• 专利附录
  • 专利说明
  • 权利要求
  • 类似技术
  • 同族专利
  • 引用文献
  • 法律状态
  • 优先权
    • 专利摘要:
    Nickel-based alloy with structural hardening, characterized in that its composition is, in weight percentages: - 18% ≤ Cr ≤ 22%, preferably 18% ≤ Cr ≤ 20%; - 18% ≤ Co ≤ 22%, preferably 19% ≤ Co ≤ 21%; - 4% ≤ Mo + W ≤ 8%, preferably 5.5% ≤ Mo + W ≤ 7.5%; - traces ≤ Zr ≤ 0.06%; - traces ≤ B ≤ 0.03%, preferably traces ≤ B ≤ 0.01%; - traces ≤ C ≤ 0.1%, preferably traces ≤ C ≤ 0.06%; - traces ≤ Fe 51%; - traces ≤ Nb ≤ 0.01%; - traces ≤ Ta 50.01%; - traces ≤ S ≤ 0.008%; - traces ≤ P ≤ 0.015%; - traces ≤ Mn ≤ 0.3%; - traces ≤ If ≤ 0,15%; - traces ≤ O ≤ 0.0025%; - traces ≤ N ≤ 0.0030%; the remainder being nickel and impurities resulting from the elaboration, the contents of Al and Ti satisfying, further conditions: - (1) Ti / Al ≤ 3; - (2) Al + 1.2 Ti ≥ 2%; - (3) (0.2 Al - 1.25) 2 - 0.5 Ti ≥ 0%; - (4) Ti + 1.5 Al ≤ 4.5%. Part made of this alloy, and its manufacturing process.
    公开号:FR3018525A1
    申请号:FR1452157
    申请日:2014-03-14
    公开日:2015-09-18
    发明作者:Coraline Crozet；Alexandre Devaux
    申请人:Aubert and Duval SA；
    IPC主号:
    专利说明:

    [0001] The invention relates to nickel-based alloys (superalloys), and more specifically those for the manufacture of parts to be used at high temperatures. Typically, this is the case of the elements of terrestrial, aeronautical and other turbines. For this type of use, a NiCo20Cr20MoTi alloy (AFNOR standard) known as "C263" is known whose composition is typically Ni, Cr (19-21%), Co (19-21%), Mo (5,6- 6.1%), Ti (1.9-2.4%), Al (<0.6%). The percentages are percentages by weight, as will be the case for all the compositions indicated thereafter. It is a structurally hardened alloy, which is provided by the presence of phase y '(Ni3Ti, Al), and which has good forgeability and weldability properties. On this last point, this is because, contrary to what is often encountered for alloys hardened by the phase y ', it is not subject to the phenomenon of cracking due to embrittlement cracking under stress. heat (in English "strain age cracking") in the welding areas. It also has good ductility in hot traction and satisfactory heat resistance. In general, its weldability / forgeability compromise is advantageous.
    [0002] However, it has the disadvantage of having a microstructural instability between 700 and 900 ° C, temperature range in which phase 1-1 can be formed at the expense of the phase y '(see reference: Zhao, Metallurgical and Materials Transactions A, 2001, vol.32A, pp1271-1282). Ductility and resilience are degraded. It is therefore not optimally suitable for uses where the parts are brought to such temperatures. Other alloys are known for such uses and do not exhibit this structural instability, but they have other disadvantages. The alloy known as INCO 617 (Ni, Cr (20-24%), Co (10-15%), Mo (8-10%), Al (0.8-1.5%), Ti (0-0.6 %)) has a good compromise weldability / forgeability, but its mechanical properties hot (especially at about 750 ° C which is a frequent use temperature for the parts to which the invention is addressed in a preferred manner) are insufficient. The alloy known as RENE 41 (Ni, Cr (18-20%), Co (10-12%), Mo (9-10.5%), Al (1.4-1.6%) , Ti (3-3.3%)), on the other hand, has good mechanical properties when hot, but its weldability / forgeability compromise is not optimal. It is the same for the alloy known as WASPALOY (Ni, Cr (18-21%), Co (12-15%), Mo (3.5-5%), Al (1.2- 1.6%), Ti (2.75-3.25%) These compromises weldability / forgeability unsatisfactory are probably due to a proportion of phase y 'too important, so there is a need for manufacturers to have alloys Ni base for use at high temperatures (typically 700-900 ° C) having both good microstructural stability at the temperatures of use, good mechanical properties at these same temperatures, and simultaneously good forgeability and good weldability allowing the manufacture said parts in the desired configurations and their integration in the devices for which they are intended.
