation cracking, require the presence of liquid films. Therefore, it allows us to exclude any mechanism occurring at the solid-state. That being said, discriminating liquation cracking from solidification cracking is not straightforward as it could be in directional solidification experiments, especially in our case because the thermal path can turn out to be very complex:

melting, re-melting, partial-re-melting, cyclic annealing etc. Again welding literature of Nibased superalloys can help us to tend towards one mechanism rather than another.

Ojo et al. [START_REF] Ojo | Study of the fusion zone and heat-affected zone microstructures in tungsten inert gas-welded INCONEL 738LC superalloy[END_REF] have reported the formation of secondary solidification constituents formed from the interdendritic liquid such as MC-type carbides, M3B2-boride and Ni7Zr2 in a fusion zone of a Tungsten Inert Gas welded Inconel 738LC superalloy. The liquation of those secondary solidification constituents was suspected to be at the origin of the microfissuring observed in the HAZ. Note that carbides were never observed in the material investigated here. Unlike in Ojo et al. in [START_REF] Ojo | Study of the fusion zone and heat-affected zone microstructures in tungsten inert gas-welded INCONEL 738LC superalloy[END_REF], a close examination of the cracked boundaries using high magnifications SEM images did not reveal the presence liquated borides. In addition, no obvious traces of γ' constitutional liquation as reported by Zhong et al. in [START_REF] Zhong | Boundary liquation and interface cracking characterization in laser deposition of Inconel 738 on directionally solidified Ni-based superalloy[END_REF] or Ojo et al in [START_REF] Ojo | Contribution of constitutional liquation of gamma prime precipitate to weld HAZ cracking of cast Inconel 738 superalloy[END_REF] was detected. Finally, the fracture surface exhibits a fully dendritic morphology. This tips the scales in favor of solidification cracking. We also suspect that the B-local enrichment intensifies when the liquid film is melted several times which is the case in the S-EBM process. Further investigations would be required to solve undoubtedly the cracking mechanism. These borides were also essentially composed of Cr and Mo and identified as M2B similarly to some of the borides analyzed in the present work. Interestingly, the size of the borides observed in [START_REF] Cadel | Atom probe tomography investigation of the microstructure of superalloys N18[END_REF] was about two orders of magnitude smaller (15 nm) than those observed in the present study (~1 μm). The presence of such Mo and Cr containing borides support the idea of local enrichments of these solutes to grain boundaries. Lemarchand et al. [START_REF] Lemarchand | Investigation of grain-boundary structuresegregation relationship in an N18 nickel-based superalloy[END_REF] emphasized such local enrichment in Cr, B and Mo along the grain boundary in the form of a continuous film wetting the grain boundaries and similar local compositional enrichment was also reported in an Astroloy [44][START_REF] Letellier | Atomic-Scale Investigation of Grain-Boundary Microchemistry in the Nickel-Based Superalloy Astroloy with a 3-Dimensional Atom-Probe[END_REF][START_REF] Letellier | Direct Observation of Boron Segregation at Grain-Boundaries in Astroloy by 3D Atomic Tomography[END_REF]. Local B enrichment in the regions that are going to become the grain boundaries upon impingement is expected to contribute to solidification cracking [START_REF] Cloots | Investigations on the microstructure and crack formation of IN738LC samples processed by selective laser melting using Gaussian and doughnut profiles[END_REF]. Indeed, even in very small amount, some elements can have a deleterious effect on the solidification cracking susceptibility of Ni-based superalloys. The literature dealing with welding and weldability of Ni-based superalloys turns out to be relevant towards the effect of minor elements such as B on the hot cracking susceptibility, see e.g. [START_REF] Grodzki | Effect of B, Zr, and C on Hot Tearing of a Directionally Solidified Nickel-Based Superalloy[END_REF][START_REF] Zhou | Effect of grain boundary fraction on castability of a directionally solidified nickel alloy[END_REF][START_REF] Cloots | Investigations on the microstructure and crack formation of IN738LC samples processed by selective laser melting using Gaussian and doughnut profiles[END_REF][START_REF] Heydari | Hot tearing in polycrystalline Ni-based IN738LC superalloy: Influence of Zr content[END_REF].

This detrimental influence was attributed to (i) its strong segregation tendency to grain boundaries [START_REF] Lemarchand | Investigation of grain-boundary structuresegregation relationship in an N18 nickel-based superalloy[END_REF][START_REF] Kontis | On the effect of boron on grain boundary character in a new polycrystalline superalloy[END_REF][START_REF] Cadel | Atom probe tomography investigation of the microstructure of superalloys N18[END_REF][START_REF] Raabe | Grain boundary segregation engineering in metallic alloys: A pathway to the design of interfaces[END_REF][START_REF] Tytko | Microstructural evolution of a Nibased superalloy (617B) at 700 degrees C studied by electron microscopy and atom probe tomography[END_REF] and (ii) its ability to lower the solid/liquid interface energy [START_REF] Lippold | Welding Metallurgy and Weldability of Nickel-Base Alloys[END_REF] promoting cracking by extensive wetting of the solid dendrites by the liquid film still present during the very last stage of solidification. As it was shown in section 3.1., the grain boundary density (HAGB) decreases along the Z direction (Build Direction), resulting in shorter grain boundary lengths per unit area. This means that the degree of compositional segregation might be reduced in presence of small grains due to the resulting higher grain boundary density and lower specific Scheil type segregation content. Conversely, in presence of wide columnar grains (i.e. low grain boundary density), the degree of segregation during the solidification is expected to increase, resulting in a higher concentration of B and consequently to an increased volume fraction of borides, as confirmed by the comparison between Figure 8a andb. This leads us to suspect that if one aims at decreasing the hot cracking susceptibility, the boron content should be reduced to limit the formation of low melting liquid film.

Possible effect of micro-segregations on the liquid film

To evaluate roughly the possible effect of micro-segregations on the solidification of the investigated material, thermodynamic calculations using the ThermoCalc Software AB [START_REF] Andersson | THERMO-CALC & DICTRA, computational tools for materials science[END_REF] relying on the data base dedicated to Ni-based superalloys (TCNI8) were performed based respectively, on the equilibrium and non-equilibrium (Scheil-Gulliver: no diffusion in the solid state and infinite diffusion in the liquid) assumptions. The calculations were run by retaining only the phases that were observed in the microstructures of the S-EBM samples, namely: γ, γ', MB, M2B, Cr5B3 (for the M5B3 borides) and the liquid phase. Even if the situation is supposed to be more complicated in the S-EBM samples because of the very complex thermal path (re-melting, re-heating...), it is rather clear that in presence of micro-segregations (Scheil-Gulliver), the solidus point is significantly lower, leading to a larger solidification range: ~600°C whereas it was estimated to be ~200°C under equilibrium conditions, see Figure 12. Due to the inability to simulate back-diffusion in the solid, the Scheil model gives a good estimation of the liquidus temperature but overestimates solute segregation and thus predicts very low solidus temperature.

