Method for producing hot work tool steel, and hot work tool steel

Abstract

Provided is a method for producing a hot work tool steel having a fine annealed structure even if the cross-sectional dimension is large. A method for producing a hot work tool steel and a hot work tool steel, the method for producing a hot work tool steel having a cross-sectional area of 200000 mm2 or more, the method comprising: a cogging forging step in which a steel block is heated to 1100-1250 DEG C and then hot-forged to obtain a steel sheet; a first finish forging step for heating the steel sheet to 850-1020 DEG C and performing upsetting forging with an upsetting forming ratio of 1.6 or more to obtain an intermediate forged material; and a second finish forging step for obtaining a forged material by subjecting the intermediate forged material to solid forging at a solid exercise molding ratio of 2.0-4.0.

Description

Technical Field

[0001] The invention relates to a manufacturing method of hot working tool steel and the hot working tool steel. Background Art

[0002] Since hot working tools are used while in contact with high-temperature or hard workpieces, they need to have toughness that can withstand impacts. In addition, as a material suitable for hot working tools (hot working tool steel), for example, SKD61 series alloy tool steel, which is a Japanese Industrial Standards (JIS) steel grade, is known. This type of hot working tool steel is usually manufactured by using a raw material consisting of a steel block or a steel sheet obtained by blanking a steel block as a starting material, subjecting it to various hot working or heat treatments to make a specified forging material, and annealing the forging material. Then, the manufactured hot working tool steel is transferred to a process of machining into the desired hot working tool shape in a low-hardness annealed state.

[0003] Hot work tool steel that is machined into the shape of a hot work tool is generally finished by adjusting the hardness to a specified use by quenching and tempering. Quenching is an operation that heats the hot work tool material to the austenite temperature range and quenches it to cause the structure to undergo a martensitic phase transformation. Therefore, the composition of the hot work tool material can be adjusted to a martensitic structure by quenching.

[0004] In order to improve the properties of the hot work tool steel, various studies have been conducted on the manufacturing method of the hot work tool steel. For example, Patent Document 1 discloses a manufacturing method of the hot work tool steel, which is characterized in that, in order to obtain excellent toughness, a series of hot forging processes are performed in at least one of three directions orthogonal to an ingot of the hot work tool steel containing 3% to 6% of Cr by mass to introduce strain, and then the ingot is subjected to soaking at a temperature of 1200° C. to 1300° C. for more than 6 hours.

[0005] It is known that the toughness of hot working tools can be improved by making the martensitic structure fine. Specifically, the original austenite grain size confirmed in the martensitic structure is made fine. Moreover, as a method for making the original austenite grain size fine, it is effective to operate the annealing structure at the time point of the hot working tool material before the quenching. For example, the applicant of the present application proposed the following method in Patent Document 2: the annealing structure is set to "the following annealing structure: for the ferrite grains in the cross section, it has the following particle size distribution, that is, in the cumulative distribution of the oversize based on the cross-sectional area of ​​the ferrite grains, the particle size when the cumulative cross-sectional area is 90% of the total cross-sectional area is less than 25μm in terms of circle equivalent diameter".

[0006] Prior art literature

[0007] Patent Literature

[0008] Patent Document 1: Japanese Patent Application Publication No. 2007-100194

[0009] Patent Document 2: International Publication No. 2015 / 182586 Summary of the invention

[0010] Problems to be solved by the invention

[0011] Patent Document 2 specifies an annealing structure for obtaining a hot working tool steel having excellent toughness. However, in order to obtain a corresponding structure, a solid forging with a processing ratio (cross-sectional area ratio) of 5s or more is required, and the size of the forged material obtained by forging is limited by the size of the steel block. Therefore, for a forged material with a large cross-sectional size, it is difficult to obtain a fine annealing structure if the processing ratio cannot be increased.

[0012] As a method for producing a forged material having a cross-sectional area larger than that of a steel block, upsetting forging described in JIS-G-7101 is known. However, upsetting is usually performed at a high temperature exceeding 1200°C, and since grain growth occurs immediately, it is difficult to obtain a fine structure. There is no description of the issues during such forging in Patent Documents 1 and 2, leaving room for research.