    [0003] For this purpose, the subject of the invention is a nickel-based alloy with a structural hardening, characterized in that its composition is, in weight percentages: the remainder being nickel and impurities resulting from the preparation, the contents of Al and Ti satisfying, further conditions: - (1) Ti / A1 3; (2) Al + 1.2 Ti 2%; - (3) (0.2 Al - 1.25) 2 - 0.5 Ti k 0; - (4) Ti + 1.5 Al 4.5%. Its phase fraction y 'is preferably between 5 and 20%. The solvate temperature of its phase y 'is preferably less than or equal to -18% Cr 22%, preferably 18% Cr 20%; - 18% Co 22%, preferably 19% Co 21%; - 4% Mo + W 8%, preferably 5.5% Mo + W 7.5%; - traces - traces - traces Zr 0.06%; B 0.03%, preferably traces B 0.01%; C 0.1%, preferably traces C 0.06%; Fe traces 1%; traces Nb 0.01%; Ta traces 0.01%; traces S 0.008%; traces P 0.015%; traces Mn 0.3%; traces If 0.15%; traces O 0.0025%; traces N 0.0030%; 980 ° C.
    [0004] The subject of the invention is also a process for manufacturing a nickel-based alloy part, characterized in that an ingot having the previously defined composition is prepared and homogenized at a temperature of at least 1150 ° C. for 24 to 72 hours, it is hot worked by forging or rolling in a supersolvus temperature range, it is dissolved at a temperature of 1100 to 1200 ° C for 1 to 4 hours, it is cooled to at least 1 ° Cirnine, for example in water, is aged at a temperature of 750 to 850 ° C for 7 to 10 hours, and is cooled, for example in calm air, or in an enclosure. The invention also relates to a nickel-based alloy part, characterized in that it has been prepared according to the preceding method. This is, for example, a turbine element land or aeronautics. As will be understood, the invention is based on an optimization of the known C263 grade, which essentially passes through a judiciously chosen balance between the contents of Al and Ti. This equilibrium will drive: The stability of the phase y 'at high temperature (700-900 ° C., in particular 750 ° C.), in order to prevent it from changing into an acicular phase q (of Ni3Ti composition, thus devoid of Al); The phase fraction y 'formed at 700-900 ° C, particularly at 750 ° C; The solvus temperature of the phase y '.
    [0005] On the remainder of the composition of the alloy, the changes relative to the known C263 are small, and it has been verified that the optimizations of the contents of Al and Ti according to the invention do not lead to a modification of the advantageous properties of the alloy that are not directly related to the phase y '. The invention will be better understood with the aid of the description which follows, given with reference to the following appended figures: FIGS. 1 to 8 which show micrographs of reference samples (FIGS. 1 and 5 to 8) and according to FIG. invention (Figures 2 to 4); Figure 9 which shows the results of tests for measuring the tensile strength Rm of these samples as a function of temperature; Figure 10 which shows the results of tests for measuring the conventional yield strength Rp0.2 of these samples as a function of temperature; FIG. 11 which shows the results of tests for measuring the elongation at break A% of these samples as a function of temperature; Figure 12 which shows the results of tests to measure the necking Z% of these samples as a function of temperature; Figure 13 shows the results of 750 ° C rupture creep tests of these samples, where the breaking stress is given as a function of the Larson-Miller parameter; FIG. 14 which shows the results of resilience tests of two samples (a reference sample and a sample according to the invention), carried out after the final heat treatment of the sample and after overaging at 750 ° C. for 3000 h. representative of what might be the metal in a use for which it is intended in a preferred manner; FIGS. 15 to 18 show a sample according to the invention and reference samples being forged. A first condition for optimizing the equilibrium between Al and Ti is that phase formation 1-1 is avoided at temperatures of use of the alloy during its preferred uses, that is to say at temperatures of 700-900 ° C, typically of the order of 750 ° C. The formation of phase 1-1 is directly related to the contents of Ti and Al present in the alloy and their ratio. It is thus necessary to determine the ranges of contents in these elements which make it possible to avoid it with 700-900 ° C, considering the remainder of the composition of the alloy. Thermodynamic calculations, carried out using the THERMOCALC software commonly used by metallurgists and which was also used as a first approach for the rest of the optimization, indicated that for C263, if the ratio Ti / A1 was lower or equal to 3, the formation of phase 1-1 was avoided, regardless of the level of Al in the alloy. It is therefore necessary to respect the condition: (1) Ti / A1 3 Another condition is that to guarantee the properties of tensile strength and creep at 700-900 ° C, the atomic percentage of phase y 'at these temperatures in the alloy must be at least 5%. Below this value, there is not sufficient structural hardening. It is believed that this condition is fulfilled when the weight percentages of