Hot working die steel, heat treatment method, and hot working die

Abstract

The present invention relates to a hot working tool steel, a heat treatment method thereof, and a hot working tool. In the hot working tool steel, the chemical composition is, by weight%, Ni: 4.5 to 12%, Al: 0.8 to 2.1%, Cu: 0.4 to 1.99%, Mn: 0.001 to 1%, C: 0.0001 to 0.1%, where the ratio of Al to Ni is 0.075 to 0.234, and the balance is Fe and inevitable impurities. The heat treatment method includes a hardening step, where when Cu is 0.4% or more and less than 0.8%, the hot working tool steel is heated to a temperature in the range of 560 to 650 °C, when Cu is 0.8% or more and less than 1.5%, the hot working tool steel is heated to a temperature in the range of 470 to 580 °C, when Cu is 1.5% or more and 1.99% or less, the hot working tool steel is heated to a temperature in the range of 450 to 520 °C, held for 1 to 240 hours, preferably 4 to 24 hours, and then cooled to room temperature.

Description

Technical Field

[0001] The present invention relates to hot working die steel, a heat treatment method, and a hot working die.

Background Art

[0002] Hot working die steel is alloy tool steel in which alloying elements such as carbon, silicon, chromium, tungsten, molybdenum, nickel, manganese, vanadium, and cobalt are selectively added to iron in different weight percentages, and is often remade into a die for forming materials during die casting, forging, and extrusion molding. In recent years, with the emergence of a forming technology for advanced high-strength steel sheets for automobiles that can simultaneously meet the requirements of weight reduction and safety, that is, hot stamping forming technology, new requirements and problems have been raised for die steel.

[0003] The thermal conductivity of general hot working die steel at room temperature is 18 to 24 W / mK. Generally, when the temperature exceeds a certain value, the thermal conductivity decreases as the temperature rises. During use, since the hot working die is subjected to high thermal and mechanical loads at high temperatures, it may cause thermal shock or thermal fatigue due to the temperature gradient. In addition, the thermal conductivity of the hot working die decreases due to high temperature, and a thermal expansion difference due to the temperature gradient occurs, and thermal fatigue cracks may occur in the hot working die, shortening its service life. Also, at high temperatures, the hardness of the precipitation carbides, which is important for ensuring wear resistance, decreases, and the wear resistance of the hot working die decreases at high temperatures.

[0004] It is well known that the higher the thermal conductivity of a material, the smaller the thermal gradient, and as a result, the smaller the surface load, the reduction of thermal shock and thermal fatigue, and thereby the extension of the service life of the hot working die. Therefore, there is a demand for hot working die steel having high thermal conductivity, high hardness, and high toughness in order to enhance the heat crack resistance and wear resistance of the hot working die and thereby extend its service life. Furthermore, when the thermal conductivity of the hot working die is high, cooling during the manufacturing process is promoted, and the cycle time of the hot working die is shortened.

[0005] U.S. Patent No. 9,689,061 discloses a high thermal conductivity alloy tool steel having the following chemical composition by weight percent: C: 0.26 - 0.55%, Cr: less than 2%, Mo: 0 - 10%, W: 0 - 15%, Mo + W: 1.8 - 15%, Ti + Zr + Hf + Nb + Ta: 0 - 3%, V: 0 - 4%, Co: 0 - 6%, Si: 0 - 1.6%, Mn: 0 - 2%, Ni: 0 - 2.99%, S: 0 - 1%, with the balance being Fe and unavoidable impurities. This patent teaches that after solution treatment and quenching, the carbides of Mo and W are replaced by the carbides of Cr, thereby improving the thermal conductivity of the alloy tool steel. However, it is difficult to control the size of Mo carbides and W carbides. Especially when the size of the primary carbides that did not dissolve after solution treatment is about 3 μm, these large carbides may become the sites where fatigue cracks occur, not only having a significant impact on the fatigue life of the tool steel, but also significantly reducing the toughness. Domestic researchers have discovered that although the thermal conductivity of this high thermal conductivity alloy tool steel is about 47 W / mK at room temperature, the thermal conductivity decreases as the temperature rises. At 300 °C, the thermal conductivity drops below 39 W / mK, and at 500 °C, the thermal conductivity is only 35 W / mK. Furthermore, when the hardness exceeds 50 HRC, the impact toughness (measured with a 7×10 mm notch specimen) is less than 210 J. When this high thermal conductivity alloy tool steel is used at high temperatures, the advantage of high thermal conductivity is lost, and an optimal combination of high thermal conductivity, high toughness, and high hardness cannot be achieved.

[0006] Chinese Patent Application Publication No. 108085587 discloses a die steel for die casting with high thermal conductivity and long service life, as well as a manufacturing method thereof. The chemical composition of this die steel for hot die casting with high thermal conductivity and long service life is, by weight percentage, C: 0.35 - 0.45%, Si: 0.20 - 0.30%, Mn: 0.30 - 0.40%, Ni: 0.50 - 1.20%, Cr: 1.5 - 2.2%, Mo: 2 - 2.6%, W: 0.0001 - 1.0%, Ti: 0 - 0.40%, V: 0.30 - 0.50%, B: 0.0001 - 0.003%, Cu: 0.005 - 0.02%, and the balance is Fe and inevitable impurities. In this patent, Mo carbides and W carbides are also used instead of Cr carbides. However, there are the following problems: 1) It is difficult to control the size of the carbides. If the carbides are too large, the toughness will decrease; 2) When Ti is added, liquid-phase TiN and excessive TiC are formed, deteriorating the toughness; 3) Since the Cr element cannot be completely precipitated from the matrix as carbides, high thermal conductivity cannot be obtained; 4) Multiple tempering processes are complex and attention needs to be paid to avoiding the secondary hardening peak because, thereby, the hardness is maximized but the toughness is minimized. In the preferred embodiment of this patent, the impact energy shown in the U-notch impact test of the example steel is 50 J or less, and the maximum thermal conductivity is 35.982 W / mK. Still, an optimal performance balance of high thermal conductivity, high toughness, and high hardness cannot be achieved.

[0007] Chinese Patent No. 104046917 discloses an ultra-high strength ferrite steel strengthened by nano-sized intermetallic compounds. The chemical composition of this ultra-high strength ferrite steel is, by weight, C: 0 to 0.2%, Cu: 0.5 to 5%, Ni: 0.01 to 4%, Mn: 0.01 to 4%, Al: 0.001 to 2%, Cr: 0 to 12%, Mo: 0 to 3%, W: 0 to 3%, Mo + W: 0.05% or more, V: 0 to 0.5%, Ti: 0 to 0.5%, Nb: 0 to 0.5%, V + Ti + Nb: 0.01% or more, Si: 0 to 1%, B: 0.0005 to 0.05%, P: 0.04% or less, S: 0.04% or less, N: 0.04% or less, O: 0.05% or less, and the balance is Fe and unavoidable impurities. This patent teaches that after rolling the steel, solution treatment and aging treatment are performed to obtain a ferrite steel mainly strengthened by nano-sized intermetallic compounds, combined with fine crystal, solid solution strengthening, and dislocation strengthening. As a result, excellent strength-toughness balance, weldability, and corrosion resistance are obtained.

