Robin Graillot-Vuilecot

Anne-Lise Thomann

Thomas Lecas

Christophe Cachoncinlle

Eric Millon

Amaël Caillard
email: amael.caillard@univ-orleans.fr

Hot target magnetron sputtering process: effect of infrared radiation on the deposition of titanium and titanium oxide thin films

Keywords: hot target, magnetron sputtering, titanium, infrared radiation

Titanium-based thin films are synthesized by the magnetron sputtering of a titanium target in pure argon or in a mixture of argon and oxygen (reactive regime) in three experimental configurations: a hot target configuration for which the target is thermally disconnected from the cooled magnetron body, a cold/conventional target configuration, and a cold target configuration with an additional infrared source introduced in the chamber to understand the role of the radiation on thin film properties. In hot configuration, the target temperature and its emissivity are measured and compared to a 0D steady state model taking into account the main different cooling mechanisms. IR radiations constitute the major part of the total energy deposited on the substrate during hot target magnetron sputtering, whereas an increase of target temperature allows a rise of sputtered Ti and Ti + species fluxes and slightly affect their kinetic energy. In reactive regime, the obtained films are constituted of polycrystalline anatase TiO2, and an increase of target/IR source temperature leads to an improvement of crystallinity. In pure argon, hot configuration seems to enhance its porosity, induces columns exhibiting sharp-features, slightly modifies its crystalline properties and helps to decrease internal compressive stress.

Introduction

Magnetron sputtering and evaporation (vacuum, arc or electron beam based) are nowadays the most widely used physical vapour deposition (PVD) processes to eject atoms from a solid target allowing the formation of nanoparticles or thin films. In vacuum evaporation process [START_REF] Donald | Vacuum Evaporation and Vacuum Deposition[END_REF], especially adapted for elements with low melting point, material is heated until atoms acquire enough energy to leave surface, and cross vacuum chamber with a very low energy (between 0.1 and 0.3 eV depending on the material) toward substrate.

Deposition rate is very sensitive to the nature of materials, making thus difficult to evaporate high melting point elements and to obtain thin films of complex stoichiometry (polycationic oxides or alloys). In sputtering techniques, target is submitted to the bombardment of ions generated in a plasma or issued from an ion beam.

This results in the ejection of energetic atoms (few 10's eV depending on sputtering conditions) that are able to form a dense thin film (compared to the thermal/vacuum evaporation technique), which can enhance the film hardness [START_REF] Kelly | Magnetron sputtering: a review of recent developments and applications[END_REF][START_REF] Helmersson | Review Ionized physical vapor deposition (IPVD): A review of technology and applications[END_REF] but induces tensile or compressive intrinsic stress depending and the deposition conditions [START_REF] Abadias | Review Article: Stress in thin films and coatings: Current status, challenges, and prospects[END_REF][START_REF] Magnfält | Compressive intrinsic stress originates in the grain boundaries of dense refractory polycrystalline thin films[END_REF][START_REF] Spaepen | Interfaces and stresses in thin film[END_REF]. During sputter deposition, substrate is submitted to action of other energetic particles (ions, electrons, high-energy neutrals) that can transfer energy and thus affect film formation (sputtering, desorption, implantation, diffusion…) and morphology [START_REF] Thornton | Influence of apparatus geometry and deposition conditions on the structure and topography of thick sputtered coatings[END_REF][START_REF] Anders | A structure zone diagram including plasma-based deposition and ion etching[END_REF][START_REF] Kersten | The energy balance at substrate surfaces during plasma processing[END_REF][START_REF] Bornholdt | Characterization of the energy flux toward the substrate during magnetron sputter deposition of ZnO thin films[END_REF] In turn, film properties such as adhesion, stress, density, grain size, orientation and defects could be modified [START_REF] Mattox | Particle bombardment effects on thin-film deposition: A review[END_REF][START_REF] Petrov | Average energy deposited per atom: A universal parameter for describing ion assisted film growth[END_REF][START_REF] Sarakinos | High power pulsed magnetron sputtering: A review on scientific and engineering state of the art[END_REF][START_REF] Cormier | Measuring the energy flux at the substrate position during magnetron sputter deposition processes[END_REF]. Even if magnetron sputtering techniques are widely used in industry, the synthesis of defect-free compounds containing labile (as fluor, lithium, sodium …) and low melting point (as gallium, indium ...) is a difficult task. Moreover, the synthesis of oxides and nitrides with appropriate stoichiometry, good crystalline quality at high deposition rate is really challenging [START_REF] Kubart | High rate reactive magnetron sputter deposition of titanium oxide[END_REF][START_REF] Toku | The effect of oxygen concentration on the low temperature deposition of TiO2 thin films[END_REF][START_REF] Musil | Low-temperature sputtering of crystalline TiO2 films[END_REF]. This leads manufacturers to mix various deposition techniques and use post annealing treatments, which complicates the process and lowers its efficiency (in terms of energy saving and environmental fingerprint). A way to overcome these issues would be to perform sputtering of a hot target heated by ion bombardment, so emitting infrared (IR) radiations and acting as a substrate front heater. The effect of target radiation has been already observed almost 20 years ago [START_REF] Rohde | The energy influx during plasma deposition of amorphous hydrogenated carbon films[END_REF]. But it was rather regarded as an undesired side effect. Whereas a cooling system is usually used to prevent the overheating of magnets and to limit the temperature of target in conventional sputtering configuration, it is possible to promote target surface heating by thermally disconnecting the rear side of target from cooled magnetron. This can even lead to a liquid target surface. Performing sputtering with a high target temperature leads to an increase of deposition rate [START_REF] Kubart | High rate reactive magnetron sputter deposition of titanium oxide[END_REF][START_REF] Toku | The effect of oxygen concentration on the low temperature deposition of TiO2 thin films[END_REF][START_REF] Musil | Low-temperature sputtering of crystalline TiO2 films[END_REF][START_REF] Rohde | The energy influx during plasma deposition of amorphous hydrogenated carbon films[END_REF][START_REF] Musil | High-rate reactive deposition of transparent SiO2 films containing low amount of Zr from molten magnetron target[END_REF][START_REF] Yang | Low substrate temperature fabrication of high-performance metal oxide thin-film by magnetron sputtering with target selfheating[END_REF][START_REF] Lapshin | Composition and Structure of Copper Oxide Films Synthesized by Reactive Magnetron Sputtering with a Hot Target[END_REF][START_REF] Ismailov | Sputtering of a hot ceramic target: Experiments with ZnO[END_REF][START_REF] Laurikaitis | Deposition of Zr-ZrOX and Y-YXOY films by reactive magnetron sputtering[END_REF][START_REF] Sidelev | High-rate magnetron sputtering with hot target[END_REF][START_REF] Sidelev | A comparative study on the properties of chromium coatings deposited by magnetron sputtering with hot and cooled target[END_REF][START_REF] Grudinin | Chromium films deposition by hot target high power pulsed magnetron sputtering: Deposition conditions and film properties[END_REF], improves crystalline quality (especially oxides) and influences texture [START_REF] Musil | High-rate reactive deposition of transparent SiO2 films containing low amount of Zr from molten magnetron target[END_REF][START_REF] Yang | Low substrate temperature fabrication of high-performance metal oxide thin-film by magnetron sputtering with target selfheating[END_REF][START_REF] Lapshin | Composition and Structure of Copper Oxide Films Synthesized by Reactive Magnetron Sputtering with a Hot Target[END_REF][START_REF] Ismailov | Sputtering of a hot ceramic target: Experiments with ZnO[END_REF][START_REF] Laurikaitis | Deposition of Zr-ZrOX and Y-YXOY films by reactive magnetron sputtering[END_REF][START_REF] Sidelev | High-rate magnetron sputtering with hot target[END_REF][START_REF] Sidelev | A comparative study on the properties of chromium coatings deposited by magnetron sputtering with hot and cooled target[END_REF][START_REF] Grudinin | Chromium films deposition by hot target high power pulsed magnetron sputtering: Deposition conditions and film properties[END_REF] without any-annealing treatments, modifies stoichiometry, especially for thin film containing volatile/labile elements [START_REF] Kawamata | K-M-S (keep-molecules sputtering) deposition of optical MgF2 thin films[END_REF], enhances reactive sputtering stability by reducing hysteresis loop [START_REF] Mercs | Hot target sputtering: A new way for high-rate deposition of stoichiometric ceramic films[END_REF], enhances thin film hardness, decreases average surface roughness [START_REF] Sidelev | High-rate magnetron sputtering with hot target[END_REF][START_REF] Sidelev | A comparative study on the properties of chromium coatings deposited by magnetron sputtering with hot and cooled target[END_REF], allows heteroepitaxy framework film growth that usually requires high substrate temperature to overcome large lattice parameter mismatch between film and substrate [START_REF] Jung | High mobility Si0.15Ge0.85 growth by using the molten target sputtering (MtS) within heteroepitaxy framework[END_REF]. This could also help igniting/sustaining magnetron discharge during the sputtering of a ferromagnetic target by reaching its target paramagnetic state above the Curie temperature [START_REF] Caillard | Effect of the target temperature during magnetron sputtering of Nickel[END_REF][START_REF] Sidelev | Hot target magnetron sputtering for ferromagnetic films deposition[END_REF][START_REF] Sidelev | Angular thickness distribution and target utilization for hot Ni target magnetron sputtering[END_REF]. This emerging technology could be very promising, but so far the temperature dependence of mechanisms related to material ejection (sputtering, evaporation), plasma properties (composition, energy distribution, density, …) and the mechanisms of film growth (under the effect of kinetic and radiative energy transfers) are not yet fully understood, and actually difficult to access and interpret. Moreover, the heating of target induced by ion bombardment cannot be adjusted and controlled properly regardless the plasma properties. The precise knowledge (and the control) of the target temperature is also experimentally challenging either by using thermocouples, thermopiles (heat flux sensor), pyrometers or by analysing thermal radiation spectrum using visible or near-IR spectrometer [START_REF] Gerasimov | Method of magnetron target temperature evaluation by analysis of thermal radiation spectrum[END_REF][START_REF] Shapovalov | Studying heating of magnetron target based on measurement of substrate temperature[END_REF].

Modelling all the physical mechanisms occurring inside target/magnetron system (IR radiation loss, sputtering at high temperature, evaporation, enthalpy modification, heat propagation, residual cooling by conduction) and at surface (ion reflection, electron secondary emission, atom ejection …) to finally estimate target temperature is an interesting approach developed by different groups [START_REF] Shapovalov | Studying heating of magnetron target based on measurement of substrate temperature[END_REF][START_REF] Karzin | Simulation of heating of the target during highpower impulse magnetron sputtering[END_REF][START_REF] Bleykher | Magnetron deposition of coatings with evaporation of the target[END_REF][START_REF] Bleykher | Energy and substance transfer in magnetron sputtering systems with liquid-phase target[END_REF][START_REF] Bleykher | Evaporation factor in productivity increase of hot target magnetron sputtering systems[END_REF].

