EP4592003A1

EP4592003A1 - Method for producing hot work tool steel, and hot work tool steel

Info

Publication number: EP4592003A1
Application number: EP23868056.5A
Authority: EP
Inventors: Yousuke Nakano; Yuya KOGA; Kouta Kataoka
Original assignee: Proterial Ltd
Current assignee: Proterial Ltd
Priority date: 2022-09-21
Filing date: 2023-09-07
Publication date: 2025-07-30
Also published as: WO2024062933A1; JPWO2024062933A1; CN119907721A; KR20250054085A

Abstract

Provided is a method for producing hot work tool steel that has a fine annealed structure even if a cross-sectional size is large. This method for producing hot work tool steel having a cross-sectional area of at least 200,000 mm² includes: a cogging step for heating a steel ingot to 1100 to 1250°C and then performing hot forging to obtain a semi-finished product; a first finishing forging step for heating the semi-finished product to 850 to 1020°C, and performing upset forging with an upset forging forming ratio of 1.6 or more to obtain an intermediate forged material; and a second finishing forging step for subjecting the intermediate forged material to solid forging with a solid forging forming ratio of 2.0 to 4.0 to obtain a forged material.

Description

Technical Field

The present invention relates to a method for producing hot work tool steel and hot work tool steel.

Description of Related Art

Hot work tools, which are used while in contact with high-temperature workpieces or hard workpieces, need to possess toughness that can withstand impact. And as a material applied to hot work tools (hot work tool steel), for example, alloy tool steel of SKD61 series, which is a JIS steel grade, is known. Such hot work tool steel is normally produced by using a steel ingot or a semi-finished product obtained by a blooming process of a steel ingot as a starting material, performing various hot working and heat treatment thereon to make a predetermined forged material, and performing annealing treatment on the forged material. And the produced hot work tool steel proceeds to a process of machining into a desired hot work tool shape in an annealed state with low hardness.
Hot work tool steel that has been machined into the shape of a hot work tool is generally subjected to finishing machining after being adjusted to a predetermined use hardness by quenching and tempering. Quenching is an operation in which the hot work tool material is heated to the austenite temperature range and then rapidly cooled to transform the structure into martensite. Therefore, the component composition of the hot work tool material is designed to be adjustable to a martensite structure by quenching.
In order to improve the properties of the aforementioned hot work tool steel, various studies have been conducted on the method for producing hot work tool steel. For example, Patent Literature 1 discloses a method for producing hot work tool steel characterized by introducing strain by performing a series of hot forging on a steel ingot of hot work tool steel containing, in mass%, 3 to 6% of Cr in at least one direction of three orthogonal directions, and then conducting soaking for 6 hours or more under temperature conditions of 1200 to 1300°C to obtain excellent toughness.
By the way, it is known that the toughness of a hot work tool can be improved by refining the martensite structure. Specifically, this means refining the prior austenite grain size that can be observed in the martensite structure. And, as a method to refine the prior austenite grain size, it is effective to manipulate the annealed structure at the stage of the hot work tool material before quenching, and for example, the applicant of the present application has proposed in Patent Literature 2 a method of making the annealed structure "an annealed structure in which ferrite grains in a cross-section have a grain size distribution in which a grain size at a cumulative cross-sectional area of 90% of a total cross-sectional area in an oversize cumulative distribution based on a cross-sectional area of the ferrite grains is 25 µm or less in an equivalent circular diameter".

