EP2805784A1 - Nickellegierung - Google Patents

Nickellegierung Download PDF

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Publication number
EP2805784A1
EP2805784A1 EP20140157622 EP14157622A EP2805784A1 EP 2805784 A1 EP2805784 A1 EP 2805784A1 EP 20140157622 EP20140157622 EP 20140157622 EP 14157622 A EP14157622 A EP 14157622A EP 2805784 A1 EP2805784 A1 EP 2805784A1
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alloy
atomic
following composition
bal
forging
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English (en)
French (fr)
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EP2805784B1 (de
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Mark Hardy
Howard Stone
Paul Mignanelli
Bryce Conduit
Gareth Conduit
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Rolls Royce PLC
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Rolls Royce PLC
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C19/00—Alloys based on nickel or cobalt
    • C22C19/03—Alloys based on nickel or cobalt based on nickel
    • C22C19/05—Alloys based on nickel or cobalt based on nickel with chromium
    • C22C19/051—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W
    • C22C19/056—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being at least 10% but less than 20%
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22F—WORKING METALLIC POWDER; MANUFACTURE OF ARTICLES FROM METALLIC POWDER; MAKING METALLIC POWDER; APPARATUS OR DEVICES SPECIALLY ADAPTED FOR METALLIC POWDER
    • B22F3/00—Manufacture of workpieces or articles from metallic powder characterised by the manner of compacting or sintering; Apparatus specially adapted therefor ; Presses and furnaces
    • B22F3/17—Manufacture of workpieces or articles from metallic powder characterised by the manner of compacting or sintering; Apparatus specially adapted therefor ; Presses and furnaces by forging
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22F—WORKING METALLIC POWDER; MANUFACTURE OF ARTICLES FROM METALLIC POWDER; MAKING METALLIC POWDER; APPARATUS OR DEVICES SPECIALLY ADAPTED FOR METALLIC POWDER
    • B22F3/00—Manufacture of workpieces or articles from metallic powder characterised by the manner of compacting or sintering; Apparatus specially adapted therefor ; Presses and furnaces
    • B22F3/24—After-treatment of workpieces or articles
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22F—WORKING METALLIC POWDER; MANUFACTURE OF ARTICLES FROM METALLIC POWDER; MAKING METALLIC POWDER; APPARATUS OR DEVICES SPECIALLY ADAPTED FOR METALLIC POWDER
    • B22F3/00—Manufacture of workpieces or articles from metallic powder characterised by the manner of compacting or sintering; Apparatus specially adapted therefor ; Presses and furnaces
    • B22F3/24—After-treatment of workpieces or articles
    • B22F2003/248—Thermal after-treatment
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22F—WORKING METALLIC POWDER; MANUFACTURE OF ARTICLES FROM METALLIC POWDER; MAKING METALLIC POWDER; APPARATUS OR DEVICES SPECIALLY ADAPTED FOR METALLIC POWDER
    • B22F2203/00—Controlling
    • B22F2203/11—Controlling temperature, temperature profile
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22F—WORKING METALLIC POWDER; MANUFACTURE OF ARTICLES FROM METALLIC POWDER; MAKING METALLIC POWDER; APPARATUS OR DEVICES SPECIALLY ADAPTED FOR METALLIC POWDER
    • B22F2301/00—Metallic composition of the powder or its coating
    • B22F2301/15—Nickel or cobalt
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22F—WORKING METALLIC POWDER; MANUFACTURE OF ARTICLES FROM METALLIC POWDER; MAKING METALLIC POWDER; APPARATUS OR DEVICES SPECIALLY ADAPTED FOR METALLIC POWDER
    • B22F2998/00—Supplementary information concerning processes or compositions relating to powder metallurgy
    • B22F2998/10—Processes characterised by the sequence of their steps

