Hot work die steel, heat treatment method thereof and hot work die

Abstract

The invention relates to hot work die steel which comprises the following chemical components in percentage by weight: 0.10% < C < = 0.35%, 5.5% < = Ni < = 8.0%, 0.5% < = Cu < = 1.5%, 0.8% < = Al < = 1.4%, Mo < = 0.2%, W < = 0.2%, V < = 0.2%, Nb < = 0.2%, 0.001% < = Mn < = 1%, Cr < = 0.4%, Si < = 0.4%, P < 0.05%, S < 0.015%, N < 0.015% and the balance of Fe and inevitable impurities.

Description

Technical Field

[0001] This invention relates to a hot work die steel, its heat treatment method, and a hot work die. Background Technology

[0002] Hot work die steel is an alloy tool steel in which different weight percentages of alloying elements such as carbon, silicon, chromium, tungsten, molybdenum, nickel, manganese, vanadium, and cobalt are selectively added to iron. It is often used in hot forging, hot extrusion, hot stamping, and die casting of aluminum and magnesium alloys. In particular, hot work die steel is often used to manufacture aluminum alloy differential pressure and low-pressure dies.

[0003] After precision machining, aluminum alloy differential pressure and low-pressure molds require sandblasting. Following sandblasting, a release agent layer several hundred micrometers thick is applied to the mold surface. Sandblasting increases the surface roughness of the mold, improves the adhesion of the release agent, and slows down its shedding during the die-casting process. The release agent serves two purposes: firstly, it facilitates the removal of aluminum alloy parts from the mold after die-casting; secondly, it isolates the molten aluminum from direct contact with the mold, preventing the aluminum from sticking to it.

[0004] During repeated die casting, the release agent may detach. For small, localized detachments, repairs can be made by spraying additional release agent. However, when a large area of ​​the release agent detaches, the remaining release agent on the mold surface must be cleaned using sandblasting, and the mold can only be reused after re-coating with the new release agent.

[0005] During the service life of a mold, the release agent on the mold surface is typically removed multiple times through sandblasting. Each sandblasting process causes the mold profile to shrink and the cavity to enlarge, leading to increased aluminum alloy usage, increased size and weight of aluminum alloy parts, and consequently increased part costs and defect rates. When the mold cavity enlarges to a certain extent and the increased part weight causes part costs to rise to a point where there is no profit, the mold must be scrapped. This is the main failure mode of aluminum alloy differential pressure and low-pressure molds.

[0006] Therefore, it is hoped that the hardness of aluminum alloy differential pressure and low-pressure molds can be increased to reduce the degree of mold surface descent during each sandblasting process, thereby increasing the mold life.

[0007] It is generally believed that increasing the carbon content (C) of mold steel can improve its hardness, but it will also reduce its toughness. During the machining and heat treatment of mold steel, if the toughness is low, it may crack during machining or during quenching in heat treatment, leading to a shortened mold life and, in severe cases, the mold may be scrapped. For example, H13 steel, when used in differential pressure or low-pressure casting molds, typically has a hardness of at most 38-42 HRC. If the hardness is further increased by increasing the C content, the mold may crack during manufacturing and use due to excessively low toughness, causing premature mold failure. Therefore, existing mold steels face the technical challenge of simultaneously improving hardness and toughness.

[0008] In addition to the expectation of higher hardness and toughness, aluminum alloy differential pressure and low-pressure molds are also expected to have better thermal conductivity.

[0009] Differential pressure and low-pressure casting processes for aluminum alloys are typically used to produce thick-walled parts, and their solidification and cooling require a long time, resulting in a slow production cycle and impacting production efficiency. Improving the efficiency of differential pressure and low-pressure casting is a pressing need in the industry. By improving the thermal conductivity of the mold and accelerating the solidification and cooling rate of the aluminum alloy during casting, the cycle time of differential pressure and low-pressure casting can be significantly reduced, improving part production efficiency. However, it is generally believed that increasing the hardness of the mold often accompanies a decrease in thermal conductivity. This is because the strengthening of mold steel is usually achieved by increasing defects, such as point defects (vacancies), line defects (dislocations), surface defects (fine-grain strengthening), and bulk defects (precipitation strengthening). These defects all increase the resistance to electron movement, thereby reducing the thermal conductivity of the mold steel.

[0010] Therefore, there is a current demand for a mold steel that can simultaneously achieve high hardness, high toughness, and high thermal conductivity, especially suitable for manufacturing aluminum alloy differential pressure and low-pressure molds.

[0011] US 2012063946 A1 discloses a hot-working tool steel. It has the following weight percentage composition: Ceq=0.20~1.2%, C=0.20~1.2%, N=0~1%, B=0~1%, Cr<1.5%, Ni=1.0~9%, Si<0.4%, Mn=0~3%, Al=0~2.5%, Mo=0~10%, W=0~15%, Ti=0~3%, Ta=0~3%, Zr=0~3%, Hf=0~3%, V=0~4%, Nb=0~3%, Cu=0~4%, Co=0~6%, S=0~1%, Se=0~1%, Te=0~1%, Bi=0~1%, As=0~1%, Sb=0~1%, Ca=0~1%, with the balance consisting of Fe and unavoidable impurities, which are in the range of %Ceq=%C+0.86*%N+1.2*%B, %Mo+1.2%·W>1.2.

[0012] This patent teaches that after solution treatment and hardening, Mo and W carbides replace Cr carbides, thereby improving the thermal conductivity of alloy tool steel. However, the size of Mo and W carbides is difficult to control, especially the size of undissolved primary carbides after solution treatment, which is ~3 μm. These large carbides become fatigue crack initiation sites, severely affecting not only the fatigue life of the tool steel but also significantly deteriorating its toughness. Domestic researchers have found that although the thermal conductivity of this high thermal conductivity alloy tool steel is ~47 W / mK at room temperature, its thermal conductivity decreases with increasing temperature. At 300℃, its thermal conductivity is below 39 W / mK, and at 500℃, its thermal conductivity is only 35 W / mK. Furthermore, when the hardness value reaches 50 HRC or above, the impact energy (7×10 mm unnotched specimen) is <210 J. When used at high temperatures, this high thermal conductivity alloy tool steel loses its advantage of high thermal conductivity and fails to achieve a good performance match of high thermal conductivity, high toughness and high hardness.

