Die for hot forging and method for manufacturing same

Abstract

Provided are: a die for hot forging, which has high high-temperature compressive strength and is capable of achieving a long die life; and a method for manufacturing the die. Provided is a hot forging die comprising a casting alloy having the following composition in mass%: 7.5-20.0% of W, 0-5.0% of Mo, 5.0-7.5% of Al, 0.5-5.0% of Cr, 1.0-12.0% of Ta, 0.01-0.15% of C, 0.03% or less of B, 0.015% or less of S, 0-0.020% in total of one or more elements selected from the group consisting of rare earth elements, Y, Ca, and Mg, 0.5% or less in total of one or more elements selected from the group consisting of Zr and Hf, 5.0% or less of Ti, 5.0% or less of Nb, 25.0% or less of Co, and the balance of Ni and unavoidable impurities, the equivalent circle diameter of the gamma'phase in the dendritic crystal trunk is 2.20 [mu] m or less, and a preferred method for manufacturing a hot forging die is also provided.

Description

Technical Field

[0001] The present invention relates to a hot forging die and a manufacturing method thereof. Background Art

[0002] In the forging of products made of heat-resistant alloys, the forging blank is heated to a specified temperature in order to reduce deformation resistance. However, since heat-resistant alloys have high strength even at high temperatures, the hot forging die used for forging them requires high mechanical strength. In addition, when the temperature of the hot forging die is lower than that of the forging blank during hot forging, the workability of the forging blank is reduced due to heat dissipation. Therefore, in the forging of products made of difficult-to-process materials such as Alloy718 and Ti alloy, the hot forging die is used in a state heated to a high temperature that is the same as or close to the heating temperature of the forging blank. Therefore, the hot forging die must have high mechanical strength at high temperatures. As a hot forging die that meets this requirement, a Ni-based super heat-resistant alloy having high high-temperature compressive strength and capable of being used in hot forging at a die temperature of more than 1000°C in the atmosphere has been proposed (for example, see patent documents 1 to 6).

[0003] It should be noted that the hot forging referred to in the present invention includes hot die forging in which the temperature of the hot forging die is brought close to the temperature of the forging material, and isothermal forging in which the temperature is set to the same temperature as the forging material.

[0004] Prior art literature

[0005] Patent Literature

[0006] Patent Document 1: Japanese Patent Application Publication No. 62-50429

[0007] Patent Document 2: Japanese Patent Application Publication No. 60-221542

[0008] Patent Document 3: Japanese Patent Application Publication No. 2016-069702

[0009] Patent Document 4: Japanese Patent Application Publication No. 2016-069703

[0010] Patent Document 5: Japanese Patent No. 6645627

[0011] Patent Document 6: U.S. Patent No. 4740354 Summary of the invention

[0012] Problem that the invention aims to solve

[0013] The above-mentioned Ni-based super heat-resistant alloy has a structure with a large amount of precipitation strengthening phase, contains a large amount of solid solution strengthening elements and has high high-temperature strength, so it can be used as a hot forging die. However, when a large number of forgings with a higher load ratio and complex shapes are manufactured at high forging temperatures, even for the above-mentioned Ni-based super heat-resistant alloy, the die life is a technical problem. In order to solve this technical problem, it is necessary to use a hot forging die composed of an alloy with higher high-temperature compressive strength.

[0014] The object of the present invention is to provide a hot forging die having high high temperature compressive strength which is particularly advantageous in adapting to a die subjected to a high load, and also to provide a method for manufacturing a hot forging die which is preferred therein.

[0015] Solutions for solving problems

[0016] The present inventors studied the above-mentioned technical problems, found a hot forging die having high high-temperature compressive strength, and completed the present invention.

[0017] That is, the present invention is a hot forging die, which is composed of a casting alloy having the following composition: in mass %, W: 7.5-20.0%, Mo: 0-5.0%, Al: 5.0-7.5%, Cr: 0.5-5.0%, Ta: 1.0-12.0%, C: 0.01-0.15%, B: less than 0.03%, S: less than 0.015%, one or more selected from rare earth elements, Y, Ca, Mg, the total of which is 0-0.020%, one or two selected from Zr and Hf elements, the total of which is less than 0.5%, Ti: less than 5.0%, Nb: less than 5.0%, Co: less than 25.0%, the remainder is Ni and unavoidable impurities, and the circle equivalent diameter of the γ' phase in the dendrite trunk is less than 2.20μm. Furthermore, the following hot forging die is preferably composed of a casting alloy having the following composition: in mass %, W: 10.0-20.0%, Mo: 0.5-5.0%, Al: 5.0-7.5%, Cr: 0.5-4.0%, Ta: 1.0-12.0%, C: 0.01-0.15%, B: 0.03% or less, S: 0.015% or less, one or more selected from rare earth elements, Y, Ca, Mg, the total of 0-0.020%, one or two selected from Zr and Hf elements, the total of 0.5% or less, Ti: 5.0% or less, Nb: 5.0% or less, Co: 20.0% or less, the balance being Ni and unavoidable impurities, and the circle equivalent diameter of the γ' phase in the dendrite trunk is 2.20μm or less.

[0018] In the present invention, the hot forging die is preferably one in which the area ratio of the γ' phase having an equivalent circle diameter of 2.00 μm or less among the γ' phase having an equivalent circle diameter of 2.20 μm or less in the dendrite trunk is 90% or more.

[0019] In the present invention, it is preferred that the hot forging die has an area ratio of the eutectic γ' phase of 4.0% or less.

[0020] In the present invention, the hot forging die preferably has a pore area ratio of 0.7% or less. More preferably, the size of each pore is 4000 μm. 2 The following hot forging dies are used.

[0021] In the present invention, a hot forging die having an average grain size of 0.5 mm or more is preferred.

[0022] In addition, the present invention preferably has a test temperature of 1100°C and a strain rate of 10 -3 A hot forging die having a 0.2% compressive yield strength at a pressure of 450 MPa or more.

