Cécile Delacour

Luc Ortega

Marc Faucher

Thierry Crozes

Thierry Fournier

Bernard Pannetier

Vincent Bouchiat
email: bouchiat@grenoble.cnrs.fr

Persistence of

Introduction

Superconductivity has been recently shown to survive down to extremely confined nanostructures such as metal monolayers 1 or clusters composed of few organic molecules 2 .

While these structures are extremely interesting for probing the ultimate limits of nanoscale superconductivity, their studies are limited to in-situ measurements. Preserving a superconducting state in ultrathin films that can be processed by state-of-the-art nanofabrication techniques and withstand multiple cooling cycles remains technologically challenging and is timely for quantum devices applications. Having the highest critical temperature among elemental superconductors, niobium is an ubiquitous material for superconducting thin films and its performance is known to increase when epitaxy conditions can be reached. Hetero-epitaxy of niobium is best 3,4 performed on sapphire substrates since both their lattice parameters and thermal dilatation coefficients match rather well 3 . Most combined structural and transport studies have involved A-plane (1000) oriented sapphire 5 as the growing substrate. However the high step-edge roughness 6 and interfacial stress found for that crystal orientation is likely to be a probable source of rapid aging. Only few studies 4,6 have described Nb growth on R-plane oriented ( ) sapphire. However, this orientation provides an ideal substrate to promote epitaxy of Nb with (001) orientation. We present here a combined structural and electronic transport analysis of the early growth stages of niobium films on that orientation which lead to a better understanding of the mechanism of superconductivity depression in ultrathin Nb films and show the existence of an intermediate metallic behaviour.

Thin-film Growth and Structural Analysis

The niobium films are grown in ultra-high vacuum using electron gun evaporation technique with the following details : R-Plane oriented, single faced polished, 2 inches sapphire substrates with reduced miscut (<10 -3 rad) (from Le Rubis S.A., and Kyocera) are cleaned using water-based surfactants in ultrasonic bath. The surface is then prepared 7,8 by annealing during 1 hour in air at 1100°C. Wafers are transferred in the deposition chamber (from Isa-Riber, Ultra High Vacuum, base pressure ~10 -10 torr). Before niobium deposition, in-situ cleaning of the surface is performed during several minutes using Argon milling (Ion tech, Inc) using settings previously reported in ref. 9 . Then, the substrate is heated at 660°C and niobium (ingot target prepared from 99.99% purity pellets from Neyco) is deposited using electron gun at 0.2 to 0.5 nm/s evaporation rate. This substrate temperature has been shown 3 to optimize the deposition conditions as it provides a trade-off between atomic mobility of niobium adatoms on the surface and oxygen contamination migrating from the sapphire that starts to occur above that temperature 5 . The vacuum is kept at residual pressures lower than 10 -9 Torr during niobium deposition. Finally, the sample is cooled below 80°C and covered by evaporation of a capping layer of 2nm-thick silicon in order to protect the film from oxidation. The thickness d of the niobium layer film is varied from 200 nm down to 2 nm (crystallographic analysis was restricted to the lowest thickness range below 10nm samples A-G).

The sapphire surface annealing that is performed prior to niobium growth has been proven critical to improve the epitaxy 3,10 . Atomic Force Microscopy analysis performed prior to growth clearly shows vicinal surfaces with atomically flat terraces separated by straight steps edges 7 of 0.3 nm height (Fig. 1a). This measured step height is in agreement with the calculated 4 distance between two adjacent ( ) lattice planes (a rhom =0.35 nm) see Fig. 1d.

After Nb/Si deposition, the surface exhibits similar features replicated on top of the capping layer, which have same height, orientation and size than the bare alumina substrate vicinal surfaces (Fig. 1b). These features along with the low measured roughness indicate a good ordering of the niobium films.

On a larger scale, X-ray scattering measurements confirm an average vicinal surface width averaged over the entire wafer of about 300 nm (Fig. 1c).

In order to probe the inner structure of the multilayer, X-ray reflectivity (XRR) measurements have been performed (Fig. 2a). All samples exhibit a layered structure in agreement with Al 2 O 3 /Nb/Si simulation model (dotted curves) from which a precise determination of the multilayer structure can be deduced (Fig. 2b). The magnitude of the measured roughness at the top and bottom niobium interfaces are deduced from the fits and

show almost constant values over the whole deposition range. While the Nb/Sapphire roughness is compatible with an atomically sharp interface, a significantly rougher Nb-Si interface is extracted from the fits (Fig. 2b). However surface roughness measured by AFM (Fig. 1b) on the top of the silicon capping layer is repeatedly found below 0.1 nm. This suggests the presence of a diffuse and smooth Nb/Si interface rather than a sharp one, which might result from a slight inter-diffusion zone of about 2 monolayers (ML). X-ray diffraction (XRD) measurements reveal that films thicker than 4 nm (samples A to C) are single crystalline with a (001) niobium bulk orientation in agreement with the models (Fig. 1d). The (110) pole figure (Fig. 2c) shows the diffracted intensities of the 4 diagonal (110)-planes of the body center cubic lattice of niobium (Eulerian angle  = 45° +/-3°) as expected for a hetero-epitaxial growth of (001) on R-plane sapphire (Fig. 1c). The 3° angle mismatch corresponds to a previously reported 4 misalignment between the ( ) sapphire and (001) niobium planes. Interestingly, this mismatch is observed even for the thinnest singlecrystalline sample (4 nm/ equivalent to 12 ML). XRD patterns of samples thinner than 4nm (samples D, E, F) exhibit no trace of the body center cubic Nb lattice, but reveal only a hexagonal lattice corresponding to the sapphire substrate, suggesting that hetero-epitaxy is not achieved if the growth is interrupted in its early stages.

Further investigations by grazing incidence X-ray diffraction confirm that the thinnest films (samples D, E) are indeed made of mosaic polycrystals, as the patterns are similar to the ones encountered for niobium powder (Fig. 2d). The presence of polycrystals most probably explains the slight discrepancy observed for two samples between XRR measurements and the fit (Fig. 2a). Indeed the model involves three layers of single crystals.

All XRR measurements agree with a structural transition from polycrystalline for the thinnest films to a single crystalline growth occurring for films thicker than 3.3 nm (10 ML).

The initial growth stage of niobium films on an atomically flat substrate follows the step-flow model 11 as previously predicted 4 . Niobium adatoms move on the terraces, rearrange themselves along the sapphire step-edges and form patches which eventually percolate, leading to a mosaic polycristal film. When the layer gets thicker, typically over 10 MLs, volume interactions become dominant with respect to step-edge interaction and force the film to be single-crystalline with a (100) niobium bulk orientation.

Electronic Transport Measurements

The sheet resistance of samples have been measured from room temperature down to 0.1 K on 5 micron-wide strip-lines over a length of typically 100 squares. Sheet resistance measured as a function of the temperature (Fig. 3a) shows superconducting transitions for all films thicker than 3 nm.

As expected, the critical temperature T c (defined by default as the temperature at which the film has lost half of its resistance) is found to decrease simultaneously to the film thickness d.

The superconducting transition widens notably for the polycrystalline films (samples D, E).

This transition widening mostly occurs at the top of the transition as a result of an increase of the superconducting fluctuations. However there is still a rather sharp drop of several orders of magnitude of the film resistance at T c (fig 3a, right) and we found no trace of residual resistance below T c that would have revealed the existence of phase fluctuations. As for the thinnest films (samples F,G), one does not observe any superconducting state at the lowest temperatures (0.1K) reached in our filtered measurement setup. This suggests that the superconductivity is suppressed for a critical thickness d c ranging between 2.8 and 3 nanometers.

For sample F however, high resolution measurements have shown a drop at 0.5K towards a partially resistive state which equals 70% of the resistance measured at 5K, (see Fig. 3c). This suggests that the sample experiences a transition towards an intermediate metallic state.

Interestingly, when one crosses the critical thickness d c , one can simultaneously observe a change of sign of the temperature coefficient of resistance (dR/dT) measured below 70K as shown in figure 4. In samples F and G which lie below d c , the conductance decreases upon cooling. For sample G, this decrease clearly follows a logarithmic dependence with temperature (Fig. 4) , which is a signature of a 2D weak localization.

It is not clear, however, for sample G, whether this transition is towards a metallic state or an insulating one, as both systems are known to exhibit negative temperature coefficient 12 .

The magnetic field dependence of the R(T) curve has been measured as well for the 3nm film (Fig. 3b, sample D) and similarly shows a rapid drop of the superconducting transition for fields above 3T together with a slight increase of the resistance upon cooling. The same curve measured for sample F (Fig. 3c) shows a very rapid suppression of the partial drop at 0.5K under a small transverse magnetic fields and a similar change of sign of the temperature coefficient for higher fields.

