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US20070227630A1 - Nickel-based alloy - Google Patents

Nickel-based alloy Download PDF

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Publication number
US20070227630A1
US20070227630A1 US11/694,204 US69420407A US2007227630A1 US 20070227630 A1 US20070227630 A1 US 20070227630A1 US 69420407 A US69420407 A US 69420407A US 2007227630 A1 US2007227630 A1 US 2007227630A1
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United States
Prior art keywords
alloy
amounts
sum
temperature
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Prior art date
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Abandoned
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US11/694,204
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English (en)
Inventor
Isabelle Augustins Lecallier
Pierre Caron
Jean-yves Guedou
Didier Locq
Loeiz Naze
Current Assignee (The listed assignees may be inaccurate. Google has not performed a legal analysis and makes no representation or warranty as to the accuracy of the list.)
Association pour la Recherche et le Developpement des Methodes et Processus Industriels
Office National dEtudes et de Recherches Aerospatiales ONERA
Safran Aircraft Engines SAS
Original Assignee
Association pour la Recherche et le Developpement des Methodes et Processus Industriels
Office National dEtudes et de Recherches Aerospatiales ONERA
SNECMA SAS
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
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Publication date
Application filed by Association pour la Recherche et le Developpement des Methodes et Processus Industriels, Office National dEtudes et de Recherches Aerospatiales ONERA, SNECMA SAS filed Critical Association pour la Recherche et le Developpement des Methodes et Processus Industriels
Assigned to ONERA, ARMINES, SNECMA reassignment ONERA ASSIGNMENT OF ASSIGNORS INTEREST (SEE DOCUMENT FOR DETAILS). Assignors: AUGUSTINS LECALLIER, ISABELLE, CARON, PIERRE, GUEDOU, JEAN-YVES, LOCQ, DIDIER, NAZE, LOEIZ
Publication of US20070227630A1 publication Critical patent/US20070227630A1/en
Abandoned legal-status Critical Current

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Classifications

    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C1/00—Making non-ferrous alloys
    • C22C1/04—Making non-ferrous alloys by powder metallurgy
    • C22C1/0433—Nickel- or cobalt-based alloys
    • 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
    • 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%
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
    • C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
    • C22F1/10—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of nickel or cobalt or alloys based thereon

