WO2016202892A1

WO2016202892A1 - Max phase ceramics and methods for producing the same

Info

Publication number: WO2016202892A1
Application number: PCT/EP2016/063810
Authority: WO
Inventors: Thomas LAPAUW; Jozef Vleugels; Konstantina LAMBRINOU
Original assignee: Katholieke Universiteit Leuven; Sck.Cen
Priority date: 2015-06-15
Filing date: 2016-06-15
Publication date: 2016-12-22
Also published as: GB201519387D0

Abstract

A method for producing a MAX phase material with general formulation M_n+1AX_n-y is disclosed. Thereby M is a transition metal, A is an A-group element and X is carbon (C) and/or nitrogen (N), where the M, A and X position can be occupied by more than 1 element with n being in the range 1 to 3, y being within the range -0.5 to 0.5. The method comprises providing a starting powder said starting powder comprising least one starting hydride powder, compacting the starting powder and producing the MAX phase material by using a pressure assisted sintering technique. Preferably Zr, Hf, Ti, Nb and Ta are used for M and Si, Sn, Pb, Bi and Al for A.

Description

MAX PHASE CERAMICS AND METHODS FOR PRODUCING THE SAME Field of the invention

The invention relates to the field of ceramics and more specifically MAX phase materials and methods for producing the same. Background of the invention

MAX phases are a group of ternary carbides and nitrides described by the general chemical formula M_n+iAX_n, where M corresponds to an early transition metal, A is an A-group element (mainly elements from groups 13-16 in the periodic table), X is C or N, and n is an integer equal to 1, 2 or 3. Based on the n-value, three stoichiometric '211', '312' and '413' types have been identified.

I n addition to the "pure" MAX phases that contain one of each of the M, A and X elements, the number of possible solid solutions is quite large. Solid solutions have been processed and characterized with substitution on M sites, e.g. (Ti,\/)₂AIC, (Ti,N b)₂AIC, (Ti,Cr)₂AIC, (Ti,Hf)₂lnC and (Ti,V)₂SC; on A sites, e.g. Ti₃(Si,Ge)C₂ and Ti₃(Sn,AI)C₂ and on X sites, e.g. Ti₂AI(C,N) and Ti₃AI(C,N)₂.

The growing interest in the MAX phases is due to their unique properties as these ternary compounds exhibit merits of both ceramics and metals, i.e. they have a good corrosion resistance, thermal stability, high flexural strength and exceptional fracture toughness for ceramic compounds. Furthermore, they are good electrical and thermal conductors and are machineable with conventional tools. This exceptional property combination results from their nano-layered crystal structure : MAX phases have a hexagonal unit cell (space group P63/mmc) that combines ΜεΧ octahedra with an intercalating A layer.

There is a need for new MAX phase materials with unique properties and methods for producing the same. Summary of the invention

It is an object of embodiments of the present invention to provide new MAX phase materials with good properties and methods for producing the same.

It is an advantage of embodiments of the present invention to provide new MAX phase materials which are suitable for advance nuclear systems.

It is an advantage of embodiments of the present invention to provide new MAX phase materials with a high fracture toughness.

It is an advantage of embodiments of the present invention to provide new MAX phase materials having a higher thermal stability among all reported Mn+lAXn phases. The above objective is accomplished by a method and device according to the present invention.

The present invention relates to a method for producing a MAX phase material with general formulation M_n+iAX_n-y, whereby M is a transition metal, A is an A-group element and X is a carbon C and/or nitrogen N, where the M, A and/or X position can be occupied by more than 1 element with n=l to 3, y=-0.5 to 0.5, said method comprising:

- providing a mixture, said mixture comprising a precursor of M, a precursor of A and a precursor of X, wherein the precursor of M is at least one hydride selected form the group consisting of M hydrides or wherein the precursor of M comprises hydride properties.

- producing the MAX phase material by using a pressure assisted sintering method. The mixture may be a starting powder, wherein the starting powder is compacted by pressure assisted densification.

In other embodiments, X may be a carbon C and/or nitrogen N and/or Boron B.

The method may comprise adding another transition metal to the mixture.

The precursor of M, the precursor of A and the precursor of X of the mixture may be mixed to obtain a near stoichiometric ratio with a slight excess of A and sub- stoichiometric amount of X.

The starting powder may be mixed in an organic fluid and said starting powder in in the organic fluid is dried before compacting the powder. Mixing said powders may comprise mixing with balls.

Compacting the powder mixture may comprise pouring the powder mixture into a die and cold-compacting said mixture resulting in a powder compact.

Compacting the starting powder may comprise providing a pressure in the range of 20 MPa to 40 MPa, preferably from 25 MPa to 35 MPa on the powder.

The organic fluid may be ethanol and the powder in said organic fluid may be mixed for at least 10 h.

The pressure assisted sintering technique may comprise reactive hot pressing with a heating rate of 5°C/min to 100°C/min up to the final sintering temperature of 1700°C. Upon reaching the final sintering temperature, a pressure may be provided increasing from 5 MPa to 30 MPa, preferably from 5 MPa to 20 MPa, optionally with a dwell time of 30 min.

The at least one hydride may be a Ti, Nb, Zr, Hf and/or Ta based hydride.

At least one TiH₂, NbHo.89, ZrH₂, HfH₂ or TaHx, hydride and Al and C may be used as a starting powder mixture.

The present invention also relates to a MAX phase material obtained by a method according to an embodiment as described above, wherein the MAX phase material is in the form of M_n+iAX_n-y with n being in the range 1 to 3, y being in the range -0.5 to 0.5:

- with M selected from the group consisting of Zr and Hf or

- with M occupied by at least two elements in the form (M_x, ΜΊ-_Χ) with M and M' selected form the group consisting of Ti, Nb, Ta, Zr and Hf, with M different from M', and wherein M' partially substitutes M in a M-A-X based system as a starting mixture; and with x= 0.01 to 1

- with A selected from the group consisting of Si, Sn, Pb and Bi or A occupied by at least two elements in the form (A, Al) selected from the group consisting of Si, Sn, Pb and Bi and

- X selected from the group consisting of B, C, and N or X occupied by at least two elements in the form (C, N).

M may comprise Nb. The MAX material may be a ternary carbide of the form (Ti_v,Nbw,Zrx,Hfy,Ta_z)n+i(AI,A)(C,N)n-y with v, w, x, y, z > 0, and v+w+x+y+z being equal to 1, n being within the range 1 to 3, y being in the range -0.5 to 0.5 and A Si, Sn, Pb or Bi. The present invention also relates to a MAX phase material with general formulation Mn+iAXn-y, whereby M is a transition metal, A an A-group element and X being carbon C and/or nitrogen N with n being in the range 1 to 3, y being in the range -0.5 to 0.5:,

- with M occupied by at least Zr and a 312 stoichiometry or

- with M occupied by at least Hf and a 211 or 312 stoichiometry or

- where M is occupied by at least two elements in the form of (M_x, ΜΊ-_Χ) with M and M' selected from the group consisting of Ti, Nb, Ta, Zr and Hf with M different from M', and wherein M' partially substitutes M in a M-A-X based system as a starting mixture and with x being in the range 0.01 to 1, and where the A and X position can be occupied by one or more elements.

A may be selected from the group consisting of Si, Sn, Pb and Bi or A occupied by at least two elements in the form (A, Al) selected from the group consisting of Si, Sn, Pb and Bi and X selected from the group consisting of B, C, and N or X occupied by at least two elements in the form (C, N).

M may comprise Nb.

The MAX material may be a ternary carbide of the form Zr3AIC₂.

The MAX material may be a ternary carbide of the form (Ti_v,Nbw,Zrx,Hfy,Ta_z)n+i(AI,A)(C,N)n with v, w, x, y, z > 0, and v+w+x+y+z being equal to 1, n being in the range 1 to 3 and A being an Si, Sn, Pb or Bi.

M may be occupied by at least two elements in the form of (M_x, ΜΊ-_Χ), has a 211, 312 or 413 stoichiometry.

The MAX phase material may be of the form (Nb_x, Zri-_x)_n+iAX_n-y with a 312 or 413 stoichiometry.

The MAX phase material may be a powder.

The MAX phase material may be a sintered material, a sintered bulk material and/or a spark plasma sintered material.

The MAX phase material may be a solid solution. The MAX phase material may be a pre-reacted powder compact.

The MAX phase material may be a coating material.

The MAX phase material may be a sintered or spark plasma sintered and hot pressed material.

The MAX phase material may be a sintered or spark plasma sintered and cold-pressed material.

The MAX phase material may be a powder sintered in the in the 1350°C to 1900°C temperature range under a pressure in the range of 5 M Pa to 100 M Pa, preferably 10 M Pa to 50 MPa, more preferably 28 MPa to 32 M Pa or about 30 MPa.

The MAX phase material may be a powder that is sintered in the 1450°C to 1725°C temperature range under a pressure in the range of 20 M Pa to 40 M Pa, preferably 25 MPa to 35 MPa, more preferably in the range of 28 M Pa to 32 M Pa or about 30 M Pa and a dwell time of 230 s to 2500 s, preferably 500 s to 2000 s, more preferably 1700 s to 1900 s.

The MAX phase material may be nano-layered.

The MAX phase material may comprise a partial alignment of the c-axis of the grains with the compression direction.

The MAX phase material may be a Nb- based ternary carbide , preferably of the form (N b_x,Zri-x)₄AIC3 (413 N bZ), the MAX phase material exhibiting at room temperature one of the following mechanical properties in any one of the following ranges: p_exp within 5 g/cm³ to 9 g/cm³, p_th within 5 g/cm³ to 8 g/cm³, GSL within 15 μιη to 28 μιη, GSw within 3 μηι to 6 μιη, HVio in the GPa range, E within 325 GPa to 365 GPa, o_4pt within 450 M Pa to 550 M Pa, Kic within 5 M Pa.m^1/2 to 15 M Pa.m^1/2 or a combination thereof. The MAX phase material may be a Nb-based ternary carbide, preferably of the form 413 NbZ, i.e. with a 413 stoichiometry, characterized in that said ternary carbide exhibits at room temperature one of the following mechanical properties in any one of the following ranges: p_exp within 6 g/cm³ to 7 g/cm³, p_th within 6 g/cm³ to 7 g/cm³, GSL within 21 μιη to 25 μιη, GSw within 5.5 μιη to 6.8 μιη, HVio in the GPa range, E within 340 GPa to 350 GPa, o_4pt within 500 M Pa to 515 M Pa, Kic being 10.1 M Pa.m^1/2 ± 0.3 M Pa.m^1/2 or a combination thereof. The MAX phase material may be a Nb-based ternary carbide, preferably of the form 413 NbZ, i.e. having a stoichiometry 413, characterized in that said ternary carbide exhibits at room temperature one of the following mechanical properties in any one of the following ranges: p_exp within 6.5 g/cm³ to 6.9 g/cm³, p_th within 6.8 g/cm³ to 6.9 g/cm³, GSL within 10 μιη to 30 μιη, GSw being within 6μιη to 6.6 μιη, HVio within the GPa range, E within 340 GPa to 355 GPa, o_4pt within 500 MPa to 600MPa (± 92 MPa), Kic being within 9 MPa.m^1/2 to 11 MPa.m^1/2or a combination thereof.