Al and Ti respect the relationship: (2) Al + 1, 2 Ti k 2%. As regards the forgeability properties (or hot deformability in general, for example by rolling) and weldability, we can say the following. Under the standard forging conditions at high temperature, the forging is carried out in a temperature range where there is no precipitation of phase y 'which would make the metal too hard and subject to the appearance of defects, such as cracks, during deformations. It is therefore performed at a temperature above the solvus temperature of this phase. This temperature is therefore advantageous not to be too high, for a forging is possible in industrial conditions. More precisely, the solvus temperature of the phase y 'must be as low as possible in order to avoid the precipitation of this phase during the inevitable cooling of the product during the forging. It is also necessary to take into account the fraction of phase y 'that can precipitate at high temperature. Indeed, the higher the hardening phase fraction precipitated at high temperature, the more the alloy is likely to harden during temperature variations that may occur during the forging, which can complicate the execution of the operation. This unwanted precipitation of phase y 'at this point in the product preparation also has an influence on the weldability, because of the possibility of cracking due to embrittlement under hot stress. Indeed, the larger the fraction of phase y 'precipitated in the welded zone, the greater the stresses generated by the precipitation of the phase y' in this same zone during cooling are high and promote cracking after welding. In order for the good conditions of hot formability and weldability to be simultaneously satisfied, it is therefore necessary to maintain a solvate temperature of the phase y 'of at most 980 ° C., and to limit the phase fraction y' present at 700-900 ° C at 20% (in atomic%), in particular at 750 ° C. These conditions are met if the weight contents of Ti and Al comply with both conditions: - (3) (0.2 Al - 1.25) 2- 0.5 Ti k 0%: - (4) Ti + 1.5 Al 4.5% Regarding the other elements that must or may be present, either as mandatory or optional alloying elements or as impurities to be limited, the following can be said. The preferred ranges are those where one is most assured of obtaining the cited advantages of each element without having the disadvantages. The Cr content is between 18 and 22%, preferably 18 to 20%. Cr is important to ensure resistance to corrosion and oxidation, and to establish the resistance of the alloy to the effects of the environment at high temperatures. Too high a content favors undesirable fragile phases, such as phase a, and the limit of 22% by weight is set accordingly. The content of Co is between 18 and 22%, preferably 19 to 21%. A high Co content is necessary in order to improve the forgeability of the grade by decreasing the solvus temperature of the γ phase, however, it must be limited mainly for cost reasons. The sum of the contents in Mo and W must be between 4 and 8%, preferably 5.5 to 7.5%. These two elements are substitutable for each other. The lower limit of 4% guarantees structural hardening and good creep resistance, and the upper limit of 8% prevents the formation of harmful phases. The Zr content is between traces (in other words, a lack of voluntary addition, the residual content of possible Zr resulting only from the melting of the raw materials and the elaboration, with the associated impurities) and 0.06%. .
    [0006] The content of B is between traces and 0.03%, preferably 0.003 to 0.01%. The content of C is between traces and 0.1%, preferably 0.04 to 0.06%. These last three elements form segregations at grain boundaries that contribute to heat resistance and ductility by trapping any harmful elements present, such as S. They promote creep resistance under low stress and high temperature conditions. However, if they are present in excess, they decrease the melting temperature of the segregated zones and strongly alter the forgeability. Their eventual presence must therefore be well controlled.
    [0007] It should be understood that the preferred contents of the elements just mentioned are independent of each other. In other words, an alloy which has a preferential content on one or more of them only, but not on the others, must nevertheless be considered as an advantageous variant of the invention. Concerning the elements whose contents have interest to be minimized as much as possible, one can say the following. The Fe content is limited to 1% maximum. Beyond, it may form phases harmful to the properties of the alloy. The contents of Nb and Ta are both limited to 0.01% maximum. These elements are expensive and have a strong tendency to segregate without these segregations having advantages that could offset their disadvantages (contrary to what can happen for Zr, B and C). The contents of S, P, Mn and Si must also be limited so as not to reduce the hot ductility. An excess of Si would also cause a precipitation of Laves phases during solidification, and it will be difficult to put them back in solution during subsequent heat treatments. Resilience would be degraded.