The purity of Nb films and its thickness-dependence can be estimated by measuring the residual resistivity ratio (RRR) which is obtained by dividing the room temperature resistance by the resistance R n taken just above the transition. It is known that the quantity k=(RRR-1) -1 is proportional to the density of defects [START_REF] Kittel | Introduction to Solid State Physics[END_REF] . The RRR extracted from R(T) curves show a linear dependence on the thickness d (Fig. 5a) and the quantity k diverges for the thinnest samples (Fig. 5b), revealing a high concentration of defects within the top Nb monolayers.

One can also deduce from the resistance-temperature measurement the mean free path knowing that its product with the bulk resistivity ( ) is equal [START_REF] Ashcroft | Solid State Physics[END_REF] to 3.75 10 .cm 2 . One finds 3.5 nm for the 5 nm film. One can then evaluate the superconducting coherence length in the dirty limit [START_REF] Ashcroft | Solid State Physics[END_REF] : , with Nb-bulk coherence length  0 =38 nm. One obtains =9.8 nm for the 5-nm-thick sample. The superconducting coherence length equals twice the film thickness d and clearly places our study within the two dimensional limit.

A more accurate determination of the superconducting coherence length () can be extracted [START_REF] Tinkham | Introduction to superconductivity[END_REF] by measuring the temperature dependence of the critical magnetic field  c2 (T) (Fig. 6) at small applied fields .

For that purpose, we have measured R(T) curves for increasing magnetic fields. A precise determination of the transition temperature T c is obtained by fitting the initial drop of resistance above with a model describing the superconducting fluctuations defined by the quantity ( , measured in presence of a weak magnetic field. Superconducting fluctuations have been calculated by Ullah and Dorsey [START_REF] Ullah | [END_REF] in the limit of 2D thin films 17 . Above the critical temperature, the fluctuation term follows a linear function of which allows to calculate by extrapolating the fitting line. By inverting the set of T c obtained by these fits, one can then plot the H c2 (T) curve in the low field limit (inset of Fig. 6). Since the Ginzburg-Landau theory gives : H c2 (T)= 1 , a linear fit of this curve leads to a determination of the coherence length =10.6 nm, which is in good agreement with the first estimate.

 0 852 . 0   

Physical origins of the depression of the superconductivity

The suppression of superconductivity has been studied before in several configurations, for granular 18 or amorphous films 19 and with different elemental materials 19 and alloys 20 . Its origins could be multiple as they range from transition to a localized state due to disorder [21][22][23] to quantum confinement 18,24 or to inverse proximity effect with the surrounding materials 25,26 . In a disorder induced superconducting to insulator transition 27 , R s was believed to be of the order of the resistance quantum h/(4e²)=6.5 kIn our system however the transition is observed for a well lower resistance per square (Fig. 3a) suggesting that the localization is not the dominant factor that leads to the suppression of the superconducting state.

Finkelstein's model

We analyze our results in the framework of a weak disorder model in homogeneous superconducting films established by Finkel'stein 21 . The suppression of superconductivity is driven by impurities that reinforce Coulomb and spin interactions. T c is expressed as a function of sheet resistance R S and the elastic diffusion time  :

t c c t t t t e T T 2 / 1 0 4 / 2 / / 1 4 / 2 / / 1                 (1)
with , and = 9.22 K is the Nb-bulk critical temperature. The transition occurs for a non-universal sheet resistance such as t=². Fig. 7 compares the Finkel'stein model (Eq. 1) with our experimental data for several values of the parameter.

The model seems at first sight rather consistent with our data measured above d c and leads to a fitting parameter = 9.5. However, the model is in complete disagreement with our data obtained below the critical thickness d c , as it would imply a much smoother superconducting

h e R t S  ²             0 ln c B T k h 0 c
T depression. Furthermore the model implies a critical sheet resistance R c = 900 at the transition, which is twice the experimental one measured at d c (Fig. 3a). Furthermore, in order to get a critical resistance of 500  the elastic relaxation time  one has to inject in Eq. ( 1) would have to be less than 1.5.10 -17 s which gives v F nm a value in disagreement with the mean free path evaluated above.

Inverse proximity effect induced by the capping layer.

The formation of metallic alloys and the interdiffusion of silicon or oxygen atoms at the interface is likely to be involved in the disappearance of the superconductivity by an inverse proximity effect for the thinnest films, and their influence is probably aggravated by the fact that the film cannot be grown as a single crystal for thicknesses below 4 nm. According to the simplified McMillan model 28 , we consider an inverse proximity effect in a planar thin film geometry. The critical temperature c T is depressed as a function of the normal layer thickness d N according to the relation :

d c D c c T T T T               0 0 56 . 3 (2) with ) 0 ( ) 0 ( S N N N N d    N N,
S (0) are respectively the density of states in normal (N) and superconducting (S) layers and T D =277 K is the Nb Debye temperature. We have compared our measurements with (Eq.

2) as a function of  (Fig. 8) and find the best agreement to be obtained for d N =0.55 +/-0.05 nm, considering a N N (0)/N S (0) ratio equal to 1. This value is consistent with the thickness of the Nb/Si interdiffusion layer estimated from the XRD data (Table in Fig. 2b Si-interface thickness), which could create a metallic silicide at the Nb/Si interface. Oxygen traces migrating through the silicon capping layer could also further contribute to the inverse proximity effect since niobium oxides with low stoichiometry are known to be metallic and even suspected to be a potential source of destruction of superconductivity by magnetic effects 29 .

In addition, when compared with previous studies 30 , this simple model is well consistent with our data above d c . However the drastic fall of T c for thinnest resistive samples (red stars) is not explained even by taking more realistic models 31,32 . The reason why these models cannot describe the full set of data is likely to be associated with the structural transition in the film, as the physical origin of the superconducting depression most probably varies according to the crystalline structure of the films. Additional models are therefore needed to describe the situation for the thinnest films.

Proximity effect between 2D grains within the layer.

One has shown in the growth section that the film morphology differs considerably below the critical thickness d c . Mesoscopic fluctuations of the metal properties can create huge inhomogeneities leading to normal and superconducting regions close to each other. The appropriate model to describe the electronic structure is therefore not anymore the McMillan model but rather models involving Normal/Superconducting 2D networks. Such a model implies a much more drastic effect on the superconducting depression than the one found using the McMillan model.

In the case of a transparent N/S interface one finds the Cooper limit ∝ . Spin flip interactions would even have a more detrimental effect on T c . This latter model would then involve a transition from superconductivity towards a "bad" metal which square resistance indeed does not require a square resistance close to h/4e 2 to be triggered.

Below 4 nm, films have been found to be polycrystalline (Fig 2d). Macroscopic superconducting state can be ensured by proximity effect between Nb-grains. Such 2D NS granular network systems have been investigated by Feigel'man et al. 33 but this model concerns the limit of diluted superconducting grains, which is unlikely to apply here. The approach of Spivak et al. 34 needs to be investigated in further details as it deals with a situation closer to our system. It indeed focuses on superconducting grain size (R<) and film thickness (d<) lower than the superconducting coherence length. In both models, it is worth noting that the critical resistance at the transition may appear at arbitrarily small values as it is experimentally observed here.

Applications to the realization of Quantum Devices

The reported single crystalline niobium films are stable in time over months even exposed to the ambient atmosphere, as demonstrated by the 4-nm-thick film having a critical temperature T c = 5 K that is still preserved after multiples thermal cycles and 3 years of shelf storage. The thinnest samples are also resistant to micro and nano-fabrication processes, with a superconducting transition temperature preserved above 4.2K for nanostructures made from sample C.

Elaborating time-and air-stable ultrathin superconducting films is useful for quantum device fabrication. An ultrathin film improves the sensitivity of single particle detectors (either based on bolometric, kinetic inductance, or hot spot diffusion effects). These quantum detectors are in high demand for astrophysics particle detectors or for quantum cryptography applications [START_REF] Natarajan | Quantum Commmunication and Quantum Networking[END_REF] . These films have already proven to be useful for fabricating more exploratory devices such as single electron transistors that have shown room temperature operation 36 or nanoscale magnetometers 37 allowing for increased detection sensitivity of nanomagnet magnetization 38 . Growing high quality niobium (Nb) films has also shown to increase the stability and supercurrent density of Nb-based tunnel junctions 39 . When applied to superconducting single-photon detection, ultrathin Nb-based devices have provided faster repetition times 40 than niobium nitride counterparts due to their lower kinetic inductance.

Finally the resilience of such films upon the application of in-plane magnetic field offers promising perspectives of applications including the realization of on-chip superconducting LC resonators that remain superconducting upon intense magnetic fields. This would help for the manipulation of spin qubits coupled to readout superconducting cavities (for a review see ref. 41 ).

In order to demonstrate that it is possible to produce stable nanostructures following our technique, the 3-nm-thick (9 ML) Nb-thick superconducting film has been patterned using conventional Deep-UV photolithography followed by reactive ion etching (using SF 6 ).

Critical temperature measured on micron wide lines is barely affected by the process and local anodisation induced by Atomic Force microscope based nanolithography lead to the thinnest nanometer scaled integrated on-chip nanodevices ever made so-far 36 , 37 .

