COLD WORK DIE STEEL, DIE, AND METHOD FOR PRODUCTION OF COLD WORK DIE STEEL

TECHNICAL FIELD

The present invention relates to a cold work die steel and a die, and to a method for producing a cold work die steel. Precisely, the invention relates to a die steel useful as a material for dies for use in cold/warm press-forming (punching, bending, drawing, trimming, or the like) of steel sheets for automobiles and steel sheets for household electric appliances, and to a method for producing such a die steel.

BACKGROUND ART

Dies for use in forming steel sheets for automobiles and steel sheets for household electric appliances have been desired to have a prolonged life with the increase in the strength of the steel sheets. In particular, for steel sheets for automobiles, there is increasing a great demand for high-tensile steel sheets having a tensile strength of at least about 590 MPa for automobile mileage improvement in consideration of environmental problems; and accordingly, there has occurred a problem in that the surface film of a die is damaged in early stages to cause “seizure” (phenomenon of soldering in press-forming) and the die life is thereby extremely shortened.

A die is composed of a die matrix (die steel) and a surface hardened layer (surface film) formed on its surface. The die steel for matrix is produced generally in a process including annealing, cutting, and quenching/tempering (in the present specification, particularly, quenching is referred to as solution treatment, and tempering is referred to as aging treatment) performed in this order.

As a die steel (cold work die steel), heretofore generally used are a high-C high-Cr alloy tool steel such as typically JIS SKD11, and a high-speed tool steel such as typically JIS SKH51, which has further improved abrasion resistance. In those tool steels, the hardness is increased mainly by precipitation hardening of a Cr-based carbide, or an Mo, W or V-based carbide. Furthermore, in order for enhancing both abrasion resistance and toughness, also used is a low-alloy high-speed tool steel (generally referred to as matrix high-speed steel) derived from JIS SKH51 by reducing the alloy content of C, Mo, W, V and the like therein.

For further improving the properties of cold work die steel, for example, Patent Reference 1 to Patent Reference 3 propose a technique of modifying steel-constitutive ingredients.

Patent Reference 1 is proposed for further increasing the hardness of matrix high-speed steel, and this reference describes a method for increasing the hardness (enhancing the abrasion resistance) of the steel by adding a large quantity of Nb and/or Tb thereto so as to prevent the crystal grains from growing coarsely in high-temperature quenching, in which the steel therefore accepts high-temperature quenching.

Patent Reference 2 relates to a cold work die steel having a dimensional change-resistant property and a high hardness property, mainly disclosing (a) the dimensional change due to expansion in tempering to be caused by the decomposition of residual austenite in quenching is cancelled by the dimensional change-resistant effect that results from the precipitation reinforcement of an Ni—Al intermetallic compound, and (b) the dimensional change is further suppressed by the segregation index K to be computed from predetermined ingredients of steel. FIG. 1 in Patent Reference 2 shows tempering at a temperature at which the steel may have a largest hardness.

Patent Reference 3 discloses a cold work die steel with a suitable quantity of Ni and Al added thereto and with a suitable quantity of Cu also added thereto in accordance with the afore-mentioned ingredients, for the purpose of reducing the degree of dimensional change (dimensional change) in quenching/tempering, especially for preventing the dimensional change by expansion in tempering as well as the purpose of increasing the hardness. This reference also describes controlling the content of C and Cr and finely dispersing the carbide in the microstructure, to thereby enhance the abrasion resistance of the steel.

On the other hand, Patent Reference 4 discloses a technique of a “pre-hardened steel” to be produced not by cutting and then quenching/tempering a steel as before, but by cutting a quenched/tempered steel (quenching/tempering, followed by cutting), for the purpose of reducing the die production cost. Concretely, as a steel having a high hardness but capable of exhibiting good machinability and capable of being worked by cold punching, the reference specifically discloses a pre-hardened steel in which the content of C, Si and S is suitably controlled. However, the life of the die formed of a pre-hardened steel is short, and at present, the die is not as yet in practical use.

The above-mentioned Patent References 1 to 4 are mainly for preventing the dimensional change of steel after heat treatment (after aging treatment or after tempering treatment) by controlling the steel-constituting ingredients; but Patent References 5 to 7 to be mentioned below disclose a technique of preventing dimensional change by controlling the condition in heat treatment such as quenching/tempering.

Of those, Patent Reference 5 discloses a method of preventing the dimensional change of steel after quenching/tempering by subjecting a steel to at least one pass of low-temperature tempering at 150 to 450° C. and at least one pass of high-temperature tempering at 480 to 550° C.

Patent Reference 6 discloses a method including quenching, sub-zero treatment at 0 to −200° C., and low-temperature tempering at 500° C. or lower performed in this order. In detail, a steel is subjected to a sub-zero treatment at the above-mentioned temperature to thereby control the residual austenite amount for controlling the dimensional change of the processed steel, and then this is subjected to a low-temperature tempering to realize the intended dimension of the steel.

Patent Reference 7 discloses a method of realizing a predetermined hardness of steel by controlling the steel ingredients to enhance the quenchability of steel followed by controlling the cooling rate in quenching by pearlite nose and gas cooling, thereby reducing the thermal treatment strain while maintaining the necessary hardness of steel for dies.

Patent Reference 1: JP-A-10-330894

Patent Reference 2: JP-A-2006-152356

Patent Reference 3: JP-A-2006-169624

Patent Reference 4: JP-A-2002-241894

Patent Reference 5: JP-A-9-125204

Patent Reference 6: JP-A-2001-172748

Patent Reference 7: JP-A-2002-167644

DISCLOSURE OF THE INVENTION
Problems that the Invention is to Solve

As the necessary properties thereof, a cold work die steel should have a high hardness and should be excellent in dimensional change resistance after heat treatment as mentioned above and, in addition, it should be excellent in welding repairability.

Welding repairing is mainly for revising and repairing damages (in detail, faults and depressions of a surface hardened layer) of a die to thereby regenerate and reuse the die; and for example, overlaying welding by argon welding or the like is generally employed therefor. As so mentioned in the above, due to the increase in the demand for high-tensile steel having a tensile strength of at least about 580 MPa, the life of dies has become extremely shortened; and for running cost reduction, welding repairing of dies is frequently carried out.

However, when a die coated with a hardened film is subjected to welding repairing, then the hardness around the welded part may much fluctuate to thereby easily cause cracking and seizure. In particular, the heat-affected zone (HAZ) after welding is remarkably softened (HAZ softening), thereby causing a problem in that the die life after welding repairing is shortened. HAZ softening is a phenomenon seen in a region spaced a little from the bonded part (boundary between a welding metal and a matrix, and this may be referred to as “weld melt line”); and in this region, it is considered that since the heating temperature is lower than that at the bonded part and transformation from grain-refined austenite is caused, the quenchability is lowered to increase a soft ferrite phase fraction and the side further remoter from the region may be tempered at a high temperature, whereby the hardness is lowered. FIG. 1(a) is a schematic view showing welding of matrices with a weld metal; and FIG. 1(b) graphically shows a hardness distribution in the region A in FIG. 1(a). As shown in FIG. 1(b), the HAZ hardness lowers with being spaced more from the bonded part, and the zone becomes softened. When HAZ becomes softened, then the surface hardened layer to be formed as a result of subsequent surface hardening could not sufficiently exhibit its protective effect, and the surface hardened layer may be damaged in early stages, whereby the die life becomes short.

As so mentioned in the above, welding repairing may be attained after a surface hardened film has been formed on the matrix, or may be attained before the formation of such a surface hardened film on the matrix. In particular, in press-forming a high-tensile steel die having a tensile strength of at least about 590 MPa, since it is difficult to press the steel to have a desired shape, in some cases, the steel is previously tested for press-forming and processed for welding repairing (overlaying welding), and thereafter it is actually press-formed into a desired form. In the test press-forming step, the steel is press-formed without heat treatment after welding repairing, and therefore, the HAZ softening part may often be damaged to have faults. The faults formed in such a HAZ softening part may remain in the surface film to be formed in the subsequent hardening treatment, and therefore, the remaining faults may be the starting points to cause film damages. In addition, not only the HAZ softening part but also a hardened part may be formed (see FIG. 1 and FIG. 7), and in the hardened part, breaking or cracking may often occur to cause troubles.

Accordingly, it is desired to provide a die steel excellent in welding repairability, which can prevent HAZ softening in welding repairing and which can readily realize overlaying welding at corners. However, all the above-mentioned patent references have no consideration in welding repairability, in which, therefore, the life of dies after welding repairing may be shortened.

The present invention has been made in consideration of the above-mentioned situation, and its object is to provide a cold work die steel having a high hardness, excellent in dimensional change resistance after heat treatment and having good welding repairability, and to provide a die.

Another object of the invention is to provide a method for efficiently producing a cold work die steel having a high hardness and excellent in dimensional change resistance after heat treatment.

Means for Solving the Problems

Specifically, the invention relates to the following items 1 to 12:

1. A cold work die steel comprising, by mass %,

C: 0.20 to 0.60%,

Si: 0.5 to 2.00%,

Mn: 0.1 to 2%,

Cr: 3.00 to 9.00%,

Al: 0.3 to 2.0%,

Cu: 1.00 to 5%,

Ni: 1.00 to 5%,

Mo: 0.5 to 3%, and/or W: 2% or less (including 0%),

S: 0.10% or less (not including 0%),

wherein the following requirements (1) to (3) are satisfied {wherein each square bracket [ ] means a content (%) of each element}:

[Cr]×[C]≦3.00, (1)

[Cu]/N11:0.5 to 2.2, (2)

[Mo]+0.5×[W]:0.5 to 3.0%, (3)

with the remainder being iron and unavoidable impurities.

2. The cold work die steel according to item 1, which further contains V: 0.5% or less (not including 0%).

3. The cold work die steel according to item 1 or 2, which further contains at least one element selected from the group consisting of Ti, Zr, Hf, Ta and Nb in a total amount of 0.5% or less (not including 0%).

4. The cold work die steel according to any one of items 1 to 3, which further contains Co: 10% or less (not including 0%).

5. The cold work die steel according to any one of items 1 to 4, which has a martensite transformation point (Ms point) represented by the following formula:

Ms point=550−361×[C]−39×[Mn]−35×[V]−20×[Cr]

−17×[Ni]−10×[Cu]−5×([Mo]+[W])+

15×[Co]+30×[Al]

{wherein each square bracket [ ] means a content (%) of each element}, of 170° C. or higher.

6. A die obtained by using the cold work die steel according to any one of items 1 to 5.

7. A method for producing a cold work die steel, comprising steps of:

preparing a steel satisfying the composition according to claim 1 and further satisfying the following requirement (4) {wherein each square bracket [ ] means a content (%) of each element}:

[Cu]/[C]: 4.0 to 15; (4)

and subjecting the steel to a solution treatment and an aging treatment under the condition satisfying the following formula (5):

TA−10≦T2≦TA+10 (5)

wherein,

TA=0.29×T1−2.63×[Cu]/[C]+225,

T1 means a solution treatment temperature (° C.), and

T2 means an aging temperature (° C.).

8. The production method according to item 7, wherein the steel further contains V: 0.5% or less (not including 0%).

9. The production method according to item 7 or 8, wherein the steel further contains at least one element selected from the group consisting of Ti, Zr, Hf, Ta and Nb in a total amount of 0.5% or less (not including 0%).

10. The production method according to any one of items 7 to 9, wherein the steel further contains Co: 10% or less (not including 0%).

11. The production method according to any one of items 7 to 10, wherein the steel has a martensite transformation point (Ms point) represented by the following formula:

Ms point=

550−361×[C]−39×[Mn]−35×[V]−20×[Cr]

−17×[Ni]−10×[Cu]-5×([Mo]+[W])+

15×[Co]+30×[Al]

{wherein each square bracket [ ] means a content (%) of each element}, of 170° C. or higher.