Related Art

Patent Literature

Patent Literature 1: Japanese Patent Application Laid-Open No. 2007-100194
Patent Literature 2: International Publication No. WO 2015/182586

SUMMARY

Technical Problem

Patent Literature 2 specifies the annealed structure for obtaining hot work tool steel possessing excellent toughness. However, to obtain the corresponding structure, it requires solid forging with a working ratio (cross-sectional area ratio) of 5s or more, and due to the limitations on the size of the forged material obtained by forging from the size of the steel ingot, it was difficult to obtain a fine annealed structure when the working ratio could not be increased for forged materials with large cross-sectional sizes.
As a means for producing a forged material having a cross-sectional area larger than the cross-sectional area of a steel ingot, upset forging described in JIS-G-7101 is known. However, normally, the temperature during upsetting is often implemented at high temperatures exceeding 1200°C, and since crystal grain growth occurs immediately, it has been difficult to obtain a fine structure. Regarding such challenges during forging, there are no descriptions in Patent Literature 1 and Patent Literature 2, leaving room for further study.
Therefore, the purpose of the present invention is to provide a method for producing hot work tool steel that has a fine annealed structure even if a cross-sectional size is large.

Solution to the Problem

The present inventors have arrived at the present invention by discovering conditions for obtaining a fine annealed structure through optimizing the timing and temperature of the upset forging mentioned above, as well as the forging ratio, in order to obtain hot work tool steel having a large cross-sectional area and a fine structure.
That is, an aspect of the present invention is a method for producing hot work tool steel having a cross-sectional area of 200000mm² or more, includes: a blooming forging process in which a steel ingot is heated to 1100-1250°C and then hot forged to obtain a semi-finished product; a first finish forging process in which the semi-finished product is heated to 850-1020°C and upset forged with an upset forging ratio of 1.6 or more to obtain an intermediate forged material; and a second finish forging process in which solid forging with a solid forging ratio of 2.0-4.0 is performed on the intermediate forged material to obtain a forged material.
It is preferable to have a soaking process in which the steel ingot is heated at 1250°C or higher before the blooming forging process.
Another aspect of the present invention is hot work tool steel having a cross-sectional area of 200000mm² or more, and ferrite grains in a cross-section of a structure have a grain size distribution in which a grain size at a cumulative cross-sectional area of 90% of a total cross-sectional area in an oversize cumulative distribution based on a cross-sectional area of the ferrite grains is 25µm or less in an equivalent circular diameter.

Effects

According to the present invention, hot work tool steel having a fine annealed structure may be provided even if a cross-sectional size is large.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a schematic diagram for describing the manufacturing method of the present invention.
FIG. 2 is an optical micrograph of the cross-sectional structure of the hot work tool steel (sample No. 1) of an example of the present invention.
FIG. 3 is a crystal grain boundary diagram obtained by electron backscatter diffraction (EBSD) of the hot work tool steel (sample No. 1) of the example of the present invention.
FIG. 4 is an optical micrograph of the cross-sectional structure of hot work tool steel (sample No. 2) of an example of the present invention.
FIG. 5 is a crystal grain boundary diagram obtained by electron backscatter diffraction (EBSD) of the hot work tool steel (sample No. 2) of the example of the present invention.
FIG. 6 is an optical micrograph of the cross-sectional structure of hot work tool steel (sample No. 3) of an example of the present invention.
FIG. 7 is a crystal grain boundary diagram obtained by electron backscatter diffraction (EBSD) of the hot work tool steel (sample No. 3) of the example of the present invention.
FIG. 8 is an optical micrograph of the cross-sectional structure of hot work tool steel (sample No. 11) of a comparative example.
FIG. 9 is a crystal grain boundary diagram obtained by electron backscatter diffraction (EBSD) of the hot work tool steel (sample No. 11) of the comparative example.
FIG. 10 is a diagram showing the grain size distribution of ferrite grains distributed in the cross-sectional structure of the hot work tool steel (sample No. 1, sample No. 2, sample No. 3, and No. 11) of the examples of the present invention and comparative example.