Definitions

  • the invention relates to alloys, or superalloys, based on nickel (Ni) and more particularly intended for the production of compressor or turbine disks for turbo-machines using powder metallurgy processes.
  • the turbo-machines concerned may be aeronautical (turbojet engine, turboprop engine) or ground-based (gas turbine for the production of energy).
  • compressor and turbine disks located respectively upstream and downstream of the combustion chamber of a turbojet engine, are subjected to mechanical stresses that can be attributed to tension, creep, and fatigue, at temperatures that can reach 800° C. Said disks should nevertheless have operational service lives of several thousand hours. Thus, said disks must be produced from an alloy which, at high temperatures, has high resistance to traction forces, very good creep strength, and crack propagation resistance.
  • said disks can be produced from nickel-based alloys using powder metallurgy processes, said processes limiting chemical segregation phenomena and encouraging good microstructural homogeneity of the alloy.
  • That example of an alloy falls into the category of two-phase alloys that comprise: a phase termed the gamma phase formed by a nickel-based solid solution that constitutes the matrix for the metallurgy grains, and a phase termed the gamma-prime phase, of structure that is based on the coherent intermetallic compound Ni 3 Al.
  • the gamma-prime phase forms several populations of inter- or intra-granular precipitates that appear at different stages of the thermomechanical history of the alloy and that play distinct roles in the mechanical behavior of the alloy.
  • the inter-granular precipitate population limits the growth of gamma matrix grains during recrystallization heat treatment.
  • the recrystallization heat treatment of the alloy controls the inter-granular precipitate population, and thus the size of said grains.
  • the maximum temperature reached during said heat treatment is higher (supersolvus treatment) or lower (subsolvus treatment) than the solution temperature (or solvus temperature) of the inter-granular precipitates of the gamma-prime phase.
  • two-phase alloys are thermomechanically treated to produce either a fine-grained microstructure (small grains), i.e. with a grain size of the order of 5 ⁇ m [micrometer] to 15 ⁇ m (i.e. ASTM [American Society for Testing and Materials] indices 12 to 9), or a microstructure with coarse grains, i.e. with a grain size of the order of 20 ⁇ m to 180 ⁇ m (i.e. ASTM indices 8 to 2).
  • a fine-grained microstructure small grains
  • ASTM American Society for Testing and Materials
  • the grain strength is ensured by the presence of different populations of intra-granular precipitates of the gamma-prime Ni 3 Al base phase and it is generally accepted that the high temperature tensile strength of said alloys increases with the volume fraction of the gamma-prime phase, said fraction possibly reaching 60%.
  • the N18 alloy with a volume fraction of the gamma-prime phase of about 55%, principally undergoes subsolvus treatments since a fine-grained microstructure is desirable.
  • the fatigue strength and tensile strength of said alloy are generally favored over its creep strength, because the service temperature is often less than 650° C., i.e. relatively moderate.
  • the invention aims to provide Ni-based alloys for which it is possible to carry out not only a subsolvus treatment, but also a supersolvus treatment on an industrial scale and which preferably has high-temperature mechanical characteristics, especially creep strength, that are at least equivalent to, and preferably better than those of N18 alloy.
  • alloys that essentially comprise (i.e. apart from any impurities) the following elements, in the amounts indicated as percentages by weight:
  • the Applicant has shown firstly, that said high volume fraction tends to reduce the difference between the solvus temperature of the gamma-prime phase and the melting temperature of the N18 alloy, rendering that difference too small to carry out a supersolvus treatment on an industrial scale.
  • the Applicant has shown that when the temperature is held at over 650° C. for a sufficiently long period, the elemental composition of the N18 alloy causes the development of topologically compact phases, generally denoted sigma and mu phases, which are deleterious to the high temperature behavior of a disk in operation.
  • composition of the alloys of the invention is selected so as to cause a limited volume fraction of gamma-prime phase to precipitate.
  • high tensile strength is particularly favorable to the rupture behavior of said disks as may occur during accidental overspeeding. This high strength is also an indicator of good oligocyclic fatigue properties and adequate service lives.
  • the reduction in the volume fraction of the gamma-prime phase relative to the N18 alloy is favorable to the production of disks having a coarse-grained microstructure and thus high creep strength at high temperature (i.e. for temperatures of 700° C. or more).
