WO2024058101A1 - Matrice pour le forgeage à chaud et son procédé de production - Google Patents

Matrice pour le forgeage à chaud et son procédé de production Download PDF

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WO2024058101A1
WO2024058101A1 PCT/JP2023/032994 JP2023032994W WO2024058101A1 WO 2024058101 A1 WO2024058101 A1 WO 2024058101A1 JP 2023032994 W JP2023032994 W JP 2023032994W WO 2024058101 A1 WO2024058101 A1 WO 2024058101A1
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hot forging
gamma prime
prime phase
alloy
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PCT/JP2023/032994
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Japanese (ja)
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翔悟 鈴木
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株式会社プロテリアル
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    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21D—WORKING OR PROCESSING OF SHEET METAL OR METAL TUBES, RODS OR PROFILES WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21D37/00—Tools as parts of machines covered by this subclass
    • B21D37/01—Selection of materials
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21J—FORGING; HAMMERING; PRESSING METAL; RIVETING; FORGE FURNACES
    • B21J13/00—Details of machines for forging, pressing, or hammering
    • B21J13/02—Dies or mountings therefor
    • 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
    • C22C30/00—Alloys containing less than 50% by weight of each constituent
    • 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
    • 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
    • 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
    • C22F1/16—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of other metals or alloys based thereon

Definitions

  • the present invention relates to a hot forging die and a method for manufacturing the same.
  • the forging material When forging products made of heat-resistant alloys, the forging material is heated to a predetermined temperature in order to reduce deformation resistance.
  • heat-resistant alloys have high strength even at high temperatures, the hot forging mold used for forging them requires high mechanical strength.
  • the temperature of the hot forging die if the temperature of the hot forging die is lower than that of the forging material, the workability of the forging material decreases due to heat removal, so products made of difficult-to-work materials such as Alloy 718 and Ti alloy, etc.
  • hot forging dies are heated to the same or close to the temperature at which the forging material is heated. Therefore, this hot forging die must have high mechanical strength at high temperatures.
  • hot forging As a hot forging die that satisfies this requirement, a Ni-based super heat-resistant alloy has been proposed that has high high-temperature compressive strength and can be used for hot forging at a die temperature of 1000°C or higher in the atmosphere (for example, (See Patent Documents 1 to 6).
  • hot forging as used in the present invention includes hot die forging in which the temperature of the hot forging die approaches the temperature of the forging material, and isothermal forging in which the temperature is brought to the same temperature as the forging material.
  • the above-mentioned Ni-based super heat-resistant alloy has a structure in which a large amount of precipitation-strengthening phases exist, contains a large amount of solid solution strengthening elements, and has high high-temperature strength, so it can be used as a mold for hot forging.
  • a hot forging die made of an alloy with higher high-temperature compressive strength An object of the present invention is to provide a mold for hot forging having high compressive strength at high temperature, which is particularly advantageous in application to molds subjected to high loads.
  • Another object of the present invention is to provide a method for manufacturing a hot forging die that is preferable for this purpose.
  • the present inventor studied the above-mentioned problems, found a hot forging mold having high high temperature compressive strength, and arrived at the present invention. That is, the present invention, in mass %, W: 7.5 to 20.0%, Mo: 0 to 5.0%, Al: 5.0 to 7.5%, Cr: 0.5 ⁇ 5.0%, Ta: 1.0 ⁇ 12.0%, C: 0.01 ⁇ 0.15%, B: 0.03% or less, S: 0.015% or less, rare earth elements, Y, Ca , 0 to 0.020% as a total of one or more elements selected from Mg, 0.5% or less as a total of one or two elements selected from Zr and Hf, Ti: 5.0 % or less, Nb: 5.0% or less, Co: 25.0% or less, and the remainder consists of a cast alloy having a composition of Ni and unavoidable impurities, and the equivalent circle diameter of the gamma prime phase in the dendrite core is 2.
  • W 10.0 to 20.0%
  • Mo 0.5 to 5.0%
  • Al 5.0 to 7.5%
  • Cr 0.5 to 4.
  • Ta 1.0 to 12.0%
  • C 0.01 to 0.15%
  • B 0.03% or less
  • S 0.015% or less
  • rare earth elements Y, Ca, Mg 0 to 0.020% as a total of one or more selected elements, 0.5% or less as a total of one or two selected from the elements Zr and Hf
  • Ti 5.0% or less
  • Nb 5.0% or less
  • Co 20.0% or less
  • the remainder is a cast alloy having a composition of Ni and unavoidable impurities
  • the equivalent circle diameter of the gamma prime phase in the dendrite core is 2.20 ⁇ m or less.
  • the area ratio of the gamma prime phase having an equivalent circle diameter of 2.00 ⁇ m or less to the gamma prime phase having an equivalent circle diameter of 2.20 ⁇ m or less in the dendrite core is 90% or more.
  • This is a hot forging mold.
