EP4677135A2

EP4677135A2 - Tm-diboride based coatings exhibiting enhanced oxidation stability at high temperatures

Info

Publication number: EP4677135A2
Application number: EP24711161.0A
Authority: EP
Inventors: Jürgen RAMM; Oliver HUNOLD; Ahmed Adel Ibrahim BAHR; Helmut Riedl-Tragenreif; Peter Polcik; Szilard KOLOZSVARI
Original assignee: Oerlikon Surface Solutions AG Pfaeffikon
Current assignee: Oerlikon Surface Solutions AG Pfaeffikon
Priority date: 2023-03-08
Filing date: 2024-03-08
Publication date: 2026-01-14
Also published as: WO2024184534A2; WO2024184534A3

Abstract

The present invention relates to coating films based on TM-diborides exhibiting enhanced oxidation stability at elevated temperatures and a method for producing such inventive coating films. The coating films comprising in the as deposited state a transition metal diboride, preferably a quaternary metal diboride, the metal diboride formed by at least two transition metals, silicon and boron, wherein the coating film material exhibiting a AlB₂-structure type of TM¹B₂ and having a chemical composition in atomic percentage corresponding to TM¹ _aTM² _bSi_cB_d, wherein: • TM¹ = Ti or Cr, • TM2 = Ta or Mo, • 15 ≤ a ≤ 35, preferably 20 ≤ a ≤ 30, • 5 ≤ b ≤ 15, preferably 5 ≤ b ≤ 10, • 10 ≤ c ≤ 30, preferably 10 ≤ c ≤ 20, • 35 ≤ d ≤ 65, preferably 40 ≤ d ≤ 60.

Description

TM-diboride based coatings exhibiting enhanced oxidation stability at high temperatures

The present invention relates to coatings based on TM-diborides exhibiting enhanced oxidation stability at elevated temperatures, formed by addition of TM-disilicides.

The present invention relates also to a method for producing such inventive coatings.

In the context of the present invention the acronym TM is used for abbreviating Transition Metals.

Technical Field

TM-diborides of the type TiB2±z thin films are the subject of growing research interests and are considered as potential future protective coating material applied in aerospace, microelectronic, or cutting tool industry, as it can be seen in many scientific publications, for example as reported by:

• W.G. Fahrenholtz, G.E. Hilmas in the article Ultra-high temperature ceramics: Materials for extreme environments, Ser. Mater., 129 (2017) 94-99;

• C. Mitterer in the article Borides Thin Film Technology, Journal of Solid State Chemistry, 133 (1997) 279-291 ;

• J.S. Chen, J.L. Wang in the article Diffusion Barrier Properties of Sputtered TiB2 Between Cu and Si, Journal of The Electrochemical Society, 147 (2000) 1940;

• C.S. Choi, G.A. Ruggles, A.S. Shah, G.C. Xing, C.M. Osburn, J.D. Hunn in the article Stability of TiB2 as a Diffusion Barrier on Silicon, Journal of The Electrochemical Society, 138 (1991 ) 3062-3067;

• P.H. Mayrhofer, C. Mitterer, J.G. Wen, J.E. Greene, I. Petrov in Self-organized nanocolumnar structure in the article superhard TiB2 thin films, Appl. Phys. Lett., 86 (2005) 131909;

• M. Berger, S. Hogmark in the article Evaluation of TiB2 coatings in sliding contact against aluminium, Surface and Coatings Technology, 149 (2002) 14-20;

• M.N. Polyakov, M. Morstein, X. Maeder, T. Nelis, D. Lundin, J. Wehrs, J.P. Best, T.E.J. Edwards, M. Ddbeli, J. Michler in the article Microstructure-driven strengthening of TiB2 coatings deposited by pulsed magnetron sputtering, Surface and Coatings Technology, 368 (2019) 88-96;

• C. Fuger, R. Hahn, A. Hide, P. Kutrowatz, M. Weiss, A. Limbeck, 0. Hunold, P. Polcik, H. Riedl in the article Revisiting the origins of super-hardness in TiB2+z thin films - Impact of growth conditions and anisotropy, Surface and Coatings Technology, 446 (2022) 128806.

Within the interesting family of transition metal diborides (TMB2), TiB2±z exhibits an attractive aggregate of properties with high thermal stability (TM~3225 °C), superhardness (> 40 GPa), low density, but also good thermal and electrical conductivity accompanied by chemical inertness, as it can be seen in some scientific publications, for example as reported by W.G. Fahrenholtz et al and P.H. Mayrhofer et al in the articles already mentioned above, or also as reported by R.G. Munro in the article Material Properties of Titanium Diboride, J Res Natl Inst Stand Technol, 105 (2000) 709-720.

However, still one of the major obstacles against the wide applicability of TiB2±z-based films is their limited oxidation resistance, which is only lightly above 400 °C, as it can be seen in some scientific publications, for example as reported by:

• J. Thdrnberg, B. Bakhit, J. Palisaitis, N. Hellgren, L. Hultman, G. Greczynski, P.O. A. Persson, I. Petrov, J. Rosen in the article Improved oxidation properties from a reduced B content in sputter-deposited TiBx thin films, Surface and Coatings Technology, 420 (2021 ) 127353;

• S. Dorri, J. Palisaitis, G. Greczynski, I. Petrov, J. Birch, L. Hultman, B. Bakhit in the article Oxidation kinetics of overstoichiometric TiB2 thin films grown by DC magnetron sputtering, Corrosion Science, 206 (2022) 110493;

• B. Bakhit, J. Palisaitis, J. Thdrnberg, J. Rosen, P.O. A. Persson, L. Hultman, I. Petrov, J.E. Greene, G. Greczynski, in the article Improving the high- temperature oxidation resistance of TiB2 thin films by alloying with Al, Acta Materialia, 196 (2020) 677-689.

The oxidation behavior for TiB2±z bulk and thin film materials has been extensively studied as reported for example by J. Thdrnberg et al and S. Dorri et al in the above mentioned articles, and specific morphological and kinetic-related aspects have been analyzed. According to Cai et al. in the article Unmasking the anomalous rapid oxidation of refractory TiB2 at low temperatures, Journal of the European Ceramic Society, 41 (2021 ) 5100-5108, monolithic bulk TiB2 starts to oxidize at 400 °C, whereas a rapid, anomalous oxidation sets in at around 500 °C. The reported accelerated oxidation at 500 °C is related to the formation of an outer mixed amorphous/crystalline B2Os/TiO2 scale and an inner unstable Ti-B-0 layer. Between 650 to 1000 °C - also referred as the low-temperature regime - the scale formation changes to a laminated configuration with crystalline TiO2 and a glassy amorphous B2O3 providing a certain oxidation resistance. As it has been reported for example by Cai et al. in the article already mentioned above but also reported by:

• A. Tampieri, A. Bellosi in the article Oxidation of monolithic TiB2 and of AI2O3- TiB2 composite, Journal of Materials Science, 28 (1993) 649-653;

• R.A. Andrievskii, Y.M. Shul’ga, L.S. Volkova, 1.1. Korobov, N.N. Dremova, E.N. Kabachkov, G.V. Kalinnikov, S.P. Shilkin in the article Oxidation behavior of TiB2 micro- and nanoparticles, Inorganic Materials, 52 (2016) 686-693;

• X. Huang, S. Sun, G. Tu in the article Investigation of mechanical properties and oxidation resistance of CVD TiB2 ceramic coating on molybdenum, Journal of Materials Research and Technology, 9 (2020) 282-290.

At higher temperatures, the accelerated kinetics predominates the oxidation process forming volatile B2O3 accompanied by non-protective porous TiO2 scales as reported for example by Tampieri et al. in the article mentioned above, and also reported by:

• G.B. Raju, B. Basu, A.K. Suri in the article Oxidation Kinetics and Mechanisms of Hot-Pressed TiB2-MoSi2 Composites, Journal of the American Ceramic Society, 91 (2008) 3320-3327;

• T.A. Parthasarathy, R.A. Rapp, M. Opeka, R.J. Kerens in the article A model for the oxidation of ZrB2, HfB2 and TiB2, Acta Materialia, 55 (2007) 5999-6010.