  • This creep strength associated with very good tensile and fatigue-creep crack propagation properties allows these disks to be used at temperatures that are higher than in current turbo-machines, providing access to better thermal efficiencies and a reduction in the specific consumption of the turbo-machines.
  • compositions of the alloys of the invention are such that this range spans 35° C. or more. This means that heat treatments above the solvus temperature can be carried out on an industrial scale, without risking melting the alloy.
  • this capability allows dual-structured disks to be produced.
  • a coarse-grained structure is developed in the peripheral zone of the disk where the service temperatures are the highest and where creep plays a significant role in material damage, and a small grain structure is developed in the central zone of the disk (close to the hub), which is cooler, where damage essentially results from traction forces and cyclic stresses.
  • the alloys of the invention have relatively low density, preferably 8.3 kg/dm 3 [kilograms/cubic decimeter] or less, which means that the mass of the disk and stresses resulting from centrifugal force are limited.
  • alloys of the invention provide them with good microstructural stability as regards the appearance of sigma and mu phases, which is retarded to more than 500 hours maintained at 750° C.
  • the compositions of the alloys of the invention have a limited gamma-prime phase volume fraction, preferably of 50% or less. Sufficient gamma-prime phase must nevertheless be present, so the gamma-prime phase volume fraction is preferably in the range 40% to 50%.
  • the sum of the Al, Ti, and Nb contents, as atomic percentages is 10.5% or more, and 13% or less, i.e. 10.5% ⁇ Al+Ti+Nb ⁇ 13%.
  • the elements Ti and Nb which, by being substituted for Al, are constituents of that phase are considered to be elements that are favorable to the formation of the gamma-prime phase in the same amount and they are termed gamma-prime-genic.
  • the value of the volume fraction of the gamma-prime phase is thus a function of the sum of the atomic concentrations of Al, Ti, and Nb.
  • tantalum is also a gamma-prime-genic element, but it does not appear in the composition of the alloys of the invention.
  • Ta is a high atomic mass element, which means that complex compositional adjustments have to be made to maintain the density of the alloy within reasonable limits (preferably 8.3 kg/dm 3 or less). Further, Ta is expensive and it has not been possible to establish clearly that it has any beneficial role in crack resistance. Finally, its strengthening effect on the gamma-prime phase does not appear to be greater than that of the elements Ti and Nb. It has even been shown that the strength of the alloys of the invention is at least equivalent to that of alloys containing Ta.
  • the amounts of Al, Ti, and Nb, as an atomic percentage in the alloys of the invention are such that the ratio between the sum of the amounts of Ti and Nb and the amount of Al is 0.9 or more and 1.1 or less, i.e. 0.9 ([(Ti+Nb)/Al] (1.1.
  • Ni3Ti eta phase The Ti and Nb atoms substituting for Al in the gamma-prime phase Ni3Al base strengthen it by mechanisms analogous to those of solid solution hardening. Said hardening is greater as the ratio [(Ti+Nb)/Al] rises.
  • concentration of Ti the coherent Ni3Ti eta phase precipitates in the form of elongate plates that have a deleterious effect on the mechanical behavior, especially on the ductility, of alloys containing it.
  • concentration of Nb must be limited, since an excessive Nb content is prejudicial to the crack propagation resistance in this type of alloy.
  • the amounts of W, Mo, Cr, and Co, as an atomic percentage are such that the sum of the amounts of W, Mo, Cr, and Co is 30% or more and 34% or less, and such that the sum of the amounts of W and Mo is 3% or more and 4.5% or less, i.e.: 30% ⁇ W+Mo+Cr+Co ⁇ 34%; and 3% ⁇ W+Mo ⁇ 4.5%.
  • the elements which essentially substitute for Ni in the gamma solid solution are Cr, Co, Mo, and W.
  • Cr is essential for oxidation and corrosion properties of the alloy, and it participates in hardening the gamma matrix by the solid solution effect.
  • Co improves the high-temperature creep strength of these alloys. Further, an increase in the concentration of Co within the stability limits of the structure of the gamma phase can reduce the solvus temperature of the gamma-prime phase and hence facilitate carrying out the partial or complete solution heat treatments thereof.
  • Mo and W greatly harden the gamma matrix by the solid solution effect.
  • those elements have high atomic masses and their substitution for Ni (in particular substitution of W for Ni) results in a substantial increase in the density of the alloy.
  • the amounts of Cr, Mo, Co, and W in the alloys of the invention must thus be carefully adjusted relative to one another in order to obtain the desired effects, in particular optimum hardening of the gamma matrix, without in any way risking causing the premature appearance of fragile intermetallic compound phases, namely sigma and mu. Said phases, when they develop in excessive quantities, can cause a significant reduction in the ductility and mechanical strength of the alloys.