This result means that when micro-segregations are considered, liquid films will prevail and hence still be present also at lower bulk temperatures. The argument of the solidification range, i.e. strictly speaking the difference between the solidus and liquidus temperatures, has to be considered carefully regarding the solidification cracking susceptibility. Actually, some references doubt that the segregation-dependent solidification range is a critical factor for cracking, see e.g. [START_REF] Zhang | Hot tearing of nickel-based superalloys during directional solidification[END_REF][START_REF] Rappaz | A new hot-tearing criterion[END_REF]. It is rather suggested in these papers [START_REF] Rappaz | A new hot-tearing criterion[END_REF] that the factor that matter more is the difference in temperature when the fraction of solid fs changes from ~0.95 (coalescence fraction) to 1. In the literature, it is usually called the "critical temperature range" [START_REF] Zhang | Hot tearing of nickel-based superalloys during directional solidification[END_REF]: ΔTCTR = T(fS = 1) -T(fS = 0.95). For our current alloy ΔTCTR was estimated to be ~100°C assuming thermodynamic equilibrium and ~350°C assuming Scheil-Gulliver conditions. It is also interesting to note that the variation of the fraction of solid with temperature, dfS/dT, within this critical temperature range ΔTCTR is usually used in the definition of hot tearing criteria, see e.g. [START_REF] Zhang | Hot tearing of nickel-based superalloys during directional solidification[END_REF][START_REF] Rappaz | A new hot-tearing criterion[END_REF]. dfS/dT appears to be crucial for the hot tearing resistance because it determines the strain and strain rate during solidification.

For the Scheil assumption, the evolution of the liquid composition has been calculated so as to assess the liquid B-enrichment in the last stages of solidification. Interestingly, the liquid phase was found to be significantly enriched in boron in the temperature critical range (fs ~0.95-1), see Another important point that needs to be highlighted when discussing the mechanism leading to solidification cracking is the different coalescence behavior of dendrites belonging to the same grain (intragranular coalescence) and of dendrites located on both sides of a grain boundary (intergranular coalescence). The theoretical model on the very last stage of solidification developed by Rappaz et al. [START_REF] Wang | Solidification cracking of superalloy single-and bicrystals[END_REF][START_REF] Rappaz | Last-stage solidification of alloys: Theoretical model of dendrite-arm and grain coalescence[END_REF] is particularly relevant to the discussion of the present results. In [START_REF] Rappaz | Last-stage solidification of alloys: Theoretical model of dendrite-arm and grain coalescence[END_REF], Rappaz relied on the concept of "attractive" and "repulsive" boundaries depending on the crystallographic misorientation. An attractive boundary is defined in this work as a grain boundary where the coalescence among dendrite secondary arms occurs during the last stage of solidification contrary to a repulsive one for which a liquid film remains stable at lower temperatures delaying the occurrence of dendrite secondary arms coalescence. To discriminate attractive from repulsive boundary, Rappaz et al. [START_REF] Wang | Solidification cracking of superalloy single-and bicrystals[END_REF][START_REF] Rappaz | Last-stage solidification of alloys: Theoretical model of dendrite-arm and grain coalescence[END_REF], compare the grain boundary energy, γgb to twice the solid-liquid interface energy γsl. When γgb < 2γsl, the liquid film is unstable and bridging occurs between dendrite secondary arms: "attractive boundary" (corresponds typically to dendrites belonging to the same grain).

When γgb > 2γsl, the liquid film remains stable to a lower temperature ("repulsive" boundary) corresponding to the critical coalescence undercooling, denoted ΔTcc and defined for a pure metal as: ΔTcc= (γgb-2 γsl)/ΔSf δ, where ΔSf is the entropy of fusion per unit volume and δ the thickness of the diffuse interface (~1 nm). In other words, it means that the dendrites arms within the same grain start to coalesce at a higher temperature than the dendrites arms located on both sides of a grain boundary.

This model provides a physical background to explain why cracks develop preferentially at HAGBs and that LAGBs are not affected by solidification cracks. For HAGBs, the situation is rather clear. Indeed, the grain boundary energy is likely to exhibit high values (except for CSL-Σ specific boundaries), therefore γgb-2γsl is expected to be > 0. As a result, ΔTcc is significant and consequently the liquid film remains stable at lower temperatures, see Figure 13a.

For LAGB (misorientation < 15°), the grain boundary energy drops suddenly when the misorientation decreases. Thus γgb-2γsl becomes smaller and smaller as the misorientation decreases. In this situation, ΔTcc is negligible until it is reduced to 0 for a critical value of γgb corresponding to the following case: γgb-2γsl = 0 (Figure 13b). The liquid film becomes unstable and bridging occurs. In these regions, strains could be transmitted by the bridges existing between the dendrite secondary arms avoiding the formation of cracks. In other words it means that depending on the misorientation, the behaviour of grain boundaries changes [START_REF] Wang | Liquid channel segregation and morphology and their relation with hot cracking susceptibility during columnar growth in binary alloys[END_REF]. Based on Rappaz' model [START_REF] Rappaz | Last-stage solidification of alloys: Theoretical model of dendrite-arm and grain coalescence[END_REF], it turns out that HAGBs remain wetted by the liquid film at lower temperature compared to LAGB, see Figure 13.

Stresses/strain required to trigger hot cracking

As mentioned previously, solidification cracking occurs when two conditions are combined: (i) presence of a liquid film and (ii) presence of sufficient thermal stresses. The first condition was largely discussed in section 4.2.1, the second one is discussed below.

At a macroscopic scale, the thermal stresses were associated with the solidification of the successive layers. The top layer is under tensile stress, because its shrinkage is limited by the already solidified layers. These tensile stresses in the top layers are thought to be very detrimental regarding the hot cracking issue as it would contribute to pull on the liquid film wetting the dendrites during the last stage of solidification. When the temperature reaches the coalescence temperature, bridging occurs between the dendrite secondary arms preventing liquid feeding and leading to the formation of solidification pores [START_REF] Zhang | Hot tearing of nickel-based superalloys during directional solidification[END_REF].