    [0008] The maximum permissible levels for these elements are therefore 0.008% for S, 0.015% for P, 0.3% for Mn, and 0.15% for Si. To guarantee good mechanical properties of the alloy, limit the content of O to 25 ppm maximum and the N content to 30 ppm maximum. For this purpose, evacuation under vacuum and also involving a process such as electroslag remelting (ESR) or vacuum arc remelting (VAR) is particularly recommended. But from these points of view, the alloys of the invention are not particularly distinguished from the usual C263 to which they are called to substitute.
    [0009] With regard to the process for preparing the parts, typically an ingot having the above composition is prepared by double or triple melting, thus involving at least one of the ESR and VAR processes, and homogenized at a temperature of at least 1150. For 24 to 72 hours, it is hot-worked by forging or rolling in a supersolvus temperature range, dissolved at a temperature of 1100 to 1200 ° C for 1 to 4 hours, cooled rapidly to room temperature. minus 1 ° C / min, for example in water, it is aged at 750 to 850 ° C for 7 to 10 hours, and is cooled, for example in calm air, or in an enclosure. Depending on the intended applications, variations can be made to this process, by not performing some of these steps or by adding others. They can be followed in particular by machining or any other operation of final dimensioning of the part. An elaboration of the part using a powder metallurgy process and resulting in a product having the required compositional properties would also be conceivable. Tests were performed on samples whose compositions are listed in Table 1. Ech. Ni% Cr% Co% Mo% W% B% C% Zr% AI% Ti% O ppm N ppm (1) (2) (3) (4) A 51,60 19,71 20,15 5,98 Traces 0.005 0.051 0.02 0.77 1.50 3.5 17 1.06 2.57 2.66 0.45 B 47.50 20.86 20.49 5.96 1.43 0.010 0.050 0.02 1.95 1 , 13 3.1 18 0.58 3.31 4.06 0.17 C 51.00 19.79 20.12 6.11 Traces 0.010 0.050 0.01 2.64 0.22 3.4 15 0.08 2.90 4 , 18 0.41 D 51.50 19.74 20.00 6.20 Traces Traces 0.052 0.01 0.42 2.24 3.1 22 5.33 3.11 2.87 0.24 E 50.40 19, 60 20.00 5.97 Traces 0.002 0.049 0.003 3.00 0.252 3 16 0.08 3.30 4.75 0.30 F 48.20 19.52 20.60 4.22 3.48 0.010 0.050 0.02 3.62 0.15 4 17 0.04 3.80 5.58 0.20 G 49.70 19.97 18.50 7.50 Traces 0.010 0.060 0.02 2.20 1.95 3.3 14 0.89 4.54 5.25 0.32 H 52.10 20.00 18.20 8.00 Traces Traces 0.060 0.01 1.10 0.48 3.2 16 0.44 1.68 2.13 0.82 Table 1 : Compositions of samples tested Samples A, B and C correspond to the invention, the other samples are reference alloys which do not comply with at least one of the conditions (1) to (4) previously defined s because of their contents of Al and Ti. Sample B corresponds to the version of the invention considered optimal, where the contents of all the elements are in the preferred ranges. The reference sample D corresponds to a conventional C263 type alloy which does not respect the relation (1). Sample E and sample F do not respect relationship (3). Sample G does not respect relationships (3) and (4). Sample H does not respect relationship (2). This shows that the respect of all relations (1) to (4) is necessary to obtain the desired results.