[0013] CN 111636037 B discloses an ultra-high strength ferritic steel reinforced with nano-intermetallic compounds. The chemical composition of this ultra-high strength ferritic steel, by weight percentage, is as follows: C: 0~0.2%, Cu: 0.5~5%, Ni: 0.01~4%, Mn: 0.01~4%, Al: 0.001~2%, Cr: 0~12%, Mo: 0~3%, W: 0~3%, Mo+W not less than 0.05%, V: 0~0.5%, Ti: 0~0.5%, Nb: 0~0.5%, V+Ti+Nb not less than 0.01%, Si: 0~1%, B: 0.0005~0.05%, P not more than 0.04%, S not more than 0.04%, N not more than 0.04%, O not more than 0.05%, with the balance being Fe and unavoidable impurities.

[0014] This patent teaches that solution treatment and aging treatment after rolling can produce ferritic steel primarily strengthened by nano-intermetallic compounds, combined with fine grains, solution treatment, and dislocation strengthening, resulting in excellent toughness, weldability, and corrosion resistance. However, because most of the alloying elements remain dissolved in the ferrite matrix after the final heat treatment, this ultra-high strength ferritic steel has low thermal conductivity. For example, Cr dissolves in the ferrite matrix to form Cr carbides, which not only have low thermal conductivity but also tend to coarsen during heat treatment, thus deteriorating the material's toughness. Similarly, Si, Cr, and Al elements, when dissolved in ferrite, significantly reduce the thermal conductivity of ferrite. Furthermore, the carbonitrides of microalloying elements V, Ti, and Nb not only reduce the steel's thermal conductivity but also its toughness.

[0015] CN 114807774 B discloses a hot work die steel whose chemical composition, by weight percentage, includes Ni: 4.5~12%, Al: 0.8~2.1%, with Ni:Al=0.075~0.234, Cu: 0.4~1.99%, C: 0.0001~0.1%, Mn: 0.001~1.0%, with the remainder being Fe and impurities.

[0016] This patent teaches that by controlling NiAl precipitation through a simple hardening process, ensuring that the Al content in the matrix is ​​≤0.05%, not only high hardness and high toughness are guaranteed, but the thermal conductivity of hot work die steel is also improved. The invention also teaches that by maintaining a Cu:Al ratio of 0.5 < Cu:Al < 1, preferably Cu: 0.4~1.5%, the resistance to hardening of hot work die steel can be improved. The patent further teaches that by omitting spheroidizing annealing of the die steel, lowering the solution treatment temperature (850~950℃), and designing a low carbon content, cementite will not grow during the hardening process. The patent also teaches that carbon (C) is one of the most effective and economical strengthening elements in steel and is a stabilizing austenite element. C is an interstitial solid solution element, and its strengthening effect is far greater than that of substitutional solid solution elements. C can improve the hardenability of steel, and the formed cementite or alloy carbides significantly increase the hardness of the alloy. When the C content increases, network carbides and coarse cementite will form along the grain boundaries, severely reducing the toughness of the steel. Since this patent is based on NiAl and Cu precipitation strengthening and does not rely on carbide strengthening and hardening, the C content is between 0.0001 and 0.1%.

[0017] Both CN 111636037 B and CN 114807774 B employ Cu+NiAl precipitation strengthening to improve thermal conductivity and develop new hot-work die steels. However, carbides and carbon elements both negatively impact thermal conductivity. For low-pressure and differential-pressure casting dies of aluminum alloys, a hardness of 42 HRC is clearly insufficient to improve die life. To further increase hardness, it is necessary to increase the content of Cu, Ni, and Al elements or increase carbide precipitation. It is generally believed that increasing carbon content to improve hardness will decrease toughness and thermal conductivity. If the content of Cu, Ni, and Al elements is increased while the carbon content is low (the carbon content of both patents is ≤0.2%), the steel is prone to blocky ferrite and bainite phase transformations during quenching. The coarse grain size will drastically reduce the toughness of the steel, increasing the risk of overall die cracking. Summary of the Invention

[0018] This invention is made in view of the aforementioned problems existing in the prior art. One object is to provide a hot work die steel that possesses high hardness, high toughness, and high thermal conductivity. The chemical composition of the hot work die steel, by weight percentage, is: 0.10 < C ≤ 0.35%, 5.5% ≤ Ni ≤ 8.0%, 0.5% ≤ Cu ≤ 1.5%, 0.8% ≤ Al ≤ 1.4%, Mo ≤ 0.2%, W ≤ 0.2%, V ≤ 0.2%, Nb ≤ 0.2%, 0.001% ≤ Mn ≤ 1%, Cr ≤ 0.4%, Si ≤ 0.4%, P < 0.05%, S < 0.015%, N < 0.015%, with the remainder being Fe and unavoidable impurities.

[0019] The basic inventive concept of this invention is to achieve a synergistic improvement in hardness and toughness while not significantly reducing thermal conductivity by moderately increasing the C content to greater than 0.1 wt.% and combining it with a specific heat treatment process (quenching and tempering).

[0020] Common knowledge teaches that lower carbon (C) content in martensite results in better toughness and thermal conductivity, but also lower hardness, making the steel unsuitable for certain applications. Increasing C content enhances solid solution strengthening and increases hardness, but the toughness of martensite decreases with increasing C content. Furthermore, increasing C content also increases lattice distortion in the iron matrix, exacerbating electron and phonon scattering and reducing thermal conductivity. Therefore, based on existing knowledge, to achieve both high toughness and high thermal conductivity, it is necessary to reduce the C content in martensite.

[0021] This invention discovers that increasing the carbon content to greater than 0.1 wt.% can simultaneously improve the hardness and toughness of mold steel without causing a decrease in toughness due to increased hardness, which is contrary to the teachings of existing knowledge. Furthermore, by rationally controlling the carbon content and heat treatment process, the thermal conductivity decreases only slightly while hardness increases.

[0022] Regarding the toughness of martensitic steel, existing knowledge teaches that lower carbon content results in better toughness. However, this invention discovers that lower carbon content does not necessarily lead to better martensitic toughness; rather, when the carbon content falls below 0.1 wt.%, the toughness of martensitic steel decreases significantly. On one hand, low carbon content results in a higher martensitic transformation temperature, leading to a lower dislocation density. This lower dislocation density results in lower work hardening capacity. Since the toughness of steel is positively correlated with work hardening, lower dislocation density leads to weaker work hardening capacity and poorer toughness. On the other hand, low carbon content causes austenite grains to coarsen more easily during forging. During forging, carbon atoms can, to some extent, inhibit the dynamic recrystallization of austenite and suppress grain growth. Carbon content above 0.1 wt.% can refine the original austenite grains and strengthen austenite while lowering the martensitic transformation temperature, thus promoting the refinement of lath martensite packets. These martensitic packets, acting as large-angle grain boundaries within martensite, can improve both martensite strength and toughness. When the carbon content is below 0.1 wt.%, it leads to packet coarsening, failing to inhibit crack propagation and thus reducing martensite toughness. Therefore, increasing the carbon content to above 0.1 wt.% can increase the martensite dislocation density, refine the original austenite grains, and further refine the packet size within the martensite that determines strength and toughness, thereby improving martensite toughness.