[0023] Furthermore, the present invention is a method for manufacturing a hot forging die, wherein a solution heat treatment is performed on the cast alloy having the above-mentioned composition at a temperature of 1250 to 1350° C. for 0.5 hours or more.

[0024] In the present invention, it is preferable that the method for manufacturing a hot forging die is performed by further performing an aging heat treatment in a temperature range of 800°C to 1150°C after performing the above-mentioned solution heat treatment.

[0025] Effects of the Invention

[0026] According to the present invention, a hot forging die having high high temperature compressive strength can be obtained, thereby achieving a long die life. BRIEF DESCRIPTION OF THE DRAWINGS

[0027] Figure 1 : is a figure which shows the macro structure of the ingot used in Example.

[0028] Figure 2 The figure shows the optical microscope photograph of the pores of the example of the present invention and the comparative example.

[0029] Figure 3 The figures are optical microscope photographs showing the microstructures of the examples of the present invention and the comparative examples.

[0030] Figure 4 The figures are secondary electron images or reflected electron images showing the microstructures of the examples of the present invention and the comparative examples.

[0031] Figure 5The figures are secondary electron images or reflected electron images showing the microstructures of the examples of the present invention and the comparative examples.

[0032] Figure 6 The figures are secondary electron images or reflected electron images showing the microstructures of the examples of the present invention and the comparative examples.

[0033] Figure 7 It is a graph showing the distribution and cumulative area ratio of the area ratio per circle-equivalent diameter of the γ′ phase of the examples of the present invention and the comparative examples.

[0034] Figure 8 It is a graph showing the distribution and cumulative area ratio of the area ratio per circle-equivalent diameter of the γ′ phase of the examples of the present invention and the comparative examples.

[0035] Fig. 9 It is a graph showing the distribution and cumulative area ratio of the area ratio per circle-equivalent diameter of the γ′ phase of the examples of the present invention and the comparative examples.

[0036] Fig.10 It is a graph showing the high temperature compressive strength of the examples of the present invention and the comparative examples. DETAILED DESCRIPTION

[0037] The hot forging die of the present invention will be described in detail below. First, a hot forging die Ni-based alloy as a blank of the hot forging die will be described. The unit of chemical composition is mass %.

[0038] <W>

[0039] W is dissolved in the austenite matrix (γ phase) and also in the gamma prime phase (γ' phase) based on Ni3Al as a precipitation strengthening phase, thereby improving the high temperature strength of the alloy. In addition, W forms MC carbides together with C described later, precipitates at the grain boundaries and improves the grain boundary strength, thereby improving the high temperature strength and ductility. On the other hand, W has the effect of reducing oxidation resistance and easily precipitating harmful phases such as TCP (Topologically Close Packed) phases. From the perspective of improving high temperature strength, ductility and suppressing the reduction of oxidation resistance and the precipitation of harmful phases, the W content in the Ni-based alloy of the present invention is set to 7.5-20.0%. In order to more reliably obtain the effect of W, the preferred lower limit is 10.0%, and the further preferred lower limit is 12.0%. In addition, the preferred upper limit is 16.0%, and the further preferred upper limit is 15.0%.

[0040]

[0041] Mo is dissolved in the austenite matrix like W, and is also dissolved in the γ' phase based on Ni3Al as a precipitation strengthening phase, thereby improving the high temperature strength of the alloy. On the other hand, Mo also has the effect of reducing oxidation resistance and easily precipitating harmful phases such as TCP phase. In addition, when the Mo content is too much, M6C carbides are formed together with C described later during the maintenance process at high temperature, and the solid solubility is reduced, so the high temperature strength during use is reduced. When this reduction becomes a problem, it is preferably not contained. From the perspective of improving high temperature strength and suppressing the reduction of oxidation resistance and high temperature strength during use, the Mo content in the Ni-based alloy of the present invention is set to 0 to 5.0% below the W content. In order to more reliably obtain the effect of Mo, the preferred lower limit is 0.5%, and the further preferred lower limit is 1.5%. In addition, the preferred upper limit is 4.0%, and the further preferred upper limit is 3.5%.

[0042] <Al>

[0043] Al combines with Ni to precipitate a γ' phase composed of Ni3Al, improves the high temperature strength of the alloy, forms an aluminum oxide film on the surface of the alloy, and imparts oxidation resistance to the alloy. On the other hand, when the Al content is too much, it has the effect of excessively generating a γ' phase and reducing the toughness of the alloy. From the perspective of improving oxidation resistance and high temperature strength and suppressing the reduction of toughness, the Al content in the Ni-based alloy of the present invention is set to 5.0-7.5%. In order to more reliably obtain the effect of Al, the preferred lower limit is 5.2%, and the further preferred lower limit is 5.4%. In addition, the preferred upper limit of Al is 6.7%, and the further preferred upper limit is 6.5%.

[0044] <Cr>

[0045] Cr has the effect of promoting the formation of a continuous layer of aluminum oxide on the surface or inside of the alloy and improving the oxidation resistance of the alloy. Therefore, it is necessary to contain more than 0.5% Cr. On the other hand, when the Cr content is too much, it also has the effect of easily precipitating harmful phases such as TCP phase. In particular, when the austenite matrix and γ' phase contain a large amount of elements such as W, Mo, Ta, etc. that improve the high temperature strength of the alloy, harmful phases are easily precipitated. From the perspective of improving oxidation resistance and maintaining the content of elements that improve high temperature strength at a high level while suppressing the precipitation of harmful phases, the Cr content in the present invention is set to 0.5-5.0%. In order to more reliably obtain the effect of Cr, the preferred lower limit is 1.2%. The preferred upper limit of Cr is 4.0%, the further preferred upper limit is 3.0%, and the more preferred upper limit is 2.5%.

