HK1192289A - Hot-work tool steel and a process for making a hot-work tool steel - Google Patents

Description

Hot-work tool steel and method for producing a hot-work tool steel

Technical Field

The present invention relates to a low chromium hot work tool steel and a method of manufacturing a low chromium hot work tool steel article.

Background

The term "hot working tool" applies to many different kinds of tools for metal working or forming at relatively high temperatures, for example tools for die casting, such as dies, inserts and cores (inserts and cores), inlet parts, nozzles, ejector elements (pistons), high pressure chambers, etc.; tools for extrusion processing, such as dies, die holders, bushings (lines), pressure pads and rods (pressing pads and stems), mandrels (spindles), and the like; tools for hot pressing, such as tools for hot pressing of aluminum, magnesium, copper alloys, and steel; molds for plastics, such as molds for injection molding, press molding and extrusion molding; as well as various other kinds of tools, such as tools for hot shearing, shrink-rings/collars and wear parts intended for use in high temperature machining. Low alloy hot work tool steels are used for small and medium sized tools in applications where the requirements for temper resistance and thermal fatigue are high. Temper resistance refers to the ability of a hot work tool steel to retain its hardness at high temperatures for long periods of time. Hot work tool steels have been developed for strength and hardness under prolonged exposure to high temperatures, and typically use alloys formed from large amounts of carbides.

Another type of tool steel is high speed steel, which is used in cutting tools where strength and hardness must be maintained at temperatures up to or exceeding 760 ℃. In order to reduce the amount of tungsten and chromium required, for example 18% and 4% by weight respectively, variants have been developed which use 5-10% by weight of molybdenum. High speed steel differs from hot-worked steel in composition and price, and high speed steel cannot be used as a substitute for hot-worked steel.

Disclosure of Invention

It is an object of the present invention to provide a low chromium hot-work tool steel with an improved performance profile, in particular an improved temper resistance. The steel of the invention is particularly suitable for small tools which do not require a steel composition with high hardenability for its manufacture.

This object is achieved by providing a low-chromium hot-work tool steel as defined in claim 1, i.e. a steel consisting of (in weight%):

optionally, optionally

Ni ＜3

Co ≤5

B ＜0.01

The balance of Fe except impurities.

Further objects are achieved by the low chromium hot work tool steel of the invention, which fulfils one or more of the following conditions (in weight%):

preferred embodiments of the low chromium hot work tool steel may satisfy one or more of the following conditions (in weight%):

more preferred embodiments of the low chromium hot work tool steel may satisfy one or more of the following conditions (in weight%):

n is 0.042 to 0.15, preferably 0.045 to 0.12

C+N 0.39～0.41

Cr 1.3 to 2.3 are preferably 1.4 to 2.1

Even more preferred embodiments of the low chromium hot work tool steel may fulfill one or more of the following conditions (in weight%):

according to the inventive concept, the low chromium hot work tool steel may have a composition (in weight%) according to the following examples:

optionally, optionally

The balance being Fe except for impurities, or

Optionally, optionally

The balance being Fe except for impurities, or

Optionally, optionally

B 0.001～0.01

Mo/V 1.8～2.3

Cr/V ＜2

The balance being Fe except for impurities, or

Optionally, optionally

B 0.001～0.01

Mo/V 1.8～2.3

Cr/V ＜2

The balance being Fe except for impurities, or

Optionally, optionally

B 0.001～0.005

Mo/V 1.8～2.3

Cr/V ＜2

The balance of Fe except impurities.

It is a further object of the present invention to provide an article of low chromium hot work tool steel with an improved property profile, in particular an improved temper resistance.

According to the invention, the above object is achieved by a method as defined in claim 11, i.e. a method comprising the steps of:

a) providing a low chromium hot work tool steel as defined in any one of the claims;

b) forming a steel article from the steel composition;

c) austenitizing the steel product obtained in step b) at a temperature of up to 1200 ℃ for about half an hour, followed by quenching; and is

d) Tempering the quenched steel product at a temperature of 500-700 ℃ at least twice, each for about 2 hours.

Preferred embodiments of the method are set forth in the dependent claims 12 to 15.

In creep-resistant steels with high chromium content, i.e. 9-12 wt.%, vanadium carbonitride (vanadia carbide) may already be dissolved at relatively low temperatures, i.e. 1020-1050 ℃. However, if the chromium content is low (less than about 4-5 wt%), primary (primary) vanadium carbonitrides will form in the melt and in practice they will not be possible to dissolve later.

