The present invention relates to a cast and wrought (C&W) nickel-based superalloy composition for use in high stress high temperature applications beyond 650°C, for example rotating components in a gas turbine engine and other turbomachinery. Increases in alloy performance - in terms of maximum operating temperature and maximum service life - can have a significant impact on the efficiency of the engine as well as the cost effectiveness of operating the engine.

Examples of typical compositions of C&W nickel-based superalloys which are used for high temperature and high stress applications are listed in Table 1. In development of higher strength alloys there has been a tendency to reduce workability of the alloy and increase the use of higher cost elements.

Table 1: Nominal composition in wt. % of cast & wrought gamma/gamma-prime nickel-based superalloys.

It is the aim of this invention to achieve an improved trade off between mechanical strength and hot workability. In certain cases an improved mechanical strength compared to alloys such as Waspaloy or Haynes 282 is desired, the targeted strength being that of high strength alloys such as Rene65, AD730, 720Li and TMW-4. This achieved by increasing g’ volume fraction to levels higher than Waspaloy or Haynes 282, described in relation to Figure 2. In certain cases a g’ solvus lower than current high strength alloys, Rene65, AD730, 720Li and TMW-4 is desired to maintain better hot workability. In its optimal form the alloy of the invention achieves both improved mechnical strength compared to Waspaloy and Haynes 282 and lower solvus than Rene65, AD730, 720Li and TMW-4 simultaneously.

The present invention provides a nickel-based alloy composition consisting, in weight percent, of: from 17.0% to 21.3% chromium, 7.1% or less cobalt, from 0.9% to 6.3% molybdenum, 4.9% or less tungsten, from 1.8% to 3.2% aluminium, from 1.8% to 4.0% titanium, 3.05% or less tantalum, 3.0% or less niobium, 0.1 % or less carbon, from 0.001% to 0.1% boron, from 0.001% to 0.5%. zirconium, 0.02% or less magnesium, 0.5% or less silicon, 0.1% or less yttrium, 0.1% or less lanthanum, 0.1% or less cerium, 0.003% or less sulphur, 0.25% or less manganese, 0.5% or less copper, 0.5% or less hafnium, 0.5% vanadium or less, 10.0% or less iron, the balance being nickel and incidental impurities Such an alloy has an improved trade-off between strength and gamma prime solvus temperature.

In an embodiment the following equation is satisfied in which W _Al , W _Ti , W _Nb and W _Ta are the weight percent of aluminium, titanium, niobium and tantalum in the alloy respectively

(0.6W _Ti + 0.31W _Nb + 0.15W _Ta )/W _Al £ 1.1

preferably, (0.6W _Ti + 0.31W _Nb + 0.15W _Ta )/W _Al £ 1.0

Such an alloy has reduced possibility of eta and delta phase formation particularly if produced by the cast and wrought process.

In an embodiment the following equation is satisfied in which W _Al , W _Ti , W _Nb and W _Ta are the weight percent of aluminium, titanium, niobium and tantalum in the alloy respectively

(0.6W _Ti + 0.31W _Nb + 0.15W _Ta )/W _Al ³ 0.64 preferably, (0.6W _Ti + 0.31W _Nb + 0.15W _Ta )/W _Al ³ 0.72

more preferably (0.6W _Ti + 0.31W _Nb + 0.15W _Ta )/W _Al ³ 0.76 even more preferably (0.6W _Ti + 0.31W _Nb + 0.15W _Ta )/W _Al ³ 0.93

Such and alloy has an improved mechanical strength.

In an embodiment the alloy comprises a volume fraction of gamma prime phase of 48% or less at 760°C, preferably of 44% at 760°C or less, more preferably of 40% or less at 760°C and most preferably of 36% or less at 760°C.This is effective to limit the gamma prime solvus and so leads to better hot workability.

In an embodiment the following equation is satisfied in which W _Al , W _Ti , W _Ta and W _Nb are the weight percent of aluminium, titanium, tantalum and niobium in the alloy respectively 3.4 £ 0.6W _Ti + 0.31W _Nb + 0.15W _Ta + 0.79W _Al

preferably

3.6 £ 0.6W _Ti + 0.31W _Nb + 0.15W _Ta + 0.79W _Al

Such an alloy has an acceptable level of gamma prime and so has high strength.

In an embodiment the following equation is satisfied in which W _AI , W _T , , W _T3 and W _Nb are the weight percent of aluminium, titanium, tantalum and niobium in the alloy respectively

4.6 ³ 0.6W _Ti + 0.31W _Nb + 0.15W _Ta + 0.79W _Al preferably

4.3 ³ 0.6W _Ti + 0.31W _Nb + 0.15W _Ta + 0.79W _Al more preferably

4.0 ³ 0.6W _Ti + 0.31W _Nb + 0.15W _Ta + 0.79W _Al Such an alloy has improved workability.

In an embodiment the alloy composition consists of, in weight percent, of 18.0% or more chromium. Such an alloy has improved oxidation and corrosion resistance.

In an embodiment the alloy composition consists of, in weight percent, of 19.5% or less chromium. Such an alloy is less prone to the precipitation of deleterious phases.

In an embodiment the alloy composition consists of, in weight percent, of 2.25wt% or less tantalum, preferably of 1.5wt% or less tantalum, most preferably of 0.65wt% or less tantalum. Such an alloy can have improved mechanical strength and/or allows a lower level of niobium for the same strength.

In an embodiment the alloy composition consists of, in weight percent, of 2.0% or more molybdenum, preferably 3.0% or more molybdenum. Such an alloy has an improved combination of increased creep resistance and lower density. In an embodiment the alloy composition consists of, in weight percent, of 1.9% or more titanium, preferably 2.0% or more titanium, more preferably 2.3% or more titanium, yet more preferably 2.4% or more titanium, most preferably 2.5% or more titanium. Such an alloy has lower gamma prime solvus temperature.

In an embodiment the alloy composition consists of, in weight percent, of 2.7 wt.% or less tungsten, preferably 0.7wt% or less tungsten. Such an alloy has lower density.

In an embodiment the alloy composition consists of, in weight percent, of 2.0% or less niobium. Such an alloy has better oxidation resistance.

In an embodiment the alloy composition consists of, in weight percent, of 3.5% or less titanium, preferably 3.0% or less titanium. Such an alloy has an improved oxidation resistance.

In an embodiment the alloy composition consists of, in weight percent, of 1.9% or more aluminium. Such an alloy has improved gamma prime phase stability.

In an embodiment the alloy composition consists of, in weight percent, of 2.75% or less aluminium, preferably of 2.6% or less aluminium, more preferably of 2.3% or less aluminium. Such an alloy has a better combination of strength and hot workability.

