PVD coatings with a HEA ceramic matrix with controlled precipitate structure

Abstract

The present invention discloses a PVD coating process for producing a multifunctional coating structure comprising the steps of producing a HEA ceramic matrix on a substrate and the targeted introduction of a controlled precipitate structure into the HEA ceramic matrix to generate a desired specific property of the coating structure.

Claims

1. A PVD coating process for producing a multifunctional coating structure comprising the steps of: Producing a HEA ceramic matrix on a substrate, and Targeted introduction of a controlled precipitate structure into the HEA ceramic matrix to generate a desired specific property of the coating structure, wherein the targeted introduction of the controlled precipitate structure into the HEA ceramic matrix takes place via a thermal treatment, and only less than 50 nm are heated directly in the HEA ceramic matrix layer by means of the thermal treatment and the targeted introduction of the controlled precipitate structure into the HEA ceramic matrix takes place in situ during the production of the HEA ceramic matrix, and the controlled precipitate structure is introduced into the HEA ceramic matrix by means of an instantaneous heating source for heating the surface of the HEA ceramic matrix to a temperature of between 900° C. and 1000° C.

2. The PVD coating process according to claim 1, wherein a variation of the targeted introduction of the controlled precipitate structure takes place by varying at least the treatment time or the treatment temperature.

3. The PVD coating process according to claim 1, wherein the HEA matrix is deposited on a substrate by applying a negative bias voltage to the substrate during the coating process, wherein the bias voltage is <200 V.

4. The PVD coating process according to claim 1, wherein at least the production of the HEA ceramic matrix or the controlled precipitate structure takes place within a reactive atmosphere.

5. The PVD coating process according to claim 1, wherein a sputtering technique is used as the PVD coating process.

6. The PVD coating process according to claim 1, wherein the substrate temperature during the production of at least the HEA ceramic matrix or the controlled precipitate structure is between 100° C. and 400° C.

7. The PVD coating process according to claim 1, wherein the process is conducted in a chamber designed in such a way that the coated substrate is alternately exposed to coating and the instantaneous heating source, and wherein the coating growth continues by coating deposition and instantaneous heating, in this manner forming the HEA matrix and the controlled precipitate structure in a sequential way.

Description

DETAILED DESCRIPTION

(1) FIG. 1 shows the structural evolution as a function of temperature for different alloys,

(2) FIG. 2 shows the structural evolution as a function of temperature for different alloys (a), the evolution of T.S config. as a function of temperature (b), and the hardness evolution as a function of annealing temperatures for different alloys (c),

(3) FIG. 3 shows the inventive microstructure of HEA matrix with controlled precipitates (a), an x-SEM image in BSE contrast, wherein the proposed structure is formed by thermal annealing (b), and a schematic representation of sequential coating deposition and flash heating (c).

(4) FIG. 1 shows the structural evolution as a function of temperature for different alloys, particularly FIG. 1 shows a comparison of the XRD evolution of AlVTiSiN with known metastable alloys of c-Ti—Al—N and c-Ti—Si—N.

(5) In contrast to the metastable alloy, the proposed multi-principal high entropy alloy of Ti.sub.37Al.sub.34Si.sub.12V.sub.17N does not show any phase transformations. In spite the alloy has immiscible components of AlN and Si.sub.3N.sub.4, and the cubic phase solid solution is retained after annealing of 1100° C.

(6) The alloy c-TiSiN in its deposited state as shown in FIG. 1 forms a nano-crystalline phase as evident from the broad XRD cubic phase. However, at elevated temperature the peak becomes narrower indicating a grain growth process, with precipitation of undesirable phases as marked by arrow.

(7) In contrast, for example in TiAlSiVN alloy, the nano-crystallinity is retained, and the cubic solid solution is preserved after 1100° C. annealing, as can be seen in the corresponding XRD diffractogram.

(8) FIG. 2a shows the structural evolution as a function of temperature for different alloys. The images on the left side of the XRDs are cross-sectional SEMs of fractured coating on BSE mode. Note that the as deposited cubic solid solution is retained in the alloy of Ti.sub.37Al.sub.34Si.sub.12V.sub.17N after annealing to 1100° C. but not in the case of Ti.sub.35Al.sub.65N as confirmed by the SEM image of fractured coating. In FIG. 2a the fractured image does not show any structural or contrast changes indicating the phase changes for the Ti.sub.37Al.sub.34Si.sub.12V.sub.17N alloy. In contrast, the SEM image of TiAIN coating in cross-sectional view shows a structural change from columnar to a fine granular structure with local contrast variations corresponding to phase changes.

(9) FIG. 2b evaluates the thermodynamic driving force for the cubic phase stabilization by considering the values of H.sub.mix (enthalpy of mixing), this is the energy penalty to cause mixing, and the energy gain caused by configurational entropy S conf.

(10) In FIG. 2b, the evolution of T.S configuration, and H.sub.mix as a function of temperature for different alloys were estimated. Note that for the alloy Ti.sub.37Al.sub.34Si.sub.12V.sub.17N the value of T.S.sub.conf over comes H.sub.mix there by solid solution is energetically more preferable compared to their binary nitrides, this is not the case for TiSiN, and AlVN. For the alloy of c-Al.sub.65V.sub.35N, and c-Ti.sub.35Al.sub.65N in FIG. 2b a cubic phase is formed as deposited state, and the alloy undergo decomposition as evidenced in the XRD above annealing of 900 C. The evolution of w-AlN in the XRD of c-Al.sub.65V.sub.35N, and c-Ti.sub.35Al.sub.65N represents following decomposition pathway.
c—Ti.sub.35Al.sub.65N.fwdarw.cTiN, and w—AlN.