[0013] Therefore, an object of the present invention is to provide a method for producing a hot working tool steel having a fine annealed structure even when the cross-sectional size is large.

[0014] Technical means of solving problems

[0015] The present inventors have found that conditions for obtaining a fine annealing structure can be obtained by optimizing the timing, temperature, and forging ratio of the above-mentioned upset forging in order to obtain a hot working tool steel having a large cross-sectional area and a fine structure, thereby arriving at the present invention.

[0016] That is, one embodiment of the present invention is a method for manufacturing hot working tool steel to obtain a cross-sectional area of ​​200000mm 2 The above hot working tool steel, the manufacturing method of the hot working tool steel includes: a blanking forging process, heating the steel block to 1100℃~1250℃, and then hot forging to obtain a steel sheet; a first finishing forging process, heating the steel sheet to 850℃~1020℃, and performing upsetting forging with an upsetting forming ratio of 1.6 or more to obtain an intermediate forging material; and a second finishing forging process, performing solid forging with a solid forging forming ratio of 2.0~4.0 on the intermediate forging material to obtain a forging material.

[0017] Preferably, a soaking step of heating the steel block at 1250° C. or higher is provided before the billet forging step.

[0018] Another embodiment of the present invention is a hot working tool steel having a cross-sectional area of ​​200000mm 2 The ferrite grains in the microstructure cross section have the following particle size distribution: in the cumulative distribution of the oversize material based on the cross-sectional area of ​​the ferrite grains, the particle size when the cumulative cross-sectional area is 90% of the total cross-sectional area is less than 25 μm in terms of circle equivalent diameter.

[0019] Effects of the Invention

[0020] According to the present invention, it is possible to provide a hot working tool steel having a fine annealing structure even if the cross-sectional size is large. BRIEF DESCRIPTION OF THE DRAWINGS

[0021] [ Figure 1 ] is a schematic diagram for illustrating the manufacturing method of the present invention.

[0022] [ Figure 2 ] is an optical microscope photograph of the cross-sectional structure of the hot working tool steel (sample No. 1) of the present invention.

[0023] [ Figure 3 ] is a grain boundary map of the hot work tool steel (sample No. 1) of the present invention example obtained by electron beam backscatter diffraction (EBSD).

[0024] [ Figure 4 ] is an optical microscope photograph of the cross-sectional structure of the hot working tool steel (sample No. 2) of the present invention.

[0025] [ Figure 5 ] is a grain boundary map of the hot work tool steel (sample No. 2) of the present invention example obtained by electron beam backscatter diffraction (EBSD).

[0026] [ Figure 6 ] is an optical microscope photograph of the cross-sectional structure of the hot working tool steel (sample No. 3) of the present invention.

[0027] [ Figure 7 ] is a grain boundary map of the hot work tool steel (sample No. 3) of the present invention example obtained by electron beam backscatter diffraction (EBSD).

[0028] [ Figure 8 ] is an optical microscope photograph of the cross-sectional structure of the hot work tool steel (sample No. 11) of the comparative example.

[0029] [ Fig. 9 ] is a grain boundary map of the hot work tool steel (sample No. 11) of the comparative example obtained by electron beam backscatter diffraction (EBSD).

[0030] [ Fig.10 ] is a graph showing the particle size distribution of ferrite grains distributed in the cross-sectional structure of the hot work tool steels (sample No. 1, sample No. 2, sample No. 3, sample No. 11) according to the examples of the present invention and the comparative examples. DETAILED DESCRIPTION

[0031] The present invention is characterized in that, for an intermediate forging material obtained by subjecting a steel sheet obtained by billet forging to upset forging at a predetermined temperature, a finish forging of a solid forging ratio is performed in the next step, thereby obtaining a forging material with a fine ferrite grain size. The constituent elements of the present invention are described below.