However, this oxidation sequence is reported to be different for TiB2±z-based thin film materials. Huang et al. highlighted the formation of low viscosity B2O3 (I) phase, which becomes volatile at around 720 °C and results in a rapid oxidation of their chemical vapor deposited TiB2 coating. In contrast, the oxidation of physical vapor deposited (PVD) TiB2±z below 800 °C did not exhibit the formation of B2O3 (I), and the oxide scales have been reported to be Ti-rich, as reported for example by Thdrnberg et al, and Dorri et al in the articles already mentioned above. Thbrnberg et al. indicated that the B concentration (boron concentration) plays a prominent role on the oxidation kinetics of sputtered TiB2±z films. Their sub- stoichiometric TiBi.43 films exhibited lower oxidation rates compared to B-rich films, due to the absence of the rapidly oxidizing B-rich tissue phases. Nevertheless, all their investigated TiB2±z films follow linear-rate laws with low oxidation resistance lightly over 400 °C.

Different routes have been attempted to enhance the oxidation resistance of TiB2±z coatings mainly based on alloying with elements known as strong oxide formers such as Al (aluminum) as reported for example by Bakhit et al in the above mentioned article, and Si (silicon) - both Al and Si are prone to form protective oxide scales. These alloying routes have been reported in some scientific articles, for example as reported by:

• A.H. Navidi Kashani, S. Mraz, D.M. Holzapfel, M. Hans, L. Lofler, P. Ondracka, D. Primetzhofer, J.M. Schneider in the article Synthesis and oxidation behavior of Ti0.35AI0.65By (y = 1.7-2.4) coatings, Surf. Coat. Technol., (2022) 128190;

• J. Thbrnberg, S. Mraz, J. Palisaitis, F.F. Klimashin, P. Ondracka, B. Bakhit, P. Polcik, S. Kolozsvari, L. Hultman, I. Petrov, P.O. A. Persson, J.M. Schneider, J. Rosen in the article Oxidation resistance and mechanical properties of sputter- deposited Ti0.9AI0.1 B2-y thin films, Surface and Coatings Technology, 442 (2022) 128187;

• T. Glechner, H.G. Oemer, T. Wojcik, M. Weiss, A. Limbeck, J. Ramm, P. Polcik, H. Riedl in the article Influence of Si on the oxidation behavior of TM-Si-B2±z coatings (TM = Ti, Cr, Hf, Ta, W), Surface and Coatings Technology, 434 (2022) 128178;

• B. Grancic, M. Mikula, T. Roch, P. Zeman, L. Satrapinskyy, M. Gregor, T. Plecenik, E. Dobrocka, Z. Hajovska, M. Micusik, A. Satka, M. Zahoran, A. Plecenik, P. Kus in the article Effect of Si addition on mechanical properties and high temperature oxidation resistance of Ti-B-Si hard coatings, Surface and Coatings Technology, 240 (2014) 48-54.

Bakhit et al. in the article already mentioned above reported an improvement in the oxidation resistance of sputtered Tio.68Alo.32B1.35 with retarded kinetics at 800 °C for 0.5 h, due to the formation of a dense Al-based oxide scale of 470 nm compared to a scale thickness of 1900 nm obtained for their binary TiB2.4 counterpart.

Navidi et al. in the article already mentioned above followed the Al-alloying strategy to deposit Al-rich but nearly stoichiometric (Tio.35Alo.65)B2 films revealing an outstanding oxidation resistance at 700 °C by forming thin Al-based oxide scale of only 39 ± 7 nm after 8 h. The high Al-content predominates the oxidation behavior, but at the expense of the mechanical properties with reported hardness between 9 to 24 GPa. On the other hand, the influence of Si-alloying on the oxidation resistance of several TMB2±Z- based coatings was studied by Glechner et al. as reported in the article mentioned above. For Ti-Si-B2±z coatings, the Si addition provided high-temperature oxidation resistance with strongly retarded kinetics up to 1200 °C, while the reported hardness drastically decreased upon high-Si addition to around 16 GPa. Still, the above mentioned alloying routes by Al and Si to form oxidation-resistant ternary diborides are limited by the deterioration of the coatings’ mechanical properties. For bulk refractory TMB2, a different strategy is followed to enhance the oxidation resistance. Through the addition of secondary Si-based phases (i.e. SiC or TMSi2), which minorly influence the desired mechanical properties, highly protective glassy-like borosilicate scales can be formed as mentioned for example by Raju et al in the above mentioned article, as well as in the article Microstructural characterization and isothermal oxidation behavior of hot-pressed TiB2-10wt.% TiSi2 composite, Scripta Materialia, 61 (2009) 104-107, and as for example also reported by:

• L. Silvestroni, G. Meriggi, D. Sciti in the article Oxidation behavior of ZrB2 composites doped with various transition metal silicides, Corrosion Science, 83 (2014) 281 -291 ;

• L. Silvestroni, K. Stricker, D. Sciti, H.-J. Kleebe in the article Understanding the oxidation behavior of a ZrB2-MoSi2 composite at ultra-high temperatures, Acta Materialia, 151 (2018) 216-228;

• W.G. Fahrenholtz in the article Thermodynamic Analysis of ZrB2-SiC Oxidation: Formation of a SiC-Depleted Region, Journal of the American Ceramic Society, 90 (2007) 143-148;

• A.N. Astapov, Y.S. Pogozhev, M.V. Prokofiev, A.Y. Potanin, E.A. Levashov, V.l. Vershinnikov, L.N. Rabinskiy in the article Kinetics and mechanism of the oxidation of ZrSi2-MoSi2-ZrB2 ceramics in air at temperatures up to 1400 °C, International Journal of Heat and Mass Transfer, 140 (2019) 12-20;

• X. Cao, B. Wang, X. Ma, L. Feng, X. Shen, C. Wang in the article Oxidation behavior of melt-infiltrated SiC-TiB2 ceramic composites at 500-1300 °C in air, Ceramics International, 47 (2021 ) 9881 -9887.

A similar approach has been also applied for Zr-Mo-Si-B based-coatings to provide high-temperature oxidation resistance, as reported for example by:

• P. Kiryukhantsev-Korneev, A. Sytchenko, Y. Pogozhev, S. Vorotilo, A. Orekhov, P. Loginov, E. Levashov in the article Structure and Properties of Zr-Mo-Si-B- (N) Hard Coatings Obtained by d.c. Magnetron Sputtering of ZrB2-MoSi2 Target, Materials, 14 (2021 );

• Y. Niu, H. Wang, Z. Liu, C. Hu, X. Wang, X. Zheng, C. Ding in the article Microstructure evolution of ZrB2-MoSi2 composite coatings at middle and high temperatures, Surface and Coatings Technology, 273 (2015) 30-38.

Objective of the present invention

The main objective of the present invention is to provide quaternary diboride-based thin film materials by alloying TMSi2 phases into binary TiB2±z and CrB2±z in such a manner that it is possible to achieve the challenging compromise between good mechanical properties and highest oxidation resistance.

Description of the present invention

The objective of the present invention is attained by providing a coating film material as described in claim 1 and a preferred method for producing a such coating film material as described in claim 10.

The dependent claims 2 to 9 describe preferred embodiments of the inventive coating film material, and the dependent claims 11 -20 describe preferred embodiments of the inventive method for producing a such inventive coating film material.

The coating film materials according to the present invention are coating films of the type TM¹ _a-TM²b-Sic-Bd (specifying the general nomenclature of TM¹-TM²-Si-B2±z), produced by using a physical vapor deposition process (PVD process), wherein TM¹ = Ti or Cr, and TM² = Ta or Mo, the coating film material having a chemical composition in atomic percentage corresponding to:

TM¹ _aTM²bSicBd, wherein:

• 15 < a < 35, preferably 20 < a < 30,

• 5 < b < 15, preferably 5 < b < 10,

• 10 < c < 30, preferably 10 < c < 20,

• 35 < d < 65. preferably 40 < d < 60.