(11) Thermodynamic parameters in FIG. 2b, i.e H.sub.mix of the alloys was estimated using Density functional theory calculations, and S conf. is estimated by formulism using Boltzmann equation, assuming random solid solution model.

(12) The above alloy of Ti.sub.37Al.sub.34Si.sub.12V.sub.17N is only an example, experts in the field realize similar entropy stabilized multi-principal nitrides, can be synthesized from group IV, V, and VI elements in the periodic table.

(13) The high thermal stability of this alloy is caused by the formation of entropy stabilized solid solution as shown in FIG. 2b. Only for the alloy of c-TiAlSiVN the TΔ Sconfig.mix is greater than ΔH.sub.mix, at a temperature above 1000° C. Thereby ΔG.sub.mix of the alloy is lower in the solid solution form in relation to their competing phases. Where T is temperature, S config. is the configurational entropy, G.sub.mix is the free energy of mixing and R is Gas constant.

(14) FIG. 2c shows the hardness evolution as a function of annealing temperature for different alloys. Note that the thermally stable alloy of Ti.sub.37Al.sub.34Si.sub.12V.sub.17N also shows a relatively stable hardness as a function of annealing temperature. Especially at temperature above 1000° C., where the alloy decomposition is triggered. The entropy stabilized Ti.sub.37Al.sub.34Si.sub.12V.sub.17N alloy also shows a relatively stable hardness up to annealing temperature of 1100° C. as shown in FIG. 2c

(15) The second step of the first aspect of the present invention is schematically shown in FIG. 3. The matrix can be any high entropy alloy ceramic of nitride, carbide, oxide, and boride. The precipitate composition is carefully chosen that it imparts either additional strengthening or it induces additional functionalities like lubrication etc. Also the precipitate composition and structure chosen in a way that they do not induce undesirable stresses.

(16) The composition and structure of precipitate can be similar or different, for example: H—BN in HEA nitride or HEA carbide alloy. The controlled precipitate structure can form during the as deposited state in-situ or via a post annealing route.

(17) According to a first example the proposed microstructure comprises two components as presented in FIG. 3a. The first component is used to form an entropy stabilized ceramic matrix.

(18) The second component is to form a controlled precipitates in the structure. The pathways to form the desired micro-structure is schematically presented in FIG. 3. The precipitates can be any oxides, carbides, and borides in a HEA ceramic matrix. The example path-ways to form this structure is as following.

(19) (a) By post annealing as shown in FIG. 3b: By tuning the matrix composition appropriately for metallic, and non-metallic lattice, a desired precipitation of phases is achieved in a HEA matrix. The example in FIG. 3a shows WAlTiSiN alloy, where few nm scale precipitates are achieved by thermal annealing. The size of the precipitate, chemistry, and the interface structure of the precipitate is tuned by the alloy composition, and time, and temperature of the annealing cycle.

(20) (b) To enable similar precipitates in the as-deposited coating, the deposition chamber is additionally equipped with an additional heating source that can instantaneously heat the substrate surface to a temperature between 900° C. and 1000° C. in a few seconds up to a depth of a few nanometers, hence forth called as flash heating. The source of flash heating can be a nanosecond laser, electron heating with filament etc. Energy of the flash heater is tuned such a way that only the surface few nm are heated for every exposure. The chamber is designed such a way that the coated substrates is alternatively exposed to coating, and a flash heating source as schematically shown in FIG. 3c. When the coated substrate is exposed to the flash heating source, the local heating enables decomposition of a given alloy and results in a controlled precipitate structure. The composition of the matrix and precipitate are tuned by the choice of alloy composition, the energy input and time of the flash heating cycle following the thermodynamic rule of free energy minimization combined with the controlled kinetics available during the flash heating. The coating growth continues by coating deposition, and flash heating to form a HEA matrix and controlled precipitate structure in a sequential way.

(21) In the proposed method, formation of the precipitate structure and the host matrix is guided by alloy choice.

(22) a) Alloying elements that can form entropy stabilized solid solutions are screened using the first principles calculations similar to the example of Ti.sub.37Al.sub.34Si.sub.12V.sub.17N where the solid solution has lower ΔG.sub.mix relative to their all competing states.

(23) b) Precipitate structure are formed by having alloying elements/components like BN, Al.sub.2O.sub.3, and V.sub.2O.sub.5 that do not mix with the above mentioned HEA matrix by careful thermodynamic considerations. The necessary kinetics required to form the precipitates are provided either by post annealing or by flash heating in-situ during a coating deposition. Based on the selected alloy, and precipitate combination, a reactive atmosphere is created in the chamber comprising for example nitrogen and/or oxygen, and/or CH.sub.4 separately or a mixture of them as needed.

(24) Some examples, but not limited to are precipitates of BN, Al.sub.2O.sub.3, and VC in a HEA matrix of Ti.sub.37Al.sub.34Si.sub.12V.sub.17N to enable additional functional properties of the coating.