[0032] The hot work tool steel obtained by the production method of the present invention has an annealed structure and is used after quenching and tempering, and has a component composition that can be adjusted to a martensitic structure by the quenching.

[0033] As long as the raw material exhibits a martensitic structure after quenching and tempering, the effect of refining the hot work tool steel structure obtained by the manufacturing method of the present invention can be achieved by implementing the manufacturing method of the present invention described later. Therefore, in order to achieve the above-mentioned effect of the present invention, it is not necessary to determine the component composition of the hot work tool steel. However, on the basis of establishing the absolute mechanical properties of the hot work tool, for example, the hot work tool steel of the manufacturing method of the present invention preferably has a component composition containing C: 0.3% to 0.5%, Cr: 3.0% to 6.0% in mass%, and preferably also has a component composition containing V: 0.1% to 1.5%. In addition, the hot work tool steel of the manufacturing method of the present invention is preferably applied to steel having a component composition of C: 0.3% to 0.5%, Si: 2.0% or less, Mn: 1.5% or less, P: 0.05% or less, S: 0.05% or less, Cr: 3.0% to 6.0%, Mo and W alone or in combination (Mo+1 / 2W): 0.5% to 3.5%, V: 0.1% to 1.5%, and the remainder is Fe and impurities. The reasons for limiting the above component ranges are described below.

[0034] C: 0.3 mass % to 0.5 mass % (hereinafter simply expressed as "%")

[0035] C is a basic element of hot working tool materials, which partially dissolves in the matrix to impart strength, and partially forms carbides to improve wear resistance or sintering resistance. In addition, when C and Cr, which are dissolved as interstitial atoms, are added together with substitutional atoms with high C affinity, the I (interstitial atom)-S (substitutional atom) effect is also expected; it acts as a drag force of solute atoms to increase the strength of hot working tools. However, excessive addition will lead to a decrease in toughness or hot strength. Therefore, it is preferably set to 0.3% to 0.5%.

[0036] Si: 2.0% or less

[0037] Si is a deoxidizing agent during steelmaking, but if it is too much, ferrite will be generated in the tool structure after quenching and tempering. Therefore, it is preferably set to 2.0% or less. On the other hand, Si has the effect of improving the machinability of the material. In order to obtain the above effect, less than 0.2% can be added, but it is preferably added 0.2% or more.

[0038] Mn: 1.5% or less

[0039] If Mn is too much, the viscosity of the matrix will increase and the machinability of the material will decrease. Therefore, it is preferably set to 1.5% or less. On the other hand, Mn has the effect of improving hardenability, inhibiting the formation of ferrite in the tool structure, and obtaining a moderate quenching and tempering hardness. In addition, by existing in the form of MnS as a non-metallic inclusion, it has a great effect on improving machinability. In order to obtain these effects, less than 0.1% can be added, but it is preferably added 0.1% or more.

[0040] P: 0.05% or less

[0041] P is an impurity element that is inevitably contained in various hot work tool materials. Moreover, it is an element that segregates at the prior austenite grain boundaries during heat treatment such as tempering and embrittles the grain boundaries. Therefore, in order to improve the toughness of the hot work tool, it is preferably limited to 0.05% or less.

[0042] S: 0.05% or less

[0043] S is an impurity element that is inevitably contained in various hot work tool materials. Moreover, it is an element that deteriorates the toughness of the material. Therefore, in order to improve the toughness of the hot work tool, it is preferably limited to 0.05% or less. On the other hand, S has the effect of improving machinability by bonding with the Mn and existing in the form of MnS as a non-metallic inclusion. In order to obtain the above effect, less than 0.03% may be added, but it is preferably added 0.03% or more.

[0044] Cr: 3.0%~6.0%

[0045] Cr is an element that improves hardenability and forms carbides, and is effective in strengthening the matrix or improving wear resistance. Moreover, it is a basic element of hot working tool materials that also contributes to improving temper softening resistance and high-temperature strength. However, excessive addition leads to a decrease in hardenability or high-temperature strength. Therefore, it is preferably set to 3.0% to 6.0%. Moreover, it is more preferably 5.0% or less. In the present invention, since the effect of improving toughness due to the refinement of the martensitic structure is obtained, the Cr of this effect portion can be reduced. In this case, for example, by setting Cr to 5.0% or less, the above-mentioned high-temperature strength can be further improved.