Some thermal stresses may also develop at the scale of the melt pool. Welding literature is again helpful to understand how thermal stresses developed and eventually lead to tensile stresses finally resulting in hot cracking. The work done by Feng et al. [START_REF] Feng | Thermal stress development in a nickel based superalloy during weldability test[END_REF] and Babu et al. [START_REF] Babu | Joining of nickel base superalloy single crystals[END_REF] show that the local stress fields that develop in the vicinity of the weld pool have a strong spatial dependency (position of the weld pool within in the specimen). They showed that a transverse tensile stress will lead to longitudinal weld centerline cracking whereas a longitudinal tensile stress will lead to transverse weld cracking. However, it is difficult to rationalize the nature of the cracks in our case: transverse or longitudinal centerline cracking based on those results. Indeed, the situation is much more complex in the S-EBM samples:

Feng et al. [START_REF] Feng | Thermal stress development in a nickel based superalloy during weldability test[END_REF] have investigated the stress fields around the weld pool during a single track.

Here multiple weld tracks are required to melt a layer with a given degree of overlapping between the tracks and a raster-melt scan strategy is used. The stress fields are suspected to be significantly affected by the melting strategy used. For example, some authors have adapted the melting strategy in the S-EBM process to limit the development of thermal stresses and consequently of hot cracks, see e.g. [START_REF] Ramsperger | Microstructure of the Nickel-Base Superalloy CMSX-4 Fabricated by Selective Electron Beam Melting[END_REF].

As shown in the experimental results, no cracks were detected over the first few millimeters. It was reported in [START_REF] Mercelis | Residual stresses in selective laser sintering and selective laser melting[END_REF] that the stresses can increase along the build direction and can likely contribute to generate cracking when the stresses are high enough [START_REF] Mercelis | Residual stresses in selective laser sintering and selective laser melting[END_REF]. Therefore, it was thought that any cracking phenomenon occurs until a critical Z position, i.e. a critical level of thermal stresses is achieved. Another possible explanation is the effect of grain size or grain boundary density. Indeed, the crack free region corresponds to the lower part of the samples where the density of grain boundaries (HAGB) was the highest, see Figure 4a. Those results are consistent with the fact that the hot tearing resistance increases when the grain boundary fraction increases, see [START_REF] Zhang | Effect of grain-boundary characteristics on castability of nickel-base superalloys[END_REF][START_REF] Zhou | Effect of grain boundary characteristics on hot tearing in directional solidification of superalloys[END_REF][START_REF] Zhou | Effect of grain boundary fraction on castability of a directionally solidified nickel alloy[END_REF]. The beneficial effect of having small grains towards the hot tearing resistance is attributed to the fact that the strain is accommodated by more interfaces, i.e. HAGB.

Summary and Conclusions

The main conclusions emerging from the present work are:

 Samples of a non weldable Ni-based superalloy have been produced

successfully by S-EBM without lack of fusion type defects.

 The S-EBM-samples exhibit a heterogeneous microstructure at the grain scale: gradient of columnar grain size along the building direction as well as within the grain: gradient of γ'-precipitate size along the building direction.

 Various types of defects such as gas porosity, shrinkage pores and intergranular cracks have been identified. The origin of each category of defect was determined with a specific emphasis put on the origin of the cracks.

 The mechanism at the origin of crack formation in parts made of a nonweldable Ni-based superalloy made by S-EBM requires the presence of liquid films.

 The hot cracking susceptibility has been found to depend on the nature of the grain boundary misorientation, only HAGBs were affected by cracking.

 The presence of micron sized borides with stoichiometries corresponding to MB, M2B and M5B3 as confirmed by APT suggests a significant local enrichment in B in the vicinity of the grain boundaries.

 Local compositional enrichment by minor elements, in particular B, are thought to play a significant role to maintain a liquid film at lower temperature in comparison with the theoretically assumed equilibrium solidus temperature.

 The grain boundary energy which varies significantly between LAGB and HAGB has been also pointed out as a factor that can affect the stability of the liquid film.

These results lead us to suggest a guideline to limit hot cracking in a non-weldable superalloy produced by S-EBM. It can be first suggested to reduce the boron content for limiting the formation of B-rich liquid films. Adjusting the process parameters to achieve solidification routes leading to fine grained microstructures (~100 μm) more resistant to hot cracking can also be considered to limit the development of hot cracks. The strain generated during the S-EBM process could be accommodated by a high density of interfaces while compositional segregation leading to the prevalence of a liquid film during the last stage of solidification would be reduced. Optimization of the processing parameters to produce crackfree samples with fine columnar grains, equiaxed grains or a single crystal is the topic of ongoing research in that context.

Figure 1 .Figure 2 .

 12 Figure 1. Powder characteristics: (a) typical powder morphology as-revealed by SEM under secondary electron contrast (SE); (b) IPF-EBSD map revealing that the powder particles are polycrystalline; (c) SE-SEM micrograph illustrating gas porosity within initial powder particles.

Figure 3 .Figure 6 .Figure 7 .

 367 Figure 3. Typical microstructure of an as-built specimen observed in the central part on a XZ cross section. (a) Optical micrograph showing the macrostructure. (b) EBSD-IPF maps showing the columnar grain size evolution and (c) BSE-SEM micrographs illustrating the size of the γ' precipitates at different positions along the build direction. Black dashed square boxes indicate the position where the EBSD maps were acquired. Note that EBSD maps and BSE-SEM images are not necessarily shown at the same location as the evolution of grain and precipitate size do not occur at similar position along the build direction.

Figure 8 .Figure 9 .

 89 Figure 8. (a) Low magnification BSE-SEM micrograph and (b) high magnification BSE-SEM micrograph taken at Z = 20 mm (cracked region) showing the tip of a cracked HAGB. Note that only one HAGB can be observed and that it is decorated by second phase particles pointed out by white arrows. Damage is highlighted in red. (c) BSE-SEM micrograph taken at Z = 5 mm (uncracked region). Several HAGB can be observed but not all of them are decorated by second phase particles marked by black arrows.

Figure 10 .Figure 11 .Figure 12 .

 101112 Figure 10. (a) A 3D atom probe reconstruction containing MB and M2B type borides alongside with an intergranular γ particle with iso-concentration interfaces shown at 34 at.% Ni (green) and 46 at.% B (orange). (b) Zoom on the region at the interface showing the region depicted in red in (c) with an iso-density surface with 4 Cr atoms per nm 3 . (c) Concentration profile analysis at γ/M2B interface showing concentration of Ni, Al, Co, Mo, B and Cr. Highlighted in red is the region suspected to be a "tracer" of the residual liquid film. Error bars are shown as lines filled with colour.

Figure 13 .

 13 Figure 13. Schematic illustration of the last stage of solidification during columnar dendritic solidification : (a) for a "repulsive" high angle grain boundary (HAGB) and (b) for an "attractive" low angle grain boundary (LAGB). Liquid in white, solid in gray.