[0023] Once the carbon content exceeds 0.1 wt.%, as commonly known, strength increases further with increasing carbon content, but toughness gradually decreases. This is because high carbon content not only pins dislocations, making them immobile and thus reducing toughness, but also leads to the formation of twinned martensite, which significantly reduces toughness.

[0024] Common knowledge also teaches that increasing the carbon (C) content decreases thermal conductivity. This is because C in steel primarily achieves high hardness through solid solution strengthening (C atoms dissolve into the iron crystal lattice as interstitial atoms, inducing strong local lattice distortion, thereby effectively hindering dislocation movement and significantly increasing the material's strength and hardness), dislocation strengthening (austenite is quenched into lath martensite, forming a large number of dislocations), grain refinement strengthening (austenite grains are sheared into multiple packets by martensitic phase transformation), and precipitation strengthening (C atoms precipitate in the form of carbide particles). All strengthening mechanisms introduce defects into the steel, but higher hardness introduces more defects, which significantly reduces the steel's thermal conductivity.

[0025] The thermal conductivity κ of steel is mainly determined by electronic thermal conductivity κ e and lattice thermal conductivity κ l Composition, with electronic thermal conductivity κ at room temperature e Primarily, the solid-solidified carbon atom, acting as a point defect, scatters free electrons, reducing the mean free path λ, thus affecting κ. e ∝ λ decreases. According to Matthiassen's rule, the impurity scattering rate is 1 / τ i The thermal conductivity decreases more significantly with increasing C concentration; the higher the C content, the greater the decrease.

[0026] This invention cleverly controls the distribution and phase transformation behavior of carbon (C) through targeted heat treatment steps. This effectively suppresses the adverse effect of decreased thermal conductivity caused by increased C content while simultaneously increasing hardness, thus achieving a balance between high strength and toughness. It minimizes the negative impact of increased C content. The heat treatment steps according to this invention include quenching and tempering.

[0027] A key aspect of this invention is that by using a C content greater than 0.1 wt.% combined with a quenching process to form martensite with a high dislocation density, C is promoted to accumulate on dislocations, i.e., to form a Cottrell atmosphere. This increases the concentration of C-rich regions on dislocations and essentially eliminates the content of free dissolved C in the matrix. As a result, the thermal conductivity does not decrease significantly due to the high C content in the matrix.

[0028] Another key aspect of this invention is that, during tempering, the dissolved C and C enriched on dislocations precipitate as nano-cementite, further "extracting" the dissolved and dislocation-enriched C, thereby further reducing the C content in the matrix and improving thermal conductivity compared to the quenched state. Simultaneously, due to the still high dislocation density after tempering, a large amount of C remains enriched on the dislocations, increasing the hardness of the steel of this invention.

[0029] This invention, through targeted heat treatment steps, not only preserves the strengthening effect of C, but also reverses as much as possible the negative impact of increasing C content on thermal conductivity and toughness.

[0030] Specifically, in the quenching process of this invention, the steel is first heated to the austenitic region, held at that temperature to allow C to fully dissolve into the austenite, and then rapidly cooled (water quenching, oil quenching, or gas quenching) to room temperature. A C content higher than 0.1 wt.% improves the stability of austenite (i.e., compared to a C content less than 0.1 wt.%, it lowers the martensitic transformation initiation temperature). M s (Point), and the carbon content should not exceed 0.35 wt.% to avoid the formation of excessive twinned martensite.

[0031] During the quenching process, austenite (FCC structure) transforms into martensite via a shear transformation mechanism. This transformation causes a volume expansion (approximately 4%) and generates a large number of dislocations (density up to 10). 12 ~10 14 m -2 It is far higher than the 10% of the traditional annealed state. 8 m -2 To accommodate supersaturated carbon atoms. These dislocations are not only strengthening sources (i.e., dislocation strengthening, hardness increment ΔH∝) , ρ (This is the dislocation density), and it also acts as a "trap" for C atoms, causing carbon atoms to accumulate at these dislocations.

[0032] In the quenching process of this invention, C atoms accumulate in the dislocation core region, forming a Cotillard atmosphere, which is a short-range diffusion process. The diffusion coefficient of C (DC ≈ 10) -7 m 2 (At 400℃ / s) allows C to migrate to dislocations, and the stress field of the dislocations lowers the chemical potential of C, further promoting segregation. As a result, the concentration of free dissolved C in the matrix decreases, while the concentration of C-rich regions on dislocations can reach 1wt.%~2wt.%. This enrichment of C reduces the content of free dissolved C in the matrix (from supersaturated solution at the beginning of quenching to local equilibrium). Free dissolved C is a major negative factor for thermal conductivity, and its large electron cross-section leads to κ e The thermal conductivity of the steel decreases because C atoms are "captured" by dislocations, reducing the concentration of free dissolved C in the matrix, increasing the mean free path of electrons, and thus improving the thermal conductivity of the steel.

[0033] In the heat treatment process of this invention, a C content of 0.1~0.35 wt.% combined with the quenching process promotes the refinement of lath martensite packets (crystal regions). As large-angle grain boundaries within martensite, the martensite packets (crystal regions) improve the strength and toughness of martensite.

[0034] During the tempering process of this invention, a large number of dislocations are still retained (the high dislocation density remains after tempering), which is the key to achieving high toughness and high thermal conductivity. Simultaneously, Ni and Al combine to precipitate almost completely in the form of nanoscale NiAl intermetallic compounds, maintaining a coherent interface with the iron matrix; Cu combines to precipitate almost completely in the form of nanoscale ε-Cu, also maintaining a coherent interface with the iron matrix; a C content of 0.1~0.35 wt.% causes C to combine with Fe to precipitate nano-cementite (θ-Fe3C). High dislocation density + nearly completely precipitated NiAl and Cu phases coherent with the iron matrix + pure iron matrix (because Ni, Al, and Cu are completely precipitated) + C enriched on dislocations (rather than in the octahedral interstices of martensite) + high thermal conductivity θ-Fe3C precipitate + refined martensite packets (crystal regions) achieve improved toughness while maintaining high hardness, with minimal decrease in thermal conductivity.

