EP4588590A1 - Die for hot forging and production method therefor - Google Patents

Die for hot forging and production method therefor

Info

Publication number
EP4588590A1
EP4588590A1 EP23865454.5A EP23865454A EP4588590A1 EP 4588590 A1 EP4588590 A1 EP 4588590A1 EP 23865454 A EP23865454 A EP 23865454A EP 4588590 A1 EP4588590 A1 EP 4588590A1
Authority
EP
European Patent Office
Prior art keywords
less
hot forging
alloy
gamma prime
forging die
Prior art date
Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
Pending
Application number
EP23865454.5A
Other languages
German (de)
French (fr)
Other versions
EP4588590A4 (en
Inventor
Shogo Suzuki
Current Assignee (The listed assignees may be inaccurate. Google has not performed a legal analysis and makes no representation or warranty as to the accuracy of the list.)
Proterial Ltd
Original Assignee
Proterial Ltd
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by Proterial Ltd filed Critical Proterial Ltd
Publication of EP4588590A1 publication Critical patent/EP4588590A1/en
Publication of EP4588590A4 publication Critical patent/EP4588590A4/en
Pending legal-status Critical Current

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Classifications

    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • 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
    • C22C19/057—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being less 10%
    • 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
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21K—MAKING FORGED OR PRESSED METAL PRODUCTS, e.g. HORSE-SHOES, RIVETS, BOLTS OR WHEELS
    • B21K5/00—Making tools or tool parts, e.g. pliers
    • B21K5/20—Making working faces of dies, either recessed or outstanding
    • 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
    • C22F1/002—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working by rapid cooling or quenching; cooling agents used therefor
    • 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 a hot forging die and to a method for producing the same.
  • Hot forging die must have high mechanical strength at a high temperature.
  • a hot forging die that satisfies this requirement a Ni-base heat-resistant superalloy that has superior high-temperature compressive strength and can be used for hot forging at a die temperature of 1000°C or more in the air has been proposed (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 is made close to the temperature of the forging material, and isothermal forging in which the temperature of the hot forging die is made same as the forging material.
  • the Ni-base heat-resistant superalloy has a structure in which a large amount of a precipitation strengthening phase is present, contains a large amount of a solid solution strengthening element, and has superior high-temperature strength, and thus can be used as a hot forging die.
  • a hot forging die made of an alloy having even more superior high-temperature compressive strength.
  • the present inventor has studied the above problems, made a hot forging die having superior high-temperature compressive strength, and arrived at the present invention.
  • the present invention is preferably a hot forging die including a casting alloy having a component composition consisting of, by mass%, 10.0 to 20.0% W, 0.5 to 5.0% Mo, 5.0 to 7.5% Al, 0.5 to 4.0% Cr, 1.0 to 12.0% Ta, 0.01 to 0.15% C, 0.03% or less B, 0.015% or less S, 0 to 0.020% in total of one or more selected from rare earth elements, Y, Ca, and Mg, 0.5% or less in total of one or two selected from Zr and Hf, 5.0% or less Ti, 5.0% or less Nb, 20.0% or less Co, and the balance of Ni with inevitable impurities, an equivalent circle diameter of a gamma prime phase in a dendrite core being 2.20 ⁇ m or less.
  • the present invention is preferably a hot forging die wherein an area ratio of a eutectic gamma prime phase is 4.0% or less.
  • the present invention is preferably a hot forging die wherein a pore area ratio is 0.7% or less. Then, the present invention is more preferably a hot forging die wherein a size of each pore is 4000 ⁇ m 2 or less.
  • the present invention is preferably a hot forging die wherein an average grain size is 0.5 mm or more.
  • the present invention is preferably a hot forging die wherein a 0.2% compressive proof strength at a test temperature of 1100°C and a strain rate of 10 -3 /s is 450 MPa or more.
  • the present invention is a method for producing a hot forging die, including subjecting a casting alloy having the component composition mentioned above to solution heat treatment at a temperature of 1250 to 1350°C for 0.5 hours or more.
  • the present invention is preferably a method for producing a hot forging die, further including, after the solution heat treatment, subjecting it to aging heat treatment in a temperature range of 800°C to 1150°C.
  • a hot forging die having superior high-temperature compressive strength can be obtained. This makes it possible to achieve a long die life.
  • W dissolves in an austenite matrix (y phase) and also dissolves in a gamma prime phase (y' phase) based on Ni 3 Al, which is a precipitation strengthening phase, to increase the high-temperature strength of the alloy.
  • W forms an MC carbide together with C, which will be described later, and precipitates at a grain boundary to increase grain boundary strength, thereby increasing high-temperature strength and ductility.
  • W has an effect of reducing oxidation resistance and an effect of facilitating to the precipitation of harmful phases such as TCP (Topologically Close Packed) phases.
  • the content of W in the Ni-based alloy in the present invention is 7.5 to 20.0%.
  • a preferable lower limit thereof for obtaining the effect of W more reliably is 10.0%, and a further preferable lower limit thereof is 12.0%.
  • a preferable upper limit thereof is 16.0%, and a further preferable upper limit thereof is 15.0%.
  • Mo like W, dissolves in an austenite matrix and is also dissolves in a gamma prime phase based on Ni 3 Al, which is a precipitation strengthening phase, to increase the high-temperature strength of the alloy.
  • Mo also has an effect of reducing oxidation resistance and an effect of facilitating to the precipitation of harmful phases such as TCP phases.
  • Mo forms an M 6 C carbide together with C, which will be described later, during holding at a high temperature, and the amount of solid solute decreases, and thus the high-temperature strength decreases during use. When this decrease becomes a particular problem, it is more preferable not to contain Mo.
  • 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 an effect of imparting oxidation resistance to the alloy.
  • the content of Al in the Ni-based alloy in the present invention is 5.0 to 7.5%.
  • a preferable lower limit thereof for obtaining the effect of Al more reliably is 5.2%, and a further preferable lower limit thereof is 5.4%.
  • a preferable upper limit of Al is 6.7%, and a further preferable upper limit thereof is 6.5%.
  • Ti can be contained in an amount of 5.0% or less (including 0%).
  • Ti unlike Nb, which will be described later, has the effect of finely dispersing a carbide, and thus, particularly when emphasis is placed on ductility, Ti may be selected instead of Nb.
  • a preferable lower limit thereof for obtaining the effect of Ti more reliably is 0.5%, and a further preferable lower limit thereof is 1.0%.
  • a preferable upper limit thereof is 3.5%.
  • the Ni-based alloy for a hot die in the present invention can contain Nb.
  • Nb like Ta, dissolves in a gamma prime phase consisting of Ni 3 Al in such a way as to replace an Al site and also forms an MC carbide together with C, which has been described above, and thus increases the high-temperature strength of the alloy.
  • Nb is a less expensive element than Ta, and thus, is advantageous in terms of die cost.
  • Nb, like Ta also has an effect of facilitating to the precipitation of harmful phases such as TCP phases, and an effect of reducing the toughness of the alloy because of excessive generation of a gamma prime phase.
  • Nb unlike Ta, does not have an effect of improving oxidation resistance.
  • Nb can be contained in an amount of 5.0% or less (including 0%).
  • the carbide formed by Nb is more stable at a high temperature than the carbide formed by Ti, and thus, particularly when emphasis is placed on the stability of a mechanical property at a high temperature, Nb may be selected instead of Ti.
  • a preferable lower limit thereof for obtaining the effect of Nb more reliably is 0.5%, and a further preferable lower limit thereof is 1.0%. In addition, a preferable upper limit thereof is 3.5%.
  • the Ni-based alloy for a hot die in the present invention can contain Co.
  • Co is dissolves in an austenite matrix to increase the high-temperature strength of the alloy.
  • Co also has the effect of suppressing the formation of a coarse M 6 C carbide and the effect of lowering the temperature of solution treatment, which will be described later, by lowering the solvus temperature of the gamma prime phase.
  • the content of Co is too high, the die cost increases because Co is a more expensive element than Ni.
  • the solid solution strengthening ability of Co is lower than that of W or Mo, and thus, the addition of Co is not essential when both superior high-temperature strength and phase stability can be achieved by adjusting the content of W, Mo, or the like.
  • Ni-based alloy for a hot die in the present invention Ni and inevitable impurities.
  • Ni is a main element constituting the gamma phase, and constitutes the gamma prime phase together with Al, Ta, Ti, Mo, and W.
  • examples of the inevitable impurities include P, O, N, Si, Mn, Fe, and Cu, and when an ingot is cast in a furnace usually used for a Ni-based alloy, V, Re, and Ru are also expected.
  • Ni-based alloy in the present invention can also be called a Ni-base heat-resistant alloy.
  • the Ni-based alloy for a hot die described above contains Al and Ta as essential elements, and thus, mainly consists of an austenite matrix and a gamma prime phase.
  • the Ni-based alloy for a hot die also contains C, and thus, carbides are also present.
  • a hot forging die according to the present invention is obtained by machining, for example, an ingot (an ingot of a near net shape of a die) having the composition mentioned above without undergoing hot plastic working, and thus, the Ni-based alloy for a hot die constituting the hot forging die has a dendrite structure.