  • the present invention preferably provides a hot forging die in which the area ratio of the eutectic gamma prime phase is 4.0% or less.
  • the present invention preferably provides a hot forging die having a porosity area ratio of 0.7% or less. More preferably, the hot forging die has individual porosity of 4000 ⁇ m 2 or less.
  • the present invention preferably provides a hot forging die having an average crystal grain size of 0.5 mm or more. Further, the present invention preferably provides a hot forging die having a 0.2% compression yield strength of 450 MPa or more at a test temperature of 1100° C. and a strain rate of 10 ⁇ 3 /s.
  • the present invention is a method for manufacturing a hot forging die, in which a cast alloy having the above-mentioned composition is subjected to solution heat treatment at a temperature of 1250 to 1350° C. for 0.5 hours or more. Further, the present invention is a method for manufacturing a hot forging die, preferably in which, after the solution heat treatment described above, an aging heat treatment is further performed in a temperature range of 800°C to 1150°C.
  • a hot forging mold having high high temperature compressive strength can be obtained. This makes it possible to achieve a long mold life.
  • FIG. 2 is a diagram showing optical micrographs of porosity in inventive examples and comparative examples.
  • FIG. 2 is a diagram showing optical micrographs of microstructures of inventive examples and comparative examples.
  • FIG. 2 is a diagram showing secondary electron images or backscattered electron images of microstructures of inventive examples and comparative examples.
  • FIG. 2 is a diagram showing secondary electron images or backscattered electron images of microstructures of inventive examples and comparative examples.
  • FIG. 2 is a diagram showing secondary electron images or backscattered electron images of microstructures of inventive examples and comparative examples.
  • FIG. 2 is a diagram showing secondary electron images or backscattered electron images of microstructures of inventive examples and comparative examples.
  • FIG. 3 is a diagram showing the distribution and cumulative area ratio of the gamma prime phase for each equivalent circle diameter of the present invention example and the comparative example.
  • FIG. 3 is a diagram showing the distribution and cumulative area ratio of the gamma prime phase for each equivalent circle diameter of the present invention example and the comparative example.
  • FIG. 3 is a diagram showing the distribution and cumulative area ratio of the gamma prime phase for each equivalent circle diameter of the present invention example and the comparative example.
  • FIG. 3 is a diagram showing high-temperature compressive strengths of inventive examples and comparative examples.
  • Ni-based alloy for hot forging molds which is a material for hot forging molds, will be explained.
  • the unit of chemical composition is mass %.
  • W is dissolved in solid solution in the austenite matrix ( ⁇ phase) and also in the gamma prime phase ( ⁇ ' phase) whose basic form is Ni 3 Al, which is a precipitation strengthening phase, thereby increasing the high temperature strength of the alloy. Further, W forms MC carbide together with C, which will be described later, and precipitates at grain boundaries to increase grain boundary strength, thereby increasing high-temperature strength and ductility. On the other hand, W has the effect of lowering the oxidation resistance and the effect of making it easier to precipitate harmful phases such as TCP (Topologically Close Packed) phase.
  • TCP Topicologically Close Packed
  • the content of W in the Ni-based alloy in the present invention is set to 7.5 to 20.0% from the viewpoint of increasing high-temperature strength and ductility and suppressing a decrease in oxidation resistance and precipitation of harmful phases.
  • a preferable lower limit for obtaining the effect of W more reliably is 10.0%, and a more preferable lower limit is 12.0%.
  • a preferable upper limit is 16.0%, and a more preferable upper limit is 15.0%.
  • Mo dissolves in solid solution in the austenite matrix and also in the gamma prime phase whose basic form is Ni 3 Al, which is a precipitation-strengthening phase, thereby increasing the high-temperature strength of the alloy.
  • Mo also has the effect of lowering the oxidation resistance and the effect of making harmful phases such as TCP phase more likely to precipitate.
  • Mo content is too large, M 6 C carbide is formed together with C, which will be described later, during holding at high temperatures and the amount of solid solution decreases, resulting in a decrease in high temperature strength during use. If this decrease becomes a particular problem, it is preferable not to contain it.
  • the Mo content in the Ni-based alloy in the present invention is 0 to 5.0% below the W content. shall be.
  • a preferable lower limit for obtaining the effect of Mo more reliably is 0.5%, and a more preferable lower limit is 1.5%.
  • a preferable upper limit is 4.0%, and a more preferable upper limit is 3.5%.
  • ⁇ Al> Al combines with Ni to precipitate a gamma prime phase consisting of Ni 3 Al, increases the high temperature strength of the alloy, forms an alumina film on the surface of the alloy, and has the effect of imparting oxidation resistance to the alloy.
  • the Al content is too high, gamma prime phase is excessively generated, which has the effect of lowering the toughness of the alloy.
  • the content of Al in the Ni-based alloy in the present invention is set to 5.0 to 7.5%.