Conclusion

We have presented a combined analysis of structural and electron transport properties of sub-10 nm niobium films evaporated in ultra-high vacuum on R-plane sapphire. Nb films undergo a sharp drop of the superconducting transition at 3 nm (9 ML) accompanying a structural transition from poly to single crystal occurring at a thickness slightly above that value. The Tc cut-off occurs over a narrow thickness range (1ML) obtained for a non-universal critical resistance which is more than 10 times below the quantum of resistance h/(4e 2 ). Resistance temperature measurements show a transition toward a metallic state with 2D weak localization features. Above that critical thickness, the overall stability demonstrated by both structural and electronic measurements of the ultrathin films offers a reliable starting material for the realization of quantum superconducting devices.

Acknowledgements.

The c Normalized resistance of sample F (d = 2.8 nm), measured at very low temperature. From bottom to top, the applied transversal magnetic field is respectively 0, 0.35T, 6.5 T.

Figure 4

Temperature dependence of the sheet conductance of the thinnest samples (E: d=2.98 nm; F: d=2.8nm; G: d=2 nm) shown in a semilog plot. The conductance of sample G shows a logarithmic decrease upon cooling while for sample F a partial transition occurs at T~0.5 .

Figure 1 a,

 1 Figure 1

Figure 2 `

 2 Figure 2 `

Figure 3 ab

 3 Figure 3

Figure 5 a

 5 Figure 5 a. Residual resistivity ratio (RRR) as a function of Nb films thickness (black

Figure 6

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Figure 7

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Figure 8 Figure 1

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  authors are deeply indebted to G. Battuz and A. Hadj-Azzem for substrate preparation. One of us (C.D.) thanks the D.G.A. for grant support.

	Correspondence	and	requests	for	materials	should	be	addressed	to
	V.B.(vincent.bouchiat@grenoble.cnrs.fr)


Gérard Demazeau
email: demazeau@icmcb-bordeaux.cnrs.fr

Solvothermal and hydrothermal processes: the main physico-chemical factors involved and new trends

Keywords: Hydro/solvothermal processes, History of hydrothermal processes, Key factors, New trends

Hydro-and solvothermal processes play an important role in the different scientific domains involving either basic or applied research. During the last years, such processes have been strongly developed, in particular with the elaboration of nanocrystallites. This review article presents a brief history of their development, an analysis of the key factors governing such processes, and the recent trends.

Introduction

Due to the role of water in most mechanisms involved in different natural systems (geology, chemistry, biology, etc.), the first part of this article will be devoted to a brief history of hydrothermal processes.

In the second part, an analysis of the key factors governing such hydro/ solvothermal processes will be conducted.

Taking into account the strong development of such processes during the last 20 years, the objective of the third part will be to summarize the recent trends.

The conclusion will propose some trails to improve the development of hydro/ solvothermal processes in the near future either for a better understanding of the mechanisms governing these reactions or for the adaptation of these processes to major problems in different scientific domains.

A brief history of hydrothermal processes

Hydrothermal processes: the overlap between pressure as thermodynamical parameter and the specific properties of water After the definition of pressure during the 17th century as a thermodynamical parameter, in particular by Pascal in the paper ''Traite´de l'e´quilibre des liqueurs'' published in 1663 [START_REF] Richet | Une courte histoire de la pression[END_REF], different research areas involving this parameter were developed: (i) in physics, with the studies involving the compressibility of gases and liquids, conducted by Andrews [2], Amagat [3], Cailletet and Mathias [4], (ii) in geology, with the research focusing on the synthesis of natural minerals, in particular the experiments of Hall at the end of the 18th century [5], (iii) in hydrometallurgy [START_REF] Habashi | The origins of hydrothermal reactions[END_REF] with the elaboration processes of different metals by Beketoff [START_REF] Mambote | Hydrothermal metallurgy: an overview of basic concepts and applications[END_REF] and Ipatiev [8], and then (iv) in biology at the end of the 19th century, with the discovery of living organisms in the deep sea with Regnard [START_REF] Regnard | Recherches expe´rimentales sur l'influence des tre's hautes pressions sur les organismes vivants[END_REF] and Certes [START_REF] Certes | De l'action des hautes pressions sur les phe´nome'nes de putre´faction et sur la vitalite´des micro-organismes d'eau douce et d'eau de mer[END_REF] and then the first pressure experiments dealing with the inactivation of bacteria conducted by Roger [START_REF] Roger | Action des Hautes pressions sur quelques bacte´ries[END_REF] and Hite [START_REF] Hite | The effect of pressure in the preservation of milk (West Virginia Agricultural Experiment Station[END_REF].

All these scientific domains took into account the impact of water under high pressure. Therefore, the development of hydrothermal processes resulted from the overlap between pressure, the specific properties of water in such conditions, and chemistry.

The so-called ''hydrothermal processes'' can be defined as a chemical reaction in a closed system involving one or different precursors in the presence of water at a temperature higher than its boiling temperature. Due to these conditions, pressure is involved. Consequently, the first research in geochemistry was entitled ''Minerals in hot water'' [5].

The first developments of hydrothermal processes in geochemistry

Hall, taking into account Hutton's proposition that pressure was the solution to retain CO 2 during the fusion of minerals, succeeded in transforming chalk to marble. Such experiments led to the publication in 1805 of the paper: ''Account of a series of experiments showing the effects of compression in modifying the action of heat'' (Hall 1805) [5].

Hall's experiments inspired a group of French mining engineers and geochemists: de Senarmont (1808-1862), Daubree (1814-1896), and Friedel (1832-1899) [5].

The first high-pressure equipment was developed in glass ampoules filled to 50% of their volume with water and submerged in brass gun barrels also containing water as the pressure-transmitting medium. The barrels were placed at the top of the gas furnace of a foundry (200-300 ꢀC) and covered with coal dust [13]. Due to the experimental difficulties for the management of these experiments, the gun barrels were rapidly replaced by steel bombs as reaction vessels [START_REF] Daubree | Etudes Synthe´tiques de Ge´ologie Expe´rimentale DUNOD Ed PARIS[END_REF]. Then, Friedel and Sarasin (1879) initiated extensive hydrothermal studies in particular with the synthesis of orthoclase from potassium silicate and AlCl 3 [START_REF] Friedel | [END_REF].

During approximately the same period, Schafha¨utl (1803-1890), with the preparation of quartz microcrystals through the transformation of freshly precipitated silicic acid in a Papin's digestor [16], Bunsen (1811-1899) with the use of thick-walled glass tubes allowing to reach 100-150 atmospheres (10-15 MPa) [17,18], Von Chrustoff [19], and Doelter (1890) [START_REF] Doelter | Allgemeine Chemische Mineralogie[END_REF] were involved in hydrothermal geochemistry.

All these scientific works can be considered as the early developments of the hydrothermal processes and the basis for all research conducted from the end of the 19th century to 1950.

The diffusion of hydrothermal processes in different scientific domains

From the end of the 19th century to the middle of the 20th century, different scientists developed hydrothermal reactions mainly in the crystal growth of a-quartz (Spezia [START_REF] Spezia | [END_REF][22][23], Nacken [24,25], Walker [START_REF] Walker | Hydrothermal synthesis of quartz crystal[END_REF], Wooster and Wooster [START_REF] Wooster | Preparation of synthetic quartz[END_REF], Brown et al. [28]. Hale [START_REF] Hale | The laboratory growing of quartz[END_REF], Buehler and Walker [START_REF] Buehler | Growing quartz crystals[END_REF], Shaftal and Butszov in [START_REF] Iwasaki | Historical review of quartz crystal growth[END_REF], Taki and Kunitomi [32], see also the review of Moriya and Ogawa [START_REF] Moriya | Growth history of a synthetic quartz crystal[END_REF]) and in the systems silicates-H 2 O (Katsurai [34], Nagai [35], Yamasaki and Titani [START_REF] Yamasaki | On the vapour-phase hydrolysis[END_REF], Niggli and Morey [37], Straub [38], Morey and Ingerson [START_REF] Morey | The pneumatolytic and hydrothermal alteration and synthesis of silicates[END_REF], Goranson [START_REF] Goranson | Silicate-water systems: phase equilibria in the NaAlSi O -H O and KAlSi O -3 8 2 3 8 H 2 O systems at high temperatures and pressures[END_REF], Barrer [41], Kiyoura [START_REF] Kiyoura | Hydrothermal decomposition of glass[END_REF], Bowen and Tuttle [START_REF] Bowen | The system NaAlSi O -KAlSi O -H O[END_REF], , Kennedy [START_REF] Kennedy | Pressure-volume-temperature relations in water at elevated temperatures and pressures[END_REF], Otsubo and Kato [48], Kunitomi and Saito [49], and Roy and Osborn [START_REF] Roy | Some simple aids in the hydrothermal investigation of mineral systems[END_REF]).