[0008] However, since many of the alloying elements remain in solid solution within the ferrite matrix even after the final heat treatment, the thermal conductivity of the ultra-high strength ferrite steel is low. For example, Cr, which forms Cr carbides in the ferrite matrix, not only has a low thermal conductivity but also tends to coarsen during heat treatment, deteriorating the toughness of the steel. Similarly, when Si, Cr, and Al elements dissolve into the matrix, the thermal conductivity of the ferrite significantly decreases. Furthermore, carbides and carbonitrides of trace alloying elements such as V, Ti, and Nb not only reduce the thermal conductivity of the steel but also decrease the toughness.

[0009] Chinese Patent No. 105568151 discloses an aluminum-strengthened maraging steel. The chemical composition (by weight %) of this maraging steel is as follows: C: 0.01 - 0.2%, Ni: 6 - 24%, Mo: 6% or less, Mn: 0 - 4%, Al: 0.5 - 6%, Cr: 0 - 12%, Nb: 1.5% or less, Cu: 0 - 4%, W: 0 - 3%, B: 0.0005 - 0.05%, and the balance is Fe and inevitable impurities. In this patent, Ti and Co are removed from the alloy, and a strengthening composition mainly strengthened by NiAl and combined with trace carbides and nano-sized clusters of Cu is designed, realizing an alloy with a strength exceeding 2000 MPa, an elongation exceeding 6%, and excellent fatigue and processing properties. However, in this patent, the appropriate balance of the ratios of Ni, Al, and Cu is not achieved. As a result, excessive Al dissolves in the matrix, significantly reducing the thermal conductivity of the matrix.

[0010] Chinese Patent Application Publication No. 111636037 discloses a hot working die steel. The chemical composition (by weight %) of this die steel is as follows: Cu: 2 - 8%, Ni: 0.8 - 6% (the ratio of Ni to Cu is 0.4 or more), C: 0 - 0.2%, Mo: 0 - 3%, W: 0 - 3%, Nb: 0 - 0.2%, Mn: 0 - 0.8%, Cr: 0 - 1%, Al: 0 - 3%, Ni:Al is 2 or more (preferably 2 - 2.5), and the balance is Fe and inevitable impurities. This patent compensates for NiAl precipitates and carbides as strengthening phases to improve the hardness of the hot working die steel and utilizes a large amount of Cu precipitates. At the same time, to increase the thermal conductivity of the matrix, it mainly utilizes the precipitation of Cu particles from the matrix. After solution treatment and aging treatment, the hot working die steel exhibits the following properties: a hardness of 42 HRC or more, a thermal conductivity of 35 W / mK or more, and a room temperature impact toughness of 250 J or more for a notched specimen of 7×10 mm. However, this patent does not recognize the influence of Al dissolved in the matrix on the thermal conductivity. As a result, the thermal conductivity of the hot working die steel cannot be further improved.

[0011] Chinese Patent No. 110578103 discloses a plastic die steel with high toughness, high abrasion resistance, and high corrosion resistance. By adding Cu and Al, this patent significantly improves the surface abrasion resistance and corrosion resistance of the plastic die steel, and greatly extends the service life of the steel. On the other hand, the plastic die steel does not need to meet the required performance of hot working die steel. Although Cu, Ni, and Al elements are added, the proportion is not appropriate, so the thermal conductivity has not been significantly improved. Furthermore, other elements such as Si, Cr, and V may significantly reduce the thermal conductivity, so this plastic die steel does not have high thermal conductivity.

Prior Art Documents

Patent Documents

[0012]

Patent Document 1

Patent Document 2

Patent Document 3

Patent Document 4

Patent Document 5

Patent Document 6

Summary of the Invention

Problems to be Solved by the Invention

[0013] Although various improved hot working die steels have been provided by the prior art, the demand for hot working die steels with high thermal conductivity, high hardness, and high toughness still exists.

Means for Solving the Problems

[0014] The present invention has been made in view of the problems existing in the prior art. One of its purposes is to provide a hot-work die steel capable of improving the thermal conductivity and mechanical properties by a simple hardening treatment process. In the hot-work die steel, the chemical composition (by weight %) is Ni: 4.5 to 12%, Al: 0.8 to 2.1%, Cu: 0.4 to 1.99%, Mn: 0.001 to 1%, C: 0.0001 to 0.1%, where the ratio of Al to Ni is 0.075 to 0.234, and the balance is Fe and inevitable impurities.

[0015] Preferably, the ratio of Al to Ni is 0.075 or more and less than 0.18.

[0016] Preferably, the ratio of Al to Ni is 0.18 or more and 0.234 or less.

[0017] More preferably, the ratio of Al to Ni is 0.18 or more and 0.22 or less, and Ni is 5 to 9%.

[0018] Preferably, C: 0.0001 to 0.05%.

[0019] The chemical composition of the hot-work die steel is designed such that after appropriate heat treatment, all alloy elements precipitate from the matrix in the form of Cu particles and Ni(Fe, Al) intermetallic compounds. In particular, since Al precipitates as an intermetallic compound with Ni even after hardening treatment, lattice defects in the matrix are reduced and the thermal conductivity is improved. In this application, when the amount of Al dissolved in the matrix is 0.05% or less, it is considered that Al has completely precipitated from the matrix. These precipitates exhibit excellent thermal conductivity, not only ensuring high hardness and high toughness, but also significantly improving the thermal conductivity of the hot-work die steel.

[0020] Preferably, the chemical composition of the hot working die steel is designed to improve tempering resistance. According to tests, it has been shown that the tempering resistance of the hot working die steel can be improved by adjusting the ratio of copper to aluminum to 0.5 - 1, preferably adjusting the copper content to 0.4 - 1.5%. More preferably, the ratio of copper to aluminum is 0.6 - 0.9, and even more preferably, the copper content is 0.4% - 1.2%.

[0021] Preferably, the hot working die steel of the present invention contains, by weight, 5 - 12% Ni, 0.8 - 1.5% Al (where the ratio of Al to Ni is 0.075 or more and less than 0.18), 0.4 - 1.0% Cu (where the ratio of Cu to Al is more than 0.5 and less than 0.9), 0.001 - 1% Mn, 0.0001 - 0.1% C, and the balance Fe and unavoidable impurities.

[0022] Preferably, the hot working die steel of the present invention contains, by weight, 5.5 - 11% Ni, 1.2 - 2.1% Al (where the ratio of Al to Ni is 0.18 or more and 0.234 or less), 1.0 - 1.9% Cu (where the ratio of Cu to Al is more than 0.8 and less than 1), 0.001 - 1% Mn, 0.0001 - 0.1% C, and the balance Fe and unavoidable impurities. Preferably, the hot working die steel of the present invention contains, by weight, 5.5 - 7.0% Ni, 1.0 - 1.6% Al (where the ratio of Al to Ni is 0.18 - 0.20), 0.4 - 1.0% Cu (where the ratio of Cu to Al is more than 0.5 and less than 0.9), 0.001 - 1% Mn, 0.0001 - 0.1% C, and the balance Fe and unavoidable impurities. Preferably, the hot working tool steel of the present invention contains, by weight percentage, 7.5 - 9% Ni, 1.2 - 2.1% Al (where the ratio of Al to Ni is 0.20 - 0.234), 1.0 - 1.9% Cu (where the ratio of Cu to Al is more than 0.8 and less than 1), 0.001 - 1% Mn, 0.0001 - 0.1% C, and the balance Fe and unavoidable impurities.