The TM¹ B2 material comprised in the coating film material exhibits an AIB2-structure type, which can be detected by using well-known X-Ray diffraction analysis (XRD analysis). In other words, the coating film material has an AIB2-structure type of TM¹ B2. The AIB2-structure type of TM¹ B2 can be detected by well-known XRD analysis as already mentioned above. It means that if the coating material having chemical composition TM¹ _aTM²bSicBd is examined by using XRD analysis, then if TM¹ is Ti or Cr, the coating material exhibits peaks of TiB2 or CrB2, respectively, having AIB2- structure type.

With regard to a compromise between advantageous mechanical coating properties, in particular high hardness and a high modulus of elasticity, and high oxidation resistance, preferably, the coating film material may comprise a composition in which the ratio B/(TM¹ + TM²) < 2,2 and > 1 .

Hereby, according to one preferred coating film, TM¹ = Ti, wherein the coating film material may comprise a composition in which the ratio B/(TM¹ + TM²) < 2,2 and > 1 ,2.

Likewise, according to another preferred coating film, TM¹ = Cr, wherein the coating film material may comprise a composition in which the ratio B/(TM¹ + TM²) < 1 ,8 and > 1.

AIB2-structure types of TM¹B2, shown by the film material according to the invention, are well described in the literature, for example by:

• M. Magnuson, L. Hultman, H. Hdgberg, Review of transition-metal diboride thin films, Vacuum. 196 (2022) 110567; • M. Frotscher, W. Klein, J. Bauer, C.-M. Fang, J.-F. Halet, A. Senyshyn, C. Baehtz, B. Albert, M2B5 or M2B4? A Reinvestigation of the Mo/B and W/B System, Z. Anorg. Allg. Chem. 633 (2007) 2626-2630.

One method that can be used for producing a coating film material according to the present invention preferably involves the use of a PVD process for the deposition of the inventive coating material on a substrate surface, preferably a sputtering process, in particular a magnetron sputtering process.

In the case of using the above-mentioned method, preferably at least one target (hereafter also called compound target) comprising one or more diborides of TM¹ (also referred to as a TM¹ -diborides) and one or more disilicides of TM² (also referred to as a TM²-disilicides) is used as coating material source.

Preferable the at least one TM¹ -diboride material present in the target is TiB2 or CrB2.

Preferably the at least one TM²-disilicide material present in the target is TaSi2 or MoSi2.

The above-mentioned method and preferred features of the method are not a limitation of the methods that can be used for producing coating film materials according to the present invention. Therefore, the above-mentioned method should be understood as one example of a method for producing coating film materials according to the present invention.

Hence, one preferred method for producing coating films according to the present invention, involves the production of the coating film material by using a PVD process, wherein the coating film material is deposited on a surface of a substrate, an wherein at least one target is used as coating material source, and the at least one target comprises at least one TM¹ -diboride material having a chemical composition in atomic percentage corresponding to TM¹B2, and at least one TM²-disilicide material having a chemical composition in atomic percentage corresponding to TM²Si2, wherein TM¹ = Ti or Cr, and TM² = Ta or Mo.

According to a preferred embodiment of the present invention, the inventive coating films of the type TM¹ _aTM²bSicBd are deposited on a substrate surface and afterwards subjected to an annealing process in an oxygen-containing atmosphere for producing a top oxide layer comprising Si-0 or Si-0 and Cr-O, i.e. comprising silicon oxide or comprising silicon oxide and chromium oxide. Preferably, the top oxide layer comprises silicon oxide of the type silicon dioxide (the silicon dioxide can be also for example amorphous silicon dioxide but it is not mandatory), and additionally the top oxide layer can comprise chromium oxide.

Hereby, for example, the top oxide layer (in the context of the present description also simply called oxide layer) may be formed as an oxide multilayer system having at least two layers, wherein the at least two layers are two distinct oxide layers, wherein the multilayer system preferably comprises an chromium oxide layer (the chromium oxide layer can comprise for example Cr20s and/or any other kinds of chromium oxide) as outer-layer and a silicon oxide layer as inner layer, i.e. the oxide layer is formed as an oxide multilayer system, comprising at least two distinct oxide layers, wherein:

- the multilayer system preferably comprises an outermost layer comprising chromium oxide, e.g. a chromium oxide layer, and an inner layer closer to the substrate than the outermost layer, the inner layer comprising silicon oxide, e.g. a silicon oxide layer, or

- the multilayer system preferably comprises an outermost layer comprising silicon oxide, e.g. a silicon oxide layer, and an inner layer closer to the substrate than the outermost layer, the inner layer comprising chromium oxide, e.g. a chromium oxide layer.

The chemical composition TM¹ _aTM²bSicBd of the coating film material in atomic percent is given for the as deposited coating film before annealing.

Preferably, hereby, the coating film material exhibiting a second phase of TM²Si2.

According to a preferred variant of the above-mentioned preferred embodiments, the oxide layer comprising Si-0 has a layer thickness between 100 nm and 3000 nm, preferably between 500 nm and 2000 nm, in particular between 1000 nm and 1500 nm.

Preferably the above-mentioned annealing process is an isothermal annealing treatment in ambient atmosphere in the temperature range (800 - 1200 °C, preferably 900 - 1000 °C) for at least 1 hour. This process can be modified in terms of temperature (between 800 to 1200 °C) and time to control the oxide scale thickness and composition.

Preferably, the annealing treatment conditions, in particular with respect to the annealing time, the annealing temperature and the annealing atmosphere, are purposefully selected to reduce the probability for the formation of point defects such as vacancies in the coating film material. By reducing the probability for the formation of point defects such as vacancies in the coating film material, it is possible to increase the probability for maintaining the AIB2-structure type of TM¹B2 of the coating film material. Thus, it could be observed that by carrying out an annealing process in an oxygen-containing atmosphere by using a constant annealing temperature in the temperature range from 800 up to 1200 °C for at least one hour, preferably under inert gas conditions at an inert gas pressure between 0.2 Pa and 5 Pa, the probability for the formation of point defects such as vacancies in the coating film material could be successfully reduced.

The inventive coating film materials can be used as monolayer coating layer or as one or more coating layers within a multilayer coating system.

Coating films according to the present invention are preferably provided on a substrate surface as outermost layer, for example as a monolayer or as outermost layer of a coating system comprising more than one layer, for preventing deterioration of the rest of the layers in the coating system and also for preventing deterioration of the substrate surface on which the coating system is deposited.

Preferably, hereby, the substrate is any kind of material used manufacturing components to be used in aerospace applications. For example the substrate material may comprise carbon, boron, Si-C (silicon carbide) or Si-N (silicon nitride), which all need a protection against oxidation.

However, the above-mentioned substrates should not be understood as a limitation of the present invention. In general, the inventive coating film can be suitable for any kind of substrate material requiring protection against oxidation.

Moreover, during the conduction of the PVD process a negative DC bias voltage between - 20 V and - 200 V, preferably between - 40 and - 150 V may be applied to the substrate and/or the substrate may be heated previous to the conduction of the PVD process to a substrate temperature to be maintained during the PVD process, the substrate temperature being preferably in a range between 300 °C and 600 °C, preferably between 400 °C and 500 °C.

Furthermore, the PVD process in the method according to the invention may be carried out in an inert gas atmosphere, wherein preferably argon is used as inert gas, in particular at a pressure between 0.2 Pa and 5 Pa, preferably at a pressure of at least 0,3 Pa - for example between 0.3 Pa and 2 Pa, more preferably of at least 0,4 Pa - for example between 0.4 Pa and 2 Pa or 3 Pa.

Such inventive coating systems are especially suitable for protecting substrate surfaces to be exposed to wear at elevated temperatures (i.e. temperatures in a range from 800 °C and 1400 °C, preferably from 900 °C up to 1300 °C) due to their outstanding oxidation resistance at elevated temperatures.

