Cold working die steel

Info

Publication number: 20060157163
Type: Application
Filed: Jan 17, 2006
Publication Date: Jul 20, 2006
Applicant:
Inventors: Takayuki Shimizu (Nagoya-shi), Yasutaka Ikeuchi (Nagoya-shi), Toshimitsu Fuji (Nagoya-shi)
Application Number: 11/332,910

Abstract

This invention is aimed at providing a cold working die steel successfully reduced in the hardness, and increased in readiness in the cold workability (forging, pressing and so forth) and machinability (milling, drilling, endmilling processing, grinding and lathe turning). A cold working die steel aimed at solving the above-described subject consists essentially of, in % by mass, 0.6%≦C≦1.60%, 0.10%≦Si≦1.20%, 0.10%≦Mn≦0.60%, 5.5%≦Cr≦13.0%, 0.80%≦Mo+0.5W≦2.10%, 0.10%≦V≦0.40%, 0.0002%≦O≦0.0080%, 0.001%≦A1≦0.10%, and the balance of Fe and inevitable impurities; has transformation point Ar3 in the range from 750° C. to 850° C., with both ends inclusive; has a mean circle-equivalent diameter of a carbide, which belongs to a circle-equivalent diameter range from 0.1 μm to 3 μm observed in a section of a structure obtained after spherodizing a sample that was heated at a temperature of (Ar3+50° C.) or above and 1,050° C. or below, of 0.25 μm to 0.8 μm, with both ends inclusive; has a Brinell hardness attained after the spherodizing of HB179 to HB235, with both ends inclusive; has a steel cleanliness of (dB+dC)60×400 in the group C inclusion and the group B inclusion specified by JIS G0555 of 0.05% or less; and has a K value defined as Cr(%)−6.8×C(%) of 0.1 to 3.5, both ends inclusive.

Description

Description

FIELD OF THE INVENTION

This invention relates to a cold working die steel used for cold die and structural components processed by forging or pressing in cold working; mechanical components demanding wear resistance; punch and die for cold forging; mold die for high tensile steel sheets; bending dies; cold forging dies; swaging dies; thread rolling dies; punch components; slitter knives; punch dies for lead frames; gauges; deep-drawing punches; bender punches; shear blades; benders for stainless steel; drawing dies; tools for plastic working such as heading; punches for gears; cam components; press punch dies; progressive punch dies; seal plates for soil conveyers; screw components; rotary plates for concrete spraying machines; IC molding dies; precision press mold demanding high dimensional accuracy; and dies used for the above-described applications after being surface-treated by CVD, PVD or TD.

BACKGROUND ART

SKD11, SKD12 and so forth, which are representative steel types for high-alloy cold dies specified by JIS G4404 have conventionally been used as cold die steel. These cold die steels are subjected to cold working after being subjected to hot working (forging, rolling), and then to annealing, in which the steel is heated to a temperature range from point Ar3 to Ar3+50° C. and then slowly cooled. The hardness obtained after annealing under such annealing conditions falls in the range approximately from HB241 to HB255 on the Brinell hardness basis.

Japanese Laid-Open Patent Publication “Tokkaihei” No. 6-322439 discloses that the hardness can be lowered as compared with that attained by conventional annealing, by raising the annealing temperature higher than in the conventional process. Although the document describes that the lowering of the hardness can increase the machinability, it is hardly said that the hardness of HV262 to 278 (approximately 250 to 265 on the HB basis) listed in Table 2 is small enough. Moreover, no description is made on the hardness after quench-and-temper, despite a known problem in that annealing carried out at a temperature higher than hardening temperature results in lowering in the hardness after quench and temper than usual.

In recent years, shorter delivery periods and lower costs in the fabrication process have been more strongly demanded, and there have been increased demands on cold working die steel more easily used in machinability such as post-anneal machinability after annealing or cold workability, in view of achieving near-net-shaping and shortening of the time period for fabricating the structural and mechanical components. With cold tool steel, however, it is hard to achieve a low level of hardness by annealing intrinsically because of the composition thereof.

Steel having large amounts of coarse carbide are used for applications in particular demanding wear resistance such as dies, tools and mechanical components, in view of the wear resistance, but making too much account of wear resistance inevitably results in degradation of the material characteristics, such as a lowering of the overall toughness. On the other hand, a decrease in the coarse carbide or an addition of a large amount of element enhancing free-cutting properties, such as S, aiming at improving the machinability, lowers the wear resistance required for tool steel. It is to be noted herein that the “coarse carbide (primary carbide)” refers to a carbide having a diameter of approximately 10 μm or larger on a circle-equivalent basis.

The present invention was conceived taking the above-described problems into account, and an object thereof resides in providing a cold working die steel which can be reduced in its hardness by annealing, and is improved in readiness in the cold workability (forging, pressing and so forth) and machinability (milling, drilling, endmilling processing, grinding and turning).

SUMMARY OF THE INVENTION

Aiming at solving the aforementioned problems, cold working die steel according to the first aspect of this invention consists essentially of, in % by mass and by both ends inclusive, 0.6%≦C≦1.60%, 0.10%≦Si≦1.20%, 0.10≦Mn≦0.60%, 5.5%≦Cr≦13.0%, 0.80%≦Mo+0.5W≦2.10%, 0.10%≦V≦0.40%, 0.0002%≦O≦0.0080%, 0.001%≦Al≦0.10%, and the balance of Fe and inevitable impurities; having transformation point Ar3 in the range from 750° C. to 850° C., with both ends inclusive; having a mean circle-equivalent diameter of a carbide, which belongs to a circle-equivalent diameter range from 0.1 μm to 3 μm observed in a section of a structure obtained after spherodizing a sample heated at a temperature of (Ar3+50° C.) or above and 1,050° C. or below, of 0.25 μm to 0.8 μm, with both ends inclusive; having a Brinell hardness attained after the spherodizing of HB179 to HB235, with both ends inclusive; having a steel cleanliness level of (dB+dC)60×400 in group C inclusion and group B inclusion specified by JIS G0555 of 0.05% or less; and having a K value defined as Cr(%)−6.8×C(%) of 0.1 to 3.5, with both ends inclusive.