In an embodiment the alloy composition consists of, in weight percent, of 5.3% or less cobalt, preferably 3.5% or less cobalt, more preferably 1.6% or less cobalt. Such an alloy has reduced cost.

In an embodiment the following equation is satisfied in which W _Ta and W _Co are the weight percent of tantalum and cobalt in the alloy respectively: W _Ta + 0.43 W _Co £3.05, preferably WTa + 0.43 W _Co £ 2.25, more preferably W _Ta + 0.43W _Co £1.5, most preferably W _Ta + 0.43 W _Co £0.65. Such an alloy has reduced cost.

In an embodiment the following equation is satisfied in which W _Ti , W _Nb , W _Ta and W _AI are the weight percent of titanium, niobium, tantalum and aluminium in the alloy respectively

0.6W _Ti + 0.44(W _Nb + 0.66 W _Ta ) + 0.2W _Al ³ 2.8 Such an alloy has a good mechanical strength.

In an embodiment the alloy composition consists of, in weight percent, of 0.3% or more niobium, preferably 0.4% or more niobium, more preferably 0.7% or more niobium, even more preferably 0.8% or more niobium, most preferably 1.2 % or more niobium. Such an alloy has improved strength.

In an embodiment the following equation is satisfied in which W _MO and W _W are the weight percent of molybdenum and tungsten in the alloy respectively

W _MO + 0.5 W _w ³ 3.4,

preferably W _MO + 0.5 W _w ³ 3.6

Such an alloy has good high temperature properties including tensile strength and creep resistance because of a stronger matrix due to solid solution hardening.

In an embodiment the following equation is satisfied in which W _MO and Ww are the weight percent of molybdenum and tungsten in the alloy respectively

W _MO + 0.5 W _w £ 6.3

preferably W _MO + 0.5 W _w £ 5.6

more preferably W _Mo + 0.5W _w £ 5.1

most preferably W _MO + 0.5 W _w £ 4.3

Such an alloy has good high temperature creep strength due to a strong matrix phase.

In an embodiment the alloy composition consists of a volume fraction of gamma prime phase of at least 29% at 760°C, preferably at least 31% at 760°C and most preferably at least 40% at 760°C. Such an alloy has good tensile strength characteristics.

The term“consisting of’ is used herein to indicate that 100% of the composition is being referred to and the presence of additional components is excluded so that percentages add up to 100%. Unless otherwise stated, percents are expressed in weight percent.

The invention will be more fully described, by way of example only, with reference to the accompanying drawings in which: Figure 1 shows the calculated correlation between g’ volume fraction at 760°C and g’ solvus temperature for alloys within the alloy design space listed in Table 2, dotted lines show the preferred limits for g’ solvus temperature;

Figure 2 shows the calculated correlation between the volume fraction of g’ phase at 760°C and the strength of the alloy, in terms of strength merit index. The calculated strength of high strength alloys such as Rene65, AD730, 720Li and TMW-4 is identified with dashed lines;

Figure 3 describes the relationship between alloy strength and the ratio of elements according to the formula ((0.6W _Ti +0.31W _Nb +0.15W _Ta )W _Al ). Dashed lines delineated on the graph identify the limits for strength at different preffered ranges of g’ volume fraction. The calculated strength of high strength alloys such as Rene65, AD730, 720Li and TMW-4 is identified with dashed lines;

Figure 4 is a contour plot showing the effect of elements aluminium and niobium plus tantalum plus titanium (according to the relationship 0.6W _Ti +0.31W _Nb +0.15W _Ta on volume fraction of g' at 760°C. Superimposed are the preferred limits for the ratio of the elements titanium, tantalum, niobium and aluminium according to the relationship (0.6W _Ti +0.31W _Nb +0.15W _Ta /W _Al . The shaded area defines the target compositional range for the application;

Figure 5 is a contour plot showing the effect of elements aluminium and titanium on yield strength (in terms of strength merit index), when the niobium content is fixed at 0.0 wt.% and tantalum content is fixed at 0.0 wt.%. Superimposed are the preferred limits for sum of the elements titanium, tantalum, niobium and aluminium according to the relationship

0.79W _Ai +0.6W _Ti +0.31W _Nb +0.15W _Ta which describes different levels of g’ volume fraction and also preferred limits for (0.6W _Ti +0.31W _Nb +0.15W _Ta /W _{Al .} . The shaded area defines the target compositional range for the application;

Figure 6 is a contour plot showing the effect of elements aluminium and titanium on yield strength (in terms of strength merit index), when the niobium content is fixed at 1.0 wt.% and tantalum content is fixed at 0.0 wt.%. Superimposed are the preferred limits for sum of the elements titanium, tantalum, niobium and aluminium according to the relationship

0.79W _Al +0.6W _Ti +0.31W _Nb +0.15W _Ta which describes different levels of g’ volume fraction and also preferred limits for (0.6W _Ti +0.31W _Nb +0.15W _Ta /W _{Al .} . The shaded area defines the target compositional range for the application; Figure 7 is a contour plot showing the effect of elements aluminium and titanium on yield strength (in terms of strength merit index), when the niobium content is fixed at 2.0 wt.% and tantalum content is fixed at 0.0 wt.%. Superimposed are the preferred limits for sum of the elements titanium, tantalum, niobium and aluminium according to the relationship

0.79W _Al +0.6W _Ti +0.31W _Nb +0.15W _Ta which describes different levels of g’ volume fraction and also preferred limits for (0.6W _Ti +0.31W _Nb +0.15W _Ta /W _{Al .} . The shaded area defines the target compositional range for the application;

Figure 8 is a contour plot showing the effect of elements aluminium and titanium on yield strength (in terms of strength merit index), when the niobium content is fixed at 3.0 wt.% and tantalum content is fixed at 0.0 wt.%. Superimposed are the preferred limits for sum of the elements titanium, tantalum, niobium and aluminium according to the relationship

0.79W _Al +0.6W _Ti +0.31W _Nb +0.15W _Ta which describes different levels of g’ volume fraction and also preferred limits for (0.6W _Ti +0.31W _Nb +0.15W _Ta /W _{Al .} . The shaded area defines the target compositional range for the application;

Figure 9 is a contour plot showing the effect of elements molybdenum and tungsten (according to the relationship W _Mo +0.5WW) on the solid solution index;

Figure 10 describes the effect of g’ volume fraction and the sum of the elements molybdenum and tungsten (according to the relationship W _Mo +0.5WW) on the maximum achievable chromium content for a target stability number of 0.90;

Figure 11 describes the effect of g’ volume fraction and the sum of the elements molybdenum and tungsten (according to the relationship W _Mo +0.5WW) on the maximum achievable chromium content for a target stability number of 0.89;

Figure 12 is a contour plot showing the effect of elements cobalt and tantalum on alloy cost; and

Figure 13 is a contour plot showing the effect of elements tungsten and tantalum on alloy density.