  • the hot forging die according to the present invention is characterized by using a material that has an equivalent circle diameter of the gamma prime phase in the dendrite core of 2.20 ⁇ m or less and has particularly superior high-temperature strength among such Ni-based alloys for a hot die.
  • the dendrite core as used here means, for example, a region surrounded by carbides and a eutectic gamma prime phase (in a 250 ⁇ field of view, a wide region appearing relatively pale in color that is surrounded by a eutectic gamma prime phase, which is a black, irregularly shaped region, and a white, branched carbide phase), as shown in FIG. 4 and FIG. 5 , which, when observed with a field of view area of about 40 to 600 ⁇ m 2 , does not include a locally coarse gamma prime phase such as the surrounding eutectic gamma prime phase or carbides, and in which in the field of view, the size of the gamma prime phases included therein is relatively uniform.
  • a gamma prime phase that is too small is difficult to identify, and if the gamma prime phase has a circular equivalent diameter of less than 0.05 ⁇ m, in use as a die for hot forging, the gamma prime phase easily dissolves during heating before forging, and does not affect the strength related to the effect of the present invention.
  • the high-temperature strength is higher than that of a material in which a gamma prime phase having an equivalent circle diameter of more than 2.20 ⁇ m is present in the dendrite core.
  • the equivalent circle diameter of the gamma prime phase in the dendrite core is 2.20 ⁇ m or less, even if because the solvus temperature of the eutectic gamma prime phase is higher than that in the dendrite core because of the segregation of Ta, the structure has the eutectic gamma prime phase left as cast even after solution heat treatment, the area ratio thereof in all gamma prime phases is low, and thus, the strength, which is a macroscopic property, is less affected thereby than by the refinement effect of the gamma prime phase in the dendrite core.
  • the equivalent circle diameter of the gamma prime phase in the dendrite core is preferably 2.00 ⁇ m or less, is more preferably 1.80 ⁇ m or less, and is even more preferably 1.20 ⁇ m or less.
  • the equivalent circle diameter of the gamma prime phase in the dendrite core is 2.20 ⁇ m or less
  • the area ratio of gamma prime phases having an equivalent circle diameter of 2.00 ⁇ m or less in all the gamma prime phases is 90% or more, the effect of increasing high-temperature strength through the refinement of the gamma prime phase can be more reliably obtained. It is preferable that the area ratio of gamma prime phases having an equivalent circle diameter of 1.80 ⁇ m or less be 90% or more, and it is more preferable that the area ratio of gamma prime phases having an equivalent circle diameter of 1.60 ⁇ m or less be 90% or more.
  • the gamma prime phase refers to one in a state before being used as a hot forging die, such as immediately after a heat treatment process during die production. Then, even if, for example, rafting due to continuous loading occurs and coarsening of the gamma prime phase occurs during use as a hot forging die, this does not become a major problem because the state before use is important from the viewpoint of the total die life from the start of use.
  • the equivalent circle diameter of the gamma prime phase may be 2.20 ⁇ m or more.
  • the subject region is the same as the region described above.
  • the area ratio of the eutectic gamma prime phase in the structure (observation field of view) thereof is 4.0% or less.
  • the eutectic gamma prime phase as used here means an irregularly shaped region that appears whiter (paler in color) than the surroundings thereof, when observed with an optical microscope at a field of view area of about 2.5 mm 2 (observation magnification: 200 ⁇ ), for example, as shown in FIG. 3 .
  • the eutectic gamma prime phase may mean an irregularly shaped region that appears black as compared with the surroundings thereof, when observed from a secondary electron image or a backscattered electron image with the same field of view area, for example, as shown in FIG. 4 or FIG. 5 .
  • the reason high-temperature strength increases as the area ratio of the eutectic gamma prime phase decreases is that the gamma prime phase that constitutes the eutectic gamma prime phase, which is coarse and thus has little influence on strength, is refined, and the volume fraction of the gamma prime phase, which has an influence on strength, is increased.
  • the structure after casting needs to be held at a higher temperature, and thus, this increases the amount of pore (described later), and for example, when a large quantity of a forged product is produced at a relatively low temperature and a high strain rate, even a hot forging die can be adversely affected.
  • the eutectic gamma prime phase can have a more adverse influence than pore. In this case, it is preferable to reduce the area ratio of the eutectic gamma prime phase as much as possible within a range in which the amount of pores is not excessively large.
  • a pore In a hot forging die, a pore has a small influence than a Ni-based alloy used in an engine component or the like on the service life and the safety of the component. This is because more emphasis is placed on high-temperature compressive strength than fatigue strength in a hot forging die, and the influence of pores on high-temperature compressive strength is small or almost non-existent.
  • the pore area ratio in the structure is preferably 0.7% or less.
  • the pore area ratio is more preferably 0.5% or less, and is further preferably 0.4% or less.
  • the pore area ratio as used here means the average pore area ratio of a material, for example, when a pore is observed at each of about 4 locations with a field of view area of about 2.5 mm 2 and the average thereof is taken.
  • the average grain size in the structure is 0.5 mm or more.
  • a component made of a fine grain may have a longer service life.
  • a coarse grain is more preferable from the viewpoint of suppressing creep deformation due to grain boundary diffusion.
  • the average grain size is preferably 1.5 mm or more, is more preferably 3.0 mm or more, and is even more preferably 5.0 mm or more.
  • the upper limit of the average grain size is not particularly limited, and is realistically about 20 mm.
  • the grain preferably has a columnar shape parallel to the stress direction rather than an equiaxed shape.
  • the heat treatment temperature is too high or the holding time is too long, not only will the effect be saturated, but also the treatment cost will be high, and thus, usually the heat treatment temperature may be 1350°C or less and the holding time may be 10 hours or less.
  • the cooling rate of solution heat treatment is desirably as high 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, and it is preferably an inert atmosphere or is a vacuum, and in addition, if the atmosphere is atmospheric air, an antioxidant may be applied to the surface of the material.
  • the aging heat treatment may consist of one step or a plurality of steps.
  • the soaking process of the hot forging die carried out before use may be made longer than usual, and this may also serve as an aging heat treatment process.
  • a hot forging die having superior high-temperature compressive strength such that the 0.2% compressive proof strength at a test temperature of 1100°C and a strain rate of 10 -3 /s is 450 MPa or more can be obtained.
  • a forging material is heated to a predetermined forging temperature.
  • the forging temperature varies depending on the material, and thus, the temperature is adjusted appropriately.
  • a hot forging die according to the present invention has the property of enabling isothermal forging and hot die forging even in atmospheric air at a high temperature, and thus is suitable for hot forging of a Ni-base heat-resistant superalloy, a Ti alloy, or the like, which is known as a poor workability material.
  • a representative forging temperature is in the range of 1000 to 1150°C.
  • the forging material heated in the first step is hot forged by using a preheated hot forging die (second step).
  • the hot forging in the second step is preferably die forging.
  • a hot forging die according to the present invention enables hot forging in atmospheric air at a high temperature of 1000°C or more by providing a component having an adjusted Cr content or the like, and can achieve a long die life because of the high high-temperature compressive strength thereof described above.
  • Ingots of Ni-based alloys for a hot die of alloy 1, alloy 2, and alloy 3 shown in Table 1 were produced by vacuum melting. The unit is mass%.
  • various raw materials each having a weight adjusted in such a way as to have the desired composition were melted into liquid at 1500 to 1600°C, and then cast into a ceramic mold preheated to 800 to 900°C. After casting, the alloy and the mold were allowed to cool slowly to room temperature as is, and after slow cooling, the alloy and the mold were separated from each other.
  • the weight of the ingot was about 10 kg, and the approximate shape of the portion without the riser part was a cube of 100 mm on each side.
  • the amounts of P and O contained in the following ingot were each 0.005% or less.
  • the amounts of Si and Fe are each 1.0% or less.
  • the central part of the ingot from which a test piece material, to be described later, was cut out consisted of an equiaxed grain having an average size of 0.5 mm or more.
  • the macrostructures of alloy 1, alloy 2, and alloy 3 are shown in FIG. 1a, FIG. 1b, and FIG. 1c , respectively, together with the positions of test pieces for macrostructure observation in the ingots (the positions indicated by solid lines in the upper figures) and the regions in which the average grain size was measured (the regions surrounded by dotted lines in the lower figures).
  • the average grain size of alloy 1 is 7.0 mm, that of alloy 2 is 5.1 mm, and that of alloy 3 is 5.5 mm.
  • the average grain size is defined as the equivalent circle diameter calculated from the average grain cross-sectional area obtained by dividing the area of the measurement region by the total number of grains therein (one located on the boundary line of the region was counted as 0.5, and one located outside was counted as 1).
  • the etching solution used in preparing the test pieces was ferric chloride solution. In the Inventive Examples, heat treatments, which will be described later, were applied to this material, and the average grain size does not become smaller by this heat treatment.
  • FIG. 2a to FIG. 2c Representative examples of the optical micrographs of the test pieces of alloy 1, alloy 2, and alloy 3 are shown in each of FIG. 2a to FIG. 2c , FIG. 2d to FIG. 2f , and FIG. 2g and FIG. 2h .