The MAX phase material may be a Nb-based ternary carbide, preferably of the form 413 NbZ, i.e. having a 413 stoichiometry, characterized in that said ternary carbide has lattice parameters in any one of the following ranges a_exp within 2.5 A to 3.6 A and/or Cexp within 20 A - 30 A or a combination thereof.

The MAX phase material may be a Nb-based ternary carbide, preferably of the form 413 NbZ, characterized in that said ternary carbide retains more than 80% of its stiffness at 1550°C.

The MAX phase material may be the ternary carbide (Nbo.85,Zro.is)4AIC3, characterized in that said ternary carbide has a higher thermal stability as compared to Nb₄AIC3. The MAX phase material may be a Nb-based ternary carbide, preferably of the form 413 NbZ, characterized in that said ternary carbide has a higher thermal stability among all reported M_n+iAX_n phases than reported M_n+iAX_n phases.

The MAX phase material may be a Nb-based ternary carbide, preferably of the form 413 NbZ, characterized in that said ternary carbide has a fracture toughness in the range of 8 MPa.m^1/2 to 12 MPa.m^1/2, preferably in the range of 9 MPa.m^1/2 to 11 MPa.m^1/2.

The MAX phase material may be a Nb-based ternary carbide, preferably of the form 413 NbZ, characterized in that said ternary carbide has a fracture toughness increased to 10.1 MPa.m^1/2 (±0.3 MPa.m^1/2).

The MAX phase material may be a Nb-based ternary carbide, preferably of the form 413 NbZ, characterized in that said ternary carbide has a flexural strength in the range of 400 MPa to 600 MPa. The method may comprise deforming a provided MAX phase material wherein the provided MAX phase material is compacted.

The deformation may be performed uniaxial at temperatures above 1200°C by hot pressing or spark plasma deformation.

The provided MAX phase material may be a MAX phase material as described above. The method further may comprise texturing the obtained MAX phase material using a method as described above.

The MAX phase material may comprise grains that show stronger alignment.

The MAX phase material may comprise an increased fracture toughness of 11.1 MPa.m^1/2 in comparison to the non-deformed material.

The present invention also relates to the use of a MAX phase material as described above, with M or M' comprising Zr, as fuel cladding material of light water reactors.

In another aspect the present invention provides ternary carbides with the general formulation (Nbi-_x,Zr_x)_n+iAIC_n-y, whereby x = 0.01 to 1, n = 2 to 3 and y = -0.5 to 0.5. In preferred embodiments the ternary carbide is in a solid solution.

In further preferred embodiments the ternary carbide is a powder.

In yet further preferred embodiments the ternary carbide is a pressureless sintered

(Nbi-x,Zr_x)n+iAICn-y material.

In preferred embodiments the ternary carbide is spark plasma sintered (Nbi-x,Zr_x)n+iAICn-y material.

In preferred embodiments the ternary carbide is sintered (Nbi-_x,Zr_x)_n+iAIC_n-y bulk material.

In preferred embodiments the ternary carbide is sintered (Nbi-_x,Zr_x)_n+iAIC_n-y coating material.

In preferred embodiments the ternary carbide is sintered or spark plasma sintered and hot pressed (Nbi-_x,Zr_x)_n+iAIC_n-y material.

In preferred embodiments the ternary carbide is sintered or spark plasma sintered and cold pressed (Nbi-_x,Zr_x)_n+iAIC_n-y material.

In preferred embodiments the ternary carbide is a (Nbi-_x,Zr_x)_n+iAIC_n-y powder sintered in the 1550°C to 1725°C temperature range. I n preferred embodiments the ternary carbide is a (N bi-_x,Zr_x)_n+iAIC_n-y powder sintered in the 1450°C to 1725°C temperature range under a pressure in the range of 5 M Pa to 100 M Pa, preferably within the range 10 M Pa to 50 M Pa, more preferably within the range 28 M Pa to 32 M Pa or about 30 M Pa.

I n preferred embodiments the ternary carbide is a (N bi-_x,Zr_x)n+iAICn-y powder that is sintered in the 1450°C to 1725°C temperature range under a pressure in the range of 20 M Pa to 40 MPa, preferably within the range 25 M Pa to 35M Pa, more preferably in the range of 28 MPa to 32 MPa or about 30 M Pa and a dwell time of 230 s to 2500 s, preferably between 500 s to 2000 s, more preferably between 1700 s to 1900 s.

I n preferred embodiments the ternary carbide is nano-layered.

I n preferred embodiments the ternary carbide has a partial alignment of the c-axis of the grains with the compression direction as shown by the XRD patterns in FIG. 3 of this application.

I n preferred embodiments the ternary carbide can be deformed at tem peratures above 1200°C by hot pressing or spark plasma deformation (as illustrated in FIG. 11).

I n preferred embodiments the ternary carbide is after deformation at temperatures above 1200°C by hot pressing of spark plasma deformation and that it comprises grains that show stronger alignment, as illustrated in FIG. 12.

I n preferred embodiments the ternary carbide has been deformed at temperatures above 1200°C by hot pressing of spark plasma deformation, and has an increased fracture toughness of 11.1 M Pa.m^1/2 in comparison to the non-deformed material.

I n preferred embodiments the preferably Nb-based ternary carbide exhibits at room temperature one of the following mechanical properties in any one of the following ranges: p_exp within 5 g/cm³ to 9 g/cm³, p_th within 5 g/cm³ to 8 g/cm³, GSL within 15 μιη to 28 μιη, GSw within 3 μιη to 6 μιη, HVio within the GPa range, E within 325 GPa to

365 GPa, o_4pt within 450 M Pa to 550 MPa, Kic within 5 M Pa.m^1/2 to 15 MPa.m^1/2or a combination thereof.

I n preferred embodiments the preferably Nb-based ternary carbide exhibits at room temperature one of the following mechanical properties in any one of the following ranges: p_exp within 6 g/cm³ to 7 g/cm³, p_th within 6 g/cm³ to 7 g/cm³, GSL within 21 μιη to 25 μιη, GSw within 5.5 μιη to 6.8 μιη, HVio within the GPa range, E within 340 GPA and 350 GPa, o_4pt within 500 MPa to 515 MPa, Kic being 10.1 MPa.m^1/2 ± 0.3 MPa.m^1/2 or a combination thereof.

In preferred embodiments the preferably Nb-based ternary carbide exhibits at room temperature one of the following mechanical properties in any one of the following ranges: p_exp within 6.5 g/cm³ to 6.9 g/cm³, p_th within 6.8 g/cm³ to 6.9 g/cm³, GSL within

10 μιη to 30 μιη, GSw within 6 μιη to 6.6 μιη, HVio within the GPa range, E within 340 GPa to 355 GPa, o_4pt within 500 MPa to 600 MPa (± 92 MPa), Kic within 9 MPa.m^1/2 to

11 MPa.m^1/2 or a combination thereof.

In preferred embodiments the ternary carbide has lattice parameters in any one of the following ranges : a_exp within 2.5 A to 3.6 A and/or c_exp within 20 A to 30 A or a combination thereof.

In preferred embodiments the ternary carbide retains more than 80% of its stiffness at 1550°C.

In preferred embodiments, the ternary carbide, (Nbo.85,Zro.is)4AIC3, according to any embodiments of the present invention, has a higher thermal stability as compared to Nb₄AIC₃.

Ternary carbides according to embodiments of the present invention advantageously have a higher thermal stability among all reported M_n+iAX_n phases than reported Mn+iAXn phases.

Ternary carbides according to embodiments of the present invention preferably have a fracture toughness in the range of 8 MPa.m^1/2 to 12 MPa.m^1/2, preferably in the range of 9 MPa.m^1/2 to 11 MPa.m^1/2.

Ternary carbides according to embodiments of the present invention have a fracture toughness increased to 10.1±0.3 MPa.m^1/2.

Ternary carbides according to embodiments of the present invention have a flexural strength in the range of 400 MPa to 600 MPa.

In another aspect the invention provides production of a (Nbi-_x,Zr_x)n+iAICn-y, according to embodiments of the present invention, starting from Nb and Zr hydrides. Zr could partially substitute Nb in Nb-AI-C based MAX phases. In embodiments of the invention the production of a (Nbi-_x,Zr_x)_n+iAIC_n-y, according to embodiments of the invention comprises the following steps: 1) adding the NbHo.89, ZrH₂, Al and C powders to a reactor.

In preferred embodiments the production of a (Nbi-_x,Zr_x)_n+iAIC_n-y, according to embodiments of the present invention, comprises the following steps:

- (Nbi-_x,Zr_x)n+iAIC_n-y starting powders are mixed in a near stoichiometric ratio, with a Nb/Zr ratio,

- adding a slight excess of Al and sub-stoichiometric amount of C (to prevent the formation of binary transition metal carbides),

- mixing the powders in an organic fluid,

- drying the powder,

- precompacted (with a pressure of 30 MPa), and

- producing the (Nbi-_x,Zr_x)_n+iAIC_n-y phase by reactive hot pressing sintering/spark plasma sintering.

In preferred embodiments the production of a (Nbi-_x,Zr_x)_n+iAIC_n-y, according to embodiments of the present invention comprises the following steps:

- (Nbi-_x,Zr_x)n+iAICn-y-starting powder is mixed in a near stoichiometric ratio, with a Nb/Zr ratio in the range of 100/0 to 50/50,

- adding a slight excess of Al and sub-stoichiometric amount of C to prevent the formation of binary transition metal carbides,

- mixing the powders for at least 10 h in ethanol whereby agglomerate formation is prevented, for instance, by mixing with balls,

- drying the powder,

- precompaction with a pressure in the range of 20 MPa to 40 MPa, preferably from 25 MPa to 35 MPa, 5) producing the (Nbi-_x,Zr_x)_n+iAIC_n-y phase by reactive hot pressing with a heating rate was set to 20°C/min up to the final sintering temperature of 1700°C,

- sintering whereby upon reaching the sintering temperature, increasing from 5 MPa to 20 MPa, for instance, with a dwell time of 30 min.