[0046] Mo and W alone or in combination (Mo+1 / 2W): 0.5% to 3.5%

[0047] Mo and W can be added alone or in combination to precipitate or agglomerate fine carbides through tempering to impart strength and improve softening resistance. As for the amount of addition at this time, since the atomic weight of W is about twice that of Mo, it can be specified together with the Mo equivalent of (Mo+1 / 2W) (of course, only one of them can be added, or both can be added at the same time). In addition, in order to obtain the above effect, it is preferred to add more than 0.5% calculated as the value of (Mo+1 / 2W). However, if it is too much, it will lead to a decrease in machinability or toughness, so it is preferably calculated as 3.5% or less as the value of (Mo+1 / 2W).

[0048] V: 0.1%~1.5%

[0049] V has the effect of forming carbides to strengthen the matrix or improve wear resistance and temper softening resistance. Moreover, the carbides distributed in the annealed structure act as pinning particles that inhibit the coarsening of austenite grains during quenching heating, which helps to improve toughness. In order to obtain these effects, it is preferably added at 0.1% or more. Moreover, in the present invention, it is preferred to add V in terms of further promoting the refinement of the martensitic structure. However, if it is too much, it will lead to a decrease in machinability or toughness caused by the increase of carbides themselves, so it is preferably set to 1.5% or less.

[0050] Furthermore, in addition to the above-mentioned element species, the following element species may be contained.

[0051] Ni: 0%~1.0%

[0052] Ni is an element that increases the viscosity of the matrix and reduces machinability. Therefore, it is preferably set to 1.0% or less. On the other hand, Ni is an element that suppresses the formation of ferrite in the tool structure. In addition, it is an element that is effective in the following aspects: together with C, Cr, Mn, Mo, W, etc., it imparts excellent hardenability to the tool material, and forms a martensite-based structure even when the cooling rate during quenching is slow, thereby preventing the toughness from decreasing. Furthermore, since it also improves the intrinsic toughness of the matrix, it can also be added as needed in the present invention. In the case of addition, less than 0.1% can be added, but it is preferably added at 0.1% or more.

[0053] Co: 0%~1.0%

[0054] Co reduces toughness, so it is preferably set to 1.0% or less. On the other hand, when using hot working tools, Co forms an extremely dense and well-adhesive protective oxide film on the surface when the temperature rises. The oxide film prevents metal contact with the workpiece, suppresses the temperature rise of the tool surface, and brings excellent wear resistance. Therefore, Co can be added as needed. In the case of addition, it can be less than 0.3%, but preferably 0.3% or more.

[0055] Nb: 0%~0.3%

[0056] Nb is preferably set to 0.3% or less because it causes a decrease in machinability or toughness. On the other hand, Nb has the effect of forming carbides, strengthening the matrix or improving wear resistance. In addition, it has the effect of increasing temper softening resistance and, like V, suppressing the coarsening of grains and contributing to the improvement of toughness. Therefore, Nb can be added as needed. When added, it can be less than 0.01%, but preferably 0.01% or more.

[0057] The main elements that may remain in the steel as inevitable impurities are Cu, Al, Ca, Mg, O (oxygen), N (nitrogen), etc. In the present invention, these elements are preferably as low as possible. However, on the other hand, in order to obtain additional effects such as morphological control of inclusions, or other mechanical properties, and improvement of manufacturing efficiency, a small amount may be contained. In this case, as long as it is in the range of Cu≦0.25%, Al≦0.025%, Ca≦0.01%, Mg≦0.01%, O≦0.01%, and N≦0.03%, it can be fully allowed and is the preferred upper limit of the present invention.