In the steel of the present invention, the total amount of carbon and nitrogen should be adjusted to 0.30. ltoreq. C + N.ltoreq.0.50, preferably 0.36. ltoreq. C + N.ltoreq.0.44. The nominal content should be about 0.40% by weight. Meanwhile, it is advantageous to adjust the nitrogen content to 0.015 to 0.30N, preferably 0.015 to 0.15N, and even more preferably 0.015 to 0.10N, and carbon may be preferably adjusted to at least 0.20% by weight. Preferred ranges are listed in the product claims.

When the nitrogen content is balanced to about 0.05 to 0.10 wt.%, vanadium carbonitride is produced, which is partially dissolved during the austenitizing step and then precipitated in the form of nano-sized particles during the tempering step. Vanadium carbonitride has better thermal stability than vanadium carbide and, as a result, the temper resistance of low chromium hot work tool steel articles is significantly improved. Furthermore, by tempering at least twice, the tempering curve (showing hardness as a function of tempering temperature) will have a higher secondary peak.

In the most preferred embodiment of the invention, the nitrogen content is preferably about 0.05% by weight. The above values yield better performance than larger values. During quenching, a nitrogen content of about 0.05 wt.% has a higher secondary hardening potential than a higher nitrogen content, thus giving the steel a higher hardness. However, a nitrogen content of about 0.10 wt.% has shown that its secondary hardening peak shifts to an advantageously slightly higher tempering temperature. Preferred ranges are listed in the product claims. In addition, tests and model calculations performed show that increasing the nitrogen content requires increasing the austenitizing temperature.

Chromium enhances the hardenability and corrosion resistance of the steel. Too low a content may adversely affect the corrosion resistance. Therefore, the minimum value of the chromium content in the steel is set to 1 wt%. The maximum content is set to 4 wt.% in order to avoid undesirable formation of chromium-rich carbides/carbonitrides, such as M₂₃C₆. The chromium content preferably does not exceed 3 wt% and even more preferably does not exceed 2.6 wt%. In one embodiment of the present invention, the chromium content is 1.5 to 1.7% by weight. Preferred ranges are listed in the product claims. The low chromium content delays the precipitation of chromium carbides in the microstructure, which is favorable for the more heat-stable chromium-rich carbonitrides. Thus, the recovery of the material is slowed down and the tempering resistance is improved.

In order to provide sufficient precipitation potential and thereby sufficient temper resistance and desired high temperature strength properties, the steel should contain vanadium in an amount of at least 0.8 wt.%. In order to avoid the formation of excessive M (C, N) precipitates, which increase the risk of leaving large insoluble precipitates in the matrix after heat treatment and also the risk of losing carbon and nitrogen in the matrix, the upper limit of vanadium is 1.3 wt.%. The content of vanadium is preferably 1.0 to 1.3 wt%. Preferred ranges are listed in the product claims.

In order to obtain the desired MC phase, the Cr/V ratio should preferably be less than 2, more preferably less than 1.8. The reason is that chromium can be considered harmful to the MC phase.

The silicon content of the steel should be within 0.1 to 0.5 wt.%, preferably 0.2 to 0.4 wt.%. By keeping the silicon content low, it is possible to obtain metastable M₃Initial precipitation of C carbides. These carbides will act as reservoirs of carbon for subsequent precipitation of the desired M (C, N) particles. Also, undesirable chromium-rich M at grain and lattice interfaces is avoided₂₃C₆And (4) separating out particles. Preferred ranges are listed in the product claims.

Manganese is present in order to give the steel sufficient hardenability, especially in the case of relatively low contents of chromium and molybdenum in the steel. The manganese content in the steel is between 0.5 and 2 wt.%, preferably 1.0 to 2.0 wt.%. Preferred ranges are listed in the product claims.

In order to provide secondary hardening during tempering and to increase hardenability, the molybdenum should be present in the steel in an amount of between 1.5 and 3 wt.%, preferably 2.2 to 2.8 wt.%. Preferred ranges are listed in the product claims.

Part of the molybdenum may replace tungsten in a manner known per se, but preferably the steel should not contain any intentionally added amount of tungsten, i.e. should not contain tungsten in amounts exceeding impurity levels, because the presence of the element tungsten involves certain disadvantages.