Table 1 .

 1 Chemical composition of borides as measured by atom probe tomography (at.%).

		B	Cr	Mo	Other elements
	MB	52.2	25.1	14.6	Bal.
	M2B	33.3	20.1	35.5	Bal.
	M5B3	36.6	20.3	36.8	Bal.

Table 2 .

 2 This supports the tendency to form low melting B-rich liquid films in the critical temperature range.

	Temperature (°C)	1340	1200	1100	1050	1000	800

Table 2 .

 2 Boron content in the liquid phase calculated at different temperatures using ThermoCalc based on the Scheil-Gulliver assumption. 4.2.2.3. Last-stage solidification: theoretical model of Rappaz et al.

Acknowledgements

This work was performed within the framework of the Center of Excellence of Multifunctional Architectured Materials ''CEMAM'' n°AN-10-LABX-44-01 funded by the ''Investments for the Future Program''. Dr Michel Suery and Dr. Frederic De Geuser are gratefully acknowledged respectively, for fruitful discussions and his help to build a collaboration between both groups. P. K., E. A. J., D. R. and B. G. are grateful to U. Tezins and A. Sturm for their technical support of the focus ion beam and atom probe tomography facilities at the Max-Planck-Institut für Eisenforschung.


Nathalie Prud'homme
email: nathalie.prudhomme@u-psud.fr

Vassilios Constantoudis

Asiya E Turgambaeva

Vladislav V Krisyuk

Diane Samélor

François Senocq

Constantin Vahlas

Chemical vapor deposition of Cu films from copper(I) cyclopentadienyl triethylphophine: Precursor characteristics and interplay between growth parameters and films morphology

Keywords: Copper(I) cyclopentadienyl triethylphophine, Copper, Thin films, Metal-organic chemical vapor deposition, Coalescence, Nucleation, Morphology, Image analysis

à la diffusion de documents scientifiques de niveau recherche, publiés ou non, émanant des établissements d'enseignement et de recherche français ou étrangers, des laboratoires publics ou privés.

Introduction

Copper films are widely used as interconnect material seed layers [START_REF] Zhang | Effects of 2mercaptopyridine and Janus Green B as levelers on electrical resistance of electrodeposited copper thin film for interconnects[END_REF][START_REF] Chae | Polycrystalline-silicon thin-film transistor fabricated with Cu gate controlled morphology by plating mode[END_REF][START_REF] Solanki | Atomic layer deposition of copper seed layers[END_REF][START_REF] Thompson | Vapor phase deposition of copper films with a Cu(I) beta-diketiminate precursor[END_REF], for device encapsulation [START_REF] Yan | An investigation on air stability of copper phthalocyaninebased organic thin-film transistors and device encapsulation[END_REF], or as components of antimicrobial coatings [START_REF] Aktug | Characterization and investigation of in vitro properties of antibacterial copper deposited on bioactive ZrO2 coatings on zirconium[END_REF] due to their high electrical conductivity, high ductility, and biocidal nature. Such films are conveniently deposited on flat surfaces by evaporation or sputtering. However, these methods cannot be applied to the metallization of non-line-of-sight surfaces, such as trenches and vias on Si, or the inner walls of tubings such as polymer waveguides [START_REF] Xiong | Atomic layer deposition of copper thin film and feasibility of deposition on inner walls of waveguides[END_REF], because of their weak hole-filling capability and the resulting non-uniformity. As a characteristic example, specifications concerning the metallization of the internal surface of polymer waveguides include the deposition of a film composed of a several-microns-thick, highly conductive, corrosion/oxidation-resistant metal with smooth surface morphology and adherence to the composite substrate. These specifications are not necessarily compatible, since it is difficult to preserve high smoothness while promoting adhesion [START_REF] Duguet | Metallization of polymer composites by metalorganic chemical vapor deposition of Cu: Surface functionalization driven films characteristics[END_REF][START_REF] Addou | Metallization of carbon fiber reinforced polymers: Chemical kinetics, adhesion, and properties[END_REF]. The heat sensitivity of polymers gives rise to additional constraints, namely processing at low to moderate temperatures. Alternatively, chemical vapor deposition (CVD) can be used specifically in view of the possibility of achieving conformal films [START_REF] Addou | Metallization of carbon fiber reinforced polymers: Chemical kinetics, adhesion, and properties[END_REF][START_REF] Murarka | Advanced multilayer metallization schemes with copper as interconnection metal[END_REF][START_REF] Norman | A new metal organic-chemical vapor-deposition process for selective copper metallization[END_REF] combined with relatively elevated growth rates. Despite considerable work on the CVD of Cu, the reluctance of the process to yield smooth films remains to be overcome. Indeed, in general terms, the activation energy for the heterogeneous nucleation of Cu is higher than the activation energy for the growth of 2D structures, e.g. thin films ( [START_REF] Kodas | The Chemistry of Metal CVD[END_REF] and references therein). This situation results in rough, even discontinuous films, which are inappropriate for the targeted properties. For these reasons, research since the early days of this technique has focused on monitoring Cu film morphology through the appropriate selection of process conditions [START_REF] Psarellis | Investigation of reaction mechanisms in the chemical vapor deposition of al from DMEAA[END_REF].

Numerous copper precursors contain oxygen and/or halogen atoms in their ligands, or require oxygen-containing co-reactants for the deposition of copper [START_REF] Krisyuk | CVD of pure copper films from amidinate precursor[END_REF]. However, an oxygen-free precursor can be considered as more appropriate for the processing of metallic films, because alternative, halogen-based deposition chemistries require high process temperatures and they involve corrosive by-products. Copper(I) cyclopentadienyl phosphine derivatives were among the few oxygen-and halogen-free precursors that were investigated in CVD processes [START_REF] Chichibu | Chemical vapor-deposition of Cu film on SiO2 using cyclopentadienylcoppertriethylphosphine[END_REF][START_REF] Katagiri | Metalorganic chemical vapor deposition of aluminum-copper alloys films[END_REF][START_REF] Kondoh | Interconnection formation by doping chemical-vapor-depositon aluminum with copper simultaneously: Al-Cu CVD[END_REF][START_REF] Hampden-Smith | Chemical Vapor Deposition of Copper from Copper(I)Trimethylphospine c ompounds[END_REF][START_REF] Wilkinson | Cyclopentadienyl-triethylphosphine-copper(I) and biscyclopentadienylmercury(II)[END_REF][START_REF] Whitesides | A nuclear magnetic resonance study of .sigma.cyclopentadienyl(triethylphosphine)copper(I)[END_REF][START_REF] Cotton | Systematic preparation and characterization of pentahaptocyclopentadienylcopper(I) compounds[END_REF].