In yet another aspect the present invention provides ternary carbides with the general formulation (Nbi-_x,T_x)_n+iAIC_n-y and T₂AICi-y ceramic materials with x = 0.01 to 1, n = 2 to 3, y = -0.5 to 0.5 and T be Zr and/or Hf. In preferred embodiments the ternary carbide is in a solid solution. In other preferred embodiments the ternary carbide is a powder. Embodiments of the present invention provide a ternary carbide which is a sintered (Nbi-x,Tx)n+iAIC_n-y or T2AICi-_y material.

In embodiments the ternary carbide is a spark plasma sintered (Nbi-x,Tx)_n+iAIC_n-y or T2AICi-_y material.

In embodiments the ternary carbide is a sintered (Nbi-x,Tx)_n+iAIC_n-y or T2AICi-_y bulk material.

In embodiments the ternary carbide is a sintered (Nbi-x,Tx)_n+iAIC_n-y or T2AICi-_y coating material.

In preferred embodiments the ternary carbide is a sintered or spark plasma sintered and hot pressed (Nbi ;Τχ)η+ iAICn- y or T2AICi-_y material.

In preferred embodiments the ternary carbide is sintered or spark plasma sintered and cold pressed (Nbi-x,Tx)_n+iAIC_n-y or T2AICi-_y material.

In preferred embodiments the ternary carbide is a (Nbi-x,Tx)_n+iAIC_n-y or T2AICi-_y powder sintered in the 1550°C - 1725°C temperature range.

In preferred embodiments the ternary carbide is a (Nbi-x,Tx)_n+iAIC_n-y or T2AICi-_y powder sintered in the 1350°C - 1725°C temperature range under a pressure in the range of 5 MPa - 100 MPa, preferably within 10 MPa to 50 MPa, more preferably within 28 MPa to 32 MPa or about 30 MPa.

In preferred embodiments the ternary carbide is a (Nbi-x,Tx)_n+iAIC_n-y or T2AICi-_y powder that is sintered in the 1450°C to 1725°C temperature range under a pressure in the range of 20 MPa to 40 MPa, preferably within 25 MPa to 35 MPa, more preferably in the range of 28 MPa to 32 MPa or about 30 MPa and a dwell time within a range of 230 s to 2500 s, preferably within 500 s to 2000 s, more preferably within 1700 s to 1900 s.

In preferred embodiments the ternary carbide is nano-layered.

In preferred embodiments the ternary carbide has a partial alignment of the c-axis of the grains with the compression direction as shown by the XRD patterns in Figure 3 of this application. I n preferred embodiments the ternary carbide can be deformed at tem peratures above 1200°C by hot pressing or spark plasma deformation (as illustrated in Figure 11). I n preferred embodiments the ternary carbide is after deformation at temperatures above 1200°C by hot pressing of spark plasma deformation and that it comprises grains that show stronger alignment, as illustrated in Figure 12.

I n preferred embodiments the ternary carbide has been deformed at temperatures above 1200°C by hot pressing of spark plasma deformation, and has an increased fracture toughness of 11.1 M Pa.m^1/2 in comparison to the non- deformed material. I n preferred embodiments the preferably Nb-based ternary carbide exhibits at room temperature one of the following mechanical properties in any one of the following ranges: p_exp (g/cm³) 5 to 9, p_th(g/cm³) 5 to 8, GSL^ITI) 15 to 28, GSw (μιη) 3 to 6, HVio (GPa),E (GPa) 325 to 365, o_4pt (MPa) 450 to 550, K,_c (M Pa.ml/2) 5 to 15 or a combination thereof.

I n preferred embodiments the preferably Nb-based ternary carbide exhibits at room temperature one of the following mechanical properties in any one of the following ranges: p_exp (g/cm³) 6 to 7,p_th (g/cm³) 6 to 7,GSL (μιη) 21 to 25, GSw (μιη) 5.5 to 6.8, HVio (GPa), E (GPa) 340 to 350, o_4pt (M Pa) 500 to 515, K,_c (M Pa.m^1/2) 10.1 ± 0.3 or a combination thereof.

I n preferred embodiments the preferably Nb-based ternary carbide exhibits at room temperature one of the following mechanical properties in any one of the following ranges: p_ex (g/cm³) 6.5 to 6.9, p_th (g/cm³) 6.8 to 6.9, GSL (μιη) 10 to 30, GSw (μιη) 6 to 6.6, HVio (GPa), E (GPa) 340 to 355, o_4pt (MPa) 500 to 600 ± 92, K,_c (M Pa. ml/2) 9 to 11 or a combination thereof.

The ternary carbide according to embodiments of the present invention characterized in that said ternary carbide retains more than 80% of its stiffness at 1550°C.

The ternary carbide, N bo.85,Zro.is)4AIC3, characterized in that said ternary carbide has a higher thermal stability as compared to Nb₄AIC3.

The ternary carbide according to embodiments of the present invention, characterized in that said ternary carbide has a higher thermal stability among all reported M_n+iAX_n phases than reported M_n+iAX_n phases. The ternary carbide according to embodiments of the present invention, characterized in that said ternary carbide has a fracture toughness in the range of 8 to 12 MPa.m^1/2, preferably in the range of 9 to 11 MPa.m^1/2.

The ternary carbide according to embodiments of the present invention, characterized in that said ternary carbide has a fracture toughness increased to 10.1±0.3 MPa.m^1/2. The ternary carbide according to embodiments of the present invention, wherein the ternary carbide has lattice parameters in any one of a_exp (A) 2.5 to 3.6 and/or Cexp (A) 20 to 30 or a combination thereof.

The ternary carbide according to embodiments of the present invention, characterized in that said ternary carbide has a flexural strength in the range of 400 to 600 MPa.

In another aspect the present invention provides production of (Nbi-_x,Tx)n+iAIC_n-y or T2AICi-_y according to embodiments of the present invention, starting from Nb, Zr and/or Hf hydrides. Zr could partially substitute Nb in Nb-AI-C based MAX phases. The production of a (Nbi-_x,Tx)n+iAIC_n-y or T2AICi-_y according to embodiments of the present invention, whereby the process comprises the following steps: 1) adding the NbHo.89, ZrH₂, HfH₂, Al and C powders to a reactor.

The production of a (Nbi-_x,Tx)n+iAIC_n-y or T₂AICi-y according to embodiments of the present invention, the production process comprising the following steps:

- (Nbi-_x,Tx)n+iAICn-yor T₂AICi-y starting powders are mixed in a near stoichiometric ratio, - adding a slight excess of Al and sub-stoichiometric amount of C (to prevent the formation of binary transition metal carbides),

- mixing said powders in an organic fluid,

- drying the powder,

- precompacted (with a pressure of 30 MPa), and

- producing the (Nbi-x,Tx)_n+iAIC_n-y or T₂AICi-_y phase by reactive hot pressing sintering / spark plasma sintering.

The production of (Nbi-x,Tx)_n+iAIC_n-y or T₂AICi-_y according to embodiments of the present invention, the production process comprising the following steps:

- (Nbi-x,Tx)n+iAIC_n-y or T₂AICi-_y-sta rti ng powder is mixed in a near stoichiometric ratio, - adding a slight excess of Al and sub-stoichiometric amount of C to prevent the formation of binary transition metal carbides,

- mixing said powders for at least 10 h in ethanol whereby agglomerate formation is prevented for instance by mixing with balls,

- drying the powder,

- precompaction with a pressure in the range of 20MPa to 40 MPa, preferably from 25 MPa -to35 MPa,

- producing the (Nbi-_x,Tx)n+iAIC_n-y or T2AICi-_y phase by reactive hot pressing with a heating rate was set to 5°C/min to 100°C/min up to the final sintering temperature of 1700°C,

- sintering whereby upon reaching the sintering temperature, increasing from 5 MPa to 20 MPa for instance with a dwell time of 30 min.

Embodiments of the present invention provide methods for producing new materials by reactive hot pressing and starting from at least one hydride, like e.g. starting from niobium (Nb) and zirconium (Zr) hydrides. Furthermore, this approach can be extended towards more complex solid solution systems: (Ti_v,N bw,Zrx,Hfy,Ta_z)n+i(AI,A)(C,N)n with v, w, x, y, z >0 and v+w+x+y+z = 1, n = 1 to 3 and A an Si, Sn, Pb or Bi where many new solid solutions were found stable and several new ternary MAX phase end members were discovered. For the latter, most of them were only reported as hypothetical before. Embodiments of the present invention preferably use fine Ti, Nb, Ta, Zr and Hf hydride powders in combination with a pressure assisted sintering technique. As an example of the improved properties, the influence of Zr and Hf on the mechanical properties and temperature stability can be investigated for the case of the 413 phase. Embodiments of the present invention provide high-purity solid solutions by using reactive hot pressing of NbHo.89, ZrH₂, Al, and C as starting powder mixtures. It is an advantage of embodiments of the present invention that an appreciable increase in fracture toughness was observed from 6.6 ± 0.1 MPa/m^1/2 for pure Nb₄AIC3 to 10.1 ± 0.3 MPa/m^1/2 for the (Nbo.85, Zro.i5)₄AIC3 solid solution according to embodiments of the present invention. 413 MAX phases according to embodiments of the present invention are preferably synthesized in bulk group. 413 MAX phases may be referred to as MAX phases having a 413 stoichiometry. The interest in the 413 subgroup is justified by the fact that this stoichiometry shows typically an enhanced temperature stability compared to those of the other stacking sequences in the same chemical system.

A (Nb,Zr)₄AIC3 (NZ413) structure according to embodiments of the present invention has a remarkably better temperature stability for instance being stable above 1000°C. The preferably solid solution may be prepared starting from Nb- and Zr-hydrides and their lattice parameters may be evaluated over the entire solubility range. Furthermore, this approach could be extended towards more complex solid solution systems: (Ti_v,Nb_w,Zrx, Hfy,Ta_z)n+i(AI,A)(C,N)n with v, w, x, y, z >0, v+w+x+y+z = 1, n = 1 to 3 and A an Si, Sn, Pb or Bi. Many new solid solutions were found stable and several new ternary MAX phase end members were discovered. For the latter, most of them were only reported as hypothetical before. The success of this synthesis procedure is based on the usage of fine at least one Ti, Nb, Ta, Zr and Hf hydride powder in combination with a pressure assisted sintering technique.