In order to obtain the desired precipitation order and precipitation potential of the secondary carbides, the Mo/V ratio should preferably be in the range of 1.8 to 2.3, more preferably 1.9 to 2.1. Mo is known to stabilize M₂C phase, and Mo-rich M by adjusting Mo and V contents to fall within a range of 1.8 to 2.3₂C also forms, which has a higher coarsening rate than the vanadium rich MC phase.

Nickel and cobalt are elements that may be included in the steel in amounts of up to 3% and 5% by weight, respectively. Cobalt may increase hardness at high temperatures, which may be beneficial for some applications of the steel. If cobalt is added, the effective amount is about 4 wt%. Nickel may increase the corrosion resistance, hardenability and toughness of the steel. Preferred ranges are listed in the product claims.

In principle, austenitization may be carried out at a temperature between the softening annealing temperature 820 ℃ and the maximum austenitization temperature 1200 ℃, but austenitization of the steel product is preferably carried out at a temperature of about 1050 to 1150 ℃, preferably 1080 to 1150 ℃, typically 1100 ℃. Internal tests have shown that higher austenitizing temperatures shift the temper hardness to higher temperatures, i.e. the secondary hardening peak will shift to higher temperatures, which indicates that the desired hardness will be reached at higher initial tempering temperatures. Thus, the material will get an improved tempering resistance and the working temperature of the tool can be increased.

The quenched steel product is preferably tempered at least twice, with a retention time of 2 hours at a temperature of between 500 and 700 ℃, preferably 550 to 680 ℃. In a most preferred embodiment of the steel composition, the tempering is performed at a temperature within the range of 600-650 ℃, preferably 625-650 ℃.

The nitrogen content ranging from 0.05 to 0.10 wt% can be obtained by adding nitrogen in a conventional casting method to form a melt, casting the melt to form an ingot, and homogenizing the ingot by heat treatment. The addition of nitrogen produces large nascent vanadium-rich M (C, N) precipitates, which in turn can cause non-uniform hardness of the material. However, if the nitrogen content is reduced and there is a homogenization heat treatment prior to subsequent forging, such large primary carbonitrides do not occur.

In variants of the steel, higher nitrogen contents than in the preferred embodiment are also conceivable. In this variation, the nitrogen may amount to up to 0.30 wt%. In order to obtain higher nitrogen contents, conventional casting methods are not sufficient. Alternatively, nitrogen may be added by first making a steel powder of substantially the desired composition (except for nitrogen), then nitriding this powder in a solid state by a fluid containing nitrogen, such as nitrogen gas, and then isostatically pressing the powder at a temperature of about 1150 ℃ and a pressure of about 76Mpa to form an ingot. By manufacturing the tool steel by powder metallurgy, the problem of the formation of large primary carbides is avoided.

Preferably, the ingot is forged at a temperature of about 1270 ℃ and then soft annealed at a temperature of about 820 ℃, followed by cooling to a temperature of 650 ℃ at a rate of 10 ℃ per hour and then natural cooling in air to prepare it for austenitization.

The steel of the invention has a significantly improved temper resistance, which results in a longer product life in hot working applications. As already indicated above, the nitrogen content is preferably about 0.05% by weight and the chromium content is preferably less than 3% by weight, i.e. 1.2 to 2.6 or 1.3 to 2.3.

The steel product of the invention should preferably also meet some of the following requirements:

-a good resistance to tempering,

-a good high-temperature strength,

-a good thermal conductivity of the heat-generating material,

do not have an unacceptably large coefficient of thermal expansion.

Drawings

The invention will be described in more detail below with reference to preferred embodiments and the accompanying drawings.

FIG. 1 is a graph showing hardness versus tempering temperature for an exemplary nitrogen-free low chromium hot work tool steel of the prior art.

FIG. 2 is a graph showing the hardness of prior art steels in terms of Cr15, Mo1, C0.6 and Cr15, Mo1, C0.29, N0.35 (in weight%) at different tempering temperatures.

FIG. 3 is a schematic illustrating the effect of low chromium content on M (C, N) stability in austenite.

FIG. 4 shows M₆C. Plot of mole fraction of M (C, N) and bcc matrices as a function of temperature. (equilibrium phase: Austenitic matrix)

FIG. 5 is a graph showing M (C, N) phase and metastable state M₂Graph of the amount of C as a function of temperature. (equilibrium phase: ferrite)

FIG. 6 is a graph showing hardness versus tempering temperature curves for the test alloys N0.05, N0.10, and N0.30.