This invention has major features described below.

(1) Lowering of the Hardness After Annealing

Hardness after annealing is determined mainly by state of distribution of carbide having a diameter ranging from 0.1 μm to 3 μm, with both ends inclusive, on the circle-equivalent basis (referred to as secondary carbide, hereinafter), so that the hardness can be decreased by adjusting the mean circle-equivalent diameter of the carbide to the range from 0.2 μm to 0.8 μm, with both ends inclusive. This range of the mean circle-equivalent diameter corresponds to the enlargement of the secondary carbide. In other words, an absolute amount of the secondary carbide produced by the annealing is the same, so that a larger mean circle-equivalent diameter results in a smaller hardness due to scarceness of the secondary carbide having relatively large sizes.

On the contrary, a smaller mean circle-equivalent diameter results in a larger hardness by virtue of the denseness of the secondary carbide having relatively small sizes. More specifically, adjustment of the mean circle-equivalent diameter of the secondary carbide in the above-described range makes it possible to obtain a Brinell hardness of as low as HB179 to HB235, with both ends inclusive (conventional steels showed a Brinell hardness of as large as HB241 to HB255 or around), and to distinctively improve the efficiency in machining and cold working.

(2) Control of the Size of the Secondary Carbide by Annealing Conditions

Size of the secondary carbide is controlled by spherodizing, or keeping the steel heated at a temperature of (Ar3+50° C.) or above and 1,050° C. or below. In more detail, the steel is heated at a temperature (from Ar3+50° C. to a hardening temperature (1,050° C. or around)) higher than the conventional annealing temperature (from Ar3 to Ar3+50° C.) so as to allow more secondary carbides to dissolve into the matrix. The reason why the temperature at which the steel is heated is adjusted to a temperature lower than the hardening temperature (1,050° C. or around) is to avoid lowering of the hardness after quench-and-temper.

Thereafter, the steel is gradually cooled at a cooling rate slower than 60° C./h to as low as 750° C. or below, in order to allow the dissolved secondary carbide to grow into larger grains (gradual cooling method). Other known treatment methods include the “repetitive heating-and-gradual-cooling method” in which heating and cooling are repeated at least twice within temperature ranges from 650° C. to transformation point Ar1, and from transformation point Ar3 to 1,050° C.; “prolonged annealing method” in which the steel is kept at a temperature lower than Ar3 for a long period of time; and “isothermal transformation method” in which the steel is kept at a constant temperature during the gradual cooling process in the gradual cooling method.

Any of these methods can successfully control the size of the secondary carbide. Low temperature annealing at a temperature of Ar1 or below before the spherodizing annealing makes it possible to reduce variation in the post-anneal hardness after annealing and to obtain still an even lower hardness. It is to be noted that the above-described spherodized annealing will not largely alter the size and population of the coarse carbide (primary carbide) which affects the wear resistance required for applications such as dies and tools.

Transformation point “Ar3” herein expresses the A3 transformation point (austenization temperature), wherein “r” means cooling (refroidissement). In the spherodized annealing, the temperature at which the steel is heated is necessarily set higher than that in the conventional process in view of thoroughly dissolving the secondary carbides into the matrix, and the transformation point of Ar3 is necessarily low enough so as to adjust the temperature to the quenching temperature or below. In other words, it is required that there is a large enough difference between the transformation point of Ar3 and the quenching temperature.

The steel is therefore necessarily adjusted in the composition thereof so as to make transformation point Ar3 fall within the range of 750° C. to 850° C. If the Ar3 is too high only a small difference from the normal quenching temperature (1,050° C. or around) will result, and consequently fail in raising the annealing temperature thereby resulting in only an insufficient dissolution of the secondary carbide (not so different from the conventional annealing). On the other hand, if the Ar3 is too low, a vastly longer duration time for the depositing and growing of the secondary carbide. The annealing temperature herein is determined to a desirable level (Ar3+50° C. to quenching temperature (at around 1,050° C.)), based on results of the measurements of the transformation point Ar3 using an apparatus such as DTA (differential thermal analyzer).

(3) Influence of Oxide-Base Inclusion on Machinability

Increases in groups B and C (oxide-based ones in particular) inclusions (conforming to JIS G0555) are known to decrease the machinability. These inclusions have extremely high hardness of their own, which exceeds the hardness of the matrix. By coming into contact with these inclusions, tool edges can be chipped and the lifetime of tools can be considerably decreased. Both group B and C inclusions become less affective as the contents of which reduce closer to 0, so that it is necessary, for the purpose of obtaining more sufficient machinability than conventional steel, to adjust the steel cleanliness to (dB+dC)60×400≦0.05%. This cleanliness level can be obtained by adjusting mainly the contents of O and Al within the range described later.

The following paragraphs will describe reasons for limitations of the individual numerical ranges.