CpCuPEt3 (Cp = n 5 -C5H5, Et = C2H5) was synthesized and characterized in the late fifties, but it had not been used as a metalorganic CVD (MOCVD) precursor before the 1980s [START_REF] Dupuy | Laser induced deposition of copper from (triethylphosphine)cyclopentadienylcopper(I)[END_REF] [START_REF] Hara | Epitaxial Growth of CuGaS2 by Metalorganic Chemical Vapor Deposition[END_REF].

In the 1990s, it was used in combination with alanes for the co-deposition of Cu-Al films [START_REF] Katagiri | Metalorganic chemical vapor deposition of aluminum-copper alloys films[END_REF][START_REF] Kondoh | Interconnection formation by doping chemical-vapor-depositon aluminum with copper simultaneously: Al-Cu CVD[END_REF]. CpCuPEt3-based CVD processes operate at a relatively low temperature, typically 473 K, thus allowing surface functionalization of thermally sensitive parts. In our previous investigation of the physical and thermal properties of CpCuPEt3, [START_REF] Senocq | Thermal behaviour of CpCuPEt3 in gas phase and Cu thin films processing[END_REF] we defined the law of 𝑃 𝑠𝑎𝑡 vs. sublimation temperature in the 313-343K range. We demonstrated that the compound is stable below 343K and that it provides pure copper films in the 423-563K range. We also investigated the heterogeneous thermal decomposition mechanism of CpCuPEt3 by online and in situ mass spectrometry (MS), revealing all of the changes in the composition of the gas phase [START_REF] Turgambaeva | Mass spectrometric monitoring of the gas phase during chemical vapor deposition of copper from copper cyclopentadienyl triethylphosphine[END_REF]. However, no systematic investigation of the film morphology, especially regarding smooth films, has been reported for this suitable adequate Cu precursor, especially considering the previously mentioned unfavorable competition between nucleation and growth. Moreover, the process was observed to be unstable during the course of the various deposition experiments [START_REF] Turgambaeva | Mass spectrometric monitoring of the gas phase during chemical vapor deposition of copper from copper cyclopentadienyl triethylphosphine[END_REF]. This instability was attributed to the variation of the precursor characteristics from one batch to the other, revealing the need to purify the precursor and to revisit its thermal behavior.

The present contribution aims at investigating these two aspects, namely the precursor thermal behavior and the film morphologies.The purification protocol is presented and a saturatedvapor-pressure behavior is defined as a function of temperature for the purified precursor. The thermal behavior of the latter is established and the Arrhenius plot of the process is determined.

Finally, the morphology of the deposited Cu films is systematically investigated as a function of deposition temperature and precursor concentration, taken as representatives of the process parameters.

Experimental details

Precursor

As-received CpCuPEt3 (Strem Chemicals) was purified by sublimation at 313 K and 10 -2 Torr.

The elemental composition of the precursor was determined by ICP-MS for Cu and P, and by CHN microanalysis for C and H. Its thermal stability in the condensed phase was investigated by Thermo-Gravimetric and Differential Thermal Analysis (TGA and DTA) under nitrogen flow (40 L/min), from 298 K to 873 K, with a 10 K/min ramp.

The saturated vapor pressure (Psat) of the purified compound was determined with the aim of controlling the transport conditions in the CVD experiments. Data for Psat versus temperature were obtained by using a previously described static method. [START_REF] Juarez | Chemical vapor deposition in spouted bed reactors[END_REF] The precursor was conditioned in a homemade sublimator composed of tubing and a frit. In this setup, two valves are positioned before and after the sublimator, and a third one serves as bypass. A new batch of the compound was used for each investigated sublimation temperature. For each temperature, two sequential measurements were performed with the same batch of the compound, with the aim of monitoring the evolution of Psat due to the possible degradation of the compound after the first measurement. Between the two measurements, the chamber was evacuated to the base pressure. Based on the determined temperature dependence of Psat, the flow rate QCpCuPEt3 of CpCuPEt3 in the deposition chamber can be estimated from Eq. 1 [START_REF] Vahlas | An original device for the delivery of vapors to a deposition apparatus based on the sublimation in a fluidized bed[END_REF].

𝑄 𝐶𝑝𝐶𝑢𝑃𝐸𝑡 3 (𝑠𝑐𝑐𝑚) = 𝑄 𝑁 2 ,𝑝𝑟𝑒𝑐 × 𝑃 𝑠𝑎𝑡 (𝑇 𝑠𝑎𝑡 ) 𝑃 𝑠𝑢𝑏 -𝑃 𝑠𝑎𝑡 (𝑇 𝑠𝑎𝑡 ) [1 -𝑒𝑥𝑝 (- 𝛽 𝑀 𝑄 𝑁 2 ,𝑝𝑟𝑒𝑐 )] [1]
where M is the mass of precursor introduced in the sublimator in g, QN2,prec is the carrier flow rate of nitrogen through the sublimator in standard cubic centimeters per minute (sccm), Psat(Tsat) is the saturated vapor pressure in Torr at the sublimation temperature Tsat, Psub is the pressure in the sublimator in torr, and β is mass transfer coefficient in sccm/g. β values, typically between 0.2 and 4 sccm/g, were determined by mass loss of the precursor in each set of deposition conditions.

CVD of Cu films

CVD experiments of Cu films were performed in a vertical, cylindrical, stagnant-flow, warmwall MOCVD reactor, which has been described in detail in [START_REF] Xenidou | An experimental and computational analysis of a MOCVD process for the growth of Al films using DMEAA[END_REF][START_REF] Xenidou | Reaction and transport interplay in Al MOCVD investigated through experiments and computational fluid dynamic analysis[END_REF][START_REF] Aviziotis | Experimental and Computational Investigation of Chemical Vapor Deposition of Cu from Cu Amidinate[END_REF]. The reactor is equipped with a showerhead system, ensuring homogeneous flow of the reactants on the surface of the substrate. 10 x 10 x 1 mm 3 Si (100) wafers (Siltronix or ACM) and 304L stainless steel coupons of various centimeter size were used as substrates. The Si coupons were etched in a 1 ml HF:10 ml H2O solution for 1 min, sonicated in an acetone and ethanol for 5 min, and dried in Ar flow.

The stainless steel substrates were polished with 2400 grade SiC paper: they were then ultrasonically cleaned in ethanol for 5 min and finally dried in Ar flow. Substrates were placed horizontally at different radial positions, on a 58-mm-diameter susceptor (substrate holder) facing the shower plate and heated by a resistance coil just below the surface.