To explain the modification of thin film properties, two main energetic contributions can be pointed out. IR radiations emitted by hot target have been suspected by some authors [START_REF] Cormier | IR emission from the target during plasma magnetron sputter deposition[END_REF] to play a key role (in the heating of growing film for example), but have never been quantified and clearly related to thin film quality. Second, high-energy species which are generally produced at target in conventional (cold) magnetron sputtering transfer their energy to growing film and modify its physical properties. The influence of target temperature on flux and on the energy of these species has never been measured even if research groups investigated the influence of hot target during sputtering [START_REF] Kozin | Modeling of thermal processes in magnetrons with single hot target and "sandwich-target[END_REF][START_REF] Shapovalov | Hot Target. Physicochemical Model of Reactive Sputtering[END_REF] and pointed out the fact that this hot process should induce lower energy ejected species (due to evaporation mechanism).

The present paper aims at investigating both target temperature in hot configuration using a straightforward steady state model validated by experimental measurements and the effect of high target temperature on the sputtering process on the properties of titanium-based thin films (morphology, composition and structure) in pure Ar and in an Ar/O2 gas mixture. As model material in research (largely documented in the literature) and for various applications, titanium is a good candidate for this study. Especially the relative influence of IR radiations (emitted by the target) and the energy of condensing species (only Ti atoms and Ti + ions, here) on film properties are studied.

Experimental

Sputtering apparatus

Experiments were performed in a 6-way cross-shaped HV deposition chamber pumped down to 8.10 -5 Pa using a 500 l/s turbomolecular pump. The pressure is regulated to 1 Pa during deposition with 20 sccm Ar flow rate, whereas the O2 one is fixed to 0 for non-reactive sputtering (in metallic regime) and 2 sccm (50W) or 3 sccm (100 and 150 W) for reactive sputtering (in compound/oxide regime). The Ti target (thickness of 3 mm and a radius R of 25.4 mm) without (cold configuration) or with (hot configuration) a 1.8 mm thick MACOR® 2 inches ceramic disk, is clamped on magnetron (2 inches TORUS® from K.J. Lesker) backing plate. This ceramic disk thermally disconnects the target and promotes its temperature rise. However, clamping surrounding target to ensure electrical connection between magnetron and target induces a residual cooling.

The magnetron is powered by a DC power supply (Pinnacle plus 5 kW) between 50 and 200 W in this study.

A bichromatic pyrometer (IGAR 6 from Lumasense) placed outside vacuum chamber in front of a ZnSe viewport is used to measure target surface temperature (using the Wien law) and emissivity, denoted εpyro. The distance between ZnSe viewport and target is sufficient (about 300 mm) to neglect deposition on viewport and the modification of its transmission. The pyrometer is placed at the focusable position of 350 mm from the target in normal incidence and the spot size on the target is close to 3 mm. This emissivity naturally depends on the chemical properties of the target surface (metallic, partially or fully oxidized, roughness etc.) but also on the emitted angle Ѳ, the temperature T and the wavelength λ. In this study, εpyro is defined as the emissivity in the direction of the pyrometer (normal to the target surface, Ѳ = 0°) in the 2-2.5 µm wavelength range at a temperature T and is a quasi-monochromatic directional emissivity. The precision of the measured temperature is estimated to 2%. Figure 1a displays the experimental set up during the deposition of thin films or during the measurement of the target temperature. During the cooling step of the target, the temperature measured with the pyrometer displays an inflection point corresponding to Ti solid state transition (β -Ti transforms into α -Ti below 882°C) as previously reported by Tesař et al [START_REF] Tesař | On surface temperatures during high power pulsed magnetron sputtering using a hot target[END_REF]. This transition temperature is used to calibrate directional emissivity ratio between both spectral ranges (1.5-1.6 and 2-2.5 µm) in pure Ar gas, which is found to be between 1.28 and 1.35. These values indicate that the emissivity decreases when wavelength increases, which is commonly observed for metals [START_REF] Setién-Fernández | Spectral emissivity of copper and nickel in the mid-infrared range between 250 and 900 °C[END_REF] and Ti in particular [START_REF] Touloukian | Thermophysical Properties of Matter[END_REF]. A mean value of 1.3 will be considered in this study. In hot configuration, target temperature reaches 660, 870 and 1010°C (± 50°C, high uncertainty between runs of experiments due to the progressive erosion of target, see after) for the 3 respective power values (50, 100 and 150 W) in pure Ar gas. These temperature values are measured in racetrack area (defined by r=13 mm) and thus are higher than temperature averaged on the whole surface area. The upper temperature, 1010 °C, is largely inferior to titanium melting temperature at 1 atm (1668°C). A thermocouple placed on rear side of Si sample indicates a rise of its temperature up to values ranging from 120 to 220°C in hot configuration depending on DC power, whereas temperature inferior to 50°C are obtained in cold configuration. A higher temperature on the front side of substrate is probable considering that silicon is a good thermal insulator.

The substrate holder is placed in front of target with a substrate/target distance of 45 mm and an angle of 45° as shown on Figure 1a. A heating source (covered by a titanium thick film) is installed in front of substrate holder in order to mimic (at the best) target IR radiation emission with the same distance and angle (see Figure 1a). Its temperature can be regulated between 20 and 800°C, but only two temperatures have been tested in this study (600 and 800°C). The addition of this heating source allows to work in three different deposition configurations: hot target, cold target and cold target assisted by IR radiation.

In-situ / in-operando analysis

As shown on Figure 1b, substrate holder can be removed and replaced by a 6 mm in diameter thermopile sensor (also called heat flux sensor, HFM) [START_REF] Thomann | Diagnostic system for plasma/surface energy transfer characterization[END_REF]. This latter is covered by a zynolite back paint and a copper disk (6 mm in diameter and 100 µm in thickness) is fixed on it using a thermal conductive paste. Due to the presence of cold junctions (whose temperature is fixed to 5°C) just below the active junctions of sensor, the thermopile acts as a bolometer and directly gives the value of heat flux arriving to sensor (based on Seebeck effect) expressed in mW/cm². All the energy of species impacting to the surface of sensor is converted into heat flux and so such system allows to measure the energy flux density (Ф) coming from gas conduction/convection (Фgas, neglected in this study), the plasma species (Фplasma) and IR radiations emitted by the hot parts of chamber, especially the target (ФIR,target). A more detailed list of energy contributions will be given in the results section. In this Figure 1b, surface sensor is perpendicular to the glowing target surface (denoted  geometry), but the sensor has also been placed parallel to the target (// geometry) as indicated by dashed blue arrow. Sensor has been previously calibrated under vacuum according to the NIST protocol and using a homemade blackbody (ensuring absorption and emission coefficients close to 1) [START_REF] Thomann | Diagnostic system for plasma/surface energy transfer characterization[END_REF]. In our study, the black paint is covered by a copper disk and a Ti based thin film (10 min deposition time in pure Ar), so the sensor system cannot be considered as a blackbody anymore. A part of energy is reflected at thin film surface and must be taken into account. The reflected and absorbed parts of energy are difficult to precisely estimate. Considering emitting source as diffuse grey-body, we can estimate its steady-state temperature using the Stefan-Boltzmann law from the IR radiation (emitted by the hot target ФIR,target or the IR source ФIR,sourcet and expressed in mW/cm²) collected on sensor. Moreover, we consider that sensor temperature is much smaller than the emitting source one.

Ф 𝐼𝑅 (𝑥) (T) = Aσ 𝑇 4 with 𝐴 = α ε F 𝑥 S 𝑆 𝐻𝐹𝑀 (1)
where σ is the Stefan-Boltzmann constant (5.67 10 -8 W/m².K), A is a constant (without dimension), ε is the total emissivity which corresponds to ratio between the excitance of material and that of a blackbody, both integrated on Ѳ, T and λ (Ѳ defined the angle from the normal to the target surface). ε is considered as the total emissivity even if sensor only measured radiation emitted in its direction Ѳ. α is the total absorptivity of the sensor surface and Fx is radiative view factor (geometrical factor taking into account that only a portion of target IR emission can be detected by sensor) with x indicates the corresponding geometry (x= //, , 45°).

Using a ray-trace software and assuming that the spectral radiance of target is Lambertian, F// and F have been found to be equal to 3.7 10 -3 and 1.3 10 -3, respectively. F// can also be deduced from the analytical expression given in [START_REF] Modest | Radiative Heat Transfer[END_REF]. For substrate (Figure 1a) positioned with an inclination of 45°, software gives a view factor F45° of 3.3 10 -3 . We found that the ratio F45° / F// is equal to 0.89 (superior to the expected cos(45°) value), whereas 2.5 is obtained for F45° / F ratio. This 0.89 factor (close to 1) indicates that measurement performed in // geometry gives a good estimation of the flux energy density incoming onto substrate.

According to the large thickness of disk (100 µm) covering the black paint, we consider that all un-reflected energy is absorbed, converted into heat and transferred to thermopile. In both regimes (oxide and metallic), copper disk is covered (before recording the measurement) by a Ti metallic thin film (containing residual oxygen) prior to the recording of experiment. The thickness of this thin film reaches few hundreds of nanometers in // or in 45° geometry. This thickness is high enough to consider Ti thin film as an IR transducer absorbing all IR radiation up to few µm [START_REF] Amin-Chalhoub | Thermal conductivity measurement of porous silicon by the pulsed-photothermal method[END_REF][START_REF] Amin-Chalhoub | Improvement of the identification of multiwall carbon nanotubes carpet thermal conductivity by pulsed photothermal method[END_REF]. So the absorptivity of sensor surface is not equal to the copper one but to the Ti thin film in 45° or // geometry. Considering the evolution of refractive index and of the extinction coefficient of titanium bulk in near IR between 1.5 and 2 µm wavelength, an absorptivity of 0.3-0.35 has been found for such Ti thin film (obtained from Openfilter software and from the Fresnel laws).

However, the Ti thin film is not purely metallic and contains oxygen species which increased its global absorptivity. Previously experiments have shown that an absorptivity of 0.5 for such Ti thin film was a good approximation [START_REF] Melhem | Structural, Optical, and Thermophysical Properties of Mesoporous Silicon Layers: Influence of Substrate Characteristics[END_REF][START_REF] Sparrow | Radiative Heat Transfer[END_REF]. This is the value that we took in this study. During the recording of heat flux in oxide regime, an oxide thin film with a thickness inferior to 100 nm covers the previously deposited Ti metallic thin film but the absorptivity of sensor surface is not modified considering that such additional TiOx thin layer onto Ti, is transparent to IR radiation. In this study, the absorptivity α of the Ti thin film (acting as IR transducer) is close to its emissivity. Indeed, considering local thermodynamic equilibrium, most metal can be considered as a diffuse grey body (contrary to most non-conductive materials).