All of these research activities were driven by the development of hydrothermal equipment. The development of sealed glass ampoules as ''hydrothermal reactors'' was improved in particular by Allen et al. [START_REF] Allen | The mineral sulphides of iron, experiments on the formation of the iron sulphides[END_REF]. Different apparatuses for the hydrothermal synthesis and growth of single crystals were set-up, in particular by Morey [START_REF] Morey | Hydrothermal synthesis[END_REF], Smith and Adams [START_REF] Smith | The Role of Hydrothermal Synthesis in Preparative Chemistry[END_REF], Tuttle [START_REF] Tuttle | A new hydrothermal quenching apparatus[END_REF], Litvin and Tules [START_REF] Litvin | Apparatus for Hydrothermal Synthesis and Growth of Single Crystals[END_REF], and Butusov and Bryatov [START_REF] Butusov | [END_REF].

During the second part of the 20th century, different research groups were involved in the development of hydrothermal processes either in materials chemistry (synthesis of new materials) or in materials science (crystal growtheven at the industrial scale, elaboration of small crystallites-from micro to nanoscale, the deposition of thin films), particularly by: Somiya and Roy [START_REF] Somiya | Hydrothermal synthesis of fine powders[END_REF], Roy and Tuttle [START_REF] Roy | Investigations under hydrothermal conditions[END_REF], Roy and Osborn [START_REF] Roy | Some simple aids in the hydrothermal investigation of mineral systems[END_REF], Fyfe [START_REF] Fyfe | Hydrothermal synthesis and determination of equilibrium between minerals in the subsolidus region[END_REF], Yamasaki et al. [START_REF] Yamasaki | A hydrothermal hot-pressing method: apparatus and application[END_REF][START_REF] Yamasaki | Hydrothermal dynamics on environmental problems using the aspects of earth science[END_REF], Koizumi and coworkers [START_REF] Endo | Synthesis of potassium manganese oxides under hydrothermal conditions[END_REF], Laudise and Nielsen [START_REF] Laudise | Hydrothermal crystal growth[END_REF], Somiya [START_REF] Somiya | Proceedings of the First International Symposium on Hydrothermal Reactions[END_REF][START_REF] Somiya | Historical developments of hydrothermal works in Japan, especially in ceramic science[END_REF], Yoshimura et al. [START_REF] Yoshimura | Preparation of alpha alumina fine powders by hydrothermal oxidation method[END_REF][START_REF] Yoshimura | Hydrothermal conversion of calcite crystals to hydroxyapatite[END_REF], Demianets [START_REF] Demianets | Hydrothermal synthesis of new compounds[END_REF], Demianets and Lobachev [START_REF] Demianets | Hydrothermal synthesis of crystals[END_REF], Eitel [START_REF]Silicate Science, Hydrothermal Silicate Systems[END_REF], Rabenau [START_REF] Rabenau | Methods for the study of hydrothermal crystallization[END_REF], Joubert and coworkers [START_REF] Ferrand | Growth of epitaxial substituted garnet films by hydrothermal synthesis[END_REF], Demazeau and coworkers [START_REF] Rault | Sur deux nouveaux hydrofluorures de chrome et de gallium[END_REF][START_REF] Demazeau | Sur un nouveau mate´riau pour l'enregistrement magne´tique de´rive´du dioxyde de chrome[END_REF], Whittingham and coworkers [START_REF] Reis | Low temperature hydrothermal reduction of ammonium paratungstate[END_REF], Komarneni and coworkers [START_REF] Komarneni | Hydrothermal preparation of ultrafine ferrites and their sintering[END_REF][START_REF] Kumada | Microwave hydrothermal synthesis of ABi 2 O (A = Mg, Zn)[END_REF], and Riman et al. [START_REF] Riman | Hydrothermal crystallization of ceramics[END_REF].

At the beginning of the 1990s, after the oxides ''era'' and the beginning of nonoxide ceramics and in particular the synthesis of nitrides, for example, non-aqueous solvents were developed. Therefore different appellations were used for the corresponding reactions using these solvents in a closed system such as ammonothermal [START_REF] Purdy | Ammonothermal crystal growth of sulphide materials[END_REF], glycothermal [START_REF] Inoue | Glycothermal synthesis of rare earth aluminum garnets[END_REF], alcohol-thermal [START_REF] You | Alcohol-thermal synthesis of flowerlike hollow cobalt tungstate nanostructures[END_REF]. Demazeau and coworkers [START_REF] Demazeau | Particular compositions of rare earth oxides their preparation and their use[END_REF][START_REF] Th | Solvothermal synthesis of Fe O microcrystallites for catalytic appli-3 4 cations[END_REF] proposed the word ''solvothermal'' as a generic appellation including all solvents.

Over the last 10 years, different trends have arisen from all the previous investigations, including: (i) the strong improvement of solvothermal processes in different scientific domains, in particular nano-systems, and the very prolific activity of several research groups in China, and (ii) the optimization of these processes thanks to the analysis and control of the key factors governing each step.

Key factors governing hydro/solvothermal processes

The role of the solvent in solvothermal process

The role of the solvent in the solvothermal synthesis involving oxides and sulfides in particular has been the subject of several investigations [START_REF] Biswas | Effect of the precursors and solvents on the size shape and crystal structure of manganese sulphide crystals in solvothermal synthesis[END_REF].

The solvent is involved through: (i) its physical properties (dielectric constant, polarity, density, etc.), (ii) its chemical properties (mainly the solvation and the ability or not to stabilize complexes), and (iii) its interactions with the reactants and the additives.

Role of the solvent in materials chemistry

In materials chemistry, the selection of the solvent mainly depends on the purpose of the experiment: for example, to help the formation of a specific material through the control of either the solvation of the chemical reactive species or the concentration of the reactants using in particular non-miscible mixed solvents.

In the case of the solvation, the synthesis of KTaO 3 with the perovskite structure in non-strong alkaline solutions can be chosen as an example. In hydrothermal processes using Ta O and KOH as reactants, the hydroxyl ion activity is a critical 2 5 parameter for the formation of the perovskite KTaO 3 phase. Consequently, the hydrothermal synthesis of the ferroelectric perovskite KTaO 3 requires a high concentration of KOH (=15 mol/l) [START_REF] Goh | Hydrothermal synthesis of perovskite and pyrochlore powders of potassium tantalate[END_REF].

The solvent effect can play an important role in acid and base activity through the solvation phenomenon [START_REF] Epshtein | Present ideas on the influence of the medium on acid-base equilibria, hydrogen bonds in the gas phase and in solution[END_REF][START_REF] Dyumaev | The influence of solvation on acid-base properties in various media[END_REF].

In water, the solvation of OH groups, through the formation of hydrogen bonds, leads to a partial shielding of hydroxyl group. Consequently, hydroxyl ion activity can be decreased.

In ethanol, the solvation effect is weaker than in water and then induces a higher OH activity. Therefore the chemical reactivity of KOH in ethanol would be improved. He et al. [START_REF] He | Mixed solvents: a key in solvothermal synthesis of KTaO 3[END_REF] have observed that the hydrothermal synthesis of the KTaO 3 perovskite in pure ethanol required only a concentration of KOH close to 1 mol/l (far below the 15 mol/l used in water).

In the case of non-miscible solvents, one objective is to govern the concentration of the reactants or one specific reactant in the droplets (used as tiny reactors) of one solvent, the size of such droplets being governed by the volume ratio of solvents. This approach has also been used for the preparation of KTaO 3 with the perovskite structure.

The reactants KOH and Ta O , being only soluble in water when a mixed-

2 5
solvent, H 2 O/hexane is used, the KOH concentration in water droplets (as tiny reactors) can be adjusted by the ratio between the volumes of water and hexane [START_REF] He | Mixed solvents: a key in solvothermal synthesis of KTaO 3[END_REF].

Another example is to orient a specific structural form for the final product. For this application, the control of the kinetics governing the formation of the final product is important. Using the same precursors, MnCl , 4H O, and SC(NH ) , several solvents (water, ethylenediamine, ammonia, benzene, THF) were evaluated for their ability to induce a specific structural form for MnS. The stable a form with the rocksalt structure is observed with solvents able to stabilize an intermediate complex [Mn(H 2 O) 2 6 ? , Mn(en) 3 2? , Mn(NH 3 ) 2 6 ? ]. On the contrary, metastable structures are observed when the Mn 2? complexation is not possible: the c form (wurtzite structure) with THF and the b form (zinc-blend structure) with benzene.

The stable a rocksalt structure results from the lower kinetics induced by the intermediate complex. The difference between benzene and THF has been attributed to the difference of polarity between these solvents, such a difference being able to modify the assembly of atoms [START_REF] Lu | Metastable MnS crystallites through solvothermal synthesis[END_REF].

Role of the solvent in materials science

In materials science, three main domains have been investigated: (i) the elaboration of nanoparticles with the control of their size and morphology, (ii) the solvothermal growth of single crystals, and (iii) the elaboration of nanocomposites.