  • the minor elements which are C, B, and Zr, form segregations principally at the grain boundaries, for example in the form of carbides or borides. They thus contribute to increasing the strength and ductility of alloys by modifying the chemistry of the grain boundaries, and their absence would be prejudicial.
  • an excess of those elements causes a reduction in the temperature of melting onset and causes excessive precipitation of carbides and borides, which consume the elements of the alloy and which no longer participate in hardening the alloy.
  • the concentrations of carbon, boron, and zircon are thus adjusted, in particular with non-zero minimum amounts of carbon and boron, so as to obtain good high-temperature strength and optimum ductility for alloys of the invention.
  • Hf is also present in moderate quantities, since that element improves the high-temperature inter-granular cracking resistance.
  • the invention also provides a method of fabricating a part, more particularly a turbo-machine part such as a compressor or turbine disk, wherein a blank of said part or the part itself is produced from a powder of an alloy of the invention, using a powder metallurgy technique.
  • said blank or said part undergoes recrystallization heat treatment during which the blank or part is brought either to a temperature that is below the solvus temperature of the gamma-prime phase of said alloy or to a temperature that is above the solvus temperature of the gamma-prime phase of said alloy, and lower than the melting onset temperature of said alloy, to encourage the development of a microstructure with a grain size which is adapted to the stress conditions.
  • FIG. 1 is a scanning electron microscope image showing the microstructure of alloy A, described below.
  • FIG. 2 is a scanning electron microscope image showing the microstructure of alloy C1, described below.
  • the parts produced from the alloys of the invention are preferably fabricated using powder metallurgy techniques.
  • production of a compressor or turbine disk using a powder metallurgy technique comprises the following steps:
  • the cooling rate which follows the solution treatment can control the distribution of intra-granular precipitates of gamma-prime phase.
  • One or more tempering treatments can control the size of the tertiary precipitates of gamma-prime phase and relax internal stresses which result from quenching.
  • Alloy A is alloy N18 and alloy B is sold with reference number René-88DT.
  • a partial solution treatment for the gamma-prime phase was carried out at a temperature below the solvus temperature (Tsolvus) of the gamma-prime phase (at about Tsolvus ⁇ 25° C.).
  • the rate of cooling was of the order of 100° C./minute after solution. This treatment was followed by tempering for 24 hours at 750° C. and air cooling.
  • a total gamma-prime phase solution treatment was carried out at a temperature above the gamma-prime solvus temperature (at about Tsolvus +15° C. to 20° C.).
  • the rate of cooling was of the order of 140° C./min after solution. Said treatment was followed by tempering for 8 hours at 760° C. and air cooling.
  • Tables III and IV show some results of mechanical tests carried out in tension, creep, and crack propagation respectively for alloys which received a subsolvus treatment (Table III) and a supersolvus treatment (Table IV).
  • the creep tests were carried out in air at 700° C. at an initial stress of 550 MPa (650 MPa [megapascal] for alloy C1).
  • the parameter t 0.2% is the time in hours to reach a plastic deformation of 0.2%.
  • the crack propagation tests were carried out in air and at 650° C.
  • the stress cycle was as follows: load ramp-up for 10 seconds, hold for 300 seconds at maximum load and release in 10 seconds with a load ratio (minimum load/maximum load) of 0.05.
  • the parameter V f35 is the crack propagation rate, measured at a value of delta K of 35 MPa ⁇ m 1/2 .
  • V f35 (m/cycle) A 1474 340 12.10 ⁇ 5 B 1445 610 3.10 ⁇ 5 C1 1590 3000* 2.10 ⁇ 5 C2 1635 2300 3.10 ⁇ 5 C3 1589 — — *under initial stress of 650 MPa
  • micro structural examinations were carried out on alloys A and C1 which had undergone a subsolvus treatment, to detect the appearance of topologically compact phases (i.e. fragile intermetallic compounds) after an ageing heat treatment of 500 hours at 750° C.
  • the observations were carried out by back-diffused electron scanning electron microscopy on non-attacked specimens.
  • severe ageing of 500 hours at 750° C. caused inter- and intra-granular formation of phases rich in heavy elements. These phases show up in clear contrast (white borders) at the grain boundaries in FIG. 1 .
  • These phases when formed in excessive quantities, may cause a significant reduction in the ductility and strength of the alloys.
  • alloy C1 which had undergone the same treatment of 500 hours at 750° C. showed that said phases were not formed during ageing.
  • the alloys of the invention were thus more stable than alloy A (N18) as regards the formation of fragile intermetallic compounds, which are topologically compact phases.