DESCRIPTION OF EMBODIMENTS

The feature of the present invention is that an intermediate forged material obtained by performing upset forging at a predetermined temperature on a semi-finished product obtained by blooming forging, and in the next process, a forged material with a fine ferrite grain size is obtained by implementing finish forging at a solid forging ratio. The constituent elements of the present invention are described below.
The hot work tool steel obtained by the manufacturing method of the present invention possesses an annealed structure, and is hot work tool steel that is used after quenching and tempering, and is hot work tool steel having a component composition that can be adjusted to a martensite structure by the above-mentioned quenching.
The effect of refining the structure of the hot work tool steel obtained by the manufacturing method of the present invention can be achieved by implementing the manufacturing method of the present invention described later, as long as the material is one that develops a martensite structure after quenching and tempering. Therefore, it is not needed to specify the component composition of the hot work tool steel to achieve the aforementioned effect of the present invention. However, in establishing the absolute mechanical properties of the hot work tool, for example, the hot work tool steel according to the manufacturing method of the present invention preferably has a component composition containing, in mass%, C: 0.3 to 0.5%, Cr: 3.0 to 6.0%, and more preferably further has a component composition containing V: 0.1 to 1.5%. Also, the hot work tool steel according to the manufacturing method of the present invention is preferably applied to a 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%, with the remainder being Fe and impurities. The reasons for limiting the above component ranges are described below.

C: 0.3~0.5% by mass (hereinafter, simply referred to as "%")

C is a basic element of the hot work tool material, which partially dissolves in the matrix to provide strength, and partially forms carbides to enhance wear resistance and seizure resistance. Additionally, C dissolved as an interstitial atom, when added together with a substitutional atom such as Cr that has a high affinity with C, is expected to act as I (interstitial atom)-S (substitutional atom) effect; functioning as drag resistance of a solute atom, and is expected to increase the strength of hot work tools. However, excessive addition leads to a decrease in toughness and hot strength. Therefore, it is preferable to set the content at 0.3~0.5%.

Si: 2.0% or less

Si is a deoxidizing agent during steelmaking, but excessive amounts lead to the formation of ferrite in the tool structure after quenching and tempering. Therefore, it is preferable to limit the content to 2.0% or less. On the other hand, Si has the effect of enhancing the machinability of the material. To obtain the effect, an addition of less than 0.2% may be sufficient, but an addition of 0.2% or more is preferable.

Mn: 1.5% or less

Mn, if excessive, increases the viscosity of the matrix and decreases the machinability of the material. Therefore, it is preferable to keep the content at 1.5% or less. On the other hand, Mn has the effect of enhancing hardenability, suppressing the formation of ferrite in the tool structure, and obtaining an appropriate quenching and tempering hardness. Additionally, when existing as non-metallic inclusion MnS, Mn has a significant effect on improving machinability. To obtain the effects, an addition of less than 0.1% may be sufficient, but an addition of 0.1% or more is preferable.

P: 0.05% or less

P is an impurity element that is inevitably included in various hot work tool materials. And, P is an element that segregates to prior austenite grain boundaries during heat treatment such as tempering and embrittles the grain boundaries. Therefore, to improve the toughness of hot work tools, it is preferable to restrict the content to 0.05% or less.

S: 0.05% or less

S is an impurity element that is inevitably included in various hot work tool materials. And, S is an element that deteriorates the toughness of the material. Therefore, to improve the toughness of hot work tools, it is preferable to restrict the content to 0.05% or less. On the other hand, S has the effect of improving machinability by combining with the aforementioned Mn and existing as non-metallic inclusion MnS. To obtain the effect, an addition of less than 0.03% may be sufficient, but an addition of 0.03% or more is preferable.

Cr: 3.0 to 6.0%

Cr is an element that enhances hardenability, and also forms carbides, possessing effects for strengthening the matrix and improving wear resistance. And, Cr contributes to the improvement of tempering softening resistance and high temperature strength, being a basic element of the hot work tool material. However, excessive addition leads to a decrease in hardenability and high temperature strength. Therefore, the content is preferable to be 3.0 to 6.0%, and more preferably 5.0% or less. In the present invention, since the effect of toughness improvement by refinement of martensite structure is obtained, the Cr content by the amount equivalent to the effect may be lowered. In this case, for example, by setting Cr to 5.0% or less, further improvement of the aforementioned high temperature strength can be achieved.