The control of the size and morphology of fine particles Two main routes have been developed using solvothermal reactions:

-The control of the nucleation step using the physical properties of the solvent in particular (for example its critical temperature value). Hakuta et al. [START_REF] Hakuta | Production of phosphor (YAG#Tb) fine particles by hydrothermal synthesis in supercritical water[END_REF][START_REF] Hakuta | Continuous production of phosphor YAG:Tb nanoparticles by hydrothermal synthesis in supercritical water[END_REF] reported the preparation of YAG powder doped with Tb by a hydrothermal process (water being used as solvent) at 400 ꢀC and 30 MPa. Recently Li et al. [START_REF] Li | Solvothermal synthesis and luminescent properties of YAG:Tb nano-sized powders[END_REF] described the preparation of the same Tb-doped YAG phosphor particles using a mixed solvent (ethanol/water). The experimental conditions are slightly reduced (300 ꢀC, 10 MPa) compared to the ones used in pure water. Such a difference has been attributed to the lower supercritical conditions corresponding to a mixed solvent: ethanol-water 2:1 (267 ꢀC, 7.5 MPa) compared to (374 ꢀC, 22 MPa) for pure water. The solubility of metal oxides and hydroxides decrease drastically above the critical point and consequently the nucleation of YAG crystallites takes place at a lower temperature in such a mixed solvent. This example illustrates the role of critical conditions on the solvothermal preparation of small crystallites.

The control of the crystal growth step -Two different aspects have been particularly explored.

-The formation of an intermediate compound used as template to orient the morphology of the final material.

For the solvothermal synthesis of one-dimensional metal selenide crystallites, using metal chlorides (or nitrates) and selenium powder as reactants and hydrazine hydrate as the solvent, the hypothesis that the ''NH NH '' 2 2 group could act as bidentate ligand between two different M n? cationic clusters has been postulated. Such intermediate chemical species could then induce the one-dimensional morphology of the resulting particles [START_REF] Liu | Hydrazide route to one-dimensional structural metal selenides crystals[END_REF].

The selection of the solvent with the appropriate physico-chemical properties.

-Indium-titanium oxide (ITO) has been prepared through a solvothermal process using hydrated indium nitrate [In(NO ) ÁnH O], tin (IV) chloride H(OC H ) OH] improved the crystal growth rate compared to the nucleation 4 n 2 rate and, in parallel, induced the formation of oxygen vacancies leading to more free electrons and higher conductivity [START_REF] Lee | Solvent effect on synthesis of indium-tin oxide nano-powders by a solvothermal process[END_REF]. The role of the viscosity of the solvent has also been underlined for the morphological control of rare earth oxide nanoparticles by solvothermal reaction [START_REF] Yin | Synthesis and morphological control of rare earth oxide nanoparticles by solvothermal reaction[END_REF].

The solvothermal growth of single crystals The hydro-and solvothermal growth of bulk single crystals was an important step in the development of such processes.

During the last 10 years, the solvothermal crystal growth of sulfide materials has been an important challenge due to the potential developments of this class of materials for different applications [START_REF] Plass | Quantum dot sensitization of organic-inorganic hybrid solar cells[END_REF][START_REF] Poelman | Advances in sulphide phosphors for displays and lighting[END_REF][START_REF] Chen | Nanoparticle fluorescence based technology for biological applications[END_REF].

Manganese sulphide (MnS) crystallites were prepared from MnCl , 4H O, 200 ꢀC during 10 h. The c-form of MnS was observed. The presence of the mineralizer KNO 3 improved the crystallization of the manganese sulphide. Substituting KOH for KNO 3 as the mineralizer, the a-MnS form was stabilized. The solvent CS 2 , associated with such a mineralizer, appeared to modify the structural form, crystals morphology, and to improve the crystallinity [START_REF] Zhu | Mineralizer-assisted solvothermal synthesis of manganese sulphide crystallites[END_REF].

The elaboration of nanocomposite systems Nanocomposite systems can be defined as the uniting of different nano-components.

A novel strategy has been recently proposed to synthesize metal-oxide nanotubes through the homogeneous coating of carbon nanotubes (CNTs) through a solvothermal process. For example, using Ce(NO ) Á6H O as the precursor, . The formation of a uniform CeO 2 deposit on CNTs was only observed when pyridine was chosen as the solvent. This phenomenon may be related to N present in pyridine that could be able to modify the surface of CNTs in the earliest stage in order to facilitate the CeO 2 deposition.

Some bionanocomposites were also prepared with layered substrates such as mica or different inorganic matrix with intercalation of biomolecules in the interlamellar spaces [101].

In addition, the physico-chemical properties of the solvent can be modified through (a) the use of additives, and (b) the values of the thermodynamical parameters of temperature and pressure (with temperature, the physico-chemical properties are strongly modified depending on whether it is a subcritical or supercritical domain).

Impact of additives in solvothermal processes

Additives are generally developed to modify either the properties of the solvent or one step of the process. For example, they are important for modifying:

(i) the solubility of the reactants with the use of mineralizers; (ii) the oxidation-reduction properties; (iii) the crystal growth using either capping agents, surfactants, or biomolecules to orient growth to obtain an anisotropic morphology (for example, ethylene glycol cal also play the role of capping agent for preparing hierarchical FeWO 4 microcrystals [102]).

Role of the thermodynamical parameters of temperature and pressure

The main impact of temperature in solvothermal processes

The main impact of temperature in solvothermal processes is to differentiate two different domains where the physico-chemical properties of the solvent are modified: the subcritical and the supercritical domains [103].

Temperature is a key parameter of all solvothermal processes and can be adjusted to:

(i) stabilize a specific structural form of the final material, such as in the case of the solvothermal synthesis of hybrid vanadium oxyfluorides (iv) and the elaboration of nano-composites [111].

The selection of the reactants

The choice of the reactants can affect the final product because of their different characteristics, such as:

• • •
the chemical composition of one reactant used as a specific chemical source; the structure of the reactant to orient one of the final product; and the general composition of the reaction medium (ratio of the reactants).

The role of chemical composition

Using thiosemicarbazide as the starting sulphur source, CdCl 2 as the cadmium precursor and ethylenediamine as the solvent, CdS nanowhiskers were obtained. Such elongated nano-crystallites would be generated through the attack of the thiosemicarbazide molecules by the strong nucleophilic nitrogen of ethylenediamine leading to the weakening of the C=S double bonds. The rupture of C=S being slow, the reaction rate is also slow and contributes to the lengthening of CdS crystallites [112].

The selection of Fe-oleate precursors can induce the formation of seednanocrystals able to improve the crystal growth of iron oxide nanoparticles [113].

The structural form of the reactant

In most cases, solvothermal processes involve the dissolution of the reactants. If one of the reactants remains in the solid state under solvothermal conditions (in particular for low-temperature synthesis), it becomes possible to initiate the structure of the final product through one of the insoluble reactants. Such a solvothermal route was called ''structure-inheriting solid-state reaction'' by Whittingham and coworkers [114]. The effect of the reaction time during the synthesis of CoMoO 4 through a solvothermal process (using water as solvent) was investigated at 453 K using MoO 3 and CoO as precursors.

The composition of the reaction medium

Under ambient pressure and temperature conditions, NaYF 4 is characterized by two structural forms: the cubic (fluorite type) and the hexagonal (Na 1.5 Nd 1.5 F 6 type) structures, depending on the preparation method. NaYF 4 has attracted great interest due to its potential use as host lattice for up-conversion phosphors [115,116].

Y 2 O 3 and NaF were used as the precursors and EDTA (ethylenediamine tetraacetic) as the complexing agent and water as the solvent.

Depending on the Y 3? /F -ratio, the cubic form was stabilized with a stoichiometric ratio and the hexagonal one for an excess of F -(Y 3? /F -= 1/7.5). The F -concentration played an important role on the stabilization of the structural form, the EDTA/Y 3? ratio governing mainly the morphology. In such a synthesis, F -acts both as the reactant and as the mineralizer [117].

The associated technologies

The associated technologies (electrochemistry, microwaves, mechanical mixing, ultrasounds, external magnetic field, etc.) are able to improve the kinetics of solvothermal reactions of one or two orders of magnitude [118] or to induce a specific morphology of magnetic crystallites [119]. These associated technologies have been developed in particular for the preparation of well-defined nanocrystallites.

The control of different reaction mechanisms

Over the last 15 years, different reaction mechanisms have been investigated, such as:

(i) the solvothermal pressure-relief process during which the produced gas into the high-pressure vessel is controlled

Recent trends in solvothermal processes

Over the last years, some specific trends have been defined for the improvement of solvothermal processes, in particular through the control of the reactions and the preparation of specific material either as nano-crystallites or single crystals. In most cases, these trends are triggered by the development of these processes for industrial applications.

The control of solvothermal reactions

In order to manage solvothermal processes, new solvents have been developed such as ionic solvents due to their physico-chemical properties but also as directing agents [124], organic solvents [125], or some mixed solutions [126].

In parallel, several reaction mechanisms have also been investigated, such as the controlled reduction method [127], the fluorination process [128], the hot-press method [129], the resin-assisted synthesis [130], the synthesis under external magnetic field [119], and the interdiffusion process for the preparation of nanotubes [131].