[0023] Another object of the present invention is to provide a heat treatment method for the above-described hot working die steel that can be implemented with existing heat treatment and processing equipment. This method simplifies the current heat treatment process of die steel, thereby reducing costs and enabling the hot working tool to have better mechanical properties and thermal conductivity.

[0024] The heat treatment method of the present invention includes a hardening step, where when Cu is 0.4% or more and less than 0.8%, the hot working die steel is heated to a temperature in the range of 560 to 650 °C, when Cu is 0.8% or more and less than 1.5%, the hot working die steel is heated to a temperature in the range of 470 to 580 °C, when Cu is 1.5% or more and 1.99% or less, the hot working die steel is heated to a temperature in the range of 450 to 520 °C, the steel is held for 1 to 240 hours, preferably 4 to 24 hours, and then cooled to room temperature.

[0025] On the one hand, the Cu precipitates tend to coarsen with the increase in Cu content and temperature, and as a result, the coarsening of NiAl precipitates is induced. On the other hand, the precipitated Cu particles function as nucleation sites for NiAl precipitation. At a given temperature, the lower the Cu content, the smaller the Cu precipitates and the coarsening is prevented. However, if the Cu content is too low, the precipitation of NiAl is hindered. Both factors affect the properties of the die steel. Therefore, in the present invention, by adjusting the Cu content and the hardening temperature and determining the hardening temperature based on the Cu content, in the hardening process, due to the synergistic effect of the Cu content and temperature, while preventing coarsening, the precipitation of both Cu particles and NiAl is promoted.

[0026] Preferably, when the Cu in the hot working die steel is 0.4% or more and less than 0.8%, it is heated to a temperature in the range of 560 to 620 °C, when the Cu is 0.8% or more and less than 1.5%, it is heated to a temperature in the range of 480 to 580 °C, and when the Cu is 1.5% or more and 1.99% or less, it is heated to a temperature in the range of 450 to 500 °C.

[0027] After the hardening treatment, the hot work die steel exhibits the following properties: a hardness of 36 HRC or higher, a room temperature thermal conductivity of 34 to 50 W / mK, an impact energy of 200 J or higher for a notch-free test piece (7×10×55 mm), and an Al dissolution amount in the matrix of 0.05% or less.

[0028] Preferably, before the step of hardening, a solution treatment step is performed, where the hot work die steel is heated to a temperature in the range of 850 to 950 °C and held for 0.5 to 240 hours, and then cooled to room temperature.

[0029] Preferably, before the solution treatment step, a forging step is performed, where the hot work die steel is heated to a temperature in the range of 900 to 1150 °C for forging, and then cooled to room temperature. Preferably, the forging temperature is in the range of 950 to 1100 °C, and more preferably 1000 to 1050 °C.

[0030] Preferably, the cooling in the solution treatment step and / or the forging step is first cooled to a temperature of less than 600 °C, more preferably less than 500 °C, and even more preferably less than 400 °C at a rate of 60 °C / hour or less, more preferably 30 °C / hour or less, and then air-cooled to room temperature.

[0031] In the heat treatment method of the present invention, since the steel defined in the present invention is used, the spheroidizing annealing process used in conventional die steels becomes unnecessary. Further, the solution treatment temperature is reduced from 1000 to 1050 °C to 850 to 950 °C, and the requirements for heat treatment equipment are reduced. Further, the hardening process is performed at a single temperature for a long time. This is because the steel composition specified in the present application does not contain elements that can form other alloy carbides, so the carbide formed is only cementite. Since the carbon content in the steel is low, the growth of the carbide is limited. As a result, the hardening treatment can be performed at a single temperature for a long time, so that the process is simplified, and the complicated hardening procedures that require holding at various temperatures for various times used in conventional hot work die steels become unnecessary.

[0032] Another object of the present invention is to provide a hot working die made from the above heat-treated hot working die steel and having high hardness, high thermal conductivity, and high toughness. Preferably, since the die has high tempering resistance at high temperatures (the precipitation phase is difficult to grow at high temperatures), it can prevent the performance degradation of the die during use and improve the thermal fatigue resistance. Furthermore, since the hot working die according to the present application has high thermal conductivity, it can shorten the production cycle time and improve the production efficiency. Preferably, the hot working die can be a hot stamping die for steel sheets, an aluminum alloy die-casting die, or a hot extrusion die.

[0033] The % shown for chemical elements in this specification means weight %. Unless otherwise specified, preferred embodiments can be freely combined as needed.

[0034] Furthermore, those skilled in the art will understand that any sub-range or specific value within the aforementioned range is applicable to the present invention. For example, in the range of 0.075 to 0.18 (excluding 0.18), any sub-range or specific value within that range, such as 0.075, 0.08, 0.082, 0.085, 0.088, 0.09, 0.095, 0.1, 0.13, 0.15, 0.157, 0.16, 0.165, 0.17, as well as sub-ranges such as 0.075 to 0.085, 0.08 to 0.098, 0.09 to 0.14, and 0.12 to 0.17 are included. Similarly, in the range of 0.18 to 0.234, any sub-range or specific value within that range, such as 0.18, 0.185, 0.19, 0.195, 0.198, 0.203, 0.207, 0.213, 0.215, 0.225, as well as sub-ranges such as 0.18 to 0.215, 0.18 to 0.207, 0.185 to 0.225, 0.19 to 0.225, and 0.19 to 0.22 are included. Similarly, in the range of 0.5 to 1 (excluding the end values), any sub-range or specific value within that range, such as 0.56, 0.62, 0.69, 0.73, 0.80, 0.88, 0.90, 0.92, 0.98, as well as sub-ranges such as 0.51 to 0.98, 0.56 to 0.92, 0.62 to 0.98, 0.65 to 0.92, and 0.65 to 0.9 are included. Unless otherwise specified, all ranges include the end values.

Brief Description of the Drawings

[0035]

Figure 1

Embodiments for Carrying Out the Invention

[0036] (Detailed Description of the Invention) Hereinafter, the present invention will be described in more detail with reference to exemplary embodiments.

[0037] In the hot working die steel according to the present invention, the chemical composition is, by weight%, Ni: 4.5 to 12%, Al: 0.8 to 2.1%, Cu: 0.4 to 1.99%, Mn: 0.001 to 1%, C: 0.0001 to 0.1%, where the ratio of Al to Ni is 0.075 to 0.234, and the balance is Fe and unavoidable impurities. The roles and ratios of each element in the present invention are described below.

