EP2298946A2 - Superalliage forgeable à haute résistance à base de nickel et procédé de fabrication - Google Patents

Superalliage forgeable à haute résistance à base de nickel et procédé de fabrication Download PDF

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Publication number
EP2298946A2
EP2298946A2 EP10176855A EP10176855A EP2298946A2 EP 2298946 A2 EP2298946 A2 EP 2298946A2 EP 10176855 A EP10176855 A EP 10176855A EP 10176855 A EP10176855 A EP 10176855A EP 2298946 A2 EP2298946 A2 EP 2298946A2
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Prior art keywords
mass
heat treatment
content
based wrought
wrought superalloy
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EP10176855A
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German (de)
English (en)
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EP2298946A3 (fr
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Hironori Kamoshida
Shinya Imano
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Mitsubishi Power Ltd
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Hitachi Ltd
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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%
    • 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
    • C22C19/055—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being at least 20% but less than 30%
    • 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 present invention relates to nickel (Ni) based superalloys, especially to Ni-based wrought superalloys having high creep strength for use mainly in high temperature environments (such as high temperature steam turbine components), and manufacturing methods thereof.
  • precipitation strengthened Ni-based superalloys used in such applications require good formability of large ingot and good hot-forgeability.
  • a precipitation strengthened Ni-based superalloy may be used in combination with a ferrite steel to form a high temperature steam turbine component.
  • the ferrite steel usually has a relatively small thermal expansion coefficient, and therefore the precipitation strengthened Ni-based superalloy used in combination with such ferrite steel also needs to have a low thermal expansion coefficient.
  • FENIX-700 of a Ni-based superalloy (produced by Hitachi, Ltd.) provides very high phase stability, very high strength, and excellent formability of large ingot.
  • M-252 of a Ni-based superalloy has a linear thermal expansion coefficient close to that of a ferrite steel, and can therefore be used combined with the ferrite steel.
  • JP-A Hei 10(1998)-317079 and WO 2009/028671 report a Ni-based superalloy which has a low thermal expansion coefficient and excellent high-temperature strength.
  • the high temperature creep properties of materials are swayed by various factors, among which the microstructure is one.
  • the grain boundaries of a material are prone to be the starting points of a creep rupture. Therefore, creep deformations can be suppressed by reducing grain boundaries.
  • rotor blades of some state-of-the-art gas turbine are made of a single crystalline material in order to eliminate potentially rupture-causing grain boundaries and to increase the creep strength.
  • the present invention it is possible to provide a high-strength Ni-based wrought superalloy which exhibits high creep strength owing to the coarsened grains of the polycrystalline body and that still provides sufficient ductility owing to the granular precipitations of appropriate amount and sizes precipitated along the grain boundaries of the polycrystalline body. Moreover, it is possible to provide a production method of the high-strength Ni-based wrought superalloy having such microstructures.
  • the present inventors intensively investigated preferable microstructures of an Ni-based wrought superalloy and a heat treatment for controlling the microstructures. Therethrough, the inventors have developed the following two-step solution heat treatment as a novel heat treatment for producing an Ni-based wrought superalloy having preferable microstructures.
  • a starting alloy is solution heat treated at a temperature higher than the solvus temperature of the ⁇ '-phase and higher than the solvus temperature of the carbides, which is a higher temperature than conventional, thereby coarsening the grains.
  • the alloy is further solution heat treated at a temperature higher than the solvus temperature of the ⁇ '-phase but lower than the solvus temperature of the carbides, that is an intermediate temperature range, thereby causing some of the carbides to precipitate in the form of granular precipitations along each grain boundary and in grains of the superalloy.
  • the inventors have developed the following single-step prolonged solution heat treatment as another novel heat treatment for producing an Ni-based wrought superalloy having preferable microstructures.
  • a starting alloy is solution heat treated at a temperature higher than the solvus temperature of the ⁇ '-phase but lower than the solvus temperature of the carbides, that is an intermediate temperature range, for a long time (specifically, for 24 hours or more, and more preferably for 48 hours or more).
  • a long time specifically, for 24 hours or more, and more preferably for 48 hours or more.
  • Figure 1 shows a schematic illustration in a cross-sectional view of an exemplary microstructure of an Ni-based wrought superalloy after a solution heat treatment according to the present invention.