12. A die obtained in accordance with the production method according to any one of items 7 to 11.

ADVANTAGE OF THE INVENTION

According to the cold work die steel of the invention, since the alloy ingredients are suitably controlled as in the above, the steel has a high hardness, is excellent in dimensional change resistance after heat treatment and has good welding repairability. Accordingly, the die obtained by the use of the above-mentioned cold work die steel is favorably used especially as a molding die for high-tensile steel sheets having a tensile strength of at least about 590 MPa, and the life of the die, especially the life after welding repairing thereof can be further prolonged.

In addition, in the production method of the invention, since the steel-constituting ingredients and also the condition for solution treatment and for aging treatment are suitably controlled, a cold work die steel having a high hardness and excellent in dimensional change resistance after heat treatment can be produced efficiently. Accordingly, the die obtained according to the production method of the invention is favorably used especially as a molding die for high-tensile steel sheets having a tensile strength of at least about 590 MPa, and the life of the die, especially the life after welding repairing thereof can be further prolonged.

BRIEF DESCRIPTION OF THE DRAWINGS

FIGS. 1(
a) and 1(b) are views schematically showing a condition of welding of matrices with a weld metal, in which FIG. 1(a) is a cross-sectional view of a weld part, and FIG. 1(b) is a view schematically showing a hardness distribution in the region A in

FIG. 1(
a).

FIG. 2(
a) is an optical microscopic photograph showing a condition of seizure that has occurred on the surface of a die formed of a die steel, JIS SKD11 and coated with a TiN film; FIG. 2(b) and FIG. 2(c) each are a partly-enlarged optical microscopic photograph thereof; and FIG. 2(d) is an optical microscopic photograph of a die matrix before coated with a TiN film.

FIG. 3(
a) is an outline view showing the shape of a weld test piece used in Examples; FIG. 3(b) is an enlarged cross-sectional view of a groove.

FIG. 4 is an outline view schematically showing a condition of a test piece processed by buttering.

FIG. 5 is an outline view showing the shape of a Charpy impact test piece used in Examples.

FIG. 6 is a graph showing the relationship between a ratio [Cu]/[Ni] and a HAZ softening width.

FIG. 7 is a graph showing a hardness distribution profile.

FIG. 8 is a graph showing the relationship between a ratio [Cu]/[C] and a dimensional change rate (mean value, maximum value).

FIG. 9 is a view schematically showing the influence of aging treatment on a hardness and a dimensional change (dimensional change rate).

FIG. 10 is a view schematically showing the influence of aging treatment on a dimensional change amount.

BEST MODE FOR CARRYING OUT THE INVENTION

The invention is described in detail hereinunder. In the present specification, the percentage is by mass, unless otherwise specifically indicated. All the percentage expression by mass is the same as the percentage expression by weight.

A cold work die steel of the first aspect of the invention is described in detail.

For the purpose of providing a cold work die steel having improved especially in terms of the hardness, the dimensional change resistance after heat treatment and the welding repairability among various properties necessary for cold work die steel, the present inventors have first studied conventional dies formed of JIS SKD11 or matrix high-speed steel to clarify the reason why the surface film of the die is damaged and seizure is generated.

FIG. 2(
a) is an optical microscopic photograph showing a condition of seizure that has occurred on the surface of a die formed of a die steel, JIS SKD11 and coated with a TiN film; and FIG. 2(b) and FIG. 2(c) each are a partly-enlarged optical microscopic photograph thereof. For reference, FIG. 2(d) is an optical microscopic photograph of a die matrix before coated with a TiN film. In FIG. 2(d), the part looking white is a Cr carbide. As is obvious from FIG. 2(b) and FIG. 2(c), it is seen that, in the region where the coating film peeled off, a hard and coarse Cr carbide (a carbide mainly containing Cr and Fe and having a size of from about 1 to 50 μm) precipitated on the surface, and cracks were formed starting from the carbide.

From the above-mentioned analytical result, the present inventors have considered that the starting point for seizure is the above-mentioned coarse Cr carbide, and when the formation of the carbide is prevented as much as possible (or, when the carbide is not formed), then the surface coating film could be prevented from peeling off and the life of the die may be kept long.

Based on the above finding, the present inventors have further studied. As a result, the inventors have found the fact that, in order to prevent the formation of the coarse carbide thereby improving the above-mentioned properties, it is extremely important to suitably control the C amount and, in addition, to positively add various alloy ingredients to thereby suitably control the alloy ingredient planning. In detail, the inventors have found that, for obtaining the desired properties, it is effective to positively add alloy ingredients (especially Al, Cu, Ni, Mo and W) to thereby increase the hardness by precipitation hardening of the added alloy ingredients, but not increasing the hardness by carbide control as before, and mainly for this, precipitation hardening by an Al—Ni intermetallic compound and secondary hardening by carbide formation with Mo or W and C may be utilized. The inventors have made further experiments, and have reached the constitution of the present invention.

In the present specification, “having a high hardness” means that, when samples are analyzed for determining the maximum hardness thereof according to the method described in the section of Examples to be given hereinunder, those having a maximum hardness of at least 650 HV are defined as ones having a high hardness.

In the first aspect of the invention, “dimensional change (dimensional change rate) after heat treatment” is determined as follows: Samples are analyzed for determining the dimension thereof in three directions of the thickness, the width and the length before and after aging treatment, and the dimensional change is evaluated with both of the mean value thereof and the difference between the maximum value and the minimum value thereof. For convenience in description, the former is referred to as “mean value of dimensional change rate”; and the latter is as “dimensional change rate difference”. The first aspect of the invention differs from the technique of Patent Reference 2 in that, in the present invention, the dimensional change after heat treatment is evaluated based on both the “mean value of dimensional change rate” and the “dimensional change rate difference” while in the reference, only the former (mean value of dimensional change rate) is determined. Through experimental results, the present inventors have confirmed that, for sufficiently suppressing the dimensional change after heat treatment, the reduction in the mean value of the dimensional change rate as in Patent Reference 2 is unsatisfactory and it is indispensable to reduce the dimensional change (fluctuation) in all directions of the thickness, the width and the length, and that, even though the mean value of dimensional change rate could be reduced, the dimensional change rate difference may increase in some cases (and it is vice versa) (see Examples to be given hereinunder). In the first aspect of the invention, the phrase “dimensional change after heat treatment is small (that is, excellent in dimensional change resistance)” means that, when the dimensional change before and after heat treatment is determined according to the method described in the section of Examples to be given hereinunder, the mean value of the dimensional change rate is within a range of ±0.05%, and the dimensional change rate difference is 0.08% or less.

In the present specification, “welding repairability” is evaluated with the HAZ softening width. “Excellent in welding repairability” means that, when the HAZ softening width is determined according to the method described in the section of Examples given hereinunder, it is within a range of 6.5 mm or less.

The steel ingredients in the first aspect of the invention are as described in detail hereinunder. In the steel, not only the content of the alloy elements that contribute toward precipitation hardening is controlled to fall within a predetermined range but also the balance of predetermined elements is suitably controlled as defined by the formulae (1) to (3) mentioned below, whereby the above-mentioned properties of the steel are improved. As shown in Examples to be given hereinunder, those not satisfying any of those requirements could not have the desired properties. In particular, in the invention, it is indispensable to add all of Cu, Ni and Al to the steel; and for example, steel not containing any one of these ingredients as in the above-mentioned Patent Reference 1 or Patent Reference 3, could not attain the desired effect, which the inventors have confirmed through experiments (see Examples given hereinunder).

The steel ingredients in the first aspect of the invention are described simply as follows, with reference to the relationship between the “welding repairability” (evaluated by HAZ softening width) and “dimensional change resistance before and after heat treatment” (evaluated by both the dimensional change rate in the machine direction and the dimensional change rate difference), which are the principle objects to be improved in the first aspect of the invention.

First, for enhancing the welding repairability (for reducing the HAZ softening width), principally, it is important to suitably control the uppermost limit of [Cr]×[C], the Ms point (lowermost limit), the C amount (lowermost limit), the Al amount (lowermost limit), the Ni amount (lowermost limit), [Cu]/[Ni] (uppermost limit and lowermost limit), [Mo]+0.5×[W] (lowermost limit), and the V amount (uppermost limit). Specifically, as a planning guideline for reducing the HAZ softening width, not a hardening by martensite formation but a precipitation hardening by C content reduction to be low as from about 0.2 to 0.60% and by addition of alloy ingredients (mainly Al, Cu, Ni, Mo and W) (for example, ε—Cu, Ni—Al intermetallic compound or Ni—Mo intermetallic compound) is utilized. These precipitates are fine coherent precipitates in the matrix, and they significantly increase the hardness of the steel.

In particular, Cu, Ni and Al are important as precipitation hardening elements, and they are elements that greatly contribute toward inhibition of HAZ softening. A steel to which any of these elements is not substantially added could not have a desired HAZ softening preventing effect, which the inventors have confirmed through experiments.

Further, the ratio [Cu]/[Ni] (the ratio of [Cu] to [Ni]) has a close relation to HAZ softening inhibition, and HAZ softening can be inhibited by suitably controlling the above-mentioned ratio, as confirmed by the inventors. FIG. 6 is a graph showing the influence of the ratio [Cu]/[Ni] on the HAZ softening width, in which the HAZ softening width was determined according to the method described in Examples to be given hereinunder. On the graph, the data of Nos. 7, 8 and 10 in Table 3 and the data of Nos. 31 to 35 and 37 in Table 4 mentioned below are plotted. As shown in FIG. 6, the ratio [Cu]/[Ni] has a close relation to the HAZ softening width; and it is seen that, by controlling the above ratio to fall within a range of from 0.5 to 2.2, the HAZ softening width can be controlled to fall within the range defined by the invention (6.5 mm or less).

On the other hand, for reducing the dimensional change after heat treatment as much as possible, it is important to suitably control the product of the contents of Cr and C (uppermost limit of [Cr]×[C]), the C amount (uppermost limit), the Si amount (uppermost limit), the Mn amount (uppermost limit), the Ms point (lowermost limit), the Al amount (uppermost limit), the Ni amount (uppermost limit), the Cr amount (uppermost limit), and [Mo]+0.5×[W] (uppermost limit). The invention is based on low-C, in which, therefore, the Ms point is high and the formation of residual austenite is naturally small, and in addition, the content of the alloy ingredients such as Cu, Ni and Al is suitably controlled. Accordingly, in the invention, the expansion and contraction of steel after aging treatment at about 400 to 550° C. or after surface hardening treatment can be significantly retarded. The reason for this is considered to be as follows. Namely, owing to the addition of the above-mentioned alloy ingredients, for example, ε—Cu is mainly formed within a low temperature range of from about 400 to 500° C., an Ni—(Al, Mo) intermetallic compound is mainly within a middle temperature range of from about 450 to 530° C., and an Mo—V carbide is mainly within a high temperature range of from about 500 to 550° C.; however, since the crystal structure (FCC structure) of these precipitates differs from the matrix (BCC structure), the volume of the steel contracts to thereby contribute toward dimensional change inhibition after heat treatment. In addition, in the invention, since the ingredient planning is made for retarding coarse Cr carbide precipitation as much as possible, the crystal structure is isotropic in any direction, and even in producing a large-sized and complicated die stricture, the dimensional change after heat treatment thereof can be effectively inhibited.