(4) 0.6%≦C≦1.60%

C is an essential element for raising the hardness of martensite after quenching. The element forms carbide through binding with carbide-forming elements such as Cr, Mo and V, and thereby make the crystal grain more fine. The carbide also contributes to the improvement of the wear resistance. An addition to an amount of the lower limit or more is necessary in order to realize a hardness after quench and temper of HRC55 or above. On the other hand, the element is added to as much as the upper limit or less, because excessive addition results in excessive content of the carbide, which thereby decreases the toughness.

(5) 0.10%≦Si≦1.20%

Si is added as a deoxidizing element. Because the element contributes to the increase in hardness in the high-temperature-temper, it is added to as much as the lower limit or more, so as to obtain the effect. On the other hand, the element is added to as much as the upper limit or less, because excessive addition degrades the hot workability, and the after-quench-and-temper toughness.

(6) 0.10%≦Mn≦0.60%

Mn is added as a deoxidizing element. Because the element contributes to the enhancement of the hardenability, and increasing in the hardness and strength, it is added to as much as 0.10% or more. On the other hand, the upper limit of the element is set to 0.60%, because excessive addition degrades the hot workability.

(7) 5.5%≦Cr≦13.0%

Cr dissolves into the matrix which raises the hardenability and increases the hardness, and forms carbides which increase the wear resistance. Additions to as much as the lower limit or more are necessary for obtaining these effects. On the other hand, the element is added to as much as the upper limit or less, because excessive addition results in formation of an excessive amount of carbides which decrease the toughness and machinability after quench-and-temper.

(8) 0.80%≦Mo+0.5W≦2.10%

Mo and W dissolve into the solid matrix to thereby raise the hardenability and to contribute to increase hardness, and forms carbides to which increase the wear resistance. The elements also have an effect of raising the anti-softening hardness in quench-and-temper. Addition to as much as the lower limit or more, on the Mo-equivalent basis expressed as Mo(%)+0.5W(%), is necessary for obtaining these effects. On the other hand, the element is added to as much as the upper limit or less, because excessive addition decreases the hot workability, toughness and machinability.

(9) 0.10%≦V≦0.40%

V forms a stable carbide, and thereby effectively prevents the crystal grains from coarsening. The element forms a fine carbide, and thereby contributes to increases in the wear resistance and the hardness. Addition to as much as 0.10% or more is necessary for obtaining these effects. On the other hand, the upper limit is set to 0.40%, because excessive addition increases the carbide content, and thereby decreases the machinability and hot workability.

(10) 0.0002%≦O≦0.0080%, 0.001%≦Al≦0.10%

O and Al are inevitably contained in the steel. The elements are constituent elements of group B and C inclusions, and large contents of which can decrease the toughness, so that it is necessary to suppress the contents to as much as the upper limit or less. Positive efforts of reducing these elements, although depending on balance with the production cost, makes it possible to maintain a high, stable toughness. Excessive lowering of the contents only results in an increase in the production cost and a saturation of influences exerted on the toughness, so that the elements are added to as much as the lower limit or more.

(11) Transformation Point Ar3 Adjusted to 750° C. to 850° C.

Transformation point Ar3 adjusted in the above-described range can expand the temperature range ranging from Ar3+50° C., at which spherodizing is effected. The quenching temperature (about 1,050° C.), ensures a sufficient range of annealing temperature allowing the secondary carbides to thoroughly dissolve. An excessively high Ar3, however, narrows the temperature range from Ar3+50° C. to the quenching temperature (about 1,050° C.), and fails to ensure a sufficient range of an annealing temperature. This, as a consequence, extremely reduces the amount of dissolution of the secondary carbides, and fails in obtaining low hardness required for raising machinability. An excessively low Ar3 takes a longer time to allow the secondary carbides to deposit and grow in the process of gradual cooling at a temperature as low as Ar3 or below, and this extremely raises the cost from the industrial viewpoint. It is to be understood that Ar3 can be measured typically by DTA, and defined as being obtained at a cooling rate of 5° C./h or more to 60° C./h, because Ar3 varies depending on conditions of the measurement.

(12) The mean circle-equivalent diameter of the carbide, which belongs to the circle-equivalent diameter range of 0.1 μm to 3 μm was observed in the section of the structure obtained after spherodizing a sample that was heated at a temperature of (Ar3+50° C.) or above and 1,050° C. or below, adjusted to 0.25 μm or more and 0.8 μm or less.

The mean circle-equivalent diameter of the carbide is calculated by an image analysis of a polished section of the steel structure. More specifically, a circle-equivalent diameter is calculated for every carbide grain having a circle-equivalent diameter ranging from 0.1 μm to 3 μm which can be seen in a magnified field of view under a scanning electron microscope or an optical microscope, and a mean value of which is determined as the mean circle-equivalent diameter. It is necessary that the observation in the magnified field of view is carried out at least in an area of 1 mm²or larger, at randomly selected positions excluding the surface and center portions of the material. Alternatively, it is also allowable to use 20 or more fields of view, when each field contains 20 to 50 carbide grains having a circle-equivalent diameter ranging from 0.1 μm to 3 μm. The carbide having a diameter ranging from 0.1 μm to 3 μm means a carbide (secondary carbide) contributive to the hardness.

A mean grain size of the carbide belonging this range of less than 0.25 μm results in a high hardness, and fails in obtaining the effect of increasing the machinability (this applies to the case where the conventional style of annealing was carried out). On the other hand, an excessively large mean grain size results in an extremely small number of carbide grains grown during the process of gradual cooling in the annealing, and this makes the regenerative perlite more likely to deposit in the cooling process, and conversely increases the hardness. It is therefore necessary to set the upper limit of the mean circle-equivalent diameter of carbide to 0.8 μm.