[0046] <Ta>

[0047] Ta is dissolved in the γ' phase composed of Ni3Al to replace Al sites and improve the high temperature strength of the alloy. Furthermore, the adhesion and oxidation resistance of the oxide film formed on the surface of the alloy are improved, and the oxidation resistance of the alloy is improved. In addition, Ta forms MC carbides together with C described later, precipitates at the grain boundaries and improves the grain boundary strength, thereby improving the high temperature strength and ductility. On the other hand, when the content of Ta is too much, it has the effect of easily precipitating harmful phases such as TCP phases and excessively generating γ' phases to reduce the toughness of the alloy. From the perspective of improving oxidation resistance and high temperature strength and inhibiting the precipitation of harmful phases and the reduction of toughness, the content of Ta in the present invention is set to 1.0-12.0%. In order to more reliably obtain the effect of Ta, the preferred lower limit is 2.5%, and the more preferred lower limit is 3.0%. The preferred upper limit of Ta is 10.0%, and the more preferred upper limit is 7.0%.

[0048] <C>

[0049] C forms MC carbides together with W, Mo, Ta, etc., and improves the effects of grain boundary strength and the like by precipitation at the grain boundaries, thereby improving the high temperature strength and ductility. On the other hand, when the content of C is too much, it also has the effect of reducing the high temperature strength of the alloy by the formation of coarse carbides and the formation of M6C carbides during high temperature maintenance, which leads to a significant reduction in the amount of Mo solid solution. In addition, there is also a risk of reducing fatigue strength due to the lowering of the initial melting temperature and the formation of pores during the solution treatment described later. Therefore, from the perspective of improving the high temperature strength and ductility of the alloy and suppressing the reduction of high temperature strength, etc., the content of C in the present invention is set to 0.01-0.15%. In order to more reliably obtain the effect of C, the preferred lower limit is 0.02%, and the further preferred lower limit is 0.04%. The preferred upper limit of C is 0.13%, and the further preferred upper limit is 0.12%.

[0050]

[0051] The Ni-based alloy for hot molds in the present invention can contain 0.03% or less (including 0%) of B (boron). B, like carbides, increases the strength of the grain boundaries of the alloy, and improves high-temperature strength and ductility. On the other hand, when the B content is too much, coarse borides are formed, which also has the effect of reducing the strength of the alloy. In addition, due to the formation of low-melting-point borides, pores are formed during the solution treatment described later, thereby also presenting a risk of reducing fatigue strength. Therefore, especially when it is desired to improve high-temperature strength, ductility, etc., B can be added as needed. In order to reliably obtain the effect of B, the preferred lower limit is 0.005%, and the further preferred lower limit is 0.01%. The preferred upper limit is 0.02%, and the more preferred upper limit is 0.015% or less.

[0052] <S>

[0053] <Rare earth elements, Y, Ca and Mg>

[0054] In addition, in the Ni-based alloy for hot molds of the present invention, S (sulfur) reduces the adhesion of the oxide film by segregating to the interface between the oxide film formed on the surface of the alloy and the alloy and inhibiting their chemical bonding. Therefore, the upper limit of S is limited to 0.015% or less (including 0%). In addition, it is preferred that one or more of the rare earth elements, Y, Ca and Mg elements that form sulfides with S are contained in a range of 0.020% or less in total while limiting S. Regarding these rare earth elements, Y, Ca and Mg, there is a risk of reducing fatigue strength because low melting point compounds are formed when added excessively and pores are formed during the solution treatment described later. Therefore, the upper limit of the total amount of rare earth elements, Y, Ca and Mg is set to 0.020%. It should be noted that S is a component that can be contained as an impurity, and sometimes it is greater than 0% and remains. When the S content may reach 0.0001% (1 ppm) or more, it is preferable to contain one or more of the rare earth elements, Y, Ca and Mg elements in an amount greater than or equal to S. It should be noted that in the Ni-based alloy of the present invention, the rare earth elements, Y, Ca and Mg elements may be 0%.

[0055] Among the rare earth elements, La is preferably used, which has the effect of improving oxidation resistance by a mechanism other than the harmlessness of S. From an economic point of view, Ca or Mg is preferably used. In addition, Mg has a smaller effect of reducing toughness and ductility than Ca, and can also be expected to prevent cracks during casting. Therefore, when any one of rare earth elements, Y, Ca and Mg is selected, Mg is preferably used. In the case where a sufficient effect can be obtained by adding Mg, Ca is not added. It should be noted that in order to reliably obtain the effect of Mg, it is appropriate to contain 0.0002% or more regardless of the presence or absence of S. It is preferably 0.0005% or more, and more preferably 0.0010% or more.

[0056] <Zr and Hf>

[0057] The Ni-based alloy for hot molds in the present invention can contain one or two selected from Zr and Hf in a range of 0.5% or less (including 0%) in total. These elements have the effect of improving oxidation resistance by a mechanism other than the harmlessness of S. In addition, they form MC carbides together with the aforementioned C, precipitate at the grain boundaries and improve the grain boundary strength, thereby improving high temperature strength and ductility. However, these elements also form low melting point compounds when added too much and cause pores to form during the solution treatment described later, thereby 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 can be added as needed. It should be noted that Hf can also be expected to prevent cracks during casting. Therefore, when any one of Zr and Hf is selected, Hf is preferably used. In order to reliably obtain the effects of one or two selected from Zr and Hf, the preferred lower limit of the total is 0.01%, and the further preferred lower limit is 0.02%. The preferred upper limit is 0.3%, and the more preferred upper limit is 0.2%.

[0058] <Ti>

[0059] The Ni-based alloy for hot molds in the present invention can contain Ti. Ti, like Ta, solid-dissolves in the form of replacing Al sites in the γ' phase composed of Ni3Al, and forms MC carbides together with the aforementioned C to improve the high-temperature strength of the alloy. In addition, since it is a cheaper element than Ta, it is advantageous in terms of mold cost. On the other hand, when the content of Ti is too much, it also has the effect of easily precipitating harmful phases such as TCP phases, and excessively generating γ' phases to reduce the toughness of the alloy, just like Ta. In addition, unlike Ta, Ti does not have the effect of improving oxidation resistance. In the present invention, from the perspective of suppressing excessive reduction in oxidation resistance and reducing mold cost, Ti can be contained in a range of 5.0% or less (including 0%). It should be noted that, unlike Nb described later, Ti has the effect of finely dispersing carbides, so when ductility is particularly important, it is appropriate to select Ti instead of Nb. In order to reliably obtain the effect of Ti, the preferred lower limit is 0.5%, and more preferably 1.0%. In addition, the preferred upper limit is 3.5%.