Definitions

  • This invention relates to nickel (Ni) base alloys, and particularly, though not exclusively, to alloys suitable for use in compressor and turbine discs of gas turbine engines. Such discs are critical components of gas turbine engines, and failure of such a component in operation can be catastrophic.
  • Some known nickel base alloys have compromised resistance to surface environmental degradation (oxidation and Type II hot corrosion) in order to achieve improved high temperature strength and resistance to creep strain accumulation, and in order to achieve stable bulk material microstructures (to prevent the precipitation of detrimental topologically close-packed phases).
  • Turbine discs are commonly exposed to temperatures above 650°C, and in future engine designs will be exposed to temperatures above 700°C. As disc temperatures continue to increase, oxidation and hot corrosion damage will begin to limit disc life. There is therefore a need, in the design of future disc alloys, to prioritise resistance to dwell crack growth and oxidation and hot corrosion damage ahead of other properties.
  • the invention provides a nickel base alloy, and a method of making such an alloy, as set out in the claims.
  • the aim was to produce alloys in which the disordered face-centred cubic gamma ( ⁇ ) phase is precipitation strengthened by the ordered L 1 2 gamma prime ( ⁇ ') phase.
  • the inventors have determined that the following two groups of compositions produce the required balance between high temperature proof strength, resistance to fatigue damage and creep strain accumulation, damage tolerance and oxidation / hot corrosion damage.
  • the first group (examples of which are shown in Table 1) has a bias towards providing a higher resistance to oxidation / hot corrosion damage and shows a minimum amount of strengthening precipitates that is necessary to provide the required high temperature proof strength and resistance to creep strain accumulation. This resistance to environmental damage is superior to that shown by existing powder nickel alloys that are currently used for disc rotor applications.
  • the second group shows a level of environmental resistance, which is at least equivalent or better than the existing powder nickel alloys but offers a higher level of high temperature proof strength and resistance to creep strain accumulation than the first group of compositions.
  • Group 2 in which Al can have values from about 6.5 to about 8.25 at.%, Ti can have values between about 2.25 and about 3.5 at.%, Nb can have values between about 0.75 and about 2.25 at.%, Ta can have values between about 0.75 and about 1.5 at.% and Si can have values between zero and about 1.2 at%.
  • Phase diagram modelling using JMat Pro version 6.1 and the Thermatech Nickel-8 database, indicates that Group 1 and 2 alloys are expected to have 48-50 and 53-55 mole % of ⁇ ' precipitates respectively at 800°C. The predicted content of ⁇ ' and other phases are shown in Figures 1 and 2 for alloys V207S125C and V207S135A respectively at temperatures between the respective liquidus temperatures and 600°C.
  • Table 3 also provides the predicted mole % of ⁇ ' at 800°C for the proposed compositions and other solvus data, which will be discussed later.
  • alloys which show combinations of Al, Ti, Nb, Ta and Si additions between12.5 and 13.5 atomic % offer a further balance of properties in terms of environmental resistance and high temperature strength/ resistance to primary creep. Examples of such alloys are presented in Table 4.
  • a high volume fraction of small ⁇ ' precipitates will effectively hinder the movement of dislocations and will give rise to good high temperature proof strength.
  • secondary ⁇ ' particles of 50 to 300 nm in size should be developed after quenching the alloy from the solution heat treatment temperature. It is proposed that such particle sizes can be achieved in large diameter (500-700 mm) forgings using forced compressed/fan air cooling, particularly if attention is paid to the definition of compositions in which the temperature for secondary ⁇ ' precipitation is minimised. This assertion is based on the principle that rates of diffusion are less at lower temperature so that the driving force for particle coarsening is reduced. Similarly, alloy composition will determine the ⁇ ' solvus temperature, i.e. the temperature at which all ⁇ ' particles dissolve into solution.
  • the amount of Al in the alloy has the most significant influence on the ⁇ ' solvus temperature, with higher quantities increasing the solvus temperature.
  • reducing the amount of cobalt (Co), and Ti in the composition a proportion of which partitions to the ⁇ ' phase, will reduce the ⁇ ' solvus temperature.
  • Silicon is also known to reduce ⁇ ' solvus temperature.
  • V207S125E and V207S135B the addition of about 1.15 at.% Si is expected to reduce the ⁇ ' solvus temperature of the Group 1 and 2 alloys by 7-16°C and 6-11°C respectively.