Traditionally, nickel-based superalloys have been designed through empiricism. Thus their chemical compositions have been isolated using time consuming and expensive experimental development, involving small-scale processing of limited quantities of material and subsequent characterisation of their behaviour. The alloy composition adopted is then the one found to display the best, or most desirable, combination of properties. The large number of possible alloying elements indicates that these alloys are not entirely optimised and that improved alloys are likely to exist.

In superalloys, generally additions of chromium (Cr) and aluminium (Al) are added to impart resistance to oxidation/corrosion, cobalt (Co) is added to improve resistance to sulphidisation. For creep resistance, molybdenum (Mo), tungsten (W) and cobalt (Co) are introduced, because these retard the thermally-activated processes - such as, dislocation climb - which determine the rate of creep deformation. To promote static and cyclic strength, aluminium (Al), tantalum (Ta), niobium (Nb) and titanium (Ti) are introduced as these promote the formation of the precipitate hardening phase gamma-prime (g'). This precipitate phase is coherent with the face-centered cubic (FCC) matrix phase which is referred to as gamma (g).

A modelling-based approach used for the isolation of new grades of nickel-based superalloys is described here, termed the“Alloys-By-Design” (ABD) method. This approach utilises a framework of computational materials models to estimate design relevant properties across a very broad compositional space. In principle, this alloy design tool allows the so called inverse problem to be solved; identifying optimum alloy compositions that best satisfy a specified set of design constraints.

The first step in the design process is the definition of an elemental list along with the associated upper and lower compositional limits. The compositional limits for each of the elemental additions considered in this invention - referred to as the“alloy design space” - are detailed in Table 2.

Table 2: Alloy design space studied.

The second step relies upon thermodynamic calculations used to calculate the phase diagram and thermodynamic properties for a specific alloy composition. Often this is referred to as the CALPHAD method (CALculate PHAse Diagram). These calculations are conducted at the typical service temperature for the new alloy (900°C), providing information about the phase equilibrium (micro structure). A third stage involves isolating alloy compositions which have the desired microstructural architecture. In the case of nickel based superalloys which require superior resistance to creep deformation, the creep rupture life generally improves as the volume fraction of the precipitate hardening phase g' is increased, the most beneficial range for volume fraction of g' lies between 60%-70% at 900°C (however often due to other design restraints volume fraction may be limited to lower values than this and so alloys with a g' volume fraction of 50% to 60% are included). At values above 70% volume fraction of g' a drop in creep resistance is observed.

It is also necessary that the g/g' lattice misfit should conform to a small value, either positive or negative, since coherency is otherwise lost; thus limits are placed on its magnitude. The lattice misfit d is defined as the mismatch between g and g' phases, and is determined according to

where a _g and a _g ' are the lattice parameters of the g and g' phases.

Thus the model isolates all compositions in the design space which are calculated to result in a desired volume fraction of g', which have a lattice misfit g' of less than a predetermined magnitude.

In the fourth stage, merit indices are estimated for the remaining isolated alloy compositions in the dataset. These include: creep-merit index (which describes an alloy' s creep resistance based solely on mean composition), strength-merit index (which describes an alloy’s precipitation yield strength based solely on mean composition), density, cost, stable microstructure and gamma-prime solvus temperature.

In the fifth stage, the calculated merit indices are compared with limits for required behaviour, these design constraints are considered to be the boundary conditions to the problem. All compositions which do not fulfil the boundary conditions are excluded. At this stage, the trial dataset will be reduced in size quite markedly.

The final, sixth stage involves analysing the dataset of remaining compositions. This can be done in various ways. One can sort through the database for alloys which exhibit maximal values of the merit indices - the lightest, the most creep resistant, the most oxidation resistant, and the cheapest for example. Or alternatively, one can use the database to determine the relative trade-offs in performance which arise from different combination of properties.

The seven merit indicies are now described.

The first merit index is the creep-merit index. The overarching observation is that time- dependent deformation (i.e. creep) of a nickel-based superalloy occurs by dislocation creep with the initial activity being restricted to the g phase. Thus, because the fraction of the g' phase is large, dislocation segments rapidly become pinned at the g/g' interfaces. The rate-controlling step is then the escape of trapped configurations of dislocations from g/g' interfaces, and it is the dependence of this on local chemistry - in this case composition of the g phase - which gives rise to a significant influence of alloy composition on creep properties.

A physically-based microstructure model can be invoked for the rate of accumulation of creep strain when loading is uniaxial and along the (001) crystallographic direction. The equation set is where r _m is the mobile dislocation density, F _p is the volume fraction of the g' phase, and w is width of the matrix channels. The terms s and T are the applied stress and temperature, respectively. The terms b and k are the Burgers vector and Boltzmann constant, respectively. The term is a constraint factor, which accounts

for the close proximity of the cuboidal particles in these alloys. Equation 3 describes the dislocation multiplication process which needs an estimate of the multiplication parameter C and the initial dislocation density. The term D _eff is the effective diffusivity controlling the climb processes at the particle/matrix interfaces. Note that in the above, the composition dependence arises from the two terms F _p and D _eff . Thus, provided that the microstructural architecture is assumed constant (microstructural architecture is mostly controlled by heat treatment) so that F _p is fixed, any dependence upon chemical composition arises through D _eff . For the purposes of the alloy design modelling described here, it turns out to be unnecessary to implement a full integration of Equations 2 and 3 for each prototype alloy composition. Instead, a first order merit index M _reep is employed which needs to be maximised, which is given by

where x _i is the atomic fraction of solute i in the g phase and is the appropriate interdiffusion coefficient.

The second merit index is a strength merit index. For high nickel-based superalloys, the vast majority of strength comes from the precipitate phase. Therefore, optimising alloy composition for maximal precipitate strengthening is a critical design consideration. From hardening theory a merit index for strength, M _strength , is proposed. The index considers the maximum possible precipitate strength - determined to be the point where the transition from weakly coupled to strongly coupled dislocation shearing occurs - which can be approximated using,

Where is the Taylor factor, g _APB is the anti-phase boundary (APB) energy, F _p is the volume fraction of the g' phase and b is the Burgers vector.