Mo and W, either individually or in combination, (Mo+1/2W): 0.5 to 3.5%

Mo and W can be added individually or in combination to precipitate or coagulate fine carbides during tempering to impart strength and improve softening resistance. In this case, the addition amount can be specified together as Mo equivalent of (Mo+1/2W), since W has an atomic weight approximately twice that of Mo (of course, either one can be added alone, or both can be added together). And to obtain the aforementioned effect, an addition of 0.5% or more in terms of a value of (Mo+1/2W) is preferable. However, excessive amounts lead to a decrease in machinability and toughness, so the value of (Mo+1/2W) of 3.5% or less is preferable.

V: 0.1 to 1.5%

V possesses the effect of forming carbides, enhancing the strengthening of the matrix, wear resistance, and tempering softening resistance. And, the above carbides distributed in the annealed structure work as pinning particles that suppress the coarsening of austenite crystal grains during quenching heating, contributing to the improvement of toughness. To obtain the effects, an addition of 0.1% or more is preferable. And, in the present invention, in order to further promote the refinement of the martensite structure, it is preferable to add V. However, excessive amounts lead to deterioration of machinability and a decrease in toughness due to the increase of carbides themselves, the content is preferable to be 1.5% or less.
In addition to the above element types, the following element types may also be included. Ni: 0 to 1.0%
Ni is an element that increases the viscosity of the matrix and decreases machinability. Therefore, it is preferable to limit the content to 1.0% or less. On the other hand, Ni is an element that suppresses the formation of ferrite in the tool structure. Additionally, together with C, Cr, Mn, Mo, W, etc., Ni imparts excellent hardenability to the tool material, and even when the cooling rate during quenching is moderate, Ni forms a structure primarily consisting of martensite, making Ni an effective element for preventing reduction in toughness. Furthermore, Ni also improves the intrinsic toughness of the matrix, so in the present invention, Ni may be added as needed. When added, an addition of less than 0.1% may be sufficient, but an addition of 0.1% or more is preferable.

Co: 0 to 1.0%

Co, as it decreases toughness, is preferably 1.0% or less. On the other hand, Co forms an extremely dense and adherent protective oxide film on the surface during heating when the hot work tool is in use. The oxide film prevents metal contact with the counterpart material, suppresses temperature rise on the tool surface, and provides excellent wear resistance. Therefore, Co may be added as needed. When added, the content may be less than 0.3%, but 0.3% or more is preferable.

Nb: 0 to 0.3%

Nb, as it leads to a decrease in machinability and toughness, is preferable 0.3% or less. On the other hand, Nb possesses the effect of forming carbides, and enhancing the strengthening of the matrix and wear resistance. Also, Nb possesses the effect of enhancing tempering softening resistance, and similar to V, suppressing crystal grain coarsening, and contributing to the improvement of toughness. Therefore, Nb may be added as needed. When added, the content may be less than 0.01%, but 0.01% or more is preferable.
The main elements that may remain in the steel as inevitable impurities include Cu, Al, Ca, Mg, O (oxygen), N (nitrogen), etc. In the present invention, it is preferable that the elements be as low as possible. However, on the other hand, a small amount may be contained to obtain additional effects such as controlling the form of inclusions, improving other mechanical properties, and enhancing manufacturing efficiency. In this case, the ranges of Cu≦0.25%, Al ≦0.025%, Ca≦0.01%, Mg≦0.01%, O≦0.01%, N≦0.03% are sufficiently acceptable and are the exemplarily regulatory upper limits of the present invention.
Next, the manufacturing method of the present invention will be described. Hot work tool steel having an annealed structure is normally produced by starting with a material consisting of a steel ingot or a semi-finished product obtained by a blooming process of a steel ingot as a starting material, performing various hot working and heat treatment thereon to form a predetermined forged material, and applying annealing treatment to the forged material to finish the forged material into a desired shape such as a block shape. FIG. 1 shows a schematic diagram for describing the manufacturing method of the present invention (in the schematic diagram, the "hatched" surfaces of the semi-finished product and forged material are common surfaces between each of processes). The present invention is suitable for a method for producing large-sized hot work tool steel with a cross-sectional area of 200000mm² or more. A more preferable cross-sectional area is 300000mm² or more. Also, the aspect ratio in 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 a cross-section perpendicular to the longitudinal direction of the forged material.