The steel ingredients in the first aspect of the invention are described below. C: 0.20 to 0.60%

C is an element that ensures hardness and abrasion resistance, and contributes toward reduction in the HAZ softening width. In case where a carbide film such as VC or TiC is formed on the surface of a die matrix according to a CVD method, when the C concentration is low, then there may occur a problem in that the film thickness is insufficient. Taking it into consideration, the lowermost limit of the C amount to effectively exhibit the above action is 0.20%. Preferably, the C amount is 0.22% or more. However, when it is added too much, the residual austenite may increase and the steel could not have a desired hardness in the absence of high-temperature aging treatment, and additionally, the steel may expand after aging treatment and its dimensional change may increase. Accordingly, the uppermost limit of the C amount is 0.60%. Preferably, the C amount is 0.50% or less, more preferably 0.45% or less. Si: 0.5 to 2.00%

Si is an element useful as a deoxidizing element in steel production, and this element contributes toward enhancing the hardness and securing the machinability of steel. In addition, Si is useful for prevention of tempering softening of martensite in matrix and for suppression of HAZ softening width. For effectively exhibiting these actions, the lowermost limit of the Si amount is 0.5%. However, when it is added too much, the segregation may increase and the dimensional change after heat treatment may also increase, and additionally, the toughness may lower. Accordingly, the uppermost limit thereof is 2.00%. The lowermost limit of the Si amount is preferably 1%, more preferably 1.2%. On the other hand, the uppermost limit of the Si amount is preferably 1.85%.

Mn: 0.1 to 2%

Mn is an element useful for securing quenchability. However, when it is added too much, then the Ms point remarkably lowers and the residual austenite increases, so that the steel could not have the desired hardness in the absence of high-temperature aging treatment. Taking these into consideration, the Mn content is defined to fall within the above range. The lowermost limit of the Mn amount is preferably 0.15%; and on the other hand, the uppermost limit of the Mn amount is preferably 1%, more preferably 0.5%, even more preferably 0.35%.

Cr: 3.00 to 9.00%

Cr is an element useful for ensuring a predetermined hardness. When the Cr amount is less than 3.00%, then the quenchability may be poor and bainite may be partly formed to lower the hardness, and the abrasion resistance could not be ensured. Preferably, the Cr amount is 3.5% or more, more preferably 4.0% or more. However, when it is added too much, a coarse Cr carbide may be formed in large quantities, and the steel may contract after heat treatment and the film durability may lower. Accordingly, the uppermost limit of the ingredient is 9.00%. The Cr amount is preferably 7.0% or less, more preferably 6.5% or less, even more preferably 6.0% or less.

Al: 0.3 to 2.0%

Al is an element necessary for increasing the hardness through precipitation reinforcement with an Al—Ni intermetallic compound such as Ni₃Al, and this contributes toward suppression of HAZ softening width. In addition, Al is also useful as a deoxidizing agent. Taking these into consideration, the lowermost limit of Al is 0.3%. However, when it is added too much, the segregation may increase, and the dimensional change after heat treatment (especially the dimensional change rate difference) may increase, and the toughness may lower. Accordingly, the uppermost limit of the ingredient is 2.0%. Preferably, the Al amount is from 0.50% to 1.8%, more preferably from 0.7% to 1.6%.

Cu: 1.00 to 5%

Cu is an element necessary for increasing the hardness through precipitation reinforcement with ε—Cu, and this contributes toward HAZ softening width suppression. However, when it is added too much, the steel may crack during forging. Accordingly, its uppermost limit is 5%. The Cu amount is preferably from 2.0% to 4.0%.

Ni: 1.00 to 5%

Ni is an element necessary for increasing the hardness by precipitation reinforcement with an Al—Ni intermetallic compound such as Ni₃Al, and this contributes toward HAZ softening width suppression. In addition, when combined with Cu, Ni is effective for retarding hot work brittleness owing to excessive addition of Cu and for preventing cracking during forging. However, when it is added too much, residual austenite may increase and the steel could not ensure a predetermined hardness in the absence of high-temperature aging treatment, and, in addition, the steel may expand after heat treatment. The Ni amount is preferably from 1.5% to 4.0%.

Mo: 0.5 to 3% and/or W: 2% or Less (Including 0%)

Mo and W are both elements that form an M₆C carbide and form an Ni₃Mo intermetallic compound, and contribute toward precipitation reinforcement. However, when Mo and W are added too much, the above-mentioned carbide may be formed excessively to lower the toughness and, in addition, the dimensional change after heat treatment (especially dimensional change rate difference) may increase. Therefore, the above-mentioned range is defined herein. In the invention, Mo is an indispensable ingredient and W is an optional element, but the steel may contain both of them. The Mo content is preferably from 0.5% to 3%, more preferably from 0.7% to 2.5%. The W content is preferably 2% or less, more preferably 1.5% or less.

S: 0.10% or Less (not Including 0%)

S is an element useful for ensuring machinability. However, when it is added too much, it may cause weld cracking, and therefore, its uppermost limit is 0.10%. The S amount is preferably 0.07% or less, more preferably 0.05% or less, even more preferably 0.025% or less.

Further, the first aspect of the invention must satisfy the following requirements (1) to (3) {wherein each square bracket [ ] means a content (%) of each element}.

[Cr]×[C]≦3.00, (1)

The above requirement (1) is defined for the purpose of retarding the formation of a coarse Cr carbide. When the product of [Cr] and [C] exceeds 3.00, then the dimensional change after heat treatment may increase and the durability of the surface film may worsen. The product of [Cr] and [C] is preferably 1.80 or less, more preferably 1.70 or less. Its lowermost limit is preferably smaller from the viewpoint of preventing dimensional change after heat treatment. However, from the viewpoint of effectively exhibiting the above action owing to the addition of Cr and C, the lowermost limit thereof is preferably about 0.8.

[Cu]/N11:0.5 to 2.2, (2)

The above requirement (2) is defined mainly as a parameter for utilizing the precipitation reinforcement with ε—Cu and for suppressing the HAZ softening width (see Examples to be given hereinunder). For effectively exhibiting these actions, the ratio of [Cu] to [Ni] is defined to be 0.5. However, when the ratio is too large, then it may cause cracking during forging. Therefore, its uppermost limit is 2.2. Preferably, the ratio is from 0.7 to 1.5, more preferably from 0.85 to 1.2.

[Mo]+0.5×[W]:0.5 to 3.0%, (3)

Mo and W constituting the above requirement (3) are elements to contribute toward precipitation reinforcement, as so mentioned in the above. The above requirement (3) is defined as a parameter for ensuring the hardness enhancement by precipitation reinforcement with them, and is also effective for HAZ softening width suppression. In the above requirement (3), the coefficient for [W] (0.5) is defined in consideration of the fact that the atomic weight of Mo is about ½ of that of W. For effectively exhibiting these actions, the lowermost limit of the above requirement (3) is 0.5%. However, when the amount of Mo and W is too much, then the above-mentioned carbide may be formed excessively to lower the toughness and, if so, in addition, the dimensional change after heat treatment (especially dimensional change rate difference) may increase. Accordingly, the uppermost limit of the above requirement (3) is 3.0%. Preferably, the lowermost limit of the above requirement (3) is 1.0%, more preferably 1.2%; and on the other hand, the uppermost limit thereof is preferably 2.8%.

The steel ingredients in the first aspect of the invention are as described in the above, and the remainder is iron and unavoidable impurities. The unavoidable impurities are, for example, elements that may be unavoidably contained in the steel during its production process. For example, they include P, N and O. The P amount is preferably about 0.05% or less, more preferably 0.03% or less. The N amount is preferably about 350 ppm or less, more preferably 200 ppm or less, even more preferably 150 ppm or less. The 0 amount is preferably about 50 ppm or less, more preferably 30 ppm or less, even more preferably 20 ppm or less.

In the invention, the following ingredients may further be optionally added for the purpose of improving the other properties of steel.

V: 0.5% or Less (not Including 0%)

V contributes toward increasing the hardness, as forming a carbide such as VC, and is an element effective for HAZ softening width suppression. In addition, when a diffusion hardening layer is formed on the surface of the matrix through nitridation treatment such as vapor nitridation, bath salt nitridation, or plasma nitridation, V is effective for increasing the surface hardness and for increasing the hardened layer depth. For effectively exhibiting these actions, the V amount is preferably about 0.05% or more. However, when it is added too much, the dissolved amount of C may decrease and the hardness of the martensite texture of the matrix may lower. Therefore, the uppermost limit thereof is preferably 0.5%. The V amount is more preferably 0.4% or less, even more preferably 0.30% or less. At least one element selected from the group consisting of Ti, Zr, Hf, Ta and Nb in a total amount of 0.5% or less (not including 0%)

These elements are all nitride-forming elements, and these contribute toward toughness increase owing to fine dispersion and grain refining of the nitride and AlN. For effectively exhibiting these actions, preferably, the Ti amount is about 0.01% or more, the Zr amount is about 0.02% or more, the Hf amount is about 0.04% or more, the Ta amount is about 0.04% or more, and the Nb amount is about 0.02% or more. However, when these elements are added too much, the dissolved amount of C may decrease and the martensite hardness may decrease. Therefore, the total amount of the above elements is preferably 0.5%. More preferably, the total amount of these elements is 0.4% or less, even more preferably 0.30% or less. These elements may be added singly or in combination.

Co: 10% or Less (not Including 0%)

Co is an element effective for increasing the Ms point and for reducing residual austenite, by which the hardness of the steel increases. For effectively exhibiting the above actions, the Co amount is preferably about 1% or more. However, when it is added too much, then the cost may increase. Therefore, the uppermost limit thereof is preferably 10%. The uppermost limit of the Co amount is more preferably 5.5%.

Martensite Transformation Point (Ms Point)≧170° C.

Ms point=550−361×[C]−39×[Mn]−35×[V]−20×[Cr]

−17×[Ni]−10×[Cu]−5×([Mo]+[W])+

15×[Co]+30×[Al]

{wherein each square bracket [ ] means a content (%) of each element}

In the invention, the Ms point is essentially an index of the hardness and the dimensional change resistance after heat treatment of the steel. When the Ms point is lower than 170° C., then the residual austenite may increase and the steel could not have a desired hardness in the absence of high-temperature aging treatment, and if so, in addition, the steel may expand after heat treatment. The Ms point is preferably higher, more preferably 230° C. or higher, even more preferably 235° C. or higher, still more preferably 250° C. or higher. From the viewpoint of the above-mentioned effects, the uppermost limit of the Ms point is not specifically defined. However, in consideration of the action and the effect of the addition of the above-mentioned elements to constitute the Ms point, the uppermost limit thereof is preferably around 350° C., more preferably 320° C.

The invention also includes a die obtained by the use of the above-mentioned die steel. The method of producing the die is not specifically defined. For example, there is mentioned a method including preparing the above-mentioned steel in melt, then hot-forging and annealing it (for example, keeping at about 700° C. for 7 hours, then cooling in the furnace to about 400° C. at a mean cooling rate of about 17° C./hr, and thereafter further cooling it as such) to be softened, and thereafter cutting it into a predetermined shape in a mode of rough machining, then processing it for solution treatment at a temperature of from about 950 to 1150° C., followed by aging it at about 400 to 530° C. thereby to make it have a predetermined hardness.

Next, a method for producing a cold work die steel according to the second aspect of the invention is described in detail hereinunder.

The present inventors made further investigations for providing cold work die steels that satisfy the necessary properties for them, especially those having improved properties of hardness, dimensional change resistance after heat treatment and welding repairability (for life prolongation in repairing the damage of a die by welding). As a result, the inventors have found that, when the steel ingredients are suitably controlled, then the desired object can be attained (the first aspect of the invention).

Following the first aspect of the invention, the inventors have further made various investigations for further enhancing the dimensional change resistance after heat treatment of steel, based on the constitution of the steel ingredients disclosed in the prior patent application. As a result, the inventors have found that, when the steel of the first aspect of the invention is used and when it is subjected to a solution treatment and aging treatment under suitable conditions, then a cold work die steel having further enhanced dimensional change resistance after heat treatment can be efficiently obtained.