[0038] Ni: The main role of nickel (Ni) in the present invention is to form NiAl intermetallic compounds and contribute to precipitation hardening and strengthening. Furthermore, Ni can suppress the formation of liquid-phase precipitation of Cu at grain boundaries at high temperatures, prevent hot cracking during hot deformation, and ensure the hot forming performance of the alloy. Ni improves the hardenability of the steel, suppresses pearlite transformation, and Ni enriched at grain boundaries can enhance toughness. When it is desired to mainly strengthen and harden the steel with the precipitated NiAl intermetallic compound, a high nickel content is required so that Al can be completely combined with Ni and precipitate from the matrix. Otherwise, Al remains dissolved in the matrix, and the thermal conductivity of the steel decreases significantly. For the purpose of completely precipitating Al, the excess Ni dissolved in the matrix does not cause significant lattice strain, but when Ni becomes excessive, the Ms temperature decreases, and martensite transformation does not occur during cooling, so cracking may occur due to the high thermal expansion coefficient of austenite. For the purpose of completely precipitating Al, in order to prevent the decrease in the Ms temperature caused by Ni, the nickel content in the steel is 4.5 to 12%. Preferably, the nickel content is 5 to 12%, more preferably 5 to 11%, and even more preferably 5 to 9%. Thereby, Al completely precipitates from the matrix, meeting the performance requirements of the present invention. Furthermore, excessive Ni only increases the alloy cost without improving the performance. Moreover, excessive Ni may complicate machining, such as causing "tool sticking" during turning operations.

[0039] Aluminum and nickel form NiAl intermetallic compounds during the hardening treatment process at 400 - 650 °C, contributing to precipitation strengthening and hardening. Furthermore, the precipitation of NiAl intermetallic compounds improves the purity of the matrix, and these compounds also have excellent thermal conductivity (about 90 W / mK), so both increase the thermal conductivity of the steel. To ensure that Al completely precipitates from the matrix as NiAl intermetallic compounds, when Al is in excess relative to Ni, the following problems may occur: 1. Al may not precipitate sufficiently from the matrix, potentially reducing the thermal conductivity of the steel. 2. Large AlN inclusions may form, which do not completely dissolve in austenite at high temperatures and significantly impair the toughness of the steel. 3. As a strong ferrite stabilizer, Al raises the Ac1 and Ac3 temperatures of the steel. This requires a high temperature for solution treatment to achieve austenitization, increasing manufacturing costs, energy consumption, and the requirements for heat treatment equipment. When the Al content is low, complete precipitation of Al can be ensured. However, if the precipitation of NiAl intermetallic compounds is insufficient, the hardness of the steel may not meet the required specifications. Also, when Ni is in excess and Al is insufficient, it does not significantly affect the hardness, thermal conductivity, and impact toughness of the steel, but it leads to waste of Ni and an increase in material costs. Therefore, the aluminum content in the steel is 0.8 - 2.1%, and the ratio of Al to Ni is 0.075 - 0.234. Preferably, the ratio of Al to Ni is 0.18 - 0.234, and more preferably 0.18 - 0.22, where the ratio of Ni to Al is optimal and the performance is maximized.

[0040] Cu: Pure copper is an excellent heat conductor. While the thermal conductivity of pure iron is 80 W / mK, that of pure copper is 398 W / mK. Copper has a high solubility in the face-centered cubic phase (austenite), but a very low solubility in the body-centered cubic phase (ferrite and / or martensite), so copper particles precipitate. The size of the precipitated ε-Cu particles is 3 - 10 nm. When 1 wt% of Cu is added, the hardness of the matrix increases by about 100 HV. When Cu particles precipitate from the body-centered cubic matrix (ferrite and / or martensite), the lattice strain decreases, and combined with the high thermal conductivity of the precipitated Cu particles, the thermal conductivity of the matrix increases. However, the precipitation of Cu particles occurs at a low aging temperature. When the temperature of the matrix rises (e.g., during quenching), the Cu particles grow rapidly, and the hardness and strength of the steel decrease. Furthermore, in the hot forming process (rolling, forging, etc.) of steel, Cu easily forms a liquid phase at the grain boundaries of austenite, causing thermal cracks during deformation, reducing the plastic deformability of the steel, and making processing difficult. To suppress the liquid-phase precipitation of Cu at the grain boundaries, usually, a certain amount of nickel is added. The precipitated Cu particles provide nucleation sites for the precipitation of NiAl intermetallic compounds, thereby promoting the precipitation of NiAl. In the present invention, considering the effects of the Cu content and the Al content, the copper content of the steel is 0.4 - 1.99%. To improve the aging performance, the weight ratio of Cu to Al is preferably set to 0.5 - 1 (excluding the end values), more preferably set to 0.6 - 0.99, and / or the copper content is 0.4 - 1.2%.

[0041] Mn: Manganese is an austenite-forming element, which delays the transformation from austenite to ferrite, refines the ferrite grain structure, and is effective in improving the strength and toughness of steel. Manganese dissolved in the matrix reduces the thermal conductivity of steel. When Mn replaces some of the Fe and Al atoms to form the Ni(Fe, Mn, Al) intermetallic compound, the thermal conductivity of the original Ni(Fe, Al) decreases. Furthermore, if the Mn content is too high, retained austenite may remain in the steel, resulting in a non-uniform microstructure and unstable mechanical properties. On the other hand, NiMn precipitates are formed at the grain boundaries, which may significantly impair the mechanical properties of the steel. Therefore, the Mn content in the present invention is limited to 0.001 - 1%.

[0042] C: Carbon is one of the most effective and economical strengthening elements in steel and is a stabilizing element for austenite. Carbon is an interstitial solid solution element, and its strengthening effect is much greater than that of substitutional solid solution elements. Carbon increases the hardenability of steel, and the resulting cementite or alloy carbides significantly improve the hardness of the alloy. As the carbon content increases, network carbides or coarse cementite particles may be formed along the grain boundaries, which significantly reduces the toughness of the steel. Since the present invention is based on the precipitation strengthening of NiAl and Cu rather than the strengthening and hardening of carbides, the carbon content in the steel is maintained at 0.0001 - 0.1%.

[0043] The following examples or experimental data are provided for illustrative purposes to explain the present invention, and it should be understood by those skilled in the art that the present invention is not limited to these examples or experimental data.

[0044] The present invention provides a first preferred hot working die steel containing, by weight%, Ni: 4.5 - 5.5%, Al: 0.8 - 1.3% (where the ratio of Al to Ni is 0.075 - 0.234), Cu: 0.4 - 1.99% (where the ratio of Cu to Al is 0.6 or more and 0.99 or less), Mn: 0.001 - 1%, C: 0.0001 - 0.1%, and the balance: Fe and unavoidable impurities.

[0045] The present invention provides a second preferred hot working die steel containing, by weight %, Ni: 5.5 to 7.5%, Al: 1.0 to 1.6% (the ratio of Al to Ni is 0.075 to 0.234), Cu: 1.0 to 1.5% (where the ratio of Cu to Al is more than 0.5 and less than 1), Mn: 0.001 to 1%, C: 0.0001 to 0.1%, and the balance: Fe and unavoidable impurities.