The preparation of specific materials

As specific materials, we can consider materials with a macro-or a nanosize:

-with a particular composition (for example the extension from oxides systems to fluorides

--

The solvothermal growth of single crystals for industrial development initiated important works in particular with . In this domain, the research focused on three main goals: (i) the improvement of the kinetics for the solvothermal crystal growth, (ii) a lower concentration of defects (chemical defects such as impurities or physical defects such as dislocations), and (iii) the economical cost of the solvothermal process compared to the conventional ones (in particular in the case of ZnO).

The development of specific processes (green chemistry, synthesis of prebiotic molecules, wastes retreatment, etc.)

Hydrothermal processes in green chemistry

In the future, the increase of the earth's surface temperature is one of the major concerns. This is in part due to CO 2 emission from fossil-fuel combustion. The sequestration of CO 2 could be achieved through its transformation to organic components able to be used as basic products [153]. This is why during these last years, hydrothermal processes based on CO 2 have been developed [154,155].

Another domain of application for hydrothermal processes deals with the production of bioethanol and biofuel [156].

Hydrothermal synthesis of prebiotic molecules

Hydrothermal processes have been considered a possible explanation for the origin of life [157]. Consequently, this route is continuously explored [158].

Hydrothermal post-treatments

Solvothermal post-treatments were developed either for the control of the preparation of specific materials or for the conversion of wastes.

In the first case, the control of the crystal structure of titanium dioxide through the hydrothermal treatment of titanate nanotubes under acidic conditions can be an interesting illustration [159].

Concerning the hydrothermal retreatment of wastes, this topic (initiated in particular by Yamasaki) is still under investigation [160].

Continuous solvothermal processes

The development of continuous solvothermal processes is an important challenge from an economical point of view. During the last years, different nanoparticles have been prepared by this means [161][162][163].

Conclusions

Solvothermal processes are important in different scientific areas involving both basic research and industrial applications.

In order to improve such processes in the near future, four main domains need to be investigated:

----the potential development of different solvents (including aqueous solutions) through a better understanding of their physico-chemical properties; the investigations involving the evolution of the physico-chemical properties of these solvents in the reaction conditions; a better knowledge of the in situ chemical mechanisms governing the formation of the target material; and the impact of different parameters (in particular pressure) to control the reaction mechanisms.

3 3 2 [

 2 SnCl Á5H O] as reactants and NH OH as the mineralizer. The resulting 4 2 4 precipitate, after drying, was treated under solvothermal conditions (250 ꢀC, autogeneous pressure) using different solvents: ethylene glycol (C H O ), 2 polyethylene-glycol [H(OC H ) OH], and ethanol [C H OH]. the resulting material, solvent viscosity seemed to play an important role in the size and the physico-chemical properties of the resulting ITO crystallites. The increase of the viscosity [from C H OH to 2 5

2 thiourea(

 2 CS(NH ) ), CS being the solvent and KOH or KNO as the mineralizer at

  ethanol, dimethylformamide (DMF), toluene, pyridine] were investigated for the preparation of CeO 2 nanotubes[100]

  [104]; (ii) modify the oxidation state of transition elements, for example the n? value for V n? in the case of oxyfluorides prepared in the Na-V-(O)F and K-V-(O)F systems by solvothermal methods [105]. If V 4? oxyfluorides are prepared at 100 ꢀC, only V 3? fluorides are obtained at higher temperature as 220 C; (iii) optimize the size and shape of nanoparticles by increasing the temperature or reaction time. For example, the size and the crystallinity of the resulting CdS nanorods are improved [106]. ꢀ The role of pressure in solvothermal processes The role of pressure in solvothermal processes has not yet been extensively investigated as the autogeneous pressure value commonly being used. In the solvothermal crystal growth of aquartz, a specific evaluation of the impact of pressure on the concentration of chemical defects has been investigated [107]. Due to the compensated substitution Al 3? ? Si 4? and OH -? O 2-, the IR coefficient a 3500 (corresponding to the maximum IR absorption for OH groups) is used to estimate the concentration of chemical defects. The decrease of a values with increasing pressure underlines the role of this parameter on the Al 3? insertion into the a-SiO 2 lattice. Impact of the pH of the solution The pH value can play different roles in solvothermal processes for: (i) the preparation of specific materials [108]; (ii) the stabilization of different structural forms [109]; (iii) the control of the size and morphology of crystallites: --through the modification of the decomposition of one reactant [110]; through the change of the ionization equilibrium of the additive [111];

  [120]; (ii) the solvothermal scission-template-transportation route with the formation of an intermediate compound acting as the template [121]; (iii) the extraction solvothermal method [122]; (iv) and the phase-transfer catalysis [123].


Angéline Poulon-Quintin

Cyril Faure

Lionel Teulé-Gay

Jean-Pierre Manaud

Bilayer systems of tantalum or zirconium nitrides and molybdenum for optimized diamond deposition

Thermo chemical computing validates the stability of different nitrides against Co, Mo, and methane up to 1150 K, showing the highest chemical stability against carburization for ZrN and TaN under static conditions. Single zirconium and tantalum nitrides layers have been sputtered onto WC-Co substrates as diffusion barriers and buffer layers under specific reactive sputtering conditions. To improve the nuclei density of diamond during CVD processing, a thin Mo extra layer has been added (b 500 nm). In this study, two bilayer systems have been tested: TaN-Mo and ZrN-Mo. Nano crystalline diamond has been grown under negative biased substrates. After diamond deposition, a massive carburization of molybdenum and tantalum nitride is observable whereas zirconium nitride is not. Nevertheless, a small amount of cobalt has migrated through the ZrN layer. The better efficiency of the ZrN layer to prevent diffusion of the Co element, leads to expect an increased adhesion of diamond on ZrN-Mo bilayer coating. A TEM study is done to improve understanding of phenomena occurring at the interfaces during process.

Introduction

the way through which elemental Co could reach the diamond WC-Co substrate interface. According to the phase diagram of Co-C system in the temperature range of 975-1275 K (typical diamond CVD temperature), carbon is soluble in Co up to 0.2-0.3 wt.%. During the initial stage of the CVD process, the WC-Co substrate is exposed to a hydrocarbon rich atmosphere and carbon species can diffuse into the bulk of the binder phase until the carbon solubility is exceeded. Once the carbon concentration at the substrate surface is sufficiently large to promote solid carbon condensation, the preferential formation of graphite layer is promoted by the presence of the binder. Moreover, as a further consequence of carbon solubility in the binder phase, diffusion of carbon from the deposited diamond into the metal binder can also occur. To suppress the interaction between Co and the deposited diamond film, an interlayer with low diffusion coefficients for C and Co has to be used. Moreover, an interlayer material with an intermediate thermal expansion coefficient between WC-Co and diamond could relieve the residual thermal stresses. Improvement in adhesion using various interlayers, ranging from amorphous carbon to metallic materials (Cu, Mo, Nb, Pt, Ti, Ta, W and Ag) and ceramics (SiC, Si N , TiC, TiN and WC) has already been reported [22].

Diamond films grown by Chemical Vapor Deposition (CVD) are widely used as surface overlay coating onto WC-Co cutting tools to improve their performances (endurance, surface hardness, thermal conductivity, friction, wear protection…). It is well known that the Co content has to be lower than 6 wt.% and the fraction of other carbides should be as small as possible to achieve satisfactory adhesion levels of CVD diamond films. The reduction of the Co content from the substrate causes toughness deterioration and reduces the tool lifetime. Whereas the reduction of the cemented carbide grain size (b1 μm) combined with an increase of the Co content (N10 wt.%) leads to mechanical properties close to those obtained with conventional grain size (N10 μm) and 3 wt.% Co content [1]. However the increase of the cemented carbide Co content is detrimental for the diamond adhesion. During deposition of sp 3 diamond film by CVD process onto sintered WC-Co substrates, formation of sp² graphitic species can occur because of the Co binder catalytic properties related to their valence electrons [2][3][4]. To improve the diamond adhesion with the substrate and the disagreements encountered, many pretreatment methods have been applied such as selective etching of Co [5][6][7][8][9], forming stable Co compounds [10][11][12][13], giving a suitable diffusion barrier layer on the substrates [14][15][16][17][18][19][20] or a tungsten-carbide gradient coating [21] etc. The latter three strategies aim at blocking 3 4 In a previous paper [17], we have shown that refractory nitrides like ZrN, NbN and TaN can act as efficient Co diffusion barriers for CVD diamond deposition on WC-Co substrates containing 12 wt.% of cobalt. In this study, we propose to focus on bilayer systems composed of a Co diffusion barrier (TaN or ZrN) and a thin metallic layer to improve the diamond nucleation. Amorphous Si coating can act as a promoter for diamond nucleation thanks to its good carbide former properties [23]. Unfortunately, Si is not compatible with most of transition or refractory nitrides because the lowest chemical potential values of formed silicides lead to inescapable destruction of the barrier. In this study, molybdenum is used as another candidate to improve the nuclei density of diamond during CVD processing.