  • Solution heat treatment above the ⁇ ' solvus temperature is necessary to produce the required grain size for optimising dwell crack growth resistance. If the ⁇ ' solvus is high and close to the solidus temperature of the alloy, incipient melting and grain boundary boron (B) liquation can occur during solution heat treatment, as can quench cracking of the forging from fast cooling, following solution heat treatment. Since the ⁇ ' solvus temperature determines whether the alloy is viable for high temperature discs, the alloy composition needs to be defined to minimise the ⁇ ' solvus temperature. Hence the need for careful selection of Al, Co and Ti levels.
  • Nb and Ta are important as these elements show slower rates of diffusion in Ni compared to Al and Ti, which is significant during quenching of forgings and high temperature operation in terms of reducing the rate of coarsening of secondary and tertiary ⁇ ' respectively, and in terms of resistance to oxidation damage since Al and Ti readily migrate from ⁇ ' to form oxidation products.
  • a protective chromia (Cr 2 O 3 ) scale should form as quickly as possible at temperatures above 500°C.
  • Three features of the proposed compositions facilitate this: firstly, to maximise the chromium (Cr) level in the ⁇ phase; secondly, to minimise the Co and iron (Fe) content in the ⁇ phase; and thirdly, to minimise the occurrence of rutile (TiO 2 ) by reducing the Ti content.
  • the inventors have determined that at temperatures between 500°C and 800°C, the Cr level in the ⁇ should be greater than 25 at.%, and that the combined levels of Co and Fe be below 15 at.%.
  • Figure 3 shows the predicted elemental content in the ⁇ phase for alloy V207S135B.
  • surface scales at temperatures between 650-800°C will be composed predominantly of Cr and Ti oxides, with initial transient oxides based on Cr, Ti and Ni, Fe and Co.
  • the level of Cr that can be added is limited by the propensity for topological close packed (TCP) phases such as sigma ( ⁇ ) 1 ,mu ( ⁇ ) 2 and 1 (Ni, Co,Fe) x (Cr, Mo,W) y where x and y can vary between 1 and 7
  • TCP topological close packed
  • Table 3 shows ⁇ solvus temperature for the proposed compositions.
  • the solvus temperatures for the TCP phases were predicted from phase diagram modelling using JMat Pro v6.1 and the Thermatech Nickel-8 database. They are below or around 700°C, which indicates that on the basis of thermodynamics, these TCP phases are expected to form after a long exposure. In practice, TCP phases are considered unlikely to form at these low temperatures. They are much more likely to form if the solvus temperatures were at a higher temperature such as 800 or 900°C.
  • compositions in Tables 1, 2 and 4 were evaluated by predicting the ⁇ , ⁇ and P solvus temperature for the target chemistry plus the expected deviation that would be added to create a material specification, i.e. the TCP solvus temperatures were predicted for the upper limits of the elements in the material specification. These solvus temperatures should be below 800°C. This approach ensures that all possible variations in elemental content within the material specification will be free of TCP phases. It is important that alloys are free from these phases since they often precipitate at grain boundaries, reducing grain boundary strength and corrosion/oxidation resistance. They are therefore detrimental to material properties.
  • the proposed compositions show relatively high concentrations of Cr (12.5 to 15.5 at.%) for alloys with high volume fractions of ⁇ ', the level of Cr being determined by the amount of ⁇ ', molybdenum (Mo), tungsten (W), Fe and Co present. This is made possible in the higher Cr alloys by minimising the Mo content to about 0.35 at.% and the tungsten (W) content to about 1.1 at.%. Such low levels of Mo and W will increase coherency strains that arise because the ⁇ ' phase has a larger lattice parameter than the ⁇ phase. To increase the Cr level beyond 16 at.% would require a further reduction of the ⁇ ', Mo, W, Fe and Co content.
  • Table 3 shows the predicted maximum misfit between the ⁇ and ⁇ ' phases calculated at 800°C from predictions of lattice parameter for the ⁇ (a ⁇ ) and ⁇ '(a ⁇ ' ) phases.
  • the proposed compositions show greater predicted coherency strains than RR1000, which is expected to have a value of 0.12% at 800°C (Table 3). However, the predicted values in Table 3 are not considered to be excessive and should provide some additional strengthening through coherency strain, at least at low temperatures. It is also expected that deformation behaviour in the proposed alloys will be characterised by shearing of ⁇ ' particles by either anti-phase boundary coupled super lattice dislocations or by intrinsic/extrinsic stacking faults. 4 H.A. Roth et al, (1997), Met. Trans., 28A (6), pp. 1329-1335 .
  • the Co and Fe content in the ⁇ phase has been minimised to enhance the effectiveness of Cr.