The invention will now be described in more details based on examples and with the help of the figures.

Detailed Description

During a detailed analysis a comparison has been carried out between coating films according to the present invention and coating films according to the state of the art

For analyzing and comparing the oxidation resistance, in particular in relation to the phase formation, and also for comparison of mechanical properties, TM¹ _aTM²bSicBd coatings were deposited via magnetron sputtering from targets comprising TiB2 /TMSi2 or CrB2 /TMSi2 with TM = Ta or Mo.

As example, the addition of TaSi2 resulted in the formation of a single-phased hexagonal Tio.28Tao.o7Sio.12Bo.53 coating films maintaining a high hardness of up to 36 GPa and exhibiting strongly retarded oxidation kinetics till 1000 °C, characterized by forming only a 550 nm thin oxide scale after 1 h air annealing at 800 °C.

The coating films of Tia-Mob-Sic-Bd (TiB2 alloyed with MoSi2) showed a moderate hardness of up to 27 GPa but their oxidation resistance behavior was exceptional attaining to preserve outstanding oxidation resistance up to 1200 °C. These coatings exhibited a protective Si-based oxide scales formation that seems to be responsible for this exceptional oxidation resistance behavior. For example, the coating film having chemical composition in atomic percentage measured by ToF-ERDA corresponding to Ti0.23Mo0.07Si0.i6B0.54 formed an oxide layer with thickness of 335 nm after annealing during 1 h at 1200 °C in an oxygen-containing atmosphere.

For the deposition of the different coating films to be analyzed, the coating films were deposited by alloying TM-diboride with TM-disilicide from compound targets comprising TM-diborides and TM-disilicides.

For example, TiB2±x (hereafter for simplifying nomenclature also referenced to as TiB2) was alloyed with TaSi2±y (hereafter for simplifying nomenclature also referenced to as TaSi2) by sputtering from targets with various compositions - for example with composition corresponding to a mol% ratio TiB2/TaSi2 of 90/10 and 80/20.

Also similarly, TiB2±x (hereafter for simplifying nomenclature also referenced to as TiB2) was alloyed with MoSi2±y (hereafter for simplifying nomenclature also referenced to as MoSi2) by sputtering from targets with various compositions - for example with composition corresponding to a mol% ratio TiB2/MoSi2 of 85/15, 80/20 and 70/30.

With these targets, all quaternary Ti_a-TM²b-Sic-Bd (with TM² = Ta or Mo) coating materials have been deposited in a laboratory-scaled magnetron sputtering system using 3-inch sized target materials from Plansee Composite Materials GmbH. Each of the above-mentioned targets was solely DC-sputtered at a target current of 0.5 A in pure argon atmosphere (working pressure of 0.4 Pa). Additionally, a binary TiB2.57 coating was deposited from a TiB2 target (actually a TiB2±x target but for simplifying nomenclature also referenced to as TiB2) at a pressure of 0.56 Pa.

In a similar procedure, the synthesis of the quaternary Cr_a-TM²b-Sic-Bd (with TM² = Ta or Mo) was performed.

These coatings were grown onto sapphire and single-crystalline Si-substrates (1011- oriented, 10x10x0.53 mm³ and 100-oriented, 20x7x0.38 mm³) as well as polycrystalline Al2O3-substrates (20x7x0.38 mm³).

As example, for the Ti_a-TM²b-Sic-Bd coatings Table 1 shows the chemical compositions evaluated by Time-of-Flight Elastic Recoil Detection Analysis (ToF-ERDA) and Rutherford Backscattering Spectrometry (RBS) are summarized for all grown films. For ToF-ERDA, ¹²⁷l⁸⁺ projectiles with a primary energy of 36 MeV were employed with an incident angle of 67.5° with respect to the surface normal and a recoil detection angle of 45°. RBS was carried out using 3 MeV ⁴He⁺ ions and a detection angle of 170°.

Table 1 shows the chemical composition and mechanical properties (Hardness and E- modulus) for all Ti_a-TM²b-Sic-Bd coatings in comparison to binary TiB2.57 coating.

The structural evolution of all grown films was investigated by X-ray diffraction (XRD). Figure 1 presents the X-ray diffractograms of the as-deposited Ti_a-TM²b-Sic-Bd coatings in comparison with the binary TiB2.57. Only peaks corresponding to the hexagonal-TiB2 phase (SG 191 ) - in addition to the AI2O3 substrate - can be indexed. Apart from the amorphous Tio.20Moo.11Sio.26Bo.43, all alloyed coatings exhibit a single-phased structure with broad 001 peaks as the preferred orientation. The high Si content in Tio.2oMoo.uSio.26Bo.43 leads to the amorphous character with diminished peaks. In contrast, the higher Ta-containing Tio.28Tao.o7Sio.12Bo.53 shows an increase in the predominant 001 peak intensity accompanied by a shift towards lower 29 values, suggesting the dissolution of Ta in the hexagonal phase.

Moreover, the mechanical properties (Hardness and Young’s modulus) of the as deposited coatings were investigated using an ultra-micro indentation (IIMIS) system equipped with Berkovich diamond tip, see also Table 1. The Ta-alloyed coatings maintained relatively high hardness values with an observed hardening effect by increasing the Ta-Si content from 32.8 to 36 GPa for Ti0.31Ta0.04Si0.06B0.59 and Tio.28Tao.o7Sio.12Bo.53, respectively. The increase in hardness is related to solid solution hardening with Ta. In contrast, the Mo-Si alloying route lead to decreased hardness with increasing alloying content. Generally, the alloying of TiB2±z with Si was emphasized to result in material softening, as mentioned for example by Glechner et al and Grancic et al. in the articles mentioned above.

Figure 1 shows X-ray diffractograms of (a) TiB2.57, (b) Tia-Tab-Sic-Bd and (c) Tia-Mob- Sic-Bd alloyed coatings with their stoichiometries indicated. X-ray diffraction (XRD) was performed in Bragg-Brentano configuration using a Panalytical X’pert Pro MPD system equipped with Cu-Ka radiation source (A = 1 .54 A).

To study the oxidation behavior, dynamic measurements up to 1400 °C were done using DTA/TG system (Netzsch STA 449-F1 ) at a heating rate of 10 °C/min, under flowing synthetic air (50 ml/min). Figure 2 summarizes the mass change during dynamic oxidation of Ti_a-TM²b-Sic-Bd coatings as a function of the annealing temperature in comparison to a binary TiB2.57 coating. The onset oxidation temperature for the un-alloyed TiB2.57 is observed to be around 490 °C. Above this temperature, the coating exhibits a mass increase with accelerated oxidation kinetics till it is fully oxidized at 975 °C, followed by a mass decrease above 1000 °C due to volatilization of B2O3. Alloying with Ta-Si provides a clear improvement in the oxidation resistance with a significant shift in the onset temperature up to 770 °C for Tio.28Tao.o7Sio.12Bo.53 (see Figure 2a). Additionally, the slope of the mass curve reduces significantly till 1000 °C - compared to the binary coating - indicating retarded oxidation kinetics due to formation of protective scales. Furthermore, also the Tia-Mob-Sic-Bd coatings exhibit excellent oxidation resistance, where the slope of the mass gain curves significantly flattens upon alloying (see Figure 2b). By increasing the Mo-Si content, the mass signal shows a plateau over 1000 °C indicating the formation of highly protective oxide scale for both Ti0.23Mo0.07Si0.i6B0.54 and Ti0.20Mo0.11Si0.26B0.43, respectively.

Figure 2 shows TG curves of mass change during dynamic oxidation of (a) Ti_a-Tab-Sic- Bd, and (b) Ti_a-Mob-Sic-Bd coatings in synthetic air under heating rate of 10 °C/min. The TG curve for the binary coating TiB2.57 is indicated by a dashed line.