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  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Materials Engineering (AREA)
  • Mechanical Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Physics & Mathematics (AREA)
  • Thermal Sciences (AREA)
  • Crystallography & Structural Chemistry (AREA)
  • Powder Metallurgy (AREA)
US11/694,204 2006-03-31 2007-03-30 Nickel-based alloy Abandoned US20070227630A1 (en)

Applications Claiming Priority (2)

Application Number Priority Date Filing Date Title
FR0651145 2006-03-31
FR0651145A FR2899240B1 (fr) 2006-03-31 2006-03-31 Alliage a base de nickel

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US (1) US20070227630A1 (fr)
EP (1) EP1840232B1 (fr)
JP (1) JP5398123B2 (fr)
CA (1) CA2583140C (fr)
DE (1) DE602007001092D1 (fr)
FR (1) FR2899240B1 (fr)
RU (1) RU2433197C2 (fr)

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US20150013852A1 (en) * 2012-03-27 2015-01-15 Alstom Technology Ltd Method for manufacturing components made of single crystal (sx) or directionally solidified (ds) nickelbase superalloys
US20160326613A1 (en) * 2015-05-07 2016-11-10 General Electric Company Article and method for forming an article
US9518310B2 (en) 2009-05-29 2016-12-13 General Electric Company Superalloys and components formed thereof
US20170304900A1 (en) * 2016-04-25 2017-10-26 Thomas Strangman Methods of fabricating turbine engine components
US10138534B2 (en) 2015-01-07 2018-11-27 Rolls-Royce Plc Nickel alloy
US10266919B2 (en) 2015-07-03 2019-04-23 Rolls-Royce Plc Nickel-base superalloy
US10309229B2 (en) 2014-01-09 2019-06-04 Rolls-Royce Plc Nickel based alloy composition
US10415121B2 (en) * 2016-08-05 2019-09-17 Onesubsea Ip Uk Limited Nickel alloy compositions for aggressive environments
US10519529B2 (en) 2013-11-20 2019-12-31 Questek Innovations Llc Nickel-based alloys
US10801088B2 (en) 2015-12-09 2020-10-13 General Electric Company Nickel base super alloys and methods of making the same
US11326230B2 (en) 2017-05-22 2022-05-10 Kawasaki Jukogyo Kabushiki Kaisha High temperature component and method for producing same
CN115074557A (zh) * 2022-05-16 2022-09-20 北京科技大学 一种超高塑性低屈强比的高密度镍合金及其制备方法
CN115156472A (zh) * 2022-06-27 2022-10-11 中国航发四川燃气涡轮研究院 一种高性能镍基合金变形涡轮盘锻件的制备方法