[0035] The first aspect of the present invention relates to a hot work die steel having the following chemical composition by weight percentage: 0.10 ≤ C ≤ 0.35%, 5.5 ≤ Ni ≤ 8.0%, 0.5 ≤ Cu ≤ 1.5%, 0.8 ≤ Al ≤ 1.4%, Mo ≤ 0.2%, W ≤ 0.2%, V ≤ 0.2%, Nb ≤ 0.2%, 0.001 ≤ Mn ≤ 1%, Cr ≤ 0.4%, Si ≤ 0.4%, P < 0.05%, S < 0.015%, N < 0.015%, with the remainder being Fe and unavoidable impurities.

[0036] In another preferred embodiment, 0.15% ≤ C ≤ 0.30%. This essentially avoids the formation of twinned martensite.

[0037] In another preferred embodiment, the dislocation density after tempering is ≥4.261. 14 / m 2 .

[0038] In another preferred embodiment, the martensitic packet size is ≤4.79 μm.

[0039] In another preferred embodiment, the hardness is 46.3~52.3 HRC, preferably 48.5~51.4 HRC.

[0040] In another preferred embodiment, the thermal conductivity in the temperature range of 200~420°C is ≥36.24W / m·K, preferably ≥38.23W / m·K.

[0041] In another preferred embodiment, the impact energy of the U-notch specimen is ≥12.2J, preferably ≥20.1J.

[0042] The second aspect of the present invention relates to a heat treatment method for the hot work die steel described above, comprising: a quenching step of holding the hot work die steel at 900~1000°C for 0.1 to 72 hours and then quenching it to room temperature.

[0043] In a preferred embodiment, the quenching step is followed by a tempering step, in which the temperature is held at 420-520°C for 0.1 to 144 hours, and then cooled to room temperature.

[0044] In another preferred embodiment, the properties of the heat-treated hot work die steel are as follows: hardness 46.3~52.3HRC, thermal conductivity ≥36W / m·K in the temperature range of 200~420℃, and impact energy of U-notch specimen ≥12.2J.

[0045] The third aspect of the invention relates to a hot work die made of hot work die steel obtained by the above-described heat treatment method, which can be used as a hot stamping die for steel plates, a die forging die for aluminum alloys, a hot forging die, or a hot extrusion drawing die.

[0046] In a preferred embodiment, the hot work die can be used as an aluminum alloy differential pressure or low-pressure casting die.

[0047] Furthermore, those skilled in the art will understand that any range or specific value within the aforementioned ranges is applicable to the present invention. For example, 0.075 to 0.18 (excluding 0.18) includes any sub-range and any specific value therein, such as any value of 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, 0.175, etc., and any range of 0.075 to 0.085, 0.08 to 0.098, 0.09 to 0.14, 0.12 to 0.17, etc. Similarly, 0.18 to 0.234 includes any sub-range and any specific value within it, such as any value of 0.18, 0.185, 0.19, 0.195, 0.198, 0.203, 0.207, 0.213, 0.215, 0.225, etc., and any range of 0.18 to 0.215, 0.18 to 0.207, 0.185 to 0.225, 0.19 to 0.225, 0.19 to 0.22, etc. Similarly, 0.5 to 1 (excluding the extreme values) includes any subrange and any specific value within it, such as any value of 0.56, 0.62, 0.69, 0.73, 0.80, 0.88, 0.90, 0.92, 0.98, and any range of 0.51 to 0.98, 0.56 to 0.92, 0.62 to 0.98, 0.65 to 0.92, 0.65 to 0.9. Unless otherwise specified, all ranges include the extreme values. Attached Figure Description

[0048] Figure 1 The microstructure of B1 steel after quenching is shown. Figure 2 A schematic diagram showing the measurement of the Packet intercept size for martensitic phase transformation is provided. Figure 3 This illustrates the effect of C content on the martensitic packet of the steel of the present invention; Figure 4 This shows the effect of C content on the dislocation density of the steel after quenching. Figure 5 This illustrates the effect of C content on the dislocation density of the steel after tempering. Figure 6a The cementite morphology of example steel A1 is shown; Figure 6b The cementite morphology of example steel A3 is shown; Figure 6c The cementite morphology of example steel A5 is shown. Detailed Implementation

[0049] The present invention will now be described in more detail with reference to exemplary embodiments.

[0050] The chemical composition of the hot work die steel involved in this invention, by weight percentage, is as follows: 0.10 ≤ C ≤ 0.35%, 5.5 ≤ Ni ≤ 8.0%, 0.5 ≤ Cu ≤ 1.5%, 0.8 ≤ Al ≤ 1.4%, Mo ≤ 0.2%, W ≤ 0.2%, V ≤ 0.2%, Nb ≤ 0.2%, 0.001 ≤ Mn ≤ 1%, Cr ≤ 0.4%, Si ≤ 0.4%, P < 0.05%, S < 0.015%, N < 0.015%, with the remainder being Fe and unavoidable impurities.

[0051] The functions and proportions of each element in this invention are described below.

[0052] C: Compared to traditional H13 steel, this invention, by selectively setting the C content (0.1% < C ≤ 0.35%), achieves two key improvements. First, it fosters high dislocation density martensite while largely preventing the formation of twinned martensite and refining the martensite packet structure, thus improving matrix toughness. Second, the high dislocation density martensite promotes C segregation at dislocations, thereby suppressing the negative impact of free dissolved C atoms in the matrix on thermal conductivity, ensuring that thermal conductivity does not significantly decrease with increasing C content. When the C content is below 0.1%, the dislocation density in martensite decreases significantly, leading to a decrease in matrix toughness. When the C content is above 0.1%, the matrix toughness increases significantly with the increase in dislocation density in martensite. When the carbon content is below 0.05%, the supercooled austenite cannot undergo martensitic transformation during quenching. Instead, it leads to blocky ferrite or lath-like bainitic ferrite, further reducing toughness. More preferably, when the carbon content is ≥0.15wt%, the high hardness requirement of this invention is met, resulting in higher hardness and a good balance between strength and toughness. When the carbon content exceeds 0.35%, the dislocation density and martensite packet size do not change significantly with increasing carbon content, but more twinned martensite is produced, leading to a significant decrease in toughness. Therefore, more preferably, the carbon content is ≤0.30%. In this alloy system, the formation of twinned martensite can be basically suppressed, achieving a better balance between hardness and toughness. Furthermore, increasing the carbon content further brings more negative effects, such as reduced thermal conductivity, easier formation of twinned martensite leading to reduced toughness, and precipitation of large cementite during tempering, further reducing toughness.