(13) Brinell Hardness Attained After Spherodizing of HB179 to HB235

Adjustment of the Brinell hardness after the spherodizing to HB179 to HB235 makes it possible to obtain the machinability and cold workability superior to those in the prior art. Excessive lowering of the hardness results in a high cost from the industrial point of view, so that a hardness of HB179 or more is enough for the purpose.

(14) Steel Cleanliness in Group B and C Inclusions Specified by JIS G0555 of (dB+dC)60×400≦0.05%

The class B inclusion (mainly alumina and so forth) refers to grains forming an agglomerate discontinuously arrayed in the working direction, and the group C inclusion (such as granular oxide) refers to grains irregularly dispersed without causing viscous deformation. Excessive amounts of these inclusions decreases the machinability, and a desirable level of machinability can be obtained by adjusting the cleanliness (dB+dC)60×400 of the steel with respect to group B and C inclusions, determined by the test method as specified in JIS G0555, of 0.05% or less.

(15) K Value Expressed by Cr(mass %)−6.8×C(mass %) of 0.1 or More and 3.5 or Less

K value expresses the amount of Cr solubilized in the matrix at an appropriate quenching temperature. Adjustment within the above range results in the hardness obtained after quench-and-temper almost equivalent to the previous, and therefore makes it possible to adjust the amount of deposition of carbides depending on the wear resistance, toughness and machinability required for the cold working die steel. In contrast, if the K value is out of the above range, it results in only an insufficient amount of the secondary carbide which deposits during tempering and contributes to the hardness, and consequently fails in maintaining the hardness necessary for the cold working die steel. FIG. 1 shows a relation between the C content and the Cr content. The straight lines ascending from the lower left towards the upper right in the drawing relate to the K value, and the straight lines descending from the upper left towards the lower right in the drawing relate to the L value described later.

The cold working die steel of this invention can further contain, as steel components, either one of or both of 0.0030%≦N≦0.0500% and 0.001%≦P≦0.040%.

These elements are inevitably contained in the steel. Large contents of these elements decrease the toughness, so it is preferable to control them to as much as the upper limit or below. A positive effort of reducing these elements, in a trade-off manner with the manufacturing cost, makes it possible to maintain a stable and high toughness. It is to be understood that excessive lowering in the contents only results in an increase in cost while allowing an effect on the toughness to saturate, so that it is preferable to control them to the lower limit or above.

The cold working die steel of this invention can further contain, as steel components, any one of, or two or more of steel components selected from 0.01%≦Cu≦1.0%, 0.01%≦Ni≦1.0%, 0.2%≦Co≦1.0% and 0.0003%≦B≦0.010%.

These elements have an effect of increasing the hardenability by dissolving themselves into the matrix. They have also an effect of increasing the toughness by lowering the impact transition temperature, and by consequently preventing the weldability from degrading. Co has an effect of increasing the high temperature strength. Addition to as much as the lower limit or above is preferable in view of obtaining these effects. The elements are added to as much as the upper limit or less, because excessive addition only results in a saturated effect.

The cold working die steel of this invention can further contain, as steel components, any one of, or two or more of 0.001%≦S≦0.20%, 0.005%≦Se≦0.10%, 0.005%≦Te≦0.10%, 0.0002%≦Ca≦0.010%, 0.005%≦Pb≦0.10% and 0.005%≦Bi≦0.10%.

S can be added as an element increasing the free cutting property. Depending on the contents of the carbide-forming elements such as Cr, Mo and V, addition to as much as 0.04% or more is preferable in view of obtaining the effect of increasing the free cutting property. Excessive addition of the element considerably decreases the toughness, or mechanical properties including surface roughness after discharge processing and cutting, so that the upper limit is preferably adjusted to 0.20%. For applications making a great account of machinability, the element is added considering the balance with the mechanical properties. On the other hand, for applications in which a greater account is placed on the mechanical properties rather than on the machinability, the amount of addition of S is set to 0.02% or less, and more preferably to 0.01% or less, considering the balance with the cost for manufacturing. The mechanical properties can be satisfied by adjusting the amount of the addition to 0.01% or more and 0.02% or less for the practical operation. This makes it possible to obtain the S content described in the above.

Also Se, Te, Ca, Pb and Bi can be added for the purpose of increasing machinability. Se and Te can be used as substitutive elements of S in Mn sulfide. Ca improves the machinability by forming a protective film on the surface of a tool during cutting, by forming an oxide or by dissolving itself into Mn sulfide. Pb and Bi segregate in the grain boundary, to thereby lower the grain boundary strength and to improve the machinability. The elements are necessarily added to as much as the lower limits or more in view of obtaining these effects. On the other hand, excessive addition results in degraded mechanical properties, so that the upper limits should be met.

The cold working die steel of this invention can further contain, as the steel components, any one of, or two or more of 0.01%≦Nb≦0.12%, 0.005%≦Ta≦0.10%, 0.005%≦Ti≦0.10%, 0.005%≦Zr≦0.10%, 0.005%≦Mg≦0.10% and 0.005%≦REM≦0.10%.

These elements can be added for the purpose of obtaining an effect of increasing the toughness through fineness of the carbide grains and fineness of the crystal grains. Mg and REM have an effect of increasing the toughness and machinability, through formation of oxides, which contribute to the reduction of the O content and consequently in coarse oxide grains. Addition to as much as the lower limit or above is preferable in view of obtaining these effects. On the other hand, excessive addition results in decreasing in the toughness and weldability, so that the amount of addition is preferably set to the upper limit or less. A single species or two or more species of rare earth metals may be used as REM.