[0058] Next, the manufacturing method of the present invention is described. Hot working tool steel having an annealed structure is usually made of a raw material consisting of a steel block or a steel sheet obtained by blanking a steel block as a starting material, which is subjected to various hot working or heat treatments to produce a predetermined forging material, and the forging material is annealed and finished into a desired shape such as a block shape. Figure 1 Schematic diagram for explaining the manufacturing method of the present invention is shown in FIG. (In the schematic diagram, the surface of the steel sheet or forging material with "shading" is a surface common to each process.) In addition, the present invention is suitable for a cross-sectional area of ​​200,000 mm 2 The above large hot working tool steel manufacturing method. The more preferred cross-sectional area is 300000mm 2 In addition, the aspect ratio of the cross section of the hot work tool steel is preferably 2 to 3. Here, the "cross-sectional area" in the present invention can be the area of ​​the cross section perpendicular to the longitudinal direction of the forging material.

[0059] <Blanking and Forging Process>

[0060] First, in the present invention, a blank forging process is implemented in which a steel block obtained by casting is heated to 1100°C to 1250°C and then hot forged to obtain a steel sheet. Generally, with respect to a steel block obtained by casting, the larger the size of the steel block, the slower the solidification rate and the larger the coarse dendrites grow. In addition, there are a large number of nests caused by solidification shrinkage in the final solidification portion. In order to destroy the dendritic structure and press the nests, the blank forging of the steel block is preferably carried out at a high temperature, and therefore the heating temperature in the blank forging process of the present invention is set to 1100°C to 1250°C. In addition, depending on the cross-sectional area of ​​the steel block and the cross-sectional area of ​​the hot working tool steel to be finally obtained, upsetting forging may be performed in the blank forging process.

[0061] In the present invention, it is preferred to perform annealing after the blanking forging process. Through this process, residual stress can be removed from the steel sheet after hot plastic working. The annealing temperature at this time is preferably set to 600°C or more. It is further preferred that the annealing temperature can be set to above the austenite transformation point and below 900°C. By transforming from a ferrite structure to an austenite structure at one time, the effect of adjusting the crystal grain size can be expected.

[0062] <First Finish Forging Process>

[0063] Next, the steel sheet obtained by the blank forging is heated to 850°C to 1020°C and subjected to a first finishing forging process of upsetting with an upsetting ratio of 1.6 or more to obtain an intermediate forging material. Figure 1The ratio of the length L1 of the steel sheet shown to the length L2 of the steel sheet after upsetting can be calculated by the formula 1 / (L2 / L1). By performing the first fine forging within the heating temperature range, grain growth is suppressed, and the nucleation sites during phase transformation during the heat treatment of the next step are increased by recrystallization or accumulated strain, so there is a tendency to refine the structure. The preferred upper limit of the heating temperature in the first fine forging step is 1000°C. The more preferred upper limit is 990°C. The subsequent fine forging steps are preferably carried out below the temperature at which the pinning carbides are fully present. The lower limit temperature is preferably a temperature above the austenite transformation point. In the case of SKD61, which is a representative steel type of hot working tool steel, the austenite transformation point is about 850°C. Therefore, in the present invention, the lower limit of the heating temperature in the first fine forging step is set to 850°C.

[0064] The upsetting forming ratio in the first finishing forging process is set to be greater than 1.6. If the upsetting forming ratio is small, the cross-sectional area of ​​the intermediate forging material after upsetting cannot be fully obtained. Therefore, in the second finishing forging process described later, when the solid forging forming ratio is set to 2.0 to 4.0, it is difficult to obtain a forging material with a large cross-sectional size. Although there is no upper limit for the upsetting forming ratio, if the upsetting forming ratio is increased and the cross-sectional area increases excessively, the load required for forging increases and exceeds the pressure of the forging machine. Therefore, it can be set to about 2.0, for example. After upsetting forging, the intermediate forging material having the cross-sectional area required to obtain a solid forging forming ratio of 2.0 to 4.0 is formed into a shape according to the size of the forging material obtained in the second finishing process.