In one aspect of the invention, a ternary carbide with the general formulation (Ti_v,Nbw,Zrx, Hfy,Ta_z)n+i(AI,A)(C,N)n With v, w, x, y, z >0, v+w+x+y+z = 1, n = 1 to 3 and A an Si, Sn, Pb or Bi according to embodiments of the present invention is in any of the following forms, a solid solution, a powder, a pre-reacted powder compact, a sintered (Ti_v,Nbw,Zrx, Hfy,Ta_z)n+i(AI,A)(C,N)n, a sintered and hot-pressed or spark plasma-sintered (Ti_v,Nbw,Zrx, Hfy,Ta_z)n+i(AI,A)(C,N)n material or a hot-pressed or spark plasma-sintered (Ti_v,Nbw,Zrx, Hfy,Ta_z)n+i(AI,A)(C,N)n material or a sintered

(Ti_v,Nbw,Zrx, Hfy,Ta_z)n+i(AI,A)(C,N)n material that was cold-pressed (Ti v, N bw,Zr_x, Hfy,Ta_z)_n+i( Al, A)(C, N )_n.

Embodiments of the present invention result in that up to 17 to 19%, preferably up to 18 to 19%, more preferably up to 18,2 to 18,7% and most preferably 18,5% or about 18,5% of Nb atoms in a Nb-AI-X system could be substituted by e.g. Zr preferably in the 413-structure stoichiometry. A 413 stoichiometric (Nb_x,Zri-_x)₄AIC3 (referred to as NZ 413) structure according to embodiments of the present invention advantageously comprises a remarkably better temperature stability, for instance above 1000°C. The solid solutions were prepared starting from Nb and Zr hydrides. During this procedure, the synthesis of Zr3AIC₂ was achieved. The lattice parameters of the solid solutions were evaluated over the entire solubility range. Embodiments of the present invention provide an enhanced formation and grain growth of (Nbo.85,Zro.is)4AIC3.

This approach may be extended towards more complex solid solution systems: (Ti_v,Nbw,Zrx,Hfy,Ta_z)n+i(AI,A)(C,N)n with v, w, x, y, z > 0, v+w+x+y+z = 1, n = 1 to 3 and A an Si, Sn, Pb or Bi where many new solid solutions were found stable and several new ternary MAX phase end members were discovered. Some of the newly synthesized phases according to embodiments of the present have n = 1, these Al-based 211 MAX phases are known for their excellent oxidation resistance as is stated in literature for Ti₂AIC and Cr₂AIC.

The mechanical properties of (Nbo.85,Zro.is)4AIC3, a MAX phase according to embodiments of the present invention, did not show significant change in hardness and stiffness compared to pure Nb₄AIC3: the measured fracture toughness of (Nbo.85,Zro.i5)4AIC3 was 10.1± 0.3 MPa.m^1/2 and the measured flexural strength was 507± 92 MPa. However, the temperature stability of (Nbo.85,Zro.is)4AIC3 was found to be exceptionally high and the synthesized material retained more than 80% of its Young's modulus at 1550°C.

Max phase or ceramics according to embodiments of the present invention can be used as high-temperature conductive die/punch materials; components that must perform in heavy liquid metal environments (e.g. liquid lead and lead-bismuth eutectic alloys), such as pump impellers or coatings for nuclear fuel cladding materials; turbine blades; precursor for other novel materials such as new MXenes, new ultra-high temperature ceramics with melting point close to 4000°C, etc.

Particular and preferred aspects of the invention are set out in the accompanying independent and dependent claims. Features from the dependent claims may be combined with features of the independent claims and with features of other dependent claims as appropriate and not merely as explicitly set out in the claims. These and other aspects of the invention will be apparent from and elucidated with reference to the embodiment(s) described hereinafter.

Brief description of the drawings

FIG. 1 illustrates XRD patterns of the (Nb_x,Zri-_x)2AIC samples obtained using embodiments of the present invention. The dashed identification lines correspond to the peaks of the sample with the highest Zr-content.

FIG. 2 illustrates an elemental mapping of the (Nbo.8,Zro.2)2AIC sample obtained using embodiments of the present invention, indicating the higher solubility of Zr in NZ413 than in NZ211.

FIG. 3 illustrates XRD patterns of the (Nb_x,Zri-_x)₄AIC3 samples obtained using embodiments of the present invention. The dashed identification lines correspond to the peaks of the sample with the highest Zr-content.

FIG. 4 illustrate the lattice parameters as a function of the added amount of Zr in the starting powder. The linear trend of Vegard's law is followed up to an amount of x = 18%.

FIG. 5 illustrates an elemental mapping of (a) the (Nbo.85,Zro.is)4AIC3 sample according to embodiments of the present invention displaying the small impurities AI2O3 and the Al-rich intermetallic phase at the grain boundaries and (b) the (Nbo.5,Zro.s)4AIC3 sample according to embodiments of the present invention, showing NZC particles in between NZ413 grains.

FIG. 6 illustrates scanning electron micrographs at the tip of a Vickers' indent for (a) Nb₄AIC3 and (b) (Nbo.85,Zro.is)4AIC3 MAX phases obtained by embodiments of the present invention.

FIG. 7 illustrates the Young's moduli of the Nb₄AIC3 and Nbo.85,Zro.is)4AIC3 phase according to embodiments of the present invention as compared to known commercial materials (such as e.g. Maxthal 211^® and Maxthal 312^®) and the elastic damping plotted as function of temperature. FIG. 8 illustrates the Young's modulus of the (Nbo.85,Zro.is)4AIC3 phase according to embodiments of the present invention as function of temperature compared to some refractory metals [Modified from www.plansee.com - 11/06/2015].

FIG. 9 illustrates the XRD pattern of the ceramic with Zr₄AIC3 starting composition according to embodiments of the present invention processed at 1550°C. Three phases could be identified as Zr3AIC₂, ZrC, ZrAI₂.

FIG. 10 illustrates scanning electron micrographs from the tip of a Vickers' indent for the Zr3AIC₂ phase according to embodiments of the present invention illustrating the nano-layered structure of this phase.

FIG. 11 illustrates a deformation method or procedure according to embodiments of the present invention, the procedure comprising following steps: (a) powder compaction, (b) densification of the powder by SPS, (c) loading the grinded disc in a larger die, and (d) deformation by hot compression in SPS.

FIG. 12 illustrates the results of the EBSD analysis of the deformed Nb₄AIC3 according to embodiments of the present invention.

FIG. 13 illustrates the XRD pattern of the Zr-AI-C ceramic according to embodiments of the present invention processed at 1525°C with Zr₂AIC as main component.

FIG 14 illustrates XRD patterns of Hf-AI-C ceramics according to embodiments of the present invention processed at 1550°C wherein (a) Hf₂AIC and (b) Hf3AIC₂ can be identified.

FIG. 15 illustrates the solid solubility range of Zr in Nb₄AIC3 investigated over the entire composition range.

FIG. 16 illustrates the solid solubility range of Zr in Nb₂AIC investigated over the entire composition range.

The drawings are only schematic and are non-limiting. In the drawings, the size of some of the elements may be exaggerated and not drawn on scale for illustrative purposes. Any reference signs in the claims shall not be construed as limiting the scope.

In the different drawings, the same reference signs refer to the same or analogous elements. Detailed description of illustrative embodiments

The present invention will be described with respect to particular embodiments and with reference to certain drawings but the invention is not limited thereto but only by the claims. The drawings described are only schematic and are non-limiting. In the drawings, the size of some of the elements may be exaggerated and not drawn on scale for illustrative purposes. The dimensions and the relative dimensions do not correspond to actual reductions to practice of the invention.

Furthermore, the terms first, second and the like in the description and in the claims, are used for distinguishing between similar elements and not necessarily for describing a sequence, either temporally, spatially, in ranking or in any other manner. It is to be understood that the terms so used are interchangeable under appropriate circumstances and that the embodiments of the invention described herein are capable of operation in other sequences than described or illustrated herein.

Moreover, the terms top, under and the like in the description and the claims are used for descriptive purposes and not necessarily for describing relative positions. It is to be understood that the terms so used are interchangeable under appropriate circumstances and that the embodiments of the invention described herein are capable of operation in other orientations than described or illustrated herein.

It is to be noticed that the term "comprising", used in the claims, should not be interpreted as being restricted to the means listed thereafter; it does not exclude other elements or steps. It is thus to be interpreted as specifying the presence of the stated features, integers, steps or components as referred to, but does not preclude the presence or addition of one or more other features, integers, steps or components, or groups thereof. Thus, the scope of the expression "a device comprising means A and B" should not be limited to devices consisting only of components A and B. It means that with respect to the present invention, the only relevant components of the device are A and B.

Reference throughout this specification to "one embodiment" or "an embodiment" means that a particular feature, structure or characteristic described in connection with the embodiment is included in at least one embodiment of the present invention. Thus, appearances of the phrases "in one embodiment" or "in an embodiment" in various places throughout this specification are not necessarily all referring to the same embodiment, but may. Furthermore, the particular features, structures or characteristics may be combined in any suitable manner, as would be apparent to one of ordinary skill in the art from this disclosure, in one or more embodiments.

Similarly, it should be appreciated that in the description of exemplary embodiments of the invention, various features of the invention are sometimes grouped together in a single embodiment, figure, or description thereof for the purpose of streamlining the disclosure and aiding in the understanding of one or more of the various inventive aspects. This method of disclosure, however, is not to be interpreted as reflecting an intention that the claimed invention requires more features than are expressly recited in each claim. Rather, as the following claims reflect, inventive aspects lie in less than all features of a single foregoing disclosed embodiment. Thus, the claims following the detailed description are hereby expressly incorporated into this detailed description, with each claim standing on its own as a separate embodiment of this invention.

Furthermore, while some embodiments described herein include some but not other features included in other embodiments, combinations of features of different embodiments are meant to be within the scope of the invention, and form different embodiments, as would be understood by those in the art. For example, in the following claims, any of the claimed embodiments can be used in any combination.

In the description provided herein, numerous specific details are set forth. However, it is understood that embodiments of the invention may be practiced without these specific details. In other instances, well-known methods, structures and techniques have not been shown in detail in order not to obscure an understanding of this description.

Where in embodiments of the present invention reference is made to a parameter "p within a range a to b" or "p=a-b", reference is made to p being an element of the interval [a,b]. Where in embodiments of the present invention reference is made to "A occupied by at least two elements in the form (A, Al) selected from the group consisting of Si, Sn, Pb and Bi" reference is made to the fact that A is occupied by at least two elements, one element being in general an element of the A group and the second element being Al.

Embodiments of the present invention provide ternary carbides with the general formulation (Nbi-_x,Mx)n+iAICn-_y and M_nAICi-_y ceramic materials with x = 0.01 to 1, n = 2 to 3, y = -0.5 to 0.5 and the transition metal M Zr and/or Hf.