  • a preferable lower limit for obtaining the effect of Al more reliably is 5.2%, and a more preferable lower limit is 5.4%.
  • a preferable upper limit of Al is 6.7%, and a more preferable upper limit is 6.5%.
  • ⁇ Cr> Cr promotes the formation of a continuous layer of alumina on or inside the alloy, and has the effect of improving the oxidation resistance of the alloy. Therefore, it is necessary to contain 0.5% or more of Cr. On the other hand, if the Cr content is too high, it also has the effect of making harmful phases such as TCP phase more likely to precipitate. In particular, when the austenite matrix or gamma prime phase contains a large amount of elements that improve the high-temperature strength of the alloy, such as W, Mo, and Ta, harmful phases tend to precipitate. From the viewpoint of improving oxidation resistance and suppressing the precipitation of harmful phases while maintaining the content of elements that improve high-temperature strength at a high level, the Cr content in the present invention is 0.5 to 5.0. %. The preferable lower limit for obtaining the effect of Cr more reliably is 1.2%. The upper limit of Cr is preferably 4.0%, more preferably 3.0%, and even more preferably 2.5%.
  • Ta solidly dissolves in the gamma prime phase composed of Ni 3 Al by replacing Al sites, thereby increasing the high-temperature strength of the alloy. Furthermore, it improves the adhesion and oxidation resistance of the oxide film formed on the alloy surface, thereby improving the oxidation resistance of the alloy. Further, Ta forms MC carbide together with C, which will be described later, and precipitates at grain boundaries to increase grain boundary strength, thereby increasing high-temperature strength and ductility. On the other hand, if the Ta content is too high, it tends to cause harmful phases such as TCP phase to precipitate, or excessively generates gamma prime phase, which reduces the toughness of the alloy.
  • the content of Ta in the present invention is set to 1.0 to 12.0%.
  • a preferable lower limit for obtaining the effect of Ta more reliably is 2.5%, and a more preferable lower limit is 3.0%.
  • a preferable upper limit of Ta is 10.0%, and a more preferable upper limit is 7.0%.
  • ⁇ C> C forms MC carbide together with W, Mo, Ta, etc., and precipitates at grain boundaries to increase grain boundary strength, thereby increasing high-temperature strength and ductility.
  • the C content is too large, the high temperature strength of the alloy may be reduced due to the formation of coarse carbides or the formation of M 6 C carbides during high temperature holding, resulting in a significant decrease in the amount of Mo solid solution.
  • the content of C in the present invention is set to 0.01 to 0.15%.
  • a preferable lower limit for obtaining the effect of C more reliably is 0.02%, and a more preferable lower limit is 0.04%.
  • a preferable upper limit of C is 0.13%, and a more preferable upper limit is 0.12%.
  • the Ni-based alloy for hot molds in the present invention can contain 0.03% or less (including 0%) of B (boron). Like carbides, B improves the strength of the grain boundaries of the alloy, increasing high-temperature strength and ductility. On the other hand, if the content of B is too large, coarse borides are formed, which also has the effect of reducing the strength of the alloy. Furthermore, the formation of low-melting-point borides causes porosity formation during solution treatment, which will be described later, and there is also a risk of reducing fatigue strength. Therefore, B may be added as necessary, especially when it is desired to improve high-temperature strength or ductility. A preferable lower limit for reliably obtaining the effect of B is 0.005%, and a more preferable lower limit is 0.01%. A preferable upper limit is 0.02%, and a more preferable upper limit is 0.015% or less.
  • ⁇ S> ⁇ Rare earth elements, Y, Ca and Mg>
  • S sulfur
  • the upper limit of S is regulated to 0.015% or less (including 0%).
  • one or more elements selected from rare earth elements, Y, Ca, and Mg that form sulfides with S can be contained in a total amount of 0.020% or less. preferable.
  • the upper limit of the total amount of rare earth elements, Y, Ca, and Mg is 0.020%.
  • S is a component that may be contained as an impurity, and may remain in a considerable amount exceeding 0%.
  • the S content is likely to be 0.0001% (1 ppm) or more, one or more selected from rare earth elements, Y, Ca, and Mg elements are added to the S content or more. It is better to do this.
  • the rare earth elements, Y, Ca, and Mg elements may be contained in an amount of 0%.
  • the rare earth elements it is preferable to use La, which also has the effect of improving oxidation resistance by a mechanism other than making S harmless. From an economic point of view, it is preferable to use Ca or Mg.
  • Mg has a smaller effect on reducing toughness and ductility than Ca, and can also be expected to have the effect of preventing cracking during casting, so when selecting one of the rare earth elements, Y, Ca, and Mg,
  • Mg is used. If a sufficient effect can be obtained by adding Mg, Ca is not added.