Fig. 7 is a back-scattered SEM picture showing small M (C, N) precipitates undissolved in N0.05 and spherical mixed oxide-sulfide particles.

FIG. 8 is a back-scattered SEM picture showing primary M (C, N) in alloy N0.10 that was not dissolved at the prior austenite grain boundaries.

Fig. 9 is a back-scattered SEM picture showing primary particles in N0.10 that were soft annealed.

Fig. 10 is a back-scattered SEM picture showing uniform distribution of undissolved M (C, N) in N0.30.

Fig. 11 is a back-scattered SEM picture showing some undissolved M (C, N) clusters found in N0.30.

Detailed Description

Medium alloyed (medium alloyed) hot work tool steels of molybdenum and vanadium have good resistance to thermal fatigue, softening and high temperature creep. An exemplary nominal chemical composition of such prior art steels is listed in table 1.

TABLE 1

C	Cr	Mo	V	Mn	Si	Fe
							0.38	2.6	2.3	0.9	0.75	0.3	92.8

It has been suggested that the high temperature properties of the steels in table 1 are due to precipitation of nano-sized vanadium carbide during tempering. These hard MC type carbides (2900HV) cause the material to secondary harden. Figure 1 shows the tempering curve (hardness versus tempering temperature) for an exemplary prior art tool steel. These samples were austenitized at 1030 ℃ and then tempered twice at different temperatures from 200 ℃ to 700 ℃ for 2+2 hours. It can be seen that in the range of 500-650 ℃, a significant secondary hardening peak value exists at 550 ℃. Later studies also showed that there was metastable molybdenum-rich M in the exemplary prior art tool steel during tempering at 625 deg.C₂C precipitates significantly, which promotes the secondary hardening effect.

The ability of a hot work tool steel to retain its hardness at high temperatures for long periods of time and its temper resistance, usually related to the initial tempering temperature; if the material is maintained at a temperature significantly below the initial tempering temperature, it will not soften. Softening at holding temperatures near or above the initial tempering temperature will be more pronounced.

If the secondary hardening peak can be shifted to higher temperatures, this would mean that the desired hardness (e.g., 44-46 HRC) can be achieved at higher initial tempering temperatures. Thus, the tempering resistance of the material may be improved and the working temperature of the tool may be increased.

Early studies on high chromium steels showed that when nitrogen was added to the steel, it was possible to achieve higher hardness during tempering. Samples Cr15, Mo1, C0.6 and Cr15, Mo1, C0.29, N0.35 were solution treated at 1050 ℃, followed by water quenching and cooling to liquid nitrogen, and then they were tempered at different temperatures for 2 hours. As can be seen in fig. 2, the peak hardness becomes significantly higher when nitrogen is added. As a nitrogen-containing steel, the initial hardness of martensite is low, but this steel reaches a higher hardness during tempering than a steel containing no nitrogen.

The explanation for this is that nitrogen makes the chromium more uniformly distributed in the matrix due to the increased solubility of chromium in the austenite phase. The martensite phase inherits the uniformly distributed chromium from the austenite after quenching and produces a very fine distribution of chromium nitride precipitates during tempering, thus producing a stronger hardening effect in the material.

In addition, nitrogen is used to replace some of the carbon to achieve higher hardness in the martensitic steel matrix. The addition of nitrogen initially results in a greater amount of retained austenite. However, the above mentioned austenite may subsequently be transformed into martensite by cold working, and in this way it is possible to achieve a hardness of up to 68 HRC.

The low chromium content appears to have a positive effect on the tempering resistance. Comparison of two different hot work tool steels with 1.5 wt.% and 5.0 wt.% chromium shows that the lower chromium content delays the precipitation of chromium carbides in the microstructure, contributing to the vanadium rich MC with better thermal stability. Thus, the recovery of the material is slowed down and the tempering resistance is improved.

However, studies on typical creep 9-12 wt.% chromium resistant steels (containing 0.06 wt.% N) show that the low chromium content significantly stabilizes MX (X is C + N) particles, see fig. 3. If austenitization is carried out at 1100 ℃, all M (C, N) particles will dissolve in the steel containing 10 wt.% chromium. If the chromium content is reduced to 2.5 wt.% (see the exemplary low chromium tool steel of fig. 1), a significant amount of M (C, N) will still be present in the austenite. Clearly, the result of the low chromium content is that only a small amount of interstitial species (interstitials) will dissolve into the austenite during the austenitizing treatment.