[0065] <Second Finish Forging Process>

[0066] In the manufacturing method of the present invention, a second finishing forging step is performed in which the intermediate forging material obtained through the first finishing forging step is subjected to solid forging at a solid forging forming ratio of 2.0 to 4.0 to obtain a forging material. The solid forging forming ratio is Figure 1 The ratio of the cross-sectional area A3 of the intermediate forging material shown and the cross-sectional area A4 after the second fine forging process can be obtained by the formula A3 / A4. The second fine forging can be started immediately after the first fine forging process is completed, but it can also be carried out after being put into a heating furnace for reheating and maintaining. In addition, the temperature for implementing the second fine forging is not a problem even if it is a temperature different from the temperature of the first fine forging. It is preferably carried out at the same temperature as the first fine forging temperature or at a lower temperature.

[0067] The manufacturing method of the hot working tool steel of the present invention preferably anneals the forging material after the second fine forging process at a temperature above the austenite transformation point. Through this process, the residual stress is removed from the forging material after hot plastic processing, and the hardness is sufficiently reduced, thereby making it easy to perform mechanical processing on the shape of the hot working tool such as the mold described later. Moreover, the annealing temperature at this time is preferably set to be above the austenite transformation point and below 900°C. If the annealing temperature is too high, for example, the carbides in the forging material structure will be dissolved into the matrix, and the carbides will be preferentially re-precipitated at the grain boundaries during the cooling process from the subsequent annealing temperature, which may have an adverse effect on the distribution of carbides in the final product structure.

[0068] In the manufacturing method of the present invention, it is preferred to perform a soaking process of heating the steel block at 1250°C or above before the blank forging process. Usually, when the steel solidifies, the components of the initial solidified portion become thinner, and the components are concentrated in the most solidified portion to form coarse eutectic carbides. Furthermore, the larger the size of the steel block, the slower the cooling rate of solidification, the coarser the eutectic carbides, and the worse the segregation of the components. By heating at 1250°C or above, it is expected that the eutectic carbides will be dissolved in the matrix and the components will be diffused. Although the upper limit temperature is not specifically set, if the temperature is too high, part of the steel block will dissolve to form a liquid phase, and the shape of the steel block cannot be maintained. Therefore, it is preferably set for each component of the material.

[0069] In addition, the same effect can be expected in the process of heating and holding the steel block for the above-mentioned blank forging process, so the soaking process can also serve as the above-mentioned blank forging process. However, if the temperature is low, the effect of solid solution or segregation diffusion of eutectic carbides is small, so it is preferably carried out at a heating temperature above the blank forging process.

[0070] The hot working tool steel (forging material) of the present invention obtained by the manufacturing method of the present invention has a cross section perpendicular to the longitudinal direction with an area of ​​200000 mm 2 Above (preferably 300000mm 2The ferrite grains in the cross section of the annealed structure have a large size of more than 100 μm, and the ferrite grains in the cross section of the annealed structure have the following particle size distribution: in the cumulative distribution of the oversize based on the cross-sectional area of ​​the ferrite grains, the particle size (D90) when the cumulative cross-sectional area is 90% of the total cross-sectional area is less than 25 μm in terms of the circle equivalent diameter. As a result, a fine structure with a prior austenite grain size of, for example, No. 9.0 or more after quenching and tempering can be stably obtained. The preferred upper limit of D90 is 22 μm, and the more preferred upper limit of D90 is 20 μm. Here, the ferrite grains of the present invention can be sampled from the central part of the cross section of the forging material orthogonal to the long side direction, and the 400 μm×400 μm area of ​​the surface of the sample collected parallel to the long side direction (forging direction) is measured. The reason is that in the case of a forging material with a large cross-sectional area, the grains are easily coarsened in the central part of its cross section. In addition, if the grain refinement caused by recrystallization or phase transformation is insufficient, coarse grains extending in the forging direction may remain, which can be confirmed by observing the parallel surface. Then, the measurement is performed on two or more sections of the forging material. Furthermore, the two or more sections are preferably sections including the two ends of the forging material after the ends are cut off. In terms of being unsuitable for products, the ends are useless parts that are usually cut off and removed. In addition, if the annealed structure of the present invention is a structure containing carbides, the carbides act as pinning particles that inhibit the coarsening of austenite grains during quenching and heating, further inhibiting the growth of austenite crystals. Moreover, these carbides exist in the form of undissolved carbides in the hot working tool after quenching and tempering in the structure after the original austenite grain size is refined. Here, the so-called "annealed structure" in the present invention refers to a structure obtained by annealing treatment, preferably a structure whose hardness is softened to about 150 Brinell hardness (Brinell Hardness, HBW) to 230 HBW in terms of Brining hardness, for example.