    [0010] The samples tested were made by VIM-VAR double melting (that is, as is conventional, by melting the raw materials in a vacuum induction furnace, followed by casting and solidification of an electrode, the latter being refined by vacuum reflow in an arc furnace), to obtain ingots of 200 kg. This method is commonly used for the manufacture of ingots for forming forged or laminated parts of high purity inclusionary and low levels of residual elements, especially gaseous. It is however not necessarily used to develop the alloys of the invention, if they are intended for the production of parts that do not have very high requirements on these points. In these cases, less complex conventional methods of preparation can be used, provided that they make it possible to reach the necessary low levels on certain residual elements, in particular by a suitable choice of raw materials. These ingots were homogenized at a temperature greater than 1150 ° C for 48 h, then forged into rods with a diameter of 80 mm between 1200 and 1050 ° C. The examples were then subjected to the following heat treatments: - Dissolving at 1140 ° C. +/- 10 ° C. for 2 hours, followed by quenching with water; - Aging at 800 ° C +/- 10 ° C for 8 h followed by cooling in air. This heat treatment is typical of the C263 alloy for its usual applications such as turbine elements. The THERMOCALC software does not foresee a phase 1-1 appearance for these samples in their treatment conditions, except for sample D. In fact, micrographs were made on portions of said samples having undergone 750 ° aging. C for 3000 h to simulate a use of the corresponding alloys at high temperature. Field effect electron micrographs are shown in Figures 1 (Sample D), 2 (Sample A), 3 (Sample B), 4 (Sample C), 5 (Sample E), 6 (Sample F) , 7 (sample G) and 8 (sample H).
    [0011] It is confirmed that only the sample D, representative of a conventional C263 alloy, contains a significant amount of acicular phase 1-1 (in needles). The other samples, in particular those of the invention A, B and C, do not have this phase whose invention aimed in particular to prevent the appearance during use at 700-900 ° C, typically 750 ° C. FIG. 9 shows the results of mechanical tensile tests on these same samples for the measurement of Rm, carried out between ambient and 800 ° C. FIG. 10 shows the measurement results of Rp0.2, FIG. 11 shows the measurement results of the elongation at break A%, and FIG. 12 shows the results of zero stress tests carried out in the same conditions. It turns out that the alloys B and C according to the invention have tensile results (Rm and Rp0,2) similar to those of the reference alloy D. The tensile results of the alloy A according to the invention are slightly degraded compared to those of alloy D, but remain satisfactory. And the hot ductility of alloy A is the best of all, which can be a benefit for some uses. The invention therefore makes it possible to optimally optimize or preserve all of these mechanical properties with respect to the reference alloy C263. Alloys E, F and G have very good results in traction, especially hot. But their loss of hot ductility is very important, which can be attributed to a poor balancing of the contents of Al and Ti. Alloy H is unsatisfactory in all respects at high temperatures. Figure 13 shows the results of breaking creep tests at 750 ° C: the breaking stress in MPa is given as a function of the Larson-Miller parameter (PLM) as is conventional to proceed.
    [0012] The alloys A, B, C according to the invention, and the reference alloys F and G have longer rupture times than that of the reference alloy D. This shows that, from this point of view too, the invention provides an improvement in the performance of the alloy D which is closest thereto. The alloy E has a short life because of its insufficient hot ductility, and the tests could not be prolonged beyond a PLM of 23.4. Alloy H is, again, very clearly unsatisfactory. FIG. 14 shows the results of resilience tests carried out on several specimens of the alloys A according to the invention and of reference D, firstly after heat treatment of dissolution and then aging as described above, secondly after over-aging of 3000 h at 750 ° C following the previous heat treatment, again to simulate the evolution of the alloy in use. The results are clear: the Kv resilience is practically unaffected by the over-aging of the sample A, whereas it drops very substantially for the sample D. This confirms that the phase 1-1 formed during a use high temperature of the conventional C263 alloy has a strong weakening effect, and that the invention overcomes this problem. Forging tests were also carried out under identical conditions (homogenization at more than 1150 ° C. for 48 hours and then forging at 1200 ° C.-1050 ° C. up to 80 mm diameter), and figures 15 to 18 show the results obtained. . The alloys A, B and C according to the invention, as well as the alloy H of reference, were forged without problems as would have been the alloy D: no crack appeared during the forging. Figure 15 shows alloy A being forged at about 1100 ° C and no crack is actually visible. Figure 16 shows the alloy E being forged at the same temperature, and slight cracks are visible. Figure 17 shows the alloy F being forged at the same temperature, and the cracks are much deeper than in the previous cases. Figure 18 shows the G alloy being forged at the same temperature, and again deep cracks are visible. The good forgeability of the alloys according to the invention is thus confirmed, and is attributed to a lower proportion of phase y 'than for the reference samples E, F and G.
    [0013] A preferred application of the invention is the manufacture of elements of terrestrial and aeronautical turbines, but it is, of course, not exclusive.