According to the present invention, a low chromium hot-work tool steel with enhanced temper resistance is manufactured by performing the following process steps:

a) adding nitrogen to the low chromium hot work tool steel melt composition as defined in any one of the method claims and thereby obtaining a steel composition;

b) forming a steel article from the steel composition;

c) austenitizing the steel product obtained in step b) at a temperature of up to 1200 ℃ for about half an hour, followed by quenching; and is

d) Tempering the quenched steel product at a temperature between 500 and 700 ℃ at least twice for about 2 hours each time.

These results are surprising in view of the general understanding in the art that a reduction in chromium content leads to a reduction in hardenability and difficulty in dissolving nascent M (C, N) particles.

In creep resistant steels with high chromium content (i.e. 9-12 wt.%), vanadium carbonitride may have dissolved at relatively low temperatures, i.e. 1020-1050 ℃. However, if the chromium content is low, less than about 4 to 5 wt.%, primary vanadium carbonitrides will form in the melt and they are virtually impossible to dissolve later.

The inventors have found that when the nitrogen content in low chromium steels is balanced to about 0.015-0.30 wt%, vanadium carbonitrides are generated, which partially dissolve in the austenitizing step and subsequently precipitate as nano-sized particles in the tempering step. These particles are about 1 μm to about 10 μm. In some cases, where the nitrogen content is low, typically at 0.05 wt%, the particles have an average size of less than 1 μm. Vanadium carbonitride has better thermal stability than vanadium carbide and therefore the temper resistance of low chromium hot work tool steel articles is significantly improved. Furthermore, the tempering curve (showing hardness as a function of tempering temperature) will have a higher secondary peak after at least two tempers.

In a preferred embodiment of the steel, the nitrogen content is preferably about 0.05 wt.%. The above values give better performance than larger values. A nitrogen content of about 0.05 wt.% has a greater potential for secondary hardening during quenching than larger contents.

In a preferred embodiment, the chromium content is preferably 1.5 to 1.7% by weight. The low chromium content delays the precipitation of chromium carbide in the microstructure, which is beneficial for the more thermally stable vanadium-rich carbonitride. Thus, the recovery of the material is slowed down and the tempering resistance is improved.

In principle, austenitization can be carried out at a temperature between the softening annealing temperature 820 ℃ and the maximum austenitization temperature 1200 ℃. In a preferred embodiment, i.e. a composition having a nitrogen content of about 0.05 wt.% and a chromium content of about 1.5 to 1.7 wt.%, the austenitization of the steel product is preferably carried out at a temperature of about 1050 to 1150 ℃, preferably 1100 ℃. Internal tests have shown that a higher austenitizing temperature shifts the temper hardness to a higher temperature, i.e. the secondary hardening peak will shift to a higher temperature, which means that the desired hardness will be reached at a higher initial tempering temperature. Thus, the material will have improved tempering resistance and the working temperature of the tool will be increased.

The quenched steel product is preferably tempered at a temperature of 500-700 ℃, preferably 550-680 ℃, for at least two times, the holding time being 2 hours. In a most preferred embodiment of the steel composition, the tempering is performed at a temperature of 600 to 650 ℃, preferably 625 to 650 ℃.

The nitrogen content ranging from 0.05 to 0.10 wt% can be obtained by introducing nitrogen by adding nitrogen using a conventional casting method to generate a melt, casting the melt to generate an ingot, and homogenizing the ingot by heat treatment. The addition of nitrogen produces large nascent vanadium-rich M (C, N) precipitates, which in turn make the material non-uniform in hardness. However, if the nitrogen content is reduced and the heat treatment for homogenization is performed before the subsequent forging, massive primary carbonitrides are not generated.

In a preferred embodiment of the invention, the nitrogen content is preferably about 0.05% by weight. The above values lead to better performance than larger values. A nitrogen content of about 0.05 wt.% gives the steel a higher potential for secondary hardening during quenching than a larger content, thus giving the steel a high hardness. However, it has been shown that an amount of about 0.10 wt.% shifts the secondary hardening peak to a slightly higher favorable tempering temperature. In addition, tests and modeling calculations performed indicate that increasing nitrogen content requires increasing austenitizing temperature.