[0053] Ni: The main role of Ni in this invention is to form NiAl intermetallic compounds, which play a role in precipitation strengthening and hardening. Furthermore, this invention found that when the mass fraction of Ni is between 5.5% and 8.0%, the packet size gradually decreases with increasing Ni content, meaning the packets are gradually refined, which improves the toughness of the steel matrix. When the Ni content is higher than 8.0%, the packet size hardly changes. In addition, Ni can suppress the formation of liquid phases at grain boundaries of Cu at high temperatures, which leads to hot cracking during high-temperature deformation of the alloy, thus ensuring the hot forming performance of the alloy. The alloying element Ni can improve the hardenability of steel, suppress bainite transformation, and Ni enriched at grain boundaries can improve toughness. When it is desired that the steel uses the precipitated intermetallic compound NiAl as the main strengthening and hardening phase, an excess of Ni is required; otherwise, Al cannot fully combine with Ni to precipitate from the matrix, resulting in Al dissolving in the matrix, which severely reduces the thermal conductivity of the steel. Under the condition of complete Al precipitation, excess Ni dissolved in the matrix does not lead to significant lattice distortion. However, excessive Ni will cause M s If the temperature is too low, the austenite cannot completely transform into martensite during the cooling process. The untransformed austenite is called retained austenite. The presence of retained austenite reduces the thermal fatigue life of the mold. The Ni content must ensure complete Al precipitation while preventing excessive Ni from leading to low martensite content. s Temperature must be controlled to ensure the packet size is small. In this invention, the Ni content of the steel is 5.5-8.0%.

[0054] Cu: In this invention, the main role of Cu is to promote the precipitation of NiAl particles. During tempering at 420~520℃, Cu particles precipitate before NiAl, providing nucleation sites for NiAl particle precipitation, thereby promoting the large-scale and sufficient precipitation of NiAl particles. This invention has found that Cu affects the martensitic packet size of the steel. When the mass fraction of Cu is 0.5~1.5%, the size of the martensitic packet gradually decreases with increasing Cu content, that is, the packet is gradually refined, which can improve the toughness of the steel matrix. When the Cu content exceeds 1.5%, the packet size increases with increasing Cu content. Cu particles have a low aging temperature. When a higher Cu content is added, the thermal stability of the steel deteriorates, and the hardness of the steel drops sharply at high temperatures, leading to premature mold failure. In addition, to suppress Cu liquidation at grain boundaries, a certain weight fraction of the alloying element Ni is usually added. Considering the combined effects of Cu on packet size refinement and precipitation strengthening, the Cu content of the steel in this invention is 0.5~1.5%.

[0055] Al: In this invention, Al combines with Ni during tempering to precipitate a large number of NiAl particles, resulting in a hardening effect. Simultaneously, Al almost completely precipitates from the Fe matrix of the steel, reducing lattice defects in the matrix and improving the thermal conductivity of the steel. Furthermore, during the quenching process of this invention, the steel does not undergo a bainitic phase transformation, resulting in a lath martensite structure after quenching. When the Al content exceeds 1.4%, a bainitic + martensite phase transformation occurs during quenching, increasing the packet size and reducing the toughness of the steel. Additionally, the NiAl intermetallic compound also possesses good thermal conductivity (~90 W / m·K), both of which contribute to improving the thermal conductivity of the steel.

[0056] To ensure that Al is completely precipitated from the matrix as NiAl intermetallic compounds, excessive Al relative to Ni content will cause the following problems: (1) Al cannot be completely precipitated from the matrix, thus reducing the thermal conductivity of the steel; (2) It is easy to form large AlN inclusions, which will not be completely dissolved in austenite at high temperatures, thus seriously impairing the toughness of the steel; (3) As a strong ferrite stabilizing element, Al will increase the A content of the steel. c1 and A c3 Temperature is crucial; when solution treatment is required, higher temperatures are necessary to achieve austenitization, increasing manufacturing costs, energy consumption, and the demands on heat treatment equipment. While low Al content ensures complete Al precipitation, insufficient NiAl intermetallic compounds result in steel hardness falling short of requirements. Conversely, excessive Ni leading to insufficient Al does not significantly affect the hardness, thermal conductivity, and impact strength of this invention, only wasting Ni and increasing material costs. Therefore, the Al content in the steel of this invention is 0.8-1.4%, forming NiAl intermetallic compounds with Ni for sufficient precipitation strengthening while also improving packet size, reducing bainite formation, and thus improving the steel's toughness.

[0057] Nb: Even a small amount of niobium can form dispersed carbides, nitrides, and carbonitrides to refine the grains, improving the strength and toughness of the steel. Simultaneously, even without the formation of carbonitrides, the absorptive effect of solute atoms at grain boundaries can refine the austenite grains and improve the steel's deformation capacity at high temperatures. During hardening heat treatment, it precipitates from the matrix as carbides without affecting the thermal conductivity of the matrix. In this invention, the Nb content is below 0.2%, preferably 0.01~0.1%.

[0058] Mn: Manganese is an austenite-forming element that delays the transformation of austenite to ferrite, thereby refining ferrite grains and improving the strength and toughness of steel. Mn dissolved in the matrix reduces the thermal conductivity of steel. When Mn replaces some Fe and Al atoms to form Ni(Fe,Mn,Al) intermetallic compounds, it reduces the thermal conductivity of the original Ni(Fe,Al). Furthermore, excessively high Mn content leads to several problems: firstly, residual austenite in the steel causes uneven microstructure and unstable steel properties; secondly, NiMn precipitates form at grain boundaries, severely impairing the mechanical properties of the steel. Therefore, this invention limits the Mn content to between 0.001% and 1%.