First, in the present invention, a blooming forging process is implemented, in which a steel ingot obtained by casting is heated to 1100-1250°C and then hot forged to obtain a semi-finished product. Normally, as the size of a steel ingot obtained by casting increases, the solidification rate decreases and coarse dendrites grow. Also, many shrinkage cavities due to solidification shrinkage exist in the final solidification part. Since it is preferable to perform at high temperature to destroy the dendritic structure and compress the cavities, the heating temperature in the blooming forging process of the present invention is set to 1100-1250°C. Furthermore, depending on the cross-sectional area of the steel ingot and the cross-sectional area of the hot work tool steel to be finally obtained, upset forging may be implemented during the blooming forging process.
In the present invention, it is preferable to perform annealing following the blooming forging process. By the process, residual stress can be removed from the semi-finished product after hot plastic working. At this time, it is preferable that the annealing temperature is 600°C or higher. More preferably, the annealing temperature is set to be the austenite transformation point or higher and set to be 900°C or lower. By transforming from a ferrite structure to an austenite structure once, an effect of regulating the crystal grain size can be expected.

Next, a first finish forging process is implemented, in which the semi-finished product obtained by blooming forging is heated to 850-1020°C and upset forging is performed with a forging ratio of 1.6 or more, to obtain an intermediate forged material. The forging ratio refers to the ratio of a length L₁ of the semi-finished product shown in FIG. 1 to a length L₂ of the semi-finished product after upset forging, and can be obtained by the formula 1/(L₂/L₁). By implementing the first finish forging at the heating temperature range described above, there is a tendency to be able to refine the structure because crystal grain growth is suppressed, and nucleation sites during phase transformation at the time of heat treatment in the next process increase due to recrystallization and accumulated strain. The preferable upper limit of the heating temperature in the first finish forging process is 1000°C. The more preferable upper limit is 990°C. The subsequent finish forging processes is preferably implemented at a temperature equal to or lower than the temperature at which the aforementioned pinning carbides sufficiently exist. The lower limit temperature is preferably a temperature at the austenite transformation point or higher. In the case of SKD61, which is a representative grade of hot work tool steel, since the austenite transformation point is around 850°C, the lower limit of the heating temperature in the first finish forging process in the present invention was set to 850°C.
The upset forging ratio in the first finish forging process is set to 1.6 or more. If the upset forging ratio is small, the cross-sectional area of the intermediate forged material after upsetting is unable to be sufficiently obtained, making it difficult to obtain a forged material with a large cross-sectional size when the solid forging ratio is set to 2.0-4.0 in a second finish forging process described later. Although there is no upper limit for the upset forging ratio, if the forging ratio is increased and the cross-sectional area increases too much, the load required for forging increases and exceeds the pressing force of the forging machine, so the forging ratio can be set to, for example, approximately 2.0. After upset forging, the shape is formed into an intermediate forged material having a cross-sectional area needed to obtain a solid forging ratio of 2.0-4.0, in accordance with the size of the forged material to be obtained in the second finishing process.