In specific embodiments a Zr₂AIC MAX phase can be synthesized by means of reactive hot pressing of at least a hydride, e.g. ZrH₂, Al and C as starting powder mixture. Embodiments of the present invention preferably use a relatively narrow temperature window for the synthesis of Zr₂AIC. ZrC was always present as a secondary phase by hot pressing in the 1475°C to 1575°C range. These Zr-based materials are of great interest for the nuclear industry because Zr atoms have a small cross-section for thermal neutrons. Apart from this economical argument, the fuel cladding materials of next generation (Gen-ll 1+) light water reactors (LWRs) must withstand severe operating conditions where mechanical and thermal loads are combined with high neutron irradiation doses and strongly oxidative or corrosive environments. Based on their superior properties, MAX phases are considered as candidate materials for fuel cladding applications, either in bulk form or as coatings. For the high-temperature steam environment in Gen-111+ LWRs during a loss of coolant (LOCA) accident, Zr₂AIC is one of the MAX phases of interest that could potentially surpass the performance of the commercial zircaloy dads.

Embodiments of the present invention advantageously provide a method to synthesis Zr₂AIC.

Example of a method to synthesize Zr2AIC according to embodiments of the present invention

A hydride, e.g. ZrH₂ (grain size < 6 μιτι, > 99% purity, Chemetall, Germany), Al ( < 5 μιτι, > 99% purity, AEE, US) and C ( < 5 μιτι, > 99% purity, Asbury Graphite Mills, US) powders were used as starting materials for the synthesis of the MAX phases. The powders were mixed in a Zr:AI:C molar ratio of 50:20:30 (corresponding to a 2:0.8:1.2 stoichiometry) in a Turbula multidirectional mixer. The original intent was to synthesize Zr3AIC₂, with this starting stoichiometry (equivalent to 3:1.2:1.8). Zr0₂ (Tosoh 3Y-TZP, 5 mm 0) milling balls and isopropanol were added in order to homogenize the mixing and to break-up soft agglomerates. After drying, the powder mixture was poured into a 30 mm inner-diameter graphite die and cold- compacted at 20 MPa. Subsequently, the die/punch/powder set-up was hot pressed (W100/150-2200-50 LAX, FCT Systeme, Frankenblick, Germany) in a vacuum environment at a heating rate of 25°C/min up to 1475, 1525 or 1575°C. Once the processing temperature was reached, the samples were held at temperature for 0.5 h. The initially-applied load of 7 MPa was increased to 20 MPa upon reaching the dwell temperature.

The outer ~lmm-thick layer was preferably ground off from the hot pressed disc prior to mechanical polishing for X-ray diffraction (XRD) characterization and microstructural analysis. XRD diffractograms were obtained from the polished top and cross-sectional surfaces using Cu Ka radiation in a Bruker D8 advance diffracto meter operated at 40 kV and 40 mA in the Bragg-Brentano geometry with a divergence slit of 0.4°. For room temperature characterization, 2 step intervals of 0.005° were applied from 8 to 158° with a counting time of 3 s per step. The lattice parameters of the top and cross-section of the disk-shaped sample were determined. No statistically- significant difference between those two surfaces was found and the reported value is the average of both. Rietveld refinements of the diffraction patterns were performed using the Materials Analysis Using Diffraction (MAUD) software.

Electron probe microanalysis (EPMA, JXA-8530F, JEOL Ltd., Japan) was used for microstructural and chemical analysis. The Zr:AI ratio in the MAX phase grains was determined by quantitative energy dispersive X-ray spectrometry (EDS, EDAX, US). Furthermore, the elemental distribution of Zr, Al, and C in the sintered ceramics was mapped. The beam current and accelerating voltage were fixed at 15 nA and 15 kV.

High-resolution transmission electron microscopy (HRTEM), EDS and selected area diffraction (SAED) were performed using a FEI Tecnai G2 TF20 UT equipped with a field emission gun operating at 200 kV with a point resolution of 0.19 nm. The TEM sample was prepared by embedding manually crushed powder obtained from the hot pressed samples in a Ti grid with a carbon-based glue. The sample was then mechanically polished down to 50m followed by ion milling to reach electron transparency.

Neutron powder diffraction (NPD) experiments were done on the KARL double axis diffractometer, mounted on the Israeli Research Reactor 1. The measurements were performed at room temperature (RT) with an incident neutron wavelength of 0.982(2) A. This low incident wavelength, combined with an angular step of 0.05°, generated sufficient angular range and angular resolution. A powder sample of ~5 g, taken from the sample hot pressed at 1525°C, was loaded into a cylindrical vanadium sample holder, which was used to significantly reduce coherent scattering from the holder. The results were analyzed using the Rietveld refinement method with the FullProf software package.

X-ray photoelectron spectroscopy (XPS) employing monochromatic Al Ka radiation (h =1486.6eV) was used to determine compositions in the surface region of a powder sample. Prior to the analysis, samples were sputter-cleaned in-situ with 4 keV Ar+ ions incident at an angle of 70° with respect to the surface normal for 10 min. Sputtering was performed until a steady-state (i.e., minimizing surface oxygen contaminations in the powder) was observed for the core levels. Deconvolution and quantification was performed using the CasaXPS software with elemental sensitivity factors supplied by Kratos Analytical Ltd.

The hardness was measured using a Vickers indenter (FV-700, Future-Tech Corp., Tokyo, Japan) and an indentation load of 30 N was applied for 10 s on a polished surface. The reported value is the average of 5 indents.

XRD patterns of the top surface of the samples that were reaction hot pressed in the 1475-1575°C range with a Zr:AI:C starting powder molar ratio of 50:20:30 are compared in Fig. 13. At 1475°C, the main compound was the binary carbide ZrC_x. The stoichiometry of the binary carbide can vary between ZrCo.99 and ZrCo.55. Additionally, Zr₂AIC was detected as a secondary phase together with the intermetallic Zr₂Al3. When the synthesis temperature increased to 1525°C, the amount of Zr₂AIC increased significantly, the intermetallic phase disappeared and the ZrC_x content decreased. Based on this observation, the following formation reaction may be proposed: Zr₂Al3

Embodiments of the present invention enable synthetization of the Zr₂AIC MAX phase by the reactive hot pressing of a ZrH₂, C, and Al starting powder mixture with preferably a Zr:AI:C molar ratio of 50:20:30. The optimal synthesis temperature was found to be 1525°C. Experimental investigation revealed that ZrC may be present as a secondary phase. At lower temperatures, Zr₂Al3 is identified as the Zr₂AIC-forming intermetallic. Between 1525 and 1575 °C there is a transition towards the Zr3AIC₂ phase, which appears to decompose into ZrAI₂ and ZrC starting at around 1575°C. The Zr₂AIC atomic structure revealed a 211-type atomic stacking. The a and c lattice parameters were 3.3237(2) A/3.3239(4) A and 14.5705(4) A/14.556(2) A, respectively. The Vickers hardness of the Zr₂AIC ceramic with 28 vol% ZrC, measured under a load of 30 N, was 6.4 ± 0.1 GPa. It is worthwhile mentioning that two new phases (i.e., Zr₂AIC and Zr3AIC₂) could be identified in a rather well-known system as Zr-AI-C.

Example of a method to synthesize Zr3AIC2 according to embodiments of the present invention

The starting powders were the hydride ZrH₂ (particle size <6 μιη, >99% purity, Chemetall, Germany), Al (particle size <5 μιη, >99% purity, AEE, US) and C (particle size <5 μιη, >99% purity, Asbury Graphite Mills, US). The powders were mixed in a stoichiometric ratio of 3:0.94:1.95 (equivalent to 4:1.25:2.6). The original intent was to synthesize the 413 MAX phase. As with most other Al-containing MAX phases, an excess of Al was added to compensate for its loss during processing. The sub- stoichiometric C content was chosen for two reasons: the first was to compensate for any C inward diffusion from the graphite dies/punches and the second to take into account the fact that most 413 MAX phases are C-deficient. Said otherwise, the C- content in the 413 phases is typically less than 3. The starting powders were mixed on a Turbula multidirectional mixer for 24 h in ethanol. Five millimeter diameter Zr0₂ balls (Grade TZ-3Y, Tosoh, Japan) were employed to break up agglomerates and mix the powders. After drying, the powder mixture was pre-compacted at 20 MPa in a 56 mm diameter graphite die to about a 10 mm-high disc. The latter was placed in a hot press (HP) (W100/150-2200-50 LAX, FCT Systeme, Frankenblick, Germany) and heated to the desired temperature. An optical pyrometer may be used to measure the temperature on the outer side of the graphite dies. The heating rate was set to 25°C/min up to the final sintering temperature of 1500°C. Upon reaching 1500°C, the applied pressure in the HP was increased from 5 MPa to 20 MPa followed by a dwell time of 30 min. After cooling in the HP, the disc was removed from the die and was ground to remove any outer reaction layers formed on the disc surface as a result of reaction with the graphite dies.

Figure 9 illustrates the XRD pattern of a Zr₄AIC3 ceramic according to embodiments of the present invention processed at 1550°C. Three phases could be identified as Zr3AIC₂, ZrC, ZrAI₂. Embodiments of the present invention provide synthesis of hexagonal Zr3AIC₂ of space group P63/mmc. In addition to the Zr3AIC₂ predominant phase, the produced ceramic contained ZrC and Zr-AI intermetallics as secondary phases. Figure 10 illustrates scanning electron micrographs from the tip of a Vickers' indent for Zr3AIC₂ according to embodiments of the present invention illustrating the nano-layered structure of this phase.

Example of a method to synthesize Hf-AI-C ceramics processed at 1550°C according to embodiments of the present invention wherein (a) Hf2AIC and (b) fsAld can be identified.

Figure 14 illustrates the XRD patterns of Hf-AI-C ceramics processed at 1550°C using methods according to embodiments of the present invention wherein (a) Hf₂AIC and (b) Hf3AIC₂ can be identified.

Examples of a ternary carbide (Ti_v,Nb_w,Zrx,Hfy,Ta_z)n+i(AI,A)(C,N)n MAX phases with v, w, x, y, z≥0, v+w+x+y+z = 1, n = 1-3 and A an Si, Sn, Pb or Bi according to embodiments of the present invention. Table 1 below illustrates the quaternary (M,M')₄AX3 phases, wherein the first rows are three known in the art and the last has been synthesized by methods according to the invention.