  • Fig. 1 there are also showed comparative exemplary microstructures of Ni-based wrought superalloys without the solution heat treatment of the present invention (comparative Ni-based wrought superalloys A and B).
  • a polycrystalline body of the Ni-based wrought superalloy in accordance with the present invention has a microstructure comprising coarsened grains 1 and granular carbides 3 in which the granular carbides 3 are precipitated along grain boundaries 2 between the coarsened grains 1 like an intermittent chain (a discontinuous chain). Also, the granular carbides 3 are precipitated inside the coarsened grains 1.
  • a polycrystalline body of an Ni-based wrought superalloy without the solution heat treatment of the present invention has microstructures comprising relatively small grains 1' and granular carbides 3 in that the granular carbides 3 are precipitated along grain boundaries 2 between the relatively small grains 1'.
  • a polycrystalline body of another Ni-based wrought superalloy without the solution heat treatment of the invention e.g., comparative Ni-based wrought superalloy B
  • Carbon (C) constituent has an effect of preventing exaggerated grain coarsening by forming carbide.
  • carbides are liable to precipitate in a form of a stringer and ductility of the superalloy is deteriorated in a perpendicular direction to a working direction.
  • Ti titanium
  • the C content is preferably from 0.005 to 0.2 mass%, more preferably from 0.005 to 0.15 mass%, further preferably from 0.005 to 0.08 mass%, and most preferably from 0.005 to 0.05 mass%.
  • Si and Mn constituents are used as deoxidizers when melting an alloy. It is effective even if addition of these constituents is small. However, if the excessive amounts of Si and/or Mn are added, hot-workability (hot-forgeability) in production and toughness (ductility) in use of the superalloy are deteriorated. Therefore, the Si content is preferably from 0 to 1 mass%, and the Mn content is preferably from 0 to 1 mass%.
  • the each content of Si and Mn is more preferably from 0 to 0.5 mass%, further preferably from 0 to 0.1 mass%, and most preferably from 0 to 0.01 mass%.
  • Chromium (Cr) constituent is dissolved into a matrix to make a solid solution thereby improving oxidation resistance property of the superalloy.
  • the Cr content is required to be 10 mass% or more.
  • the excessive addition of Cr deteriorates plastic workability of the superalloy.
  • the Cr content is preferably from 10 to 24 mass%, more preferably from 15 to 24 mass%, further preferably from 18 to 22 mass%, and most preferably from 19 to 21 mass%.
  • Molybdenum (Mo) and tungsten (W) constituents are important elements having an effect of lowering a thermal expansion coefficient of the superalloy, so that one or more of Mo and W is indispensable.
  • the total content expressed by "[Mo content] + 0.5x[W content]” is preferably from 5 to 17 mass%. If the total content of "[Mo content] + 0.5x[W content]” is less than 5 mass%, the above effect is not obtainable and if the total content of "[Mo content] + 0.5x[W content]" exceeds 17 mass%, plastic workability of the superalloy is deteriorated.
  • the total content of "[Mo content] + 0.5x[W content]” is more preferably from 5 to 15 mass%, and further preferably from 5 to 12 mass%. Moreover, if the content ratio of W to Mo is high, a LAVES phase is prone to generate, thereby deteriorating ductility or hot-workability of the superalloy. Thus, a single addition of Mo is preferable, and its content is preferably from 8 to 12 mass%, more preferably from 9 to 11 mass%.
  • Aluminum (Al) constituent forms an intermetallic compound (Ni 3 Al), which is called ⁇ ' (gamma prime) phase, when the superalloy is subjected to an artificial aging heat treatment, thereby improving high temperature strength of the superalloy.
  • the Al content should be 1 mass% or more. However, if the Al content exceeds 2 mass%, hot-workability of the superalloy is deteriorated.
  • the Al content is preferably from 1 to 2 mass%, more preferably from 1 to 1.8 mass%.
  • Titanium (Ti) constituent forms a ⁇ ' phase (in this case, Ni 3 (Al,Ti)) as a precipitation strengthening phase together with the Al constituent.
  • the ⁇ ' phase formed with Ni, Al and Ti (Ni 3 (Al,Ti)) exhibits more excellent high temperature strength as compared with the aforementioned ⁇ ' phase of Ni 3 Al.