In the manufacturing method of the present invention, a second finish forging process is implemented, in which solid forging with a solid forging ratio of 2.0-4.0 is performed on the intermediate forged material obtained through the first finish forging process to obtain a forged material. The solid forging ratio refers to the ratio of a cross-sectional area A₃ of the intermediate forged material shown in FIG. 1 to a cross-sectional area A₄ after the second finish forging process, and can be calculated by the formula A₃/A₄. The second finish forging may be started immediately after the completion of the first finish forging process, or may be implemented after the workpiece is put into a heating furnace and reheated and held therein. The temperature at which the second finish forging is implemented may be different from the temperature of the first finish forging without causing any problems. Preferably, it is good to implement the second finish forging at the same temperature as the temperature of the first finish forging or at a lower temperature.
The method for producing hot work tool steel of the present invention preferably includes performing annealing at a temperature at the austenite transformation point or higher on the forged material after the second finish forging process. The process removes residual stress from the forged material after hot plastic working, and by sufficiently lowering the hardness, facilitates machining into the shape of hot work tools such as dies described later. At this time, the annealing temperature is preferably at the austenite transformation point or higher and 900°C or lower. If the annealing temperature is too high, for example, the dissolution of carbides in the forged material structure into the matrix progresses, and during cooling from the annealing temperature, carbides may preferentially re-precipitate at crystal grain boundaries, potentially adversely affecting the carbide distribution in the final product structure.
In the manufacturing method of the present invention, it is preferable to implement soaking, which involves heating the steel ingot at 1250°C or higher, before the blooming forging process. Normally, when steel solidifies, the initial solidification part becomes depleted in components, while the final solidification part becomes enriched in components and coarse eutectic carbides are formed. Furthermore, as the steel ingot size increases and the cooling rate of solidification decreases, the eutectic carbides become coarser and component segregation worsens. By heating at 1250°C or higher, it is expected to have the effect of dissolving the eutectic carbides into the matrix and diffusing the components. Although there is no specific upper limit on the temperature, if the temperature is too high, a part of the steel ingot may melt and form a liquid phase, making it difficult to maintain the shape of the steel ingot. Therefore, it is preferable to set the temperature for each of components of the material.
Furthermore, since similar effects can be expected during the process of heating and holding the steel ingot for the aforementioned blooming forging process, the soaking process can also be combined therewith. However, since the effect of dissolution of eutectic carbides and diffusion of segregation is small at low temperatures, it is preferable to implement the soaking process at a temperature at the heating temperature of the blooming forging process or higher.
The hot work tool steel (forged material) of the present invention obtained by the manufacturing method of the present invention has an area of a cross-section perpendicular to the longitudinal direction of 200000mm² or more (preferably 300000mm² or more), which is a large size, while the ferrite grains in the cross-sectional annealed structure have a grain size distribution in which a grain size (D90) at the cumulative cross-sectional area of 90% of the total cross-sectional area in an oversize cumulative distribution based on the cross-sectional area of the ferrite grains is 25 µm or less in an equivalent circular diameter. This enables stable obtaining of a fine structure with a prior austenite grain size of, for example, No.9.0 or more after quenching and tempering. The preferable upper limit of D90 is 22µm, and the more preferable upper limit of D90 is 20µm. Here, the ferrite grains of the present invention can be measured by taking a sample from the center portion of the cross-section perpendicular to the longitudinal direction of the forged material, and measuring a 400µm×400µm region on the plane parallel to the longitudinal direction (forging direction) of the collected sample. This is because in the case of a forged material with a large cross-sectional area, crystal grains tend to coarsen in the center portion of the cross-section. Also, if recrystallization or phase transformation does not sufficiently refine the crystal grains, coarse grains elongated in the forging direction may remain, which can be confirmed by observing the parallel plane. And, the above measurement is performed on two or more cross-sections of the forged material. Furthermore, it is preferable that the two or more cross-sections include the cross-sections at both end portions of the forged material after cutting off the end portions. The above end portions are unneeded parts that are normally cut off and removed because of being unsuitable for products. Moreover, when the annealed structure according to the present invention includes carbides, the carbides act as pinning particles that suppress coarsening of austenite crystal grains during quenching heating, further suppressing the growth of austenite crystals. And the carbides exist as undissolved carbides in the hot work tool after quenching and tempering, in the structure in which the prior austenite grain size has been refined. Here, "annealed structure" in the present invention refers to the structure obtained by annealing treatment, preferably a structure softened to a hardness of, for example, about 150-230HBW on the Brinell hardness scale.
The hot work tool steel obtained by the manufacturing method of the present invention is prepared into a martensite structure having a predetermined hardness by quenching and tempering, and is finished into a hot work tool product. And, the hot work tool steel is finished into the shape of a hot work tool by various machining such as cutting and drilling. The timing of the above processing is preferably performed in a state where the hardness of the material is low (annealed state) before quenching and tempering. In this case, finishing processing may be performed after the quenching and tempering. Also, depending on the case, processing may be performed in a pre-hardened state after the quenching and tempering is performed.
The quenching temperature and the tempering temperature differ depending on the component composition of the material and the target hardness, etc., but it is preferable that the quenching temperature is generally about 1000-1100°C, and the tempering temperature is generally about 500-650°C. Since high quenching temperatures cause the aforementioned pinning carbides to dissolve into the matrix and crystal grain growth to occur, it is preferable to set a temperature at which the pinning carbides sufficiently remain. For example, in the case of SKD61, which is a representative grade of hot work tool steel, the quenching temperature is about 1000-1030°C, and the tempering temperature is about 550-650°C. It is preferable that the hardness after quenching and tempering is 50HRC or less, and more preferably 48HRC or less. Example
Using a melting furnace, 10t steel ingots having the component compositions shown in No.1-No.3 of Table 1 were melted. After performing soaking to maintain the steel ingots at a temperature of 1250°C or higher, blooming processing was conducted at 1160°C, followed by annealing at 870°C to obtain semi-finished products with dimensions of 820mm thickness × 820mm width × 1900mm length. The processes shown in No.1-No.3 of Table 2 were applied to the semi-finished products, and the obtained forged materials were annealed at 870°C to obtain hot work tool steel as examples of the present invention.