  • the Ni-based alloy for hot molds in the present invention may contain one or two selected from Zr and Hf in a total amount of 0.5% or less (including 0%). These elements have the effect of improving oxidation resistance by a mechanism other than making S harmless, and also form MC carbides together with the aforementioned C, precipitate at grain boundaries, and increase grain boundary strength. Increase strength and ductility. However, if these elements are added in excess, they will form low melting point compounds and cause porosity formation during solution treatment, which will be described later, and there is a risk of reducing fatigue strength. Therefore, especially when it is desired to improve oxidation resistance, high temperature strength, etc., one or two selected from Zr and Hf may be added as necessary.
  • Hf can also be expected to have the effect of preventing cracking during casting, when selecting either Zr or Hf, it is preferable to use Hf.
  • a preferable lower limit of the total is 0.01%, and a more preferable lower limit is 0.02%.
  • a preferable upper limit is 0.3%, and a more preferable upper limit is 0.2%.
  • the Ni-based alloy for hot molds in the present invention can contain Ti.
  • Ti dissolves in solid solution in the gamma prime phase consisting of Ni 3 Al by substituting Al sites, and also forms MC carbide together with the aforementioned C, thereby increasing the high temperature strength of the alloy.
  • it is a cheaper element than Ta, it is advantageous in terms of mold cost.
  • the content of Ti is too high, like Ta, it has the effect of making it easier to precipitate harmful phases such as TCP phase, and producing excessive gamma prime phase, which reduces the toughness of the alloy.
  • Ti does not have the effect of improving oxidation resistance.
  • Ti can be contained in a range of 5.0% or less (including 0%) from the viewpoint of reducing mold cost while suppressing excessive deterioration of oxidation resistance.
  • Ti has the effect of finely dispersing carbides, so when placing emphasis on ductility, it is better to select Ti instead of Nb.
  • a preferable lower limit for reliably obtaining the effect of Ti is 0.5%, more preferably 1.0%. Further, a preferable upper limit is 3.5%.
  • Ni-based alloy for hot molds in the present invention can contain Nb.
  • Nb dissolves in the gamma prime phase composed of Ni 3 Al by substituting Al sites, and also forms MC carbide together with the aforementioned C, thereby increasing the high temperature strength of the alloy.
  • it is a cheaper element than Ta, it is advantageous in terms of mold cost.
  • the content of Nb is too high, like Ta, it has the effect of making it easier to precipitate harmful phases such as TCP phase, and excessively forming gamma prime phase, which reduces the toughness of the alloy.
  • Nb does not have the effect of improving oxidation resistance.
  • Nb can be contained in a range of 5.0% or less (including 0%) from the viewpoint of reducing mold cost while suppressing excessive deterioration of oxidation resistance.
  • carbides formed by Nb are more stable at high temperatures than carbides formed by Ti, so if the stability of mechanical properties at high temperatures is particularly important, it is better to select Nb instead of Ti.
  • a preferable lower limit for reliably obtaining the effect of Nb is 0.5%, and a more preferable lower limit is 1.0%. Further, a preferable upper limit is 3.5%.
  • the Ni-based alloy for hot molds in the present invention can contain Co.
  • Co dissolves in the austenite matrix and increases the high temperature strength of the alloy. It also has the effect of suppressing the formation of coarse M 6 C carbides, and the effect of lowering the temperature of the solution treatment described later by lowering the solid solution temperature of the gamma prime phase.
  • the Co content is too large, the mold cost will increase because Co is a more expensive element than Ni. Since the solid solution strengthening ability of Co is lower than that of W or Mo, the addition of Co is not essential if both high high temperature strength and phase stability can be achieved by adjusting the content of W, Mo, etc.
  • Co can be contained in a range of 25.0% or less (including 0%) from the viewpoint of increasing high-temperature strength and suppressing an excessive increase in mold cost.
  • a preferable lower limit for reliably obtaining the effect of Co is 2.0%, more preferably 3.0%.
  • a preferable upper limit is 20.0%, more preferably 15.0%, and even more preferably 10.0%.
  • Ni-based alloy for hot molds of the present invention Elements other than the above-mentioned elements in the Ni-based alloy for hot molds of the present invention are Ni and inevitable impurities.
  • Ni is a main element constituting the gamma phase, and together with Al, Ta, Ti, Mo, and W constitutes the gamma prime phase.
  • unavoidable impurities include P, O, N, Si, Mn, Fe, Cu, etc., and if the ingot is cast in a furnace normally used for Ni-based alloys, V, Re, and Ru are also expected. Ru.
  • P, O, and N may be contained as long as they are each 0.005% or less, and Si, Mn, Fe, Cu, V, Re, and Ru are each contained in 1.0% or less, preferably 0. It may be contained as long as it is .5% or less.
  • the Ni-based alloy in the present invention can also be called a Ni-based heat-resistant alloy.
  • the aforementioned Ni-based alloy for hot molds contains Al and Ta as essential elements, and therefore mainly consists of an austenite matrix and a gamma prime phase. Furthermore, since it also contains C, carbides are also present.