  • the Ti content should be 0.5 mass% or more. However, if the Ti content exceeds 3.5 mass%, the ⁇ ' phase of Ni 3 (Al,Ti) becomes unstable, resulting in that a transformation from the ⁇ ' phase to an ⁇ (eta) phase is liable to occur, thereby deteriorating high temperature strength and hot-workability.
  • the Ti content is preferably from 0.5 to 3.5 mass%, more preferably from 1 to 3 mass%, further preferably from 1.2 to 2.5 mass%, and most preferably 1.2 to 1.8 mass%.
  • a content balance between the Al and the Ti constitutions is important in the inventive superalloy.
  • the value (the ratio expressed by "[Al content]/([A1 content] + 0.56x[Ti content])" is preferably from 0.45 to 0.7, and more preferably from 0.45 to 0.60.
  • Fe iron
  • Fe has an effect of improving hot-workability of the superalloy; thereby it may be added as occasion demands. If the Fe content exceeds 10 mass%, the thermal expansion coefficient of the superalloy becomes large, and oxidation resistance is deteriorated.
  • the Fe content is preferably from 0 to 10 mass%, more preferably from 0 to 5 mass%, and further preferably from 0 to 2 mass%.
  • B and Zr constituents strengthen grain boundaries of the superalloy, thereby improving ductility of the superalloy at a high temperature. Therefore, one or more of B and Zr are added to the superalloy. However, excessive additions of B and Zr deteriorate the superalloy in hot-workability, respectively.
  • the B content is preferably from 0.002 to 0.02 mass%, and the Zr content is preferably from 0.01 to 0.2 mass%.
  • the residuals of the inventive superalloy other than the above additive elements are nickel (Ni) constituent and inevitable impurities. If the Ni content is less than 48 mass%, high temperature strength of the superalloy is insufficient. On the contrary, if the Ni content exceeds 80 mass%, ductility of the superalloy is deteriorated. Therefore, the Ni content is preferably from 48 to 80 mass%, more preferably from 50 to 75 mass%, and further preferably from 54 to 72 mass%.
  • the inventive superalloy may contain other elements than those mentioned above, so long as they are in small amounts and essentially do not adversely affect characteristics of the superalloy.
  • the following elements are such other elements: 0.05 mass% or less of P (phosphorus); 0.01 mass% or less of S (sulfur); 1 mass% or less of Nb (niobium); 20 mass% or less of Co (cobalt), 5 mass% or less of Cu (copper); 0.01 mass% or less of Mg (magnesium); 0.01 mass% or less of Ca (calcium); 0.02 mass% or less of O (oxygen); 0.05 mass% or less of N (nitrogen); and 0.1 mass% or less of REMs (rare earth metals).
  • the Nb content is more preferably 0.8 mass% or less
  • the Co content is more preferably 5 mass% or less.
  • the Ni-based wrought superalloy according to the present invention which has a polycrystalline body including a plurality of grains, has an average grain size of 72 ⁇ m or larger and 289 ⁇ m or smaller.
  • the Ni-based wrought superalloy of the invention has a uniform grain size distribution; i.e., the grains of Ni-based wrought superalloy of the invention are generally of approximately the same size.
  • the grain size within a range from 0.99 to 5.0 specified in grain size number (GS No.) of JIS G 0551 (Methods of Austenite Grain Size Test for Steel) is suitable. That is, the average grain size of 72 ⁇ m as a lower limit is GS No. of 5.0, and the average grain size of 289 ⁇ m as an upper limit is GS No. of 0.99.
  • Ni-based wrought superalloy having an average grain size of larger than 289 ⁇ m will have very poor ductility even if the superalloy has a grain boundary structure according to the invention.
  • such Ni-based wrought superalloys (having an average grain size of larger than 289 ⁇ m) will have poor ultrasonic transmittance; thus, when a large component is formed using such a superalloy, the defect detectability by ultrasonic testing will be poor.
  • an average length of the granular precipitations is preferably from 0.5 to 2.5 ⁇ m, and more preferably from 0.5 to 2.5 ⁇ m.
  • the average length of the granular precipitations means an average length of them along the grain boundaries in an arbitrary cross-sectional view of the polycrystalline body.