An EQP energy resolved mass spectrometer (Hiden Analytical, EQP1000) is used to estimate the amount and the energy of species interacting with the substrate. Only the energy of Ti sputtered atoms and of Ti + ions have been considered in this study. The 200 µm orifice is used and the EQP faces target center at a distance of 100 mm. The distribution of Ti atoms has been obtained after their ionisation (electron bombardment) in EQP internal ioniser (RGA configuration). During ionisation, the initial energy of Ti atoms are of course modified, but we assume that this modification is similar for all energy distributions taken into account that these distributions are close to each other. In the mass spectrometer, ioniser is surrounded by a cage biased to 3.0 V (in our experiments) giving energy to ions and ensuring their transport up to the detector, so 3.0 eV has been subtracted to the distributions. An auto-tune procedure has been performed in order to adjust the different internal parameters of EQP and to optimize signal intensity.

A quartz microbalance (8 mm in diameter, 6 MHz quartz covered by gold) is placed in front of the center of Ti target (at a 120 mm distance) using water cooled crystal holder in order to measure deposition rate through a STM-2 transducer (INFICON).

Ex-situ analysis

Eight metallic and eight oxide thin films were deposited on (100) Si substrates at three DC power values (50, 100 and 150 W) in cold configuration, with or without the presence of IR source, and in hot configuration as reported in Table 1. The samples are denoted Ti_X_Y_Z or TiOx_X_Y_Z (for metallic and reactive sputtering, respectively) where X stands for sputtering configuration (cold or hot), Y magnetron power and Z indicates if IR source was (or not) ignited and its corresponding temperature. Cathode voltage is reported for most conditions.

Rutherford backscattering spectroscopy (RBS) was used to compare the evolution of Ti and O amount in thin films deposited on Si substrates in the three configurations. The energy of probing 4 He particles and scattering angle were fixed to 2 MeV and 165°, respectively. Film cross-section (for morphology, and thickness analysis) was observed by scanning electron microscopy (SEM, Zeiss Supra 40). Structural characterizations were performed by X-ray diffraction in Bragg-Brentano geometry on a diffractometer (Bruker D8 Discover) using Cu K radiation (0.15406 nm). XRD patterns were compared to JCPDS files corresponding to anatase TiO2 (00-021-1272 JCPDS files), rutile TiO2, hexagonal α-Ti (00-044-1294) and cubic β-Ti (01-074-7075). Results obtained on the 16 thin films will be presented and discussed. ** Temperature measured with bichromatic pyrometer focused on racetrack area (r=13 mm from target center).

Results and discussion

Target temperature evolution in hot configuration in metallic regime

This section focuses on the estimation of target temperature in hot configuration. First, a simple model based on energy balance at target is presented. Spatial distributions of the target surface temperature and directional emissivity are measured using bichromatic pyrometer. Last, mean target surface temperatures measured at steady state are compared to values calculated from model.

The control of process requires to precisely know target temperature as function of the power injected into the discharge. This is the reason why the measured temperature of target (T) is compared to values calculated from a 0D stationary model based on Bleykher et al [START_REF] Bleykher | Magnetron deposition of coatings with evaporation of the target[END_REF][START_REF] Bleykher | Energy and substance transfer in magnetron sputtering systems with liquid-phase target[END_REF][START_REF] Bleykher | Evaporation factor in productivity increase of hot target magnetron sputtering systems[END_REF]. This model takes into account all contributions: power delivered on target by ion bombardment (Wion), and power evacuated by heat radiation (Wrad), evaporation (Wevap) and sputtering processes (Wsputt), power due to the flux of reflected ions and to the gas convection being negligible. Target temperature is derived from the energy balance equation in stationary operation mode where Wion Wrad, Wevap, Wsputt represent the respective powers (expressed in W):

𝑾 𝒊𝒐𝒏 = 𝑾 𝒓𝒂𝒅 + 𝑾 𝒆𝒗𝒂𝒑 + 𝑾 𝒔𝒑𝒖𝒕𝒕 (2)
Wion and Wsputt can be estimated as function of discharge parameters whereas Wrad, Wevap mainly depends on temperature, material emissivity, saturated vapour pressure, the specific heat of sublimation and the specific heat of condensed matter.

𝑾 𝒊𝒐𝒏 = 𝑺 𝑬 𝒊𝒐𝒏 𝑭 𝒊𝒐𝒏 = 𝑷(𝟏 -𝜸)(𝟏 -𝜷) (3)
Where S corresponds to target surface (πR² in cm²), Fion is the ion flux density impinging onto target (ion/cm²/s), Eion is mean energy of the ions (J/ion), γ is secondary electron emission (estimated to 5% here for titanium metal and for argon ion energies in the range of 200-500 eV [START_REF] Anders | High power impulse magnetron sputtering: Current-voltagetime characteristics indicate the onset of sustained self-sputtering[END_REF]52]) and considered unchanged with temperature, 𝜷 represents the energy loss of argon ions in cathodic sheath (estimated to 40% for an argon pressure of 1 Pa [START_REF] Bleykher | Energy and substance transfer in magnetron sputtering systems with liquid-phase target[END_REF] and consistent with collisional plasma sheath model [START_REF] Sheridan | Collisional plasma sheath model[END_REF]). In our study, we consider a homogeneous ion flux over the whole target surface. But this is not the case in reality and the distribution of the current density should be taken into account in order to have a more precise calculation of the target temperature. P is the DC power delivered by power supply.

𝑾 𝒓𝒂𝒅 = 𝑺 𝜺 𝝈 𝑻 𝟒 ( 4
)
Where ε is the total emissivity of target surface (integrated on all values of λ, Ѳ and T), σ is Stefan -Boltzmann constant. Losses at the rear side of target are neglected, we only consider the front side of the target (S).

𝑾 𝒆𝒗𝒂𝒑 = 𝑺 𝑭 𝒆𝒗𝒂𝒑 𝜟𝑯 (5)
Where

𝑭 𝒆𝒗𝒂𝒑 = √ 𝟏 𝟐𝝅𝒎𝒌𝑻 (𝑷 𝒔𝒂𝒕 -𝑷′), 𝜟𝑯 = 𝑳 𝟎 + 𝟑 𝟐 𝒌𝑻 -𝒄 𝒑 𝑻 and 𝑷 𝒔𝒂𝒕 = 𝒆𝒙𝒑(𝟐. 𝟑𝟎𝟑(𝑪 - 𝑫 𝑻 ))
Where Fevap is the flux density of metal atoms leaving target (at/m².s) according to Hertz-Knudsen law, 𝜟𝑯 is the difference of enthalpy between the two phases (vapour and condensed state) (J/at), 𝟑 𝟐 𝒌𝑻 is the mean energy of atoms leaving target (J/at) for monatomic gas, k is the Boltzmann constant, m is the mass of evaporated atom (kg), L0 is standard heat of sublimation per atom (7.23 10 -19 J/at) at 0 K, cp is the specific heat of condensed matter (in 4.3 10 -29 J/at.K) close to 3k according to Dulong et Petit law, Psat is saturated vapour pressure at surface temperature Т (Pa), P' is vapour pressure at target surface temperature Т including the pressure of ejected metal and working gas (both neglected here compared to Psat), the 2.303 multiplier is used to convert the base-10 logarithm to the natural one, C and D are metal coefficients determining the saturated vapor pressure of target material in the Hertz-Knudsen equation. Different values of these coefficients can be found in literature [START_REF] Blocher | Vapor Pressure of Titanium[END_REF][START_REF] Dushman | Scientific Foundations of Vacuum Technique[END_REF]. In this study, mean values of 10 and 23000 have been taken for C and D, respectively.

𝑾 𝒔𝒑𝒖𝒕𝒕 = 𝑺 𝑭 𝒔𝒑𝒖𝒕𝒕 𝑬 𝒔𝒑𝒖𝒕𝒕 = 𝑺 𝒀 𝑭 𝒊𝒐𝒏 𝑬 𝒔𝒑𝒖𝒕𝒕 (6)
Where From equation ( 2), the temperature of Ti target can be estimated assuming that there are no conduction losses between target and magnetron (rear side and target area in contact to the clamping ring). Such conduction losses could induce a decrease of steady state temperature and eventually modify its uniformity (thermal gradient along the radius). In hot configuration, since thermal conduction is intentionally impeded at rear side, the first assumption sounds reasonable. On the contrary, it seems likely that thermal losses occur at clamping ring. Figure 2a displays the measured temperature of Ti target surface for 100 and 200 W (in pure Ar) along diameter (origin at target center). The temperature measured at target center is 860 °C for 100 W, and a temperature decrease of 150°C along a 8 mm distance is observed to target edge (corresponding to thermal gradient of 2 10 4 K/m). The mean temperature calculated on the whole circular surface is 788°C as reported on graph (dashed line). The temperature decrease is equal to 300°C for 200 W (thermal gradient of 4x10 4 K/m) with a temperature of 1130°C at center and a mean one of 1058°C. In the case of highly eroded targets, we have observed that temperature gradient between center and outer limits becomes higher due to the decrease of thermal conduction in the racetrack area. Thus, the temperature of target centre is higher for an used target than for a fresh one.

The directional emissivity in the 2-2.5 µm spectral range εpyro is reported on Figure 2b along target diameter in pure Ar. εpyro is minimum in the racetrack area between |6| and |16| mm and reaches minimum values of 0.14 at 200 W and 0.17 at 100 W, but largely increases in the outer limit of target to reach values above 0.3. A small dump (at 0.2 value) in the target is observed at center for 100W. These spatial evolutions of εpyro are due to differences in the physical state of target: very smooth surface in sputtered area, but rough and partly oxidized surface at center and around edge. The emissivity averaged over the whole target surface < εpyro> (dashed lines on graph 2b) is found to decrease from 0.26 to 0.22 between 100 and 200 W, as reported on the graph. This could appear surprising at first glance (ε generally increases with temperature for metals [START_REF] Setién-Fernández | Spectral emissivity of copper and nickel in the mid-infrared range between 250 and 900 °C[END_REF]), but is related to the physical state of target that may not be the same, depending on sputtering power. Figure 3a is obtained by progressively increasing (by steps) power between 50 and 115 W, whereas on Figure 3b power remains unchanged over time. It is clearly seen that when plasma is lighted on, or when power is increased, the target temperature rises until it reaches a steady state. For both figures, the steady state temperature evolution is extrapolated by using the empirical function T(t)=𝑇 ∞ (1 -𝑒 -𝑡/𝜏 ) where the two fitted parameters are τ, a time constant depending on the emissivity and the specific heat of the target (in particular), and 𝑇 ∞ , the achieved steady state temperature, as displayed in dashed line on Figure 3a and3b. The temperature evolution obtained with the model is reported on Figure 3c (for a total emissivity ε equal to 0.30, 0.38 and 0.40) and compared to the steady state mean temperature T measured with pyrometer. This mean temperature over the whole target has been evaluated considering a temperature decrease of 7±2% from the temperature measured at r ~ 13 mm. This is the percentage that we observed when comparing Figure 2a and3b. Figure 3c shows that model with ε = 0.38 and experimental temperature measured by pyrometer are in good accordance for power up to 100 W. This total emissivity value is higher than <εpyro> value measured by the pyrometer (0.22 at 200 W and 0.26 at 100 W) which can be due to the angular and wavelength dependence of emissivity. This shows that target is not a real diffuse grey-body as we previously assume. But the difference can also come from the emission of IR radiation at the rear side of target which would modify equation ( 4) by increasing the real emitting surface and so decreasing ε value. For 150 and 200 W, the experimental temperatures are above the modelling curve obtained for ε = 0.38 which confirms that emissivity decreases when its temperature increases or when its surface roughness decreases (both parameters being connected to the power). In addition to the temperature, this model allows to quantify the contribution of each energy loss mechanism. For a power ranging from 50 to 200 W, the mean target temperature is inferior to 1100°C and the IR radiation power density Wrad represents 99% of the total incoming power density Win.