The sublimated CpCuPEt3 precursor was thermally regulated in the evaporator. Two nitrogen (Air Liquide, 99.999 %) lines were used: one for the transport of the precursor and one for the dilution of the input gas. Hydrogen was used as reducing gas-reactant, flowing at 25 sccm. The total flow rate was maintained at 325 sccm. 19 independent experiments were performed at 1333 Pa, in the temperature range 431-523K. Table 1 lists the ranges of the investigated operating conditions. Precursor flow rate / sccm 0.004 -0.070

Characterizations

We defined the deposition rates as the mass gained per unit surface over the experiment duration (µg.h -1 .cm -2 ). We estimated the rates by weighting the substrates before and after the deposition experiments in a Sartorius-Genius microbalance, with a ±1 µg precision, corresponding to approximately four monolayers of Cu on the entire coupon surface. We probed the elemental composition of the films by electron probe microanalysis (EPMA) with a CAMECA SX50 instrument and performed a structural investigation by grazing incidence X-Ray diffraction (GIXRD) with a Seifert XRD 3000 TT diffractometer using CuKα radiation and 2 degrees incidence angle. We investigated the morphology of the films by scanning electron microscopy (SEM) with a LEO 435 VP instrument operating at 15 kV and the topography of the films by white light interferometry with a MetroProTM Zygo New View 100 optical interferometer. We analyzed the top-down micrographs obtained by SEM measurements by using the software application nanoTOPO-SEM TM (version 11, Nanometrisis) that quantifies different characteristics of the film surface morphologies.

Results and discussion

Precursor

Sublimation of the as-received brown CpCuPEt3 powder yielded only white, translucent crystalline particles, in agreement with the expected appearance of the pure compound. [START_REF] Dupuy | Laser induced deposition of copper from (triethylphosphine)cyclopentadienylcopper(I)[END_REF] Table 2 presents the elemental molar ratios of P, C, and H to Cu for the as-received and purified compounds. The corresponding ratios for the stoichiometric, pure CpCuPEt3 are given for comparison. There were no significant differences, either between the as-received and the purified compound or between the compounds and the theoretical values. It is worth noting, however, that the content of P and C in the purified compound was closer to the theoretical values than that in the as-received compound. The difficulty in determining the oxygen content prevented us from identifying the nature of the impurities present in the as-received compound.

However, the minor deviation from the theoretical value for both as-received and purified samples indicates that the impurities present in the former concern the surface and not the bulk of the individual grains. Considering that the solid-to-gas transformation during the sublimation process takes place through the surface of the grains, such a surficially modified as-received compound is expected to present a lower saturated vapor pressure than the purified compound.

Moreover, we observed a progressive change in the color of the purified compound with time, when exposed in ambient conditions: it changed from translucent white to yellow to brownish after some days. This result renders the long term stability of CpCuPEt3 in ambient conditions questionable, especially in terms of the surface area of its grains. The potential instability, in turn, is expected to shift the effective vapor pressure of CpCuPEt3 toward lower values without, however, affecting the thermal decomposition in CVD conditions and, subsequently, the characteristics of the obtained Cu films. We conclude that the precursor remains stable during the sublimation process up to this temperature, at which point its degradation is initiated. The obtained sublimation law for the purified CpCuPEt3 is:

logPsat (Torr) = 6.493 -2272/T (K) (2)
and is illustrated in figure 2 by the straight line. The gray squares in Figure 2 represent the previously reported raw Psat vs. temperature points for the as-received compound, which resulted in a reported sublimation law of logPsat (Torr) = 9.671-3455/T(K) in the temperature range 313 K-353 K. [START_REF] Senocq | Thermal behaviour of CpCuPEt3 in gas phase and Cu thin films processing[END_REF] These points correspond to systematically lower Psat values than the ones of the purified compound. The observed Psat drop at 353 K was attributed to a noticeable decomposition of the precursor. The Psat of non-purified CpCuPEt3 was previously reported as 2.10 -3 Torr at 293 K, [START_REF] Dupuy | Laser induced deposition of copper from (triethylphosphine)cyclopentadienylcopper(I)[END_REF] but information on the adopted experimental protocol was not provided. This value is reported as a gray circle in Figure 2. For this sublimation temperature, the reported laws for the as-received and purified compounds yield Psat values equal to 8.10 -3

Torr and 55.10 -3 Torr, respectively: i.e., an order of magnitude higher for the purified compounds. We conclude that the purification of the precursor by sublimation enhances the volatility and the stability of CpCuPEt3.

Fig. 2: Saturated vapor pressure of the purified CpCuPEt3 as a function of inverse temperature for two sequential measurements (empty and grey diamonds for the first and second measurements, respectively). The straight line is the obtained sublimation law for the pure precursor (equation 1). Literature results, illustrated by grey squares [START_REF] Senocq | Thermal behaviour of CpCuPEt3 in gas phase and Cu thin films processing[END_REF] and the grey circle [START_REF] Dupuy | Laser induced deposition of copper from (triethylphosphine)cyclopentadienylcopper(I)[END_REF], are also reported for comparison.

Depending on the adopted operating conditions for the MOCVD experiments, the flow rate of the precursor, estimated from the Eq. 1, varies between 0.004 and 0.070 sccm. The temperature range where Psat of the purified precursor shows a stable, linear (in Arrhenius plot coordinates)

behavior indicates that it is sufficient to ensure precursor transport to feed the MOCVD process.

3.2. Growth rate and crystalline structure EPMA analysis of the films revealed pure Cu without impurities from heteroatoms such as phosphorus or carbon. For all experiments, we observed a 5-to 30-min nucleation delay. We evaluated it by visual inspection of the illuminated substrate surface through two windows mounted on the deposition chamber and defined it as the time needed for the surface color to change from shiny grey to brick red through a sharp color transition. In view of the observed time scale (min), this observation allows the convenient and rather precise determination of the incubation time [START_REF] Aviziotis | Investigation of the kinetics of the chemical vapor deposition of aluminum from dimethylethylamine alane: experiments and computations[END_REF]. The deposition time reported in Table 1 takes into account the proper nucleation delay for each experiment. This phenomenon was also probed in a specific experiment, which consisted of monitoring the gas phase by continuous online MS [START_REF] Turgambaeva | Mass spectrometric monitoring of the gas phase during chemical vapor deposition of copper from copper cyclopentadienyl triethylphosphine[END_REF]. Indeed, at 493 K, 11 min after the beginning of the precursor feeding, a sharp decrease in the intensity of the precursor and an increase in the intensity of the Cp2 product peaks were observed, indicating the beginning of deposition. The optical tracking of the color change of the substrates revealed the onset of copper-film growth at the same time. We conclude that the observed induction period is due to the delay in nucleation on the substrate at this temperature. During the MOCVD experiments, we did not find a correlation between the value of the nucleation delay and either the deposition temperature or the flow rate QCpCuPEt3.