[0060] <Nb>

[0061] The Ni-based alloy for hot molds in the present invention can contain Nb. Nb, like Ta, solid-dissolves in the form of replacing Al sites in the γ' phase composed of Ni3Al, and forms MC carbides together with the aforementioned C to improve the high-temperature strength of the alloy. In addition, since it is a cheaper element than Ta, it is advantageous in terms of mold cost. On the other hand, when the content of Nb is too much, it also has the effect of easily precipitating harmful phases such as TCP phases, and excessively generating γ' phases to reduce the toughness of the alloy, just like Ta. In addition, unlike Ta, Nb does not have the effect of improving oxidation resistance. In the present invention, from the perspective of suppressing excessive reduction in oxidation resistance and reducing mold cost, Nb can be contained in a range of 5.0% or less (including 0%). It should be noted that the carbides formed by Nb at high temperatures are more stable than those formed by Ti, so when the stability of mechanical properties at high temperatures is particularly important, it is appropriate to select Nb instead of Ti. In order to reliably obtain the effect of Nb, the preferred lower limit is 0.5%, and the more preferred lower limit is 1.0%. In addition, the preferred upper limit is 3.5%.

[0062] <Co>

[0063] The Ni-based alloy for hot molds in the present invention can contain Co. Co dissolves in the austenite matrix and improves the high temperature strength of the alloy. In addition, it also has the effect of suppressing the formation of coarse M6C carbides and the effect of reducing the temperature of the solution treatment described later by lowering the solid solution temperature of the γ' phase. On the other hand, when the content of Co is too much, since Co is an element more expensive than Ni, the mold cost will increase. The solid solution strengthening ability of Co is lower than that of W and Mo, so when high high temperature strength and phase stability are taken into account by adjusting the content of W, Mo, etc., the addition of Co is not necessary. In the case where the cost increase can be allowed, the solid solution temperature of the γ' phase must be lowered due to the limitation of the capacity of the heating furnace used in the solution heat treatment, etc., it can be added as needed. In the present invention, from the perspective of improving the high temperature strength and suppressing the excessive increase in mold cost, Co can be contained in a range of 25.0% or less (including 0%). In order to reliably obtain the effect of Co, the preferred lower limit is 2.0%, and more preferably 3.0%. In addition, the preferred upper limit is 20.0%, more preferably 15.0%, and more preferably 10.0%.

[0064] <Margin>

[0065] In the Ni-based alloy for hot molds of the present invention, Ni and inevitable impurities are included in addition to the aforementioned elements. In the Ni-based alloy for hot molds of the present invention, Ni is the main element constituting the γ phase, and constitutes the γ' phase together with Al, Ta, Ti, Mo, and W. In addition, as inevitable impurities, P, O, N, Si, Mn, Fe, Cu, etc. are envisioned, and in the case of casting ingots in a furnace commonly used for Ni-based alloys, V, Re, and Ru are also envisioned. P, O, and N can be contained as long as they are less than 0.005% respectively, and Si, Mn, Fe, Cu, V, Re, and Ru can be contained as long as they are less than 1.0% respectively, preferably less than 0.5% respectively. In addition, the Ni-based alloy of the present invention can also be called a Ni-based heat-resistant alloy.

[0066] <γ' phase>

[0067] The aforementioned Ni-based alloy for hot forging molds contains Al and Ta as essential elements, and therefore is mainly composed of an austenite matrix and a γ' phase. In addition, since it also contains C, carbides are also present. The hot forging mold of the present invention is obtained by machining an ingot (ingot of a near-net-shape mold) having the above-mentioned composition but not undergoing hot plastic processing, so the Ni-based alloy for hot forging molds constituting it has a dendritic structure. Furthermore, the hot forging mold of the present invention is characterized in that a billet having an equivalent circular diameter of 2.20 μm or less in the γ' phase in the dendrite stem having high high-temperature strength in the Ni-based alloy for hot forging mold is used.

[0068] The dendrite stems mentioned here are, for example, Figure 4 and Figure 5 In the area surrounded by carbides and eutectic γ' phase (in the 250x field of view, the light-colored broad area surrounded by the black irregular-shaped area, i.e., the eutectic γ' phase and the white dendritic carbide phase), the 2 When observed under the field of view area of ​​left and right, the area in which the size of the γ' phase contained in the field of view is relatively uniform, excluding the surrounding eutectic γ' phase and other locally coarse γ' phases and carbides. Furthermore, the γ' phase in the dendrite trunk mentioned here refers to the γ' phase existing in the area captured in the field of view where the size of the γ' phase is finer in the dendrite trunk and the number of γ' phase particles is more than 150 and less than 1500 (the "quadrilateral surrounded" area in the figure) when the dendrite trunk is observed under the same field of view area. It should be noted that, here, the γ' phase is regarded as particles with a circular equivalent diameter of more than 0.05μm, that is, the γ' phase. Too small γ' phase is difficult to identify. In addition, this is because, if the γ' phase has a circular equivalent diameter of less than 0.05μm, it is easy to dissolve in the heating before forging when used as a hot forging mold, and will not affect the strength involved in the effect of the present invention.

[0069] Furthermore, by setting the equivalent circular diameter of the γ' phase in the dendrite trunk to 2.20 μm or less, even if a certain amount of eutectic γ' phase is present in the structure, or even if the composition is the same as that of the Ni-based alloy, the high temperature strength is improved compared to the billet in which the γ' phase with a circular equivalent diameter greater than 2.20 μm is present in the dendrite trunk. This is because if the equivalent circular diameter of the γ' phase in the dendrite trunk is 2.20 μm or less, even if the eutectic γ' phase has a higher solid solution temperature than the solid solution temperature in the dendrite trunk due to the segregation of Ta and remains in the state of casting after the solution heat treatment, the area ratio of the γ' phase in the whole is low, so the influence on the strength belonging to the macroscopic characteristics is small compared to the influence of the refinement effect of the γ' phase in the dendrite trunk. In order to more reliably obtain the effect of improving the high temperature strength, the equivalent circular diameter of the γ' phase in the dendrite trunk is preferably 2.00 μm or less, more preferably 1.80 μm or less, and further preferably 1.20 μm or less.