To gain a more detailed understanding on the oxide scale formation process, the morphology of selected oxidized samples was investigated using transmission electron microscopy (TEM FEI TECNAI F20). Figure 3 presents the cross-sectional TEM analysis for Tio.28Tao.o7Sio.12Bo.53 after 1 h oxidation in ambient air at 800 °C. The bright- field image shows an oxide scale of 535 nm on top of unoxidized intact coating (see Figure 3a). In more detail, the formed oxide scale is composed of two layers with a distinct interface: an outer dense glassy amorphous layer with a thickness of 140 nm, and an inner layer composed of mixed equiaxed and columnar crystallites (see Figure 3b). The inner scale exhibits a relatively dense morphology with small, globular crystallites near to the coating-oxide interface, while more columnar structures with larger grains predominate the upper interface. The corresponding EELS maps (see Figure 3d and the respective elements) clearly reveal that the outer oxide scale is Si- rich with small amounts of boron, while the inner crystalline scale mainly consists of Ti and Ta-based oxides (Ta is also confirmed by EDX, not shown) with no boron detected. The formation of an outer dense Si-rich borosilicate scale is the key for the excellent oxidation resistance with retarded kinetics. Furthermore, the SAED pattern presented in Figure 3c reveals an initiation of phase separation processes between the TiB2- based matrix and TaSi2 after annealing at 800 °C.

Figure 3: TEM analysis of Tio.28Tao.o7Sio.12Bo.53 coating oxidized in ambient air at 800 °C for 1 h. (a) BF image of the whole coating with the substrate at the bottom and oxide scale on top. (b) magnified area for oxide scale with coating interface, (c) SAED image for the area indicated in (a), (d) STEM image and corresponding EELS maps for the area illustrated in (b).

Figure 4 depicts the TEM analysis for the air-annealed Ti0.23Mo0.07Si0.i6B0.54 after 1 h at 1200 °C. The unoxidized coating exhibits globular morphology with clear indications for recrystallization processes as bulk diffusion was already activated at 1200 °C - evidence for reaching about 0.4 of the melting temperature (compare with data reported by Shewmon in Transformations in metals, McGraw-Hill, 1969) - see Figure 4a.

Moreover, the BF-image clearly shows a dense oxide scale of 440 nm which is amorphous according to the SAED analysis (see Figure 4b and ci). This coating experienced a separation of the MoSi2 phase as indicated in SAED image (Figure 4c2) after the annealing at 1200 °C. The EELS maps in Figure 4d reveal that the oxide scale is based only on Si with no competing boron. However, boron-rich pockets can be observed at the coating-oxide interface due to formation of MoB phase according to reaction (1 ). The formation of MoB was reported by Silvestroni et al. for ZrB2/MoSi2 bulk system at 1200 °C in the first article of Silvestroni et al. mentioned above. The MoSi2 phase beneficially acts as an active reservoir for selective oxidation of Si and the formation of the highly dense and protective Si-based scale, while suppressing the detrimental gaseous B2O3 phase.

2 • MoSi₂ + B2O3 + 5/2 • O2 = 4*SiO₂ + 2 • MoB (1 )

Figure 4 shows TEM analysis of Ti0.23Mo0.07Si0.-i6B0.54 coating oxidized in ambient air at 1200 °C for 1 h. (a) BF image of the whole coating with the substrate at the bottom and oxide scale on top. (b) magnified area for oxide scale with coating interface. (c1 ) and (c2) SAED patterns for areas indicated in (a), (d) elemental EELS maps for the area illustrated in (b).

To present the best compromise between mechanical properties and oxidation resistance, we summarized the as-deposited hardness values with respect to the oxidation temperature for selected Ti_a-TM²b-Sic-Bd coatings and compared them with TiB2±z-based literature data (see Figure 5) to demonstrate the outstanding properties of our invention

Figure 5 shows the as-deposited hardness of diverse alloyed TiB2±z coatings in relation to their oxidation temperature Tox. The obtained scale thickness (at Tox) for each coating is indicated in relation to the reported oxidation time. The as-deposited coating thicknesses are: 4.9 pm for Tio.28Tao.o7Sio.12Bo.53, 3.5 pm for Ti0.23Mo0.07Si0.i6B0.54, 4.9 pm for Tio.20Moo.11Sio.i6Bo.54, 400 nm for TiBi.43, 980 nm for Tio.9Alo.1B1 3, 1.3 pm for (Tio.35Alo.65)B2, ~1.5 pm for (Tio.68Alo.32)Bi.35,and ~1.4 pm for Tio.13Sio.41 Bo.46.

Figure 6 shows the structural analysis at different temperatures up to 1100 °C by XRD for the coatings. The TiB2 peaks are marked with black dots, the MoSi2 peaks with grey boxes.

Further experiments have been carried out with respect to Cr_a-TM²b-Sic-Bd coatings. The coatings were synthesized in a laboratory-scale magnetron sputtering system using 3-inch sized CrB2-based targets alloyed with TMSi2 secondary phases (Plansee Composite Materials GmbH). The employed targets are CrB2/CrSi2 (90/10 and 80/20 mol.%) as well as CrB2/MoSi2 (90/10, 80/20, and 70/30 mol.%). The depositions were carried out in a pure argon atmosphere (working pressure of 0.4 Pa) by sputtering each of the mentioned targets solely in a DC mode at 0.4 A (200 W). The base pressure before all depositions was maintained below 10’⁴ Pa. Additionally, a binary CrBi.s coating was grown from a pure CrB2 target at a working pressure of 0.56 Pa. For all depositions, the substrate temperature was maintained at 550 °C, whereas the rotating substrate holder (0.25 Hz) was mounted parallel to the target at a distance of 90 mm. Furthermore, a substrate DC bias potential was applied between -40 V and -150 V. The coatings were deposited on single crystalline Si (100-oriented, 20x7x0.38 mm³), single crystalline sapphire (101 1 -oriented, 10x10x0.53 mm³), and poly-crystalline AI2O3 (20x7x0.38 mm³) substrates. Prior to the depositions, the bare substrates were cleaned in an ultrasonic bath using acetone and ethanol. Afterwards, a plasma cleaning step of the substrates is carried out for 10 min in pure Ar at 5 Pa.

The chemical composition of the coatings was determined by ion beam analysis techniques using Time-of-Flight Elastic Recoil Detection Analysis (ToF-ERDA) and Rutherford Backscattering Spectrometry (RBS) at the 5 MV Pelletron Tandem accelerator laboratory at Uppsala University. For ToF-ERDA, 12718+ projectiles with a primary energy of 36 MeV were employed with an incident angle of 67.5° with respect to the surface normal while recoils were detected at an angle of 45° with respect to the incident beam. RBS was carried out using 3 MeV 4He⁺ ions and a detection angle of 170°. The analysis of the ToF-ERDA experimental data was performed using the Contes software, while the RBS data were analyzed using the SIMNRA software. The RBS-analysis provides primarily information on the concentration ratio of transition metals as well as confirms constant concentrations over a large depth range (> 1 pm), whereas ToF-ERDA provides accurate information on light species such as boron and can confirm the absence of contaminations. In combination, high accuracy without the need for standards can be achieved. The total systematic and statistical uncertainties were estimated to be at most 5-8% of the deduced value for the major constituents and below 15% for the oxygen traces. Furthermore, complimentary chemical analyses were done using liquid inductively coupled plasma optical emission spectroscopy (ICP- OES). The coatings on AI2O3 substrates were acid digested in 0.5 mL HNO3 and 0.5 ml HF in falcon tubes for 10 min without heating. The analyses were carried out with an iCAP 6500 RAD (Thermo Fisher, USA), with an ASX-520 autosampler (CETAC Technologies, USA) using an HF resistant sample introduction kit consisting of a Miramist nebulizer (Burger Research, USA), an alumina injector tube and a PTFE spray chamber. For each investigated element at least two non-interfered emission lines with sufficient sensitivity were selected, further instrumental parameters were used as recommended by the manufacturer. Background corrected emission signals were quantified using matrix-adjusted external calibration standards.