Next, the cold working die steel according to the second aspect may contain, as the steel components, 0.60%≦C≦0.80%, 0.10%≦Si≦1.20%, 0.10%≦Mn≦0.60%, 5.5%≦Cr≦8.5%, 0.80%≦Mo+0.5W≦2.10%, 0.10%≦V≦0.40%, 0.0002%≦Al≦0.0080%. In other words, further limitations are added to C, Si, Cr and Mo, out of the steel components according to the first aspect.

It is essential for cold working die steel die applications, that the toughness and fine machining are particularly required to decrease the coarse carbide. More specifically, it is necessary to avoid as much as possible, the formation of coarse carbides mainly composed of M₇C₃(where, M represents Cr, Mo or V), by adjusting the contents of C, Si, Cr and Mo within the above-described ranges. The amount of coarse carbide corresponds to 0.01 to 5% in % by mass. Assuming now L value as an index expressing the amount of coarse carbide as Cr(mass %)+15.5×C(mass %), the amount of coarse carbide corresponds to 14.9≦(L value)≦21.0 (see FIG. 1).

Next, the cold working die steel according to the third aspect may contain, as the steel components, 0.90%≦C≦1.10%, 0.8%≦Si≦1.20%, 0.10%≦Mn≦0.60%, 7.0%≦Cr≦9.0%, 1.50%≦Mo+0.5W≦2.10%, 0.10%≦V≦0.40%, 0.0002%≦Al≦0.0080%. In other words, further limitations are added to C, Si, Cr and Mo, out of the steel components according to the first aspect.

It is necessary for cold working die steel die applications, that the wear resistance and the toughness must be well-balanced, to ensure a certain amount of coarse carbides. More specifically, the coarse carbides mainly composed of M₇C₃, are formed by adjusting the contents of C, Si, Cr and Mo within the above-described ranges. The amount of coarse carbides corresponds from 5 to 10% in % by mass, and to 21.0≦(L value)≦27.0 (see FIG. 1).

Next, the cold working die steel according to the fourth aspect may contain, as the steel components, 1.40%≦C≦1.60%, 0.10%≦Si≦0.40%, 0.10%≦≦Mn≦0.60%, 11.0%≦Cr≦13.0%, 0.80%≦Mo+0.5W≦1.20%, 0.10%≦V≦0.40%, 0.0002%≦Al≦0.0080% In other words, further limitations are added to C, Si, Cr and Mo, out of the steel components according to the first aspect.

It is necessary for cold working die steel die applications, where the wear resistance is particularly required, to contain a large amount of coarse carbide. More specifically, large amounts of coarse carbides mainly composed of M₇C₃are formed by adjusting the contents of C, Si, Cr and Mo within the above-described ranges. The amount of coarse carbide corresponds to 10 to 15% in % by mass, and to 27.0≦(L value)≦37.8 (see FIG. 1).

BRIEF DESCRIPTION OF THE DRAWING

FIG. 1 is a drawing expressing relations between C content and Cr content (K value, L value).

BEST MODES FOR CARRYING OUT THE INVENTION

First, 200 kg of each of the inventive and comparative steel materials having compositions listed in Table 1 were melted in a vacuum induction furnace, made into an ingot, and the obtained steel ingot was hot-forged so as to produce a 70 mm×70 mm square rod. The rod was then annealed at the temperature listed in Table 2 (cooling rate: 18° C./h).

Of the compositions of comparative steels listed in Table 1, those departing from the compositional ranges specified by this invention are indicated by a downward arrow (↓) if they came short of the lower limits, by an upward arrow (↑) if they exceeded the upper limits.

[Table 1]