[0070] In addition, regarding the above-mentioned "the equivalent circular diameter of the γ' phase in the dendrite trunk is 2.20 μm or less", if the area ratio of the γ' phase with an equivalent circular diameter of 2.00 μm or less in the γ' phase is 90% or more of the entire γ' phase, the high-temperature strength improvement effect brought about by the refinement of the γ' phase can be obtained more reliably. In addition, it is preferred that the area ratio of the γ' phase with an equivalent circular diameter of 1.80 μm or less is 90% or more, and it is more preferred that the area ratio of the γ' phase with an equivalent circular diameter of 1.60 μm or less is 90% or more.

[0071] It should be noted that in order to form such a structure, it is necessary to maintain the structure after casting at a high temperature. However, in the Ni-based alloys envisioned for use in engine parts, etc., the pores (described later) formed by the initial melting of grain boundary strengthening elements such as C or inevitable impurities during the high-temperature holding process will lead to a decrease in fatigue strength and ductility, which becomes a technical problem. However, in the hot forging die of the present invention, the main emphasis is placed on high-temperature compressive strength, and the requirements for fatigue strength and ductility are relatively low. Therefore, even if there are a certain amount of pores formed by high-temperature holding or pores during manufacturing that are inevitably present regardless of whether or not high-temperature holding is performed, it will not become a major technical problem.

[0072] In addition, the γ' phase involved in the requirements of the present invention refers to the γ' phase in the state before being used in the hot forging die, such as after the heat treatment process during the mold manufacturing. And even if the γ' phase is coarsened during use as a hot forging die, such as rafting caused by continuous load, the state before use is important from the perspective of the total mold life from the beginning of use, so it will not be a big problem. In addition, regarding the area where almost no load is applied during use, such as the area where the equivalent stress does not exceed 100MPa, it will not be a problem from the perspective of mold life, so the equivalent circle diameter of the γ' phase can be 2.20μm or more. How to set such an area is determined according to the use environment and the requirements of the mold life. It should be noted that for other organizational factors such as the eutectic γ' phase and pores described later, the target area is the same as the above content.

[0073] <Eutectic γ' phase>

[0074] In the hot forging die of the present invention, when the area ratio of the eutectic γ' phase in the structure (observation field) is 4.0% or less, higher high temperature strength can be obtained. It should be noted that the eutectic γ' phase referred to here refers to, for example Figure 3 Shown at about 2.5mm 2 When an optical microscope is used to observe the field of view (observation magnification 200 times), an irregularly shaped area that looks whiter (appears lighter in color) than the surrounding area. If the etching is insufficient and it is difficult to distinguish under an optical microscope, it can also refer to, for example Figure 4 and Figure 5 The irregularly shaped area shown in the figure appears darker than the surrounding area when observed by secondary electron image or reflected electron image under the same viewing area. The reason why the high temperature strength increases when the area ratio of the eutectic γ' phase decreases is that the γ' phase that constitutes the eutectic γ' phase, which has little effect on the strength due to its coarseness, is refined, and the volume fraction of the γ' phase that affects the strength becomes higher. However, in order to set the area ratio of the eutectic γ' phase to less than 4.0%, it is necessary to maintain the structure after casting at a higher temperature, so the amount of pores (described later) will increase, and even hot forging dies may have adverse effects when a large number of forgings are manufactured at a lower temperature and a high strain rate. It should be noted that when a higher load is repeatedly applied, the eutectic γ' phase is more likely to have an adverse effect than the pores. In this case, it is preferred to reduce the area ratio of the eutectic γ' phase as much as possible within the range where the amount of pores does not increase excessively. Therefore, based on the use temperature of the mold and the degree of load applied to the mold, the area ratio of the eutectic γ' phase can be set to less than 4.0% as needed. The area ratio of the eutectic γ' phase is preferably set to 3.0% or less, and more preferably set to 2.0% or less.

[0075] <Porosity>

[0076] In hot forging dies, the effect of pores on the life and safety of parts is smaller than that of Ni-based alloys used in engine parts, etc. This is because in hot forging dies, high temperature compressive strength is more important than fatigue strength, and the effect of pores on high temperature compressive strength is small or almost no effect. However, even hot forging dies may have adverse effects when certain fatigue strength is required or when a large number of forgings are manufactured at relatively low temperatures and high strain rates. In this case, the area ratio of pores occupied in the structure is preferably 0.7% or less. It is more preferably 0.5% or less, and even more preferably 0.4% or less. It should be noted that the porosity area ratio mentioned here refers to the average porosity area ratio of the billet. For example, at 2.5 mm 2 Observe 4 holes in the left and right field of view and take the average value.

[0077] In addition, if the largest pores among the pores and thus the pores most likely to reduce fatigue strength are present on the surface, the reduction in fatigue strength may be equal to or exceed the degree assumed based on the area ratio. Therefore, the size of each pore is preferably 4000 μm 2 More preferably 3500 μm 2 the following.

[0078] <Average grain diameter>

[0079] In the hot forging die of the present invention, when the average grain diameter in the organization is 0.5 mm or more, a higher creep strength can be obtained. If it is a component that emphasizes fatigue strength, sometimes the life of the component composed of fine grains will be longer. However, in the hot forging die that mainly emphasizes high-temperature compressive strength, coarse grains are preferred from the perspective of suppressing creep deformation caused by grain boundary diffusion. In order to reliably obtain the effect of improving the life of the die, the average grain diameter is preferably 1.5 mm or more, more preferably 3.0 mm or more, and further preferably 5.0 mm or more. It should be noted that there is no particular limit to the upper limit of the average grain diameter, but in reality it is about 20 mm. In addition, from the same perspective, the shape of the grains is more preferably a columnar shape parallel to the stress direction than an equiaxed shape.