From Equation 5 it is apparent that fault energies in the g' phase - for example, the anti- phase boundary APB energy - have a significant influence on the deformation behaviour of nickel-based superalloys. Increasing the APB energy has been found to improve mechanical properties including, tensile strength and resistance to creep deformation. The APB energy was studied for a number of Ni-Al-X systems using density functional theory. From this work the effect of ternary elements on the APB energy of the g' phase was calculated, linear superposition of the effect for each ternary addition was assumed when considering complex multicomponent systems, resulting in the following equation,

where,x _Cr , x _MO , x _W , x _Ta , x _Nb and x _Ti represent the concentrations, in atomic percent, of chromium, molybdenum, tungsten, tantalum, niobium and titanium in the g' phase, respectively. The composition of the g' phase is determined from phase equilibrium calculations.

The third merit index is density. The density, r, was calculated using a simple rule of mixtures and a correctional factor, where, r, is the density for a given element and x, is the atomic fraction of the alloy element.

The fourth merit index is cost. In order to estimate the cost of each alloy a simple rule of mixtures was applied, where the weight fraction of the alloy element, x _i was multiplied by the current (2017) raw material cost for the alloying element, c _i .

The estimates assume that processing costs are identical for all alloys, i.e. that the product yield is not affected by composition.

A fifth merit index is based upon rejection of candidate alloys on the basis of unsuitable microstructural architecture made on the basis of susceptibility to TCP phases. To do this use is made of the d-orbital energy levels of the alloying elements (referred as Md) to determine the total effective Md level according to

where the x, represents the mole fraction of the element i in the alloy. Higher values of Md are indicative of higher probability of TCP formation.

A sixth merit index is the gamma-prime solvus temperature. The gamma-prime solvus is defined as the temperature where the volume fraction of gamma-prime tends to zero. This is determined using thermodynamic calculations - as previously described above in the second step of the Alloys-by-Design method. The phase diagram and thermodynamic properties for a specific alloy composition is calculated and used to find the temperature at which this phase transition occurs.

A seventh merit index is solid solution merit index. Solid solution hardening occurs in the (FCC) matrix phase which is referred to as gamma (b ), in particular this hardening mechanism is important at high temperatures for high strength and creep resistance. A model which assumes superposition of individual solute atoms on the strengthening of the matrix phase is employed. The solid solution strengthening coefficients, k _i , for the elements considered in the design space: aluminium, cobalt, chromium, molybdenum, niobium, tantalum, titanium and tungsten are 225, 39.4, 337, 1015, 1183, 1191, 775 and 977 MPa/at.% ^1/2 , respectively (H. Roth, C. Davis, and R. Thomson: Metallurgical and Materials Transactions A, 1997, vol. 28, pp. 1329— 1335). The solid-solution index is calculated based upon the equilibrium composition of the matrix phase using the following equation,

where, M _{solid-solution} is the solid solution merit index and x _i is the concentration of element i in the g matrix phase.

The ABD method described above was used to isolate the inventive alloy composition. The design intent for this alloy was develop a cast & wrought (C&W) type alloy with an improved balance of high strength and good hot workability. This is achieved in combination with a reduction in alloy cost, a high level of oxidation and corrosion resistance, good microstructural stability and good alloy density. The balance of properties for the new alloy provides substantial benefit over other C&W alloys described in the prior art, particularly for use in structural applications where the stresses are high and the temperatures are in excess of 650°C or greater.

The material properties - determined using the ABD method - for the nominal compositions of a number of commercial C&W alloys used high strength and high temperature applications, listed in Table 1, are listed in Table 3. The design of the new alloy was considered in relation to the predicted properties listed for these alloys.

The rationale for the design of the new alloy is now described. Table 3: Calculated phase fractions, misfit and merit indices made with the“ Alloys -by-De sign” software. Results for nickel-based superalloys listed in Table 1.

In nickel based superalloys used in the cast & wrought (C&W) form there is a balance between the ability to hot work (‘hot workability’) and the strength of the alloy. Figure 1 shows that generally as the main strengthening phase, g’, is increased there is a correlation of increasing g’ solvus temperature. Typically the alloys of this invention are hot worked near or above solvus temperature. Therefore a higher g’ solvus makes hot working more difficult as this makes the material more resistant to deformation at high temperature. Moreover, as hot working temperature increases thermally induced stresses which accumulate on cool down after hot working may be increased, this can lead to cracking upon cooling of the hot worked material, for example, quench cracking or strain-age cracking mechanisms.

In Figure 2 it is seen that increasing strength - determined in terms of strength merit index - at the desired operating temperature of the alloy (approximately 700-750°C) correlates with an increase in g’ solvus temperature. Hence from Figure 1 and Figure 2 it is demonstrated that there is complex trade-off between the ability to manufacture the alloy through hot working processes and the overall strength during operation.

Figure 1 and Figure 2 show that although the trade-off exists there are alloy compositions within the alloy design space defined in Table 2 which provide an improvement in the combination of hot workability and high temperature strength beyond the current alloys listed in Table 1. It is the aim of this invention to achieve an improved mechanical strength compared to alloys such as Waspaloy or Haynes 282, preferably the targeted strength being that of high strength alloys such as Rene65, AD730, 720Li and TMW-4. This is achieved by increasing g’ volume fraction to levels higher than Waspaloy or Haynes 282, described in relation to Figure 2. In combination with increased mechanical strength a solvus lower than current high strength alloys, Rene65, AD730, 720Li and TMW-4 is desired to maintain better hot workability.

Figure 1 shows correlation between g’ volume fraction and g’ solvus temperature for the composition space listed in Table 2. It is seen that for a given g’ volume fraction there is a spread in g’ solvus temperature of up to approximately 100°C. For an alloy with a good combination of high temperature strength and hot workability it is preferred to achieve a maximal value of g’ volume fraction and a minimal g’ solvus temperature. Alloys Rene 65 and AD730 have a solvus of 1100°C. In the present invention it is desired to have a solvus of 1080°C or less to ensure improved formability, hence g’ volume fraction is limited to 48% at 760 °C. It described later with reference to Figures 10-11 that in the invention it is desirable that g’ volume fraction is limited to less 44% or less to ensure a combination of strength, oxidation resistance and alloy microstructural stability. In certain cases it is desirable that the g’ solvus is limited to 1050°C to further improve hot workability and more preferably 1025°C for even better hot workability. Thus volume fraction of g’ is preferably limited to less than 40%, more preferably 36%. This may come at the expense of lower strength.

Figure 2 describes the trade-off between the volume fraction of g’ phase and the strength of the alloy, in terms of strength merit index. It is seen that as volume fraction of g’ increases there is a general trend for increased alloy strength. However for any given volume fraction of g’ there is the possibility to modify strength through selection of optimal alloy chemistry, described later with reference to Figure 3 and Figures 5-8. The strength levels of a number of high strength C&W superalloys is identified on Figure 2. It is desired that the alloy of the present invention has a strength of at least equivalent to AD730 or Rene65. To achieve this a minimum of 29% g’ volume fraction at 760 °C is required. Preferably an alloy with strength equivalent to alloy 720Li is desired, more preferably an alloy with a strength equivalent to alloy TMW-4 is desired. Therefore it is more preferable to have at least a minimum of 31% g’ and even more preferable to have a g’ volume fraction of 40% or greater.