  • analysis results of the pore area ratio of each of the test pieces of alloy 1 to alloy 3 are shown in Table 2.
  • the amounts of pores are increased by the solution heat treatment, but in Inventive Examples No. 1 and No. 3, the increase is only slight, and the average for each of the 4 locations is 0.4% or less.
  • a slight increase was found in each of Inventive Examples No. 2 and No. 4 in which the solution heat treatment temperature was high, the average for the 4 locations in No. 2 is 0.8%, and the average for the 4 locations in No. 4 is 1.6%.
  • aging heat treatment involving holding at 1100°C for 4 h followed by air-cooling was applied to each of Inventive Examples No. 1 to No. 5, then a mirror-polished surface was prepared in the same manner as described above, and then the polished surface was etched with an etching solution consisting of 50 ml of ethanol, 50 ml of 35%, in % by mass, concentrated hydrochloric acid, and 2.6 g of cupric chloride.
  • a pore is formed by initial melting at a high temperature, and thus, there was no pore change due to the aging heat treatment.
  • FIG. 3a to FIG. 3c Representative examples of the optical micrographs of the test pieces of alloy 1, alloy 2, and alloy 3 are shown in FIG. 3a to FIG. 3c , FIG. 3d to FIG. 3f , and FIG. 3g and FIG. 3h , respectively.
  • the secondary electron images or backscattered electron images of alloy 1, alloy 2, and alloy 3 are shown in FIG. 4a to FIG. 4c , FIG. 5d to FIG. 5f , and FIG. 6g and FIG. 6h .
  • the number of gamma prime phase particles in all images was 150 or more and 1500 or less.
  • Test results of the compression test pieces of Inventive Examples No. 1 to No. 5 and Comparative Examples No. 21 to No. 23 are shown in Table 4.
  • the high-temperature compressive strength of alloy 1 is shown in FIG. 10 . From Table 4, it can be seen that Inventive Examples No. 1, No. 2, and No. 4 all have a 0.2% compressive proof strength of 450 MPa or more and have excellent high-temperature compressive strength. Even in Inventive Example No. 5, the average value of the results of compression carried out twice is 450 MPa or more.
  • Inventive Example No. 3 has a 0.2% compressive proof strength of 450 MPa or less, and thus, it may not be suitable for an application at a high temperature such as 1100°C, but has higher strength than Comparative Example No. 22 and provides a longer die service life than the same, and a sufficient die service life can be achieved.
  • the pore area ratio and the eutectic gamma prime phase area ratio of Inventive Example No. 3 are lower than those of the other Inventive Examples, and thus, Inventive Example No. 3 is suitable for an application in which a high load is repeatedly applied at a relatively low temperature such as 1050°C.
  • a hot forging die according to the present invention has superior high-temperature compressive strength and can achieve a long die service life. It can be seen that a hot forging die according to the present invention having the properties described above is suitable for hot die forging and for isothermal forging.

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  • Mechanical Engineering (AREA)
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Abstract

Provided are a hot forging die that has superior high-temperature compressive strength and can achieve a long die service life, and a method for producing the same. Provided are a hot forging die comprising a casting alloy having a component composition of, in mass%, W: 7.5 to 20.0%, Mo: 0 to 5.0%, Al: 5.0 to 7.5%, Cr: 0.5 to 5.0%, Ta: 1.0 to 12.0%, C: 0.01 to 0.15%, B: 0.03% or less, S: 0.015% or less, 0 to 0.020% in total of one or two or more selected from rare earth elements, Y, Ca, and Mg, 0.5% or less in total of one or two selected from Zr and Hf, Ti: 5.0% or less, Nb: 5.0% or less, and Co: 25.0% or less, and the balance of Ni with inevitable impurities, and having an equivalent circle diameter of a gamma prime phase in a dendrite core of 2.20 µm or less, and a method for producing a hot forging die, which is preferable for the above.

Description

    TECHNICAL FIELD
  • The present invention relates to a hot forging die and to a method for producing the same.
    • BACKGROUND ART
  • In forging a product made of a heat-resistant alloy, the forging material is heated to a predetermined temperature in order to reduce deformation resistance. However, the heat-resistant alloy has high strength even at a high temperature, and thus, a hot forging die used for forging the heat-resistant alloy is required to have high mechanical strength. In addition, in hot forging, when the temperature of a hot forging die is lower than that of the forging material, the workability of the forging material is decreased by die chilling, and thus, in forging a product made of a poor workability material such as Alloy 718 or a Ti alloy, a hot forging die is used in the state of being heated to the same temperature as the temperature at which the forging material is heated or to a high temperature close thereto. Therefore, this hot forging die must have high mechanical strength at a high temperature. As a hot forging die that satisfies this requirement, a Ni-base heat-resistant superalloy that has superior high-temperature compressive strength and can be used for hot forging at a die temperature of 1000°C or more in the air has been proposed (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 is made close to the temperature of the forging material, and isothermal forging in which the temperature of the hot forging die is made same as the forging material.
    • REFERENCE DOCUMENT LIST PATENT DOCUMENTS
    • Patent Document 1: JP S62-50429 A
    • Patent Document 2: JP S60-221542 A
    • Patent Document 3: JP 2016-069702 A
    • Patent Document 4: JP 2016-069703 A
    • Patent Document 5: JP 6645627 B
    • Patent Document 6: US 4740354 B
    SUMMARY OF THE INVENTION PROBLEM TO BE SOLVED BY THE INVENTION
  • The Ni-base heat-resistant superalloy has a structure in which a large amount of a precipitation strengthening phase is present, contains a large amount of a solid solution strengthening element, and has superior high-temperature strength, and thus can be used as a hot forging die. However, for example, when a forged product having a complex shape that is to be subjected to a higher load than usual at a high forging temperature is produced in a large amount, die life is a problem even with the Ni-base heat-resistant superalloy. In order to solve this problem, it is necessary to use a hot forging die made of an alloy having even more superior high-temperature compressive strength.
  • An object of the present invention is to provide a hot forging die having superior high-temperature compressive strength that is particularly advantageous in application to a die to be exposed to a high load. Another object of the present invention is to provide a method for producing a hot forging die, which is preferable for the above.
    • MEANS FOR SOLVING THE PROBLEM
  • The present inventor has studied the above problems, made a hot forging die having superior high-temperature compressive strength, and arrived at the present invention.
  • That is, the present invention is a hot forging die including a casting alloy having a component composition consisting of, by mass%, 7.5 to 20.0% W, 0 to 5.0% Mo, 5.0 to 7.5% Al, 0.5 to 5.0% Cr, 1.0 to 12.0% Ta, 0.01 to 0.15% C, 0.03% or less B, 0.015% or less S, 0 to 0.020% in total of one or more selected from rare earth elements, Y, Ca, and Mg, 0.5% or less in total of one or two selected from Zr and Hf, 5.0% or less Ti, 5.0% or less Nb, 25.0% or less Co, and the balance of Ni with inevitable impurities, an equivalent circle diameter of a gamma prime phase in a dendrite core being 2.20 µm or less. The present invention is preferably a hot forging die including a casting alloy having a component composition consisting of, by mass%, 10.0 to 20.0% W, 0.5 to 5.0% Mo, 5.0 to 7.5% Al, 0.5 to 4.0% Cr, 1.0 to 12.0% Ta, 0.01 to 0.15% C, 0.03% or less B, 0.015% or less S, 0 to 0.020% in total of one or more selected from rare earth elements, Y, Ca, and Mg, 0.5% or less in total of one or two selected from Zr and Hf, 5.0% or less Ti, 5.0% or less Nb, 20.0% or less Co, and the balance of Ni with inevitable impurities, an equivalent circle diameter of a gamma prime phase in a dendrite core being 2.20 µm or less.
  • In addition, the present invention is preferably a hot forging die in which an area ratio of gamma prime phases having an equivalent circle diameter of 2.00 µm or less in gamma prime phases having an equivalent circle diameter of 2.20 µm or less in a dendrite core is 90% or more.
  • In addition, the present invention is preferably a hot forging die wherein an area ratio of a eutectic gamma prime phase is 4.0% or less.
  • In addition, the present invention is preferably a hot forging die wherein a pore area ratio is 0.7% or less. Then, the present invention is more preferably a hot forging die wherein a size of each pore is 4000 µm2 or less.
  • In addition, the present invention is preferably a hot forging die wherein an average grain size is 0.5 mm or more.
  • In addition, the present invention is preferably a hot forging die wherein a 0.2% compressive proof strength at a test temperature of 1100°C and a strain rate of 10-3/s is 450 MPa or more.
  • Then, the present invention is a method for producing a hot forging die, including subjecting a casting alloy having the component composition mentioned above to solution heat treatment at a temperature of 1250 to 1350°C for 0.5 hours or more.
  • In addition, the present invention is preferably a method for producing a hot forging die, further including, after the solution heat treatment, subjecting it to aging heat treatment in a temperature range of 800°C to 1150°C.
    • EFFECTS OF THE INVENTION
  • According to the present invention, a hot forging die having superior high-temperature compressive strength can be obtained. This makes it possible to achieve a long die life.