[0080] The hot forging die described above can provide a hot forging die with high high temperature compressive strength. The test temperature is 1100°C and the strain rate is 10 -3 The 0.2% compressive yield strength at 1.343 W / s is preferably 450 MPa or more.

[0081] <Method for manufacturing hot forging die>

[0082] The hot forging die blank of the present invention can be obtained by casting. It should be noted that although the blank can also be made by sintering the alloy powder, it is disadvantageous in terms of manufacturing cost and particle size control. The casting method is not particularly limited, but vacuum casting is preferred from the perspective of controlling the amount of elements.

[0083] As a manufacturing method for a hot forging mold composed of the above-mentioned structure, although methods such as accelerating the cooling rate after casting can also be considered, from the perspective of manufacturing costs and process stability, it is most preferred to implement a process of cooling after solution heat treatment of the cast billet at a temperature of 1250 to 1350°C for more than 0.5 hours. If the heat treatment temperature is low, in addition, if the holding time is short, it may lead to insufficient solid solution of the γ' phase, which may lead to insufficient uniform heat in the billet. In addition, since the homogenization of W, Mo, and Ta segregated during the casting process can also be expected, the heat treatment temperature is set to above 1250°C and the holding time is set to above 0.5 hours. On the other hand, even if the heat treatment temperature is high and the holding time is too long, not only will the effect be saturated, but the processing cost will also increase, so it is usually set to below 1350°C or within 10 hours.

[0084] In addition, by slowing down the heating rate of solution heat treatment or keeping a certain time at a temperature lower than the solution heat treatment temperature during the heating process, it is possible to expect the reduction of the amount of pores caused by the homogenization of the element that is the cause of initial melting. In addition, the cooling rate of solution heat treatment is preferably as fast as possible without applying excessive thermal stress to the blank. It should be noted that the atmosphere in solution heat treatment is not particularly limited, but is preferably inert or vacuum, and in addition, if it is under the atmospheric atmosphere, it is advisable to apply an antioxidant to the blank surface.

[0085] Furthermore, by appropriately performing machining or the like before, after, or during the above-mentioned solution heat treatment, it is possible to obtain the hot forging die of the present invention made of the cast alloy having a desired shape.

[0086] It should be noted that by performing aging heat treatment in the temperature range of 800°C to 1150°C after solution heat treatment, the high temperature strength can be further improved by controlling the size of the γ' phase. The aging heat treatment can be a single process or multiple processes. When the temperature used as a hot forging die is approximately the same as the aging heat treatment temperature, the soaking process of the hot forging die implemented before use can be made longer than usual, and it can also serve as an aging heat treatment process.

[0087] Furthermore, by applying an antioxidant to at least one of the molding surface or the side surface of the hot forging die before use, the oxidation resistance of the hot forging die can be further improved.

[0088] By the above-mentioned method for manufacturing a hot forging die, a hot forging die having high high temperature compressive strength can be obtained. The test temperature is 1100°C and the strain rate is 10 -3 The 0.2% compressive yield strength at a pressure of 1000 rpm and a stress of 1000 rpm is 450 MPa or more.

[0089] <Manufacturing method of forged products>

[0090] A typical process for producing a forged product using a hot forging die using the Ni-based alloy for hot die of the present invention will be described.

[0091] First, as the first step, the forging blank is heated to a predetermined forging temperature. The forging temperature varies depending on the material, so the temperature is adjusted appropriately. The hot forging die of the present invention has the characteristics of being able to perform isothermal forging and hot die forging even in an atmosphere at high temperatures, and is therefore suitable for hot forging of Ni-based super heat-resistant alloys, Ti alloys, and the like, which are known as difficult-to-process materials. As a representative forging temperature, it is in the range of 1000 to 1150°C.

[0092] Then, the forging blank heated in the first step is hot forged with a preheated hot forging die (second step). In the case of the hot die forging and the isothermal forging, the hot forging in the second step is preferably die forging. In addition, the hot forging die of the present invention is set to have a composition adjusted with the Cr content, etc., so that hot forging in the atmosphere at a high temperature of more than 1000°C can be performed, and as described above, it has high high temperature compressive strength, so that a long die life can be achieved.

[0093] Example

[0094] The present invention is described in more detail by the following examples. Ingots of hot mold Ni-based alloys of alloys 1, 2 and 3 shown in Table 1 are manufactured by vacuum melting. The unit is mass %. During the melting, various raw materials whose weights are adjusted in a manner to achieve the target composition are made into liquid at 1500-1600°C and then cast in a ceramic mold preheated to 800-900°C. After casting, the alloy and the mold are directly and slowly cooled to room temperature, and the alloy is separated from the mold after slow cooling. The weight of the ingot is about 10 kg, and the approximate shape of the part excluding the riser is a cube with each side of 100 mm. It should be noted that the P and O contained in the following ingots are respectively less than 0.005%. In addition, Si and Fe are respectively less than 1.0%.

[0095] It should be noted that the center portion of the ingot from which the billet for the test piece described later was cut was composed of equiaxed grains having an average size of 0.5 mm or more. Figure 1(a) shows the macrostructure of alloy 1, (b) shows the macrostructure of alloy 2, and (c) shows the macrostructure of alloy 3, and at the same time shows the position of the test piece for macrostructure observation in the ingot (the position shown by the solid line in the upper figure) and the area for measuring the average grain diameter (the area surrounded by the dotted line in the lower figure). Regarding the average grain diameter, alloy 1 is 7.0 mm, alloy 2 is 5.1 mm, and alloy 3 is 5.5 mm. The average grain diameter is defined as the circle equivalent diameter calculated by dividing the area of ​​the measurement area by the total number of grains therein (0.5 grains located on the boundary line of the area, otherwise 1 grain). The etching solution used to make the test piece is ferric chloride. It should be noted that in the present invention, although the blank is subjected to the heat treatment described later, the heat treatment will not reduce the average grain diameter.