  • the Co content has been reduced compared to those levels of Co (18.5-20.5 wt.%) in existing disc alloys such as RR1000, ME3, LSHR and Alloy 10, such that Fe + Co ⁇ 8.5 at . % in which Fe can have values between about 0 and about 5.5 at.% and Co can have values between about 2.0 and about 8.5 at.%.
  • the lower Co content also has benefits in reducing the propensity for ⁇ phase precipitation and in reducing elemental costs. It is recognised, however, that Co in significant quantities (> 20 at.%) is beneficial in lowering stacking fault energy of the ⁇ phase and in promoting annealing twins.
  • Additional improvements can be made to the oxidation and hot corrosion resistance by further promoting chromia scale formation, by adding a sufficient quantity either of Si, to produce a silica (SiO 2 ) film, and/or of Mn, to produce a spinel (MnCr 2 O 4 ) film beneath the chromia scale. It is predicted that Si and Mn partition between ⁇ and ⁇ ' phases, residing predominantly in the ⁇ phase above 500 °C. At such temperatures, nickel alloys begin to show signs of oxidation damage.
  • Figure 4 shows the predicted partitioning of Si in the ⁇ and ⁇ ' phases as a function of temperature for alloy V207S125E.
  • Figure 5 (for alloy V207S125A) shows that the majority of Mn is present in the ⁇ phase.
  • Nb, Ta and Hf are added to develop stable primary MC carbides (where M can represent Ti, Ta, Nb, or Hf) ( Figure 6 ).
  • Primary carbides based on Ti are not stable and, during prolonged exposure to temperatures above 700°C, decompose to M 23 C 6 , i.e. MC + ⁇ ⁇ M 23 ⁇ C 6 + ⁇
  • M 23 C 6 carbides form as films or elongated particles on grain boundaries and can reduce creep stress rupture life if extensive films decorate grain boundaries. It is understood that the formation of M 23 C 6 carbides removes Cr from the ⁇ phase adjacent to the grain boundary, and therefore reduces the resistance to oxidation in this region. If thermal and fatigue loading conditions do not give rise to fatigue cracks, then Cr from near-surface M 23 C 6 carbides can diffuse along grain boundaries towards the surface, leaving voids. These voids are a form of internal oxidation damage, which can reduce the resistance of the alloy to fatigue crack nucleation. To optimise resistance to surface oxidation resistance, it is proposed that a further group of compositions be defined, i.e.
  • Al + Ti + Nb + Ta + 0.3 ⁇ Si 12.5 at . %
  • Group 3 such that Al is about 7.0-7.5 at.%, Ti can have values between about 1.5 and about 2.0 at.%, Nb can have values between about 2.0 and about 2.5 at.%, Ta can have values between about 1.0 and about 1.5 at.% and Si can have values between zero and about 1.2 at%. These alloys will show stable MC carbide and will be free of M 23 C 6 carbide.
  • the level of carbon (C) in the compositions is between about 0.1 and about 0.29 at.% (0.02 - 0.06 wt.%).
  • a value of about 0.03 wt.% minimises internal oxidation damage from decomposition of M 23 C 6 carbides.
  • this level of C is not as effective as 0.05 wt.% C in controlling grain growth through grain boundary pinning during super-solvus solution heat treatment. It is understood that the higher concentration of C produces a smaller average grain size and a narrow grain size distribution, with lower As Large As grain sizes. This is significant as yield stress and fatigue endurance at intermediate temperatures ( ⁇ 650°C) are highly sensitive to grain size.
  • Zr zirconium
  • B boron
  • the high temperature stabilisation heat treatment that is necessary for precipitating the required decoration of M 23 C 6 carbides on grain boundaries will also significantly coarsen tertiary ⁇ ' particles. Such coarsening will reduce the resistance of the alloy to primary creep and lower the high temperature yield stress. This is considered to be 7 H.-J. Jou et al, (2012), Superalloys 2012, (Ed. by E.S. Huron et al), The Minerals, Metals & Materials Society, Warrendale, Pennsylvania, USA, pp. 893-902 . beneficial for material ahead of the crack tip in order to improve resistance to intergranular dwell crack growth, as it results in relaxation of crack tip stresses during the dwell period.
  • finer tertiary ⁇ ' particles are required for good resistance to primary creep and high elevated temperature yield strength elsewhere in the alloy. These fine particles can be precipitated from a supplementary heat treatment, after the stabilisation heat treatment, at temperatures of between 800 and 850°C for 2 to 8 hours.
  • Nb is detrimental to dwell crack growth as a result of the oxidation of large blocky MC carbides and delta ( ⁇ ), Ni 3 Nb, phase, which reside on grain boundaries and form brittle Nb 2 O 5 . It is also understood that a small fraction of the available Nb partitions to the ⁇ phase and may segregate to grain boundaries in material ahead of a growing crack as a result of chromium 9 depletion from the ⁇ phase as chromia forms from exposure to oxygen.