Figure 3 describes the relationship between alloy strength and the ratio of elements according to the formula (0.6W _Ti + 0.31W _Nb + 0.15W _Ta )W _A ) . Dashed lines delineated on the graph identify the limits for strength at different ranges of g’ volume fraction described in the previous sections. The substitution of the elements titanium, niobium and tantalum for aluminium in the g’ phase increases strengthening in the alloy by increasing APB energy, see Equation 6. However, when the value is too high there is a possibility of formation of unwanted phases such as Eta and Delta phase. To avoid precipitation of these unwanted phases for alloys produced via the C&W manufacturing route, the ratio of elements according to the formula (0.6W _Ti +0.31W _Nb +0.15W _Ta /W _Al ) is desirably less than 1.1, more preferably 1.0. This is particularly relevant due to the elemental segregation which occurs during ingot production as elemental segregation can locally enrich some areas in the alloy with a higher ratio, hence it is kept at lower levels than would be achievable if the alloy was produced by a different method, such as powder metallurgy. Figure 3 suggests that when g’ volume faction is less than 44% if the ratio is 0.64 or higher the strength merit index is equal to that of Rene65 and AD730 so this is a desirable minimum level of the ratio. Figure 3also shows that when g’ is the range for better hot workability (29-36%) if the ratio is 0.72 or greater a strength equivalent to Rene65 and AD730 is achieved so this is a desirable minimum level of the ratio. If the ratio is 0.76 or greater a strength level equivalent to alloy 720Li is achieved when g’ is the range for better hot workability (29-40%) so this is a desirable minimum level of the ratio. When g’ is the most preferred range for even better hot workability (29-36%) - a ratio of 0.93 or greater achieves strength level equivalent to alloy 720Li, providing the most preferred balance of strength and hot workability so this is a desirable minimum level of the ratio.

Figure 4 describes the relationship between alloy g’ fraction (as contours) and alloy composition based upon aluminium plotted versus the sum of elements titanium, tantalum and niobium (according to the relationship 0.6W _Ti +0.31W _Nb +0.15W _Ta ). The additions of titanium, niobium and tantalum have been given a weighting factor to convert the weight percent addition to an“aluminium equivalent”. This allows for direct comparison of elements which have very different densities. For example titanium has a density of 4.5g/cm ³ compared to aluminium with density of 2.7g/cm ³ , thus a factor of 0.6 is applied (i.e. 2.7/4.5=0.6). Similar to titanium, a constant is added to convert the elemental additions of niobium (8.57 g/cm ³ ) and tantalum (16.4 g/cm ³ ) to an “aluminium equivalent”, thus, niobium and tantalum have correctional factors (determined from their density relative to aluminium) of 0.31 and 0.15 respectively. From Figure 4 it can be determined that the additions of aluminium, titanium, tantalum and niobium determine the volume fraction g' in accordance with the following equation

f(g') = 0.6 W _Ti + 0.31W _Nb + 0.15 W _Ta + 0.79 W _Al wheref(g') is a numerical value which must be greater than 3.4 to achieve a gamma-prime fraction of 29% or greater and W _Ti , W _Nb , W _Ta , and W _AI are the weight percent of titanium, niobium, tantalum and aluminium in the alloy respectively. Preferably to achieve a gamma- prime volume fraction greater than 31% the numerical range for f(g’) is preferably 3.6. A combination of high strength and a high oxidation and corrosion resistance is achieved by including chromium contents of greater than 17.0 wt.%. This is higher than the high strength alloys (AD730, 720Li and TMW-4) listed in Table 3 providing an improvement in oxidation compared to these alloys. A high chromium content is desirably achieved whilst maintaining alloy stability i.e. substantially free from TCP phases. Therefore it is beneficial to have a gamma-prime fraction of less than 44% (described later with reference to Figures 10-11). To achieve a gamma-prime fraction of 44% or less the numerical value for f(y’) must be 4.6 or less. Thus the numerical range for f(y’) is preferably in the range between 3.6 and 4.6. For better hot workability it is more preferable that the alloy has a g’ solvus temperature of less than 1050°C, limiting the g’ to 40% or less to achieve that the numerical value for f(g’) must be 4.3 or less. Even more preferably the alloy g’ solvus temperature should be less than 1025°C. This limits the g’ fraction to 36% or less. To achieve that the numerical value for f(g’) must be 4.0 or less.

The y’ volume fraction at 760 °C is measured experimentally by the following procedure. After a substantially long thermal exposure at 760°C the specimen (e.g. 1cm ³ ) is quenched in water and a section is taken through the material and polished using conventional/standard metallurgical preparation techniques for scanning electron microscopy. Once prepared the g/g’ microstructure should be observable in a scanning electron microscope, particles of diameter 30nm or greater should be observable. A minimum of 10 images are taken which provide a statistically representative dataset, the images should cover an area of at least lmm ² . The 2-dimensional images which reveals the g/g’ microstructure should be processed to identify the gamma-prime phase, the area fraction of the g’ phase should be measured. The area fraction of the phase is taken to be the volume fraction of g’.

From Figure 4 it is determined that the minimum required aluminium content to achieve the combination of strength, processability and gamma-prime phase stability- achieved when (0.6W _Ti +0.31W _Nb +0.15W _Ta /W _Al )<=1.1 - is greater than or equal to 1.8 wt.%. Preferably when (0.6W _Ti +0.31W _Nb +0.15W _Ta /W _Al )<=1.0 - for better y’ phase stability - the minimum aluminium content should be greater than 1.9 wt.%. The maximum aluminium content in the alloy is limited to 3.2 wt.%, when g’ volume fraction is 44% and

(0.6W _Ti +0.31W _Nb +0.15W _Ta /W _Al )>=0.64. Preferably when g' is limited to less than 40% and (0.6W _TI +0.3 1 W _Nb +0. 15W _Ta /W _Al )>=0.76 the aluminium content should be limited to 2.75 wt.% Most preferably when g’ is less than 36% it is better to have a (0.6W _TI +0.3 lW _Nb +0.15W _Ta /W _Al )>= 0.72, therefore an aluminium content of less than 2.6 wt.% is desirable for an improved balance of strength and hot workability. Even more preferably for a better combination of strength and hot workability (0.6W _TI +0.3 1 W _Nb +0. 15 W _Ta /W _Al )>= 0.93 and the aluminium in the alloy should be limited to 2.3 wt.%.