In variants of the steel, higher nitrogen contents than indicated in the preferred embodiment are also conceivable. In the above variant, the nitrogen may amount to up to 0.30% by weight. To obtain higher nitrogen contents, conventional casting methods are not sufficient. Alternatively, the nitrogen is thus preferably added by first making a steel powder of essentially the desired composition except for nitrogen, followed by nitriding this solid powder by nitrogen gas, and then isostatically pressing the powder at a temperature of about 1150 ℃ and a pressure of about 76MPa to produce an ingot. The problem of the generation of primary carbides is avoided by manufacturing the tool steel by powder metallurgy.

The ingot is preferably forged at a temperature of about 1270 ℃ and then soft annealed at a temperature of about 820 ℃, followed by cooling to 650 ℃ at a rate of 10 ℃ per hour, and then naturally cooled in air to prepare it for austenitization.

Example 1

In table 2 below, the chemical composition of three different alloys, N0.05, N0.10 and N0.30, is listed in weight%. N0.05 refers to a material containing 0.05 wt% of nitrogen, and the like. Note that these are the actual compositions in the experimental ingots.

The aim is to keep the contents of all alloying elements, except carbon and nitrogen, constant. The chromium is also slightly reduced compared to the standard low chromium hot work tool steel in table 1. The molybdenum content decreased slightly while the manganese content increased. The aim is to keep the total amount of these elements fixed at around 0.40 wt.% for carbon and nitrogen, which is relatively successfully achieved.

TABLE 2

Material	C	N	Cr	Mo	V	Mn	Si	Fe
									N0.05	0.38	0.05	1.70	2.77	1.20	1.09	0.30	92.5
N0.10	0.27	0.10	1.53	2.32	1.20	1.85	0.26	92.5
									N0.30	0.08	0.32	1.51	2.20	1.20	1.88	0.29	92.5

The tempering phase mainly involves metastable phases, and previous electron microscopy studies have shown that these metastable phases are present in standard low-chromium hot-work tool steels in the tempering temperature interval, i.e. 400-700 ℃. These carbide phases are mainly vanadium-rich MC (FCC) and molybdenum-rich M₂C (HCP). A certain amount of chromium-rich M is also found in standard low-chromium hot-work tool steels₇C₃。

The following calculations were made to determine whether these alloys containing nitrogen are likely to harden if enough alloying elements are able to dissolve into the austenitic matrix at the austenitizing temperature to cause martensite to form during quenching. The temperature interval of interest is therefore between the softening annealing temperature (820 ℃) and the set maximum austenitizing temperature (1200 ℃) that is practically available.

The results of these balance calculations are shown in fig. 4. Shown here is M₆C. Mole fraction of M (C, N) and bcc matrices as a function of temperature. The equilibrium phase is austenite. The solid curve represents N0.05, the dashed curve represents N0.10, and the dotted curve represents N0.30. Note the high content of M (C, N) in the N0.30 alloy even up to 1200 ℃. As expected, the bcc phase is unstable above 850 ℃. It was interestingly found that the slope of the equilibrium curve representing the amount of M (C, N) decreases with increasing nitrogen content. This indicates that it is more difficult to dissolve M (C, N) in N0.30 than N0.05. Thus, it is expected that the contents of carbon, nitrogen and vanadium are smaller in the N0.30 matrix than in the N0.05 matrix after austenitizing at 1100 ℃.

Due to molybdenum-rich M₆Phase C only dissolves carbon but not nitrogen, this phase has a lower carbon content in N0.10 and N0.30, so M₆The amount of C decreases with decreasing carbon content. It should also be noted that M is at the austenitizing temperature used₆All C dissolved.

To estimate the potential for secondary precipitation in N0.05, N0.10 and N0.30, calculations were performed only for the tempering temperature region. The resulting equilibrium can be at most indicative of what phases will be present in the material after a sufficiently long time. Previous work has shown that in practice some auto-tempering will take place in standard low chromium hot work tool steels. This means that M₃C (cementite) precipitates after the austenitizing process.

The results calculated in the tempering temperature region are shown in fig. 5. The solid curve represents N0.05, the dashed curve represents N0.10, and the dotted curve represents N0.30. The secondary hardening is usually carried out within a temperature range of 500 to 650 ℃ and there is no great difference between N0.05 and N0.10 in terms of the amount of M (C, N) in the above temperature range. On the other hand, N0.30 has a high and almost constant M (C, N) content, which is probably due to the high vanadium and nitrogen content.