[0071] The hot work tool steel obtained by the manufacturing method of the present invention is prepared into a martensitic structure having a predetermined hardness by quenching and tempering, and adjusted to a product of a hot work tool. Moreover, the hot work tool steel is adjusted to the shape of a hot work tool by various mechanical processes such as cutting or punching. The timing of the processing is preferably carried out in a state where the hardness of the material before quenching and tempering is low (annealing state). In this case, finishing can also be performed after the quenching and tempering. In addition, depending on the situation, processing can also be performed in a pre-hardened state after the quenching and tempering.

[0072] The quenching and tempering temperatures vary depending on the composition of the raw materials or the target hardness, but preferably the quenching temperature is about 1000°C to 1100°C, and the tempering temperature is about 500°C to 650°C. If the quenching temperature is high, the pinning carbides will dissolve in the matrix and cause grain growth, so it is preferably set to a temperature at which the pinning carbides remain sufficiently. For example, in the case of SKD61, which is a representative steel grade for hot working tool steel, the quenching temperature is about 1000°C to 1030°C, and the tempering temperature is about 550°C to 650°C. The quenching and tempering hardness is preferably set to 50 Rockwell hardness (HRC) or less. More preferably, it is 48HRC or less.

[0073] Example

[0074] A melting furnace was used to melt a 10 t steel block having the composition shown in No. 1 to No. 3 of Table 1. The steel block was subjected to soaking at a temperature of 1250° C. or higher, then subjected to blanking at 1160° C., and then subjected to annealing at 870° C. to produce a steel sheet having a thickness of 820 mm×a width of 820 mm×a length of 1900 mm. The forged material obtained by subjecting the steel sheet to the steps shown in No. 1 to No. 3 of Table 2 was annealed at 870° C. to obtain a hot work tool steel as an example of the present invention.

[0075] In addition, a melting furnace was used to melt a 10 t steel block having the component composition shown in No. 11 of Table 1. The steel block was subjected to soaking at a temperature of 1250° C. or higher, then subjected to blanking at 1160° C., and then subjected to annealing at 870° C. to produce a steel sheet having a thickness of 700 mm×a width of 1090 mm×a length of 1680 mm. The forged material obtained by subjecting the steel sheet to the process shown in No. 11 of Table 2 was annealed at 870° C. to obtain a hot work tool steel as a comparative example.

[0076] [Table 1]

[0077] (mass%)

[0078] No. C Si Mn P S Cr Mo V <![CDATA[Fe ※ ]]> 1 0.37 0.45 0.64 ＜0.01 ＜0.01 4.1 2.3 0.72 Bal. 2 0.37 0.45 0.64 ＜0.01 ＜0.01 4.1 2.3 0.72 Bal. 3 0.37 0.45 0.64 ＜0.01 ＜0.01 4.1 2.3 0.72 Bal. 11 0.37 0.47 0.65 ＜0.01 ＜0.01 4.1 2.3 0.74 Bal.