For the alloys of the present invention high levels of tensile and creep properties are needed at temperatures of 700°C or greater. However in combination with these characteristics it is desirable that the alloy has a high resistance to oxidation. Oxidation can be source of failure in some application areas of the invention. In particular high temperature high stress applications in a gas turbine oxidation assisted cracking mechanisms may occur, for example dwell fatigue in turbine disc alloys. This cracking mechanism is accelerated by rapid oxidation - and the formation of non-protective oxides, such as niobium of titanium oxides - in the alloy whilst operating in extreme environments. Formation of protective oxides - by improving alloys oxidation resistance - can help arrest crack growth. To reduce susceptibility to oxidation cracking the present invention aims to increase levels of elements which improve oxidation resistance, mainly chromium, and limit the use of elements which may negatively influence oxidation, mainly titanium and niobium which can form non-protective oxides.

To provide a resistance to oxidation a high level of chromium is preferred. High levels of chromium also improve resistance to hot corrosion. A chromium level of at least 17.0 wt.% is needed. This allows for a better oxidation and corrosion resistance than other high strength alloys (TMW-4, Alloy 720Li, AD730, Rene 65). However as chromium is added the propensity for an alloy to form deleterious TCP phases, such as sigma (s) phase is increased. Limiting or stopping the precipitation of TCP phase formation is beneficial as these phases lead to deterioration in material properties during high temperature operation. The property trade-off between oxidation resistance and high temperature strength are described later with reference to Figures 10-11.

Titanium is known to degrade the oxidation performance of chromia forming superalloys due to formation of titanium oxides that limit the chromium based oxide-scales protective capability and accelerate oxidation kinetics, (see S. Cruchley, et al., Chromia layer growth on a Ni-based super alloy: Sub-parabolic kinetics and the role of titanium, Corrosion Science 75:58-66, 2013) In the present invention titanium is preferably limited to 3.5 wt.% titanium to have a better resistance to oxidation. An even lower level of titanium is preferred to have a better level of oxidation resistance. Preferably titanium is limited to 3.0 wt% as this will allow for oxidation equivalent to that of Waspaloy which is known to have very good oxidation properties. Waspaloy is also known to have good resistance to oxidation assisted cracking mechanisms such as dwell fatigue.

Similarly to protect against dwell fatigue the niobium content of the alloy should be limited to 3.0 wt.% or less. Niobium is known to accelerate the growth of oxidation assisted cracks if it forms niobium oxide which has a high Pilling-Bedworth ratio, meaning the volumetric expansion of the niobium oxide is high causing stresses to build up in the alloy as a result of the volume increase. In an embodiment particularly suitable for use in oxidising environments niobium is less than 2.0 wt.% as alloys containing less than 2.0 wt% do not show much formation of niobium oxides when exposed in oxidising environments meaning there will be no negative influence of alloy ( see Nemeth A.A.N et al., On the Influence ofNb/Ti Ratio on Environmentally Assisted Crack Growth in High-Strength Nickel Based Superalloys, Metal. And Mat Trans A 49:3923-3937, 2018 ).

The main alloying additions used to increase alloy strength - determined in terms of the strength merit index - are the gamma-prime (g’) forming elements, aluminium, titanium, niobium and tantalum. Figures 5-8 show the influence of the elements aluminium, titanium and niobium on strength for alloys.

From Figures 5-8 a relationship between elemental additions (based upon titanium, aluminium and niobium) and target alloy strength was derived. The relationship - which describes a 3 -dimensional surface - was determined in the following way. A relationship between alloy strength and aluminium and titanium was determined from Figures 5-8 for the strength target of 1450MPa or greater, this was determined for the contour lines for these values. The translation of the line as a function of different niobium contents was determined from Figures 5-8. From the process described it is determined that additions of aluminium, titanium and niobium adhere to the following equation (note a factor of 0.44 is applied for niobium instead of 0.31 determined based its influence on g’ formation because it has a significant influence on APB energy strengthening, see Equation 6. From Equation 6 the APB strengthening factor for niobium is 21.4, the factor applied to titanium is 15, therefore niobium 1.42 is times more potent in strengthening. This factor is multiplied by the aluminium equivalent of niobium (0.31), thus a factor of 0.44 is derived (i.e. 1.42 x 0.31 = 0.44)).

f (strength) = 0.6W _Ti +0.44W _Nb +0.2 W _Al

where f( strength) is a numerical value and W _Ti , W _Nb and W _Al are the weight percent of titanium, niobium, aluminium in the alloy respectively. In order to produce an alloy with a strength merit index greater than or equal to 1450 MPa the numerical value for f( strength) should be greater than or equal to 2.8.

At the maximum niobium content of 3.0 wt.% a minimum of 1.8 wt.% titanium is required (Figure 8) when g’ volume faction is 44% or less (f( g')£4.6) to achieve a strength merit index of 1450MPa or greater. Thus the alloy has a maximum niobium content of 3.0%. Preferably g’ volume fraction is limited to 40% (f( g')£ 4.3 ) and even more preferably less than 36% (f( g')£4.0 ), therefore it is preferable to have a titanium content of greater than 1.9 wt.% and 2.0 wt.% respectively (Figure 8). At the preferred niobium content of 2.0 wt.% a minimum of 2.3 wt.% titanium is required (Figure 7) to achieve a g’ volume faction is 44% or less ( f(g'£)4.6 ) to achieve a strength merit index of 1450MPa or greater. Preferably g’ volume fraction is limited to 40% (f( g')£ 4.3 ) and even more preferably less than 36% (f( g')£4.0 ), therefore it is preferable to have a titanium content of greater than 2.4 wt.% and 2.5 wt.% respectively (Figure 7).

At the maximum titanium content of 3.5 wt.% it is difficult to achieve the strength merit index of 1450MPa when the alloy does not include niobium (Figure 5) (though such an alloy does have exceptional hot workability due to low g' solvus). When the alloy is free from niobium a minimum of 4.0wt.% titanium achieves a strength merit index of 1450MPa (Figure 5). Therefore the maximum amount of titanium is 4.0%. Aadditions of niobium can be used to achieve the strength merit index of 1450MPa when titanium is 3.5wt.% or less. Niobium additions equivalent to 0.5wt.% titanium are required to achieve the strength merit index. Therefore a minimum of 0.4 wt.% of niobium is desired in the alloy to achieve the desired strength merit index of 1450MPa, (determined in the following way, 0.5wt.% titanium divided by factor of 0.6 times 0.44 is equivalent to 0.4 wt.% niobium). More preferably titanium is limited to 3.0 wt.% meaning that 1.0 wt.% equivalent of niobium is required, this is 0.7 wt.% niobium, to achieve a target strength merit index of 1450MPa. Most preferably titanium is limited to 3.0 wt.% and a strength merit index of 1550MPa is achieved, in this instance a niobium equivalent of 1.6 wt.% titanium is required, this is equal to a niobium content of 1.2wt.% or greater.