Also, using a melting furnace, a 10t steel ingot having the component composition shown in No.11 of Table 1 was melted. After performing soaking to maintain the steel ingot at a temperature of 1250°C or higher, blooming processing was conducted at 1160°C, followed by annealing at 870°C to obtain a semi-finished product with a dimension of 700mm thickness × 1090mm width × 1680mm length. The process shown in No.11 of Table 2 was applied to the semi-finished product, and the obtained forged material was annealed at 870°C to obtain hot work tool steel as a comparative example.

[Table 1]

(mass%)
No.	C	Si	Mn	P	S	Cr	Mo	V	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.

contain inevitable impurities

[Table 2]

No.	First finish forging process			Second finish forging process			Remark
No.	Temperature	Upset forging ratio	Cross-sectional size	Temperature	Solid forging ratio	Cross-sectional size	Remark
1	980°C	2.0	900mm thickness	980°C	3.1	360mm thickness	Examples of the present invention
			900mm thickness			880mm width (cross-sectional area: 316800 mm²)
			1100mm width			880mm width (cross-sectional area: 316800 mm²)
2	980°C	2.0	740mm thickness	980°C	3.1	300mm thickness
			740mm thickness			780mm width (cross-sectional area: 234000 mm²)
			990mm width			780mm width (cross-sectional area: 234000 mm²)
3	980°C	2.0	980mm thickness	980°C	3.1	390mm thickness
			980mm thickness			880mm width (cross-sectional area: 334200 mm²)
			1090mm width			880mm width (cross-sectional area: 334200 mm²)
11	-			900°C	2.4	360mm thickness	Comparative Example
11	-			900°C	2.4	880mm width (cross-sectional area: 316800 mm²)	Comparative Example