    • BRIEF DESCRIPTION OF THE DRAWINGS
    • FIG. 1 is a figure showing the macrostructures of the ingots used in Examples.
    • FIG. 2 is a figure showing optical micrographs of porosities in Inventive Examples and Comparative Examples.
    • FIG. 3 is a figure showing optical micrographs of microstructures in Inventive Examples and Comparative Examples.
    • FIG. 4 is a figure showing secondary electron images or backscattered electron images of microstructures in Inventive Examples and a Comparative Example.
    • FIG. 5 is a figure showing secondary electron images or backscattered electron images of microstructures in Inventive Examples and a Comparative Example.
    • FIG. 6 is a figure showing secondary electron images or backscattered electron images of microstructures in an Inventive Example and a Comparative Example.
    • FIG. 7 is a diagram showing the distribution of the area ratio of the gamma prime phase for each equivalent circle diameter in each of Inventive Examples and a Comparative Example and the cumulative area ratio thereof.
    • FIG. 8 is a diagram showing the distribution of the area ratio of the gamma prime phase for each equivalent circle diameter in each of Inventive Examples and a Comparative Example and the cumulative area ratio thereof.
    • FIG. 9 is a diagram showing the distribution of the area ratio of the gamma prime phase for each equivalent circle diameter in each of an Inventive Example and a Comparative Example and the cumulative area ratio thereof.
    • FIG. 10 is a diagram showing high-temperature compressive strength in Inventive Examples and a Comparative Example.
    MODE FOR CARRYING OUT THE INVENTION
  • Hereinafter, a hot forging die according to the present invention will be described in detail. First, a Ni-based alloy for a hot die, which is a material for a hot forging die, will be described. The unit of chemical composition is mass%.
    • W
  • W dissolves in an austenite matrix (y phase) and also dissolves in a gamma prime phase (y' phase) based on Ni3Al, which is a precipitation strengthening phase, to increase the high-temperature strength of the alloy. In addition, W forms an MC carbide together with C, which will be described later, and precipitates at a grain boundary to increase grain boundary strength, thereby increasing high-temperature strength and ductility. On the other hand, W has an effect of reducing oxidation resistance and an effect of facilitating to the precipitation of harmful phases such as TCP (Topologically Close Packed) phases. From the viewpoint of increasing high-temperature strength and ductility and suppressing the reduction of the oxidation resistance and the precipitation of harmful phases, the content of W in the Ni-based alloy in the present invention is 7.5 to 20.0%. A preferable lower limit thereof for obtaining the effect of W more reliably is 10.0%, and a further preferable lower limit thereof is 12.0%. In addition, a preferable upper limit thereof is 16.0%, and a further preferable upper limit thereof is 15.0%.
    • Mo
  • Mo, like W, dissolves in an austenite matrix and is also dissolves in a gamma prime phase based on Ni3Al, which is a precipitation strengthening phase, to increase the high-temperature strength of the alloy. On the other hand, Mo also has an effect of reducing oxidation resistance and an effect of facilitating to the precipitation of harmful phases such as TCP phases. In addition, if the Mo content is too high, Mo forms an M6C carbide together with C, which will be described later, during holding at a high temperature, and the amount of solid solute decreases, and thus the high-temperature strength decreases during use. When this decrease becomes a particular problem, it is more preferable not to contain Mo. From the viewpoint of increasing high-temperature strength and suppressing the reduction of oxidation resistance and the decrease of high-temperature strength during use, the content of Mo in the Ni-based alloy in the present invention is 0 to 5.0%, which is below the content of W. A preferable lower limit thereof for obtaining the effect of Mo more reliably is 0.5%, and a further preferable lower limit thereof is 1.5%. In addition, a preferable upper limit thereof is 4.0%, and a further preferable upper limit thereof is 3.5%.
    • Al
  • Al combines with Ni to precipitate a gamma prime phase consisting of Ni3Al, increases the high-temperature strength of the alloy, forms an alumina film on the surface of the alloy, and has an effect of imparting oxidation resistance to the alloy. On the other hand, if the content of Al is too high, Al has an effect of reducing the toughness of the alloy because of excessive generation of a gamma prime phase. From the viewpoint of increasing oxidation resistance and high-temperature strength and suppressing the reduction of toughness, the content of Al in the Ni-based alloy in the present invention is 5.0 to 7.5%. A preferable lower limit thereof for obtaining the effect of Al more reliably is 5.2%, and a further preferable lower limit thereof is 5.4%. In addition, a preferable upper limit of Al is 6.7%, and a further preferable upper limit thereof is 6.5%.
    • Cr
  • Cr promotes the formation of a continuous layer of alumina on the surface or inside the alloy and has an effect of improving the oxidation resistance of the alloy. Because of this, it is necessary to contain 0.5% or more of Cr. On the other hand, if the content of Cr is too high, Cr also has an effect of facilitating to the precipitation of harmful phases such as TCP phases. In particular, when the austenite matrix or the gamma prime phase contains a large amount of an element that improves the high-temperature strength of the alloy, such as W, Mo, or Ta, a harmful phase easily precipitates. From the viewpoint of improving oxidation resistance and suppressing the precipitation of harmful phases while maintaining the content of an element that improves high-temperature strength at a high level, the content of Cr in the present invention is 0.5 to 5.0%. A preferable lower limit thereof for obtaining the effect of Cr more reliably is 1.2%. A preferable upper limit of Cr is 4.0%, a further preferable upper limit thereof is 3.0%, and a more preferable upper limit thereof is 2.5%.
    • Ta
  • Ta dissolves in a gamma prime phase consisting of Ni3Al in such a way as to replace an Al site, to increase the high-temperature strength of the alloy. Furthermore, Ta increases the adhesiveness and the oxidation resistance of the oxide film formed on the alloy surface to improve the oxidation resistance of the alloy. In addition, Ta forms an MC carbide together with C, which will be described later, and increases high-temperature strength and ductility by precipitating it at a grain boundary to increase grain boundary strength. On the other hand, if the content of Ta is too high, Ta has an effect of facilitating to the precipitation of harmful phases such as TCP phases, and an effect of reducing the toughness of the alloy because of excessive generation of a gamma prime phase. From the viewpoint of increasing oxidation resistance and high-temperature strength and suppressing the precipitation of harmful phases and the reduction of toughness, the content of Ta in the present invention is 1.0 to 12.0%. A preferable lower limit thereof for obtaining the effect of Ta more reliably is 2.5%, and a further preferable lower limit thereof is 3.0%. A preferable upper limit of Ta is 10.0%, and a further preferable upper limit thereof is 7.0%.
    • C
  • C forms an MC carbide together with W, Mo, Ta, or the like and increases high-temperature strength and ductility by precipitating it at a grain boundary to increase grain boundary strength. On the other hand, if the content of C is too high, C also has an effect of reducing the high-temperature strength of the alloy by the remarkable decrease in the amount of solute Mo due to the formation of an M6C carbide during holding at a high temperature or the formation of coarse carbides. In addition, there is also the risk of reducing fatigue strength by lowering the initial melting temperature and causing pore formation during solution treatment, which will be described later. Because of these, from the viewpoint of increasing the high-temperature strength and the ductility of the alloy and suppressing a decrease in high-temperature strength or the like, the content of C in the present invention is 0.01 to 0.15%. A preferable lower limit thereof for obtaining the effect of C more reliably is 0.02%, and a further preferable lower limit thereof is 0.04%. A preferable upper limit of C is 0.13%, and a further preferable upper limit thereof is 0.12%.
    • B
  • The Ni-based alloy for a hot die in the present invention can contain 0.03% or less (including 0%) of B (boron). B, like a carbide, improves the strength of a grain boundary of the alloy to increase high-temperature strength and ductility. On the other hand, if the content of B is too high, B also has an effect of reducing the strength of the alloy because of the formation of a coarse boride. In addition, the formation of a low-melting boride causes pore formation during solution treatment, which will be described later, and thus there is also the risk of reducing fatigue strength. Because of these, B may be added as necessary, particularly, for example, when it is desired to improve high-temperature strength or ductility. A preferable lower limit thereof for obtaining the effect of B more reliably is 0.005%, and a further preferable lower limit thereof is 0.01%. A preferable upper limit thereof is 0.02%, and a more preferable upper limit thereof is 0.015% or less.
    • S Rare earth elements, Y, Ca, and Mg
  • In addition, in the Ni-based alloy for a hot die in the present invention, S (sulfur) segregates at the interface between an oxide film formed on the alloy surface and the alloy and inhibits the chemical bonding therebetween to lower the adhesiveness of the oxide film. Because of this, the upper limit of S is regulated to 0.015% or less (including 0%). Then, it is preferable to contain one or two or more selected from rare earth elements, Y, Ca, and Mg elements that form a sulfide with S in a total amount of 0.020% or less while regulating S. If these rare earth elements, Y, Ca, and Mg are added in excess, these form a low-melting compound and cause pore formation during solution treatment, which will be described later, and thus, there is a risk of reducing fatigue strength. Because of this, the upper limit of the total amount of rare earth elements, Y, Ca, and Mg is 0.020%. S is a component that can be contained as an impurity, and it may remain in a considerable amount exceeding 0%. When the content of S is likely to be 0.0001% (1 ppm) or more, one, two, or more selected from rare earth elements, Y, Ca, and Mg elements may be contained in an amount equal to or greater than the content of S. In the Ni-based alloy of the present invention, the rare earth elements, Y, Ca, and Mg elements may be contained in an amount of 0%.