  • the average length of the granular precipitations along the grain boundaries means an average length of the grain boundary covered by one granular precipitation.
  • the Ni-based wrought superalloys of the invention preferably have an average length of the granular precipitations of 2.5 ⁇ m or less, and more preferably 1.5 ⁇ m or less.
  • the granular precipitations comprise mainly: at least chromium carbides; and molybdenum carbides and/or tungsten carbides, and may include titanium carbides.
  • the amount of precipitates is also important to achieve sufficient grain boundary hardening (improvement of grain boundary connectivity). That is, too small an amount of precipitates will not provide sufficient grain boundary hardening even if the average length of the granular precipitations falls within the above-described range.
  • a ratio of a total length of the granular precipitations to that of the grain boundaries (covering ratio) in an arbitrary cross-sectional view of the polycrystalline body is 50% or more.
  • the superalloy of the invention preferably have three or more granular precipitations having the above-described average length per 10 ⁇ m grain boundary length, more preferably four or more, and further preferably five or more.
  • Manufacturing method of Ni-based wrought superalloy according to the present invention is characterized most by a heat treatment process (solution heat treatment in particular).
  • a heat treatment process solution heat treatment in particular.
  • the other processes There are no particular limitations in the other processes, and it is possible to utilize conventional ones.
  • the heat treatments in the production method of the invention will be described in detail.
  • a first-step solution heat treatment is conducted at a temperature from 1100 to 1160°C.
  • This first-step solution heat treatment accelerates grain growth in a relatively short period of time. Heat treatments at temperatures higher than 1160°C will result in too high a grain growth rate, thus making it difficult to control the average grain size to a level of 289 ⁇ m or less.
  • Ni-based wrought superalloys having an average grain size larger than 289 ⁇ m after this first-step solution heat treatment cannot provide sufficient ductility even by subjecting such an superalloy to successive heat treatments including a second-step solution heat treatment described later.
  • the defect detectability by ultrasonic testing will be poor.
  • a second-step solution heat treatment is conducted at a temperature from 980 to 1080°C.
  • Thermodynamic calculations show that, of the ⁇ '-phase solvus temperatures for Ni-based wrought superalloys of the invention, the highest one is about 980°C. Therefore, this second-step solution heat treatment needs to be conducted at least 980°C or higher.
  • carbides that would otherwise precipitate along the grain boundaries will be dissolved in the matrix of superalloy.
  • the carbides are possible to precipitate thermodynamically.
  • a supersaturation for precipitating carbides is small (i.e., a nucleation frequency of carbide nucleus is small) in this temperature region, a small number of carbide particles precipitate along the grain boundaries and grow.
  • some of the carbides are precipitated inside the grains of superalloy during the second-step solution heat treatment.
  • the second-step solution heat treatment has an effect preventing a lot of carbide fine particles from being precipitated in one breath; that is, the continuous film-like carbide precipitation along each grain boundary can be suppressed during a successive artificial aging heat treatment.
  • a single-step prolonged solution heat treatment as another novel heat treatment, in which a starting alloy is heat treated within a temperature range from 980 to 1080°C for a long time (specifically, for 24 hours or more, and more preferably for 48 hours or more).
  • the temperature range from 980 to 1080°C is an intermediate temperature that is higher than the solvus temperature of the ⁇ '-phase but lower than the solvus temperature of the carbides, as mentioned above.
  • this single-step prolonged solution heat treatment there is obtained a microstructure in that: grains of the superalloy are coarsened; and granular carbides are precipitated along each grain boundary like an intermittent chain.
  • the creep ductility of the Ni-based wrought superalloy can be improved.
  • the single-step prolonged solution heat treatment requires longer period of time than the two-step solution heat treatment.
  • the single-step prolonged solution heat treatment can also prevent a lot of carbide fine particles from being precipitated along the grain boundaries because of a temperature range in that: the carbides are not dissolved into the matrix of superalloy; and the supersaturation of the carbides is small.
  • it since there is no need to change a temperature during the single-step prolonged solution heat treatment, it is able to suppress a temperature distribution inside a product being heat-treated; therefore being suitable for forming grains with more uniform sizes.
  • An artificial aging heat treatment is conducted after the solution heat treatment.