Effect of oxygen incorporation on the target temperature and emissivity

In reactive mode, O2 introduction modifies the properties of target surface, its emissivity and so its cooling by radiation can drastically change. Figure 4 displays cathode voltage evolution, its temperature (in the racetrack) and its emissivity εpyro (in the 2-2.5 µm range) measured by the pyrometer facing the target at a DC power of 50 W for increasing and decreasing oxygen flow rates (0.1 sccm min -1 ). Note that all temperature values and emissivities reported on this Figure are instantaneous measurements and do not correspond to steady state values for a given oxygen flow rate in steady state. The absolute value of voltage increases from 275 to 370 V when cathode goes through the metal/compound transition. Meantime, εpyro emissivity increases from 0.19 in pure argon at 50 W (which is slightly higher that the directional one obtained on Figure 2b at 100 W) to 0.30 corresponding to a 1.6 increased factor. Since oxides exhibit higher emissivity than metals, this trend was expected. At 100 W, pyrometer gives an emissivity of 0.17 in racetrack position whereas the total emissivity deduced from model is about 0.38 corresponding to a 2.2 increased factor. Consequently, the total emissivity in oxide regime is probably higher than 0.30. The spread between the way in and the way back result from the different dynamics of oxidation and thermal balance of the target, both processes evolving with its proper time constant. The minimum observed on the temperature curve around 1 sccm on way back, which corresponds strictly to the bump in emissivity at same flow, is interpreted as resulting from the very slow thermal response of target. Moreover, it is well-known that de-oxidation process does not occur in the same way than oxidation route, especially when competing with sputtering. Consequently, target surface state (oxidation degree, roughness etc.) at a given oxygen flow rate may not be the same on way in and back.

In addition, the target temperature and emissivity exhibit opposite trend. This is easily explained by the fact that reached temperature depends on the ability of target to be cooled by IR radiations, which is directly connected to the emissivity value. The target temperature decreases from 640°C to 580°C (by a 1.1 factor) when turning from metallic to oxide mode, whereas the emissivity increases from 0.19 to 0.30 (by a 1.6 factor). This indicates that the power released by radiation is rather constant (1.1 4 ~1.5) at 50 W. The above results show how it is challenging to control target surface temperature in hot configuration.

Actually, the obtained value at steady state is mainly determined by the absorbed power by target and its emissivity, which is spatially inhomogeneous and target-history dependent. Nevertheless, the experimental set-up developed in this work and the simple model used allows to get a good estimation of the target temperature.

Energy distributions of Ti atoms and Ti + ions by mass spectrometry

The energy distributions of Ti atoms and Ti + ions are measured and compared in hot/cold configuration on Figure 5 using mass spectrometry in pure argon. Only the distributions obtained at 150 W are presented. The maximum of the atomic distributions (a) is obtained for about 3.5 eV (~ 0.72 Us) and seems to be unchanged between both configurations. Absolute intensity is increased in hot configuration, confirming the increase of deposition rate when target temperature rises. We do not observe any energy shift (to lower and higher value) or the appearance of Ti population group at low energy, which could have corresponded to evaporated atoms (3/2 kT = 0.15 eV for T=1150°C). Nevertheless, this does not rule out the existence of evaporation process since low intensity / low energy contributions may not be detected. Other techniques, as laser induced fluorescence, could be more relevant for this observation. As shown on Figure 5a, Thompson distribution largely overestimates the measured distribution for high energies due to collisions in gas phase at 1 Pa.

Stepanova et al [START_REF] Stepanova | Anisotropic energies of sputtered atoms under oblique ion incidence[END_REF] corrected this distribution by introducing a multiplicative factor (1-E/Emax). Both distributions on Figure 6a are well fitted by the Stepanova function with a cutting energy Emax of 25 eV for cold configuration and 20 eV for the hot one. The tail of both distributions at high energy can be also adjusted by a Maxwell-Boltzmann distribution with a kT constant of 4.5 and 4 eV for cold and hot configurations, respectively. These results show that a high target temperature slightly decreases the kT constant (or the cutting energy Emax) which leads to decrease the proportion of high energy sputtered atoms and so to a lower mean energy of the sputtered Ti species. This mean energy is equal to 6.9 eV for the distribution obtained in cold configuration and 6.35 eV for the one obtained in hot configuration. This decrease may be related to the decrease of the binding energy in materials when their temperature increases or other physical mechanisms as the ejection of clusters and the modification of cascade collisions process inside the materials. On Figure 5b, the Ti + ions energy distributions is composed of a low energy population between 0 and 4 eV and a high energy population up to 30 eV. As for the Ti distribution, absolute intensity is increased in hot regime, especially between 5 and 20 eV; but the amount of high energy ions seems to decrease in hot regime. The high energy tails are found to correspond to kT constant of 3 and 2.5 eV for respective cold and hot regime, which confirms the trend evidenced on Ti atom distributions. density coming from the gas (Фgas) and density coming from the energetic plasma species (Фplasma ()).

The measurement of energy flux density, ФIR,target and Фplasma, has been performed in all the conditions described in Table 1 with the HFM surface perpendicular to target surface ( geometry) and for few conditions with the HFM surface parallel to target surface (// geometry). Figure 7a displays the density evolution of energy flux due to plasma species in  geometry, denoted Фplasma (), as a function of the power applied to the cathode in pure Argon or in an Ar/O2 mixture. Фplasma, is proportional to the power delivered by the power supply from 0 to 150 W and the function expression is reported on the graph. The flux is 25% lower in oxide regime (light blue and orange symbols) than in metallic one (blue and red). In addition, hot configuration induces a small flux reduction of about 10% as compared to cold configuration in both cases (metallic and oxide regimes).

This may be related to the decreased amount of high energy species (see EQP results) and to the production of low energy species. geometries. The good agreement, which can be noticed between the absorbed IR radiations emanated from the hot target (red symbols) or the IR source (blue symbols), proves that this setup efficiently mimics the radiation coming from the hot target at a given temperature and will permit to isolate the effect of IR radiation on the substrate.

In // geometry, fitting curves are obtained with A equal to 0.045 and 0.075 for metal and oxide regime, respectively (where A is given by equation ( 1)). Considering an α absorptivity of 0.5 for sensor surface and a view factor of 3.4 10 -3 in // geometry, a ε emissivity of 0.36 in metallic and 0.6 in oxide regime is deduced.

Both emissivities are largely superior to the directional one measured by the bichromatic pyrometer but are consistent with the total one of target surface deduced from the model for metallic regime (0.38) and to the expected range in oxide regime. In  geometries, A values obtained from the fitting curves (Figure 7b and7c) is equal to 17 10 -3 and 6.5 10 -3 for metallic and oxide regime, respectively. If a value of 0.5 is used for absorptivity, a ε total emissivity of 0.36 in metallic is deduced, which is in accordance with the measurement performed in // geometry. In oxide regime, the thickness of the Ti film, deposited prior the TiOx one, is too thin to guarantee that it acts as a good IR absorber and the copper absorptivity is much lower that the titanium one. Indeed α is inferior to 0.1 for Cu in the 1-5 µm range using the Openfilter software based on refractive index and the extinction coefficient values from [START_REF] Edward | Handbook of optical constant and solids[END_REF].

These measurements performed in both geometries give the upper and lower limits of energy flux density on the substrate in the deposition configuration (45° from the target surface). For metallic and oxide mode and assuming again that emission is Lambertian, the power density received on the substrate Фir,target (45°) is obtained by multiplying the power density obtained in // geometry Фir,target (//) by the F45°/F// ratio (0.9). Фir,target (45°) can also be obtained by multiplying the density obtained in  geometry Фir,target () by the F45°/F (2.5) in metallic regime. Results is given in dashed line on Figure 7b and7c. The comparison of graphs (a) and (b) confirmed that the energy flux density coming from the IR radiations emitted by the hot target is much higher than the one brought by plasma species in metallic conditions. For instance, at the substrate position (45°) and 100W corresponding to 870°C in metallic regime, the plasma contribution is around 75 mW/cm 2 and that of the IR radiation close to 400 mW/cm². At this power, 85% of the energy flux density arriving on the growing Ti thin film is due to IR radiation in metallic mode and this percentage increases with power/temperature. In oxide regime and for the same power, the plasma contribution is equal to 55 mW/cm² (22 x 2.5) and the IR radiation are also close to 400 mW/cm², which means that 88% of the energy flux density arriving on the growing TiOx thin film is due to IR radiation in oxide mode. We must keep in mind that all these energy flux densities correspond to the absorbed part by sensor but not to the energy flux densities emanating from target.

Sputtering process and deposition rate

The thickness of each thin film has been measured on SEM images and the deposition rate, expressed in nm/min, has been calculated and reported in Figure 8 in pure Ar (a) and Ar/O2 mixture (b) . These values (Δ symbols) are compared to the deposition rate obtained with quartz microbalance (o symbols). The deposition rate obtained using the microbalance, expressed in µg/min, are reported in nm/min on Figure 8a assuming a Ti density of 4.5 g/cm 3 and in µg/cm² on Figure 8b as TiOx density is not known. In our setup, the quartz microbalance cannot be placed at the same location as the substrate during deposition (120 mm from target instead of 45 mm in deposition configuration) and so the values of the deposition rates measured by microbalance and by SEM cannot be directly compared. However, the evolutions with respect of the DC power applied to the cathode and differences between metallic / oxide modes and cold / hot configurations can be discussed. Clearly, all the trends displayed in Figure 8 are in good agreement for both measurement techniques, which indicates that the density of films does not significantly evolve in the investigated ranges of parameters.