Specifically, the production method of the second aspect of the invention is characterized in that, of the first aspect of the invention, a favorable production condition for efficiently producing a cold work die steel having further enhanced dimensional change resistance after heat treatment is specifically defined. In detail, the method is characterized in that the solution treatment temperature and the aging temperature are defined by the parameter (ratio of Cu to C by mass) that most contributes toward the dimensional change resistance after heat treatment. According to the production method of the second aspect of the invention, a cold work die steel having further better dimensional change resistance after heat treatment than before can be obtained through only one-pass tempering treatment (aging treatment) like before, even in the absence of any specific heat treatment, for example, such as “at least one-pass two-stage tempering treatment” as in Patent Reference 5 or sub-zero treatment as in Patent Reference 6, and therefore the producibility of the method is extremely excellent.

First described is the process from the first aspect of the invention to the second aspect of the invention for the production method for such a cold work die steel.

The present inventors have first studied conventional dies formed of JIS SKD11 or matrix high-speed steel to clarify the reason why the surface film of the die is damaged and seizure is generated. As a result, the inventors have found that in the region where the film has peeled off, a rigid and coarse Cr carbide (a carbide of mainly Cr or Fe, having a size of from about 1 to 50 μm or so) is precipitated on the surface, and cracks are formed starting form the carbide.

The above is the process by which the inventors have reached the first aspect of the invention. After that, the inventors have made further investigations for providing a production method having high producibility, which is capable of easily producing a cold work die steel having much more excellent dimensional change resistance after heat treatment, by only one-pass solution treatment/aging treatment like before, without the necessity of any specific heat treatment. As a result, the inventors have found that, in solution treatment and aging treatment of the above-mentioned steel, when the treatment temperature (solution treatment temperature and aging temperature) is suitably defined in relation to the “ratio of Cu to C by mass” that most contributes toward the dimensional change resistance after heat treatment, as in Examples to be given hereinunder, then the intended object can be attained; and the inventors have thus completed the production method of the second aspect of the invention.

Concretely, the inventors have confirmed that, when the solution treatment temperature (° C.) is represented by T1, the aging treatment temperature (° C.) is represented by T2, the ratio by mass of Cu to C is represented by [Cu]/[C], and the numeral value represented by the following formula:

0.29×T1−2.63×[Cu]/[C]+225

is indicated by TA, then a process of solution treatment and aging treatment within a range where T2 satisfies the following formula (5) (that is, within a range of TA±10° C.):

TA−10≦T2≦TA+10 (5),

gives a steel having extremely excellent dimensional change resistance in that both the mean dimensional change rate and the maximum dimensional change rate (their details are described below) after heat treatment thereof satisfy the range of the second aspect of the invention (see Table 7 in Examples to be given hereinunder).

In the present specification, “solution treatment” as referred to has the same meaning as a quenching treatment; and “aging treatment” has the same meaning as a tempering treatment.

In the present specification, “high hardness” means that the hardness as determined according to the method described in the section of Examples to be given hereinunder is at least 650 HV.

In the first aspect of the invention, “dimensional change after heat treatment (dimensional change rate)” is determined as follows: Samples are analyzed for determining the dimension thereof in three directions of the thickness (Δx), the width (Δy) and the length (Δz) before and after aging treatment, and the dimensional change is evaluated with both of the mean value [Δx+Δy+Δz)/3] and the maximum value (absolute value) of the above Δx, Δy and Δz. For convenience in description, the former is referred to as “mean value of dimensional change rate, or mean dimensional change rate”; and the latter is as “maximum value of dimensional change rate or maximum dimensional change rate”. The second aspect of the invention differs from the technique of Patent Reference 2 in that, in second aspect of the present invention, the dimensional change after heat treatment is evaluated based on both the “mean value of dimensional change rate” and the “maximum value of dimensional change rate” while in the reference, only the former (mean value of dimensional change rate) is determined. Through experimental results, the present inventors have confirmed that, for sufficiently suppressing the dimensional change after heat treatment, the reduction in the mean value of the dimensional change rate as in Patent Reference 2 is unsatisfactory and it is indispensable to reduce the dimensional change (fluctuation) in all directions of the thickness, the width and the length, and that, even though the mean value of dimensional change rate could be reduced, the dimensional change rate difference may increase in some cases (and it is vice versa) (see Examples to be given hereinunder). In the second aspect of the invention, the phrase “dimensional change after heat treatment is small (that is, excellent in dimensional change resistance)” means that, when the dimensional change before and after heat treatment is determined according to the method described in the section of Examples to be given hereinunder, the mean value of the dimensional change rate is within a range of ±0.03%, and the maximum value (absolute value) of the dimensional change rate is 0.05% or less.

The evaluation standard (method and its level) in the second aspect of the invention mentioned in the above differs from that in the above-mentioned first aspect of the invention, in the following points.

First, both in the first aspect of the invention and in the second aspect of the invention, “mean value of dimensional change rate” is employed as the evaluation standard for dimensional change after heat treatment. However, in the first aspect of the invention, the pass line is ±0.05%, while in the second aspect of the invention, the pass line is ±0.03% and is severer than in the first aspect of the invention.

Further, in the first aspect of the invention, employed is the “dimensional change rate difference”, or that is the difference (absolute value) between the maximum value and the minimum value of the above-mentioned Δx, Δy and Δz; while in the second aspect of the invention, employed is “maximum value of dimensional change rate” as so mentioned in the above. This is based on the recognition that, “for providing a steel having more excellent dimensional change resistance than in the first aspect of the invention, the part in which the dimensional change (fluctuation) after heat treatment is the largest (maximum value) must be reduced as much as possible”, and on the basis of this, “the maximum value of dimensional change rate” is employed in the second aspect of the invention in addition to the “dimensional change rate difference” described in the section of the first aspect of the invention. As shown in Examples to be given hereinunder, some steels that satisfy the “dimensional change rate difference” defined in the first aspect of the invention could not satisfy the “maximum value of dimensional change rate” defined in the second aspect of the invention (see Examples to be given hereinunder). These steels could not be said to be “steels having excellent dimensional change resistance after heat treatment” in the second aspect of the invention.

The steel ingredients in the second aspect of the invention are as described in detail hereinunder. In the steel, not only the content of the alloy elements that contribute toward precipitation hardening is controlled to fall within a predetermined range but also the balance of predetermined elements is suitably controlled as defined by the formulae (1) to (4) mentioned below, whereby the above-mentioned properties of the steel are improved. As shown in Examples to be given hereinunder, those not satisfying any of those requirements could not have the desired properties. In particular, in second aspect of the invention, it is indispensable to add all of Cu, Ni and Al to the steel; and for example, steel not containing any one of these ingredients as in the above-mentioned Patent Reference 1 or Patent Reference 3, could not attain the desired effect, which the inventors have confirmed through experiments.

In particular, in the second aspect of the invention, for the purpose of minimizing the dimensional change after heat treatment as much as possible, it is important to suitably control not only the ratio by mass of [Cu] to [C] constituting the above-mentioned formula (5) but also the product of the content of Cr and C (uppermost limit of [Cr]×[C]), the C amount (uppermost limit), the Si amount (uppermost limit), the Mn amount (uppermost limit), the Ms point (lowermost limit), the Al amount (uppermost limit), the Ni amount (uppermost limit), the Cr amount (uppermost limit), and [Mo]+0.5×[W] (uppermost limit). The invention is based on low-C, in which, therefore, the Ms point is high and the formation of residual austenite is naturally small, and in addition, the content of the alloy ingredients such as Cu, Ni and Al is suitably controlled. Accordingly, in the invention, the expansion and contraction of steel after aging treatment at about 400 to 550° C. or after surface hardening treatment can be significantly retarded. The reason for this is considered to be as follows. Namely, owing to the addition of the above-mentioned alloy ingredients, for example, ε—Cu is mainly formed within a low temperature range of from about 400 to 500° C., an Ni—(Al, Mo) intermetallic compound is mainly within a middle temperature range of from about 450 to 530° C., and an Mo—V carbide is mainly within a high temperature range of from about 500 to 550° C.; however, since the crystal structure (FCC structure) of these precipitates differs from the matrix (BCC structure), the volume of the steel contracts to thereby contribute toward dimensional change inhibition after heat treatment. In addition, in the invention, since the ingredient planning is made for retarding coarse Cr carbide precipitation as much as possible, the crystal structure is isotropic in any direction, and even in producing a large-sized and complicated die stricture, the dimensional change after heat treatment thereof can be effectively inhibited.

In the second aspect of the invention, mainly, the uppermost limit of [Cr]×[C], the Ms point (lowermost limit), the C amount (lowermost limit), the Al amount (lowermost limit), the Ni amount (lowermost limit), [Cu]/[Ni] (uppermost limit and lowermost limit), [Mo]+0.5×[W] (lowermost limit) and the V amount (uppermost limit) are suitably controlled for enhancing the welding repairability (for suppressing the HAZ softening width). Specifically, as a planning guideline for reducing the HAZ softening width, not a hardening by martensite formation but a precipitation hardening by C content reduction to be low as from about 0.2 to 0.60% and by addition of alloy ingredients (mainly Al, Cu, Ni, Mo and W) (for example, ε—Cu, Ni—Al intermetallic compound or Ni—Mo intermetallic compound) is utilized. These precipitates are fine coherent precipitates in the matrix, and they significantly increase the hardness of the steel.

In particular, Cu, Ni and Al are important as precipitation hardening elements, and are elements that greatly contribute toward inhibition of HAZ softening. Steel to which any of these elements is not substantially added could not have a desired HAZ softening preventing effect, which the inventors have confirmed through experiments.

Further, the ratio [Cu]/[Ni] (the ratio of [Cu] to [Ni]) has a close relation to HAZ softening inhibition, and by suitably controlling the above-mentioned ratio, HAZ softening can be inhibited, as confirmed by the inventors.

The steel ingredients in the second aspect of the invention are described below.

C: 0.20 to 0.60%

Si: 0.5 to 2.00%

Mn: 0.1 to 2%

Cr: 3.00 to 9.00%

Al: 0.3 to 2.0%

Cu: 1.00 to 5%

Ni: 1.00 to 5%

Mo: 0.5 to 3% and/or W: 2% or Less (Including 0%)

S: 0.10% or Less (not Including 0%)

Further, the invention must satisfy the following requirements (1) to (4) {wherein each square bracket [ ] means a content (%) of each element}.

[Cr]×[C]≦3.00, (1)

[Cu]/N11:0.5 to 2.2, (2)

[Mo]+0.5×[W]:0.5 to 3.0%, (3)

Mo and W constituting the above requirement (3) are elements to contribute toward precipitation reinforcement, as so mentioned in the above. The above requirement (3) is defined as a parameter for ensuring the hardness enhancement by precipitation reinforcement with them, and is also effective for HAZ softening width suppression. In the above requirement (3), the coefficient for [W] (0.5) is defined in consideration of the fact that the atomic weight of Mo is about ½ of that of W. For effectively exhibiting these actions, the lowermost limit of the above requirement (4) is 0.5%. However, when the amount of Mo and W is too much, then the above-mentioned carbide may be formed excessively to lower the toughness and, if so, in addition, the dimensional change after heat treatment (especially dimensional change rate difference) may increase. Accordingly, the uppermost limit of the above requirement (3) is 3.0%. Preferably, the lowermost limit of the above requirement (3) is 1.0%, more preferably 1.2%; and on the other hand, the uppermost limit thereof is preferably 2.8%.

[Cu]/[C]:4.0 to 15; (4)

The above requirement (4) is positioned as a parameter mainly for shifting the peak of the hardness after heat treatment (after aging treatment) to a lower temperature side, and this is to ensure the dimensional change resistance after heat treatment. In general, it is said that the expansion-caused dimensional change after aging treatment (tempering) is caused by disappearance (decomposition) of the residual austenite during solution treatment (quenching) (for example, see FIG. 9 to be mentioned hereinunder); however, the inventors have found that, when the ratio by mass of Cu having an action of shifting the peak of the hardness after aging to a lower temperature side to C having a close correlation to residual austenite (ratio of [Cu]/[C]) is suitably controlled as in the above requirement (4), then the dimensional change after heat treatment can be remarkably retarded.