[0079] ※Contains inevitable impurities

[0080] [Table 2]

[0081]

[0082] Next, the cross-sectional structures of the manufactured hot working tool steels (samples) No. 1 to hot working tool steels (samples) No. 3 and hot working tool steel (sample) No. 11 were observed. The cross section for the structural observation was set as the cross section obtained by cutting off the useless parts from the two ends of the steel material. Then, two plate-shaped test pieces were collected by slicing at a position of about 15 mm from the two end faces of the sample in a manner including the obtained cross section. A test piece of about 10 mm to 20 mm was cut out from the center of the plate-shaped test piece, and the surface parallel to the forging direction (i.e., the length direction of the sample) was observed. For the observation surface, the structure was etched with a picric acid alcohol solution, and then observed with an optical microscope (magnification 200 times), and the observed cross-sectional area was set to 0.16 mm 2 (400 μm×400 μm) As a result of the observation, the cross-sectional structures of Hot Work Tool Steel No. 1 and Hot Work Tool Steel No. 11 were almost entirely occupied by the ferrite phase, and the ferrite grains occupied more than 99% of the area of ​​the observed cross-sectional surface.

[0083] Next, the cross-sectional structures of Samples No. 1 to No. 3 and Sample No. 11 were observed to confirm the distribution of ferrite grains in the structures. 2 The cross-sectional structure of the sample was analyzed by analyzing the electron backscatter diffraction (EBSD) pattern at a magnification of 200 times to obtain a grain boundary map divided by large-angle grain boundaries with an azimuth difference of more than 15°. The EBSD pattern was analyzed using an EBSD device (measurement interval 1.0 μm) attached to a scanning electron microscope (Carl Zeiss ULTRA 55). The optical microscope photograph of one end face of each sample as the measurement result and the grain boundary map are shown in Figure 1. Figure 2 to Figure 5 In addition, according to the above method, the grain size of each grain is read from the grain boundary diagram, converted into a circle equivalent diameter (area circle equivalent diameter), and the grain size distribution of the ferrite grains based on the circle equivalent diameter is confirmed. The grain size distribution of one end surface of Sample No. 1 and Sample No. 11 is shown in Figure 6 In addition, Figure 6In the graph, the vertical axis is the cumulative cross-sectional area (%) of the grains, and the horizontal axis is the equivalent circular diameter of the grains. According to the measurement results, it can be confirmed that the cumulative cross-sectional area of ​​the hot working tool steels No. 1, No. 2, and No. 3 manufactured in the present invention is 18 μm, 18 μm, and 19 μm at the one end surface, respectively, and 18 μm, 17 μm, and 15 μm at the other end surface. On the other hand, the cumulative cross-sectional area of ​​the hot working tool steel No. 11 obtained by the manufacturing method of the comparative example is 30 μm at the one end surface, and 24 μm at the other end surface. From the above, it can be confirmed that the hot working tool steel of the present invention example has a finer structure than the comparative example.

[0084] Description of Figure Numbers

[0085] 1: Steel sheet after forging

[0086] 2: Intermediate forging material after upsetting in the first fine forging

[0087] 3: Intermediate forging material after the first fine forging

[0088] 4: Forging material after the second fine forging

Claims

1. A method for manufacturing hot working tool steel to obtain a cross-sectional area of ​​200000mm 2 The hot working tool steel described above, wherein the manufacturing method of the hot working tool steel comprises: a blank forging step, wherein a steel block is heated to 1100°C to 1250°C and then hot forged to obtain a steel sheet; a first finishing forging step, wherein the steel sheet is heated to 850°C to 1020°C and then upsetting forged with an upsetting forming ratio of 1.6 or more to obtain an intermediate forging material; and a second finishing forging step, wherein the intermediate forging material is subjected to solid forging with a solid forging forming ratio of 2.0 to 4.0 to obtain a cross section orthogonal to the long side direction with an area of ​​200,000 mm 2 Forging materials above. 2 . The method for producing hot work tool steel according to claim 1 , further comprising a soaking step of heating the steel block at 1250° C. or higher before the billet forging step.

3. A hot working tool steel, the area of ​​the cross section perpendicular to the long side direction is 200000mm 2 As described above, the ferrite grains in the cross section have the following particle size distribution: in the cumulative distribution of the oversize based on the cross-sectional area of ​​the ferrite grains, the particle size when the cumulative cross-sectional area is 90% of the total cross-sectional area is 25 μm or less in terms of circle equivalent diameter.

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