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JP6131186B2 (ja) * 2010-07-09 2017-05-17 ゼネラル・エレクトリック・カンパニイ ニッケル基合金、その加工、及びそれから形成した構成部品
RU2477199C1 (ru) * 2011-12-14 2013-03-10 Общество с ограниченной ответственностью "КОММЕТПРОМ" (ООО "КОММЕТПРОМ" "COMMETPROM") Деталь рабочего колеса и способ ее изготовления
CN104561662A (zh) * 2014-11-17 2015-04-29 江苏环亚电热仪表有限公司 一种粉末合金及其生产工艺
RU2676121C2 (ru) * 2016-12-28 2018-12-26 Открытое акционерное общество "Всероссийский институт легких сплавов" (ОАО "ВИЛС") Порошковые жаропрочные сплавы для изготовления биметаллических изделий и составной диск, изготовленный из этих сплавов
US10577679B1 (en) 2018-12-04 2020-03-03 General Electric Company Gamma prime strengthened nickel superalloy for additive manufacturing
FR3098849B1 (fr) 2019-07-16 2022-10-14 Safran Aircraft Engines Carter amélioré de module d’aéronef
FR3104613B1 (fr) 2019-12-11 2021-12-10 Safran Superalliage a base de nickel
FR3130294B1 (fr) 2021-12-15 2025-05-16 Safran Alliage à base de nickel
FR3130293A1 (fr) 2021-12-15 2023-06-16 Safran Alliage à base de nickel comprenant du tantale
FR3130292B1 (fr) 2021-12-15 2024-06-14 Safran Alliage à base de nickel exempt de cobalt
FR3133623A1 (fr) 2022-03-17 2023-09-22 Safran Superalliage à base de nickel
CN116875844B (zh) * 2023-07-28 2024-02-09 北京钢研高纳科技股份有限公司 一种盘轴一体涡轮盘及其制备方法

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Cited By (17)

* Cited by examiner, † Cited by third party
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US9518310B2 (en) 2009-05-29 2016-12-13 General Electric Company Superalloys and components formed thereof
US9670571B2 (en) * 2012-03-27 2017-06-06 Ansaldo Energia Ip Uk Limited Method for manufacturing components made of single crystal (SX) or directionally solidified (DS) nickelbase superalloys
US20150013852A1 (en) * 2012-03-27 2015-01-15 Alstom Technology Ltd Method for manufacturing components made of single crystal (sx) or directionally solidified (ds) nickelbase superalloys
US10519529B2 (en) 2013-11-20 2019-12-31 Questek Innovations Llc Nickel-based alloys
US10309229B2 (en) 2014-01-09 2019-06-04 Rolls-Royce Plc Nickel based alloy composition
US10138534B2 (en) 2015-01-07 2018-11-27 Rolls-Royce Plc Nickel alloy
US20160326613A1 (en) * 2015-05-07 2016-11-10 General Electric Company Article and method for forming an article
US10266919B2 (en) 2015-07-03 2019-04-23 Rolls-Royce Plc Nickel-base superalloy
US10422024B2 (en) 2015-07-03 2019-09-24 Rolls-Royce Plc Nickel-base superalloy
US10801088B2 (en) 2015-12-09 2020-10-13 General Electric Company Nickel base super alloys and methods of making the same
US20170304900A1 (en) * 2016-04-25 2017-10-26 Thomas Strangman Methods of fabricating turbine engine components
US10722946B2 (en) * 2016-04-25 2020-07-28 Thomas Strangman Methods of fabricating turbine engine components
US10415121B2 (en) * 2016-08-05 2019-09-17 Onesubsea Ip Uk Limited Nickel alloy compositions for aggressive environments
US11326230B2 (en) 2017-05-22 2022-05-10 Kawasaki Jukogyo Kabushiki Kaisha High temperature component and method for producing same
US11773470B2 (en) 2017-05-22 2023-10-03 Kawasaki Jukogyo Kabushiki Kaisha High temperature component and method for producing same
CN115074557A (zh) * 2022-05-16 2022-09-20 北京科技大学 一种超高塑性低屈强比的高密度镍合金及其制备方法
CN115156472A (zh) * 2022-06-27 2022-10-11 中国航发四川燃气涡轮研究院 一种高性能镍基合金变形涡轮盘锻件的制备方法

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FR2899240A1 (fr) 2007-10-05
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DE602007001092D1 (de) 2009-06-25
JP2007277721A (ja) 2007-10-25
EP1840232A1 (fr) 2007-10-03
FR2899240B1 (fr) 2008-06-27
EP1840232B1 (fr) 2009-05-13
RU2433197C2 (ru) 2011-11-10
RU2007111861A (ru) 2008-10-10
JP5398123B2 (ja) 2014-01-29

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