FIG. 1 is a graph showing an influence of the ratio [Cu]/[C] on the dimensional change, in which the dimensional change rate (mean value and maximum value) is determined according to the method described in Examples to be given hereinunder. On the graph, the data of No. 44 (steel A), 52 (steel C), 56 (steel D), 70 (steel J) and 73 (steel K) in Table 7 given below are plotted. These steels contain C, Si, Mn, Cr, Al, Cu, Ni, Mo and W nearly on the same level. As in FIG. 8, the ratio of [Cu]/[C] has a close relation to the dimensional change rate. It is seen that, by controlling the ratio to fall within a range of from 4.0 to 15, the dimensional change rate can be controlled to fall within the range defined in the second aspect of the invention (the mean value of the dimensional change rate is within a range of ±0.03° A), and the maximum value of the dimensional change rate is 0.05% or less).

When the ratio of [Cu]/[C] is less than 4.0, then the aging temperature at which the hardness reaches the peak is considerably higher than the temperature at which the residual austenite begins to decompose, and therefore the expansion after aging treatment increases, while on the other hand, when the ratio is more than 15, then the aging treatment does not follow contraction (to cancel the expansion after solution heat treatment); and anyhow, the predetermined dimensional change resistance could not be obtained in such cases. The above ratio is preferably from 5.0 to 13, more preferably from 6.0 to 12.

The steel ingredients in the second aspect of the invention are as described in the above, and the remainder is iron and unavoidable impurities. The unavoidable impurities are, for example, elements that may unavoidably mix in the steel during its production process. For example, they include P, N and O. The P amount is preferably about 0.05% or less, more preferably 0.03% or less. The N amount is preferably about 350 ppm or less, more preferably 200 ppm or less, even more preferably 150 ppm or less. The 0 amount is preferably about 50 ppm or less, more preferably 30 ppm or less, even more preferably 20 ppm or less.

In the invention, the following ingredients may be added for the purpose of improving the other properties of steel.

V: 0.5% or Less (not Including 0%)

V is effective for increasing the surface hardness and for increasing the hardened layer depth. For effectively exhibiting these actions, the V amount is preferably about 0.05% or more. However, when it is added too much, the dissolved amount of C may decrease and the hardness of the martensite texture of the matrix may lower. Therefore, the uppermost limit thereof is preferably 0.5%. The V amount is more preferably 0.4% or less, even more preferably 0.30% or less.

At Least One Element Selected from the Group Consisting of Ti, Zr, Hf, Ta and Nb in a Total Amount of 0.5% or Less (not Including 0%)

Co: 10% or Less (not Including 0%)

Martensite Transformation Point (Ms point)≧170° C.

Ms point=550−361×[C]−39×[Mn]−35×[V]−20×[Cr]

−17×[Ni]−10×[Cu]−5×([Mo]+[W])+

15×[Co]+30×[Al]

{wherein each square bracket [ ] means a content (%) of each element}

Next described is a method for producing a die steel according to the second aspect of the invention.

The production method of the invention includes steps of: preparing a steel satisfying the above-mentioned requirements, and subjecting the steel to a solution treatment and an aging treatment under the condition satisfying the following formula (5):

TA−10≦T2≦TA+10 (5)

wherein,

TA=0.29×T1−2.63×[Cu]/[C]+225,

T1 means the solution treatment temperature (° C.), and

T2 means the aging temperature (° C.).

Concretely, a steel satisfying the above-mentioned requirements is prepared in melt, then hot-forged and annealed (for example, kept at about 700° C. for 7 hours, then cooled in the furnace to about 400° C. at a mean cooling rate of about 17° C./hr, and thereafter kept cooled as such) to be softened, and thereafter cut into a predetermined shape in a mode of rough machining, then subjected to solution heat treatment and aging treatment under the condition of the above formula (5).

As so mentioned in the above, in the second aspect of the invention, the steel ingredients are so planned that the residual austenite amount during solution treatment is small. Further as in the above formula (5), when the ratio by mass of Cu to C ([Cu]/[C]) is controlled in relation to the solution treatment temperature T1 and the aging temperature T2, then the steel can be so controlled that its hardness could be the peak after aging but before expansion by residual austenite decomposition, and therefore the steel may satisfy both dimensional change resistance after heat treatment and hardness. In general, the process for die steel production comprises solution treatment at a temperature of from about 950 to 1150° C. followed by aging treatment at a temperature of from about 400 to 530° C., to thereby give a desired hardness to the steel. However, the present inventors have confirmed through experiments that even in solution treatment followed by aging treatment within the above-mentioned range, the steel produced could not often have a desired hardness and its dimensional change after heat treatment could not be sufficiently retarded (see Examples to be given hereinunder), and therefore, the inventors have specifically defined the above-mentioned formula (5).

The mechanism of the invention is compared with the method of the above-mentioned Patent Reference 2 (corresponding to conventional high-C high-Cr steel). In Patent Reference 2, as shown in FIG. 2 (corresponding to FIG. 1 in Patent Reference 2), the steel is tempered so that the dimensional change in tempering could be 0 at the time at which the residual austenite has been decomposed in some degree; while in the invention, the steel is tempered at a temperature at which the residual austenite is not as yet decomposed or at which it has just been decomposed; and in this point, the two differ from each other. Specifically, in the invention, the steel is aged almost at a lower temperature as compared with conventional high-C high-Cr steel (concretely, almost at a low temperature of about 500° C. or lower). According to the invention, the steel is aged not in a region in which the dimensional change after heat treatment is great (A in FIG. 3) as in Patent Reference 2 but in a region in which much stable residual austenite may be formed (B in FIG. 3), and therefore, it may be considered that, as compared with Patent Reference 2, the invention could produce a steel having reduced dimensional change or fluctuation. In case where steel is aged at such a relatively low temperature, the stability of the residual austenite therein may increase and the residual austenite changes little with time, and therefore, effectively, the dimensional change with time after heat treatment of the steel is also reduced.

The aging temperature T2 is preferably TA±5° C. with reference to the above formula.

The solution treatment temperature T1 may be lower than the temperature generally employed in die steel production. Accordingly, the steel of the invention may be free from deformation in thermal treatment. Concretely, the temperature is preferably within a range of from 900 to 1150° C.

In the invention, the solution treatment temperature and the aging temperature may be suitably controlled as in the above; and the time for the treatment is not specifically defined. The treatment may be effected under the condition generally employed for ordinary die steel production. Briefly, the solution treatment time (heating time) may be approximately from 1 to 5 hours or so, and the aging time (soaking time) may be controlled to be from 2 to 8 hours or so.

EXAMPLES

The invention is described more concretely with reference to the following Examples, by which, naturally, the invention should not be limited. Within the range within which the invention may attain the object mentioned in the above and to be mentioned below, the invention may be modified and changed suitably, and any of such modification and change shall be within the technical scope of the invention.

Examples of the first aspect of the invention are described below.

Various steel samples (Nos.) shown in Table 1 and Table 2 were used. In a vacuum induction melting furnace, 150 kg of an ingot was prepared in melt, then heated at about 900 to 1150° C., and forged into two sheets of 400 mmT×750 mmW×about 2000 mmL. Next, these were gradually cooled at a mean cooling rate of about 60° C./hr. After cooled to a temperature not higher than 100° C., these were again heated up to a temperature of about 850° C., and then gradually cooled at a mean cooling rate of about 50° C./hr (annealing).

The annealed samples prepared in the manner as above were tested for the following (1) to (4).

(1) Hardness Test (Determination of Maximum Hardness):

A test piece having a size of about 20 mmT×20 mmW×15 mmL was cut out of the above-mentioned annealed sample, and this was used as a hardness test piece. This was subjected to the following heat treatment cycle.

Solution heat treatment (quenching): heating at about 1020 to 1030° C. for 120 minutes→cooling by aeration→aging treatment (tempering): soaking at about 400 to 560° C. for about 3 hours→spontaneous cooling.

The hardness of the sample for which the tempering temperature was changed within the range of from about 400 to 560° C. as in the above, was measured with a Vickers' hardness meter (AKASHI's Model AVK, under load of 5 kg), and the maximum hardness (HV) of the sample was recorded. In this Example, those having a maximum hardness of at least 650 HV were evaluated as good (O).

(2) Dimensional Change Test (Determination of Mean Value of Dimensional Change Rate and Dimensional Change Rate Difference):

A test piece having a size of about 40 mmT×70 mmW×100 mmL was cut out of the above-mentioned annealed sample, and this was used as a dimensional change test piece. This was subjected to the same solution treatment as in the above (1) hardness test, and then this was tempered at the temperature at which it had the maximum hardness. Next, the “mean value of the dimensional change rate” and the “dimensional change rate difference” were determined in the manner mentioned below; and according to the standards mentioned below, the samples were evaluated. Those having passed the two were evaluated as good (O), as having excellent dimensional change resistance after heat treatment (pass-line level).

(2-1) Determination of Mean Value of Dimensional Change Rate:

The above-mentioned dimensional change test piece (after annealing but before solution treatment) and tempered test piece were measured in three directions of thickness, width and length. The difference in the thickness, the difference in the width and the difference in the length before and after the heat treatment were determined, and their mean value (percentage) was taken as “mean value of dimensional change rate”. In this Example, those having “mean value of dimensional change rate” to fall within ±0.05% were evaluated as good (O), and those over ±0.05% were evaluated as poor (x).

(2-2) Determination of Dimensional Change Rate Difference:

The above-mentioned dimensional change test piece (after annealing but before solution treatment) and tempered test piece were measured in three directions of thickness, width and length. The difference in the thickness, the difference in the width and the difference in the length before and after the heat treatment were determined. Of the data, the difference between the maximum value and the minimum value (percentage) was taken as “dimensional change rate difference”. Those having a dimensional change rate difference of 0.08% or less were evaluated as good (O), and those over 0.08% were evaluated as poor (x).

(3) Weld Test (Determination of Limiting Pre-Heating Temperature, and HAZ Softening Width):

A test piece having a size of about 40 mmT×45 mmW×75 mmL was cut out of the above-mentioned annealed sample, and this was used as a weld test piece. This was subjected to the same solution treatment and tempering treatment as in the above (2) dimensional change test.

Next, the thus-obtained tempered sample was worked into a plate of FIG. 3(a). The plate of FIG. 3(a) had a groove as shown in FIG. 3(b). Next, using a TIG wire (Eutectic of Japan's “TIG-Tectic 5HSS”, φ 2.4 mm) having the composition shown in Table 3 (remainder: iron and unavoidable impurities, unit: % by mass), the groove of the above plate was processed for overlying welding in the manner mentioned below.

Welding Condition:

Current: 150 A

Voltage: 11 V

Welding speed: 9.5 to 14 cm/mm

Interpass temperature: not higher than the preheating temperature

Heating: 7.1 to 10.4 kJ/cm

Preheating: absent or present (100° C., 200° C., 300° C. or 400° C.)

Regarding No. 22 and No. 23 in Table 2 (both are simulated steel samples of conventional high-C high-Cr steel), a welding material was buttered on the groove face as in FIG. 4, for the purpose of preventing the welding influence on the matrix component. For the buttering, used was a buttering TIG wire (KOBELCO's “TGS-50”) having the composition shown below, and the buttering was one-layer welding. The welding condition was the same as above.

Composition of Buttering TIG Wire: 0.09% C-0.93% Si-1.95% Mn-0.009% P-0.01% S (Remainder: Iron and Unavoidable Impurities, Unit: % by Mass)

The preheating condition was changed as in the above, whereupon the lowermost temperature at which both the welding metal (DEPO) and the HAZ part did not crack (limiting pre-heating temperature) was measured. Samples having a lower limiting pre-heating temperature are more hardly crackable. In this Example, those having a limiting pre-heating temperature of 200° C. or lower were evaluated as good (O), and those with higher than 200° C. were evaluated as poor (x).