  • the hot forging mold of the present invention is obtained by machining, for example, an ingot (an ingot of a near net shape mold) having the above-mentioned composition without undergoing hot plastic working. Therefore, the Ni-based alloy for hot molds constituting this has a dendrite structure.
  • the hot forging mold of the present invention is characterized by a material in which the equivalent circle diameter of the gamma prime phase in the dendrite core is 2.20 ⁇ m or less, which has particularly high high temperature strength among Ni-based alloys for hot forging. is that it is using .
  • the dendrite core referred to here is, for example, a region surrounded by carbide or eutectic gamma prime phase, as shown in FIGS.
  • the gamma prime phase in the dendrite core referred to here means that when the dendrite core is observed with the same viewing area, the size of the gamma prime phase is smaller than that of the dendrite core, and the size of the gamma prime phase is smaller than that of the gamma prime phase. It means a gamma prime phase existing in a field of view where the number of particles is 150 or more and 1500 or less (the "corner-encircled" region in the figure). Note that here, the gamma prime phase refers to a gamma prime phase that is a particle having an equivalent circle diameter of 0.05 ⁇ m or more. Gamma prime phases that are too small are difficult to identify.
  • the gamma plum phase has a circular equivalent diameter of less than 0.05 ⁇ m, when used as a hot forging mold, it will easily form a solid solution during heating before forging, which will affect the strength related to the effect of the present invention. This is to prevent
  • the equivalent circle diameter of the gamma prime phase in the dendrite core is 2.20 ⁇ m or less, even if the structure has a certain amount of eutectic gamma prime phase, or a Ni-based alloy with the same composition. Even if there is a gamma prime phase with an equivalent circle diameter exceeding 2.20 ⁇ m, the high temperature strength will be higher than that of a material in which the gamma prime phase exists in the dendrite core.
  • the solution heat treatment is necessary because the solid solution temperature of the eutectic gamma prime phase is higher than that in the dendrite core due to the segregation of Ta. Even if the structure remains as it was during casting, its area ratio is low in the entire gamma prime phase, so the strength, which is a macroscopic property, is affected by the fine structure of the gamma prime phase in the dendrite core. This is because its influence is smaller than that of the oxidation effect.
  • the equivalent circle diameter of the gamma prime phase in the dendrite core is preferably 2.00 ⁇ m or less, more preferably 1.80 ⁇ m or less, and 1. More preferably, it is 20 ⁇ m or less.
  • the equivalent circle diameter of the gamma prime phase in the dendrite core is 2.20 ⁇ m or less
  • the gamma prime phase with the equivalent circle diameter of 2.00 ⁇ m or less If the area ratio of the gamma prime phase is 90% or more of the entire gamma prime phase, the effect of improving high temperature strength due to the refinement of the gamma prime phase can be more reliably obtained.
  • the area ratio of the gamma prime phase with an equivalent circle diameter of 1.80 ⁇ m or less is preferably 90% or more, and the area ratio of the gamma prime phase with an equivalent circle diameter of 1.60 ⁇ m or less is 90% or more. It is more preferable that
  • the gamma prime phase refers to a state before being used as a hot forging mold, such as immediately after a heat treatment process during mold manufacturing. Even if, for example, rafting occurs due to continuous loading and coarsening of the gamma prime phase occurs during use as a hot forging mold, it is important to consider this from the perspective of the total mold life from the start of use. Since the condition before use is important, this is not a major problem. Furthermore, in areas where almost no load is applied during use, for example, where the equivalent stress does not exceed 100 MPa, there is no problem from the perspective of mold life, so the equivalent circle diameter of the gamma prime phase is 2.20 ⁇ m or more. Also good. How to set such a region is determined depending on the usage environment and mold life requirements. Note that the target regions are the same as those described above for other texture factors such as eutectic gamma prime phase and porosity, which will be described later.
  • ⁇ Eutectic gamma prime phase> In the hot forging die of the present invention, higher high-temperature strength can be obtained when the area ratio of the eutectic gamma prime phase in the structure (observation field) is 4.0% or less.
  • the eutectic gamma prime phase referred to here is, for example, as shown in Figure 3, when observed with an optical microscope at a viewing area of approximately 2.5 mm 2 (observation magnification: 200x), compared to the surrounding area. It refers to an irregularly shaped area that appears whiter (lighter in color).
  • etching is insufficient and it is difficult to distinguish by optical microscopy, for example, when observed with a secondary electron image or backscattered electron image with the same field of view as shown in FIGS. 4 and 5, It may also refer to an irregularly shaped area that appears black compared to its surroundings.
  • the reason why high-temperature strength increases as the area ratio of the eutectic gamma prime phase decreases is that the gamma prime phase that makes up the eutectic gamma prime phase, which was coarse and had little effect on strength, becomes finer and the gamma, which affects strength, becomes finer. This is because the volume fraction of the prime phase becomes high.
  • the structure after casting must be maintained at a higher temperature, which increases the amount of porosity (described later).