[0046] The present invention provides a third preferred hot working die steel containing, by weight %, Ni: 7.5 to 12%, Al: 1.2 to 2.1% (the ratio of Al to Ni is 0.075 to 0.234), Cu: 1.5 to 1.99 (where the ratio of Cu to Al is more than 0.5 and less than 1), Mn: 0.001 to 1%, C: 0.0001 to 0.1%, and the balance: Fe and unavoidable impurities.

[0047] According to the compositions of the example steels DC1 to DC8 and the comparative example steels CG1 to CG3 shown in Table 1, the steel was melted into ingots, and then hot worked at an initial temperature of about 1050°C and a final temperature of about 950°C to obtain 80×80 mm 2 square billets. Thereafter, the billets were homogenized at 1030°C for 5 hours and then air-cooled to room temperature.

[0048] Table 1 shows the compositions (weight percent, wt%) of the sample steels and the comparative example steels

[0049]

Table 1

[0050] 80×80 mm 2 The square billets of 80×80 mm were cut into 7.2×10×55 mm and diameter 12.7×2.2 mm test pieces and subjected to heat treatment. The heat treatment was a solution treatment at 950°C for 1 hour followed by rapid cooling. Then, a hardening treatment was performed at a temperature of 450 to 650°C for a predetermined time, and thereafter the test pieces were air-cooled. The specific heat treatment parameters of the example steels and the comparative example steels are summarized in Table 2.

[0051] It should be understood that the temperature and time of the hardening treatment are not limited to the specific temperature and time in the embodiments and can be selected from the following ranges: when Cu is 0.4% or more and less than 0.8%, the hot work die steel is heated to a temperature within the range of 560 to 650 °C; when Cu is 0.8% or more and less than 1.5%, the hot work die steel is heated to a temperature within the range of 470 to 580 °C; when Cu is 1.5% or more and 1.99% or less, the hot work die steel is heated to a temperature within the range of 450 to 520 °C. The heat treatment method of the present invention may optionally include other processes such as forging. Further, the temperature and time of the solution treatment and / or forging process can be selected according to actual needs.

[0052] Table 2: Parameters for solution treatment and hardening treatment of sample steel and comparative steel Steel type

[0053]

Table 2

[0054] Hardness measurement: After solution treatment and hardening treatment, the sample is polished with sandpaper to a shiny finish, and a hardness test is performed using a hardness tester (Rockwell hardness). As shown in Table 3, the hardness value H S after solution treatment and the hardness value H H after hardening treatment are obtained.

[0055] Impact energy measurement: A 7.2 × 10 × 55 mm test piece after heat treatment according to Table 2 is mechanically polished according to the notched impact test piece standard of the North American Die Casting Association to make a 7 × 10 × 55 mm notched impact test piece. Next, a pendulum notched room temperature test piece impact test of 450 J is performed to obtain the impact energy of the test piece as shown in Table 3.

[0056] Measurement of Al content: Using energy-dispersive X-ray spectroscopy equipped with a transmission electron microscope, the Al content in the microregions of a 7.2×10×55 mm sample after heat treatment is analyzed according to Table 2. For each sample, at least 5 locations are selected for component testing, and the Al contents at different locations are averaged to obtain the Al content in the matrix. The measurement results are shown in Table 3.

[0057] Measurement of thermal conductivity: The thermal conductivity λ is calculated using the formula λ = α×cp×ρ×100. Here, the thermal diffusivity α is in units of cm 2 / s, the specific heat capacity cp is in units of J / (g·K), and the density ρ is in units of g / (cm 3 ). The resulting unit is W / (cm·K)×100 and is converted to W / (m·K). After performing the heat treatment shown in Table 2, a cylindrical sample with dimensions of diameter 12.7×2.2 mm is polished to a diameter of 12.7×2.0 mm using 1000-grit sandpaper, and the thermal conductivity is measured using a DLF2800 flash thermal conductivity measurement device.

[0058] The measurement process is as follows: The sample is heated from 25°C to 100°C at a rate of 5 K / min, stabilized there for 10 minutes, and then the first measurement is taken. Then, the sample is stabilized for another 10 minutes and the second measurement is taken, and then stabilized for another 10 minutes and the third measurement is taken. Next, the sample is heated to 200°C at a rate of 5 K / min, and three measurements are taken in the same procedure as above. Then, the temperature is sequentially increased to 300°C, 400°C, and 500°C, and three measurements are taken at each temperature (corresponding to a stabilization period of 30 minutes at each measurement temperature). After the measurement is completed, the sample is cooled to room temperature.

[0059] Based on these measurements, the thermal diffusivity, specific heat capacity, and density at each temperature are obtained. Next, as shown in Table 4, the thermal conductivity is calculated using the data of thermal diffusivity, specific heat capacity, and density.

[0060] Deterioration test: Deterioration refers to the phenomenon where the strengthening effect of the second phase in steel weakens due to the coarsening of the second phase at different holding temperatures and holding times. The degree of deterioration can be represented by the rate of decrease in hardness. For samples S4 and S7, after the solution treatment process shown in Table 2, the following hardening treatment was performed: The samples were held at 520 °C for 4 hours, 24 hours, and 48 hours respectively, then the hardness was measured, and the rate of decrease in hardness was calculated according to the aforementioned measurement method (shown in Table 5).

[0061] Table 3: Hardness values, impact energies, and Al contents in the matrix of each sample of the present invention

[0062]

Table 3

[0063] Table 4: Thermal conductivities of each sample of the present invention, W / (mK)

[0064]

Table 4

[0065] Table 5: Hardness of Examples S4 and S7 of the present invention after the deterioration test

[0066]

Table 5

[0067] Hardness analysis The hardness values of samples S1 - S8 and CS1 - CS6 after solution treatment were all less than 35 HRC. This is because during solution treatment, the precipitation phases (such as NiAl and Cu) in the microstructure dissolved into the matrix and no longer played a role in strengthening and hardening.

[0068] After hardening treatment, the hardness values of Examples S1 - S8 and Comparative Examples CS1 - CS6 all increased. Specifically, the hardness values (H H) was 24% to 51.3% higher than the hardness value (H after solution treatment S ) and exceeded 36 HRC. In contrast, the maximum increase in hardness in Comparative Examples CS1 to CS6 was only 16.2%, and the hardness values were in the range of 33.4 HRC to 35.2 HRC, not meeting the required specifications.

[0069] The increase in hardness was due to the precipitation of NiAl intermetallic compounds and Cu particles during the hardening process in both Examples S1 to S8 and Comparative Examples CS1 to CS6. These precipitates serve to strengthen and harden the material, thereby increasing the strength and hardness of the steel to a certain extent.

[0070] Compared with Sample S1, Sample S8 had a higher Ni content, and the ratio of Al to Ni was 0.09. After undergoing the same heat treatment process as Sample S1, the measured hardness H H was slightly lower than that of S1 but still met the technical requirements of the present invention. This indicates that moderately excessive Ni does not significantly reduce the hardness.