    权利要求:
    Claims (6)
    [0001]
    1. A nickel-based alloy with a structural hardening, characterized in that its remainder being nickel and impurities resulting from the production, the contents of Al and Ti satisfying, further conditions: - (1) Ti / Al 3; - (2) Al + 1.2 Ti 2%; - (3) (0.2 Al - 1.25) 2- 0.5 Ti k 0%; (4) Ti + 1.5 Al 4.5%.
    [0002]
    2. Alloy according to claim 1, characterized in that its phase fraction y 'is between 5 and 20%.
    [0003]
    3.- alloy according to one of claims 1 or 2, characterized in that the solvus temperature of its phase y 'is less than or equal to 980 ° C. 30
    [0004]
    4. A process for manufacturing a nickel-based alloy part, characterized in that an ingot having the composition according to claim 1 is prepared and homogenized at a temperature of at least 1150 ° C. for 24 to 24 hours. 72 h, hot worked by forging or rolling in a supersolvus temperature range, solution at a temperature of 1100 to 1200 ° C for 1 to 4 hours, cooled to at least 1 ° C / min, for example in water, it is aged at a temperature of 750 to 850 ° C for 7 to 10 hours, and is cooled, for example in calm air or in a chamber. composition is, in weight percentages: 18% Cr 22%, preferably 18% Cr 20%; 18% Co 22%, preferably 19% Co 21%; 4% Mo + W 8%, preferably 5.5% Mo + W 7.5%; traces Zr 0.06%; traces B 0.03%, preferably traces B 0.01%; traces C 0.1%, preferably traces C 0.06%; Fe traces 1%; traces Nb 0.01%; Ta traces 0.01%; traces S 0.008%; traces P 0.015%; traces Mn 0.3%; traces If 0.15%; traces O 0.0025%; traces N 0.0030%;
    [0005]
    5.- nickel-base alloy part, characterized in that it was prepared according to the method of claim 4.
    [0006]
    6. Part according to claim 5, characterized in that it is a turbine element land or aeronautics.
    类似技术:
    公开号 | 公开日 | 专利标题
    JP6057363B1|2017-01-11|Method for producing Ni-base superalloy
    EP2467505B1|2013-06-19|Nickel-based superalloy and articles made from said alloy
    EP3117017B1|2019-05-08|Precipitation hardening nickel alloy, part made of said alloy, and manufacturing method thereof
    EP2393951B1|2021-06-16|Method for producing a piece made from a superalloy based on nickel and corresponding piece
    CA2583140C|2015-03-17|Nickel-based alloy
    JP4261562B2|2009-04-30|Ni-Fe based forged superalloy excellent in high temperature strength and high temperature ductility, its manufacturing method, and steam turbine rotor
    JP5478601B2|2014-04-23|Ni-based forged alloy and gas turbine using the same
    JP5663530B2|2015-02-04|Rhenium-free single crystal superalloy for turbine blade and vane applications
    JP6252704B2|2017-12-27|Method for producing Ni-base superalloy
    EP2435591A1|2012-04-04|Near-beta titanium alloy for high strength applications and methods for manufacturing the same
    US9828657B2|2017-11-28|Ni-base super alloy
    JP4417977B2|2010-02-17|Gas turbine blade and method for manufacturing the same
    FR2935396A1|2010-03-05|PROCESS FOR THE PREPARATION OF A NICKEL - BASED SUPERALLIATION WORKPIECE AND PIECE THUS OBTAINED
    JP5395516B2|2014-01-22|Nickel-based alloy for steam turbine turbine rotor and steam turbine turbine rotor
    KR102329565B1|2021-11-22|High-temperature, damage-resistant superalloys, articles of manufacture made from superalloys, and processes for making alloys
    JP2007009279A|2007-01-18|Ni-Fe-BASE ALLOY, AND METHOD FOR MANUFACTURING Ni-Fe-BASE ALLOY MATERIAL
    JP6293682B2|2018-03-14|High strength Ni-base superalloy
    FR2928661A1|2009-09-18|NI-BASED ALLOY FOR STEAM TURBINE ROTOR AND STEAM TURBINE ROTOR