  • the eutectic gamma prime phase may have a more detrimental effect than porosity.
  • the area ratio of the eutectic gamma prime phase may be set to 4.0% or less, if necessary, based on the usage temperature of the mold and the degree of load on the mold.
  • the area ratio of the eutectic gamma prime phase is preferably 3.0% or less, more preferably 2.0% or less.
  • ⁇ Porosity> In hot forging molds, the influence of porosity on the lifespan and safety of the parts is small compared to Ni-based alloys used in engine parts and the like. This is because high-temperature compressive strength is more important than fatigue strength in hot forging dies, but the effect of porosity on high-temperature compressive strength is small or almost non-existent. However, in cases where a large amount of forged products are manufactured at relatively low temperatures and high strain rates, which require a certain degree of fatigue strength, even hot forging dies may have an adverse effect. In such a case, it is preferable that the porosity area ratio in the tissue is 0.7% or less. More preferably it is 0.5% or less, and still more preferably 0.4% or less.
  • the porosity area ratio here refers to the average porosity area ratio of the material, for example, when porosity is observed at about four locations with a viewing area of about 2.5 mm 2 and the average is taken. Furthermore, if porosity, which is the largest among porosity and therefore has the greatest potential to reduce fatigue strength, exists on the surface, fatigue strength may decrease more than expected from the area ratio. Therefore, it is preferable that the size of each porosity is 4000 ⁇ m 2 or less. More preferably, it is 3500 ⁇ m 2 or less.
  • the average crystal grain size in the structure is 0.5 mm or more.
  • a component made of fine crystal grains may have a longer lifespan.
  • coarse grains are preferable from the viewpoint of suppressing creep deformation due to grain boundary diffusion.
  • the average crystal grain size is preferably 1.5 mm or more, more preferably 3.0 mm or more, and still more preferably 5.0 mm or more.
  • the upper limit of the average crystal grain size is not particularly limited, but realistically, it is approximately 20 mm. Further, from the same viewpoint, it is preferable that the crystal grains have a columnar shape parallel to the stress direction rather than an equiaxed shape.
  • the hot forging die has a high high temperature compressive strength such that the 0.2% compressive yield strength at a test temperature of 1100°C and a strain rate of 10 -3 /s is 450 MPa or more.
  • a forging mold can be obtained.
  • the hot forging die material of the present invention can be obtained by casting. Although it is possible to produce the material by sintering alloy powder, this is disadvantageous in terms of manufacturing costs and particle size control.
  • the casting method is not particularly limited, but vacuum casting is preferred from the viewpoint of controlling the amount of elements.
  • As a method for manufacturing a hot forging die having the above-mentioned structure it is possible to increase the cooling rate after casting, but it is possible to increase the cooling rate after casting.
  • the step of cooling after performing the above solution heat treatment is most preferable from the viewpoint of manufacturing cost and process stability.
  • the heat treatment temperature is set to 1250° C. or higher and the holding time is set to 0.5 hours or longer.
  • the heat treatment temperature is set to 1250° C. or higher and the holding time is set to 0.5 hours or longer.
  • the heat treatment temperature is too high and the holding time is too long, the effect will not only be saturated but also the treatment cost will be high, so it is usually sufficient to keep the heat treatment temperature below 1350°C or within 10 hours.
  • the cooling rate of the solution heat treatment is desirably as fast as possible within a range that does not apply excessive thermal stress to the material.
  • the atmosphere during the solution heat treatment is not particularly limited, but it is preferably inert or vacuum, and if it is in the air, it is preferable to apply an antioxidant to the surface of the material. Then, by appropriately performing machining or the like before, during, or after the above-described solution heat treatment, it is possible to obtain the hot forging mold of the present invention made of a cast alloy having a desired shape.
  • the aging heat treatment may be performed in one step or in multiple steps. If the temperature used as a hot forging die and the aging heat treatment temperature are approximately the same, the soaking process of the hot forging die before use is made longer than usual, and the aging heat treatment process May also serve as By applying an antioxidant to at least one of the molding surface or side surface of the hot forging die before use, the oxidation resistance of the hot forging die can be further improved.
  • a hot forging die having a high high temperature compressive strength such that the 0.2% compressive yield strength at a test temperature of 1100°C and a strain rate of 10 -3 /s is 450 MPa or more.
  • a forging mold can be obtained.
  • a typical process for manufacturing a forged product using a hot forging die using the Ni-based alloy for hot die of the present invention will be described.
  • the forging material is heated to a predetermined forging temperature.
  • the forging temperature varies depending on the material, so adjust the temperature accordingly.
  • the hot forging mold of the present invention has characteristics that enable isothermal forging and hot die forging even in high-temperature atmospheric environments, so it is suitable for use with Ni-based super heat-resistant alloys, Ti alloys, etc., which are known as difficult-to-work materials.
  • Suitable for hot forging. Typical forging temperatures range from 1000 to 1150°C.