  • the artificial aging heat treatment there is no particular limitation in the artificial aging heat treatment, then conventional artificial aging heat treatment can be carried out. Through intensive investigations in terms of creep strength and ductility, it was found that the following artificial aging heat treatment was preferable.
  • a first-step aging heat treatment is conducted at a temperature from 820 to 880°C.
  • a second-step aging heat treatment is conducted at a temperature from 600 to 800°C. This two-step artificial aging heat treatment can improve both creep strength and creep ductility of the superalloy.
  • an Ni-based wrought superalloy according to the present invention has good high-temperature mechanical properties
  • the superalloy can be used desirably as a material for high-temperature components of a coal fired thermal power plant.
  • Figure 2 shows schematic diagrams of a 700°C-class coal fired thermal power plant and exemplary high-temperature components used therein. As shown in Fig.
  • a 700°C-class coal fired thermal power plant has a power generating system, in which a high-temperature steam (e.g., 700 to 750°C) heated with a boiler 10 is supplied stepwise to a high-pressure turbine 21, a middle-pressure turbine 22 and a low-pressure turbine 23 that comprise a steam turbine 20, thereby an electric generator 30 connected with a steam turbine shaft is rotated.
  • a high-temperature steam e.g., 700 to 750°C
  • the Ni-based wrought superalloy of the invention can be suitably used for high-temperature components exposed directly to the high-temperature steam and applied large mechanical stress such as boiler tubes 11, high-pressure turbine blades 24 and casing bolts 25.
  • the boiler tubes 11 are connected each other usually by welding to assemble the boiler 10. At this time, the boiler tubes 11 are required to be in a soften state in order to prevent from cracking during the welding process. Therefore, when assembling the boiler 10, it is preferable to use the boiler tubes 11 subjected the solution heat treatment according to the invention but not subjected an artificial aging heat treatment. When such the boiler tubes 11 connected by welding are used in a 700°C-class coal fired thermal power plant, the boiler tubes 11 are subjected to a substantial aging heat treatment during use. Thereby, the ⁇ '-phase precipitations are generated in a matrix of the superalloy, leading to good high-temperature mechanical properties. On the other hand, with respect to the high-pressure turbine blades 24 and the casing bolts 25, it is preferable to use the components subjected both the solution heat treatment according to the invention and the above-mentioned artificial aging heat treatment.
  • a starting superalloy was prepared by a double melt process, which consists of a vacuum induction melting (VIM) and an electroslag remelting (ESR), followed by a hot forging process.
  • VIM vacuum induction melting
  • ESR electroslag remelting
  • Table 1 Compositions of starting superalloys.
  • Ni-based wrought superalloys (Examples 1 to 18, and Comparative examples 1 to 7) were prepared by subjecting the above starting superalloys to various heat treatments, respectively.
  • Cross sectional microstructure observation and creep test were carried out to each of the Ni-based wrought superalloys prepared.
  • Table 2 shows conditions of the heat treatments (solution heat treatment and artificial aging heat treatment) and experimental results (microstructure observation and creep properties) in Examples 1 to 6;
  • Table 3 shows those in Examples 7 to 16; and
  • Table 4 shows those in Comparative examples 1 to 6. Meanwhile, the creep test was conducted under a condition of 45 kgf/mm 2 of tensile stress at 700°C through Tables 2 to 4.
  • Tables 2 to 4 show, for each superalloy, the heat treatment condition and the resulting superalloy properties: the average grain size (unit: ⁇ m) and the grain size number defined in JIS; the average length of precipitations along the grain boundaries (unit: ⁇ m); the average number of precipitations per 10- ⁇ m grain boundary length; the covering ratio; and the creep properties (time to rupture (unit: hour), creep elongation, and reduction of area).
  • the continuous film-like carbide precipitations were precipitated (see, e.g., the comparative Ni-based wrought superalloy B shown in Fig. 1 )
  • the average length of precipitations was measured and the average number of precipitations was counted.
  • Comparative example 1 was heat-treated under the conventional condition. As shown in Table 4, Comparative example 1 has very high creep ductility but a relatively shorter time to rupture (260 hours or so), i.e., Comparative example 1 was desired to have longer time to rupture (larger creep strength). Then, in Comparative examples 2 to 4, in order to increase the average grain size, temperatures of the solution heat treatments were increased. As a result, the creep strength was increased but the creep ductility was drastically degraded. Furthermore, each of Comparative examples 2 to 4 had a microstructure in which a lot of carbide fine particles were precipitated and joined together along the grain boundaries like a continuous film covering each coarsened grain.