  • Oxygen diffusion along grain boundaries is accelerated as a result of stress, particularly in material ahead of a crack tip during dwell fatigue cycles.
  • the formation of Nb 2 O 5 is particularly detrimental as it produces a large volume change, as indicated by the Pilling-Bedworth Ratio of 2.5 10 , and readily cracks or spalls.
  • Nb up to about 1.7 wt.%
  • Tables 1, 2 and 4 higher levels of up to about 2.5 at.% (or about 4 wt.%) Nb are proposed in the Group 3 alloys that contain lower concentrations of Ti and show no M 23 C 6 carbide. These alloys are optimised for oxidation resistance and not necessarily resistance to intergranular dwell crack growth.
  • Titanium is beneficial to nickel alloys strengthened by ⁇ ' as it supplements Al in the ordered L 1 2 gamma prime particles and gives rise to high values of antiphase boundary (APB) energy when pairs of dislocations shear ⁇ ' particles.
  • APB antiphase boundary
  • Ti also gives rise to TiO 2 (rutile) nodules that form above Cr 2 O 3 (chromia) nodules in the surface oxide scale.
  • the source of Ti for the surface rutile nodules is considered to be ⁇ ', and with the loss of Al from ⁇ ' for sub-surface alumina "fingers", a region free of ⁇ ' is produced during prolonged high temperature exposure.
  • this ⁇ ' free region shows significantly reduced proof strength compared to the base alloy and is likely to crack under conditions that lead to the accumulation of inelastic strain.
  • Ti levels have been minimised in the proposed alloys, particularly the Group 3 compositions. 11 J. Telesman et al, (2004), Superalloys 2004, (Ed. by K.A. Green et al), The Minerals, Metals & Materials Society, Warrendale, Pennsylvania, USA, pp. 215-224 .
  • Levels of trace elements S and P should be minimised to promote good grain boundary strength and mechanical integrity of oxide scales. It is understood that levels of S and P of less than 5 and 20 ppm respectively are achievable in large production size batches of material. However, it is anticipated that the benefits of the invention would still be achievable provided the level of S is less than 20 ppm, and of P less than 60 ppm, although in these circumstances the resistance of the alloys to cracking from oxidation would be inferior.
  • alloys according to the invention will be produced using powder metallurgy technology, such that small powder particles ( ⁇ 53 ⁇ m in size) from inert gas atomisation are consolidated in a stainless steel container using hot isostatic pressing or hot compaction and then extruded or otherwise hot worked to produce fine grain size billet. Increments would be cut from these billets and forged under isothermal conditions. Appropriate forging temperatures, strains and strain rates would be used to achieve the preferred average grain size of ASTM 8 to 6 (22-45 ⁇ m) following solution heat treatment above the ⁇ ' solvus temperature.
  • the proposed alloys are expected to show the following material properties compared to the existing nickel alloy RR1000, with the same grain size, and taking account of differences in density (8.21 g.cm -3 for RR1000; 8.28 - 8.5 g.cm -3 for the proposed compositions at ambient temperature).
  • the invention therefore provides a range of nickel base alloys particularly suitable to produce forgings for disc rotor applications. Components manufactured from these alloys will have a balance of material properties that will allow them to be used at significantly higher temperatures. In contrast to known alloys, the alloys according to the invention achieve a better balance between resistance to environmental degradation and high temperature mechanical properties such as proof strength, resistance to creep strain accumulation, dwell fatigue and damage tolerance. This permits the alloys according to the invention to be used for components operating at temperatures up to 800°C, in contrast to known alloys which are limited to temperatures of 700 - 750°C.
  • compositions i) definition of compositions; ii) the process routes for billet and forgings; and iii) the heat treatment of the forgings.
  • the alloys according to the invention are particularly suitable for disc rotor applications in gas turbine engines, it will be appreciated that they may also be used in other applications.
  • they would be especially suitable for use in combustor or turbine casings, which would benefit from the expected improvements in material properties, notably the improved proof strength and resistance to creep strain accumulation.
  • the temperature of the static components of the combustor and turbine will necessarily also increase.
  • Such components could be produced by powder metallurgy given the highly alloyed compositions and the ability to produce compacts that are close to the component geometry, thereby reducing the amount of material required and the time required to machine the component.