Tantalum is an optional element present in an amount of 3.05 wt.% or less. This limit is to control alloy cost (see Figure 12). In an embodiment of the invention tantalum is added up to a level 3.05 wt.% in substitution for niobium, a substitutional factor of 0.66Ta is derived based on the density of Tantalum (.15/.31) and the strengthening coefficient (27.1/21.4). When Tantalum is at 3.05 wt.% and the alloy is free from niobium it is still possible to achieve an alloy with high strength merit index. Preferably tantalum is limited to 2.25wt.%, more preferably 1.5 wt.% and most preferably 0.65 wt.% to control cost, described later with reference to Figure 12.

f( strength) = 0.6 W _Ti + 0.44(W _Nb + 0.66 W _Ta ) + 0.2 W _Al

In order to produce an alloy with a strength merit index greater than or equal to 1450 MPa the numerical value for f( strength) should be greater than or equal to 2.8.

Along with increased strengthening contribution from precipitate hardening resulting from the gamma-prime phase (calculated in terms of strength merit index) it is desirable to design an alloy which has a strong gamma matrix phase. The strength of the gamma matrix can be calculated in terms of a solid solution merit index. The solid solution strengthening is particularly important for imparting high temperature strength (both tensile and creep strength) in the alloy. The solid solution strengthening of the gamma phase of the alloy is mainly dependent upon the additions of elements tungsten and molybdenum, as the coefficients for these elements is large and they strongly partition to the gamma-phase, unlike niobium and tantalum which also have large strengthening coefficients but they have limited partitioning to the gamma phase-see Equation 10. Figure 9 shows the relationship between molybdenum, tungsten and solid solution index. A coefficient of 0.5 is applied to the tungsten content as it is approximately twice the density of molybdenum; this factor accounts for differences in density. To provide equivalent or better creep strength than alloys in Table 3 while maintaining an improved level of workability it is desirable for the solid solution index to be greater than 97MPa (equivalent to TMW-4). From Figure 10 it is possible to determine a relationship between additions of tungsten and molybdenum and solid solution index

f (solid solution) = W _Mo + 0.5W _W where f(solid solution) is a numerical value which must be 3.4 or greater and W _MO , and Ww are the weight percent of molybdenum and tungsten in the alloy respectively in order to achieve an alloy with a solid solution merit index greater than or equal to 95 MPa and this is desirable. Desirably /(solid solution) is 3.6 or greater, achieving a solid solution merit index of greater than 100MPa.

Figures 10 describes the effect of g’ volume fraction and the sum of the elements molybdenum and tungsten (according to the relationship W _MO +0.5W _W ) at the maximum achievable chromium content for a target stability number of 0.9. The correctional factor of 0.5 is applied to tungsten as it has a density approximately twice that of molybdenum, this factor accounts for the substantial difference in density of the elements.

Based upon the strength requirements - in terms of strength merit index - for the alloy the g’ volume fraction is greater than 29%. To achieve creep resistance and high temperature strength it is necessary to have a high solid solution index, a minimum value of (W _MO +0.5W _W ) of 3.4 wt.% is required. Based upon these minimum requirements and the need for a chromium content of at least 17.0 wt.% it is seen to achieve the target stability number (Md<=0.90) along with the aforementioned desirable properties the g’ volume fraction must be limited to 44%. At a level of g’ volume fraction (29%) the W _MO +0.5W _W is limited to 6.3 wt.%. Thus the maximum Mo content is 6.3% or less. The maximum chromium content of the alloy is limited to less than 21.3 wt.% to ensure a balance of strength (g’ volume fraction and high temperature strength (W _MO +0.5W _W ), and oxidation resistance, Figure 10. Preferably the chromium content of the alloy is greater than 18.0 wt.%, this provides an even better oxidation and corrosion resistance in the alloy. For this level of chromium the volume fraction of g’ phase is preferably limited to 40% and the sum of the elements molybdenum and tungsten - according to the relationship Mo+0.5W - is limited to 5.6 wt.%

Preferably, a stability target of less than 0.89 is desired in order to ensure a better level of microstructural stability and avoid TCP formation. The alloys with the most desirable stability on Table 3 have a stability number of 0.89 or less. Figure 11 describes the effect of g’ volume fraction and the sum of the elements molybdenum and tungsten (according to the relationship W _MO +0.5W _W ) on the maximum achievable chromium content for a target stability number of 0.89. When the stability number is 0.89 it is preferred that the chromium content of the alloy is limited to less than 19.5 wt.%. The maximum volume fraction of g’ phase is preferably limited to 39% or less. Preferably the sum of the elements molybdenum and tungsten - according to the relationship W _MO +0.5W _W - is limited to 5.1 wt.%.When the preferred chromium content of greater than 18 wt.% is included in the alloy. For this level of chromium the volume fraction of g’ phase should be limited to 36% and the sum of the elements molybdenum and tungsten - according to the relationship W _MO +0.5W _W - is limited to 4.3 wt.%

For the cast & wrought alloys defined in Table 1 alloy AD730 is the lowest cost alloy which is suitable for high stress applications at operation temperatures of 700°C or higher. The cost of the alloy - based on current elemental prices (2019) - is estimated to be $14.6/kg. The most prominent cast & wrought alloy for these applications is alloy 718, an example of alloy 718 composition is listed (Ni-19Cr-0.8Ti-0.5Al-3Mo-29.7Fe-5Nb-0.06C) - the estimated cost of this alloy is $9.9/kg. However, IN718 is restricted to an operational temperature of less than 650°C in high stress applications. In the present invention a target cost of 14$/kg which is lower than AD730 which is significantly lower than AD730 is desired. The main elements which affect cost in the composition space defined in Table 2 are tantalum and cobalt, the effect of the elements on alloy cost is shown in Figure 12. Therefore a maximum cobalt content of 7.1 wt.% is desired, preferably the alloy cost is limited to 13 $/kg so it is preferable to limit cobalt to less than 5.3 wt.%. more preferably the alloy cost is limited to 12 $/kg so it is preferable to limit cobalt to less than 3.5 wt.%, most preferably alloy cost is limited to less than 1 l$/kg so it is most preferable to limit cobalt to less than 1.6wt.%.