TABLE 1 Components (wt %) No. C Si Mn P S Cu Ni Cr Mo W 1 Comparative steel 0.71 0.38 ↑0.93 0.028 ↑1.32 0.11 7.03 0.84 2 Inventive steel 1.23 0.53 0.34 0.06 0.43 10.32 1.44 3 Comparative steel 0.71 0.38 0.37 0.028 0.002 0.09 ↑1.55 7.03 0.84 4 Comparative steel 0.66 0.42 0.32 0.013 0.001 0.23 0.53 10.5 0.32 5 Inventive steel 0.61 0.95 0.47 0.004 0.063 0.05 0.13 5.7 2.01 6 Inventive steel 0.78 0.75 0.21 0.035 0.035 0.85 0.53 5.81 2.09 7 Inventive steel 0.93 0.23 0.15 0.021 0.001 0.34 0.43 6.53 0.82 8 Inventive steel 0.79 0.35 0.32 0.003 0.23 0.03 8.82 0.93 9 Inventive steel 0.62 0.11 0.11 0.028 0.058 0.53 0.84 7.53 1.56 10 Inventive steel 0.72 0.69 0.55 0.013 0.008 0.13 0.08 7.03 1.03 11 Inventive steel 0.78 1.17 0.59 0.018 0.17 0.25 0.31 6.53 0.81 0.02 12 Inventive steel 0.71 1.18 0.58 0.022 0.008 0.023 0.89 8.2 0.88 0.83 13 Inventive steel 0.83 0.11 0.42 0.006 0.004 0.14 0.67 5.96 1.65 14 Comparative steel 0.93 0.89 0.31 0.022 ↑0.25 0.08 0.33 ↑13.94 1.83 15 Comparative steel 0.99 0.89 0.52 0.012 0.003 0.35 0.46 6.34 1.93 16 Inventive steel 0.94 0.93 0.41 0.011 0.001 0.02 0.04 6.63 1.84 17 Inventive steel 1.19 0.83 0.22 0.001 0.14 0.75 0.23 8.32 1.74 0.11 18 Inventive steel 1.07 0.85 0.39 0.013 0.003 0.05 0.13 10.71 1.63 19 Inventive steel 0.81 0.93 0.18 0.032 0.002 0.13 0.45 8.83 1.99 0.03 20 Inventive steel 0.92 1.05 0.59 0.021 0.095 0.34 0.63 8.21 2.03 21 Inventive steel 0.88 1.17 0.11 0.009 0.103 0.23 0.09 7.58 1.54 22 Inventive steel 0.99 1.08 0.58 0.013 0.003 0.02 0.75 9.03 2.08 23 Inventive steel 0.99 0.42 0.11 0.023 0.38 0.08 0.09 8.25 0.97 24 Inventive steel 0.86 0.93 0.11 0.025 0.001 0.011 0.98 9.32 1.55 25 Comparative steel ↑1.83 0.32 0.53 0.001 0.09 0.31 11.95 ↓0.43 26 Comparative steel 1.53 0.13 0.43 0.011 0.001 0.21 0.33 9.31 0.88 0.31 27 Inventive steel 1.23 0.29 0.46 0.018 0.001 0.04 0.08 8.53 0.83 0.22 28 Inventive steel 1.58 0.11 0.23 0.005 0.11 0.32 0.32 11.44 0.95 29 Inventive steel 1.56 0.18 0.12 0.029 0.002 0.03 0.98 12.98 0.93 0.31 30 Inventive steel 1.44 0.21 0.18 0.003 0.08 0.09 0.75 12.84 1.18 31 Inventive steel 1.08 0.31 0.22 0.002 0.002 0.21 0.63 10.73 1.09 32 Inventive steel 1.43 0.37 0.34 0.009 0.083 0.85 0.15 12.59 1.05 0.09 33 Inventive steel 1.23 0.23 0.57 0.028 0.052 0.43 0.44 10.39 1.02 34 Inventive steel 1.45 0.97 0.36 0.038 0.034 0.88 0.1 12.78 0.76 1.33 35 Inventive steel 1.3 0.29 0.43 0.024 0.27 0.91 0.06 12.01 1.85 No. V Al O N K value L value Others 1 0.35 0.028 0.0009 2.202 18.035 2 0.15 0.093 0.0053 1.956 29.385 3 ↑1.34 0.028 0.0009 0.019 2.202 18.035 4 0.11 0.033 0.0021 0.011 ↑6.012 20.73 5 0.21 0.008 0.0023 0.015 1.552 15.155 6 0.39 0.03 0.0053 0.004 0.506 17.9 7 0.11 0.09 0.0003 0.0033 0.206 20.945 8 0.12 0.053 0.0009 0.0085 3.448 21.065 9 0.11 0.08 0.0004 0.009 3.314 17.14 10 0.14 0.032 0.0063 0.013 2.134 18.19 11 0.33 0.021 0.0078 0.024 1.226 18.62 12 0.13 0.012 0.0013 0.012 3.372 19.205 B = 0.0032 13 0.22 0.096 0.0043 0.048 0.316 18.825 Zr = 0.022, Ta = 0.07, Mg = 0.09 14 0.22 0.053 0.0031 0.011 ↑7.616 28.355 15 0.32 0.023 0.0039 0.016 ↓−0.392 21.685 16 0.31 0.011 0.0019 0.009 0.238 21.2 17 0.21 0.022 0.0043 0.0395 0.228 26.765 18 0.25 0.034 0.0041 0.0312 3.434 27.295 19 0.33 0.048 0.0077 0.0213 3.322 21.385 20 0.29 0.028 0.0063 0.0183 1.954 22.47 21 0.19 0.035 0.0031 0.018 1.596 21.22 22 0.37 0.059 0.0019 0.031 2.298 24.375 23 0.18 0.001 0.0002 0.0035 1.518 23.595 Mo = 0.02, Ti = 0.043, REM = 0.05 24 0.39 0.032 0.0079 0.01 3.472 22.65 Ca = 0.0053, Te = 0.026, Bi = 0.024 25 0.34 0.099 0.0009 ↓−0.494 40.315 26 0.25 0.008 0.0019 0.011 ↓−1.094 33.025 27 0.27 0.005 0.0023 0.014 0.166 27.595 28 0.16 0.008 0.0003 0.0185 0.696 35.93 29 0.33 0.005 0.0053 0.0045 2.372 37.16 30 0.23 0.093 0.0005 0.0093 3.048 35.16 31 0.38 0.053 0.0078 0.023 3.386 27.47 32 0.11 0.083 0.0015 0.0075 2.866 34.755 33 0.32 0.053 0.0034 0.0053 2.026 29.455 34 0.25 0.012 0.0035 0.011 2.92 35.255 Co = 0.36 35 0.33 0.004 0.0031 0.013 3.17 32.16 Pb = 0.09, Se = 0.09

The inventive steels and comparative steels were subjected to the tests and evaluations below. Results are shown in Table 2.

(a) Surface Analysis and Evaluation of Carbide

The polished surface of each steel was subjected to image analysis and the mean grain size of the carbide was measured. The image analysis was made on an image observed under a SEM, wherein a total of 1 mm²area was observed at an appropriate magnification ranging from 500× to 5,000×. The circle-equivalent diameters were calculated for all carbide grains having a diameter ranging from 0.1 μm to 3.0 μm seen in the field of view, and thereby obtained an average mean value. The polished surface herein was etched with a picric acid-ethanol solution to a depth allowing observation of the carbide grains having a diameter of about 0.1 μm, without causing dropping of the carbide grains.