[0071] Similarly, the comparison between Sample S2 and Comparative Sample CS5 also reflects the influence of excessive Ni. After undergoing the same heat treatment process, Sample S2 had a solution treatment hardness (H S ) of 31.8 HRC and a hardened hardness (H H ) of 47.5 HRC, showing a 49.3% increase with respect to H S . In contrast, Comparative Sample CS5 had H S of 31.8 HRC and H H of 34.1 HRC, and H Sshowed only a 6.7% increase compared to this. This is because, compared with sample S2, the comparative sample CS5 had a lower Al content with respect to Ni, resulting in a significant excess of Ni (the ratio of Al to Ni was only 0.069). All of the Al could form NiAl intermetallic compounds with Ni (microanalysis revealed that the Al content in the matrix of the comparative sample CS5 was 0.03%), but the amount of precipitated NiAl intermetallic compounds was limited, and thus the strengthening and hardening effects of the NiAl intermetallic compounds were also limited. Furthermore, since Ni is a substitutional solid solution element, the solid solution strengthening effect is very weak, and the effect of dissolved Ni on increasing hardness is very limited. Due to these two factors, H H decreased and did not meet the usage requirements. Furthermore, due to the low Al content with respect to Ni, excess Ni that could not contribute to strengthening and hardening was generated, leading to a waste of resources. Based on these findings, in order to precipitate sufficient NiAl intermetallic compounds, it is desirable that the ratio of Al to Ni is greater than at least 0.069.

[0072] In the present invention, it was found that when the Ni content was too high and the ratio of Al to Ni was less than 0.075 (such as the comparative sample CS5), it was not possible to obtain steel with desired properties. Therefore, in order to meet the technical requirements of the present invention, the ratio of Al to Ni must be at least 0.075 or more.

[0073] After undergoing the same heat treatment process, sample S3 showed an H H of 42.3 HRC, while the comparative sample CS6 showed an H His shown and is approximately 18% lower than the hardness of sample S3. This is because, compared with sample S3, comparative sample CS6 has an excessive Al content with respect to Ni. As a result, the precipitation of Al from the matrix is incomplete, the amount of NiAl precipitates decreases, the strengthening effect and the hardening effect become weak, and finally the hardness does not meet the usage requirements. As a result of microanalysis, it was revealed that the Al content in the matrix of comparative sample CS6 is 0.25%, and it was confirmed that Al did not completely precipitate from the matrix. Comparative sample CS4 showed an even lower hardness compared with CS6. As a result of microanalysis, the Al content in the matrix of comparative sample CS4 (the ratio of Al to Ni is 0.31) is 1.14%, and a large amount of Al did not precipitate from the matrix. As a result, it was found that the final hardness does not meet the usage requirements. From the above, in order to completely precipitate Al, it is desirable that the ratio of Al to Ni is less than 0.27.

[0074] In samples S1 to S8, the ratio of Al to Ni was in the range of 0.075 to 0.234. By microanalysis of the heat-treated samples S1 to S8, the Al content in the matrix of samples S1 to S8 was less than 0.05%, and as shown in FIG. 1, it was revealed that Ni(Fe, Al), Ni(Fe, Mn, Al), and Cu particles were precipitated in the matrix. From these results, when the ratio of Al to Ni is in the range of 0.075 to 0.234, it shows that Al can completely precipitate from the matrix in the form of NiAl intermetallic compound, and there are sufficient precipitates to bring about the precipitation strengthening and hardening effects. As a result, it can be seen that the hardened hardness value (H H ) is 36 HRC or more.

[0075] Therefore, when the composition of the steel satisfies a predetermined range, by ensuring that the ratio of Al to Ni is in the range of 0.075 to 0.234, it is possible to ensure that the steel has a desired high hardness. Preferably, Ni:Al is 0.18 to 0.234, more preferably Ni:Al is 0.18 to 0.22, and even more preferably, the Ni content is 5 to 9%. In one preferred embodiment, the lower the Ni content, the more equivalent or better performance can be obtained.

[0076] Impact energy The impact toughness of Samples S1 to S8 exceeded 200 J, and the corresponding hardness values (36 HRC or higher) also met the usage requirements. In the case of Samples S1 to S8, the precipitation of Cu and NiAl contributed to strengthening, while it was not conducive to the generation and expansion of cracks at the interface with the matrix, so it was possible to have both high toughness (exceeding 200 J) and high hardness (36 HRC or higher).

[0077] In contrast, the impact toughness of Comparative Sample CS6 was only 189 J, and the hardness was only 34.6 HRC, which did not meet both the hardness and toughness requirements necessary for use. The impact toughness of Comparative Samples CS1 to CS5 exceeded 200 J, but the corresponding hardness values were less than 36 HRC, and they could not simultaneously meet the requirements of high hardness and high toughness.

[0078] Thermal conductivity Table 4 shows the thermal conductivities calculated at each test temperature for Examples S1 to S8 and Comparative Examples CS1 to CS6.

[0079] The thermal conductivities of Examples S1 to S8 are greater than 36 W / mK in the range of 25 to 500 °C. In particular, at 300 °C, the thermal conductivities of Examples S1 to S8 exceed 40 W / mK, and the thermal conductivities of Examples S2 to S6 reach 45 W / mK or more.

[0080] Since the comparison results of the thermal conductivity among all samples are similar at various temperatures, only the thermal conductivity at 300 °C will be described here as an example. Samples S1 and S8, which only differ in Ni content, exhibit similar thermal conductivities after the same heat treatment process. This indicates that even when Ni is contained in excess to some extent, it does not cause a significant change in the thermal conductivity with respect to the Al content.

[0081] The thermal conductivities of Comparative Examples CS4 and CS6 (the ratios of Al to Ni are 0.31 and 0.27, respectively) are lower than those of all Examples S1 to S8, and the values are only 30.2 W / mK and 31.3 W / mK, respectively. This is because after the hardening treatment, about 1.14% and 0.25% of Al remains dissolved in the matrix of Comparative Examples CS4 and CS6, causing severe lattice distortion and reducing the thermal conductivity. As a result, Comparative Examples CS4 and CS6 do not exhibit both high hardness and high thermal conductivity.

[0082] Comparative Example CS5 shows a much higher thermal conductivity than CS4 and CS6. This is because all of the Al in CS5 (the ratio of Al to Ni is 0.069) has completely precipitated from the matrix, and excess Ni (with respect to the Al content) does not have a significant effect on the thermal conductivity of the steel (as seen in Examples S1 and S8). However, as described above, due to the limited precipitation of the NiAl intermetallic compound, the achieved hardness is low and does not meet the required specifications. Therefore, in Comparative Example CS5, both high hardness and high thermal conductivity cannot be achieved.