For determining the hardness distribution in the cross section of the test piece processed for overlying welding at the above-mentioned limiting pre-heating temperature, the hardness of the sample was measured from the weld melt line (bond) position at a ¼ site of the plate thickness to the position spaced by 30 mm from it continuously at a pitch of 1 mm. The distance from the welding metal center part to the position at which the hardness lowered to at most 600 HV was referred to as “HAZ softening width”. For reference, the region for determination of the HAZ softening width is illustrated in above-mentioned FIG. 1. In this Example, the samples having a HAZ softening width of 6.5 mm or less were evaluated as good (O) in point of the welding repairability; and those with more than 6.5 mm were evaluated as poor (x).

(4) Toughness Test:

The above-mentioned annealed sample was processed for heat treatment as follows.

- Solution treatment (quenching): heating at about 1020 to 1030° C. for 120 minutes→cooling by aeration→aging treatment (tempering): soaking at about 400 to 560° C. for about 3 hours→cooling by aeration or spontaneously cooling.

Next, a test piece having a 10-mmR V notch was cut out to be a test piece for toughness determination (Charpy impact test piece), as in FIG. 5. The test piece was tested in a Charpy impact test, in which the absorption energy at room temperature of test piece was determined. Three Charpy impact test pieces were taken from one sample, and their data were averaged to be the Charpy impact value of the sample. In this Example, those having a Charpy impact value of 15 J or less were evaluated as “excellent in toughness”.

The results are shown in Tables 4 and 5.

TABLE 1

Remainder: iron and unavoidable impurities (unit: mass %)

No.
C
Si
Mn
Cr
Al
Cu
Ni
Mo
W
N
O

1
0.25
1.32
0.28
4.95
1.09
3.01
2.95
1.20
0.02
0.0148
0.0013

2
0.25
1.35
0.25
4.98
1.50
3.00
2.97
1.21
0.02
0.0145
0.0012

3
0.21
1.80
0.25
4.95
1.03
3.02
2.98
1.21
0.02
0.0134
0.0013

4
0.26
1.32
0.28
4.90
1.09
3.01
2.95
2.30
1.00
0.0129
0.0013

5
0.25
1.39
0.28
4.96
1.01
3.50
3.00
1.80
0
0.0141
0.0014

6
0.24
1.35
0.29
4.97
1.02
1.55
2.99
1.21
0.02
0.0129
0.0013

7
0.24
1.35
0.29
4.97
1.02
2.12
2.99
1.21
0.02
0.0133
0.0013

8
0.25
1.33
0.29
4.99
1.02
2.50
3.00
1.20
0.02
0.0140
0.0013

9
0.27
1.32
0.28
4.98
1.02
1.53
2.99
2.21
0.55
0.0137
0.0014

10
0.24
1.33
0.28
4.99
1.02
3.01
2.03
1.22
0.02
0.0131
0.0013

11
0.24
1.35
0.29
4.97
1.02
1.55
2.99
1.19
0.02
0.0148
0.0013

12
0.25
1.36
0.28
6.80
1.01
3.02
2.97
1.20
0.02
0.0131
0.0015

13
0.25
1.34
0.26
4.98
1.01
3.02
2.95
1.22
0.02
0.0138
0.0013

14
0.25
1.33
0.28
4.99
1.02
3.01
2.97
1.20
0.02
0.0131
0.0014

15
0.25
1.33
0.28
4.98
1.01
3.00
2.98
1.24
0.02
0.0135
0.0015

16
0.25
1.39
0.27
4.97
1.00
3.03
2.97
1.20
0.02
0.0142
0.0013

17
0.40
1.35
0.25
4.45
1.03
3.00
2.98
1.21
0.02
0.0140
0.0013

18
0.42
1.35
0.26
4.10
1.49
3.01
2.98
1.22
0.02
0.0141
0.0013

19
0.41
1.38
0.25
4.10
1.03
3.01
2.98
1.21
0.02
0.0138
0.0013

20
0.44
1.33
0.15
3.99
1.03
3.00
2.99
1.23
0.00
0.0133
0.0013

21
0.40
1.35
0.25
4.45
1.03
3.02
2.98
1.21
0.00
0.0129
0.0013

Remainder: iron and unavoidable

impurities (unit: mass %)
[Cr] ×
[Cu]/
[Mo] +

No.
P
S
V
Ti
Nb
Zr
Hf
Ta
Co
[C]
[Ni]
[W]/2
Ms

1
0.018
0.004
0.20
0
0
0
0
0
0
1.24
1.02
1.21
289

2
0.019
0.003
0.20
0
0
0
0
0
0
1.25
1.01
1.22
302

3
0.019
0.003
0.21
0
0
0
0
0
0
1.04
1.01
1.22
302

4
0.018
0.004
0.20
0
0
0
0
0
0
1.27
1.02
2.80
279

5
0.020
0.005
0.20
0
0
0
0
0
0
1.24
1.17
1.80
278

6
0.019
0.004
0.21
0
0
0
0
0
0
1.19
0.52
1.22
303

7
0.018
0.004
0.20
0
0
0
0
0
0
1.19
0.71
1.22
298

8
0.019
0.004
0.21
0
0
0
0
0
0
1.25
0.83
1.21
290

9
0.019
0.004
0.20
0
0
0
0
0
0
1.34
0.51
2.49
287

10
0.018
0.004
0.20
0
0
0
0
0
0
1.20
1.48
1.23
305

11
0.019
0.004
0
0
0
0
0
0
0
1.19
0.52
1.20
311

12
0.019
0.004
0.20
0
0
0
0
0
0
1.70
1.02
1.21
249

13
0.019
0.005
0.20
0
0
0
0
0
0
1.25
1.02
1.23
287

14
0.018
0.004
0.21
0
0.1
0
0
0
0
1.25
1.01
1.21
286

15
0.019
0.004
0.20
0
0
0.1
0
0
0
1.25
1.01
1.25
286

16
0.020
0.004
0.21
0
0
0
0.1
0.1
0
1.24
1.02
1.21
286

17
0.019
0.004
0.20
0
0
0
0
0
0
1.78
1.01
1.22
244

18
0.018
0.004
0.22
0
0
0
0
0
0
1.72
1.01
1.23
256

19
0.019
0.004
0.20
0
0
0
0
0
0
1.68
1.01
1.22
247

20
0.017
0.004
0.41
0
0
0
0
0
0
1.76
1.00
1.23
235

21
0.019
0.004
0.20
0
0
0
0
0
5
1.78
1.01
1.21
319

TABLE 2

Remainder: iron and unavoidable impurities (unit: mass %)

No.
C
Si
Mn
Cr
Al
Cu
Ni
Mo
W
N
O
P

22
1.49
0.35
0.42
12.10
0.05
0.05
0.08
1.04
0.35
0.0130
0.0015
0.018

23
1.01
1.06
0.60
8.38
0.33
0.40
0.44
0.91
0.39
0.0068
0.0007
0.019

24
0.19
1.32
0.28
4.93
1.02
3.01
2.97
1.20
0.02
0.0140
0.0013
0.018

25
0.61
1.32
0.28
4.95
1.02
3.01
2.95
1.20
0.02
0.0131
0.0013
0.020

26
0.25
2.02
0.28
4.95
1.01
3.02
2.97
1.21
0.02
0.0135
0.0015
0.019

27
0.25
1.32
2.10
4.96
1.01
3.01
2.98
1.22
0.02
0.0142
0.0013
0.019

28
0.24
1.35
0.28
4.94
1.05
3.03
2.97
1.21
0.02
0.0140
0.0014
0.019

29
0.25
1.35
0.28
4.90
0.48
3.01
2.98
1.23
0.02
0.0141
0.0015
0.019

30
0.25
1.36
0.29
4.97
2.05
3.02
2.98
1.22
0.02
0.0135
0.0013
0.018

31
0.25
1.34
0.29
4.96
1.01
3.01
0.95
1.21
0.02
0.0133
0.0013
0.018

32
0.24
1.33
0.28
4.96
1.02
3.02
6.10
1.23
0.02
0.0140
0.0015
0.018

33
0.25
1.33
0.26
4.93
1.02
0.98
2.98
1.21
0.02
0.0131
0.0013
0.019

34
0.24
1.32
0.27
4.97
1.09
0.05
2.97
1.21
0.02
0.0129
0.0014
0.018

35
0.24
1.32
0.27
4.98
1.09
3.01
0.05
1.20
0.02
0.0135
0.0014
0.018

36
0.24
1.32
0.27
4.93
0.05
3.01
2.97
1.20
0.02
0.0141
0.0013
0.018

37
0.25
1.33
0.26
4.98
1.02
3.03
1.32
1.21
0.02
0.0135
0.0015
0.019

38
0.25
1.39
0.28
2.99
1.01
3.01
2.99
1.23
0.02
0.0135
0.0013
0.018

39
0.25
1.35
0.28
9.01
1.01
3.02
2.97
1.21
0.02
0.0132
0.0014
0.020

40
0.26
1.32
0.27
4.94
1.02
3.01
2.98
0.20
0.20
0.0140
0.0015
0.019

41
0.24
1.34
0.25
4.95
1.01
3.03
2.99
2.15
2.00
0.0141
0.0013
0.019

42
0.25
1.33
0.28
4.97
1.03
3.02
2.95
1.20
0.02
0.0138
0.0014
0.019

43
0.25
1.32
0.26
4.95
1.04
3.01
2.96
1.20
0.02
0.0133
0.0013
0.019

Remainder: iron and unavoidable

impurities (unit: mass %)
[Cr] ×
[Cu]/
[Mo] +

No.
S
V
Ti
Nb
Zr
Hf
Ta
Co
[C]
[Ni]
[W]/2
Ms

22
0.005
0.25
0
0
0
0
0
0
18.03
0.63
1.22
−261

23
0.007
0.09
0
0.1
0
0
0
0
8.46
0.91
1.11
−16

24
0.004
0.20
0
0
0
0
0
0
0.94
1.01
1.21
309

25
0.004
0.20
0
0
0
0
0
0
3.02
1.02
1.21
157

26
0.004
0.21
0
0
0
0
0
0
1.24
1.02
1.22
286

27
0.004
0.20
0
0
0
0
0
0
1.24
1.01
1.23
215

28
0.110
0.21
0
0
0
0
0
0
1.19
1.02
1.22
291

29
0.004
0.20
0
0
0
0
0
0
1.23
1.01
1.24
271

30
0.004
0.22
0
0
0
0
0
0
1.24
1.01
1.23
316

31
0.005
0.20
0
0
0
0
0
0
1.24
3.17
1.22
320

32
0.004
0.20
0
0
0
0
0
0
1.19
0.50
1.24
237

33
0.004
0.21
0
0
0
0
0
0
1.23
0.33
1.22
308

34
0.004
0.20
0
0
0
0
0
0
1.19
0.02
1.22
322

35
0.004
0.20
0
0
0
0
0
0
1.20
60.20
1.21
342

36
0.004
0.20
0
0
0
0
0
0
1.18
1.01
1.21
262

37
0.004
0.20
0
0
0
0
0
0
1.25
2.30
1.22
315

38
0.004
0.20
0
0
0
0
0
0
0.75
1.01
1.24
325

39
0.004
0.22
0
0
0
0
0
0
2.25
1.02
1.22
204

40
0.003
0.20
0
0
0
0
0
0
1.28
1.01
0.30
288

41
0.004
0.20
0
0
0
0
0
0
1.19
1.01
3.15
281

42
0.005
0.60
0
0
0
0
0
0
1.24
1.02
1.21
273

43
0.004
0.20
0.40
0
0
0
0
0
1.24
1.02
1.21
288

TABLE 3

Remainder: iron and unavoidable impurities (unit: mass %)