  • the forged material heated in the first step is hot forged using a preheated hot forging die (second step).
  • the second step of hot forging is preferably die forging.
  • the hot forging mold of the present invention is capable of hot forging in the atmosphere at a high temperature of 1000°C or higher by adjusting the Cr content, etc. Due to its compressive strength, long mold life can be achieved.
  • Ingots of Ni-based alloys for hot molds of Alloy 1, Alloy 2, and Alloy 3 shown in Table 1 were manufactured by vacuum melting. The unit is mass%. In melting, various raw materials whose weights were adjusted to have the desired composition were made into liquids at 1500 to 1600°C, and then cast into ceramic molds preheated to 800 to 900°C. After casting, the alloy and mold were allowed to cool down to room temperature, and after cooling, the alloy and mold were separated. The weight of the ingot is about 10 kg, and the approximate shape of the part without the riser part is a cube with each side of 100 mm. In addition, P and O contained in the following ingot were each 0.005% or less.
  • the center of the ingot from which the test piece material to be described later was cut consists of equiaxed grains with an average size of 0.5 mm or more.
  • Figure 1 shows the macrostructures of Alloy 1 in (a), Alloy 2 in (b), and Alloy 3 in (c), the position of the specimen for macrostructure observation in the ingot (the position indicated by the solid line in the above figure). and the area where the average grain size was measured (the area surrounded by the dotted line in the figure below).
  • the average grain size of Alloy 1 is 7.0 mm
  • Alloy 2 is 5.1 mm
  • Alloy 3 is 5.5 mm.
  • the definition of average grain size is obtained by dividing the area of the measurement region by the total number of grains within it (those located on the boundary line of the region are counted as 0.5, and those located outside are counted as 1). It is the equivalent circle diameter calculated from the average grain cross-sectional area.
  • the etching solution used in preparing the test pieces was ferric chloride.
  • this material is subjected to heat treatment to be described later, but the average crystal grain size is not reduced by this heat treatment.
  • the present invention simulates the structure of a hot forging die by cutting out a 10 x 10 x 8 mm rectangular parallelepiped from near the center of the ingot and subjecting some of the rectangular parallelepiped to solution heat treatment for the purpose of observing porosity.
  • a test piece of the example and a test piece of the comparative example were prepared.
  • Invention Example No. 1 was a test piece in which Alloy 1 was subjected to solution heat treatment at 1300°C for 2 hours and then air cooled. 1, and a test piece that was kept at 1325°C for 2 hours and then air-cooled was named Invention Example No. It was set as 2.
  • a test piece that was cut out and not subjected to heat treatment was used as Comparative Example No. It was set at 21.
  • inventive example No. 1 to No. 5 was kept at 1100°C for 4 hours and then subjected to air-cooling aging heat treatment, and a mirror-polished surface was prepared in the same manner as described above.
  • the polished surface was etched with an etching solution containing 2.6 g of dicopper. Note that since porosity is formed by initial melting at high temperatures, there was no change in porosity due to aging heat treatment.
  • Optical micrographs were taken at 200x magnification for the etched surfaces of all test pieces. The photographs were taken at two nearby locations. This optical micrograph was then binarized and analyzed for the eutectic gamma prime phase using image processing software ImageJ. Further, secondary electron images or backscattered electron images of the microstructure and the gamma prime phase in the dendrite core were taken using a scanning electron microscope (SEM) on the etched surfaces of all test pieces.
  • SEM scanning electron microscope
  • the equivalent circle diameters of the gamma prime phase in the dendrite cores of No. 5 were all 1.20 ⁇ m or less.
  • the area ratio of the gamma prime phase with an equivalent circle diameter of 2.00 ⁇ m or less to the gamma prime phase with a circle equivalent diameter of 2.20 ⁇ m or less in the dendrite core is 100%.
  • this invention example No. 1 and no. 3 and no. The area ratio of the eutectic gamma prime phase in No. 5 is 1.0% or more, but in Invention Example No. 2 and no.
  • the area ratio of the eutectic gamma prime phase in No. 4 was 1.0% or less.
  • Alloy 2 has a lower area ratio of the eutectic gamma prime phase than Alloy 1.
  • a bar with a diameter of 10 mm is cut from the center of the ingot, some of the bars are heat-treated, and then a material for collecting a test piece with a diameter of 8 mm and a height of 12 mm is cut out, and the surface is finished to the equivalent of No. 1000.
  • the heat treatment conditions were the same solution heat treatment and aging heat treatment as the test pieces for microstructure observation, and therefore, these test pieces were the same as those of the above-mentioned invention example No. 1 to No. 5 and Comparative Example No. 21 to No. This corresponds to No. 23. Note that for each of the examples of the present invention, two compression test pieces were produced.