In metallic mode and on the explored power range, the deposition rate linearly increases with rising power for cold and hot configuration in metallic mode and only for cold configuration in oxide mode. In this last case, in hot configuration, d exponentially rises (b, red marks). Such trends and behaviour in oxide have been already observed in literature and Musil et al [START_REF] Musil | High-rate reactive deposition of transparent SiO2 films containing low amount of Zr from molten magnetron target[END_REF] use two straight lines (representing the sputtering and the evaporation mechanisms) as fitting curves in oxide regime. Moreover, the values of deposition rate are found higher in the hot configuration than in the cold one. At 150 W, the deposition rates calculated from SEM cross sections reach 70 nm/min and 14 nm/min for respective metallic and oxide modes in hot configuration, whereas 50 and 7 nm/min are obtained in cold configuration. This corresponds to an increased factor of 1.4 and 2. This first value for metallic regime is relatively low compared to previous results obtained by Bleykher et al for Titanium [START_REF] Bleykher | Evaporation factor in productivity increase of hot target magnetron sputtering systems[END_REF], but the specific power was much higher in their study (up to 40 W/cm²) than one used here (up to 7.5 W/cm²). Both behaviours (enhanced deposition rate in hot configuration and higher increase in oxide regime)

are not fully explained at that time. At such low specific power, evaporation mechanism is negligible according [START_REF] Kozin | Modeling of thermal processes in magnetrons with single hot target and "sandwich-target[END_REF] and the EQP measurement supposing that low energy species are detectable. Our model indicates that the flux of evaporated atoms is 6 10 14 at/cm².s, whereas the one of sputtered atoms is 10 17 at/cm².s at 150 W (and 1010°C), as shown on Figure 8c. As sputtering is apparently the only ejection mechanism, this deposition rate increase could be related to a decrease of the binding energy or apparition/modification of other physical mechanisms. The addition of ceramic disk in hot configuration leads to an increase of the distance between the magnets and the target surface which modifies the discharge parameters (an increase of the cathode voltage in hot configuration is observed in Table 1 or for Ni in [START_REF] Caillard | Effect of the target temperature during magnetron sputtering of Nickel[END_REF]) and may change the distribution of magnetic field lines. This can also be an explanation of the change of the deposition rate. Finally, the increase of deposition rate could also be related to the modification of sputtering yield between hexagonal α and cubic β Ti phases.

Finally, on Figure 8c is reported the deposition rate measured on SEM cross sections depending on the temperature of the additional IR source for metallic and oxide modes. The graph shows that the temperature of IR source does not influence deposition rate without (blue) or with O2 (orange). This result proves that IR radiations emitted by the cathode would certainly have no influence on condensation mechanism and that the enhancement of deposition rate in hot configuration (see Figure 8a andb) is rather due to a modification of atom ejection efficiency from hot target. A mean value of 45 nm/min is found for Ti thin film and 2.4 nm/min for TiOx at 100 W. The lower value observed for TiOx films (more than a ten factor) was expected and is typical of deposition carried out in oxide mode. corresponds to the metallic thin films obtained in pure Ar in the same 3 configurations. These 6 thin films, whatever the configuration and the processing gas are, present a columnar morphology which is typical for thin film obtained by DC magnetron sputtering at this pressure (1 Pa) with tilted columns coming from the 45° deposition angle. No clear differences between the 6 first images are observed even if the structure appears more disordered (higher porosity, non-straight columns) in metallic thin films than in TiOx ones. The hot configuration or the cold one with IR radiation does not modify thin film morphology for these deposition parameters and more specifically for 100 W. However, for higher power, here 150W, the morphology of the metallic thin film depends on the configuration as shown on Figures 9g) and9h). Ti_Hot_150 thin film (h) seems to display a more ordered structure (than that of all other deposits and notably the Ti_Cold_150) and its columns appear faceted and sharp-ended, whereas "smooth" ones are observed in the other deposition conditions. This change in column features from domed tops to sharp ends is typical of growth mode evolution from zone 1 to zone T [START_REF] Petrov | Microstructural evolution during film growth[END_REF][START_REF] Barna | Fundamental structure forming phenomena of polycrystalline films and the structure zone models[END_REF] or zone 1a to zone 1c [START_REF] Mahieu | Biaxial alignment in sputter deposited thin films[END_REF] according to the well-known structure zone models. This is usually related to an increase of the deposition temperature and is a proof that hot target configuration influences the mobility of adatoms on the surface.

Morphology of thin films

However, this modification in the structure zone models generally goes with a thin film densification as it has been reported by Sidelev et al for chromium thin films [START_REF] Sidelev | A comparative study on the properties of chromium coatings deposited by magnetron sputtering with hot and cooled target[END_REF][START_REF] Grudinin | Chromium films deposition by hot target high power pulsed magnetron sputtering: Deposition conditions and film properties[END_REF][START_REF] Caillard | Effect of the target temperature during magnetron sputtering of Nickel[END_REF] and a V-shaped column appearance. This is not the case here where thin film porosity seems to be enhanced in hot configuration. This higher porosity, typical of thin film obtained by vacuum evaporation, could be related to the decrease of the energy of condensing species evidenced by mass spectrometry even if evaporation mechanism can be neglected in our study taken into account the relative low specific power and low temperature.

Oxygen content in films

Figure 10 displays the RBS spectra for titanium (a) and titanium oxide (b) thin films obtained at 100 W in the three configurations: hot (red), cold (blue) and cold with IR source (orange). On these graphs, the Ti peaks (between the 900 th and 1250 th channels) become larger in hot configuration (red) especially in Ar/O2 mixture, whereas deposition time was unchanged. This confirms that larger thicknesses, and consequently deposition rates, are obtained in hot configuration, and that IR radiations do not impact condensation process. The small widening of the Ti peak observed in the case of cold configuration and IR source in metal mode (Figure 10b, orange curve) is not significant. On Figure 10a (TiOx thin films), it can be seen that the maximum intensity of Ti peak decreases from 0.5 to 0.45 (as shown in the insert) when comparing the cold configuration to the two other configurations (hot and cold with presence of the IR source). This indicates that the deposition assisted by IR radiations (coming from the heated target or the IR source) enhances oxygen incorporation into thin film. The O peak intensity (between the 300 and 700 channels) is not clearly modified but this is not surprising due to the low sensibility of RBS signal at low channels. On Figure 10b, the highest Ti peak corresponds to the thin film realized with IR source. If we assume, as suggested by the results obtained in Ar/O2 mixture, that the incorporation of O is furthered in hot titanium, this trend can be explained. In pure Ar sputtering, residual oxygen is mainly consumed by the hottest surface in the vacuum experiment, i.e. the surface of the IR source that has been coated with Ti during previous runs of experiments. Additional IR source helps to consume residual oxygen into the chamber and so acts as an efficient O trapping device, in this case.

In order to obtain the absolute values of oxygen amount in these thin films, the spectra have been fitted using the Simnra® software, with Ziegler/Biersack stopping power data and Rutherford cross sections. In oxide mode, stoichiometric TiO2 thin films are obtained in hot configuration or with the presence of the IR source, whereas in cold configuration under-stoichiometric thin films are synthesized (typically TiO1.9). In pure Ar, when no oxygen is introduced in vacuum chamber, the oxygen percentage in the Ti-based thin film reaches large values: 25% in cold and hot configuration for 100 W. This high level of oxygen contamination inside Ti films is correlated to the very high affinity of this metal with oxygen. Of course, this percentage highly depends on the power applied on magnetron: 50% for 50 W, 25% for 100 W and 20% for 150W, since increasing the sputtered atom flux leads to higher Ti/O ratios inside films. As mentioned above, the addition of IR source allows to decrease percentage down to 15% for 100 W. All thin films obtained in pure argon are denoted "metallic Ti" film for convenience in the following section even if oxygen is present in these films. 3.8. Structural properties of films.

Films obtained in oxide mode

By XRD, we have evidenced that the obtained TiOx films are either amorphous or constituted by anatase TiO2, no rutile or brookite phases being observed whatever deposition conditions. In literature, the stabilization of anatase or rutile phase during magnetron sputter deposition at low substrate temperature has been reported to widely depends on deposition conditions: total pressure, oxygen percentage in gas phase, target to substrate distance etc. [START_REF] Toku | The effect of oxygen concentration on the low temperature deposition of TiO2 thin films[END_REF][START_REF] Musil | Low-temperature sputtering of crystalline TiO2 films[END_REF][START_REF] Mráz | Structure evolution of magnetron sputtered TiO2 thin films[END_REF]. Anatase phase is usually observed when energy transferred to the substrate during thin film growth is low [START_REF] Thomann | Diagnostic system for plasma/surface energy transfer characterization[END_REF]. In the present work, we focused our study on the comparison between cold and hot target configurations.

XRD analysis shows that all TiOx thin films obtained at 50 W with a thickness of 400 nm are amorphous closer to the ones corresponding bulk -Ti. Indeed, the "c/a" ratio is slightly inferior to 1.60 for Ti_Cold_100_IR and Ti_Hot_100 thin films.

Thus, it may be concluded from this part that the IR radiations coming from hot target or the IR source have an effect on the structural properties of Ti films. While the Ti_Cold_50 film is amorphous, Ti_Hot_50 and

Conclusion

We have investigated the sputtering of titanium in hot target configuration in pure Ar and in an Ar/O2 mixture by comparing three configurations: cold target, hot target (insertion of a ceramic disk between the target and the cooled magnetron) and cold target assisted by IR radiations. The effect of target surface temperature (up to about 1000°C) has been evaluated on sputtering process and on properties of Ti, and TiOx thin films for relatively low power (inferior to 150 W). This study confirms the trends reported in literature: the sputtering of a hot target induces an increase of deposition rate. In the present work, we found a 1.4 factor in metallic mode and a 2 factor in oxide mode for 150 W for which target temperature (measurement and model) reaches about 1000°C. The introduction of O2 in sputtering apparatus induces an increase of target emissivity and so a slight decrease of temperature.

No evidence of evaporation process, even at the highest temperature, has been found by mass spectrometry. A slight decrease of high energy Ti and Ti + species proportion has been detected using spectrometry leading to a decrease of their mean energy when incoming onto the substrate. The influence of the target temperatureespecially in reactive (oxygen) sputteringon the formation of oxygen radicals and negative oxygen ions near the target surface has not been studied but could be of interest. This suggests a modification of the sputtering process efficiency of heated material, leading to a rise of the sputtering yield. This can be related to the binding energy that is known to be lowered at high temperature and may also be the reason for the decrease of the high energy populations of the ejected species. However, this parameter cannot justify alone the deposition rate increase and other ones are probably affected by high temperature. Energy flux measurements have shown that the energy brought by plasma species slightly decreases with the rise of temperature. The important result is that the IR radiation emitted from hot target is the main part (about 80%) of the global energy transferred to the substrate in hot configuration. The influence of this contribution that has been suspected by several authors has clearly been found in the present work to determine final film properties.