  • As 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 the viewpoint of economy, it is preferable to use Ca or Mg. In addition, Mg has a smaller effect of reducing toughness and ductility than Ca and can additionally also be expected to have an effect of preventing cracking during casting, and thus, when one of the rare earth elements, Y, Ca, and Mg is selected, it is preferable to use Mg. When a sufficient effect can be obtained by adding Mg, Ca is not added. In order to reliably obtain the effect of Mg, Mg may be contained in an amount of 0.0002% or more, regardless of the presence or absence of S. Mg is contained in an amount of preferably 0.0005% or more, and further preferably 0.0010% or more.
    • Zr and Hf
  • The Ni-based alloy for a hot die in the present invention can 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 an MC carbide together with C, which has been described above, and increase high-temperature strength and ductility by precipitating it at a grain boundary to increase grain boundary strength. However, if these elements are added in excess, these also form a low-melting compound and cause pore formation during solution treatment, which will be described later, and thus, there is the risk of reducing fatigue strength. Because of this, particularly, for example, when it is desired to improve oxidation resistance, high-temperature strength, or the like, 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, and thus, when either Zr or Hf is selected, it is preferable to use Hf. A preferable lower limit of the total amount for obtaining the effect of one or two selected from Zr and Hf reliably is 0.01%, and a further preferable lower limit thereof is 0.02%. A preferable upper limit thereof is 0.3%, and a more preferable upper limit thereof is 0.2%.
    • Ti
  • The Ni-based alloy for a hot die in the present invention can contain Ti. Ti, like Ta, dissolves in a gamma prime phase consisting of Ni3Al in such a way as to replace an Al site and also forms an MC carbide together with C, which has been described above, and thus increases the high-temperature strength of the alloy. In addition, Ti is a less expensive element than Ta, and thus is advantageous in terms of die cost. On the other hand, if the content of Ti is too high, Ti, like Ta, also has an effect of facilitating to the precipitation of harmful phases such as TCP phases, and an effect of reducing the toughness of the alloy because of excessive generation of a gamma prime phase. Additionally, Ti, unlike Ta, does not have an effect of improving oxidation resistance. In the present invention, from the viewpoint of lowering die cost while suppressing excessive decrease in oxidation resistance, Ti can be contained in an amount of 5.0% or less (including 0%). Ti, unlike Nb, which will be described later, has the effect of finely dispersing a carbide, and thus, particularly when emphasis is placed on ductility, Ti may be selected instead of Nb. A preferable lower limit thereof for obtaining the effect of Ti more reliably is 0.5%, and a further preferable lower limit thereof is 1.0%. In addition, a preferable upper limit thereof is 3.5%.
    • Nb
  • The Ni-based alloy for a hot die in the present invention can contain Nb. Nb, like Ta, dissolves in a gamma prime phase consisting of Ni3Al in such a way as to replace an Al site and also forms an MC carbide together with C, which has been described above, and thus increases the high-temperature strength of the alloy. In addition, Nb is a less expensive element than Ta, and thus, is advantageous in terms of die cost. On the other hand, if the content of Nb is too high, Nb, like Ta, also has an effect of facilitating to the precipitation of harmful phases such as TCP phases, and an effect of reducing the toughness of the alloy because of excessive generation of a gamma prime phase. Additionally, Nb, unlike Ta, does not have an effect of improving oxidation resistance. In the present invention, from the viewpoint of lowering die cost while suppressing an excessive decrease in oxidation resistance, Nb can be contained in an amount of 5.0% or less (including 0%). The carbide formed by Nb is more stable at a high temperature than the carbide formed by Ti, and thus, particularly when emphasis is placed on the stability of a mechanical property at a high temperature, Nb may be selected instead of Ti. A preferable lower limit thereof for obtaining the effect of Nb more reliably is 0.5%, and a further preferable lower limit thereof is 1.0%. In addition, a preferable upper limit thereof is 3.5%.
    • Co
  • The Ni-based alloy for a hot die in the present invention can contain Co. Co is dissolves in an austenite matrix to increase the high-temperature strength of the alloy. In addition, Co also has the effect of suppressing the formation of a coarse M6C carbide and the effect of lowering the temperature of solution treatment, which will be described later, by lowering the solvus temperature of the gamma prime phase. On the other hand, if the content of Co is too high, the die cost increases because Co is a more expensive element than Ni. The solid solution strengthening ability of Co is lower than that of W or Mo, and thus, the addition of Co is not essential when both superior high-temperature strength and phase stability can be achieved by adjusting the content of W, Mo, or the like. Co may be added as necessary, for example, when an increase in cost is acceptable or when the solvus temperature of the gamma prime phase has to be lowered because of a limitation on the capacity of the heating furnace used for solution heat treatment. In the present invention, from the viewpoint of increasing high-temperature strength and suppressing an excessive increase in die cost, Co can be contained in an amount of 25.0% or less (including 0%). A preferable lower limit thereof for obtaining the effects of Co more reliably is 2.0%, and a further preferable lower limit thereof is 3.0%. In addition, a preferable upper limit thereof is 20.0%, is more preferably 15.0%, and is even more preferably 10.0%.
    • Balance
  • The elements other than the elements described above in the Ni-based alloy for a hot die in the present invention are Ni and inevitable impurities. In the Ni-based alloy for a hot die in the present invention, Ni is a main element constituting the gamma phase, and constitutes the gamma prime phase together with Al, Ta, Ti, Mo, and W. In addition, examples of the inevitable impurities include P, O, N, Si, Mn, Fe, and Cu, and when an ingot is cast in a furnace usually used for a Ni-based alloy, V, Re, and Ru are also expected. P, O, and N may be contained as long as these are each contained in an amount of 0.005% or less, and in addition, Si, Mn, Fe, Cu, V, Re, and Ru may be contained as long as these are each contained in an amount of 1.0% or less, preferably 0.5% or less. In addition, the Ni-based alloy in the present invention can also be called a Ni-base heat-resistant alloy.
    • Gamma prime phase
  • The Ni-based alloy for a hot die described above contains Al and Ta as essential elements, and thus, mainly consists of an austenite matrix and a gamma prime phase. In addition, the Ni-based alloy for a hot die also contains C, and thus, carbides are also present. A hot forging die according to the present invention is obtained by machining, for example, an ingot (an ingot of a near net shape of a die) having the composition mentioned above without undergoing hot plastic working, and thus, the Ni-based alloy for a hot die constituting the hot forging die has a dendrite structure. Then, the hot forging die according to the present invention is characterized by using a material that has an equivalent circle diameter of the gamma prime phase in the dendrite core of 2.20 µm or less and has particularly superior high-temperature strength among such Ni-based alloys for a hot die.
  • The dendrite core as used here means, for example, a region surrounded by carbides and a eutectic gamma prime phase (in a 250× field of view, a wide region appearing relatively pale in color that is surrounded by a eutectic gamma prime phase, which is a black, irregularly shaped region, and a white, branched carbide phase), as shown in FIG. 4 and FIG. 5, which, when observed with a field of view area of about 40 to 600 µm2, does not include a locally coarse gamma prime phase such as the surrounding eutectic gamma prime phase or carbides, and in which in the field of view, the size of the gamma prime phases included therein is relatively uniform. Furthermore, the gamma prime phase in the dendrite core as used here means a gamma prime phase in which when the dendrite core is observed with the same field of view area, the size of the gamma prime phase is smaller in the dendrite core, and which is present in a region captured in a field of view in which the number of gamma prime phase particles is 150 or more and 1500 or less (a "rectangular box" region in the figure). Here, the gamma prime phase is defined as a gamma prime phase that is a particle having an equivalent circle diameter of 0.05 µm or more. This is because a gamma prime phase that is too small is difficult to identify, and if the gamma prime phase has a circular equivalent diameter of less than 0.05 µm, in use as a die for hot forging, the gamma prime phase easily dissolves during heating before forging, and does not affect the strength related to the effect of the present invention.
  • Then, by setting the equivalent circle diameter of the gamma prime phase in the dendrite core to 2.20 µm or less, even if the structure has a certain amount of a eutectic gamma prime phase, or even if the Ni-based alloy has the same composition, the high-temperature strength is higher than that of a material in which a gamma prime phase having an equivalent circle diameter of more than 2.20 µm is present in the dendrite core. This is because if the equivalent circle diameter of the gamma prime phase in the dendrite core is 2.20 µm or less, even if because the solvus temperature of the eutectic gamma prime phase is higher than that in the dendrite core because of the segregation of Ta, the structure has the eutectic gamma prime phase left as cast even after solution heat treatment, the area ratio thereof in all gamma prime phases is low, and thus, the strength, which is a macroscopic property, is less affected thereby than by the refinement effect of the gamma prime phase in the dendrite core. In order to obtain the effect of increasing high-temperature strength more reliably, the equivalent circle diameter of the gamma prime phase in the dendrite core is preferably 2.00 µm or less, is more preferably 1.80 µm or less, and is even more preferably 1.20 µm or less.