[0059] Cr: When Cr is dissolved in the matrix, it reduces the thermal conductivity of the matrix. Only when all Cr in the matrix precipitates as carbides can the damage to thermal conductivity be reduced, which is impossible to achieve under realistic conditions. Furthermore, when the alloy contains Cr, it dissolves in Mo and W carbides during the formation of these carbides, disrupting the phonon order of the carbides and thus reducing their thermal conductivity. In this invention, Mo and W carbides are used instead of Cr carbides. Therefore, this invention does not require the inclusion of Cr. However, since it is impossible to completely eliminate Cr during smelting, Cr is considered an unavoidable impurity element in this invention, and its content is required to be ≤0.4%.

[0060] Si: The use of scrap steel is a common practice in the smelting of special steels. While scrap steel, as a raw material, offers advantages in cost-effectiveness and resource recycling, its quality is not always stable, especially given the presence of elements such as silicon (Si), which is an unavoidable impurity during the smelting process. Typically, a Si content of 0.4% is considered acceptable in hot work die steels; therefore, the Si content in the steel of this invention does not exceed 0.4%, and more preferably does not exceed 0.2 wt.%.

[0061] Impurity elements such as P, S, and N: Under normal circumstances, phosphorus is a harmful element in steel, increasing its cold brittleness, worsening weldability, reducing plasticity, and deteriorating cold bending performance. The steel in this invention requires P to be below 0.05%. Sulfur is also generally a harmful element, causing hot brittleness in steel and reducing its ductility and weldability. The steel in this invention requires S to be below 0.015%. Nitrogen is an interstitial solid solution element that can significantly improve the strength of steel. It is also an austenite stabilizing element, expanding the austenite region and lowering the Ac3 temperature. N readily combines with strong nitride-forming elements such as Al to form large nitrides, reducing the toughness of the steel. The steel in this invention requires N to be below 0.015%.

[0062] The following embodiments or experimental data are intended to illustrate the present invention by way of example. Those skilled in the art should understand that the present invention is not limited to these embodiments or experimental data.

[0063] Steel ingots were smelted according to the compositions of the example steels A1-A6 and comparative steels B1-B4 of the present invention shown in Table 1, and forged at 1200°C into 80×80×200mm ingots. 3 After the billet is homogenized at 1200 ℃ for 5 hours, it is air-cooled to room temperature. Then, under laboratory conditions, it is held at 1200 ℃ for 30 min and forged to 13 mm, and then air-cooled to room temperature.

[0064] Table 1. Composition (wt.%) of the exemplary steel and comparative steel of this invention. The billet is cut into samples of 11×11×55mm and φ12.7×2.2mm, and then heat-treated using the heat treatment method of the present invention. In a preferred embodiment, the heat treatment method includes the following steps: After holding at 950℃ in an Ar protective atmosphere furnace for 1 hour, the product is removed and oil-quenched; then, after holding at 500℃ for 12 hours, it is air-cooled to room temperature.

[0065] After the quenching step, the metallographic structure of the comparison steel B1 was first observed. The metallographic structure of the comparison steel B1 after oil quenching is as follows: Figure 1 As shown. By Figure 1 It is evident that at low C content (C=0.02%), martensitic structure cannot be obtained after quenching. Its microstructure is blocky transformed ferrite with irregular grain boundaries, low dislocation density, and poor toughness.

[0066] Measurement of packet intercept after quenching After the quenching step, oil-quenched samples of example steels A1-A6 and comparative steels B2-B4 were mechanically ground and surface-polished with argon ions before undergoing EBSD experiments. The EBSD experiment yielded the packet morphology and dimensions after the austenite underwent martensitic transformation. The packet intercepts of the A1-A6 example steels and the B2-B4 comparative steels were statistically analyzed using the intercept method.

[0067] The packet intercept for martensitic phase transformation is measured according to the intercept method in the standard GB / T 6394-2017, "Method for Determination of Average Grain Size of Metals". The specific steps are as follows: Use the EBSD data processing software AztecCrystal to process the data into an IPF (inverse pole figure) + grain boundary diagram. Define grain boundaries with an angle >50° as Packet grain boundaries. Save the processed EBSD data as an image. Use the image processing software nanomeasurer to measure the length, width, and height of the image according to the scale. Figure 2 In the example, six straight lines are drawn along the length, width, and diagonal of the image. Besides... Figure 2Besides the line drawing method shown, you can also draw lines on the image in any other direction, with a minimum of 5 lines. Figure 2 For the six lines in the example, the number of intercept points where each line intersects with Packet boundaries of different orientations is counted. The Packet intercept is equal to the length of the line divided by the number of intercept points. The average and standard deviation of the Packet intercepts obtained by counting the number of Packet intercepts of the six lines are the Packet mean intercept. Figure 2 The data for packet intercept statistics are shown in Table 2.

[0068] Table 2 Packet intercept statistics for example steel A3 The statistical results are shown in Table 3 and Figure 3 As shown. From Figure 3 It can be observed that as the carbon content increases, the size of the martensite packet gradually decreases; the packet size at a carbon content of 0.35% is half the size of the packet at a carbon content of 0.05%. However, when the carbon content increases from 0.35% to 0.40%, the packet size no longer changes with the carbon content and remains essentially unchanged.

[0069] Increasing the carbon content beyond 0.35% not only fails to change the packet size of the martensitic transformation, but the excess carbon also causes twinned martensite transformation during the austenite-to-martensite transformation, in addition to dislocation martensite (lath martensite). Twinned martensite significantly reduces the toughness of the material.

[0070] Table 3. Martensitic packet intercept dimensions in example and comparative steels Measurement of dislocation density after quenching and tempering The dislocation density of steel was measured using an X-ray diffractometer. During the experiment, the test angle ranged from 35° to 130°, with a step size of 1.5° / min. After obtaining the diffraction data, noise reduction was performed using Jade9 software. The diffraction peaks (110), (200), (211), and (220) were then selected to fit the microstrain. e Fitting micro-strain e Then, its dislocation density can be calculated.

[0071] Microstrain of example steels A1-A6 and B2-B4 was measured using X-ray diffraction (XRD). e Then, Formula 1 is used to calculate the dislocation density. (Formula 1) in ρ Dislocation density, in meters (m)-2 , b This is the Burgers vector. For martensitic steel, b It is 0.234nm.

[0072] The dislocation densities of example steels A1-A6 and B2-B4, calculated using Formula 1, are shown in Table 4. Figure 4 As shown, when the C content is between 0.05% and 0.35%, the dislocation density gradually increases with increasing C content. When the C content reaches 0.38% and 0.40%, the dislocation density no longer increases and remains essentially the same as the dislocation density at a C content of 0.35%. Therefore, it can be seen that once the C content exceeds 0.35%, further increasing the C content will neither refine the packet intercept nor increase the dislocation density.