From Figure 12 it is determined that the additions of cobalt and tantalum adhere to the following equation

f(cost ) = W _Ta + 0.43W _Co

where f(cost) is a numerical value and W _Co , is the weight percent of cobalt in the alloy. In order to produce an alloy with a cost less than or equal to 14$/kg the numerical value for f(cost) should be less than or equal to 3.05. Preferably the numerical value for f(cost) should be less than 2.25, more preferably 1.5 and most preferably less than 0.65 to produce an alloy with an elemental cost of less than 13$/kg, 12$/kg and 11 $/kg respectively.

Iron may be added to the alloy to reduce cost and increasing the ability for the alloy to be recycled. Additions of iron may result in increased microstructural instability. Limiting iron additions to a level of 10.0 wt.% produces a good balance of low cost, improved recyclability and microstructural stability, more preferably iron is limited to less than 6 wt.%, more preferably iron ranging between 1.0 wt.% and 5.0 wt.% is desirable as this provides the best balance between cost, recyclability and alloy performance.

The elements tungsten and tantalum can significantly increase strength and creep resistance in the alloy. However as these elements have a density much greater than nickel they also can increase alloy density. It is important to add these elements in a manner which balances the increased strength against increasing density. Figure 13 shows the effect of these elements on alloy density for alloys containing between 29-44% volume fraction of g’ phase. A target density of 8.6 g/cm ³ or less is desired such that density is not increased over the alloys listed in Table 3. Therefore tungsten additions are limited to 4.9 wt.% or less. As mentioned tungsten also has a role in providing strength and creep resistance through solid solution strengthening, it is determined previously that Mo+0.5W>3.4wt.% to achieve the necessary solid solution strengthening. Based upon the maximum tungsten addition of 4.9 wt.% molybdenum additions of at least 0.9 wt.% are used. More preferably tungsten is limited to 2.7 wt.% to achieve a density of less than 8.5 g/cm ² and most preferably tungsten is limited to less than 0.7 wt.% to achieve a density of 8.4 g/cm ² . Therefore it is preferable to have a minimum molybdenum content of at least 2.0wt.%, most preferably 3.0 wt.%.

Additions of carbon, boron and zirconium are required in order to provide strength to grain boundaries. This is particularly beneficial for the creep and fatigue properties of the alloy. Carbon concentrations should be limited to 0.1 wt.% or less, more preferably carbon is imited to 0.07 wt.% or less as this improves forgeability of the alloy. The boron concentration should range between 0.001 and 0.1 wt.%, preferably less than 0.03 wt.% as boron separated to the liquid phase during solidification and may lead to liquation cracking during welding of the alloy, more preferably less than 0.02 wt.% to ensure good weldability. The zirconium concentrations should range between 0.001 wt.% and 0.5 wt.%, preferably less than 0.01 wt.%, more preferably less than 0.006 wt.%. Magnesium additions can be used to improve the high temperature ductility of the wrought alloys improving the hot workability and also creep rupture life. Additions of magnesium up to 200PPM are desirable, preferably magnesium additions between 10PPM and 100PPM are desired in the alloy for improved hot workability and creep rupture properties. It is beneficial that when the alloy is produced, it is substantially free from incidental impurities. These impurities may include the elements sulphur (S), manganese (Mn) and copper (Cu). The element sulphur should remain below 0.003 wt.% (30 PPM in terms of mass). Manganese is an incidental impurity which is limited to 0.25 wt.%, preferably this limited to less than 0.1 wt.%. Copper (Cu) is an incidental impurity which is preferably limited to 0.5 wt.%. The presence of Sulphur above 0.003 wt.%, can lead to embrittlement of the alloy and sulphur also segregates to alloy/oxide interfaces formed during oxidation, preferably sulphur levels of less than less than 0.001 wt.%. Vanadium is an incidental impurity, vanadium negatively influences the oxidation behaviour of the alloy and is which is preferably limited to 0.5 wt.%, preferably less than 0.3 wt.% and most preferably this limited to less than 0.1 wt.%. This segregation may lead to increased spallation of protective oxide scales. If the concentrations of these incidental impurities exceed the specified levels, issues surrounding product yield and deterioration of the material properties of the alloy is expected.

Additions of hafnium (Hf) of up to 0.5 wt.%, or more preferably up to 0.2wt.% are beneficial for tying up incidental impurities in the alloy and also for providing strength. Hafnium is a strong carbide former it can provide additional grain boundary strengthening.

Additions of the so called ‘reactive-elements’, Yttrium(Y), Lanthanum (La) and Cerium (Ce) may be beneficial up to levels of 0.1 wt.% to improve the adhesion of protective oxide layers, such as Cr ₂ O ₃ . These reactive elements can‘mop-up’ tramp elements, for example sulphur, which segregates to the alloy oxide interface weakening the bond between oxide and substrate leading to oxide spallation. Additions of Silicon (Si) up to 0.5 wt.% may be beneficial, it has been shown that additions of silicon to nickel based superalloys at levels up to 0.5 wt.% are beneficial for oxidation properties. In particular silicon segregates to the alloy/oxide interface and improves cohesion of the oxide to the substrate. This reduces spallation of the oxide, hence, improving oxidation resistance. Based upon the description of the invention presented in this section the broad range for the invention is listed in Table 4. A preferable range is also given in Table 4.

Table 4: Compositional range in wt. % for the newly design alloy.

The following Section describes example compositions for the present invention. The calculated properties for these new alloys are listed. The rationale for the design of these alloys is now described.

Three examples of the invention are described in Table 5. The predicted properties of the alloys are listed in Table 6. The alloys listed have the benefit of an improved combination of hot workability determined by a lower ɤ’ solvus temperature and high strength - in terms of strength merit index - over the prior art alloys listed in Table 3, particularly when compared to other high strength alloys with a strength merit index of 1450MPa or greater. These improvements in material performance are attained whilst maintaining a lower elemental cost than the alloys which have a strength index of greater than 1450MPa.

Table 5 Examples of alloys in the present invention compared with example alloys listed in Table 1

Inspection between the alloys of the invention and the prior art shows that the alloys of the invention have a much lower level of cobalt. Cobalt is typically added to the prior art alloys as an element to lower gamma-prime solvus. Most of the prior art will identify their preferred range of cobalt based on that technical effect i.e. increase cobalt will reduce solvus and therefore improve workability. The alloys described here on the other hand balance other elements, such as gamma-prime forming elements- to achieve a low gamma-prime solvus with a low cobalt content. This provides an alloy with good hot workability and a low elemental cost.

Table 6: Calculated phase fractions and merit indices made with the“Alloys-by-Design” software.

Results for compositions listed in Table 5 and example alloys listed in Table 1.