Experimental

Deposition of nitrides layers

The substrates were WC-Co 12 wt.% plates 16 × 16 mm² and 5 mm thick. Prior to deposition, WC-Co subtrates have been blasted with alumina micro particles under the same conditions (air flow absolute pressure : 2 × 10 5 Pa; distance 10 cm), then cleaned in acetone for 1 h in an ultrasonic bath and finished under propan-2-ol vapor for 1 min. The roughness R a measured (White light confocal profilometer VEECO) was lower than 500 nm.

Nitride layers were deposited by reactive magnetron sputtering using Zr, Ta and Mo targets. The sputtering plant was Plassys MP700 equipped with a 75 mm in diameter target. Targets used were disks of pure Ta (99.98%), Zr (99.8%) and Mo (99.98%) from Cerac. Ultra pure nitrogen (Dinal Alphagaz 2; N99.999%) was added as reactive gas to argon gas (Dinal Alphagaz 2) purified through a R.D. Mathis purifier (titanium getter at 1075 K). The ultimate pressure for all deposits was in the range of 2-3×10 -5 Pa. During deposit process, substrates were set on a heating plate with electronic temperature monitoring. The temperature of the substrates fixed at 673 K [17], was controlled with a thermocouple placed in contact with their surface. The distance between the target and the substrate surface was set on 70 mm for all depositions to assure a uniform thickness of deposited layers in respect of a fair deposition rate. Deposition rates were accurately extrapolated from nitrides layers thicknesses deposited on silicon wafers after phases identification by X-ray diffraction. The deposition parameters are presented in Table 1.

Diamond deposition

Diamond was deposited in a microwave enhanced plasma chemical vapor deposition reactor (silica tube crossing orthogonally the wave guide powered with a 1.2 kW) using a hydrogen-methane mixture. The total pressure was held at 40 hPa. Substrates were biased under -200 V. During the deposition time (6 h maximum), the partial pressure of methane was held at 0.4 hPa as described elsewhere [17]. The temperature of the sample surface was set at 1150 K (measured with a Raytek biband pyrometer).

Material characterization

X-ray diffraction was used to determine the deposited compounds. Experiments were carried out with a Philips PW1820 diffractometer with θ-2θ geometry and a Cu K α radiation proportional detector. Morphology observations were carried out with a SEM Jeol JMS 6360A microscope equipped with a Jeol EDS system. TEM observations were carried out with a Jeol 2200FS equipped with a Jeol EDS and STEM system. Auger Electron Spectroscopy profiles were carried out with a VG Instrument microlab 310F. Etching was performed with an argon ion beam under 4 kV voltage and an ionic current of 0.5 mA. The spot surface was 3 μm² and the theoretical etching rate was 0.2 nm s -1 .

Results

Bilayer morphology before diamond growth

For Zr and Ta, binary phase diagrams show mainly the existence of two compounds M 2 N and MN with M Zr or Ta. Due to its lower enthalpy of formation M 2 N compounds should be preferentially synthesized under low temperature and pressure conditions. Thanks to the reactive magnetron sputtering process, the formation of the MN phase is possible when the partial pressure of nitrogen is increased. X-ray diffraction spectra (Fig. 1) show the presence of the hexagonal TaN (JCPDS 39-1484) and cubic ZrN (JCPDS 35-0753) compounds as well as the cubic Mo (JCPDS 03-065-7442) phase. TaN is obtained with 10% of nitrogen in sputtering gas whereas ZrN is obtained with only 6%. TaN layer is textured {110} and Mo layer {110}. ZrN layer is less textured with the presence of two main orientations {200} and {111}. AES in depth profiles (Fig. 2) show that after the bilayer deposit, no trace of Co is observable inside the diffusion layer. This result is in agreement with the thermochemical calculation from Thermodata Software [24]. No diffusion of Mo element inside the TaN or ZrN diffusion barriers is also noticeable even if the process occurs at 673 K. TEM observations of the TaN-Mo and ZrN-Mo bilayer coatings show that the materials deposited are very dense (Fig. 3). The morphology of the grains is more columnar for Mo and more of a structure of transition with fiber grains densely joined for TaN and ZrN layers. Average grain size width is 200 nm in diameter for Mo, 150 nm for TaN and 100 nm for ZrN with respectively a height of 0.4, 1.2 and 1 μm.

Thermochemical computation

Computing calculation with Thermodata Software has shown that for a temperature range of 1043 to 1143 K, a total pressure of 40 hPa and a partial pressure of 0.6 hPa of CH 4 , the most stable nitrides against methane carburization are TaN and ZrN [17]. The stability of the nucleation layer under diamond deposition condition has been computed considering the enthalpy of formation of metallic element and all related carbides available in the data base for the following reaction:

xM ðsÞ þ yCH 4ðgÞ →M x C yðsÞ þ 2yH 2ðgÞ :
Comparison between Mo, Nb and Ti has been done. Calculation results in static conditions at 1150 K for the expected reactions at the interface between the nucleation layer and the diffusion layer, are presented in Table 2. These three elements are good carbide formers as expected for a nucleation layer material. However, their stability with the materials selected for diffusion barriers varies. For titanium, nitridation can occur as consequence with possible reaction with TaN. When Ti is carburized, the same phenomenom might occur. Concerning Nb, the possibility to form Nb 2 C during carburization limits its used to ZrN diffusion barrier. Whereas the choice of Mo can be applied to both (TaN and ZrN) and seems to be a good candidate for the nucleation layer.

Characterization of the bilayers after diamond deposition

Diamond has been deposited at 1150 K on samples layered with TaN-Mo and ZrN-Mo. X-ray diffraction spectra (Fig. 4) show that ZrN layer has been preserved during diamond process whereas the hexagonal TaN is transformed into the cubic TaC phase (JCPDS 65-8264) less textured with the presence of four main orientations: {111}, {200}, {220} and {311}. For both diffusion layer types, Mo is carburized as witnessed by the presence of Mo 2 C (JCPDS 03-65-8766). Some small peaks seem to confirm the presence of Co after CVD process. However independently of the bilayer system observed, no molybdenum layer is still observable. On TEM pictures, the presence of TaC and ZrN layers is obvious but not for Mo layer. However close to the diffusion barrier where Mo layer would have been localized some circle grains are noticeable and dispersed into the diamond layer (Fig. 5). Based on EDS quantitative analyses and XRD results, these phases have been identified as Co 6 Mo 6 C 2 (JCPDS 80-6339) and Co 2 Mo 6 C (JCPDS 89-4885) (Fig. 4). Raman analyses show that sp 2 graphite formation is not prevented whatever the bilayer system studied (Fig. 6) with amorphous sp² carbon species observable (result confirms by TEM diffraction study). Concerning the sp 3 species formed, presence of nano-grained diamond is established. The calculated ratio values between sp 3 and sp² species are similar (XPS quantification); respectively 0.8 for TaN-Mo and 0.9 for ZrN-Mo after 6 h of diamond deposit process. An estimation of the nucleation density after 3 h of diamond growth is 1.7 × 10 6 nucleus/cm² for TaN-Mo and 1.3 × 10 6 nucleus/cm² for ZrN-Mo.

Discussion

The results clearly demonstrate that as intermediate layers of Mo, TaN or ZrN on WC-Co substrates show different effects on subsequent diamond growth. Based on the comparison with results obtained without the Mo layer used on the top of TaN or ZrN intermediate layers, the diamond nucleation has been significantly promoted. For TaN and ZrN, the increase is time 20. The consequences on the adhesion of the diamond film have to be quantified. We have already showed [17] that an intermediate layer (TaN or ZrN) can effectively prevent the formation of detrimental graphite at film substrate interface thanks to their Co diffusion barrier properties. The low nucleation density of diamond observed and the diamond grain morphology (micrometric crystallites) have led to weak bonding strength between diamond film and these interlayers. Whereas the use of bias during diamond deposition process has led to the formation of diamond nano-grains (NCD) as already be observed by other teams [25][26][27]. The Mo layer modification of morphology has also to be correlated with the use of a negative bias which leads to an ionic bombardment of the coated surface. Some authors have noticed modification of the surface roughness when they use Si wafer as substrate and the smoothing of edges [25]. They suppose that the polarization led to the formation of an ion flow deteriorating the plane faces increasing their roughness. These new defects are expected to be the new nucleation sites for diamond. They have shown the formation of NCD diamond. In our case, we need to decorrelate the beneficial effect of the Mo layer on nucleation from the bias effect. The printing effect (i.e. surface etching of the clusters) of the bias can explain the NCD diamond formation. However, the beneficial effect of Mo thin film on improved sp 3 formation can be related to the following two important factors. First, as already observed with the use of Al nucleation layer coated on WC-Co substrate [28], Mo reacts with the binder phase Co during CVD process with the formation of Co 6 Mo 6 C 2 and Co 2 Mo 6 C phases. Thereby this is supposed to decrease the graphite formation during CVD process. This function is especially important for diamond synthesis when Co binder is present even if a diffusion barrier such as TaN or ZrN is used. EDS-STEM study shows that whatever the diffusion barrier type (0.4 μm thickness both), Co can diffuse and is localized close to the top near surface inside diamond. With TaN-Mo bilayer system, after diamond CVD process, a lot of no-alloyed Co is observable (blue color on overlay view, Fig. 5). With ZrN-Mo bilayer system, after diamond CVD process, Co is alloyed with Mo and C (pink color on overlay view, Fig. 5) and few no-alloyed Co is observable. We can suppose than an optimization of the Mo thickness can be done to improve the efficiency of the TaN-Mo system and reduce the formation of graphite due to the Co binder close to the top interface with the consequence of the adhesion property degradation.