[0083] From the above, first, it is concluded that if Al does not completely precipitate from the matrix, the thermal conductivity of the matrix will decrease significantly. Based on the hardness and thermal conductivity data of various samples, it is clear that high thermal conductivity can be obtained by ensuring that Al completely precipitates from the matrix during the hardening process. Second, excessive Ni does not have a significant impact on the thermal conductivity. Therefore, on the condition that Al completely precipitates from the matrix and a sufficient amount of precipitates are formed, the Ni content can be adjusted according to the desired properties. To achieve both high thermal conductivity and high hardness, it is necessary to control the ratio of Al to Ni within the range of 0.075 to 0.234. As a result, Al completely precipitates from the matrix in the form of NiAl intermetallic compounds, and at the same time, sufficient precipitates are obtained to realize precipitation strengthening and hardening. As a result, the hardness value (H H ) after the hardening treatment exceeds 36 HRC, and the thermal conductivity exceeds 36 W / mK. Preferably, the ratio of Al to Ni is from 0.18 to 0.234, the Ni content is preferably from 5% to 9%, and more preferably, the ratio of Al to Ni is 0.18 to 0.22. Compared with the case where the ratio of Al to Ni is less than 0.18 and not less than 0.075, when the ratio of Al to Ni is 0.18 to 0.22, the same thermal conductivity can be obtained with a relatively low Ni content, Ni is saved, the cost is reduced, and the possibility of problems occurring during machining is reduced.

[0084] Hardening treatment To ensure the desired hardness and thermal conductivity of the steel, when the ratio of Al to Ni is in the range of 0.075 to 0.234, the selection of the hardening temperature is important. For example, the comparative sample CS1 has the same composition as sample S1 (the ratio of Al to Ni is 0.20 and the Cu content is 0.6%), but it was maintained at a lower temperature (460 °C) during the hardening process. On the one hand, when the hardening temperature is low, the precipitation of Cu becomes insufficient, and the precipitation of NiAl cannot be effectively promoted. On the other hand, when the hardening temperature is low, sufficient thermodynamic driving force for the precipitation of NiAl cannot be provided. Due to the combined effect of these two factors, the precipitation of NiAl becomes insufficient, and about 0.21% of Al remains in the matrix of the comparative sample CS1 after hardening. As a result, the hardening effect is limited, and the hardness of the obtained steel is reduced. On the other hand, Al in the matrix causes lattice strain and significantly reduces the thermal conductivity, which is only 32.3 W / mK at 300 °C.

[0085] The comparative sample CS2 has the same composition as sample S3 (the ratio of Al to Ni is 0.18 and the Cu content is 1.0%), but it was maintained at a high temperature (620 °C) during the hardening process. The high temperature promotes the precipitation and rapid coarsening of Cu, and also promotes the precipitation and coarsening of NiAl. Although Al can be completely precipitated (the Al content in the matrix of the comparative sample CS2 is 0.02%), due to the coarsening of Cu particles and NiAl precipitates, their strengthening effect and hardening effect are reduced, and the hardness of the comparative sample CS2 is significantly reduced compared to sample S3.

[0086] The comparative sample CS3 has the same composition as sample S4 (the ratio of Al to Ni is 0.21 and the Cu content is 1.7%), but it was maintained at a high temperature (650 °C) during the hardening process. Since the Cu content is high and the hardening temperature is also high, the precipitation and coarsening of Cu particles are promoted at high temperatures. This promotes the complete precipitation of Al as NiAl precipitates (the Al content in the matrix of the comparative sample CS3 is 0.02%), but also causes the coarsening of NiAl precipitates, resulting in a reduction in the strengthening effect and hardening effect. Therefore, the hardness of the comparative sample CS3 is reduced, and H H is only 33.6 HRC.

[0087] It is worth noting that the thermal conductivities of the comparative sample CS2 and the comparative sample CS3 are as high as 47 W / mK and 50 W / mK, respectively. This is because in both cases, Al has completely precipitated from the matrix. On the one hand, the coarsening of the Cu particles and NiAl precipitates does not change the thermal conductivity. On the other hand, at high temperatures, the dislocations in the matrix are significantly recovered, improving the thermal conductivity of the matrix. However, as mentioned above, since the hardening temperature is too high, the hardness of both samples drops below 35 HRC and they cannot meet the usage requirements. In the comparative sample CS2 and the comparative sample CS3, it is impossible to achieve both high hardness and high thermal conductivity simultaneously.

[0088] From the above, it can be seen that when the composition and element ratio of the steel for this application meet the specified requirements, if the hardening temperature is not properly selected, it is impossible to achieve both high hardness and high thermal conductivity in the steel. Combining the data of various samples and comparative samples, it can be seen that when the Cu content is low, the precipitated Cu particles can only limitedly promote the precipitation of NiAl. In this case, to ensure sufficient precipitation of NiAl from the matrix, a higher hardening temperature that provides a greater thermodynamic driving force for NiAl precipitation is required to obtain high hardness (see Sample S3 and Sample S5). Similarly, when the Cu content is high, the precipitated Cu particles can effectively promote the precipitation of NiAl. In this case, in order to ensure the precipitation strengthening and hardening effects, it is necessary to use a lower hardening temperature to prevent the coarsening of the Cu particles and NiAl precipitates (see Sample S4).

[0089] Therefore, in order for the steel to have the desired hardness and thermal conductivity, it is necessary to select an appropriate hardening temperature according to the range of Cu, promote the precipitation of the NiAl intermetallic compound by Cu, and prevent the coarsening of the precipitated Cu particles and NiAl intermetallic compound during the hardening process. Based on experimental data, when Cu is 0.4% or more and less than 0.8%, the hardening temperature ranges from 560 to 650 °C; when Cu is 0.8% or more and less than 1.5%, the hardening temperature ranges from 470 to 580 °C; when Cu is 1.5% or more and 1.99% or less, the hardening temperature is defined to range from 450 to 520 °C.

[0090] Summarizing the above experimental results, when the composition and element ratio of the steel satisfy a predetermined range, by setting the ratio of Al to Ni in the range of 0.075 to 0.234 and selecting an appropriate hardening temperature, it can be seen that steel with high hardness, high toughness, and high thermal conductivity can be obtained. Summarizing the main reasons, it is as follows: 1. Since Ni is a parent ferroalloy element, when dissolved in the iron matrix in an appropriate amount, it does not reduce the thermal conductivity of the iron matrix; 2. During the hardening process, the Al element combines with the Ni element to form a NiAl intermetallic compound represented by the chemical formula Ni(Fe,Al). The Al element dissolved in the matrix significantly reduces the thermal conductivity of the matrix, and the expected hardening effect cannot be obtained. By setting the ratio of Al to Ni in the range of 0.075 to 0.234, it is ensured that the Al element is completely precipitated from the iron matrix after hardening. Furthermore, even if the Ni content is moderately increased, it will not significantly improve or reduce the technical effects of the present invention such as hardness, impact toughness, and thermal conductivity. Therefore, preferably, the ratio of Al to Ni is 0.18 to 0.234, and Ni: 4.5 to 11%. More preferably, the ratio of Al to Ni is 0.18 to 0.22, and Ni: 5 to 9%. In a preferred scheme, equivalent or better technical effects can be obtained with a lower Ni content; 3. The precipitation of Cu provides nucleation sites for the precipitation of NiAl, thereby promoting the precipitation of NiAl; 4. Since a large driving force is required for the nucleation and growth of NiAl and Cu, by selecting an appropriate hardening temperature in the range of 450 - 650 °C according to the Cu content, while promoting the precipitation of the precipitate phase, it is possible to prevent the precipitate phase from becoming overly coarsened, whereby both the hardness and the thermal conductivity can meet the usage requirements. For example, in the case of sample S6 with a low Cu content, by applying a high hardening temperature to impart a large thermodynamic driving force to the NiAl precipitation, the complete precipitation of the NiAl intermetallic compound can be promoted without relying on the precipitation of Cu, and the hardness and thermal conductivity can be ensured; 5. The high thermal conductivity of the Ni(Fe, Al) precipitate, the high thermal conductivity of the Cu precipitate, and the clean matrix all contribute to the high thermal conductivity of the steel of the present invention.