Table 1

The (Nb_x,Zri-x)₂AIC (NZ211) and (Nb_x,Zri-x)₄AIC3 phases (NZ413) according to embodiments of the present invention can be synthesized starting from at least one hydride, e.g. NbHo.89 (particle size < 40 μιη) and ZrH₂ (particle size < 6 μιη, > 99% purity, Chemetall), Al (particle size < 5 μιη, > 99% purity, AEE) and C (particle size < 5 μιη, > 99% purity, Asbury) powders. NbHo.89 was produced starting from coarse Nb (particle size < 300 μιη, > 97% purity, CBMM) powder. The powder was hydrogenated at 800°C in a dynamic H₂ atmosphere and once cooled down, planetary ball-milled (PM 400, Retch) and sieved. The (Nb_x,Zri-_x)₂AIC-starting powders were mixed in a near stoichiometric ratio, with a Nb/Zr ratio of 100/0 and 80/20. The (Nb_x,Zri-x)₄AIC3-starting powders were mixed in a near stoichiometric ratio, with a Nb/Zr ratio of 100/0, 95/5, 90/10, 85/15, 80/20 and 50/50. In both cases, a slight excess of Al and a sub- stoichiometric amount of C were added to prevent the formation of binary transition metal carbides. Mixing was done in a Turbula shaker-mixer for 24 h in ethanol, using 5 mm-diameter Zr0₂ balls to break up possible agglomerates. After drying, the powder is preferable compacted in a graphite die with 56 mm-diameter and precompacted with a pressure of 30 MPa. Subsequently, the (Nb_x,Zri-x)n+iAIC_n phases can be produced by reactive hot-pressing (W100/150-2200-50 LAX, FCT Systeme, Frankenblick, Germany). Temperature can be controlled by an optical pyrometer measuring on the side surface of the die. The heating rate was set to 20°C/min up to the final sintering temperature of 1600°C and 1700°C for NZ211 and NZ413, respectively. Upon reaching the sintering temperature, the pressure was increased from 5 MPa to 20 MPa with a dwell time of 30 min. After sintering, the discs were grinded to remove the outer reaction layer.

The phase assembly of the grinded discs can be determined by X-ray diffraction (XRD). For the XRD analysis, Cu Ka radiation was used at 40 kV and 40 mA (Seifert 3003 diffractometer). The XRD pattern was measured with a step size of 0.02° and a time of 2 s per step. The lattice parameters of the constituent phases can be calculated by Rietveld refinement using the Topas Academic software. The microstructure was examined by scanning electron microscopy (SEM, XL30-FEG, FEI, Netherlands) equipped with an energy dispersive X-ray spectrometer (EDS, EDAX). The distribution of Nb, Zr, Al, C and O in the sintered materials can be mapped by electron probe microanalysis (EPMA, JXA-8530F, JEOL Ltd., Japan). The beam current and accelerating voltage were fixed at 15 nA and 15 kV. The chemical composition of the phases was quantitatively determined by using EPMA and performing point analyses on 5 different spots on the sample.

The density, p_exp, can be determined based on the Archimedes principle. The length GSL and the width GSW of the grains can be calculated by the Image-pro plus software, according to the linear intercept method. The hardness can be measured using Vickers indentation (FV-700, Future-Tech Corp., Tokyo, Japan), HVio, using a load of 98.1 N with 10 s dwell time. Flexural strength, o4pt, and fracture toughness, Kic, were determined for single phase Nb₄AIC3 and (Nbo.85,Zro.is)4AIC3. Rectangular bars with dimensions 4x3x45 mm³ can be machined out of the hot-pressed discs by electric discharge machining and subsequently grinded to their final dimensions. The flexural strength can be determined by 4-point bending in accordance with ASTM standard C1161-13. The fracture toughness can be measured using the single edge V-notch beam (SEVNB) technique, where the notch had a depth of ±1 mm and a radius of ±20 μιη. The elastic properties at room temperature can be determined using the impulse excitation technique (IET, IMCE, Belgium), according to ASTM standard C1259- 08. The temperature dependence of the Young's modulus and the internal friction can be measured using the IMCE HTVP 1750 IET set-up, equipped with automated impulse excitation and vibration detection devices [G. Roebben, B. Bollen, A. Brebels, J. Van Humbeeck, and 0. Van der Biest, "Impulse excitation apparatus to measure resonant frequencies, elastic moduli, and internal friction at room and high temperature, " Review of Scientific Instruments, vol. 68, pp. 4511-4515, 1997.]. The measurement can be performed in continuous vacuum with a heating rate of 5°C/min up to 1400°C and 1550°C, for Nb₄AIC3 and (Nbo.85,Zro.is)4AC3, respectively. Figure 8 illustrates the Young's modulus of (Nbo.85,Zro.i5)4AIC3 as function of temperature compared to some refractory metals [Modified from www.plansee.com - 11/06/2015].

The XRD patterns of the (Nb_x,Zri-_x)2AIC samples according to embodiments of the present invention are shown in Figure 1. The phase identification lines correspond to the sample with the largest amount of Zr. The peaks of the Nb₂AIC sample show a shift towards a higher 2 theta. This indicates an increase in lattice parameters due to the presence of Zr. The lattice parameters, a and c, optimized during the Rietveld refinement are compared to the values in literature. The result is presented in Table 2. The obtained dimensions are coherent with the reported ones. Looking at the pattern, three phases can clearly be identified: NZ211, NZ413 and (Nbi-_x,Zr_x)C (referred to as NZC). The latter phase is more pronounced in the sample with Nb/Zr equal to 80/20, based on the (111) peak and the (220) peak around 35° and 58° 2 theta, respectively. The NZ413 phase is only present in this sample. This phase can also be identified in the elemental mapping shown in Figure 2. This mapping reveals the lower Zr content in NZ211 compared to NZ413. The quantitative analysis results in a Nb/Zr ratio of 88/12 for the former and 82/18 for the latter phase. Furthermore, the presence of two Al- rich phases, AI2O3 and an Al-Nb intermetallic phase, is observed.

Figure 3 shows the patterns of the (Nb_x,Zri-_x)₄AIC3 sample. Slightly preferred grain orientation is observed when the pattern is compared to the non-oriented powder reference. The increased relative intensity of the (001) peaks indicate an alignment of the c-axis of the grains with the compression direction. This phenomenon is observed in other pressure-assisted sintering techniques, i.e. in spark plasma- sintered samples. Besides, similar to the NZ211 samples, the Zr atoms cause an increase in lattice parameters, which results in a shift towards lower 2 theta angles for the reflections of the same planes. The lattice parameters a and c are plotted as function of the Zr-content in the starting powder in Figure 4. For low Zr-content, a linear trend can be established which corresponds to Vegard's law for solid solutions. The deflection of the linear trend occurs around a Nb/Zr ratio of 82/18. The lattice parameters for this phase and Nb₄AIC3 are added in Table 2. A further increase of the Zr-content initiates the formation of NZC. Figure 5 shows the elemental mapping of the samples with Nb/Zr ratio of 85/15 (a) and 50/50 (b). The former one illustrates the homogeneous distribution of Zr in the grains and the presence of an Al-rich intermetallic phase at the grain boundaries. Figure 5b depicts the NZC phase with a higher Zr-content compared to the NZ413. This indicates that the Zr-atoms are initially dissolved into the MAX phases and at higher Zr-content, above their solubility limit in MAX, they initiate the formation of NZC, where Nb and Zr can form an unlimited range of solid solutions.

The reason for the increase in lattice parameters is most likely found in the difference in covalent radii of Nb and Zr, which are 1.456 A and 1.597 A, respectively. The large Zr-atoms causes an expansion of the volume of the unit cell.

The explanation for the difference in solubility between the NZ211 and NZ413 phases can be found along the same line. Figure 15 illustrates the solid solubility of zirconium (Zr) in a Nb₄AIC3 host lattice, whereas Figure 16 illustrates the solid solubility of zirconium (Zr) in a Nb₂AIC lattice. Etzkorn et al. in "Ta3AIC₂ and Ta₄AIC3 - Single- Crystal Investigations of Two New Ternary Carbides of Tantalum Synthesized by the Molten Metal Technique," Inorganic Chemistry, vol. 46, pp. 1410-1418, 2007, investigated the bonding lengths in the Ta-AI-C system and compared these for the different M_n+iAX_n-stackings with n varying from 1 till∞. The M-A length did not vary much with changing n, while an increase in the M-M bonding length is observed comparing a 211-stacking and a 413-stacking. Due to this increase in bonding length, more space is available to accommodate additional Zr atoms.

Another finding is the presence of NZ413 in the (Nbo.8,Zro.₂)₂AIC sample at 1600°C. It suggests a lower formation temperature of NZ413 caused by the presence of the Zr solute. This statement is supported by the larger grains in the (Nbo.85,Zro.i5)4AIC3 sample compared to the Nb₄AIC3 sample (as presented in Table 3). A similar effect of solute elements is observed for the case of binary carbides. The addition of VC to WC-hard metals resulted in a larger grain size at the same temperature. Furthermore, a beneficial effect of solute binary carbides is found on the densification behavior of ZrC. It is reasonable to assume that this enhanced kinetics are also valid for ternary carbides. The positive effect of larger grains on the mechanical properties is studied in the next section.

Table 2 below illustrates the refined lattice parameters for the end member phases NZ211 and NZ413 and Zr3AIC₂ compared to the values reported in literature.

Table 2

The mechanical properties of Nb₄AIC3 and (N bo.85,Zro.i5)₄AIC3 at room temperature are presented in Table 3. The lower experimental density compared to the theoretical one is caused by secondary phases, such as AI2O3 and Al-rich intermetallics, in both materials. These impurities are depicted in Figure 5a for the case of (N bo.85,Zro.i5)₄AIC3. However, these hard and brittle particles do not cause a dramatic change in hardness or toughness compared to the properties reported in literature.

Figure 6 shows the tip of a Vicker's indent on both materials and the deformed nano-laminated grains, characteristic for MAX phases. Dissimilar to Wan et al. in ""A New Method to Improve the High-Temperature Mechanical Properties of Ti3SiC2 by Substituting Ti with Zr, Hf, or Nb," Journal of the American Ceramic Society, vol. 93, pp. 1749-1753, 2010.", the Zr-substitution did not cause a significant change in hardness or stiffness. Regarding the strength and fracture toughness, a clear difference between Nb₄AIC3 and (Nbo.85,Zro.is)4AIC3 is observed. The lower strength and higher toughness of the latter can be related with the larger grains. The Hall-Petch relationship is considered valid for MAX phases and results in a higher strength for smaller grains. The opposite is true for the toughness; a larger grain size is usually associated with a higher fracture toughness. This observation was made by Hu et al. in" "In Situ Reaction Synthesis and Mechanical Properties of V2AIC," Journal of the American Ceramic Society, vol. 91, pp. 4029-4035, 2008." for V₂AIC and was confirmed for other MAX phases.