[0073] Table 4 Dislocation density of example steel and comparison steel after quenching In the tempering process, the example steels A1-A6 and the comparative steels B2-B4 were held at 500°C for 12 hours and then air-cooled to room temperature.

[0074] After the tempering step, their dislocation densities after tempering were analyzed using XRD, as shown in Table 5 and Figure 5 As shown in Table 5, the control steel B2 with a C content of 0.05% exhibits the lowest dislocation density after tempering. The example steels A1-A6 and control steels B2-B4 with C contents of 0.1-0.35% have approximately the same dislocation density. This indicates that when the C content exceeds 0.1%, the decrease in dislocation density after tempering is relatively small, thus ensuring that most of the C remains concentrated on dislocations, reducing its impact on thermal conductivity.

[0075] Table 5 Dislocation density of example and comparison steels after tempering After the tempering step, the cementite morphology in example steels A1, A3, and A5 is as follows: Figure 6a , Figure 6b and Figure 6cAs shown. The size distribution of cementite particles in the samples was statistically analyzed using a Nano Measurer 1.2: Twenty representative fields of view (approximately 1 μm × 1 μm, with clear and non-overlapping cementite morphology) were selected from the TEM bright-field image. After scale calibration, the diameter of the cementite particles was manually measured. At least 200 cementite particles were counted for each process, and parameters such as average diameter, number density, and size distribution were obtained. Since cementite is not always perfectly round, some particles are rod-shaped. Considering the maximum influence of cementite, the length direction is used as the maximum diameter. The average diameter of cementite in example steel A1 is 88.6 nm, in example steel A3 it is 106.7 nm, and in example steel A5 it is 105.7 nm. Nanoscale cementite can provide a certain strengthening effect without reducing the toughness of the steel of this invention. It can also reduce the C content in the matrix and improve thermal conductivity.

[0076] Measurement of hardness and toughness According to GB / T 229-2020 "Charpy Pendulum Impact Test Method for Metallic Materials", the heat-treated impact specimen blank of 11×11×55mm was machined into a standard impact specimen of 10×10×55mm. The U-shaped notch was machined using a curve grinding process.

[0077] In addition, after the quenched and tempered samples were polished to a bright finish with sandpaper, hardness tests were performed using a hardness tester (Rockwell hardness) to obtain hardness values.

[0078] The hardness and impact energy of the example steels A1-A6 and the comparative steels B1-B4 after being held at 500℃ for 12 hours and then air-cooled to room temperature are shown in Table 6.

[0079] Table 6. Hardness (HRC) and impact energy (J) of example steels A1-A6 and comparative steels B1-B4 As shown in Table 6, for example steels A1 to A6, the hardness increases with increasing C content, reflecting the strengthening effect of C. It is noteworthy that the impact energy of the U-notch specimen of comparative steel B1 (0.02C) is only 2.5 ± 1.3 J. This is because at low C content (C = 0.02%), martensite cannot be obtained after quenching. Its microstructure is blocky ferrite with irregular grain boundaries, low dislocation density, and poor toughness. For the comparative steel B2 (0.05C), although its toughness is significantly improved compared to the comparative steel B1 (0.02C) due to the formation of martensite, its low dislocation density results in poor toughness and low hardness. Furthermore, the toughness of the comparative steel B3 (0.38C) drops below 12J, failing to meet the performance improvement requirements. Therefore, as shown in Table 6, a mold steel meeting both hardness and toughness performance requirements can be obtained when the C content is greater than 0.1 wt.% and not higher than 0.35 wt.%. Table 6 also shows that at a C content... As the carbon content increases from 0.30 wt.% to 0.35 wt.%, the impact energy of the U-notch specimen decreases rapidly from 20.1 ± 2.8 J to 12.2 ± 2.7 J. Furthermore, as the carbon content increases from 0.11 wt.% to 0.15 wt.%, the hardness increases rapidly from 46.3 HRC to 48.5 HRC. Therefore, it is preferable to have a carbon content between 0.15 wt.% and 0.30 wt.%, within which a die steel with a hardness higher than 48.5 HRC and an impact energy of over 20.1 ± 2.8 J for the U-notch specimen can be obtained.

[0080] Measurement of thermal conductivity After heat treatment, the blank sample with dimensions φ12.7×2.2mm was machined into a standard sample with dimensions φ12×2mm. The thermal diffusivity was measured according to GB / T 22588-2008, "Measurement of Thermal Diffusivity or Thermal Conductivity by Flash Method". The density and specific heat of the example steel and comparison steel within the range of 200~420℃ were calculated using Thermo-Calc. Thermal diffusivity × specific heat × density equals thermal conductivity.

[0081] Thermal conductivity measurement: thermal conductivity λ = α × c p × ρ ×100, thermal diffusivity α The unit is cm 2 / s, specific heat capacity c p The unit for density is J / (gK), and the unit for stoichiometry is g / (cm³). 3The unit for direct calculation is W / (cmK)×100, resulting in W / (mK). A heat-treated cylindrical sample of φ12.7×2.2 mm was ground into a φ12.7×2.0 mm sample using 1000-grit sandpaper, and then its thermal conductivity was measured using a DLF2800 flash thermal conductivity meter. The measurement process was as follows: the temperature was increased from 25℃ to 100℃ at a rate of 5 K / min, stabilized at 100℃ for an estimated 10 minutes, and then the first measurement was performed; then it was stabilized for another 10 minutes, and the second measurement was performed; then it was stabilized for another 10 minutes, and the third measurement was performed. Afterwards, the temperature was increased to 200℃ at a rate of 5 K / min, and the same process was repeated three times. Then, the temperature was increased sequentially to 300℃, 400℃, and 500℃, and the same process was repeated three times (equivalent to holding at each measurement temperature for 30 minutes). After the measurements were completed, the sample was cooled to room temperature. The thermal diffusivity and specific heat capacity data at each temperature were obtained based on the measurements. Then the thermal conductivity is calculated from the thermal diffusivity, specific heat capacity, and density.

[0082] Table 7 shows the highest thermal conductivity of example steels A1-A6 and comparison steels B1-B4 in the temperature range of 200~420℃ after being held at 500℃ for 12 hours and then air-cooled to room temperature.