  • In addition, regarding the statement "the equivalent circle diameter of the gamma prime phase in the dendrite core is 2.20 µm or less", if the area ratio of gamma prime phases having an equivalent circle diameter of 2.00 µm or less in all the gamma prime phases is 90% or more, the effect of increasing high-temperature strength through the refinement of the gamma prime phase can be more reliably obtained. It is preferable that the area ratio of gamma prime phases having an equivalent circle diameter of 1.80 µm or less be 90% or more, and it is more preferable that the area ratio of gamma prime phases having an equivalent circle diameter of 1.60 µm or less be 90% or more.
  • In order to form such a structure, it is necessary to hold the casting structure at a high temperature, and in a Ni-based alloy intended for use in an engine component or the like, the reduction in fatigue strength or ductility due to a pore (described later) formed by the initial melting of a grain boundary strengthening element such as C or an inevitable impurity during holding at a high temperature becomes a problem. However, in a hot forging die according to the present invention, emphasis is mainly placed on high-temperature compressive strength, and the requirement for fatigue strength or ductility is relatively low, and thus, even if a pore formed by holding at a high temperature or a pore during production that is unavoidably present regardless of the presence or absence of holding at a high temperature is present in a small amount, this does not become a major problem.
  • In addition, the gamma prime phase according to the requirement of the present invention refers to one in a state before being used as a hot forging die, such as immediately after a heat treatment process during die production. Then, even if, for example, rafting due to continuous loading occurs and coarsening of the gamma prime phase occurs during use as a hot forging die, this does not become a major problem because the state before use is important from the viewpoint of the total die life from the start of use. In addition, in a region in which almost no load is applied during use, for example, in which the equivalent stress does not exceed 100 MPa, no problem occurs from the viewpoint of die life, and thus, the equivalent circle diameter of the gamma prime phase may be 2.20 µm or more. How such a region should be set is determined depending on the usage environment and the die life requirement. Also, for another structure factor such as a eutectic gamma prime phase or a pore, which will be described later, the subject region is the same as the region described above.
    • Eutectic gamma prime phase
  • In a hot forging die according to the present invention, even more superior high-temperature strength can be obtained if the area ratio of the eutectic gamma prime phase in the structure (observation field of view) thereof is 4.0% or less. The eutectic gamma prime phase as used here means an irregularly shaped region that appears whiter (paler in color) than the surroundings thereof, when observed with an optical microscope at a field of view area of about 2.5 mm2 (observation magnification: 200×), for example, as shown in FIG. 3. If etching is insufficient and discrimination by observation with an optical microscope is difficult, the eutectic gamma prime phase may mean an irregularly shaped region that appears black as compared with the surroundings thereof, when observed from a secondary electron image or a backscattered electron image with the same field of view area, for example, as shown in FIG. 4 or FIG. 5. The reason high-temperature strength increases as the area ratio of the eutectic gamma prime phase decreases is that the gamma prime phase that constitutes the eutectic gamma prime phase, which is coarse and thus has little influence on strength, is refined, and the volume fraction of the gamma prime phase, which has an influence on strength, is increased. However, in order to adjust the area ratio of the eutectic gamma prime phase to be 4.0% or less, the structure after casting needs to be held at a higher temperature, and thus, this increases the amount of pore (described later), and for example, when a large quantity of a forged product is produced at a relatively low temperature and a high strain rate, even a hot forging die can be adversely affected. Besides, when a higher load is repeatedly applied, the eutectic gamma prime phase can have a more adverse influence than pore. In this case, it is preferable to reduce the area ratio of the eutectic gamma prime phase as much as possible within a range in which the amount of pores is not excessively large. Because of these, the area ratio of the eutectic gamma prime phase may be 4.0% or less, as necessary, based on the operating temperature of the die or the degree of loading on the die. The area ratio of the eutectic gamma prime phase is preferably 3.0% or less, and more preferably 2.0% or less.
    • Pores
  • In a hot forging die, a pore has a small influence than a Ni-based alloy used in an engine component or the like on the service life and the safety of the component. This is because more emphasis is placed on high-temperature compressive strength than fatigue strength in a hot forging die, and the influence of pores on high-temperature compressive strength is small or almost non-existent. However, for example, when a large quantity of a forged product is produced at a relatively low temperature and a high strain rate, which also requires a certain degree of fatigue strength, even a hot forging die can be adversely affected. In such a case, the pore area ratio in the structure is preferably 0.7% or less. The pore area ratio is more preferably 0.5% or less, and is further preferably 0.4% or less. The pore area ratio as used here means the average pore area ratio of a material, for example, when a pore is observed at each of about 4 locations with a field of view area of about 2.5 mm2 and the average thereof is taken.
  • In addition, if a pore that is the largest among pores and thus has the greatest potential to reduce fatigue strength is present on the surface, fatigue strength can reduce more than expected from the area ratio. Because of this, the size of each pore is preferably 4000 µm2 or less. The size thereof is more preferably 3500 µm2 or less.
    • Average grain size
  • In a hot forging die according to the present invention, higher creep strength can be obtained if the average grain size in the structure is 0.5 mm or more. For a component that places emphasis on fatigue strength, a component made of a fine grain may have a longer service life. However, in a hot forging die that mainly places emphasis on high-temperature compressive strength, a coarse grain is more preferable from the viewpoint of suppressing creep deformation due to grain boundary diffusion. In order to obtain the effect of improving die service life reliably, the average grain size is preferably 1.5 mm or more, is more preferably 3.0 mm or more, and is even more preferably 5.0 mm or more. The upper limit of the average grain size is not particularly limited, and is realistically about 20 mm. In addition, from the same viewpoint, the grain preferably has a columnar shape parallel to the stress direction rather than an equiaxed shape.
  • According to a hot forging die described above, the hot forging die having a superior high-temperature compressive strength preferably such that the 0.2% compressive proof strength at a test temperature of 1100°C and a strain rate of 10-3/s is 450 MPa or more can be obtained.
    • Method for producing hot forging die
  • A material for a hot forging die according the present invention can be obtained by casting. This material can also be prepared by sintering an alloy powder, but this is disadvantageous from the viewpoint of production cost and grain size control. The casting method is not particularly limited, and vacuum casting is preferable from the viewpoint of controlling the amount of an element.
  • As a method for producing a hot forging die having the structure mentioned above, although a method involving increasing the cooling rate after casting is conceivable, the process of subjecting the material after casting to solution heat treatment at a temperature of 1250 to 1350°C for 0.5 hours or more and then cooling it is most preferable in terms of production cost and process stability. If the heat treatment temperature is low or the holding time is short, there is a possibility that solutioning of the gamma prime phase will be insufficient or that soaking within the material will be insufficient. In addition, the homogenization of W, Mo, and Ta that have segregated during casting can also be expected, and thus, the heat treatment temperature is 1250°C or more and the holding time is 0.5 hours or more. On the other hand, if the heat treatment temperature is too high or the holding time is too long, not only will the effect be saturated, but also the treatment cost will be high, and thus, usually the heat treatment temperature may be 1350°C or less and the holding time may be 10 hours or less.
  • In addition, by slowing down the heating rate of solution heat treatment or holding the temperature lower than the solution heat treatment temperature for a certain period of time during the heating, a reduction in the amount of a pore due to homogenization of an element that causes initial melting can be expected. In addition, the cooling rate of solution heat treatment is desirably as high 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, and it is preferably an inert atmosphere or is a vacuum, and in addition, if the atmosphere is atmospheric air, an antioxidant may be applied to the surface of the material.
  • Then, by appropriately applying machining or the like, for example, before, during, or after the solution heat treatment mentioned above, it is possible to obtain a hot forging die according to the present invention made of a casted alloy having a desired shape.
  • After the solution heat treatment, when aging heat treatment in the temperature range of 800°C to 1150°C is carried out, high-temperature strength can be further increased by controlling the size of the gamma prime phase, etc. The aging heat treatment may consist of one step or a plurality of steps. When 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 carried out before use may be made longer than usual, and this may also serve as an aging heat treatment process.
  • Then, by applying an antioxidant to at least one surface of a forming surface or a side surface of a die for hot forging before use, the oxidation resistance of a die for hot forging can be further improved.
  • By using the method for producing a hot forging die described above, a hot forging die having superior high-temperature compressive strength such that the 0.2% compressive proof strength at a test temperature of 1100°C and a strain rate of 10-3/s is 450 MPa or more can be obtained.
    • Method for producing forged product
  • Representative steps for producing a forged product by using a hot forging die forging that uses the Ni-based alloy for a hot die according to the present invention will be described.
  • First, as a first step, a forging material is heated to a predetermined forging temperature. The forging temperature varies depending on the material, and thus, the temperature is adjusted appropriately. A hot forging die according to the present invention has the property of enabling isothermal forging and hot die forging even in atmospheric air at a high temperature, and thus is suitable for hot forging of a Ni-base heat-resistant superalloy, a Ti alloy, or the like, which is known as a poor workability material. A representative forging temperature is in the range of 1000 to 1150°C.