Higher carbon content in N0.05 produced more M than N0.10₂Phase C and matrixAnd (4) balancing. While there are much fewer M in N0.30₂C。

Based on the foregoing calculations, it should be possible to estimate the potential for secondary precipitation in these alloys after austenitization at a particular temperature. The potential depends on the M (C, N) phase and the M between the metastable equilibrium at the tempering temperature and the equilibrium at the austenitizing temperature₂Difference in C phase content. In table 3, these differences are expressed in terms of the secondary precipitation potential of three different alloys. These values are given in mole percent.

TABLE 3

The results in table 3 show that N0.05 has the best hardening response due to the low content of M (C, N) phase at 1100 ℃, i.e. a large amount of alloying elements can dissolve into the austenitic matrix. This also indicates that N0.05 has the best potential for good secondary hardening during tempering at 625 ℃.

Example 2

Both N0.05 and N0.10 alloys are conventionally cast into 50kg ingots. The first trial selected N0.10 and the ingot was not homogenized prior to the forging process. The second test, N0.05, was homogenized at 1300 ℃ for up to 15 hours before forging. The third alloy, N0.30, has a nitrogen content too high to be produced by conventional casting means. Such alloys are therefore manufactured using powder metallurgy. First a steel powder is produced, which powder is then nitrided in the solid state by means of pressurized nitrogen. The powder was then Hot Isostatic Pressed (HIP) at 1150 ℃ with a pressure of 76 MPa.

All three ingots were forged at 1270 ℃ and then the samples were cut to a size of 15x15x8 mm. These samples were heat treated at 820 ℃ by a first softening annealing, the sequence of post-annealing cooling being 10 ℃ to 650 ℃ per hour, and then naturally cooled in air. After the softening annealing, N0.05 was austenitized at 1100 ℃ for 30 minutes. To compensate for the poor precipitation potential, N0.10 was austenitized at 1150 ℃ for 30 minutes, and N0.30 was austenitized at 1200 ℃ for 30 minutes. Nine samples from each of these three alloys were tempered at the following temperatures: 450 ℃, 525 ℃, 550 ℃, 575 ℃, 600 ℃, 625 ℃, 650 ℃, 675 ℃ and 700 ℃. The soaking time was 2 hours, and the treatment was a double tempering treatment, i.e., the total tempering time was 4 hours. After heat treatment, the hardness of these samples was measured. Scanning Electron Microscopy (SEM) was used to further study the morphology, distribution and size of undissolved particles in the sample. The SEM instrument used was FEI Quanta 600F.

Hardness measurement

The results of the hardness measurements are shown in fig. 6. It can be seen that all three alloys have a secondary hardening peak in the temperature range of 500 to 600 ℃. All tempering treatments were carried out for 2+2 hours. N0.05 had the highest hardness (53HRC) in the quenched state, while N0.10 and N0.30 had somewhat lower hardness. However, all three alloys are considered hardenable. The hardness curve for N0.05 is very similar to that of a standard low chrome hot work tool steel, which as shown in fig. 1 has a maximum value of about 54 HRC.

The secondary hardening peak of N0.10 appears to shift somewhat to higher temperatures with a peak in hardness at 600 ℃. The hardness peaks for N0.05 and N0.30 were at 550 ℃.

Scanning electron microscope

In conventionally cast N0.05, the alloy with the lowest nitrogen content, the undissolved M (C, N) particles have an average size of less than 1 μ M. This is comparable to the normally undissolved carbides in steel. Another phase that is easily found in N0.05 is a mixture of aluminum oxide and manganese sulfide, see fig. 7, which is an SEM image (back scattering) showing small undissolved M (C, N) precipitates 2 and spherical oxide-sulfide mixed particles 1 in N0.05. The sample was austenitized at 1100 ℃ for 30 minutes and tempered at 625 ℃ for 2+2 hours.

The reason for the presence of many non-metallic inclusions in N0.05 (and N0.10) is that all test ingots were manufactured and cast in an open environment.

After 30 minutes of austenitizing treatment at 1150 ℃ and 2+2 hours of tempering treatment at 625 ℃, the most common size of M (C, N) particles in N0.10 is 5-10 μ M equivalent ring Diameter (ECD). Larger, primary carbides 3 (precipitated in the melt) are often found at the prior austenite grain boundaries, see FIG. 8, which is a back-scattered SEM image showing undissolved, primary M (C, N) at the prior austenite grain boundaries in alloy N0.10. The sample was austenitized at 1150 ℃ for 30 minutes and tempered at 625 ℃ for 2+2 hours.