C
Si
Mn
V
Ni
Cr
Mo
Cu
Al
P
S
V

0.25
0.40
1.40
0.2
3.0
5.0
2.3
3.0
1.0
0.02
0.005
0.2

TABLE 4

Maximum
HAZ softening
Limiting pre-heating
Mean value of dimensional
Dimensional change rate
Charpy impact

hardness
width
temperature
change rate
difference
value

No.
HV
mm
° C.
%
%
J

1
685
5.1
100
0.01
0.04
22

2
702
5.3
100
−0.01
0.05
20

3
710
4.8
25
0
0.06
19

4
715
4.5
100
0.02
0.05
17

5
713
4.6
100
−0.02
0.04
18

6
655
6.4
100
0.02
0.07
27

7
663
5.9
100
0.01
0.05
24

8
679
5.6
100
0
0.05
22

9
712
6.2
100
0.02
0.06
17

10
682
5.6
100
0
0.04
20

11
660
6.4
100
0.01
0.04
25

12
688
5.0
100
−0.05
0.08
22

13
683
5.1
100
0.01
0.04
35

14
683
5.0
100
0.01
0.04
33

15
685
5.1
100
0.01
0.04
33

16
680
5.0
100
0.01
0.04
32

17
710
5.2
200
0.03
0.05
17

18
720
5.3
200
0.03
0.04
16

19
685
5.1
200
0.03
0.05
22

20
680
5.3
200
0.04
0.05
21

21
722
4.9
200
0
0.03
25

TABLE 5

Maximum
HAZ softening
Limiting pre-heating
Mean value of dimensional
Dimensional change rate
Charpy impact

hardness
width
temperature
change rate
difference
value

No.
HV
mm
° C.
%
%
J

22
690
11.0
400
0.06
0.15
10

23
720
10.5
400
0.01
0.12
13

24
645
7.0
25
−0.02
0.04
35

25
715
5.2
300
0.06
0.10
14

26
725
4.5
100
0.05
0.09
12

27
722
5.1
100
0.06
0.08
15

28
688
5.3
300
0.01
0.04
19

29
638
8.4
100
0.04
0.05
33

30
700
4.6
100
−0.02
0.09
19

31
640
7.5
100
0.03
0.04
20

32
648
6.3
100
0.06
0.08
40

33
645
6.6
100
0.05
0.04
37

34
620
7.5
100
0.06
0.03
30

35
625
7.3
100
0.06
0.04
29

36
630
7.0
100
0.06
0.03
33

37
661
7.0
100
0.01
0.03
19

38
647
6.5
100
0.02
0.05
39

39
683
4.9
100
−0.06
0.09
21

40
629
6.9
100
0.05
0.05
35

41
723
5.0
100
0.03
0.10
13

42
640
7.0
100
0.02
0.05
35

43
630
7.7
100
0.02
0.04
33

Table 4 and Table 5 bring about the following discussion.

Nos. 1 to 21 in Table 4 are the data of the samples Nos. 1 to 21 in Table 1 satisfying all the requirements of the invention. These all have high hardness and are excellent in the dimensional change resistance after heat treatment and in the welding repairability, and in addition, these are highly tough and are good in that their limiting preheating temperature is 200° C. or less.

As opposed to these, Nos. 22 to 43 in Table 5 are the data of the samples Nos. 22 to 43 in Table 2 not satisfying any of the requirements defined in the invention, and therefore have the following drawbacks.

Nos. 22 and 23 in Table 5 are the data of Nos. 22 and 23 in Table 2 that are simulated steel samples of conventional high-C high-Cr steel. These had increased HAZ softening width and increased dimensional change, since the product of [Cr] and [C] was large and the Ms point was low. These steels had more increased hardness when the tempering temperature was lower. Therefore, the tempering temperature for these steels was 510° C., and the properties thereof were measured.

No. 24 in Table 5 is the data of No. 24 in Table 2, in which the C amount was small. The sample had lowered hardness and increased HAZ softening width.

No. 25 in Table 5 is the data of No. 25 in Table 2, in which the C amount was large, the product of [Cr] and [C] was large and the Ms point was low. The sample was poor in the dimensional change resistance after heat treatment.

No. 26 in Table 5 is the data of No. 26 in Table 2, in which the Si amount was large. The sample was good in point of the mean dimensional change rate after heat treatment, but its dimensional change rate difference was large.

No. 27 in Table 5 is the data of No. 27 in Table 2, in which the Mn amount was large and the Ms point was low. The mean dimensional change rate after heat treatment was large.

No. 28 in Table 5 is the data of No. 28 in Table 2, in which the S amount was large. The limiting preheating temperature was high and the sample has a risk of weld cracking.

No. 29 in Table 5 is the data of No. 29 in Table 2, in which the Al amount was small. The hardness was low and the HAZ softening width increased.

No. 30 in Table 5 is the data of No. 30 in Table 2, in which the Al amount was large. The mean dimensional change rate after heat treatment was not so large, but the dimensional change rate difference was large.

No. 31 in Table 5 is the data of No. 31 in Table 2, in which the Ni amount is small and the ratio of [Cu]/[Ni] was large. The hardness was low and the HAZ softening width increased.

No. 32 in Table 5 is the data of No. 32 in Table 2, in which the Ni amount was large. The hardness was low and the mean dimensional change rate after heat treatment increased.

No. 33 in Table 5 is the data of No. 33 in Table 2, in which the Cu amount was small and the ratio [Cu]/[Ni] was small. The hardness was low and the HAZ softening width increased.

No. 34 in Table 5 is the data of No. 34 in Table 2, which is a simulated sample of substantially Cu-free steel. In this, the Cu amount was 0.05% and was extremely small, and the ratio of [Cu]/[Ni] was small. The hardness was low and the HAZ softening width increased. In addition, the mean dimensional change rate after heat treatment increased.

No. 35 in Table 5 is the data of No. 35 in Table 2, which is a simulated sample of substantially Ni-free steel. In this, the Ni amount was 0.05% and was extremely small, and the ratio of [Cu]/[Ni] was small. The hardness was low and the HAZ softening width increased. In addition, the mean dimensional change rate after heat treatment increased.

No. 36 in Table 5 is the data of No. 36 in Table 2, which is a simulated sample of substantially Al-free steel. In this, the Al amount was 0.05% and was extremely small. The hardness was low and the HAZ softening width increased. In addition, the mean dimensional change rate after heat treatment increased.

No. 37 in Table 5 is the data of No. 37 in Table 2, in which the Cu amount and the Ni amount satisfy the range of the invention, but the ratio of [Cu]/[Ni] was small. The HAZ softening width increased.

No. 38 in Table 5 is the data of No. 38 in Table 2, in which the Cr amount was small. The hardness was low.

No. 39 in Table 5 is the data of No. 39 in Table 2, in which the Cr amount was large. The dimensional change resistance after heat treatment was poor.

No. 40 in Table 5 is the data of No. 40 in Table 2, in which the total amount of [Mo]+0.5×[W] was small. The hardness was low and the HAZ softening width increased.

No. 41 in Table 5 is the data of No. 41 in Table 2, in which the total amount of [Mo]+0.5×[W] was large. The mean dimensional change rate after heat treatment was not so large, but the dimensional change rate difference was large.

No. 42 in Table 5 is the data of No. 42 in Table 2, in which the Ti amount was large. The hardness was low and the HAZ softening width increased.

For reference, the hardness distribution profile of the samples obtained according to the above-mentioned method is shown in FIG. 7. In the drawing, the steel of the invention (black square) is No. 4 in Table 1; and the conventional SKD11 steel (black lozenge) is No. 22 in Table 2. As in FIG. 7, the steel of the invention may greatly prevent HAZ softening after welding, as compared with the conventional steel.

Next described are Examples of the second aspect of the invention.

Various steel samples A to K shown in Table 6 were used. In a vacuum induction melting furnace, 150 kg of an ingot was prepared in melt, then heated at about 900 to 1150° C., and forged into two sheets of 400 mmT×750 mmW×about 2000 mmL. Next, these were gradually cooled at a mean cooling rate of about 60° C./hr. After cooled to a temperature not higher than 100° C., these were again heated up to a temperature of about 850° C., and then gradually cooled at a mean cooling rate of about 50° C./hr (annealing).

The annealed samples prepared in the manner as above were tested for the following (1) and (2).

(1) Hardness Test (Determination of Maximum Hardness):

A test piece having a size of about 20 mmT×20 mmW×15 mmL was cut out of the above-mentioned annealed sample, and this was used as a hardness test piece. This was processed for solution treatment→cooling by aeration→aging treatment under the condition shown in Table 2, and then left cooled. In every case, the solution treatment time was about 120 minutes, and the aging time was about 3 hours.

After aged, the hardness (HV) of the sample was measured with a Vickers' hardness meter (AKASHI's Model AVK, under load of 5 kg). In this Example, those having a hardness of at least 650 HV were evaluated as good (O).

(2) Dimensional Change Test (Determination of Mean Value of Dimensional Change Rate and Maximum Value of Dimensional Change Rate):

A test piece having a size of about 40 mmT×70 mmW×100 mmL was cut out of the above-mentioned annealed sample, and this was used as a dimensional change test piece. This was processed for solution treatment→cooling by fan aeration→aging treatment under the condition shown in Table 2, and then left cooled. Next, the “mean value of the dimensional change rate” and the “maximum value of the dimensional change rate” were determined in the manner mentioned below; and according to the standards mentioned below, the samples were evaluated. Those having passed the two were evaluated as good (O), as having excellent dimensional change resistance after heat treatment (pass-line level).

(2-1) Determination of Mean Value of Dimensional Change Rate (Mean Dimensional Change Rate):

The above-mentioned dimensional change test piece (after annealing but before solution treatment) and tempered test piece were measured in three directions of thickness, width and length. The difference in the thickness, the difference in the width and the difference in the length before and after the heat treatment were determined, and their mean value (percentage) was taken as the “mean value of dimensional change rate”. In this Example, those having “mean value of dimensional change rate” to fall within ±0.03% were evaluated as good (O), and those over ±0.03% were evaluated as poor (x).

(2-2) Determination of Maximum Value of Dimensional Change Rate (Maximum Dimensional Change Rate):

The above-mentioned dimensional change test piece (after annealing but before solution treatment) and tempered test piece were measured in three directions of thickness, width and length. The difference in the thickness, the difference in the width and the difference in the length before and after the heat treatment were determined. Of the data, the absolute value (percentage) of the maximum value was taken as the “maximum value of dimensional change rate”. Those having a maximum dimensional change rate of 0.05% or less were evaluated as good (O), and those over 0.05% were evaluated as poor (x).

(2-3) Determination of Dimensional Change Rate Difference:

For reference, the “dimensional change rate difference” in the description relating to the first aspect of the invention was also determined. Concretely, the above-mentioned dimensional change test piece (after annealing but before solution treatment) and tempered test piece were measured in three directions of thickness, width and length. The difference in the thickness, the difference in the width and the difference in the length before and after the heat treatment were determined. Of the data, the difference between the maximum value and the minimum value (percentage) was taken as the “dimensional change rate difference”. Those having a dimensional change rate difference of 0.08% or less were evaluated as good (O), and those over 0.08% were evaluated as poor (x).

The results are shown in Table 7.