  • Comparative example 5 had a microstructure in that granular carbides were precipitated along each grain boundary like an intermittent chain. However, the average grain size became so large that the creep ductility of the superalloy was deteriorated. Although Comparative example 6 also had a microstructure similar to Comparative example 5, the covering ratio of the granular precipitations to the grain boundaries became so small that the time to rupture (i.e., creep strength) was insufficient.
  • the invented and comparative superalloys had a slower creep strain rate than the conventional superalloy.
  • the comparative superalloys had very poor creep ductility (the creep elongation is as small as about 1/4 that of the conventional superalloy).
  • the Ni-based wrought superalloys formed according to the above-described invented method exhibited both long time to rupture (creep strength) due to the coarsened grains and necessary-and-sufficient creep ductility.
  • the time to rupture was 1.3-1.7 times that of Comparative example 1; and the creep ductility was 0.5-1.1 times that of Comparative example 1.

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EP10176855A 2009-09-15 2010-09-15 Superalliage forgeable à haute résistance à base de nickel et procédé de fabrication Withdrawn EP2298946A3 (fr)

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JP2009212632 2009-09-15
JP2010205741A JP5657964B2 (ja) 2009-09-15 2010-09-14 高強度Ｎｉ基鍛造超合金及びその製造方法

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EP2664686A1 (fr) * 2012-04-10 2013-11-20 Hitachi Ltd. Produit de tuyauterie à haute température et son procédé de production
EP2703507A1 (fr) * 2012-08-30 2014-03-05 Hitachi Ltd. Alliage à base de Ni et pale de turbine à gaz et turbine à gaz l'utilisant
RU2538054C1 (ru) * 2014-02-19 2015-01-10 Открытое акционерное общество Научно-производственное объединение "Центральный научно-исследовательский институт технологии машиностроения" ОАО НПО "ЦНИИТМАШ" Жаропрочный сплав на основе никеля для изготовления лопаток газотурбинных установок
RU2542194C1 (ru) * 2014-02-19 2015-02-20 Открытое акционерное общество Научно-производственное объединение "Центральный научно-исследовательский институт технологии машиностроения" ОАО НПО "ЦНИИТМАШ" Жаропрочный сплав на основе никеля для литья рабочих лопаток газотурбинных установок
RU2542195C1 (ru) * 2014-02-19 2015-02-20 Открытое акционерное общество Научно-производственное объединение "Центральный научно-исследовательский институт технологии машиностроения" ОАО НПО "ЦНИИТМАШ" Жаропрочный сплав на основе никеля для литья сопловых лопаток с равноосной структурой газотурбинных установок
EP2860272A4 (fr) * 2012-06-07 2016-02-24 Nippon Steel & Sumitomo Metal Corp ALLIAGE À BASE DE Ni
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EP3444366A4 (fr) * 2016-04-05 2019-08-28 Mitsubishi Heavy Industries Aero Engines, Ltd. Alliage à base de nickel, aube de turbine, et procédé pour la production d'article moulé par injection de l'alliage à base de nickel
CN110268078A (zh) * 2016-10-12 2019-09-20 Crs 控股公司 高温耐损伤超合金、由该合金制造的制品和制造该合金的方法
CN110300811A (zh) * 2017-02-17 2019-10-01 株式会社日本制钢所 Ni基合金、燃气轮机材料和制造Ni基合金的方法
CN110468304A (zh) * 2019-08-26 2019-11-19 飞而康快速制造科技有限责任公司 一种镍基合金及其制备方法
US10975700B2 (en) 2016-03-31 2021-04-13 Mitsubishi Heavy Industries, Ltd. Turbine blade designing method, turbine blade manufacturing method, and turbine blade
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US11634792B2 (en) 2017-07-28 2023-04-25 Alloyed Limited Nickel-based alloy
CN110468304A (zh) * 2019-08-26 2019-11-19 飞而康快速制造科技有限责任公司 一种镍基合金及其制备方法

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