Figure 3 presents the temperature dependence of the growth rate (GR), evaluated on silicon substrates. In this diagram, the mass gain per unit surface and time is divided by the precursor flow rate QCpCuPEt3 (µg.h -1 .cm -2 .sccm -1 ). Presenting the growth rate in this manner overcomes the difference among the different experiments concerning the precursor feeding. Only one region is notable in this semilogarithmic plot of GR/QCpCuPEt3 vs. reciprocal temperature. In the temperature range of 423 K-523 K, the film growth rate is almost constant with a slight increase with increasing temperature. The apparent activation energy of the global deposition mechanism, obtained from the slope of the linear plot, is low (roughly 10 kJ/mol). Krisyuk et al. reported a similar behavior for copper deposition, namely a weak dependence of the GR in the temperature range of 513 K-573 K, corresponding to a diffusion-controlled regime. [START_REF] Krisyuk | CVD of pure copper films from a novel amidinate precursor[END_REF] They also observed a sharp linear increase of the GR at temperatures up to 513 K, indicative of a surface-reaction-controlled regime with an activation energy of 65 kJ/mol. In another study on MOCVD of Cu, Addou et al. reported two regions in the Arrhenius plot of the GR on Si substrate: a diffusion-limited regime with a minor variation of the GR (468 K-493 K) and a surface-reaction-limited region at lower temperatures, for which they report an activation energy equal to 50 kJ/mol. [START_REF] Addou | Metallization of carbon fiber reinforced polymers: Chemical kinetics, adhesion, and properties[END_REF] Interestingly, the activation energy in the diffusion-limited regime, estimated from their reported Arrhenius plot (Figure 1 in [START_REF] Addou | Metallization of carbon fiber reinforced polymers: Chemical kinetics, adhesion, and properties[END_REF]), is 12 kJ/mol: i.e., almost identical to the present one. Beach et al. [START_REF] Beach | Low temperature chemical vapor deposition of high purity copper from an organometallic source[END_REF] studied copper deposition from CpCuPEt3 on Cr/Cu-seeded Si substrate and determined an activation energy of 75 kJ/mol in the temperature range 423-483 K. From these different reports, we can conclude that the very low value of activation energy obtained in the present work corresponds to a diffusion-limited regime.

Moreover, we observed that an increase of the precursor concentration in the input gasthrough the increase of the sublimation temperature and/or the increase of the nitrogen flow rate through the evaporatorresults in an increase in the growth rate. This result applies also to a diffusion-controlled regime in these conditions.

We also observed a decrease in the GR with decreasing deposition time when all the other process parameters were kept constant. This is in agreement with the results obtained by online MS monitoring of the gas phase during the deposition experiments. [START_REF] Turgambaeva | Mass spectrometric monitoring of the gas phase during chemical vapor deposition of copper from copper cyclopentadienyl triethylphosphine[END_REF] During deposition at 493 K, MS monitoring revealed a monotonic increase in the intensity of the precursor and a decrease in the intensity of the deposition reaction products with deposition time at low deposition temperature. This confirmed a gradual decrease in the growth rate. Such non-steadystate behavior may be attributed to the saturation of the growing surface with adsorbed phosphine, which hinders precursor supply to the deposition zone. Indeed, Holl et al. [START_REF] Holl | Surface reactivity of alkylgold(I) complexes -substrate -selective chemical-vapor-deposition of gold from RAuP(CH3)3 (R=Ch2Ch(3), Ch3) at remarkably low-temperatures[END_REF] reported that PEt3 is well adsorbed on the copper surface in the investigated temperature range, and CVD from phosphine-containing precursors should be performed at a higher temperature to desorb the surface-adsorbed PEt3. The phenomena observed in our experiments agree with the growth-rate behavior. The desorption of the cumulated decomposition products, such as the available surface-nucleation sites. This, in turn, allows for a corresponding increase in the Cu growth rate.

Surface morphology

Figure 4 presents SEM surface micrographs of Cu films deposited at different temperatures and precursor feeding rates, from 431-523 K and 0.006-0.070 sccm, respectively. Further information on the deposition conditions of the films shown in Figure 4 is presented in Table 3. For all the investigated deposition conditions, the films presented a nodular smooth morphology that is characteristic of a Wolmer-Weber type island growth. The grain size was approximately 40-60 nm for the film deposited with QCpCuPEt3=0.006 sccm (fig. 4a and4b); it increased with increasing CpCuPet3 flow rate and attained 90-110 nm for the film deposited at 493 K and 0.02 sccm (fig. 4f). XRD analysis revealed that the size of Cu crystallites, estimated using Sherrer's equation from the (111) diffraction peak, also increased with increasing precursor feeding rate. Indeed, the crystallite size was 20 nm and 40 nm with 0.01 sccm and 0.07 sccm CpCuPet3 flow rate, respectively. The simultaneous increase of the size of crystallites and grains with increasing precursor flow rate resulted in the coalescence of the grains (fig. 4g and 4h) and, consequently, yielded films with improved continuity. On the other hand, the grain size decreased slightly with increasing substrate temperature (fig 4e , 4d and4c).

When the deposition time was longer than 210 min (Table 3 and 3. We used top-down SEM images to calculate the coverage of silicon substrates by copper and the period (if any) of the morphology of deposited material. Figure 6 shows the two main tools we used to analyze the SEM images: the histogram of pixel intensities of images (Fig. 6a) and the radially averaged Fourier spectrum (Fig. 6b). Not surprisingly, the micrographs are characterized by bimodal histograms with two peaks corresponding to the dark regions of the substrate and the brighter pixels of the deposited material, respectively. The bimodal character dictated the definition of the threshold at the minimum between the peaks (see Fig. 6a where the threshold is defined at ~0.27). Additionally, we used Fourier analysis to identify and measure the periodicity of the deposited films. Figure 6b shows a typical Fourier spectrum (amplitude) of a top-down SEM image, in which the prominent peak justifies the presence of periodicity in the deposited morphology, with a period defined by the inverse of the peak frequency. Using these methods, we measured the coverage and period of the deposited copper films. The results are shown in Figure 7 for the majority of the images displayed in figure 4 to illustrate the impact of temperature and precursor flow rate. The results show that when the deposition temperature decreased from 493 K (Fig. 4d) to 431 K (Fig. 4e), the film coverage shrank while the period of their morphology elevated from 210 nm to 350 nm. The latter behavior is consistent with our expectation that the decrease in the deposition temperature of constraints the desorption of decomposition products, namely triethylphosphine, resulting in less dense morphologies with a greater period. This is also consistent with the observations by Shima et al. that a lower deposition temperature leads to a larger nucleate density, which suggests that Cu CVD conditions follow the general nucleation behavior as physical phenomena: the CVD reaction is not the limiting step to promote nucleation. [START_REF] Shima | Comparative Study on Cu-CVD Nucleation Using beta-diketonato and Amidinato Precursors for Sub-10-nm-Thick Continuous Film Growth[END_REF] A second observation from the results in Figure 7 pertains to the effects of precursor flow rate.