Next, the cross-sectional structures of the manufactured hot work tool steel (samples) No.1-No.3 and No.11 were observed. The cross-section observed for the structure was obtained by cutting unneeded parts from both end portions of the steel material. Then, two plate-like test pieces were collected by slicing at positions approximately 15mm from both end faces of the sample to include the obtained cross-section. Test pieces of approximately 10mm to 20mm were cut out from the center portion of the plate-like test pieces, and the surface parallel to the forging direction (i.e., the longitudinal direction of the sample) was observed. After etching the structure of the observation surface with a picric acid alcohol solution, the observation was conducted using an optical microscope (200x magnification), and the observed cross-sectional area was 0.16mm² (400µm×400µm). As a result of the observation, the cross-sectional structures of hot work tool steel No.1 and No.11 were occupied almost entirely by the ferrite phase, with ferrite grains occupying 99 area% or more of the observed cross-sections.
Next, the cross-sectional structures of samples No.1-No.3 and No.11 were observed to confirm the distribution status of ferrite grains in the structures. First, for the cross-sectional structure with a cross-sectional area of 0.16mm², an EBSD pattern at 200x magnification was analyzed to obtain a crystal grain boundary diagram delineated by high-angle grain boundaries with an orientation difference of 15° or more. For the EBSD pattern analysis, an EBSD device (measurement interval 1.0µm) attached to a scanning electron microscope (Carl Zeiss ULTRA55) was used. The optical micrographs and crystal grain boundary diagrams of one end face of each of the samples as measurement results are shown in FIGS. 2-5. Additionally, following the aforementioned procedure, the grain size of individual crystal grains was read from the crystal grain boundary diagrams and converted to an equivalent circular diameter (area equivalent circular diameter), and the grain size distribution of ferrite grains based on the equivalent circular diameter was confirmed. The grain size distribution at one end face of samples No.1 and No.11 is shown in FIG. 6. In FIG. 6, the longitudinal axis represents the cumulative cross-sectional area (%) of crystal grains, and the horizontal axis represents the equivalent circular diameter of crystal grains. From the measurement results, it was confirmed that the equivalent circular diameters of the cumulative cross-sectional area at 90% (d90) of the total cross-sectional area of hot work tool steel No.1, No.2, and No.3 manufactured by the present invention was 18µm, 18µm, and 19µm respectively at one end face, and 18µm, 17µm, and 15µm respectively at the other end face. On the other hand, the equivalent circular diameter of the cumulative cross-sectional area at 90% (d90) of the total cross-sectional area of hot work tool steel No.11 obtained by the manufacturing method of the comparative example was 30µm at one end face and 24µm at the other end face. From the above, it was confirmed that the hot work tool steel of the examples of the present invention possesses a finer structure than the comparative example.

Reference Signs List

1 semi-finished product after blooming forging
2 intermediate forged material after upset forging during first finish forging
3 intermediate forged material after first finish forging
4 forged material after second finish forging

Claims

A method for producing hot work tool steel having a cross-sectional area of 200000mm² or more, comprising: a blooming forging process of heating a steel ingot to 1100-1250°C and then performing hot forging to obtain a semi-finished product; a first finish forging process of heating the semi-finished product to 850-1020°C and performing upset forging with an upset forging ratio of 1.6 or more to obtain an intermediate forged material; and a second finish forging process of performing solid forging with a solid forging ratio of 2.0-4.0 on the intermediate forged material to obtain a forged material having an area of a cross-section perpendicular to a longitudinal direction of 200000mm² or more.
The method for producing hot work tool steel according to claim 1, comprising a soaking process of heating the steel ingot to 1250°C or higher before the blooming forging process.
Hot work tool steel, having an area of a cross-section perpendicular to a longitudinal direction of 200000mm² or more, and ferrite grains in the cross-section having a grain size distribution in which a grain size at a cumulative cross-sectional area of 90% of a total cross-sectional area in an oversize cumulative distribution based on a cross-sectional area of the ferrite grains is 25 µm or less in an equivalent circular diameter.