Second, the fine-sized films have high surface energy and chemical reactivity. In this study, Mo layer still not exist and due to the CVD process conditions, Mo is localized inside circular grains close to the interface increasing the surface energy and chemical reactivity. At the initial stage of CVD process, the Mo grains are subjected to a carburization process and are transformed into carbide phase like Mo 2 C. EDS-STEM analysis shows that for ZrN-Mo bilayer system, a significant quantity of Mo is not combined with Co as for TaN-Mo bilayer system. XRD spectrum confirms also the presence of a significant peak representative of the Mo 2 C phase for the ZrN-Mo bilayer system. These carbide particles can act as active sites for diamond nucleation producing a high nucleation density which should promote adhesion strength at diamond/substrate interface due to the increased contact area observable on TEM pictures (Fig. 5) and lowered interfacial voids [29,30]. Due to the high nucleation density promoted by the presence of Mo, nano-scaled diamond starts to nucleate at the early stage of process. This mechanism is confirmed by the experimental results obtained at different deposition periods (to be published). The formation of the carbide phase grains mixed with nano-scaled diamond grains close to the TaC or ZrN interfaces is comparable to a layer with a thickness around 2 μm. Such layer in contact with the outer-most diamond film can provide a stronger chemical bonding in comparison with the conventional mechanical interlocking as well as a continuous interface. Similar effects have been observed with other carbide metal formers employed as continue interlayers for diamond deposition such as Ti [31], Al [18] and Mo [START_REF] Huang | Raman analyses of the multilayer materials after diamond deposition: TaN-Mo and ZrN-Mo systems[END_REF].

During the CVD process, formation of graphite or diamond necessitates a sufficiently large carbon concentration at the bilayer surface to promote solid carbon condensation. Both system studied consumes one part of carbon to carburize Mo and Mo-Co compounds. Methane carburization of these layers is a gas to solid reaction whose rate is expected to be determined by carbon diffusion into the near surface formed carbide. However due to the Mo nucleation layer, diamond nucleation and growth have started and nucleus can also act as localized sources of carbon. For TaN-Mo bilayer, diffusion of carbon from the deposited diamond (or graphite) and from the gas solid reaction occurs leading to the full carburization of TaN. Whereas for ZrN-Mo bilayer a supposed lower diffusion rate of carbon leads to the interlayer phase stability [17]. This mechanism modification combined with the presence of Mo can explain the difference of behavior of Co binder localization when compared to our previous study [17]. Previously using a single diffusion barrier layer, we have reported for TaN a very low Co diffusion compared to ZrN. We have supposed that a progressive chemical transformation of the nitride into carbide from the surface to the interface instead a massive transformation, kept close to the interface the Co species without their diffusion to the surface where diamond nucleation occurs. The presence of Mo, its carburization and its affinity for Co (solid solution and binary compounds formation) as well as the carburization of the binary compounds formed may change the thermodynamic equilibrium of the system close to the interface with diamond. Moreover the efficiency of diamond nucleation thanks to Mo-phased grains combined with the effect of bias printing is supposed to increase the concentration of carbon reactive species close to the TaN top surface changing the velocity of the previously supposed progressive chemical transformation of the nitride into carbide from the surface to the interface. If now we suppose a massive transformation of TaN into TaC then Co species can diffuse to the surface where diamond nucleation occurs and partially reacts with the Mo-Co compounds. Co massive presence close to the TaC top surface is then explained. For the ZrN-Mo bilayer system, no significant modification of the mechanism is noticed excepted for the diamond nucleation phenom-enon. Because XPS quantification of the ratio between sp 3 and sp² after 6 h of diamond deposit shows no significant differences between the two bilayer systems, we can suppose that only the adhesion would be affected due to the presence of the Co binder close to the interface for TaN-Mo bilayer with the consequence of an increase of the sp² graphite quantity at the interface.

Conclusion

The efficiency of the different bilayer systems (TaN-Mo and ZrN-Mo selected using thermochemical computation) with the aim to increase the diamond film growth has been investigated. TaN and ZrN thin film pre-coated on WC-Co substrates as Co diffusion barrier leads to a lower diamond nucleation density compared to the same system with a Mo top layer. After diamond deposition, a massive carburization of molybdenum and tantalum nitride is observable whereas zirconium nitride is not. Nevertheless, a small amount of cobalt has migrated through the ZrN layer and a more significant through the TaN layer. An optimization of the Mo layer thickness could be done to reduce this phenomenon. The benefic effect of using a negative bias during CVD process to form nano-sized diamond is validated. The better efficiency of the ZrN layer to prevent diffusion of the Co element when a bilayer system is used, leads to expect an increased adhesion of diamond on ZrN-Mo bilayer coating.

Fig. 1 .Fig. 2 .

 12 Fig. 1. X-ray diffraction patterns of TaN-Mo and ZrN-Mo bilayers deposited onto WC-Co substrates.

Fig. 3 .

 3 Fig. 3. TEM pictures (BF) of TaN-Mo and ZrN-Mo bilayers deposited onto WC-Co substrates. a) TaN fiber grain structure, b) Mo columnar grain structure, and c) ZrN fiber grain structure.

Fig. 4 .

 4 Fig. 4. X-ray diffraction patterns after diamond deposition process (6 h) onto TaN-Mo and ZrN-Mo bilayers deposited onto WC-Co substrates.

Fig. 5 .

 5 Fig. 5. TEM pictures and their corresponding EDS analysis of the multilayer materials after diamond deposition: TaN-Mo and ZrN-Mo systems.

Table 1

 1 Reactive sputtering deposition conditions of TaN, ZrN and Mo layers at 673 K.

	Sample	Total pressure	Partial N 2 pressure	N 2 percentage	Power density	Target voltage	Deposition rate	Layer thickness
		(Pa)	(Pa)	(%)	(W cm -2 )	(V)	(nm min -1 )	(μm)
	TaN	0.5	0.05	10	2.3	-160	10.0	1
	ZrN	0.5	0.03	6	4.0	-280	16.5	1
	Mo	0.5	-	-	2.3	-190	20.0	0.2/0.4

Table 2

 2 Calculated values of free enthalpy for the phases susceptible to be formed during diamond process (Thermodata Software) at 1150 K.

	Interface	Chemical reaction	ΔrG°1150 K
			(kJ/mol)
	TaN/Ti	TaN + Ti → Ta + TiN	-69.4
	TaN/TiC	TaN + TiC → TaC + TiN	-38.7
	ZrN/Ti	ZrN + Ti → Zr + TiN	28.5
	ZrN/TiC	ZrN + TiC → ZrC + TiN	13.3
	TaN/Nb	TaN + Nb → 1/2Ta 2 N + 1/2Nb 2 N	6.3
	TaN/NbC	TaN + NbC → TaC + NbN	28.0
	TaN/Nb 2 C	TaN + 1/2Nb 2 C → 1/2Ta 2 C + NbN	23.8
	ZrN/Nb	ZrN + 2Nb → Zr + Nb 2 N	109.3
	ZrN/NbC	ZrN + NbC → ZrC + NbN	66.9
	ZrN/Nb 2 C	ZrN + NbC → ZrC + 1/2Nb 2 N + 1/4N 2	103.6
	ZrN/Nb	ZrN + 2Nb → Zr + Nb 2 N	109.3
	TaN/Mo	TaN + Mo → 1/2Ta 2 N + 1/2Mo 2 N	169.3
	TaN/MoC	TaN + 2MoC → TaC + Mo 2 C + 1/2N 2	22.6
	TaN/Mo 2 C	TaN + Mo 2 C → TaC + Mo + 1/2N 2	75.2
	ZrN/Mo	ZrN + 2Mo → Zr + Mo 2 N	435.3
	ZrN/MoC	ZrN + 2MoC → ZrC + Mo 2 C + 1/2N 2	61.5
	ZrN/Mo 2 C	ZrN + Mo 2 C → ZrC + Mo + 1/2N 2	114.1

Acknowledgments

The author would like to thank "Le Conseil Regional d'Aquitaine" for its financial support, Dr. Michel Lahaye (CeCaMa, Université de Bordeaux1) for AES characterization and Dr. Christine Labrugère for XPS quantifications.