We observed a noticeable increase in period from 0.02 sccm (Fig. 4f) to 0.04 sccm (Fig. 4g) and 0.07 sccm (Fig. 4h and Fig. 5c), whereas the coverage remained almost similar with a slight increase from 0.78 to 0.84 and 0.86, respectively. The fact that the morphologies of figures 4g and 4h become broader with much wider local features and uniformity points to the predominance of coalescence over nucleation. The morphologies of figures 4g and 4h (channellike structure) were also observed by Hagen et al., who attributed these features to a coalescence becoming less complete because of the decrease of curvature gradients when the deposition is further continued. [START_REF] Hagen | Island Coalescence during Film Growth: An Underestimated Limitation of Cu ALD[END_REF] Shima et al. also observed an increased coalescence with increasing amounts of precursor in the gas phase and predicted that a higher partial pressure of the precursor would result in a a larger Cu nucleation density, which is confirmed by our observations. [START_REF] Shima | Comparative Study on Cu-CVD Nucleation Using beta-diketonato and Amidinato Precursors for Sub-10-nm-Thick Continuous Film Growth[END_REF] The methodology in our study provides a quantitative assessment of the contribution and interplay of the two fundamental mechanisms of nucleation and coalescence. These results show that a high precursor flow rate QCpCuPEt3 > 0.02 sccm) allows the predominance of coalescence effects over nucleation processes. The channel-like morphology (fig 4g and4h) obtained in these conditions should impart a lower value of resistivity than that imparted by the island morphology. Indeed, the electrical resistivity of copper film prepared by ALD was 17 µΩ.cm for island morphology, as compared with the much lower value of 7.6

µΩ.cm for channel-like morphology and 1.7 µΩ.cm for the bulk, monocrystalline Cu. [START_REF] Hagen | Island Coalescence during Film Growth: An Underestimated Limitation of Cu ALD[END_REF]

Conclusions

The MOCVD of Cu films from CpCuPEt3 was investigated to determine precursor characteristics and the interplay between growth parameters and film morphology. We show here that it is necessary to purify the compound in order to obtain the best stability and volatility For all the investigated conditions, the films present nodular, smooth morphology, which is characteristic of a Wolmer-Weber type island growth. From the SEM images, we observed a simultaneous increase in the size of crystallites and grains with increasing precursor flow rate, resulting in the coalescence of the grains and, consequently, in improved film continuity. We also observed a slight decrease of the grainsize with increasing deposition temperature. Image analysis of SEM top-down micrographs revealed that when the deposition temperature decreased from 493 K to 431 K, the film coverage shrank while the period of their morphology increased from 210 nm to 350 nm. The latter behavior meets our expectation that the decrease of deposition temperature limits the desorption of decomposition products, namely triethylphosphine, resulting in less dense morphologies. We observed a noticeable increase in period with an increas in theprecursor flow rate from 0.02 sccm to 0.04 sccm to 0.07 sccm, whereas the coverage remained almost similar with a slight increase from 0.78 to 0.84 to 0.86, respectively. This indicates that the channel-like morphologies became broader with much wider local features and uniformity, underpinning the dominance of coalescence effects over nucleation processes.

With a specially developed methodology of image treatment, we quantified the period (density) and coverage to assess the contribution of nucleation and coalescence, respectively, as well as the interplay of these two fundamental mechanisms. Subsequently, we addressed the link between structure (coverage) and function (electrical conductivity). From a more general perspective, this work strengthens the correlations among processing, structure, and function by means of precursor insight and advanced and targeted morphological characterization of structure.

Figure 1

 1 Figure 1 presents the TGA and DTA profiles of the sublimated compound. Weight loss occured

Fig. 1 :

 1 Fig. 1: TGA and DTA profiles of the sublimated CpCuPEt3 under N2 flow at atmospheric

Figure 2

 2 Figure 2 presents the experimentally determined values of the Psat vs. temperature for the

Fig. 3 .

 3 Fig. 3. Plot of the MOCVD growth rate divided by the estimated flow rate of CpCuPEt3

Table 3 . 4 .Fig. 4 .

 344 Fig. 4. Top-down view SEM images showing the morphology of the copper films grown on Si

Fig. 5 .

 5 Fig. 5. Higher magnification top-down view SEM images (a) Fig. 4a, (b) Fig. 4e, and (c) Fig.

Fig. 6 .

 6 Fig. 6. Histogram of the pixel intensities of a typical top-down SEM micrograph of the copper

Fig. 7 .

 7 Fig. 7. Coverage (left axis) and period (right axis) of Cu films, measured using the methodology

  up to 363 K. The Psat values obtained from the purified precursor obey the logPsat (Pa) = 4.368 -2272/T (K) relation; they are convenient to ensure precursor transport in MOCVD process. The obtained Cu films are pure. The Arrhenius plot of the Cu growth rate yields an activation energy of 10 kJ/mol, showing weak dependence on the deposition temperature. This observation, in conjunction with literature results, indicates a diffusion-controlled regime in these deposition conditions.

Table 1 .

 1 Investigated operating conditions.

	Total reactor pressure / Pa	1333
	Time / min	90 -420
	Deposition temperature / K	431 -523
	Sublimation vessel temperature / K	326 -373

Table 2 .

 2 Elemental molar ratios of P, C, and H to Cu for the as-received and purified CpCuPEt3.

	Molar ratio P/Cu C/Cu H/Cu
	As received 0.492 2.110 0.318
	Sublimed 0.491 2.089 0.313
	Stoichiometric, pure CpCuPEt3 0.487 2.079 0.317

The corresponding ratios for the stoichiometric, pure compound are given for comparison.

Acknowledgements

We are indebted to Y. Thebault and to P. de Parseval for FEG-SEM and EPMA analysis, respectively. This work was partly supported by CNRS and RFBR through grants PRC 1986/2018 and 18-53-15005, respectively.