The growth morphology of the as-deposited coatings was investigated using scanning electron microscopy (SEM, Zeiss Sigma 500VP). The crystal structure was investigated by X-ray diffraction (XRD) in Bragg-Brentano configuration using a Panalytical Xpert Pro MPD system equipped with Cu-Ka radiation source (wavelength A = 1 .54 A, operated at 45 kV and 40 mA). In order to track the phase evolution of selected coatings at elevated temperatures, in-situ XRD measurements were carried out in Bragg-Brentano configuration using an Anton Paar high-temperature furnace chamber (HTK 1200N). The XRD measurements were performed in vacuum at room temperature and between 500 °C and 1100 °C every 50 °C step.

The oxidation behaviour of the coatings was investigated using a thermogravimetric analyzer (TGA) (Netzsch STA 449 F1 ) equipped with a Rhodium furnace. The dynamic measurements were carried out up to 1400 °C with a heating rate of 10 °C/min, under a flowing stream of synthetic air (50 ml/min) and helium (20 ml/min). For these experiments coated AI2O3 substrates (pre-weighted before deposition) have been used. For selected coatings, further isothermal annealing tests were done using a conventional furnace in ambient air at 1200 °C up to 30 h.

To investigate the morphology of the partially oxidized coatings, cross sections were prepared using a FIB-SEM dual-beam system (ThermoFisher Scientific Scios 2). The cross sections were milled using Ga⁺ ion beam currents of 7 - 15 nA for rough milling and 1 nA for fine milling at an acceleration voltage of 30 kV. Moreover, the growth morphology for selected oxidized samples was further investigated by transmission electron microscopy (TEM, FEI TECNAI F20, equipped with a field emission gun and operated at 200 kV acceleration voltage). Additionally, energy dispersive X-ray spectroscopy (EDX) and electron energy-loss spectroscopy (EELS) mappings, as well as line scans of the partially oxidized coatings, have been conducted.

The three-dimensional distribution of elements was investigated by atom probe tomography (APT) for a selected Cr-Mo-Si-B2-z coating in the as-deposited and annealed states. The APT samples were prepared using a focused ion beam (FIB) (Scios 2 DualBeam system, ThermoFischer Scientific). Lift-outs were extracted from the film surfaces using ion beam current of 3 - 5 nA for rough milling and 1 nA for fine milling. The acceleration voltage was set to 30 kV. Moreover, the final sharpening of the tips was performed at 50 pA, with a subsequent clean-up step at 28 pA and 5 kV. The APT analysis was done using a Cameca LEAP 4000X HR equipped with a 355 nm UV laser and a reflection lens. The measurements were done in a pulsed laser mode employing a laser pulse energy of 50 - 90 pJ. The target evaporation rate was 1 % and pulse rate was set to 200 kHz. Data analysis was done using a Matlab toolbox.

The hardness of the coatings was investigated using an ultra-micro indentation (UM IS) system equipped with Berkovich diamond tip. For each sample, 31 indents were done in a load-controlled mode with indentation loads varied between 3 to 45 mN and consequently evaluated based on the Oliver and Pharr method. The indentation depths were kept below 10 % of the coating thickness to reduce the substrate interference.

Table 2 shows the chemical composition for all Cr_a-TM²b-Sic-Bd coatings in comparison to binary CrBi.s coating and five different ternary Cr_a-Sic-Bd-coatings (samples 2 to 6). Figure 7 shows the cross-sectional SEM images of the as-deposited CrBi.s and alloyed Cr-(Mo)-Si-B2-z coatings with their corresponding chemical compositions. The binary CrBi.s film exhibits a columnar microstructure with extended grains along the growth direction. The addition of Si in the ternary Cr-Si-B coatings results in grain refinement with finer microstructures and relatively smooth surfaces. The Mo-containing Cr-Mo- Si-B2-z coatings show fine-structured dense morphologies which tend to grow featureless by increasing the Mo-Si content. The thickness of the as-deposited coatings ranges between 2 to 2.5 pm for the ternary Cr-Si-B2-z, while the quaternary Cr-Mo-Si-B2-z exhibit thicknesses between 3 and 4.5 pm. The chemical analysis by ToF-ERDA reveals for all coatings sub-stoichiometric boron compositions with B/TM <2, see Table 2. Generally, sputter-deposited CrB2-z films typically exhibit a lower B/Cr ratio compared to the employed target compositions, also well described in the literature with ratios between 1 to 1.5 while using stoichiometric targets. Furthermore, the boron content in the alloyed coatings decreases with increasing Si content. In a Cr- Si-B2 system, the Si tends to occupy both Or- and B-sites in their corresponding sublattices. However, the system is prone to form Si grain boundary segregates upon reaching a solubility limit of about 4 at. %. For all coatings the oxygen content was below 1 .6 at. %.

Figure 8 depicts the structural analysis by XRD for all Cr-(Mo)-Si-B2-z coatings as well as the binary CrBi.s on polycrystalline AI2O3 substrates at -40 V substrate bias potential. The AI2O3 substrate peaks are marked with black dotted vertical lines. The binary CrBi.s coating obtains a 001 preferred orientation in its hexagonal AIB2 structure type (SG 191 ). The addition of Si in Cr0.38Si0.08B0.54 results in a change of the preferred orientation from 001 to 100, while maintaining the single hexagonal structure with no other additional phases. However, a further increase in Si content results in amorphization of the Cro.37Sio.i6Bo.47 coating with diminished peak intensities. The Si addition results in a slight shift of the low-intensity (001 ) peak to higher 29 indicating a decrease in the c lattice parameter compared to the binary coating. Furthermore, the quaternary Mo-containing Cr-Mo-Si-B2-z coatings (see XRD patterns within Figure 8) exhibit only peaks of the single hexagonal phase. For the low-alloyed Cr0.38Mo0.02Si0.02B0.58 coating, a preferred 001 orientation can be observed. However, the peaks diminished by increasing the Mo-Si content within the coatings obtaining clear indications of grain refinement for Cr0.31Mo0.07Si0.15B0.47 (as also depicted in 7). Even higher Mo-Si contents lead to an X-ray amorphous character, see Cro.26Moo.iiSio.24Bo.39 in Figure 8. The incorporation of Si is reported to typically segregate along grain boundary regions hence causing grain/columnar refinement resulting in the formation of amorphous structures at high alloying contents.

Figure 9 presents (a) the hardness values of all as-deposited coatings as a function of the Si content as well as (b) the applied substrate bias potential during film growth. The indicated sample numbers correspond to the detailed chemical composition of the coatings included in Table 2. The highest hardness values were obtained for the low alloyed Mo-containing Cr0.38Mo0.02Si0.02B0.58 followed by the binary CrBi.s, both predominantly exhibiting a preferred 001 orientation in the single-phase hexagonal structure (see Figure 8). Moreover, the other alloyed coatings show a decrease in hardness by increasing Si content (see Figure 9a), where the lowest hardness value of 20.3 ± 0.9 is recorded for Cro.26Moo.11Sio.24Bo.39 exhibiting an amorphous character (as proven by XRD analysis, see Figure 8). Nevertheless, the applied substrate bias potential shows a clear influence especially on the hardness of the Mo-containing Cr- Mo-Si-B2-z coatings, wherein the hardness increased from 21.1 ± 1.1 GPa to 25.8 ± 1.4 GPa as the bias rises from -40 V to -150 V (Figure 9b). The increase in hardness is related to the enhanced degree of crystallinity and the reduced Si content within the coatings at higher bias potentials as can be seen in Figure 10 showing X-ray diffractograms of (a) Cr-Si-B2-z and (b) Cr-Mo-Si-B2-z coatings deposited at different substrate bias potential. The increase in bias potential is accompanied by increased Ar ion irradiation during film growth, hence enhancing adatom mobility as well as preferential re-sputtering of the weakly bonded Si. This effect beneficially promotes the formation of solid solutions of (Cr,Mo)B2 within the Mo-containing Cr-Mo-Si-B2-z coatings indicated by the pronounced hexagonal peaks during structural analysis (see Figure 10). However, for the ternary Cr-Si-B2-z coatings, the increased bias potential kept the hardness more or less unchanged, suggesting no strong influence on the Si distribution within the growing film. In general, Si alloying in transition metal diborides is expected to reduce hardness. Nevertheless, in case of the Mo-containing coatings, the high bias promoted the formation of single phased solid solutions of (Cr,Mo)B2 resulting in high hardness values up to ~26 GPa - even though the 001 orientation is not predominant.