Fig. 9 is a detailed SEM micrograph of nascent M (C, N) particles 4 in N0.10. They were automatically discovered in the SEM using INCA feature software from Oxford Instruments. Their clear edges indicate that they have precipitated from the melt. The white area in the figure is M rich in molybdenum₆And C, particles 5. Note that the sample in this case is soft annealed N0.10.

In powder metallurgically manufactured N0.30, the undissolved M (C, N) particles 6 have a size distribution (ECD) between 1 and 5 μ M, the most common being 2 μ M, and therefore the particles are small, despite the high nitrogen content. These particles are uniformly distributed in the microstructure, see fig. 10. However, as shown in fig. 11, some clusters 7 of M (C, N) were found.

The chemical composition of the undissolved M (C, N) phase particles in all three alloys was measured by EDS and the results are listed in table 4, which shows the chemical composition of the M (C, N) particles in alloys N0.05, N0.10 and N0.30. The above values are given in mole percent. Note that although the accuracy of EDS is not very high for light elements such as carbon and nitrogen, it can be seen that the balance of carbon and nitrogen in the M (C, N) phase is predictable based on the nominal composition. The values given in the table are the values given in the INCA program (Oxford Instruments). Part of the iron recorded may come from the surrounding matrix, especially for alloy N0.05.

TABLE 4

Industrial applicability

The method and low chromium hot work tool steel of the present invention are suitable where it is desired to have a hot work steel tool that can be used at elevated temperatures for extended periods of time.

Claims (15)

1. A low chromium hot work tool steel consisting of (in weight%):

optionally, optionally

Ni ＜3

Co ≤5

B ＜0.01

The balance of Fe except impurities.

2. A low chromium hot work tool steel according to claim 1, which fulfils one or more of the following conditions (in weight%):

3. a low chromium hot work tool steel according to claim 1 or 2, which fulfils one or more of the following conditions (in weight%):

4. a low chromium hot work tool steel according to any one of the preceding claims, which fulfils one or more of the following conditions (in weight%):

n is 0.042 to 0.15, preferably 0.045 to 0.12

C+N 0.39～0.41

Cr 1.3 to 2.3 are preferably 1.4 to 2.1.

5. A low chromium hot work tool steel according to any one of the preceding claims, which fulfils one or more of the following conditions (in weight%):

6. a low chromium hot work tool steel according to claim 1, having the composition (in weight%):

optionally, optionally

The balance of Fe except impurities.

7. A low chromium hot work tool steel according to claim 1, having the composition (in weight%):

optionally, optionally

The balance of Fe except impurities.

8. A low chromium hot work tool steel according to claim 1, having the composition (in weight%):

optionally, optionally

B 0.001～0.01

Mo/V 1.8～2.3

Cr/V ＜2

The balance of Fe except impurities.

9. A low chromium hot work tool steel according to claim 1, having the composition (in weight%):

optionally, optionally

B 0.001～0.01

Mo/V 1.8～2.3

Cr/V ＜2

The balance of Fe except impurities.

10. A low chromium hot work tool steel according to claim 1, having the composition (in weight%):

optionally, optionally

B 0.001～0.005

Mo/V 1.8～2.3

Cr/V ＜2

The balance of Fe except impurities.

11. A method of making a low chromium hot work tool steel article with enhanced temper resistance comprising:

a) providing a steel as defined in any one of the preceding claims;

b) forming a steel article from the steel;

c) austenitizing the steel product obtained in step b) at a temperature of up to 1200 ℃ for about half an hour, followed by quenching; and is

d) Tempering the quenched steel product at a temperature of 500-700 ℃ at least twice, each for 2 hours.

12. A method according to claim 11, comprising austenitizing the steel article at a temperature of 1050 to 1150 ℃, preferably 1080 to 1150 ℃.

13. A method according to any one of claims 11 or 12, comprising tempering the quenched steel article at a temperature of 550 to 680 ℃, preferably 600 to 650 ℃, even more preferably 625 to 650 ℃.

14. A method according to any one of claims 11 to 13 further comprising adding nitrogen by first producing a steel powder of substantially the desired composition other than nitrogen, subsequently nitriding the solid powder with nitrogen to provide the desired composition, and then hot pressing the powder to form an ingot.

15. A method according to any one of claims 11 to 14 further comprising the step of homogenising, forging and soft annealing the ingot prior to austenitising.