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  • Engineering & Computer Science (AREA)
  • Mechanical Engineering (AREA)
  • Chemical & Material Sciences (AREA)
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EP14157622.3A 2013-05-24 2014-03-04 Nickellegierung Active EP2805784B1 (de)

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Application Number Priority Date Filing Date Title
GBGB1309404.0A GB201309404D0 (en) 2013-05-24 2013-05-24 A nickel alloy
GBGB1318622.6A GB201318622D0 (en) 2013-05-24 2013-10-22 A nickel alloy

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EP2805784A1 true EP2805784A1 (de) 2014-11-26
EP2805784B1 EP2805784B1 (de) 2015-06-03

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EP3121298A1 (de) * 2015-07-20 2017-01-25 Rolls-Royce plc Ni-basierte legierung für bauanwendungen
RU2678352C1 (ru) * 2018-05-15 2019-01-28 Акционерное общество "Научно-производственное объединение "Центральный научно-исследовательский институт технологии машиностроения", АО "НПО "ЦНИИТМАШ" Жаропрочный сплав на основе никеля для литья рабочих лопаток газотурбинных установок
KR20190017664A (ko) * 2017-08-10 2019-02-20 미츠비시 히타치 파워 시스템즈 가부시키가이샤 Ni기 합금 부재의 제조 방법
US10358701B2 (en) 2015-04-01 2019-07-23 Oxford University Innovation Limited Nickel-based alloy
US10370740B2 (en) 2015-07-03 2019-08-06 Oxford University Innovation Limited Nickel-based alloy
EP3572541A1 (de) * 2018-05-23 2019-11-27 Rolls-Royce plc Superlegierung auf nickelbasis
CN112840054A (zh) * 2018-10-10 2021-05-25 西门子能源全球有限两合公司 基于镍的合金
EP4212639A1 (de) * 2022-01-18 2023-07-19 Garrett Transportation I Inc. Nickelbasierte legierung und turbinenrad damit