All films are found columnar whatever the deposition configuration. Concerning the thin films obtained in Ar/O2 gas, no clear influence of IR radiation has been obtained on TiOx thin film microstructure, oxides being slightly denser than metals, as usually reported. The enhancement of oxygen incorporation in thin film is observed in oxide regime from RBS measurement. Working in hot target configuration seems to help synthesizing a nearly-stoichiometric titanium dioxide. This is consistent with an enhancement of the crystalline state of TiO2 anatase films characterized by crystallites randomly distributed on substrate (even if an enhancement of the ( 112) and (211) peaks is observed). Concerning the Ti-based film obtained in pure argon, no oxide phase is detected on these films, whereas a large amount of oxygen is revealed by RBS analyses. It comes that an amorphous TiOx phase is certainly present in addition to metallic Ti. Raman spectroscopy analysis could have been of interest to precise such microstructure. Hot configuration seems to enhance porosity of Ti-based film only at 150 W and induces columns exhibiting sharp-features. Moreover, crystallinity and crystallographic orientation are slightly improved and modified with a decrease of the compressive stress inside the film.

Finally, the IR radiations emitted in hot target magnetron sputtering are thought to be responsible for the modifications of thin film properties as compared to cold configuration.

Figure 1 .

 1 Figure 1. Experimental set up during deposition (a) or during in-situ diagnostic measurements (b) performed

  𝒀 = 𝟑 𝟒𝝅²𝑼𝒔 (𝟏 -𝜶)𝑼 𝒔 (𝟎. 𝟏𝟓 + 𝟎. 𝟏𝟑 𝑴 𝑻𝒊 𝑴 𝑨𝒓 )( 𝟒𝑴 𝑻𝒊 𝑴 𝑨𝒓 (𝑴 𝑻𝒊 +𝑴 𝑨𝒓 ) 𝟐 ), 𝑭 𝒊𝒐𝒏 = 𝑰𝒅 𝒒𝑺 (𝟏 -𝜸) and 𝑬 𝒔𝒑𝒖𝒕𝒕 = 𝑼 𝒔 /𝟐 Where Fsputt is the atomic flux density of sputtered atoms, Esputt is the mean energy of sputtered atoms, Y is the sputtering yield, MTi is the molar weight of Ti atom, MAr is the molar weight of Ar atoms, Id is the discharge current, U is the cathode voltage, Id is the cathode current, Us is the Ti binding energy (4.85 eV).

Figure 2 .

 2 Figure 2. (a) Evolution of the target temperature and (b) emissivity εpyro measured using a bichromatic

Figure 3 .

 3 Figure 3. Time evolution of the Ti target temperature (in the racetrack area, r ~ 13 mm) for powers ranging

Figure 4 .

 4 Figure 4. Target temperature, cathode voltage and emissivity εpyro as a function of the O2 flow rate for 50 W.

Figure 5 .Figure 6 .

 56 Figure 5. Ti (a) and Ti + (b) energy distributions obtained using EQP mass spectrometer in cold (blue) and hot

Figure 7 (

 7 Figure 7(b) and 7(c) shows the density evolution of the energy flux coming from the hot target ФIR,target or from

Figure 7 .

 7 Figure 7. (a) Energy density flux Фplasma () as function of the power applied to the cathode and (b,c) energy

Figure 8 .

 8 Figure 8. Deposition rate determined from SEM cross-section (Δ) and quartz microbalance (QMB)

Figure 9

 9 Figure 9 compares the SEM cross section images corresponding to oxide thin films obtained in an Ar/O2 gas

Figure 9 .

 9 Figure 9. SEM cross section images corresponding to the 3 oxide thin films (TiOx_Cold_100 (a),

Figure 10 .

 10 Figure 10. RBS spectra for Ti (a) and TiOx (b) thin films obtained in hot configuration (red), cold

  whatever the configuration. The diffractogram of TiOx_Hot_50 film corresponding to a target temperature of 660°C is shown on Figure11a. JCPDS data of tetragonal anatase TiO2 (# 00-021-1272) are reported on Figure11a(black dashed bars). When increasing power to 100W, the structure still remains amorphous as observed on the TiOx_Cold_100 (Fig 11a in blue). At this power, IR radiation emitted by target or by IR source seems to promote crystallinity, since the (101) diffraction line of TiO2 anatase phase is detected at 25.28°. By increasing power, the formation of polycrystalline tetragonal anatase phase is promoted as several peaks corresponding to the (004), (112), (200) (105) and (211) lattice planes are clearly visible on Figure11afor the film obtained at 150W. However, this conclusion has to be handled with care, because the thickness of TiOx_hot_150 is more than twice that of other samples. It is well established in magnetron sputtering technique that the crystallinity of a film is directly related to its thickness, reason why the study of effect of deposition conditions can only safely be performed on films of similar depth[START_REF] Mahieu | Biaxial alignment in sputter deposited thin films[END_REF]. To overcome this issue, a second set of 3 deposits TiOx_Cold_100, TiOx_Cold_100_IR and TiOx_hot_100 with the same thickness (1 µm) were performed by adjusting the deposition time (respectively fixed to 360, 360 and 180 min). Figure11bgives the evolution of the main (101), (004), (112), (105) and (211) diffraction lines. The 3 films present a powder-like diffractogram of the TiO2 anatase phase for which crystallites are randomly distributed on substrate; one can simply notice that the relative intensities of these 3 peaks are modified in hot target and cold + IR source configurations : the intensity of the (112) and (211) peaks is enhanced with the presence of IR radiation.

Ti_Cold_50_IRFigure 11 .

 11 Figure 11. X-ray diffractograms of (a) TiOx thin films with various thicknesses, (b) TiOx films with same

Table 1 .

 1 Deposition conditions for the 16 thin films obtained in the 3 configurations in pure Ar and in Ar/O2.

	Name	Power	Cathode	Target	IR source	Dep. time
		(W)	voltage	temp.	temp.(°C)	(min)
			(V)*	(°C)**
		In pure Ar (metallic regime)
	Ti_Cold_50	50	270	<250	-	10
	Ti_Cold_100	100	285	<250	-	10
	Ti_Cold_150	150	-	<250	-	10
	Ti_Cold_100_IR600°C	100	295	<250	600	10
	Ti_Cold_100_IR800°C	100	295	<250	800	10
	Ti_Hot_50	50	300	600	-	10
	Ti_Hot_100	100	330	870	-	10
	Ti_Hot_150	150	391	1010	-	10
		In Ar/O2 (oxide/reactive regime)
	TiOx_Cold_50	50	334	<250	-	180
	TiOx_Cold_100	100	373	<250	-	80
	TiOx_Cold_150	150	-	<250	-	80
	TiOx_Cold_100_IR600°C 200	369	<250	600	80
	TiOx_Cold_100_IR800°C 200	367	<250	800	80
	TiOx_Hot_50	50	384	600	-	180
	TiOx_Hot_100	100	422	800	-	80
	TiOx_Hot_150	150	476	1000	-	80
	* Voltage and power values correspond to the ones given by power supply.

Acknowledgements

The first author thanks the "Region Centre Val de Loire" for financial supports during his PhD. The authors thank Thierry SAUVAGE, Olivier WENDLING and Aurélien BELLAMY from the CEMHTI laboratory for RBS measurements on the Pelletron accelerator. All the authors would like to thank S. KONSTANTINIDIS, research associate of the National Fund for Scientific Research (Belgium) and N. SEMMAR, Professor at the University of Orléans, for their useful discussions about energy flux transfer (especially radiative transfer).

Energy flux density

The energy flux density measured by a thermopile sensor is given in Figure 6a for the TiOx_Cold_150 (blue) and TiOx_Hot_150 (red) deposition conditions with thermopile surface perpendicular to target surface ( geometry). Between 0 and 750 s, plasma is running, but the shutter placed in front of target is closed. Thus, the increase observed is only due to IR radiation from the heated shutter (up to 35-70 mW/cm² at 750 s for the cold-hot configuration). When shutter is opened (750 s), a large increase of energy flux density is observed in both cases. These values reach about 70 and 190 mW/cm² for cold and hot target configuration (blue and red, respectively). These overall energies correspond to the sum of several energy contributions coming from:

(1) the kinetic energy of species sputtered and condensed on sensor and energy eventually released/absorbed during crystallization of these species or during surface chemical reactions.

(2) the kinetic energy of eventual species evaporated and condensed on sensor and energy eventually released/absorbed during the crystallization of these species or during surface chemical reactions.

(3) the conduction/convection energy coming from gas atoms (acting as cooling), (4) the kinetic energy of high energy plasma species (neutralized and reflected species, negative and positive ions, secondary target electrons etc.)

(5) the eventual IR radiation emitted by the discharge [START_REF] Spaepen | Interfaces and stresses in thin film[END_REF] the IR radiation emitted by all vacuum parts placed in front of the sensor and especially the target.

As already explained in previous studies [START_REF] Caillard | Effect of the target temperature during magnetron sputtering of Nickel[END_REF][START_REF] Cormier | IR emission from the target during plasma magnetron sputter deposition[END_REF], the energy contribution corresponding to IR radiation can be measured by analysing the time recording of thermopile. When power is switched off (1321 s), all energy contributions coming from (1), ( 4) and ( 5) disappear in few ms and energy flux density quickly drops as shown on inset of Figure 6b. The (3) energy contribution measured just before starting the experiment (plasma off and experiment at room temperature) is equal to 8.2 mW//cm² in this work performed at 1 Pa [START_REF] Thomann | Energy flux measurements during magnetron sputter deposition processes[END_REF] and we assume that this value is unchanged during the experiment. After subtracting this low value, the energy flux density measured just after plasma extinction (1321 s) results from the addition of the (2) contribution corresponding to the eventual evaporated species and the (6) contribution corresponding to the IR radiation emitted by the heated target (all other vacuum parts are much colder). But the (2) energy contribution can be neglected in the study since EQP measurements have shown that no significant low energy population has been detected at 150 W.

Films obtained in metallic mode

The XRD patterns of these films present the diffraction lines of hexagonal α-Ti phase (Figure 11c). The film obtained at the lowest power in cold configuration (Ti_Cold_50) is amorphous, while all other films exhibit diffraction lines corresponding to -Ti bulk reference material (JCPDS #00-044-1294). Globally, same conclusions than before can be suggested: increasing power or adding IR radiations promotes crystallinity and slightly modifies the relative intensities of peaks. The crystallographic orientation of deposited Ti films changes from the (100) to (002) plane exaltation for 100 W. Inverse trend (from (002) to (100)) is generally observed when changing the DC process to a high power impulse magnetron sputtering (HiPIMS) process [START_REF] Jing | Titanium film deposition by highpower impulse magnetron sputtering: Influence of pulse duration[END_REF][START_REF] Britun | Ion density evolution in a high-power sputtering discharge with bipolar pulsing[END_REF]. Inverse changes found between these papers and our work show that increasing energy impinging onto the substrate by an increase of IR radiation or by an increase of ion energy does not play the same role and does not alter film structure in the same way. At the same time, XRD line broadening remains nearly unaffected, implying negligible changes in average grain size.