[0091] Deterioration analysis As shown in Table 5, in the case of sample S7 (the ratio of Cu to Al is 1.6), based on the Cu content, a hardening temperature of 520 °C was selected. When the hardening time was 4 hours, the obtained hardened hardness value H H was 42.3 HRC. When the hardening time was extended to 24 hours and 48 hours, H H decreased significantly, and the deterioration degrees were 16.3% and 24.8% respectively. The precipitation temperature of Cu is slightly lower than that of NiAl and is more susceptible to deterioration. When the ratio of Cu to Al is greater than 1, the coarsening of a large number of Cu particles causes NiAl to also coarsen, and the coarsening of the second-phase particles Cu and NiAl leads to the deterioration of the steel.

[0092] On the other hand, in the case of sample S4 (the ratio of Cu to Al is 0.94), the hardened hardness value H H obtained when the hardening time was 4 hours was 50.2 HRC. Although H H decreases with the extension of the hardening time, H HIt was still high (45.7 HRC) and met the usage requirements, and the deterioration phenomenon was not significant. On the other hand, for sample S4, in the present invention, the overall Cu content was reduced to 0.4 - 1.99% (such as S4 and S7 with Al being 1.7% and 1.6% respectively). By selecting an appropriate hardening temperature according to the Cu content, the rapid growth of the precipitated Cu particles was prevented, and the rapid growth of the precipitated NiAl was promoted, so the deterioration resistance was good. On the other hand, by controlling the ratio of Cu to Al in the range of 0.5 - 1 (excluding the end values), the NiAl precipitation phase and the Cu particles suppressed each other's deterioration process, and the degree of deterioration was controlled within 10% even after a long-term hardening treatment, effectively improving the deterioration resistance of the steel.

[0093] From the above, in the heat treatment of the hot working tool steel of the present invention, alloying elements dissolved in the solid solution, especially the Al element, can be completely precipitated from the matrix. In addition, precipitates such as Cu particles and NiAl intermetallic compounds not only have good thermal conductivity but also play a strengthening role, so both the thermal conductivity and hardness of the hot working tool steel increase after the hardening treatment. Furthermore, the manufacturing process of the hot working tool steel of the present invention is simple, thus reducing the manufacturing cost and enabling production with existing heat treatment and processing equipment.

[0094] The hot working tool steel of the present invention can be used for hot stamping dies for steel plates, aluminum alloy die-castings, hot extrusion drawing dies, etc.

[0095] The above embodiments and experimental data are for illustrative purposes of the present invention. It is understood by those skilled in the art that the present invention is not limited to these embodiments and various modifications can be made without departing from the scope of the present invention.

Claims

1. A hot working tool steel, the chemical composition of which, by weight percentage, is: Ni: 4.5 to 12%, Al: 0.8 to 2.1%, Cu: 0.4 to 1.99%, Mn: 0.001 to 1%, C: 0.0001 to 0.1%, where the ratio of Al to Ni is 0.075 to 0.234, and the balance is Fe and unavoidable impurities, the hot working tool steel.

2. The hot working tool steel according to claim 1, wherein the ratio of Al to Ni is 0.075 or more and less than 0.

18.

3. The hot working tool steel according to claim 1, wherein the ratio of Al to Ni is 0.18 or more and 0.234 or less.

4. The hot working tool steel according to claim 3, wherein the ratio of Al to Ni is 0.18 or more and 0.22 or less, and Ni is 5 to 9%.

5. The hot working tool steel according to any one of claims 1 to 4, wherein the ratio of Cu to Al is more than 0.5 and less than 1, and / or Cu is 0.4 to 1.5%.

6. The hot working tool steel according to claim 5, wherein the ratio of Cu to Al is 0.6 or more and 0.9 or less, and / or Cu is 0.4 to 1.2%.

7. The hot working tool steel according to any one of claims 1 to 4 and 6, wherein C is 0.0001 to 0.05%.

8. A hardening step, wherein when Cu is 0.4% or more and less than 0.8%, the hot working tool steel is heated to a temperature in the range of 560 to 650 °C, when Cu is 0.8% or more and less than 1.5%, the hot working tool steel is heated to a temperature in the range of 470 to 580 °C, when Cu is 1.5% or more and 1.99% or less, the hot working tool steel is heated to a temperature in the range of 450 to 520 °C, held for 1 to 240 hours, and then cooled to room temperature, the step. The heat treatment method of the hot working tool steel according to any one of claims 1 to 7, including the step.

9. When Cu is 0.4% or more and less than 0.8%, the hot working tool steel is heated to a temperature in the range of 560 to 620 °C, when Cu is 0.8% or more and less than 1.5%, the hot working tool steel is heated to a temperature in the range of 480 to 580 °C, when Cu is 1.5% or more and 1.99% or less, the hot working tool steel is heated to a temperature in the range of 450 to 500 °C, held for 4 to 24 hours, and then cooled to room temperature. The heat treatment method according to claim 8, including the step.

10. A solution heat treatment step, wherein, before the hardening step, the hot working tool steel is heated to a temperature in the range of 850 to 950 °C and held for 0.5 to 240 hours, and then cooled to room temperature. The heat treatment method according to claim 8 or 9, further comprising this.

11. A forging step, wherein, before the solution heat treatment step, the hot working tool steel is heated to a temperature in the range of 900 to 1150 °C for forging, and then cooled to room temperature. The heat treatment method according to claim 10, further comprising this.

12. The heat treatment method according to claim 10 or 11, wherein the cooling in the solution heat treatment step and / or the forging step is first slowly cooled to less than 600 °C at a rate of 60 °C / hour or less and then air-cooled to room temperature.

13. The heat treated hot working tool steel exhibits the following properties: hardness of 36 HRC or more, room temperature thermal conductivity of 34 to 50 W / mK, notched impact energy of 200 J or more, and Al dissolution amount in the matrix of 0.05% or less. The heat treatment method according to claim 8 or 9.

14. A hot working tool made of hot working tool steel obtained by the heat treatment method according to any one of claims 8 to 13, wherein the hot working tool can be used as a hot stamping die for steel sheets, an aluminum alloy die casting die, or a hot extrusion drawing die. A hot working tool, characterized in that.

Images