  • Then, the forging material heated in the first step is hot forged by using a preheated hot forging die (second step). In the case of a hot die forging or an isothermal forging, the hot forging in the second step is preferably die forging. In addition, a hot forging die according to the present invention enables hot forging in atmospheric air at a high temperature of 1000°C or more by providing a component having an adjusted Cr content or the like, and can achieve a long die life because of the high high-temperature compressive strength thereof described above.
    • EXAMPLES
  • The present invention will be described in more detail with reference to the following Examples. Ingots of Ni-based alloys for a hot die of alloy 1, alloy 2, and alloy 3 shown in Table 1 were produced by vacuum melting. The unit is mass%. In melting, various raw materials each having a weight adjusted in such a way as to have the desired composition were melted into liquid at 1500 to 1600°C, and then cast into a ceramic mold preheated to 800 to 900°C. After casting, the alloy and the mold were allowed to cool slowly to room temperature as is, and after slow cooling, the alloy and the mold were separated from each other. The weight of the ingot was about 10 kg, and the approximate shape of the portion without the riser part was a cube of 100 mm on each side. The amounts of P and O contained in the following ingot were each 0.005% or less. In addition, the amounts of Si and Fe are each 1.0% or less.
  • The central part of the ingot from which a test piece material, to be described later, was cut out consisted of an equiaxed grain having an average size of 0.5 mm or more. The macrostructures of alloy 1, alloy 2, and alloy 3 are shown in FIG. 1a, FIG. 1b, and FIG. 1c, respectively, together with the positions of test pieces for macrostructure observation in the ingots (the positions indicated by solid lines in the upper figures) and the regions in which the average grain size was measured (the regions surrounded by dotted lines in the lower figures). The average grain size of alloy 1 is 7.0 mm, that of alloy 2 is 5.1 mm, and that of alloy 3 is 5.5 mm. The average grain size is defined as the equivalent circle diameter calculated from the average grain cross-sectional area obtained by dividing the area of the measurement region by the total number of grains therein (one located on the boundary line of the region was counted as 0.5, and one located outside was counted as 1). The etching solution used in preparing the test pieces was ferric chloride solution. In the Inventive Examples, heat treatments, which will be described later, were applied to this material, and the average grain size does not become smaller by this heat treatment. Table 1
    Alloy 1 (mass %)
    W Mo Al Cr Ta Nb Ti C B
    14.6 3.0 5.6 2.1 4.1 1.2 - 0.044 0.014
    Co Mg Ca La Y Zr Hf S Balance
    5.9 0.0017 - - - - - 0.0004 Ni
    Alloy 2 (mass %)
    W Mo Al Cr Ta Nb Ti C B
    15.0 2.9 5.7 2.1 4.1 1.2 - 0.046 0.009
    Co Mg Ca La Y Zr Hf S Balance
    10.1 0.0010 - - - - - 0.0006 Ni
    Alloy 3 (mass %)
    W Mo Al Cr Ta Nb Ti C B
    13.3 2.6 5.9 1.6 3.2 - 1.5 0.056 0.017
    Co Mg Ca La Y Zr Hf S Balance
    12.3 0.0009 - - - - 0.14 0.0005 Ni
    * The symbol "-" means not added.
    * The "Ni" under "Balance" includes inevitable impurities.
  • For the purpose of pore observation, 10 × 10 × 8 mm rectangular parallelepipeds were cut out from near the center of the ingots, and solution heat treatment was applied to some of the rectangular parallelepipeds to prepare test pieces of the Inventive Examples simulating the structure of a hot forging die and test pieces of Comparative Examples. The test piece in which the solution heat treatment conditions for alloy 1 were held at 1300°C for 2 h followed by air cooling was designated as Inventive Example No. 1, and the test piece in which the solution heat treatment conditions for alloy 1 were held at 1325°C for 2 h followed by air cooling was designated as Inventive Example No. 2. In addition, the test piece that was cut out and not subjected to heat treatment was designated as Comparative Example No. 21. Furthermore, the test piece in which the solution heat treatment conditions for alloy 2 were held at 1300°C for 2 h followed by air cooling was designated as Inventive Example No. 3, and the test piece in which the solution heat treatment conditions were held 1325°C for 2 h followed by air cooling was designated as Inventive Example No. 4, and the test piece that was cut out and not subjected to heat treatment was designated as Comparative Example No. 22. Additionally, the test piece in which the solution heat treatment conditions for alloy 3 were held at 1300°C for 2 h was designated as Inventive Example No. 5, and the test piece that was cut out and not subjected to heat treatment was designated as Comparative Example No. 23.
  • All test pieces were cut such that a 10 × 10 mm plane passing through the center of the 8 mm side served as an observation surface, and the cut surfaces were mirror-polished by buffing using a diamond paste. Then, 200× optical micrographs of the mirror-polished surfaces of all test pieces were taken, and the images thereof were binarized and then subjected to pore analysis using image processing software ImageJ (provided by the US National Institutes of Health (NIH)). The photographs were taken at 4 locations in close proximity. In addition, the field of view area of the 200× optical micrographs was about 2.5 mm2.
  • Representative examples of the optical micrographs of the test pieces of alloy 1, alloy 2, and alloy 3 are shown in each of FIG. 2a to FIG. 2c, FIG. 2d to FIG. 2f, and FIG. 2g and FIG. 2h. In addition, analysis results of the pore area ratio of each of the test pieces of alloy 1 to alloy 3 are shown in Table 2. The amounts of pores are increased by the solution heat treatment, but in Inventive Examples No. 1 and No. 3, the increase is only slight, and the average for each of the 4 locations is 0.4% or less. A slight increase was found in each of Inventive Examples No. 2 and No. 4 in which the solution heat treatment temperature was high, the average for the 4 locations in No. 2 is 0.8%, and the average for the 4 locations in No. 4 is 1.6%. On the other hand, in Inventive Example No. 5 of alloy 3, the amount of pores increased even at a relatively low solution temperature. The largest pore size is about 3400 µm2 in Inventive Example No. 1, and about 1100 µm2 in Inventive Example No. 3. In No. 2, No. 4, and No. 5, the largest pore size is 4000 µm2 or more. Table 2
    No. Alloy Solution heat treatment Pore area ratio [%]
    1 Alloy 1 1300°C/2 h/air cooling 0.14 0.25 0.38 0.10
    2 1325°C/2 h/air cooling 1.48 0.44 0.47 0.68
    3 Alloy 2 1300°C/2 h/air cooling 0.15 0.13 0.09 0.43
    4 1325°C/2 h/air cooling 0.23 0.16 4.26 1.68
    5 Alloy 3 1300°C/2 h/air cooling 0.33 1.91 0.74 0.78
    21 Alloy 1 No heat treatment 0.14 0.08 0.14 0.09
    22 Alloy 2 0.08 0.20 0.13 0.35
    23 Alloy 3 0.30 0.44 0.59 0.26
  • In addition, for observation of microstructures other than pores, aging heat treatment involving holding at 1100°C for 4 h followed by air-cooling was applied to each of Inventive Examples No. 1 to No. 5, then a mirror-polished surface was prepared in the same manner as described above, and then the polished surface was etched with an etching solution consisting of 50 ml of ethanol, 50 ml of 35%, in % by mass, concentrated hydrochloric acid, and 2.6 g of cupric chloride. A pore is formed by initial melting at a high temperature, and thus, there was no pore change due to the aging heat treatment.
  • Then, 200× optical micrographs of the etched surfaces of all test pieces were taken. The photographs were taken at 2 locations in close proximity. Then, these optical micrographs were binarized and then the eutectic gamma prime phases were analyzed using the image processing software ImageJ.
  • In addition, secondary electron images or backscattered electron images of the microstructures and the gamma prime phases in the dendrite cores were taken by using a scanning electron microscope (SEM) on the etched surfaces of all test pieces. The above-mentioned analysis of the gamma prime phases in the dendrite cores targeted 15000× images (field of view area of about 50 µm2) in Inventive Examples No. 1 to No. 5 and 5000× images (field of view area of about 500 µm2) in Comparative Examples No. 21 to No. 23. Then, these secondary electron images or backscattered electron images were binarized, and then the gamma prime phases in the dendrite cores were analyzed using image processing software ImageJ.