    EP1211336B1|2007-05-30|Nickel based superalloy for single crystal turbine blades of industrial turbines having a high resistance to hot corrosion
    WO2021116607A1|2021-06-17|Nickel-based superalloy
    FR2929293A1|2009-10-02|NI-BASED ALLOY FOR STEAM TURBINE ROTOR AND STEAM TURBINE ROTOR COMPRISING SAME
    FR2739870A1|1997-04-18|REFRACTORY, TOUGH AND WELDABLE ALLOY
    同族专利:
    公开号 | 公开日
    RU2016136763A3|2018-10-24|
    EP3117017B1|2019-05-08|
    CA2942604A1|2015-09-17|
    US20170002449A1|2017-01-05|
    RU2016136763A|2018-03-16|
    CN106133161A|2016-11-16|
    WO2015136094A1|2015-09-17|
    FR3018525B1|2017-05-26|
    PL3117017T3|2019-11-29|
    EP3117017A1|2017-01-18|
    JP2017514998A|2017-06-08|
    引用文献:
    公开号 | 申请日 | 公开日 | 申请人 | 专利标题
    JPS61235529A|1985-04-10|1986-10-20|Hitachi Zosen Corp|Material for roll used in continuous casting equipment|
    JP2013095949A|2011-10-31|2013-05-20|Nippon Steel & Sumitomo Metal Corp|Austenitic heat resistant alloy|
    JP3132602B2|1991-09-28|2001-02-05|大同特殊鋼株式会社|Manufacturing method of friction welding valve|
    JP4697357B1|2009-12-10|2011-06-08|住友金属工業株式会社|Austenitic heat-resistant alloy|
    EP2860272B1|2012-06-07|2017-10-04|Nippon Steel & Sumitomo Metal Corporation|Ni-BASED ALLOY|RU2685908C1|2018-09-20|2019-04-23|Федеральное государственное унитарное предприятие "Всероссийский научно-исследовательский институт авиационных материалов" |Nickel-based heat-resistant cast alloy and article made therefrom|
    CN109967674B|2019-03-22|2020-12-08|上海电气上重铸锻有限公司|Manufacturing method of high-temperature alloy forging for nuclear power steam generator|
    WO2020195049A1|2019-03-26|2020-10-01|日立金属株式会社|Method for producing ni-based super-heat-resistant alloy, and ni-based super-heat-resistant alloy|
    CN110616354B|2019-11-12|2022-03-04|湖南人文科技学院|Nickel-based high-temperature alloy powder for laser near-net shaping and preparation method and application thereof|
    法律状态:
    2016-03-21| PLFP| Fee payment|Year of fee payment: 3 |
    2017-03-17| PLFP| Fee payment|Year of fee payment: 4 |
    2018-03-23| PLFP| Fee payment|Year of fee payment: 5 |
    2020-03-19| PLFP| Fee payment|Year of fee payment: 7 |
    2021-03-23| PLFP| Fee payment|Year of fee payment: 8 |
    优先权:
    申请号 | 申请日 | 专利标题
    FR1452157A|FR3018525B1|2014-03-14|2014-03-14|NICKEL ALLOY HAVING A STRUCTURAL CURING, PIECE THEREOF AND METHOD FOR MANUFACTURING THE SAME.|FR1452157A| FR3018525B1|2014-03-14|2014-03-14|NICKEL ALLOY HAVING A STRUCTURAL CURING, PIECE THEREOF AND METHOD FOR MANUFACTURING THE SAME.|
    US15/125,836| US20170002449A1|2014-03-14|2015-03-13|Precipitation hardening nickel-base alloy, part made of said alloy, and manufacturing method thereof|
    PL15709520T| PL3117017T3|2014-03-14|2015-03-13|Precipitation hardening nickel alloy, part made of said alloy, and manufacturing method thereof|
    CN201580014356.3A| CN106133161A|2014-03-14|2015-03-13|Parts that the nickel alloy of precipitation-hardening, described alloy are made and manufacture method thereof|
    JP2016574490A| JP2017514998A|2014-03-14|2015-03-13|Precipitation hardening nickel alloy, parts made of said alloy, and method for producing the same|
    CA2942604A| CA2942604A1|2014-03-14|2015-03-13|Precipitation hardening nickel alloy, part made of said alloy, and manufacturing method thereof|
    EP15709520.9A| EP3117017B1|2014-03-14|2015-03-13|Precipitation hardening nickel alloy, part made of said alloy, and manufacturing method thereof|
    RU2016136763A| RU2016136763A3|2014-03-14|2015-03-13|
    PCT/EP2015/055346| WO2015136094A1|2014-03-14|2015-03-13|Precipitation hardening nickel alloy, part made of said alloy, and manufacturing method thereof|
    [返回顶部]