(b) Amounts of Group B and C Inclusions

According to the test method specified by JIS G0555, (dB+dC)60×400 was measured.

(c) Machinability Test

A test piece was cut out from each of the manufactured inventive steels and comparative steels, and subjected to the machinability test.

Conditions for Endmill Processing Test

Tool: carbide solid endmill (φ=10 mm), six-flute

Cutting speed: 120 m/min

Feed: 0.06 mm/rev

Cutting width: 0.5 mm, 10 mm in height

Lubricant: air blow

Cutting length: up to 60,000 mm

Judgment: marked with ∘ if the tool caused no breakage, and marked with x if the tool caused breakage or sparking during the cutting.

(c) Cold Workability

Test pieces measuring 12×18 mm were cut out from each of the inventive steels and the comparative steels, and pressed by a single stroke to 60% height of the test piece using a 600-t hydraulic press machine. Ten pressed test pieces of the individual steels were observed, and the number of test pieces causing breakage was found.

(e) Maximum Hardness

Hardness was measured using a Rockwell C scale, under varied annealing conditions in the quench-and-temper.

(f) Specific Wear

Specific wear was measured using a pin-on-disk friction and wear tester. Two 8-mm-diameter pins were cut out from the individual manufactured inventive steels and the comparative steels. A disk was cut out from S45C. The inventive steels and the comparative steels were subjected to quench-and-temper under which the maximum hardness can be attained. Test conditions included a slipping rate of 1.6 m/s, and a press load of 10.5 kgf, with no lubricant. The weight of the pins was measured before and after the test, and loss of weight by wear was measured. The specific wear of the inventive steels and the comparative steels was evaluated, assuming that the weight loss by the wear of comparative steel was 1 as one.

(g) Charpy Impact Value

A 10R-notched Charpy test piece was cut out from each of the manufactured inventive steels and the comparative steels. The direction of the test pieces was aligned to the longitudinal direction of the material. Under annealing conditions yielding the maximum hardness, the test was carried out according to the method described in JIS Z2242. The test was carried out under room temperature.

[Table 2]

TABLE 2 Mean circle- Charpy Ar3 equivalent diameter Annealing After-SA Maximum Cold impact temperature of the carbide Inclusion temperature hardness hardness work- Machin- Specific value No. (° C.) (μm) dB + dC (° C.) (HB) (HRC) ability ability wear (J/cm2) 1 Comparative steel ↓ 740 ↓ 0.21 0.005 ↓ 750 ↑ 269 60.3 x8 x 1 53 2 Inventive steel 793 0.71 0.002 980 204 63.2 ∘ ∘ 0.32 21 3 Comparative steel ↓ 733 ↑ 1.53 ↑ 0.41 ↑ 1140 ↓ 170 60.4 ∘ x 1.03 46 4 Comparative steel 831 0.53 0.003 950 204 54.3 ∘ ∘ 5.93 53 5 Inventive steel 805 0.45 0.001 970 212 61.3 ∘ ∘ 0.92 63 6 Inventive steel 822 0.58 0.034 1020 208 63.3 ∘ ∘ 0.93 68 7 Inventive steel 803 0.38 0.027 960 198 63.5 ∘ ∘ 0.89 73 8 Inventive steel 755 0.35 0.046 960 199 62.9 ∘ ∘ 0.93 58 9 Inventive steel 790 0.26 0.003 970 203 64.2 ∘ ∘ 0.89 66 10 Inventive steel 809 0.44 0.002 970 203 63.8 ∘ ∘ 0.88 67 11 Inventive steel 833 0.58 0.011 990 185 62.87 ∘ ∘ 0.84 93 12 Inventive steel 813 0.65 0.001 1030 204 62.7 ∘ ∘ 0.93 81 13 Inventive steel 836 0.28 0.006 1030 233 62.9 ∘ ∘ 0.99 88 14 Comparative steel ↑ 876 0.39 ↑ 0.39 970 226 51.1 ∘ x 6.33 67 15 Comparative steel 843 0.67 0.008 950 193 53.2 ∘ ∘ 6.29 93 16 Inventive steel 848 0.52 0.003 990 198 62.5 ∘ ∘ 0.63 43 17 Inventive steel 845 0.78 0.048 990 183 64.3 ∘ ∘ 0.55 44 18 Inventive steel 833 0.47 0.036 990 195 64.1 ∘ ∘ 0.59 48 19 Inventive steel 828 0.36 0.021 930 199 64.8 ∘ ∘ 0.63 53 20 Inventive steel 773 0.31 0.012 840 225 63.8 ∘ ∘ 0.73 41 21 Inventive steel 819 0.63 0.001 930 203 63.7 ∘ ∘ 0.72 39 22 Inventive steel 835 0.65 0.003 980 211 63.2 ∘ ∘ 0.72 46 23 Inventive steel 832 0.71 0.007 890 191 62.8 ∘ ∘ 0.73 52 24 Inventive steel 825 0.68 0.001 930 200 62.7 ∘ ∘ 0.72 37 25 Comparative steel ↑ 883 ↓ 0.11 0.002 ↓ 900 ↑ 266 48.7 x10 x 3.14 11 26 Comparative steel 817 0.73 0.004 980 213 51.3 ∘ ∘ 3.82 13 27 Inventive steel 795 0.71 0.002 890 200 62.7 ∘ ∘ 0.32 19 28 Inventive steel 808 0.68 0.043 1000 201 63.3 ∘ ∘ 0.31 23 29 Inventive steel 813 0.57 0.007 1000 205 63.2 ∘ ∘ 0.46 22 30 Inventive steel 803 0.55 0.011 940 209 64.7 ∘ ∘ 0.49 28 31 Inventive steel 755 0.49 0.028 920 205 62.1 ∘ ∘ 0.28 27 32 Inventive steel 761 0.68 0.032 990 217 61.8 ∘ ∘ 0.38 17 33 Inventive steel 783 0.73 0.003 910 218 63.3 ∘ ∘ 0.37 19 34 Inventive steel 796 0.41 0.006 870 231 63.6 ∘ ∘ 0.31 20 35 Inventive steel 807 0.29 0.007 860 223 64.4 ∘ ∘ 0.45 29