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DE102019201095A1 (de) * 2019-01-29 2020-07-30 Friedrich-Alexander-Universität Erlangen-Nürnberg Nickelbasislegierung für Hochtemperaturanwendungen und Verfahren
US11466344B2 (en) 2019-03-06 2022-10-11 Energy, United States Department Of High-performance corrosion-resistant high-entropy alloys
CN114480920B (zh) * 2021-12-31 2022-09-02 中南大学 一种3d打印用镍基高温合金粉末及其制备方法和应用
CN114672696B (zh) * 2022-03-21 2023-03-14 钢铁研究总院有限公司 一种Ni-Co基高温合金及其制备方法和应用

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EP0520464A1 (de) * 1991-06-27 1992-12-30 Mitsubishi Materials Corporation Hitzebeständige Legierung auf Nickelbasis
EP0561179A2 (de) * 1992-03-18 1993-09-22 Westinghouse Electric Corporation Legierung für eine Gasturbinenschaufel
EP1433865A1 (de) * 2002-12-17 2004-06-30 Hitachi, Ltd. Hochfeste Superlegierung auf Nickelbasis und Gasturbinenschaufeln
EP2520678A2 (de) * 2011-05-04 2012-11-07 General Electric Company Legierungen auf Nickelbasis

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Publication number Priority date Publication date Assignee Title
CA1242596A (en) * 1985-03-01 1988-10-04 Michael F.X. Gigliotti, Jr. Nickel-base superalloys especially useful as compatible protective environmental coatings for advanced superalloys
EP0520464A1 (de) * 1991-06-27 1992-12-30 Mitsubishi Materials Corporation Hitzebeständige Legierung auf Nickelbasis
EP0561179A2 (de) * 1992-03-18 1993-09-22 Westinghouse Electric Corporation Legierung für eine Gasturbinenschaufel
EP1433865A1 (de) * 2002-12-17 2004-06-30 Hitachi, Ltd. Hochfeste Superlegierung auf Nickelbasis und Gasturbinenschaufeln
EP2520678A2 (de) * 2011-05-04 2012-11-07 General Electric Company Legierungen auf Nickelbasis

Cited By (12)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US10358701B2 (en) 2015-04-01 2019-07-23 Oxford University Innovation Limited Nickel-based alloy
US10370740B2 (en) 2015-07-03 2019-08-06 Oxford University Innovation Limited Nickel-based alloy
EP3121298A1 (de) * 2015-07-20 2017-01-25 Rolls-Royce plc Ni-basierte legierung für bauanwendungen
US10287654B2 (en) 2015-07-20 2019-05-14 Rolls-Royce Plc Ni-base alloy for structural applications
KR20190017664A (ko) * 2017-08-10 2019-02-20 미츠비시 히타치 파워 시스템즈 가부시키가이샤 Ni기 합금 부재의 제조 방법
KR102078922B1 (ko) 2017-08-10 2020-02-19 미츠비시 히타치 파워 시스템즈 가부시키가이샤 Ni기 합금 부재의 제조 방법
RU2678352C1 (ru) * 2018-05-15 2019-01-28 Акционерное общество "Научно-производственное объединение "Центральный научно-исследовательский институт технологии машиностроения", АО "НПО "ЦНИИТМАШ" Жаропрочный сплав на основе никеля для литья рабочих лопаток газотурбинных установок
EP3572541A1 (de) * 2018-05-23 2019-11-27 Rolls-Royce plc Superlegierung auf nickelbasis
US11085103B2 (en) 2018-05-23 2021-08-10 Rolls-Royce Plc Nickel-base superalloy
CN112840054A (zh) * 2018-10-10 2021-05-25 西门子能源全球有限两合公司 基于镍的合金
US11441208B2 (en) 2018-10-10 2022-09-13 Siemens Energy Global GmbH & Co. KG Nickel based alloy
EP4212639A1 (de) * 2022-01-18 2023-07-19 Garrett Transportation I Inc. Nickelbasierte legierung und turbinenrad damit

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GB201318622D0 (en) 2013-12-04
EP2805784B1 (de) 2015-06-03
US20140348689A1 (en) 2014-11-27
GB201309404D0 (en) 2013-07-10

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