To study the oxidation behavior of the Cr-(Mo)-Si-B2-z coatings, dynamic oxidation analyses have been performed on coated polycrystalline AI2O3 substrates in a TG system at a heating rate of 10 °C/min up to 1400 °C in synthetic air. Figure 11 summarizes the mass gain (in percentage of the actual coating mass) during dynamic oxidation of all alloyed Cr-(Mo)-Si-B2-z coatings as a function of the annealing temperature. The mass signal for the binary CrBi.s coating is added for comparison and indicated by a black dashed line. The binary CrBi.s exhibits an onset oxidation temperature of 600 °C, proven by a previously constant mass signal indicating no pronounced oxidation process below this temperature. The mass signal starts to increase sharply above 600 °C in a step-wise behavior until 1 100 °C indicating the complete oxidation of the coating material followed by a mass loss due to the formation of volatile oxides such as CrOs and B2O3. The addition of Si clearly improves the oxidation resistance (as shown in Figure 1 1 a), where the slopes of the ternary coatings significantly reduce upon alloying compared to the binary CrBi.s. The high Si- containing Cro.37Sio.i6Bo.47 exhibits nearly constant mass signal up to 1200 °C followed by a slight increase, indicating the formation of a highly protective Si-based scale. Moreover, the quaternary alloyed Cr-Mo-Si-B2-z coatings exhibit outstanding high- temperature oxidation resistance clearly outperforming the ternary counterparts (see Figure 11 b). The mass signal completely flattens for both Cr0.31Mo0.07Si0.15B0.47 and Cro.26Moo.11 Sio.24Bo.39 with no mass gains or losses up to 1300 °C. This is clearly related to the formation of protective scales inhibiting the oxygen inward diffusion.

To further investigate the detailed oxidation mechanisms and especially the oxide scale formation of the ternary and quaternary coatings, we selected two coatings for more detailed analysis. Therefore, Cro.37Sio.i6Bo.47 and Cr0.31Mo0.07Si0.15B0.47 have been oxidized at 1200 °C in ambient air and subsequently analyzed concerning their phase formation and changes in the morphology (by SEM, TEM, and in-situ XRD).

In Figure 12, the TEM analysis for the oxidized Cro.37Sio.i6Bo.47 coating after 1 h at 1200 °C in ambient air is presented. Figure 12a shows a high-angle annular dark-field (HAADF) image of the coating cross section with the substrate at the bottom and scale on the top. Figure 12b shows a magnified area of top oxide scale. Figure 12c shows a magnified area indicated in Figure 12a and Figure 12d show EDX elemental maps of the area illustrated in Figure 12b. Figure 12e shows an EELS linescan over the coating cross section as indicated in by a white line in Figure 12a.

The high-angle annular dark-field (HAADF) image of the coating cross-section hereby shows a globular morphology being a clear indication for recrystallization due to reaching temperatures above 0.5 TM. In addition, a high degree of porosity is visible in Figure 12a. Based on the mass contrast, the HAADF image also indicates phase separations, see the bright and dark regions. The corresponding EELS linescan (see Figure 12e) proves the bright regions to be Cr- and B-rich, whereas the dark regions are enriched in Si. The phase separation of Si from the CrB2-z matrix is in good agreement with the solubility limit of around 4 at. %. At this high temperature (1200 °C), the segregated Si has high mobility and tends to diffuse towards the surface leaving voids within the coating, see Figure 12a. On the very top, a thin double-layered oxide is formed with a convoluted surface structure. The elemental EDX maps of the surface near regions (see Figure 12d) reveal the outer layer to be a Cr-based oxide with a thickness of around 200 nm, while the inner oxide layer is a thin Si-based scale of only 50 nm. The formation of this Si-0 layer at the scale-coating interface is the key to the high-temperature oxidation resistance (see Figure 11 ), as it acts as an efficient barrier against oxygen inward diffusion. This barrier also slows down the outward diffusion of Cr, which typically has the tendency to form volatile oxides in these temperature regimes. It must be noted that the EDX analysis is not accurate regarding the quantification of light elements which results in an overestimation of the B content.

Figure 13 presents the TEM analysis of the oxidized Cr0.31Mo0.07Si0.15B0.47 coating at 1200 °C for 1 h. Figure 13 a shows a High-angle annular dark-field (HAADF) image of the coating cross section with the substrate at the bottom and scale on the top. Figure 12b shows a magnified area of top oxide scale. Figure 13c an EDX elemental maps of the area illustrated in Figure 13b.

The cross-sectional HAADF image shows that the unaffected coating is recrystallized and exhibits a globular morphology with relatively large grains. The different phases can be identified based on their mass contrast where brighter grains indicate elements with high atomic number. This coating forms a separate phase of MoSi2 in addition to the main phase of hexagonal solid solution CrMoB2-z (see also the structural analysis in Figure 14). The formed oxide scale on top is composed of two layers: an outer crystalline layer of Cr20s and an inner thin layer of amorphous SiO2. Additionally, the observed bright grains at the scale-coating interface are rich in Mo and B. The formation of a Si-depleted MoBx phase at the scale near region can be related to the outward diffusion of Cr and Si to form the corresponding top oxide scales. Moreover, a small Si region which is depleted in oxygen is observed right beneath the scale. Compared to the ternary coating, the Cr0.31Mo0.07Si0.15B0.47 exhibits higher stability and a lower degree of porosity and voids at the same temperature. This positive effect is attributed to the presence of Mo which tends to form locking phases with both Si and B (MoSi2 and MoBx), and hence, retarding the mobility of the available Si within the microstructure. The formation of MoSi2 and MoBx is in good agreement with previously reported bulk ZrB2/MoSi2 also obtaining beneficial influence on the high-temperature oxidation resistance. The MoSi2 phase contributes to the formation of SiO2 as an oxidation product, while suppressing the formation of detrimental volatile B2O3 by forming a MoB secondary phase.

To further describe the phase evolution during annealing, in-situ XRD analysis of powdered Cr0.31Mo0.07Si0.15B0.47 coatings have been conducted in vacuum. As depicted in Figure 14, the Cr0.31Mo0.07Si0.15B0.47 coating exhibits a single phased (AIB2-structure) solid solution up to ~650 °C. At 700 °C, peaks of the tetragonal MoSi2 phase start to evolve indicating the initiation of the phase formation. By increasing the temperature, the MoSi2 peaks become very pronounced, while the intensity of the main peaks of the AIB2-structure significantly increases due to the recrystallization effect at elevated temperatures as described previously in the TEM analysis (see Figure 13).

The high-temperature phase evolution in Cr0.31Mo0.07Si0.15B0.47 is corroborated by APT analysis with local chemical compositions at the nm-scale. The reconstructions of atomic positions and concentration profiles for Cr0.31Mo0.07Si0.15B0.47 in the as- deposited state and after annealing at 800 °C are shown in Figure 15. Hereby, Figure 15a shows reconstructions of Cr, B, Si and Mo for the as-deposited Cr0.31Mo0.07Si0.15B0.47, Figure 15b shows a concentration profile of the cylindrical region indicated in Figure 15a. Figure 15c shows reconstructions of Cr, B, Si and Mo for the annealed Cr0.31Mo0.07Si0.15B0.47 at 800 °C, and Figure 15d a concentration profile of the cylindrical region indicated in Figure 15c.