[0096] [Table 1]

[0097] <Alloy 1>

[0098] (quality%)

[0099] W <![CDATA[M0]]> Al Cr Ta Nb Ti C B 14.6 3.0 5.6 2.1 4.1 1.2 - 0.044 0.014 C0 Mg Ca La Y Zr Hf S margin 5.9 0.0017 - - - - - 0.0004 Ni

[0100] <Alloy 2>

[0101] (quality%)

[0102] W <![CDATA[M0]]> 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 margin 10.1 0.0010 - - - - - 0.0006 Ni

[0103] <Alloy 3>

[0104] (quality%)

[0105] W <![CDATA[M0]]> Al Cr Ta Nb Ti C B 13.3 2.6 5.9 1.6 3.2 - 1.5 0.056 0.017 <![CDATA[C0]]> Mg Ca La Y Zr Hf S margin 12.3 0.0009 - - - - 0.14 0.0005 Ni

[0106] ※The symbol "-" means no addition.

[0107] ※The balance of Ni includes unavoidable impurities.

[0108] In order to observe the pores, a 10×10×8 mm cuboid was cut out from the center of the above-mentioned ingot, and a part of the cuboid was subjected to solution heat treatment to produce test pieces of the present invention example and the comparative example test pieces simulating the structure of the hot forging die. For the solution heat treatment conditions of alloy 1, the test piece kept at 1300°C for 2 hours and then air-cooled was designated as the present invention example No. 1, and the test piece kept at 1325°C for 2 hours and then air-cooled was designated as the present invention example No. 2. In addition, the test piece that was not heat-treated after cutting was designated as comparative example No. 21. In addition, for the solution heat treatment conditions of alloy 2, the test piece kept at 1300°C for 2 hours and then air-cooled was designated as the present invention example No. 3, the test piece kept at 1325°C for 2 hours and then air-cooled was designated as the present invention example No. 4, and the test piece that was not heat-treated after cutting was designated as comparative example No. 22. Regarding the solution heat treatment conditions of Alloy 3, a test piece held at 1300° C. for 2 hours and then air-cooled was designated as Inventive Example No. 5, and a test piece not subjected to heat treatment after cutting was designated as Comparative Example No. 23.

[0109] For all test pieces, the test pieces were cut so that the 10×10mm surface passing through the center of the 8mm side became the observation surface, and the cross-section was mirror-polished by polishing with diamond paste. Optical microscope photos of the mirror-polished surfaces of all test pieces were taken at 200 times, and the images were binarized and analyzed for pores using image processing software ImageJ (provided by the National Institutes of Health (NIH) of the United States). It should be noted that the photos were taken at 4 adjacent positions. In addition, the field of view area of ​​the 200-fold optical microscope photo is approximately 2.5mm 2 .

[0110] exist Figure 2 Representative examples of optical microscope photographs of each test piece of alloy 1 are shown in (a) to (c), alloy 2 is shown in (d) to (f), and alloy 3 is shown in (g) to (h). In addition, the analysis results of the area ratio of pores in each test piece of alloys 1 to 3 are shown in Table 2. Although the solution heat treatment leads to an increase in the amount of pores, it is only slightly increased in Examples No. 1 and No. 3 of the present invention, and the average values ​​at 4 locations are less than 0.4%. In Examples No. 2 and No. 4 of the present invention where the solution heat treatment temperature is higher, it increases slightly, and the average value of the 4 locations of No. 2 is 0.8%, and the average value of the 4 locations of No. 4 is 1.6%. On the other hand, in Example No. 5 of the present invention of alloy 3, the amount of pores increases even at a lower solution temperature. It should be noted that the size of the maximum pore is approximately 3400μm in Example No. 1 of the present invention. 2 , in Example No. 3 of the present invention, it is about 1100 μm 2. 4000μm in No.2, No.4 and No.5 2 above.

[0111] [Table 2]

[0112]

[0113] In addition, in order to observe the microstructure other than the pores, the inventive examples No. 1 to No. 5 were subjected to an aging heat treatment of maintaining at 1100° C. for 4 hours and then air-cooling, and then a mirror-polished surface was prepared by the same method as the above method, and then the polished surface was etched using an etching solution consisting of 50 ml of ethanol, 50 ml of 35% by mass concentrated hydrochloric acid, and 2.6 g of cupric chloride. It should be noted that since the pores are formed by initial melting at high temperature, the aging heat treatment does not change the pores.

[0114] Optical microscope photographs at 200 times magnification were taken for the etched surfaces of all test pieces. It should be noted that the photographs were taken at two adjacent locations. The optical microscope photographs were binarized and then analyzed for the eutectic γ' phase using the image processing software ImageJ.

[0115] In addition, the etched surface of all the test pieces was photographed using a scanning electron microscope (SEM) to obtain secondary electron images or reflected electron images of the microstructure and the γ' phase in the dendrite trunk. It should be noted that in the analysis of the γ' phase in the dendrite trunk, in Examples No. 1 to No. 5 of the present invention, images with a magnification of 15,000 (viewing area of ​​about 50 μm 2 ) as the object, in Comparative Examples No. 21 to No. 23, the image was magnified at 5000 times (viewing area of ​​about 500 μm 2 Then, the secondary electron image or the reflected electron image was binarized and the γ' phase in the dendrite trunk was analyzed using the image processing software ImageJ.

[0116] exist Figure 3 Representative examples of optical microscope photographs of test pieces of alloy 1 are shown in (a) to (c), alloy 2 is shown in (d) to (f), and alloy 3 is shown in (g) and (h). Figure 4 (a) to (c) show the alloy 1, and Figure 5 (d) to (f) show the alloy 2, Figure 6 (g) and (h) show secondary electron images or reflected electron images of alloy 3. Figures 4 to 6 In the images of the γ' phase analysis object, the number of particles of the γ' phase in all images is 150 or more and 1500 or less. Figure 7 Alloy 1 is shown, and Figure 8 The alloy 2 shows Fig. 9 The distribution and cumulative area ratio of the area ratio per circle equivalent diameter of the γ' phase of alloy 3 are shown. In addition, Table 3 shows the area ratio of the eutectic γ' phase of alloys 1, 2, and 3. It should be noted that in the figures and tables, the gamma prime phase is recorded as "γ' phase".