To evaluate the high-temperature stability, the normalized Young's modulus and internal friction as function of temperature were recorded. The result is shown in Figure 7. The bump observed around 700°C in the damping curve of Nb₄AIC3 is caused by the softening of the Al-rich intermetallic phase. This is most probably the reason for the slightly higher damping compared to the value reported in literature. Nevertheless, a shift toward higher temperatures is found for the damping intensity of (Nbo.85,Zro.i5)4AIC3 compared to Nb₄AIC3. The former retains more than 80% of its stiffness at 1550°C. These results are in line with the findings on the improved high- temperature stability of (Ti,Zr)3(Si,AI)C₂. Table 3

Moreover, the stability of the normalized Young's modulus and the damping are comparable to the properties of Zr(AI,Si)C₂-_x based ceramics. The reason for this improved behavior is unclear at the moment; still, (Nbo.85,Zro.is)4AIC3 has the highest thermal stability amongst all reported M_n+iAX_n phases.

Based on these findings, it is clear that the addition of a fourth solute element can improve the material performance. Further investigation of the mechanical properties of the reported MAX phase solid solutions are definitely of interest.

Embodiments of the present invention provide, (Nbi-_x,Zr_x)n+iAICn-y ceramics, where x = 0.01 to 1, n = 2 to 3 and y = -0.5 to 0.5, which can be prepared starting from Nb and Zr hydrides. Zr could partially substitute Nb in Nb-AI-C based MAX phases. It appears that Nb₄AIC3 can accommodate more Zr atoms (± 18%). The properties of (Nbo.85,Zro.i5)4AIC3 were compared to the ones of pure Nb₄AIC3. No significant change in hardness or stiffness was observed. The flexural strength decreased to 507±92 MPa and the fracture toughness increased to 10.1±0.3 MPa.m^1/2. These two trends can be correlated to the larger grain size of the solid solution compared to Nb₄AIC3, as indicated by the length (GSL) and width (GS_w)of the grains in Table 3. Finally, the temperature dependence of the elastic properties was evaluated by IET. More than 80% of the stiffness is retained at 1550°C concomitant with a low elastic damping. Based on this result, this solid solution has a comparable temperature stability to the more brittle Zr-AI-C phases. Example of a method to texture dense M_n+iAX_n ceramics by spark plasma deformation according to embodiments of the present invention

A schematic illustration of a method to texture dense ceramics according to embodiments of the present invention is provided in Figure 11. Primarily (99%) dense discs of 40 mm in diameter were prepared by spark plasma sintering of commercially- available Maxthal 312^® and Maxthal 211^® raw powders (Sandvik, Sweden). The powders were first cold pre-compacted and then spark plasma sintered at a heating rate of 100°C/min with a dwell time of 5 min under 30 MPa at the sintering temperature, Td, of 1350°C and 1300°C for Maxthal 312^® and Maxthal 211^®, respectively. The dense discs (with a 40 mm diameter) were ground plane-parallel to a thickness of 8 mm. The discs were deformed in a second SPS run by placing the 40-mm diameter discs in a 56-mm graphite die, pre-loading them at 70 MPa and heating to the sintering temperature. Non-deformed, 56-mm diameter discs of each material were prepared according to the first SPS cycle described above for comparison. Spark plasma sintering (FCT-Systeme, HP D 25, Frankenblick, Germany) was performed in vacuum (-100 Pa) and temperature control was achieved by focusing an optical pyrometer at ~2 mm to 3 mm above the middle of the disc surface.

XRD patterns of the raw powders, the textured side surface (TSS) and the textured top surface (TTS) of the deformed ceramics were evaluated. The analysis of the phase purity of the commercial powders detected the presence of TiC in Maxthal 312^® as minor (9 wt%) secondary phase. The TiC amount in the bulk ceramics remained constant during the different processing steps. The Maxthal 211^® powder was found to be a mixture of three phases, two of which were MAX phases, i.e., 62 wt% of Ti₂AIC, 30 wt% Ti₃AIC₂ and 8 wt% of Ti₂AI₅. This ratio of (Ti₂AIC)/(Ti₃AIC₂) = 2 was maintained throughout the different processing steps.

Looking at the relative intensities of the different plane reflections, one may observe a shift towards a preferred crystallographic orientation from the starting powder to the spark plasma sintered disc and then to the deformed material, the shift being most explicit for Ti3SiC₂. The relative intensities of the peaks in the XRD patterns of the raw powders correspond well to those of the reference JCPDS files. In the non- textured Ti₃SiC₂ (JCPDS 00-059-0189), the (104) peak has the highest intensity. However, in the deformed Ti3SiC₂, the strongest reflection (008) corresponds to the (0 0 Indirection, which is strongly aligned parallel to the compression axis, resulting in a Lotgering factor of 0.52.

In the case of Maxthal 211^®, the change in relative peak intensity was less apparent, due to the close proximity of the (008) and (104) T13AIC2 peaks with the (1 0 3) and (0 0 6) Ti₂AIC peaks. Nevertheless, the increase in the intensity of the (0 0 2) peak of both T13AIC2 and Ti₂AIC phases are associated with the alignment of the MAX phase grains parallel to the (001)- direction. The MAX phase grains clearly align with their c-axis, i.e., the (OOl)-direction, parallel to the compression axis. The calculated Lotgering factors are 0.51 for T13AIC2 and 0.49 for Ti₂AIC, respectively. With the exception of some isolated pores, the materials remain dense after deformation. This fact, combined with the observed grain alignment, indicates that the deformation originates mainly from grain reorientation in the bulk ceramics. SEM images of deformed grains in the Maxthal 211^® ceramic were also provided. For example, the grain in the centre is distorted, while other grains tend to delaminate and breakup in smaller domains. All these observations are in good agreement with the damage mechanisms proposed in literature., i.e., grain bending, rotating and breakup. More specifically, the latter mechanism was rather frequently observed and could be attributed to the high shear stresses generated during compression. These stresses result in local delamination and fragmentation of the elongated grains, which are then freer to rotate. A similar phenomenon was observed during the formation of MAX phase coatings using high-velocity particles: grain alignment was observed with the c- axis perpendicular to the substrate during high velocity oxy-fuel spraying of Maxthal 211^® and aerosol deposition of Maxthal 312^®.

The deformation mechanisms accountable for high- temperature plasticity are a topic of discussion. As described above, a transition in elastic response is observed with increasing the temperature. The temperature at which this transition takes place is material-dependent and differs for each MAX phase. Evidence of this transition may be found in the evolution of the stiffness as a function of temperature for both dense, deformed ceramics, together with the recorded upper piston displacement during SPS deformation. The upper piston displacement curves were corrected for the linear thermal expansion of both SPS setup and dense MAX materials. The Young's modulus follows a linear decrease up to approximately 1000°C. The decrease rate of the normalized stiffness (E/ERT) of the two ceramics was comparable and in good agreement with the results obtained by RUS. Around 1200°C, a steeper decline of the Maxthal 211^® stiffness was observed and concomitantly the disc started to deform. This phenomenon was also observed for Maxthal 312^® around 1250°C, which was close to the softening temperature (1300 °C) reported in literature. This suggests that the softening of MAX phases is accompanied by an increase in deformability. Furthermore, the softening temperature is an indication of the maximum service temperature up to which these two MAX ceramics could be used for structural (load-bearing) applications.

The fracture toughness was used to assess the effect of deformation on the material properties. The measured fracture toughness of non-deformed Maxthal 211^® and Maxthal 312^® was 6.0 ± 0.3 MPa m^1/2 and 5.1 ± 0.1 MPa m^1/2. After deformation, the fracture toughness was measured both parallel and perpendicular to the c-axis and an improvement was observed in both directions. The highest fracture toughness was recorded when the notch was parallel to the c-axis, i.e., 7.9 ± 0.1 MPa m^1/2 and 6.0 ± 0.2 MPa m^1/2 for Maxthal 211^® and Maxthal 312^®, respectively. The fracture toughness values measured across the disc cross-sections were 6.5 ± 0.1 MPa m^1/2 and 5.3 ± 0.1 MPa m^1/2 for Maxthal 211^® and Maxthal 312^®, respectively. It can, therefore, be stated that deformation does not affect the exceptional fracture toughness of these ceramics adversely. The grain alignment even enhances the resistance to crack propagation. This result might also be relevant for different deformation-induced textured MAX phases, e.g., in coatings.

In conclusion, embodiments of the present invention provide on a new method to texture MAX phase ceramics by means of high-temperature uniaxial deformation of dense bulk materials in a spark plasma sintering setup. The grains were aligned with their c-axis parallel to the deformation (compression) direction. Moreover, the Lotgering factor was calculated to be (a) 0.52 for Ti3SiC₂ in deformed Maxthal 312^® and (b) 0.51 and 0.49 for Ti₃AIC₂ and Ti₂AIC, respectively, in deformed Maxthal 211^®. The ternary carbides were deformable above -1200 °C for Maxthal 211^® and ~1250°C for Maxthal^® 312, as deduced from the evolution of the Young's modulus with temperature. Texturing of the two MAX phase ceramics improved the fracture toughness in both directions, with the largest increase measured parallel to the c-axis. For example, the fracture tough- ness of Ti₂AIC increased from 6.0 MPa m^1/2 in the as- sintered ceramic to 7.9 MPa m^1/2 and 6.5 MPa m^1/2 parallel and perpendicular to the loading direction, respectively, in the deformed ceramic.

Claims

A method for producing a MAX phase material with general formulation Mn₊iAXn-y, whereby M is a transition metal, A is an A-group element and X is carbon (C) and/or nitrogen (N), where the M, A and/or X position can be occupied by more than 1 element with n within the range 1 to 3, y within the range -0.5 to 0.5, said method comprising:

- producing the MAX phase material by using a pressure assisted sintering method. The method according to claim 1, wherein the mixture is a starting powder, wherein the starting powder is compacted by pressure assisted densification. The method of any of previous claims, further comprising adding another transition metal to the mixture.

The method of any of previous claims, wherein the precursor of M, the precursor of A and the precursor of X of the mixture are mixed to obtain a near stoichiometric ratio with a slight excess of A and sub-stoichiometric amount of X.

The method according to any of claims 2 to 4, wherein the starting powder is mixed in an organic fluid and said starting powder in in the organic fluid is dried before compacting the powder.

The method according to any of claims 4 or 5, wherein mixing said powders comprises mixing with balls.

The method according to any of claims 2 to 6, wherein compacting the powder mixture comprises pouring the powder mixture into a die and cold-compacting said mixture resulting in a powder compact.

The method according to any of claims 2 to 7, wherein compacting the starting powder comprises providing a pressure in the range of 20 MPa to 40 MPa, preferably within the range 25 MPa to 35 MPa on the powder.

9. - The method according to any of claims 5 to 8, wherein the organic fluid is ethanol and the powder in said organic fluid is mixed for at least 10 h.