As-deposited Cr0.31Mo0.07Si0.15B0.47 exhibits a homogenous composition with random elemental distribution for Cr, B and Mo in the entire volume (see Figures 15a and b). Nevertheless, chemical modulations in the nm-scale can be observed for Si which shows variation between 8 and 14 at. % (see Figure 15b), indicating the tendency for clustering in specific regions within the microstructure (e.g. at grain boundaries). At 800 °C, the Si clustering is significantly pronounced, as segregations enriched with Si and Mo can be seen in Figure 15c. Cr and B depletions are observed within the segregated regions, while the concentrations of Si and Mo within the segregations reach up to 48 and 22 at. %, respectively (see Figure 15d). It can be inferred that these segregated regions are MoSi2 domains based on the previously discussed XRD (Figure 14). However, the Mo concentration in the matrix remains around 5 at. %, suggesting the dissolution of Mo in the hexagonal (Cr,Mo)B2-z phase without being fully consumed in the MoSi2 formation.

To investigate the long-term stability of the Cro.31 Mo0.07Si0.15B0.47 coating, we performed isothermal annealing in ambient air at 1200 °C up to 30 h. Figure 16 shows the cross- sectional SEM images at different time periods. These cross-sections have been prepared by FIB milling. The white dashed lines indicate both the coating-scale and substrate-coating interfaces, while the grey dashed line indicates the interface between the oxide scale and the FIB deposited protective layer. After 3 h, the Cro.31 Mo0.07Si0.15B0.47 coating forms a top oxide scale of ~ 900 nm with an amorphous character accompanied by a thin crystalline Cr-based oxide (similar to the one obtained after 1 h). The scale thickness increases after 10 h to around 2.4 pm. After 30 h, the oxides scale is more uniform and the thickness stays at about 2.5 pm with negligible increase compared to the thickness obtained after 10 h, demonstrating the strongly retarded oxidation kinetics. The formed scale is dense and mainly amorphous, with a top bright appearing nanocrystalline oxide layer (highlighted by small arrows in Figure 16). The underlying coating is intact with some porosities.

Figure 17 shows a more detailed analysis (cross section and EDX line scan) of the Cro.31 Mo0.07Si0.15B0.47 oxidized coating after 10h. The formed oxide scale is Si-rich with small amounts of Cr confirming the layered structure of crystalline Cr20s on top of an amorphous SiO2. Beneath this scale, a region depleted in Cr and enriched in both Si and Mo is recognizable, indicating the formation of MoSi2. Moreover, a Si-depleted zone can be observed at the substrate near region. The outstanding oxidation resistance after 30 h at 1200 °C is attributed to the beneficial phase formation of MoSi2 acting as a locking phase for the highly mobile Si, and reservoir for the SiO2 formation.

Figure 18 shows isothermal oxidation tests (mass decrease over time) at 800 °C for 1 h for four different coating materials, one ternary TiSicBd composition and three different quaternary TM¹ _aTM²bSicBd, coating films together with TEM images of the coating materials. Figure 19 shows isothermal oxidation tests (mass decrease over time) at 1200 °C for 1 h for two quaternary TM¹ _aTM²bSicBd, coating films together with TEM images of the coating materials.

The present invention provides quaternary TM¹ _aTM²bSicBd, coating films that demonstrate to have a high potential as coating materials for high temperature applications, owing to their superior oxidation resistance compared to their binary and ternary counterparts.

Claims

1. A coating film comprising in the as deposited state a transition metal diboride, preferably a quaternary metal diboride, the metal diboride formed by at least two transition metals, silicon and boron, characterized in that the coating film material exhibiting a AIB2-structure type of TM¹B2 and having a chemical composition in atomic percentage corresponding to TM¹ _aTM²bSicBd, wherein:

• TM¹ = Ti or Cr

• TM² = Ta or Mo

• 15 < a < 35, preferably 20 < a < 30,

• 5 < b < 15, preferably 5 < b < 10,

• 10 < c < 30, preferably 10 < c < 20,

• 35 < d < 65, preferably 40 < d < 60.

2. The coating film according to claim 1 , characterized in that the ratio B/(TM¹ + TM²) < 2,2 and > 1.

3. The coating film according to claim 1 , characterized in that TM¹ = Ti and the ratio B/(TM¹+ TM²) < 2,2 and > 1 ,2.

4. The coating film according to claim 1 , characterized in that TM¹ = Cr and the ratio B/(TM¹ + TM²) < 1 ,8 and > 1 .

5. The coating film according to any of the previous claims, characterized in that the outermost surface of the coating film comprises an oxide layer resulted from an oxidation process of the coating film, wherein said oxide layer comprises silicon oxide or comprises silicon oxide and chromium oxide.

6. The coating film according to any of the previous claims, characterized in that the oxide layer has a layer thickness between 100 nm and 3000 nm, preferably between 500 nm and 2000 nm, in particular between 1000 nm and 1500 nm.

7. The coating film according to claim 5 or 6, characterized in that the oxide layer comprises silicon dioxide or silicon dioxide and chromium oxide.

8. The coating film according to claims 5 to 7, characterized in that the oxide layer is formed as an oxide multilayer system, comprising at least two distinct oxide layers, wherein:

9. The coating film according to any of the previous claims, characterized in that the coating film material exhibiting a second phase of TM²Si2.

10. A method for producing a coating film according to the previous claims, wherein the coating film material is produced by using a PVD process, wherein the coating film material is deposited on a surface of a substrate, characterized in that as coating material source at least one target comprising at least one TM¹ -diboride material having a chemical composition in atomic percentage corresponding to TM¹ B2, and at least one TM²-disilicide material having a chemical composition in atomic percentage corresponding to TM²Si2, wherein:

• TM¹ = Ti or Cr

• TM² = Ta or Mo

11 . The method according to claim 10 for producing coating films according to any of the previous claims 1 to 9, characterized in that the PVD process is formed as a sputtering process, preferably as a magnetron sputtering process.

12. The method according to claim 10 or 11 for producing coating films according to any of the previous claims 5 to 9, characterized in that the coating film material after deposition on the substrate surface is subjected to an annealing process in an oxygencontaining atmosphere for producing the oxide layer.

13. The method according to claim 12, characterized in that the annealing process is conducted as an isothermal annealing treatment in ambient atmosphere at an annealing temperature in the temperature range from 800 up to 1200 °C, preferably in the temperature range from 900 up to 1000 °C, for at least 1 hour.

14. The method according to claim 12 or 13, characterized in that the annealing temperature is maintained constant or varied between 800 and 1200 °C, preferably maintained constant or varied between 900 and 1000 °C, for controlling the oxide scale thickness and/or the oxide scale composition.

15. The method according to any of claims 12 to 14, characterized in that the time of the duration of the annealing treatment is selected for controlling the oxide scale thickness and/or the oxide scale composition.

16. The method according to any of the claims 10 to 15, characterized in that the annealing treatment conditions, in particular with respect to the annealing time, the annealing temperature and the annealing atmosphere, are purposefully selected to reduce the probability for the formation of point defects such as vacancies in the coating film material.

17. The method according to any of claims 10 to 16, characterized in that the coating film is deposited on a substrate, wherein the substrate is any kind of material used manufacturing components to be used in aerospace applications, in particular a material comprising carbon, boron, Si-C or Si-N, which all need a protection against oxidation.

18. The method according to any of claims 10 to 17, characterized in that for the conduction of the PVD process a negative DC bias voltage between -20 V and -200 V, preferably between -40 and -150 V, is used.

19. The method according to any of claims 10 to 18, characterized in that the PVD process is carried out in an inert gas atmosphere, wherein preferably argon is used as inert gas, in particular at a pressure between 0.2 Pa and 5 Pa, preferably at a pressure of at least 0,3 Pa, more preferably of at least 0,4 Pa.

20. The method according to any of claims 10 to 19, characterized in that the the PVD process is carried out at a substrate temperature between 300 °C and 600 °C, preferably between 400 °C and 500 °C.