On Figure 11c, it is important to note that no oxide phase is detected on these films, whereas a large amount of oxygen is revealed by RBS analyses. It comes that an amorphous TiOx phase is certainly present in addition to metallic Ti. This is confirmed by the increase of the background signal intensity observed on XRD patterns at low angle diffraction, which can be related to diffusion process by amorphous matter.

Depending on experimental growth conditions, a shift of the angular positions of the (100) and (002) Ti planes, can be noticed. As reference, the (100) and (002) peak positions (dash bars) of bulk -Ti are given in Figure 11c ; they are located at 35.09° and 38.42°, respectively, and the corresponding "a" and "c" cell parameters are 0.2950 nm and 0.4683 nm (JCPDS # 00-044-1294). The "a" and "c" lattice parameters determined from the (100) and (002) peak positions on the XRD patterns (Figure 11c) are in the ranges 0.2911 -0.2928 nm and 0.4682 -0.4704 nm, respectively. Compared to reference values, the "a" parameter is lower while the "c" parameter is higher, meaning that compressive stresses exists in crystallites. In addition, the "c/a" ratios, which range from 1.59 to 1.62 in films, are higher than that of bulk -Ti (1.587). When increasing power up to 100 W, the positions of the (100) and the (002) are shifted to the corresponding positions of the bulk -Ti, especially for hot and cold+IR configuration. This means that the 2 cell parameters and the "c/a" ratio become


Prerna Chauhan
email: prerna.chauhan@savba.sk

S Hasenöhrl

A Minj

M P Chauvat

P Ruterana

J Kuzmík

Growth evolution of N-polar indium-rich InAlN layer on c-sapphire via strain relaxation by ultrathin AlON interlayer

Keywords:

X-ray spectroscopy High-angle annular dark-field

Introduction

N-polar (In,Ga)N-based heterostructures are attracting a remarkable amount of attention in the past years. The most promising application potential is expected for the high electron mobility transistors (HEMTs) [1], tunnel junction diodes [2] and optoelectronic devices, like solar cells [3] and light-emitting diodes (LEDs) [4]. On the contrary, investigation of its counterpart InAlN mainly revolved around achieving the lattice matched (LM) conditions to GaN buffer layer for the application in HEMTs. In such applications, it is used as a thin barrier layer for Ga-polar materials [5], and as a strong back-barrier for N-polar [6]. Other previously reported devices involving similar heterostructures were distributed Bragg reflectors [7], and intersubband photodetectors [8]. Recently [9,10], we demonstrated the epitaxial growth of metal-polar and N-polar InAlN buffer layers by organometallic chemical vapor deposition (OMCVD) that opened the possibility of growing InN channel transistor heterostructures on indium-rich InAlN buffer layer for THz applications, and optoelectronic devices as well [3,4]. Although significant advances have been made on devices utilizing LM InAlN/GaN heterostructure on metal-polar and N-polar materials, only a little attention has been given to N-polar indium-rich InAlN.

The vast majority of studies have explored strain-relaxation and indium-incorporation mechanisms on metal-polar [11] and N-polar [12] InGaN/GaN systems, covering the indium content range from low to moderate. These studies relate the strain relaxation mechanism and/or the associated effect of indium-incorporation to the formation of the misfit dislocations via slip system (mainly present in the low-indium content regime), alloy ordering, formation of V-defects, indium-rich interfacial regions and stacking faults.

On the other hand, understanding the growth of InAlN requires focused attention as it faces a separate set of challenges when compared to its counterpart InGaN. For example, it remains unclear why the most fundamental relaxation mechanism proceeding via slip system is not often triggered in strained InAlN [13]. Even, the intrinsic degradation utilizing the spontaneous formation of V-pits not associated with underlying threading dislocations is reported to occur in coherently grown mismatched and LM InAlN on GaN [14]. Furthermore, theoretical investigations also predict strong alloy immiscibility, i.e., a phase separation over a broad range of indium content of InAlN [15]. These issues indicate that stringent OMCVD growth conditions are required for the fabrication of high-quality indium-rich InAlN layer. Consequently, for its more profound understanding, a study of the growth evolution of indium-rich InAlN layer is necessary.

Recently, we reported the growth of a thick N-polar N layer directly on c-plane sapphire exhibiting a highly efficient red luminescence [10]. However, the evolutionary process of Indium-incorporation along the growth direction [0 0 0 1] of N-polar indium-rich InAlN layer and the local growth dy namics by which indium is preferentially incorporated during the subsequent InAlN layer growth on sapphire are presently unidentified. Our current work provides a significant insight into the growth evolution of N-polar indium-rich InAlN layer on sapphire. We systematically investigated the surface features by Kelvin probe force microscopy (KPFM). Subsequently, to achieve in-depth understanding, the cross-sectional analysis by high-resolution transmission electron microscopy (HRTEM) and scanning transmission electron microscopy (STEM) were employed to elucidate the effect of polarity vs. strain in indium-incorporation in the InAlN layer.

Experiments

Epi-ready c-plane or (0 0 0 1) sapphire ( ) substrate was nitridated under the ammonia ( ) flow of 3000 sccm at 1011 °C for 300 s and followed by the growth of AlN and InAlN layer at 730 °C for 30 s and 8440 s, respectively. The preprocessing and growth were performed by AIXTRON 3X2″ flip-top (FT) close coupled showerhead (CCS) reactor OMCVD system. Precursors used for In, Al, and N were trimethylindium (TMIn), trimethylaluminum (TMAl) and , respectively. Other details of the growth can be found elsewhere [10]. The resultant N-polar relaxed 334 nm thick layer was investigated in this study. Surface structural and electrical properties were studied using atomic force microscopy (AFM). The tapping surface imaging mode was utilized for quantitative statistical analysis of characteristics related to surface topography, compositional inhomogeneity, surface-exposed defects such as dislocations and/or V-pits and for any other anomalies. The nanoscale surface potential mapping was implemented by employing KPFM in the amplitude-modulation mode. In this mode, the surface topography simultaneously with the KPFM map was acquired. The KPFM measurements were done under a constant-amplitude mode at the first harmonic of the AFM probe cantilever (~55 kHz) using a tip AC bias of 3 V serving as the modulating electrostatic signal at ~346 kHz, the second harmonic of the cantilever. For the measurements, the sample was grounded, the tip-sample separation was close to 10 nm, the scan speed was set to 0.5 lines/second and Pt/Ir-coated Si-AFM probes with the tip radius of 25 nm were used. All the AFM based measurements were performed on CSI Nano-Observer AFM under ambient conditions.

The samples for TEM observations were prepared using a focused ion beam (FIB) machine. The weak beam analysis was carried out to analyze the crystallographic quality of the layer using a JEOL 2010 microscope. The energy dispersive X-ray spectroscopy (EDX) and high-angle annular dark-field (HAADF) investigations took place in a double Cs corrected JEOL 200F STEM. This allowed a spatially-resolved chemical analysis using EDX. In HAADF, the contrast is mainly given by the atomic number (Z-contrast) and is employed to distinguish between In (high Z = bright contrast) and Al (low Z = dark contrast).

Results

In Fig. 1 extended defects such as dislocations and V-pits [17,18]. In our case, the pits (~2.6-6.6 × ) or dislocations present at the periphery of grains (in Fig. 1(a)) appear with a contrast in KPFM map corresponding to a higher work function but strictly within the confinement of the pit. This, however, cannot be related to possible fixed charges or trapped electrons at the dislocations since a charged dislocation would result in the dispersive nature of the potential or CPD change around it of the type ~-with r being the radial distance from the dislocation core and R the radius of the space charge region containing dislocation core [19].This type of behavior is usually observed in InAl(Ga)N and even in InGaN where the V-shaped pits with high electron affinity over their semi-polar surfaces contribute to the locally enhanced work function [20].

In addition to grains and pits, localized growth features with an average diameter of 153 ± 26.3 nm and a peak height of 22 ± 7.9 nm (indicated by arrows in Fig. 1(c) and (d)) surrounded by partially-visible trenches or pits were also observed. These growth anomalies with a surface density of ~1 × estimated over 5 × 5 are associated with a much lower work function changes of only 30 meV with respect to the background (Fig. 1(d)). This drop in the work function could correspond to a local increase in indium content facilitated by increased growth rate or could also indicate the change in or even a reversal of surface polarity. The ab-initio calculations by Strak et al. [16] predict a lower magnitude of work function in metal-polar III-nitride surfaces in comparison to N-polar surfaces. Since the work function drop measured by KPFM is considerably smaller, several other factors contributing to the work function including charged defects cannot be fully excluded. Therefore, in order to have a more in-depth insight, these anomalies were further investigated by HRTEM and EDX analysis, which will be discussed in the next section.

Figs. 2 and 3 display the microstructure of the investigated In-AlN/sapphire heterostructure. Fig. 2(a) shows a weak beam dark-field image recorded using g = 0002, and Fig. 2(b) shows the corresponding bright-field image. The surface morphology is not planar due to the brighter feature in Fig. 2(a), which extends from the interface up to the surface. This feature originates most probably from a lighter material at the heterointerface level as demonstrated by the corresponding whiter contrast in Fig. 2(b). All along the interface, higher con trast is present, which decreases gradually within the first 40 nm of layer growth. Moreover, few screw-type threading dislocations (TD) are also visible in the InAlN layer. In Fig. 3, using g = 10-10, it is challenging to attain the weak beam conditions and the predominant contrast indicates mainly the presence of strain in the basal planes all along the layer without any possibility to distinguish individual defects, such as dislocations, stacking faults, or inversion domain boundaries [21,22].

This illustration gets more evident in the HAADF image (Fig. 4), where a similar feature is analyzed along a different projection direction and exhibits higher growth rate than the adjacent areas. The rapidly growing small features with the surface density ~1 × as discussed in previous KPFM section are the metal-polarity patches [23], which originate from Al-rich areas just at the interface (as shown in Figs. 234) in the N-polar layer [10]. Moreover, all over the heterointerface, irregular distribution of the darker contrast is clearly visible with an extension of around 40 nm. This indicates that an irregular Al-rich interfacial layer (dark contrast) due to the irregular chemical distribution may lead to the strain relaxation followed by an increase of the indium-incorporation in the subsequent homogeneous InAlN layer. This assertion is further supported by the bright contrast of homogeneous area which depicts more indium content. This is consistent with the transition and homogeneous regions observed in AES depth profile shown in Ref. [10]. The origin of the columnar formation is due to the faster growth of In/Al inversion domains (IDs) in N-polar matrix [24]. However, a possible contribution of screw dislocations generated at the interface during nitridation [25] cannot be completely excluded as these areas are highly disordered.