10. - The method according to any of previous claims, wherein the pressure assisted sintering technique comprises reactive hot pressing with a heating rate of 5°C/min to 100°C/min up to the final sintering temperature of 1700°C.

11. - The method according to claim 10, wherein upon reaching the final sintering temperature, a pressure is provided increasing from 5 MPa to 30 MPa, preferably from 5 MPa to 20 MPa.

12. - The method according to claim 11, wherein a dwell time of 30 minutes is applied.

13.- The method according to any of previous claims, wherein the at least one hydride is a Ti, Nb, Zr, Hf and/or Ta based hydride.

14. - The method according to claim 13, wherein at least one TiH₂, NbHo.89, ZrH₂, HfH₂ or TaHx, hydride and Al and C are used as a starting powder mixture.

15. - A MAX phase material obtained by a method according to any of previous claims, wherein the MAX phase material is in the form of M_n+iAX_n-y with n within the range

1 to 3, y within the range -0.5 to 0.5:

- with M selected from the group consisting of Zr and Hf or

- with M occupied by at least two elements in the form (M_x, ΜΊ-_Χ) with M and M' selected form the group consisting of Ti, Nb, Ta, Zr and Hf, with M different from M', and wherein M' partially substitutes M in a M-A-X based system as a starting mixture; and with x within the range 0.01 to 1

16. - The MAX phase material according to claim 15, wherein M comprises Nb.

17. - The MAX phase material according to claim 15 or 16, wherein the MAX material is a ternary carbide of the form (Ti_v,Nbw,Zrx,Hfy,Ta_z)n+i(AI,A)(C,N)n-y with v, w, x, y, z > 0, and v+w+x+y+z being equal to 1, n being in the range 1 to 3, y being in the range -0.5 to 0.5 and A being Si, Sn, Pb or Bi.

18. - A MAX phase material with general formulation M_n+iAX_n-y, whereby M is a transition metal, A an A-group element and X being carbon C and/or nitrogen N with n being within the range 1 to 3, y being within the range -0.5 to 0.5:,

- with M occupied by at least Zr and a 312 stoichiometry or

- with M occupied by at least Hf and a 211 or 312 stoichiometry or

- where M is occupied by at least two elements in the form of (M_x, ΜΊ-_Χ) with M and M' selected from the group consisting of Ti, Nb, Ta, Zr and Hf with M different from M', and wherein M' partially substitutes M in a M-A-X based system as a starting mixture and with x= 0.01-1, and where the A and X position can be occupied by one or more elements.

19. -The MAX phase material according to claim 18, with A selected from the group consisting of Si, Sn, Pb and Bi or A occupied by at least two elements in the form (A, Al) selected from the group consisting of Si, Sn, Pb and Bi and X selected from the group consisting of B, C, and N or X occupied by at least two elements in the form (C, N).

20. - The MAX phase material according to claim 18 or 19, wherein M comprises Nb.

21. - The MAX phase material of claim 18 or 19, wherein the MAX material is a ternary carbide of the form Zr3AIC₂.

22. - The MAX phase material of claim 18, wherein the MAX material is a ternary carbide of the form (Ti_v,Nbw,Zrx,Hfy,Ta_z)n+i(AI,A)(C,N)n with v, w, x, y, z > 0, and v+w+x+y+z being equal to 1, n being within the range 1 to 3 and A an Si, Sn, Pb or Bi.

23. - The MAX phase material of claims 18 to 20 or 22, wherein the MAX phase material, where M is occupied by at least two elements in the form of (M_x, ΜΊ-_Χ), has a 211, 312 or 413 stoichiometry.

24. - The MAX phase material of claim 23, wherein the MAX phase material is of the form (Nb_x, Zri-_x)n+iAX_n-y with a 312 or 413 stoichiometry.

25. - The MAX phase material according to any of claims 18 to 24, whereby the MAX phase material is a powder.

26. - The MAX phase material according to any of claims 18 to 24, whereby the MAX phase material is a sintered material, a sintered bulk material and/or a spark plasma sintered material.

27. - The MAX phase material according to any of claims 18 to 24, whereby the MAX phase material is a solid solution.

28. - The MAX phase material according to any of claims 18 to 24, whereby the MAX phase material is a pre-reacted powder compact.

29. - The MAX phase material according to any of claims 18 to 24, whereby the MAX phase material is a coating material.

30.- The MAX phase material according to any of claims 18 to 24, whereby the MAX phase material is a sintered or spark plasma sintered and hot pressed material.

31. - The MAX phase material according to any of claims 18 to 24, whereby the MAX phase material is a sintered or spark plasma sintered and cold-pressed material.

32. - The MAX phase material according to any of claims 18 to 24, whereby the MAX phase material is a powder sintered in the in the 1350°C to 1900°C temperature range under a pressure in the range of 5 MPa to 100 MPa, preferably within the range 10 MPa to 50 MPa, more preferably within the range 28 MPa to 32 MPa or about 30 MPa.

33. - The MAX phase material according to any of claims 18 to 24, whereby the MAX phase material is a powder that is sintered in the 1450°C to 1725°C temperature range under a pressure in the range of 20 MPa to 40 MPa, preferably within the range 25 MPa to 35 MPa, more preferably in the range of 28 MPa to 32 MPa or about 30 MPa and a dwell time within the range 230 s to 2500 s, preferably within the range 500 s to 2000 s, more preferably within the range 1700 s to 1900 s.

34.- The MAX phase material according to any of claims 18 to 33, wherein the MAX phase material is nano-layered.

35.- The MAX phase material according to any of claims 18 to 34, characterized in that the MAX phase material comprises a partial alignment of the c-axis of the grains with the compression direction.

36. - The MAX phase material according to any of claims 18 to 35, when the MAX phase material is a Nb- based ternary carbide , preferably of the form (Nb_x,Zri-_x)₄AIC3 (413 N bZ), the MAX phase material exhibits at room temperature one of the following mechanical properties in any one of the following ranges: p_exp within the range 5 g/cm³ to 9 g/cm³, p_th within the range 5 g/cm³ to 8 g/cm³, GSL within the range 15 μιη to 28 μιη, GSw within the range 3 μιη to 6 μιη, HVio within the GPa range, E within the range 325 GPa to 365 GPa, o_4pt within the range 450 M Pa to 550 M Pa, Kic within the range 5 M Pa.m^1/2 to 15 M Pa.m^1/2 or a combination thereof.

37. - The MAX phase material according to any of claims 18 to 35, wherein the MAX phase material is a Nb-based ternary carbide, preferably of the form 413 NbZ, characterized in that said ternary carbide exhibits at room tem perature one of the following mechanical properties in any one of the following ranges: p_exp within the range 6 g/cm³ to 7 g/cm³, p_th within the range 6 g/cm³ to 7 g/cm³, GSL within the range 21 μιη to 25 μιη, GSw within the range 5.5 μιη to 6.8 μιη, HVio within the GPa range, E within the range 340 GPa to 350 GPa, o_4pt within the range 500 M Pa to 515 M Pa, Kic being 10.1 ± 0.3 M Pa.m^1/2or a combination thereof.

38. - The MAX phase material according to any of claims 18 to 35, wherein the MAX phase material is a Nb-based ternary carbide, preferably of the form 413 NbZ, characterized in that said ternary carbide exhibits at room tem perature one of the following mechanical properties in any one of the following ranges: p_ex within the range 6.5 g/cm³ to 6.9 g/cm³, p_th within the range 6.8 g/cm³ to 6.9 g/cm³, GSL within the range ΙΟμιη to 30μιη, GSw within the range 6 μιη to 6.6 μιη, HVio within the GPa range, E within the range 340 GPa to 355 GPa, o_4pt within the range 500 M Pa to 600 M Pa (± 92M Pa), Kic being within the range 9 M Pa.m^1/2 to 11 M Pa.m^1/2 or a combination thereof.

39. - The MAX phase material according to any of claims 18 to 35, wherein the MAX phase material is a Nb-based ternary carbide, preferably of the form 413 NbZ, characterized in that said ternary carbide has lattice parameters in any one of the following ranges : a_ex within the range 2.5 A to 3.6 A and/or c_ex within the range 20 A to 30 A or a combination thereof

40.- The MAX phase material according to any of claims 18 to 35, wherein the MAX phase material is a Nb-based ternary carbide, preferably of the form 413 NbZ, characterized in that said ternary carbide retains more than 80% of its stiffness at 1550°C.

41.- The MAX phase material according to any of claims 18 to 35, wherein the MAX phase material is the ternary carbide (Nbo.85,Zro.is)4AIC3, characterized in that said ternary carbide has a higher thermal stability as compared to Nb₄AIC3.

42. - The MAX phase material according to any of claims 18 to 35, wherein the MAX phase material is a Nb-based ternary carbide, preferably of the form 413 NbZ, characterized in that said ternary carbide has a higher thermal stability among all reported M_n+iAX_n phases than reported M_n+iAX_n phases.

43. - The MAX phase material according to any of claims 18 to 35, wherein the MAX phase material is a Nb-based ternary carbide, preferably of the form 413 NbZ, characterized in that said ternary carbide has a fracture toughness in the range of 8 MPa.m^1/2 to 12 MPa.m^1/2, preferably in the range of 9 MPa.m^1/2 to 11 MPa.m^1/2.

44. - The MAX phase material according to any of claims 18 to 35, wherein the MAX phase material is a Nb-based ternary carbide, preferably of the form 413 NbZ, characterized in that said ternary carbide has a fracture toughness increased to 10.1±0.3 MPa.m^1/2.

45.- The MAX phase material according to any of claims 18 to 35, wherein the MAX phase material is a Nb-based ternary carbide, preferably of the form 413 NbZ, characterized in that said ternary carbide has a flexural strength in the range of 400 to 600 MPa.

46. - A method for texturing MAX phase materials, said method comprising deforming a provided MAX phase material wherein the provided MAX phase material is compacted.

47. -The method of claim 46, wherein the deformation is performed uniaxial at temperatures above 1200°C by hot pressing or spark plasma deformation.

48. - The method of any of claims 46 or 47, wherein the provided MAX phase material is a MAX phase material according to any of claims 18 to 45.

49. -The method of any of claims 1 to 14 further comprising texturing the obtained MAX phase material according to the method of claims 46 to 48.

50. - A MAX phase material obtained using a method according to any of claims 46 to

49, wherein the MAX phase material comprises grains that show stronger alignment.

51. - The MAX phase material of claim 50, wherein the MAX phase material comprises an increased fracture toughness of 11.1 MPa.m^1/2 in comparison to the non- deformed material.

52. - Use of a MAX phase material according to any of claims 18 to 45, with M or M' comprising Zr, as fuel cladding material of light water reactors.