  • Representative examples of the optical micrographs of the test pieces of alloy 1, alloy 2, and alloy 3 are shown in FIG. 3a to FIG. 3c, FIG. 3d to FIG. 3f, and FIG. 3g and FIG. 3h, respectively. In addition, the secondary electron images or backscattered electron images of alloy 1, alloy 2, and alloy 3 are shown in FIG. 4a to FIG. 4c, FIG. 5d to FIG. 5f, and FIG. 6g and FIG. 6h. In the images shown in FIG. 4 to FIG. 6, which were targeted for gamma prime phase analysis, the number of gamma prime phase particles in all images was 150 or more and 1500 or less. In addition, the distribution of the area ratio of the gamma prime phase for each circle equivalent diameter in alloy 1, alloy 2, and alloy 3 and the cumulative area ratio thereof are shown in FIG. 7, FIG. 8, and FIG. 9. In addition, the eutectic gamma prime phase area ratios in alloy 1, alloy 2, and alloy 3 are shown in Table 3. In the figures and tables, the gamma prime phase is expressed as "y' phase." Table 3
    No. Alloy Solution heat treatment + aging heat treatment Eutectic γ' phase area ratio [%]
    1 Alloy 1 1300°C/2 h/air cooling + 1100°C/4 h/air cooling 5.42 6.42
    2 1325°C/2 h/air cooling + 1100°C/4 h/air cooling 0.26 0.17
    3 Alloy 2 1300°C/2 h/air cooling + 1100°C/4 h/air cooling 1.60 2.41
    4 1325°C/2 h/air cooling + 1100°C/4 h/air cooling 0.16 0.11
    5 Alloy 3 1300°C/2 h/air cooling + 1100°C/4 h/air cooling 9.27 10.05
    21 Alloy 1 No heat treatment 4.35 5.25
    22 Alloy 2 4.14 3.74
    23 Alloy 3 9.41 10.56
  • All of Inventive Examples No. 1 to No. 5 similarly had an equivalent circle diameter of the gamma prime phase in the dendrite core of 1.20 µm or less. In other words, the area ratio of gamma prime phases having an equivalent circle diameter of 2.00 µm or less in gamma prime phases having an equivalent circle diameter of 2.20 µm or less in the dendrite core is 100%. In addition, the area ratio of the eutectic gamma prime phase in each of Inventive Examples No. 1, No. 3, and No. 5 was 1.0% or more, whereas the area ratio of the eutectic gamma prime phase in each of Inventive Examples No. 2 and No. 4 was 1.0% or less. In addition, alloy 2 has a lower area ratio of the eutectic gamma prime phase than alloy 1.
  • Next, bars having a diameter of 10 mm were cut out from the central part of the ingots, heat treatment was applied to some of the bars, then a material for collecting a test piece having a diameter of 8 mm and a height of 12 mm was cut out, and the surface thereof was finished equivalent to #1000 to prepare compression test pieces of Inventive Examples simulating the structure of a hot forging die and compression test pieces of Comparative Examples. The heat treatment conditions were the same solution heat treatment and aging heat treatment as for the test pieces for microstructure observation, and thus, these test pieces correspond to Inventive Examples No. 1 to No. 5 and Comparative Examples No. 21 to No. 23 described above. For each of the Inventive Examples, two compression test pieces were prepared.
  • The compression test conditions are a test temperature of 1100°C, a strain rate of 10-3/s, and a compression reduction ratio of 10%, and the atmosphere was atmospheric air. Then, the high-temperature compressive strength was evaluated by deriving the 0.2% compressive proof strength from the stress-strain curve obtained from the compression test. This compression test tests whether a test piece has sufficient compressive strength even at a high temperature as a hot forging die under a particularly high load, and if the 0.2% compressive proof strength is 450 MPa or more at a test temperature of 1100°C, intended for isothermal forging, the test piece can be deemed to have sufficient strength. The 0.2% compressive proof strength is further preferably 500 MPa or more.
  • Test results of the compression test pieces of Inventive Examples No. 1 to No. 5 and Comparative Examples No. 21 to No. 23 are shown in Table 4. In addition, as a representative example, the high-temperature compressive strength of alloy 1 is shown in FIG. 10. From Table 4, it can be seen that Inventive Examples No. 1, No. 2, and No. 4 all have a 0.2% compressive proof strength of 450 MPa or more and have excellent high-temperature compressive strength. Even in Inventive Example No. 5, the average value of the results of compression carried out twice is 450 MPa or more. Inventive Example No. 3 has a 0.2% compressive proof strength of 450 MPa or less, and thus, it may not be suitable for an application at a high temperature such as 1100°C, but has higher strength than Comparative Example No. 22 and provides a longer die service life than the same, and a sufficient die service life can be achieved. In addition, as shown in Table 2 and Table 3, the pore area ratio and the eutectic gamma prime phase area ratio of Inventive Example No. 3 are lower than those of the other Inventive Examples, and thus, Inventive Example No. 3 is suitable for an application in which a high load is repeatedly applied at a relatively low temperature such as 1050°C. Table 4
    No. Alloy Solution heat treatment + aging heat treatment 0.2% compressive proof strength [MPa]
    1 Alloy 1 1300°C/2 h/air cooling + 1100°C/4 h/air cooling 530 477
    2 1325°C/2 h/air cooling + 1100°C/4 h/air cooling 544 509
    3 Alloy 2 1300°C/2 h/air cooling + 1100°C/4 h/air cooling 387 399
    4 1325°C/2 h/air cooling + 1100°C/4 h/air cooling 513 450
    5 Alloy 3 300°C/2 h/air cooling + 1100°C/4 h/air cooling 439 528
    21 Alloy 1 No heat treatment 404
    22 Alloy 2 378 378
    23 Alloy 3 378 361
  • From these results, it can be seen that a hot forging die according to the present invention has superior high-temperature compressive strength and can achieve a long die service life. It can be seen that a hot forging die according to the present invention having the properties described above is suitable for hot die forging and for isothermal forging.

Claims (10)

  1. A hot forging die comprising a casting alloy having a component composition consisting of, by mass%,
    7.5 to 20.0% W,
    0 to 5.0% Mo,
    5.0 to 7.5% Al,
    0.5 to 5.0% Cr,
    1.0 to 12.0% Ta,
    0.01 to 0.15% C,
    0.03% or less B,
    0.015% or less S,
    0 to 0.020% in total of one or more selected from rare earth elements, Y, Ca, and Mg,
    0.5% or less in total of one or two selected from Zr and Hf,
    5.0% or less Ti,
    5.0% or less Nb,
    25.0% or less Co, and
    the balance of Ni with inevitable impurities,
    wherein an equivalent circle diameter of a gamma prime phase in a dendrite core is 2.20 µm or less.
  2. The hot forging die according to claim 1, comprising a casting alloy having a component composition consisting of, by mass%,
    W: 10.0 to 20.0%,
    Mo: 0.5 to 5.0%,
    Al: 5.0 to 7.5%,
    Cr: 0.5 to 4.0%,
    Ta: 1.0 to 12.0%,
    C: 0.01 to 0.15%,
    B: 0.03% or less,
    S: 0.015% or less,
    0 to 0.020% in total of one or two or more selected from rare earth elements, Y, Ca, and Mg,
    0.5% or less in total of one or two selected from Zr and Hf,
    Ti: 5.0% or less,
    Nb: 5.0% or less,
    Co: 20.0% or less, and
    the balance of Ni with inevitable impurities,
    wherein an equivalent circle diameter of a gamma prime phase in a dendrite core is 2.20 µm or less.
  3. The hot forging die according to claim 1 or 2, wherein an area ratio of gamma prime phases having an equivalent circle diameter of 2.00 µm or less in gamma prime phases having an equivalent circle diameter of 2.20 µm or less in a dendrite core is 90% or more.
  4. The hot forging die according to claim 1 or 2, wherein an area ratio of a eutectic gamma prime phase is 4.0% or less.
  5. The hot forging die according to claim 1 or 2, wherein a pore area ratio is 0.7% or less.
  6. The hot forging die according to claim 5, wherein a size of each pore is 4000 µm2 or less.
  7. The hot forging die according to claim 1 or 2, wherein an average grain size is 0.5 mm or more.
  8. The hot forging die according to claim 1 or 2, wherein a 0.2% compressive proof strength at a test temperature of 1100°C and a strain rate of 10-3/s is 450 MPa or more.
  9. A method for producing a hot forging die, comprising subjecting a casting alloy having the component composition according to claim 1 or 2 to solution heat treatment at a temperature of 1250 to 1350°C for 0.5 hours or more.
  10. The method for producing a hot forging die according to claim 9, further comprising, after the solution heat treatment, subjecting it to aging heat treatment in a temperature range of 800°C to 1150°C.
EP23865454.5A 2022-09-14 2023-09-11 DIE FOR HOT FORGING AND MANUFACTURING METHOD FOR IT Pending EP4588590A4 (en)

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JPS6250429A (en) 1985-08-30 1987-03-05 Hitachi Metals Ltd Nickel-base casting alloy for hot forging die
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JP2016069702A (en) 2014-09-30 2016-05-09 日立金属株式会社 Method for producing nickel-base cast alloy
JP6476704B2 (en) 2014-09-30 2019-03-06 日立金属株式会社 Nickel base casting alloy and hot forging die
KR102193336B1 (en) * 2017-11-17 2020-12-22 미츠비시 파워 가부시키가이샤 Ni-based forged alloy material and turbine high-temperature member using the same
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WO2020059846A1 (en) * 2018-09-21 2020-03-26 日立金属株式会社 Ni-based alloy for hot die, and hot forging die obtained using same
JP7452172B2 (en) * 2019-03-29 2024-03-19 株式会社プロテリアル Method for manufacturing hot forged materials
JP2020196951A (en) * 2019-05-28 2020-12-10 日立金属株式会社 Ni-BASED ALLOY FOR HOT DIE, HOT FORGING DIE USING THE Ni-BASED ALLOY FOR HOT DIE, AND METHOD FOR MANUFACTURING FORGED PRODUCT USING THE HOT FORGING DIE
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