TABLE 6

Remainder: iron and unavoidable impurities (unit: mass %)

Steel
C
Si
Mn
Cr
Al
Cu
Ni
Mo
W
N
O
P

A
0.25
1.32
0.28
4.95
1.09
3.01
2.95
1.20
0.02
0.0148
0.0013
0.018

B
0.25
1.39
0.28
4.96
1.01
3.00
2.98
1.80
0
0.0141
0.0014
0.020

C
0.24
1.35
0.29
4.97
1.02
1.55
2.99
1.21
0.02
0.0129
0.0013
0.019

D
0.25
1.33
0.29
4.99
1.02
2.50
3.00
1.20
0.02
0.0140
0.0013
0.019

E
0.25
1.39
0.28
4.96
1.01
3.50
3.00
1.80
0
0.0141
0.0014
0.020

F
0.40
1.35
0.25
4.45
1.03
3.00
2.98
1.21
0.02
0.0140
0.0013
0.019

G
0.45
1.35
0.25
4.45
0.35
2.00
1.00
0.9
0
0.0140
0.0013
0.019

H
1.49
0.35
0.42
12.10
0.05
0.05
0.08
1.04
0.35
0.0130
0.0015
0.018

I
1.01
1.06
0.60
8.38
0.33
0.40
0.44
0.91
0.39
0.0068
0.0007
0.019

J
0.25
1.33
0.26
4.93
1.02
0.98
2.98
1.21
0.02
0.0131
0.0013
0.019

K
0.25
1.33
0.26
4.93
1.02
4.00
2.98
1.21
0.02
0.0131
0.0013
0.019

Remainder: iron and unavoidable

impurities (unit: mass %)
[Cr] ×
[Cu]/
[Mo] +

[Cu]/

Steel
S
V
Ti
Nb
Zr
Hf
Ta
Co
[C]
[Ni]
[W]/2
Ms
[C]

A
0.004
0
0
0
0
0
0
0
1.24
1.02
1.21
296
12.04

B
0.005
0.20
0
0
0
0
0
0
1.24
1.01
1.80
283
12.00

C
0.004
0.21
0
0
0
0
0
0
1.19
0.52
1.22
303
6.46

D
0.004
0.21
0
0
0
0
0
0
1.25
0.83
1.21
290
10.00

E
0.005
0.20
0
0
0
0
0
0
1.24
1.17
1.80
278
14.00

F
0.004
0.20
0
0
0
0
0
0.5
1.78
1.01
1.22
251
7.50

G
0.004
0.20
0
0
0
0
0
0
2.00
2.00
0.90
251
4.44

H
0.005
0.25
0
0
0
0
0
0
18.03
0.63
1.22
−261
0.03

I
0.007
0.09
0
0.1
0
0
0
0
8.46
0.91
1.11
−16
0.40

J
0.004
0.21
0
0
0
0
0
0
1.23
0.33
1.22
308
3.92

K
0.004
0.21
0
0
0
0
0
0
1.23
1.34
1.22
278
16.00

TABLE 7

Aging temperature T2 (° C.)

Solution

Range of

Treatment

the

Dimensional change rate (%)

Temperature

invention
Measured
Hardness
Mean
Maximum

No.
Steel
T1 (° C.)
TA*
(TA ± 10)
value
(HV)
value
Value
Difference

44
A
1020
489
479-499
490
675
0.010
0.030
0.040

45

950
469
459-479
470
678
0.000
0.020
0.040

46

1020
489
479-499
520
640
0.035
0.060
0.050

47

1020
489
479-499
460
642
0.035
0.060
0.050

48
B
1020
489
479-499
490
685
0.010
0.030
0.040

49

950
469
459-479
470
683
0.000
0.020
0.040

50

1020
489
479-499
520
652
0.040
0.065
0.050

51

1020
489
479-499
460
649
0.035
0.060
0.050

52
C
1020
504
494-514
500
655
0.020
0.050
0.060

53

950
484
474-494
485
654
0.010
0.045
0.070

54

1020
504
494-514
530
625
0.030
0.065
0.070

55

1020
504
494-514
460
620
0.035
0.070
0.070

56
D
1020
495
485-505
495
679
0.000
0.025
0.050

57

950
474
464-484
475
675
−0.010
0.035
0.050

58

1020
495
485-505
530
670
0.030
0.055
0.050

59
E
1020
484
474-494
485
713
−0.020
0.040
0.040

60

950
464
454-474
465
710
−0.030
0.050
0.035

61

950
464
454-474
480
680
−0.050
0.070
0.035

62
F
1020
501
491-511
500
710
0.030
0.050
0.040

63

950
481
471-491
480
720
0.010
0.030
0.035

64

1020
501
491-511
520
675
0.050
0.070
0.040

65
G
1020
509
499-519
510
700
0.030
0.050
0.040

66

950
489
479-499
490
695
0.020
0.040
0.040

67

1020
509
499-519
530
690
0.060
0.080
0.040

68
H
1020
521
511-531
520
690
0.060
0.135
0.150

69
I
1020
520
510-530
520
720
0.010
0.070
0.120

70
J
1020
510
500-520
510
645
0.030
0.060
0.060

71

950
490
480-500
490
640
0.025
0.055
0.060

72
K
1020
479
469-489
480
720
0.025
0.055
0.060

73

950
458
448-468
460
715
0.020
0.065
0.060

TA* = 0.29 × T1 − 2.63 × [Cu]/[C] + 225

Table 7 brings about the following discussion.

First, in Nos. 44 to 47 in Table 7, the steel A in Table 6 of which the steel ingredients satisfy the requirements of the invention was processed at different solution treatment temperature T1 and aging temperature T2, and the samples were analyzed for their physical properties.

Of those, No. 44 and No. 45 are the samples of the invention for which the aging temperature T2 satisfies the range of the invention (TA±10° C.). These both had high hardness and were excellent in the dimensional change resistance after heat treatment (not only in point of the dimensional change rate difference but also in all the mean dimensional change rate and the maximum dimensional change rate).

As opposed to these, No. 46 is a comparative sample for which the aging temperature T2 is higher than the range of the invention; No. 47 is a comparative sample for which the aging temperature T2 is lower than the range of the invention. Both the two had low hardness and, in addition, though they were good in point of the dimensional change rate difference after heat treatment but were not good in point of the mean dimensional change rate and the maximum dimensional change rate.

In Nos. 48 to 51 in Table 7, the steel B in Table 6 of which the steel ingredients satisfy the requirements of the invention was processed at different solution treatment temperature T1 and aging temperature T2, and the samples were analyzed for their physical properties.

Of those, No. 48 and No. 49 are the samples of the invention for which the aging temperature T2 satisfies the range of the invention (TA±10° C.). These both had high hardness and were excellent in the dimensional change resistance after heat treatment.

As opposed to these, No. 50 is a comparative sample for which the aging temperature T2 is higher than the range of the invention; No. 51 is a comparative sample for which the aging temperature T2 is lower than the range of the invention. Though both the two were good in point of the dimensional change rate difference after heat treatment but were not good in point of the mean dimensional change rate and the maximum dimensional change rate. In addition, the hardness of No. 51 was low.

In Nos. 52 to 55 in Table 7, the steel C in Table 6 of which the steel ingredients satisfy the requirements of the invention was processed at different solution treatment temperature T1 and aging temperature T2, and the samples were analyzed for their physical properties.

Of those, No. 52 and No. 53 are the samples of the invention for which the aging temperature T2 satisfies the range of the invention (TA±10° C.). These both had high hardness and were excellent in the dimensional change resistance after heat treatment.

As opposed to these, No. 54 is a comparative sample for which the aging temperature T2 is higher than the range of the invention; No. 55 is a comparative sample for which the aging temperature T2 is lower than the range of the invention. The hardness was low, and the maximum dimensional change rate after heat treatment increased. In addition, the mean dimensional change rate of No. 54 after heat treatment increased.

In Nos. 56 to 58 in Table 7, the steel D in Table 6 of which the steel ingredients satisfy the requirements of the invention was processed at different solution treatment temperature T1 and aging temperature T2, and the samples were analyzed for their physical properties.

Of those, No. 56 and No. 57 are the samples of the invention for which the aging temperature T2 satisfies the range of the invention (TA±10° C.). These both had high hardness and were excellent in the dimensional change resistance after heat treatment.

As opposed to these, No. 58 is a comparative sample for which the aging temperature T2 is higher than the range of the invention. The maximum dimensional change rate after heat treatment increased.

In Nos. 59 to 61 in Table 7, the steel E in Table 6 of which the steel ingredients satisfy the requirements of the invention was processed at different solution treatment temperature T1 and aging temperature T2, and the samples were analyzed for their physical properties.

Of those, No. 59 and No. 60 are the samples of the invention for which the aging temperature T2 satisfies the range of the invention (TA±10° C.). These both had high hardness and were excellent in the dimensional change resistance after heat treatment.

As opposed to these, No. 61 is a comparative sample for which the aging temperature T2 is higher than the range of the invention. The sample was good in point of the dimensional change rate difference after heat treatment, but the mean dimensional change rate and the maximum dimensional change rate thereof increased.

In Nos. 62 to 64 in Table 7, the steel F in Table 6 of which the steel ingredients satisfy the requirements of the invention was processed at different solution treatment temperature T1 and aging temperature T2, and the samples were analyzed for their physical properties.

Of those, No. 62 and No. 63 are the samples of the invention for which the aging temperature T2 satisfies the range of the invention (TA±10° C.). These both had high hardness and were excellent in the dimensional change resistance after heat treatment.

As opposed to these, No. 64 is a comparative sample for which the aging temperature T2 is higher than the range of the invention. The sample was good in point of the dimensional change rate difference after heat treatment, but the mean dimensional change rate and the maximum dimensional change rate thereof increased.

In Nos. 65 to 67 in Table 7, the steel G in Table 6 of which the steel ingredients satisfy the requirements of the invention was processed at different solution treatment temperature T1 and aging temperature T2, and the samples were analyzed for their physical properties.

Of those, No. 65 and No. 66 are the samples of the invention for which the aging temperature T2 satisfies the range of the invention (TA±10° C.). These both had high hardness and were excellent in the dimensional change resistance after heat treatment.

As opposed to these, No. 67 is a comparative sample for which the aging temperature T2 is higher than the range of the invention. The sample was good in point of the dimensional change rate difference after heat treatment, but the mean dimensional change rate and the maximum dimensional change rate thereof increased.

The following samples (Nos.) satisfy the requirements of the invention for the solution treatment temperature and the aging temperature, but these does not satisfy in point of the steel ingredients. Therefore, these comparative samples have some drawbacks.

No. 68 and No. 69 are both samples of steel H and steel I in Table 6 that are simulated steel samples of conventional high-C high Cr steel. The product of [Cr] and [C] is large, the ratio of [Cu] to [C] is small and the Ms point is low. After heat treatment, the mean dimensional change rate, the maximum dimensional change rate and the dimensional percentage difference all increased. These steels had more increased hardness when the tempering temperature was lower. Therefore, the tempering temperature for these steels was 510° C., and the properties thereof were measured.

No. 70 and No. 71 are both samples of steel J in Table 6 in which the Cu amount is small, the ratio of [Cu]/[Ni] and the ratio of [Cu]/[C] are both small. The hardness was low, and the maximum dimensional change rate increased.

No. 72 and No. 73 are both samples of steel K in Table 6 in which the ratio of [Cu]/[C] is large. The maximum dimensional change rate of the two increased.

In this Example, the time-dependent change in the dimensional change rate is not shown. However, when steel is processed for solution heat treatment followed by aging treatment under the condition satisfying the requirements in the invention, then it is expected that the thus-processed samples could keep high hardness and good dimensional change resistance, and the time-dependent change in the dimensional change rate thereof could be retarded.

The invention has been described in detail with reference to some specific embodiments; however, it is obvious to anyone skilled in the art that the invention can be changed and modified in any desired manner not overstepping the spirit and the scope of the invention.

The present application is based on a Japanese patent application (Japanese Patent Application 2006-283038) filed on Oct. 17, 2006, a Japanese patent application (Japanese Patent Application 2006-294528) filed on Oct. 30, 2006, and a Japanese patent application (Japanese Patent Application 2007-047490) filed on Feb. 27, 2007; and the entire contents thereof are hereby incorporated in this description by reference.

All the references cited herein are incorporated as a whole by reference.

INDUSTRIAL APPLICABILITY

Number	Date	Country	Kind
2006-283038	Oct 2006	JP	national
2006-294528	Oct 2006	JP	national
2007-047490	Feb 2007	JP	national

COLD WORK DIE STEEL, DIE, AND METHOD FOR PRODUCTION OF COLD WORK DIE STEEL

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