To analyze the growth evolution, a detailed investigation of the elemental distribution along the direction of growth and local preferential incorporation centers is essential. Fig. 5 shows a STEM-HAADF micrograph of InAlN layer. The indium-molar fractions of the layer were measured at six different places. Out of six EDX line scans, the first scan is parallel to the InAlN/sapphire interface and the remaining scans are along the direction of growth. To be more specific, the second and fourth EDX line scans are along the periphery of the columnar structure and the third is at its center. The EDX line scans in Fig. 5 show the probe position during EDX acquisition, where five points were ana lyzed and the corresponding molar fraction is shown. While being metal-polar predominantly observed the indium-molar fraction in columnar structure varies from 0.60 to 0.66 except for one point at periphery where the indium-molar fraction of 0.74 is observed. This could be ascribed to an enhanced incorporation of indium and formation of InN chains in the vicinity of dislocation core.

On the other hand, the indium-molar fraction in the N-polar region varies from 0.60 to 0.71. Close to interface the indium-molar fraction is close to 0.64 ± 0.04 as derived from XRD, which averages over the whole layer and all compositional fluctuations influence the FWHM. STEM-EDX analysis indicates that the range of the indium-molar fraction varies from 0.60 at interface to 0.74 close to surface. This EDX data confirm the local compositional inhomo geneities that lead to the broad width (FWHM = 0.20°) of 0002 diffraction peak of InAlN layer as shown in Ref. [10]. As demonstrated in Figure 5, the observed indium-molar fraction varied from 0.60 to 0.74 in InAlN layer. These inhomogeneities are presumably composed of nanoscopic indium-rich phases, unlike the macroscopic phases that are volumetrically large enough to be detected in 2θ/ω scan of XRD. For a more comprehensive analysis of the growth evolution, we widened up the investigations and focused on the InAlN/sapphire interface. Fig. 6(a) shows the HAADF image of the heterointerface. At the top of sapphire, a dark interfacial area of ~2 nm thickness is marked by the white arrows. The displayed contrast validates that the mean atomic number is smaller than that of InAlN and sapphire. This is further confirmed by the EDX analysis of the area as indicated by the green rectangle in Fig. 6(a). The corresponding spectrum displayed in Fig. 6(b) shows the oxygen (O) peak along with the expected nitrogen (N) and aluminum (Al) peaks. Thus, our analysis reveals the existence of crystalline AlON interlayer formed during 300 s long nitridation, which is described as the N-rich growth. Sapphire ( substrate is the source of Al and O [10]. The subsequent growth of low-temperature (LT) AlN interlayer weakens the diffusion of O atoms into AlN interlayer and the following InAlN layer.

In the HAADF image shown in Fig. 7(a), this intermediate layer is highlighted by two arrows where the white arrow underlines the sapphire surface, and the red arrow points to the bottom of the InAlN layer. In the corresponding filtered image (Fig. 7(b)), the interface dislocations are clearly visible and the misfit dislocations discernible at the termination of each atomic plane. At the AlON/sapphire interface, nine misfit dislocations are present and separated by an average distance of 2 nm as marked by the white asterisks in the observed image, whereas only three misfit dislocations (marked by the black asterisks) are present at the InAlN/AlON interface or at the bottom of the InAlN layer. This image shows the 8:10 relationship between AlON and sapphire, meaning that roughly eight Al-Al atomic distances in AlON correspond to ten Al-Al atomic distances in sapphire. Another existing relationship is 14:16. This is a clear indication that most of the plastic strain corresponding to the large lattice mismatch (~22%) between sapphire and layer is released at interface via the formation of misfit dislocations which is consistent with the first-principles calculations [26] and our previous XRD analysis [10].

Discussion

Before starting the discussion, we would like to briefly summarize our main experimental results. HAADF and EDX data show the presence of an intermediate layer ~2 nm, which is a mixture of Al-O-N. It can be seen that the interfacial layer is crystalline and has the same crystal structure as the top InAlN, and it carries out the strain relaxation with the two series of misfit dislocations. Most of the interface strain is released at the interface of AlON/sapphire with the formation of misfit dislocations.

The InAlN/sapphire interface is highly complex and can be envisioned as follows. In the initial stage of nitridation of sapphire ( ) at 1011 °C a two- dimensional layer of aluminum-oxynitride (

) forms [27]. That results in the N-polar AlN due to the dissociation of AlON. With the increasing duration of nitridation, the layer transforms into three-dimensional layer through the formation of Al-polar AlN islands. The layer grown after the nitridation is mainly but exhibits columnar structures via the formation of inversion domains at the sites of Al-polar AlN islands in AlON layer, as illustrated in Fig. 8(a). The islands as the inversion domains are the fast-growing metal-polar area inside the N-polar matrix and form highly disordered columnar structure (Fig. 8(b)). Moreover, the growth of InAlN layer under N-rich growth condition preserves the N-polarity established by AlON interlayer.

At a low growth pressure, the enhanced surface mobility of indiumand Al-adatoms and their stronger bonding with nitrogen surface atoms of the N-polar layer contribute to the enhanced indium-incorporation [28]. This indicates that after partial strain relaxation at the InAlN/sapphire interface via the AlON layer, in 40 nm wide transition region another factor that influences the indium-molar fraction of the layers is the polarity of InAlN layer. The strain relaxation and N-polarity accomplished the compositional grading in transition region, and finally a relaxed InAlN layer was formed but with local compositional fluctuation in the indium-molar fraction. Although, the weak interaction between the impinging indium-and Al-adatoms with the indium-and Al-surface atoms of metal-polar layer leads to the lesser indium-incorporation [28]. In this case under the used growth temperature of 730 °C, the indium-desorption has presumably influence on the layer growth.

The surface topography (Fig. 1) of InAlN layer confirms the Volmer-Weber growth mode that describes the nucleation and subsequent growth of distinct islands at the preferential sites provided by nitridation and/or AlN thin layer on sapphire. As suggested by Nix [29] and Sheldon et al. [30], coalesced crystallites impose displacements or tensile elastic strain at the grain boundary. Usually, the mode of relaxation is controlled by temperature, compositional grading or surfactants [31]. Thus, it is proposed that the strain relaxation mech anism is changed in the InAlN layer grown on sapphire because of the different surface kinetics of indium as a surfactant. The indium-surfactant effect on nitrides during OMCVD growth was investigated earlier theoretically [32] and experimentally [33]. When indium as a surfactant is involved in growth, an increase of the surface diffusion length that enhances the step-flow growth mode is expected. In the presence of fast-growing metal-polar inversion domains and/or screw dislocations at the interface, this may lead to spiral structure and consequently, wavy surface morphology that helps to release the elastic strain energy [31].

In our study, we observed that the compositional grading and surfactant effect are working concurrently in the transition region at the InAlN/sapphire interface. The nature of the strain during the subsequent layer growth substantially depends on the mobility of depositing atoms. As indium is larger in size than the host aluminum atom and sufficiently mobile during the growth in the used growth conditions here, it will bind to the energetically favorable tensile strained locations. The previous theoretical works [34] have shown that screw and mixed type dislocations are the preferential incorporation sites for indium. Grain boundaries, pits and dislocations act as atomic sinks and are required to relax the strain in the layer. To accomplish this relaxation, the arriving atoms should diffuse [29,30]. The atomic process by which the high-mobility indium atoms as a surfactant may preferentially incorporate into grain boundaries and pits during the subsequent growth of the InAlN layer is illustrated in Fig. 9 and maybe envisioned as follows. The tensile strain at grain boundaries in the initial stage of growth is due to the coalescence of the AlN low mobility islands. The high-mobility of indium atoms arriving at the growing surface facilitates the tensile strain relaxation in the transition region of InAlN layer growth. This elastic relaxation must be accompanied by the formation of pits [35] and/or surface roughening/undulation [36]. This relaxation promotes plastic relaxation via facilitating indium-incorporation in the layer. Once the layer is relaxed, the subsequent layer will also be relaxed. The InAlN growth temperature (730 °C) is high enough to desorb the accumulated indium and improve crystalline layer quality.

Conclusion

For indium-incorporation, besides the growth temperature, lattice latching and strain in the layer all work together, and the latter two are interconnected. The only solution to overcome these problems is the engineering of strain relaxation. This work evidences the formation of ~2 nm AlON ultrathin layer, which relaxes the interface strain in two steps at the InAlN/sapphire interface. The AlON layer leads to strain relaxation as well as the N-polar layer that further assists the indium-incorporation and provides a uniform InAlN layer. In this work, we have provided unambiguous evidence of the local fluctuation in the indium-molar fraction on nano-length scale by using HAADF-STEM analysis. The high-mobility larger-size indium atoms as a surfactant incorporate at the preferential sites, like grain boundaries, pits, and any high-potential energy sites during the growth of InAlN layer to reduce the energy and make it more stable.

Fig. 1 .

 1 Fig. 1. (a) and (c) Topography and (b) and (d) their corresponding KPFM images of the N-polar InAlN layer over the scan area of 5x5 .

Fig. 2 .

 2 Fig. 2. Microstructural features as pointed out by conventional TEM in N-polar InAlN layer: (a) Weak beam image recorded using g = 0002, a bright feature that stems from the interface to the surface. (b) The corresponding bright field shows surface pinholes as well as light contrast at the bottom of the bright feature of (a).

Fig. 3 .

 3 Fig. 3. Microstructural features as pointed out by conventional TEM in N-polar InAlN layer: (a) Weak beam image with g = 10-10. (b) The corresponding bright field image.

Fig. 4 .Fig. 5 .

 45 Fig. 4. HAADF image recorded along < > zone axis of N-polar InAlN layer on sapphire: the red arrow points to a columnar structure originated at an Al rich area (black feature). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

Fig. 6 .Fig. 7 .

 67 Fig. 6. (a) Cross-sectional HAADF image of the InAlN/sapphire interface recorded along < > zone axis of InAlN. A two-dimensional continuous and crystalline interlayer in darker contrast is visible on top of the sapphire. A green rectangle marked in the interlayer shows the location of EDX measurement area and (b) the corresponding EDX spectrum showing the only contributions of Al, O and N in the area. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

Fig. 8 .

 8 Fig. 8. Schematic illustration of In/Al-polar inversion domain formation mechanism in the N-polar InAlN layer. The colors and arrows indicate the polarity of the layer: blue color and upwards arrow-In/Al-polar; red color and downwards arrow-N-polar. (a) Sapphire ( ) annealing in 300 s results in evolution of stable Al-polar AlN islands along with layer. (b) During the further growth of InAlN layer, the islands promote metal-polar columnar structure, while the layer decomposes to give N-polar In-AlN. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

Fig. 9 .

 9 Fig. 9. Illustration of atomic processes during the growth of InAlN layer by which the high surface mobility indium-atoms may be preferentially incorporated into (1) grain boundaries via the surface diffusion, (2) pits or V-defects, and (3) dislocations. The blue color balls represent Al in AlON layer, while the red color balls represent In in InAlN layer. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

Acknowledgements

The authors would like to acknowledge the financial support granted by the Slovak Agency for R&D in the framework of project no. APVV-15-0031. The TEM samples were prepared using a FIB facility which has been acquired through the EQUIPEX GENESIS ANR-11-EQPX-0020 in the "Investissements d'avenir" national program with the contribution of the Region Normandie through the FEDER system.