[0117] [Table 3]

[0118]

[0119] In Examples No. 1 to No. 5 of the present invention, the equivalent circular diameter of the γ' phase in the dendrite trunk is also less than 1.20 μm. In other words, the area ratio of the γ' phase with a circular equivalent diameter of less than 2.00 μm in the γ' phase with a circular equivalent diameter of less than 2.20 μm in the dendrite trunk is 100%. In addition, the area ratio of the eutectic γ' phase in Examples No. 1, No. 3 and No. 5 of the present invention is more than 1.0%, while the area ratio of the eutectic γ' phase in Examples No. 2 and No. 4 of the present invention is less than 1.0%. In addition, the area ratio of the eutectic γ' phase of Alloy 2 is lower than that of Alloy 1.

[0120] Next, a bar with a diameter of 10 mm was cut from the center of the ingot, and a portion of the bar was heat treated. Then, a blank with a diameter of 8 mm and a height of 12 mm was cut out for test piece extraction, and the surface was finished to the equivalent of No. 1000 to produce compression test pieces of the present invention example and the comparative example that simulated the organization of the hot forging die. The heat treatment conditions were the same solution heat treatment and aging heat treatment as the test pieces for microstructure observation, so these test pieces corresponded to the above-mentioned present invention examples No. 1 to No. 5 and comparative examples No. 21 to No. 23. It should be noted that two compression test pieces were made for each of the present invention examples.

[0121] The compression test conditions are test temperature 1100°C, strain rate 10 -3 / s, compression rate 10%, atmosphere is air. In addition, the high temperature compressive strength is evaluated by deriving the 0.2% compressive yield strength from the stress-strain curve obtained from the compression test. This compression test is a test to test whether the hot forging die has sufficient compressive strength at high temperatures, especially under high loads. At the test temperature of 1100°C, which is assumed to be constant temperature forging, a 0.2% compressive yield strength of 450MPa or more can be said to have sufficient strength. More preferably, it is 500MPa or more.

[0122] Table 4 shows the test results of the compression test pieces of the present invention examples No. 1 to No. 5 and the comparative examples No. 21 to No. 23. Fig.10The high temperature compressive strength of alloy 1 is shown. As can be seen from Table 4, the 0.2% compressive yield strength of Examples No. 1, No. 2 and No. 4 of the present invention is all above 450 MPa, and has excellent high temperature compressive strength. The average value of the compression results of Example No. 5 of the present invention after two implementations is also above 450 MPa. Example No. 3 of the present invention is below 450 MPa, so it may not be suitable for use at high temperatures such as 1100°C, but its strength is higher than that of Comparative Example No. 22, and the mold life is longer than this, and sufficient mold life can be achieved. In addition, as shown in Tables 2 and 3, the area ratio of pores and eutectic γ' phase is lower than that of other examples of the present invention, so it is suitable for uses where high loads are repeatedly applied at relatively low temperatures such as 1050°C.

[0123] [Table 4]

[0124]

[0125] The above results show that the hot forging die of the present invention has high high temperature compressive strength and can achieve a long die life. It is known that the hot forging die of the present invention having the above characteristics is suitable for hot die forging and isothermal forging.

Claims

1. A hot forging die, characterized in that: It is composed of a casting alloy having the following composition: in mass %, W: 7.5-20.0%, Mo: 0-5.0%, Al: 5.0-7.5%, Cr: 0.5-5.0%, Ta: 1.0-12.0%, C: 0.01-0.15%, B: less than 0.03%, S: less than 0.015%, one or more selected from rare earth elements, Y, Ca, Mg, the total of which is 0-0.020%, one or two selected from Zr and Hf elements, the total of which is less than 0.5%, Ti: less than 5.0%, Nb: less than 5.0%, Co: less than 25.0%, the balance is Ni and unavoidable impurities, and the circle equivalent diameter of the γ' phase in the dendrite trunk is less than 2.20μm.

2. The hot forging die according to claim 1, characterized in that: It is composed of a casting alloy having the following composition: in mass %, W: 10.0-20.0%, Mo: 0.5-5.0%, Al: 5.0-7.5%, Cr: 0.5-4.0%, Ta: 1.0-12.0%, C: 0.01-0.15%, B: less than 0.03%, S: less than 0.015%, one or more selected from rare earth elements, Y, Ca, Mg, the total of 0-0.020%, one or two selected from Zr and Hf elements, the total of 0.5% or less, Ti: less than 5.0%, Nb: less than 5.0%, Co: less than 20.0%, the balance is Ni and unavoidable impurities, and the circle equivalent diameter of the γ' phase in the dendrite trunk is less than 2.20μm.

3. The hot forging die according to claim 1 or 2, wherein: The area ratio of the γ′ phase having an equivalent circle diameter of 2.00 μm or less in the γ′ phase having an equivalent circle diameter of 2.20 μm or less in the dendrite trunk is 90% or more.

4. The hot forging die according to claim 1 or 2, characterized in that: The area ratio of the eutectic γ' phase is 4.0% or less.

5. The hot forging die according to claim 1 or 2, characterized in that: The pore area ratio is 0.7% or less.

6. The hot forging die according to claim 5, characterized in that: The size of each pore is 4000 μm 2 the following.

7. The hot forging die according to claim 1 or 2, characterized in that: The average grain diameter is 0.5 mm or more.

8. The hot forging die according to claim 1 or 2, characterized in that: Test temperature 1100℃, strain rate 10 -3 The 0.2% compressive yield strength at a pressure of 1000 rpm and a stress of 1000 rpm is 450 MPa or more.

9. A method for manufacturing a hot forging die, characterized in that: The casting alloy having the component composition according to claim 1 or 2 is subjected to a solution heat treatment at a temperature of 1250 to 1350° C. for 0.5 hours or more.

10. The method for manufacturing a hot forging die according to claim 9, characterized in that: After the solution heat treatment, an aging heat treatment is further performed in a temperature range of 800°C to 1150°C.

Images