  • the compression test conditions were a test temperature of 1100° C., a strain rate of 10 ⁇ 3 /s, and a compression ratio of 10%, and the atmosphere was air. Then, the high temperature compressive strength was evaluated by deriving the 0.2% compressive yield strength from the stress-strain curve obtained from the compression test. This compression test tests whether the mold has sufficient compressive strength even at high temperatures as a mold for hot forging under particularly high loads. If the 0.2% compressive yield strength is 450 MPa or more, it can be said that the material has sufficient strength. More preferably, it is 500 MPa or more. Table 4 shows invention example No. 1 to No. 5 and Comparative Example No. 21 to No. The test results of No. 23 compression test pieces are shown. Further, as a representative example, FIG.
  • the hot forging die of the present invention has high high temperature compressive strength and can achieve a long die life. It can be seen that the hot forging die of the present invention having the above-described characteristics is suitable for hot die forging and isothermal forging.

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Abstract

La présente invention concerne une matrice pour le forgeage à chaud qui présente une résistance à la compression élevée à haute température et peut avoir une durée de vie de matrice prolongée et un procédé de production de la matrice pour le forgeage à chaud. La matrice pour le forgeage à chaud est constituée d'un alliage de coulée ayant une composition comprenant, en termes de % en masse, de 7,5 à 20,0 % de W, de 0 à 5,0 % de Mo, de 5,0 à 7,5 % d'Al, de 0,5 à 5,0 % de Cr, de 1,0 à 12,0 % de Ta, de 0,01 à 0,15 % de C, jusqu'à 0,03 % de B, jusqu'à 0,015 % de S, de 0 à 0,020 % d'un ou plusieurs éléments, au total, choisis parmi des éléments de terres rares, Y, Ca et Mg, jusqu'à 0,5 % d'un ou les deux éléments, au total, choisis parmi Zr et Hf, jusqu'à 5,0 % de Ti, jusqu'à 5,0 % de Nb, et jusqu'à 25,0 % de Co, le complément comprenant du Ni et des impuretés inévitables. Dans la matrice pour le forgeage à chaud, le noyau dendritique contient une phase gamma prime ayant un diamètre circulaire équivalent inférieur ou égale à 2,20 µm. Le procédé de production est approprié pour la production de la matrice pour le forgeage à chaud.
PCT/JP2023/032994 2022-09-14 2023-09-11 Matrice pour le forgeage à chaud et son procédé de production WO2024058101A1 (fr)

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WO2020059846A1 (fr) * 2018-09-21 2020-03-26 日立金属株式会社 Alliage à base de ni pour matrice de formage à chaud, et matrice de forgeage à chaud obtenue à l'aide de celui-ci
CN111575535A (zh) * 2020-05-14 2020-08-25 张家港广大特材股份有限公司 一种镍基高温合金及其制备方法
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JP2020196951A (ja) * 2019-05-28 2020-12-10 日立金属株式会社 熱間金型用Ｎｉ基合金及びそれを用いた熱間鍛造用金型及びそれを用いた鍛造製品の製造方法
JP2020196047A (ja) * 2019-05-28 2020-12-10 日立金属株式会社 鍛造製品の製造方法
CN114032420A (zh) * 2021-11-10 2022-02-11 中国航发北京航空材料研究院 一种高性能铸造高温合金

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Publication number Priority date Publication date Assignee Title
JPS49128818A (fr) * 1973-04-02 1974-12-10
JP2002146460A (ja) * 2000-08-30 2002-05-22 National Institute For Materials Science ニッケル基単結晶超合金、その製造方法およびガスタービン高温部品
WO2019097663A1 (fr) * 2017-11-17 2019-05-23 三菱日立パワーシステムズ株式会社 Matériau d'alliage corroyé à base de ni et élément de turbine à température élevée utilisant ledit matériau d'alliage
WO2019106922A1 (fr) * 2017-11-29 2019-06-06 日立金属株式会社 ALLIAGE À BASE DE Ni POUR MATRICE DE FORMAGE À CHAUD, ET MATRICE DE FORGEAGE À CHAUD L'UTILISANT
WO2020059846A1 (fr) * 2018-09-21 2020-03-26 日立金属株式会社 Alliage à base de ni pour matrice de formage à chaud, et matrice de forgeage à chaud obtenue à l'aide de celui-ci
JP2020163470A (ja) * 2019-03-29 2020-10-08 日立金属株式会社 熱間鍛造材の製造方法
JP2020196951A (ja) * 2019-05-28 2020-12-10 日立金属株式会社 熱間金型用Ｎｉ基合金及びそれを用いた熱間鍛造用金型及びそれを用いた鍛造製品の製造方法
JP2020196047A (ja) * 2019-05-28 2020-12-10 日立金属株式会社 鍛造製品の製造方法
CN111575535A (zh) * 2020-05-14 2020-08-25 张家港广大特材股份有限公司 一种镍基高温合金及其制备方法
CN114032420A (zh) * 2021-11-10 2022-02-11 中国航发北京航空材料研究院 一种高性能铸造高温合金

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