[0083] Table 7. Maximum thermal conductivity (W / m·K) of example steels A1-A6 and comparative steels B1-B4 in the range of 200~420℃. As shown in Table 7, the comparative steel B1 exhibits the highest thermal conductivity, but the lowest hardness and toughness, due to its phase transformation into blocky ferrite rather than martensite after quenching, resulting in a low ferrite dislocation density and near-complete precipitation of C into cementite after tempering. The comparative steel B2 (0.05C) has a thermal conductivity of 42.23 W / m·K, essentially the same as the example steel A1 (0.11C) at 42.21 W / m·K. This phenomenon contradicts existing knowledge. It is generally believed that the C content in the matrix has a negative effect on thermal conductivity; the lower the C content, the higher the thermal conductivity. This invention has found that the thermal conductivity remains unchanged as the C content increases from 0.05 wt.% to 0.11 wt.%. This is because, in both the comparative steel B2 (0.05C) and A1 steel (0.11C), the dislocation density after quenching and tempering allows most of the carbon (C) to accumulate in dislocations, forming Cocteau gas clusters (see Tables 4 and 5 above). Very little C dissolves in the octahedral interstitial sites of martensite, thus ensuring high thermal conductivity. When the C content exceeds 0.1 wt.% (see example steels A1 to A6 in Table 7), although the thermal conductivity decreases with increasing C content, the dislocation density also increases with increasing C content (see Table 5), resulting in a higher amount of C that can accumulate on dislocations. Therefore, the decrease in thermal conductivity is slower. Example steel A6 and comparative steel B3 have similar dislocation densities, so the amount of C accumulated on dislocations is similar, but comparative steel B3 has 0.03% more C than example steel A6. This excess C can only diffuse into the octahedral interstitial sites of martensite, and it also promotes the formation of twinned martensite, leading to a decrease in the thermal conductivity and toughness of B3 steel. The above reflects the relationship between dislocation density and thermal conductivity in mold steel. A high dislocation density allows more carbon-oriented dislocations to accumulate, forming Cobb-shaped atmospheres and reducing the negative impact of increased carbon content on thermal conductivity. However, due to increased defects (increased dislocation density), thermal conductivity also tends to decrease, albeit slowly.

[0084] Analysis of heat treatment process It is understood that the quenching and tempering steps are not limited to the specific temperatures and durations mentioned above, and the quenching method is not limited to oil quenching, but can be selected from the following ranges: The hot work die steel is held at 900~1000℃ for 0.1 to 72 hours, then quenched to room temperature; then held at 420~520℃ for 0.1 to 144 hours, and then cooled to room temperature. The quenching medium can be water, brine, oil, or polymer solution.

[0085] The austenitizing holding temperature before quenching of mold steel must ensure complete austenitization (minimum temperature) while preventing overheating / burning that could lead to coarsening of the original grains and carbides. The heating time is also related to the size of the steel. Larger and thicker steel pieces require longer holding times.

[0086] The tempering temperature and time depend on the final usage condition of the mold. For NiAl+Cu precipitation-strengthened aging steel, precipitation is slower than that of carbides, so there is no need for multiple short tempering sessions as with traditional mold steel. Continuous tempering can be used, which involves holding the steel at a specific temperature for several hours and then air-cooling it.

[0087] The following uses example steel A1 to illustrate the effect of quenching temperature-time and tempering temperature-time on the steel of the present invention (950°C has been shown in the aforementioned preferred embodiments).

[0088] Table 8 shows the heat treatment process for example steel A1. Table 9 shows the martensitic packet size and dislocation density of example steel A1 after different heat treatments, quenching, and tempering. Table 10 Hardness, toughness, and thermal conductivity of A1 steel after tempering As shown in Tables 9 and 10 above, the selection of quenching and tempering process parameters has little impact on the hardness, toughness, and thermal conductivity of the material. Mold steel with satisfactory performance can be obtained within the range of quenching and tempering process parameters specified in this invention.

[0089]

Claims

1. A hot work die steel having a chemical composition in weight percent as follows: 0.10% < C < 0.35%, 5.5% < Ni < 8.0%, 0.5% < Cu < 1.5%, 0.8% < Al < 1.4%, Mo < 0.2%, W < 0.2%, V < 0.2%, Nb < 0.2%, 0.001% < Mn < 1%, Cr < 0.4%, Si < 0.4%, P < 0.05%, S < 0.015%, N < 0.015%, and the balance of Fe and unavoidable impurities.

2. The hot work die steel of claim 1, wherein, 0.15%≤C≤0.30%。 3. The hot work die steel of claim 1, wherein, Dislocation density after tempering ≥ 4.261 14 / m 2 .

4. The hot work die steel of claim 1, wherein, Martensite packet size < 4.79 μm.

5. The hot work die steel of claim 1, wherein, Hardness 46.3 ~ 52.3 HRC.

6. The hot work die steel of claim 1, wherein, Hardness 48.5 ~ 51.4 HRC.

7. The hot work die steel of claim 1, wherein, Thermal conductivity > 36.24 W / m-K in the temperature range of 200 ~ 420°C.

8. The hot work die steel of claim 1, wherein, Thermal conductivity > 38.23 W / m-K in the temperature range of 200 ~ 420°C.

9. The hot work die steel of claim 1, wherein, Impact energy of U-notch specimen > 12.2 J.

10. The hot work die steel of claim 1, wherein, Impact energy of U-notch specimen > 20.1 J.

11. A heat treatment method for the hot work die steel according to any one of claims 1-10, comprising: Quenching step, the hot work die steel is held at 900 ~ 1000°C for 0.1 to 72 hours, and then quenched to room temperature.

12. The heat treatment method of claim 11, further comprising a tempering step after the quenching step, held at 420 ~ 520°C for 0.1 to 144 hours, and then cooled to room temperature.

13. The heat treatment method according to claim 11, wherein The properties of the heat treated hot work die steel are as follows: hardness 46.3 ~ 52.3 HRC, thermal conductivity > 36.24 W / m-K in the temperature range of 200 ~ 420°C, and impact energy of U-notch specimen > 12.2 J.

14. A hot work die made of the hot work die steel obtained by the heat treatment method according to any one of claims 11 to 13, characterized by, The hot work die can be used as a steel sheet hot stamping forming die, an aluminum alloy die casting die, a hot swaging die, or a hot extrusion drawing hot work die.

15. The hot work die of claim 14 wherein, The hot work die can be used as an aluminum alloy differential pressure or low pressure casting die. The hot work die can be used as an aluminum alloy differential pressure or low pressure casting die.

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