As is known from Table 2, comparative steel 1, having a composition departing from the compositional ranges specified by this invention, showed an extremely lowered Ar3 temperature. The steel annealed under conventional annealing conditions failed to fully solubilize the carbide into solid, and failed in allowing the carbide to grow larger under gradual cooling, so that the carbide became smaller in size, and the steel became harder. The steel was consequently poor in the cold workability, showing cracks in 8 out of ten test pieces.

Comparative steel 3, having a composition departing from the compositional ranges specified by this invention, showed an extremely lowered Ar3 temperature. Too high a temperature in the quench-and-temper resulted in an extremely lowered hardness, increased ductility of the material, and conversely degraded machinability, although the cold workability was judged as desirable by virtue of a large carbide size and a considerably lowered hardness.

Comparative steel 4 showed a K value largely departing from the inventive range, despite having a composition within the compositional ranges specified by this invention. The steel therefore failed in achieving a maximum hardness of as large as HRC60 or above after annealing by the quench-and-temper, and failed in obtaining a level of hardness required for cold working die steel. The low hardness also resulted in a large specific wear. Comparative steels 14, 15, 25 and 26, again having K values departing from the inventive range, also resulted in small maximum hardness and larger specific wear.

Claims

1. A cold working die steel consisting essentially of, in % by mass, 0.6%≦C≦1.60%, 0.10%≦Si≦1.20%, 0.10%≦Mn≦0.60%, 5.5%≦Cr≦13.0%, 0.80%≦Mo+0.5W≦2.10%, 0.10%≦V≦0.40%, 0.0002%≦O≦0.0080%, 0.001%≦Al≦0.10%, and the balance of Fe and inevitable impurities;

having transformation point Ar3 in the range of from 750° C. to 850° C., both ends inclusive;

having a mean circle-equivalent diameter of a carbide, which belongs to a circle-equivalent diameter range from 0.1 μm to 3 μm observed in a section of a structure obtained after spherodizing a sample that was heated at a temperature of (Ar3+50° C.) or above and 1,050° C. or below, of 0.25 μm to 0.8 μm, with both ends inclusive;

having a Brinell hardness attained after the spherodizing of HB179 to HB235, with both ends inclusive;

having a steel cleanliness of (dB+dC)60×400 in the group C inclusion and the group B inclusion specified by JIS G0555 of 0.05% or less; and

having a K value defined as Cr(mass %)−6.8×C(mass %) of 0.1 to 3.5, with both ends inclusive.

2. The cold working die steel as claimed in claim 1, wherein the steel components are 0.60%≦C≦0.80%, 0.10%≦Si≦1.20%, 0.10%≦Mn≦0.60%, 5.5%≦Cr≦8.5%, 0.80%≦Mo+0.5W≦2.10%, 0.10%≦V≦0.40%, 0.0002%≦O≦0.0080% and 0.001%≦Al≦0.10%.

3. The cold working die steel as claimed in claim 1, wherein the steel components are 0.90%≦C≦1.10%, 0.8%≦Si≦1.20%, 0.10%≦Mn≦0.60%, 7.0%≦Cr≦9.0%, 1.50%≦Mo+0.5W≦2.10%, 0.10%≦V≦0.40%, 0.0002%≦O≦0.0080% and 0.001%≦Al≦0.10%.

4. The cold working die steel as claimed in claim 1, wherein the steel components are 1.40%≦C≦1.60%, 0.10%≦Si≦0.40%, 0.10%≦Mn≦0.60%, 11.0%≦Cr≦13.0%, 0.80%≦Mo+0.5W≦1.20%, 0.10%≦V≦0.40%, 0.0002%≦O≦0.0080% and 0.001%≦Al≦0.10%.

5. The cold working die steel as claimed claim 1, further containing, as the steel component, either one of or both of 0.0030%≦N≦0.0500% and 0.001%≦P≦0.040%.

6. The cold working die steel as claimed in claim 1, further containing any one of, or two or more of steel components selected from 0.01%≦Cu≦1.0%, 0.01%≦Ni≦1.0%, 0.2%≦Co≦1.0% and 0.0003%≦B≦0.010%.

7. The cold working die steel as claimed claim 1, further containing any one of, or two or more of steel components selected from 0.001%≦S≦0.20%, 0.005%≦Se≦0.10%, 0.005%≦Te≦0.10%, 0.0002%≦Ca≦0.010%, 0.005%≦Pb≦0.10% and 0.005%≦Bi≦0.10%.

8. The cold working die steel as claimed claim 1, further containing any one of, or two or more of steel components selected from 0.01%≦Nb≦0.12%, 0.005%≦Ta≦0.10%, 0.005%≦Ti≦0.10%, 0.005%≦Zr≦0.10%, 0.005%≦Mg≦0.10% and 0.005%≦REM≦0.10%.

9. The cold working die steel as claimed in claim 1, wherein the spherodizing comprises keeping the steel heated, and cooling the steel to as low as 750° C. or lower at a cooling rate slower than 60° C./h.