JP2002080930A - Wear resistant steel having excellent toughness and delayed fracture resistance and its production method - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide wear resistant steel having excellent toughness and delayed fracture resistance. SOLUTION: This wear resistant steel having excellent toughness and delayed fracture resistance has a composition containing, by weight, 0.05 to 0.4% C, 0.1 to 0.8% Si, 0.5 to 2.0% Mn, 0.05 to 2.0% Cr, 0.005 to 0.5% Ti, 0.0005 to 0.005% B, 0.005 to 0.10% Al and <=0.005% N, and the balance iron with inevitable impurities, the surface layer part has a martensitic structure, the internal part has a mixed structure of a martensitic structure and a lower bainitic structure or a lower bainitic single phase structure, and the old austenitic grain expanding degree expressed by the ratio of the old austenitic grain size (dL) in the rolling direction to the old austenitic grain size (dZ) in the thickness direction, i.e., (dL/dZ) is >=2.

Description

DETAILED DESCRIPTION OF THE INVENTION

[0001]

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to an abrasion-resistant steel material having excellent toughness and delayed fracture resistance used for industrial machines, transportation equipment and the like used in the fields of construction, civil engineering, mining, and the like, and a method for producing the same. About.

[0002]

2. Description of the Related Art The life of an earth and sand abrasion portion of an industrial machine or a transportation device (eg, a power shovel, a bulldozer, a large dump truck, etc.) used in the fields of construction, excavation of civil engineering and mining, etc. is governed by the amount of wear. Therefore, it is required that the steel material has excellent wear resistance. Therefore, in order to improve the wear resistance, the hardness of the steel material is increased by changing the structure of the steel material to a martensite structure.

[0003] The hardness of a steel material having a martensite structure is uniquely determined by the C content. Therefore, the hardness of the steel material can be improved by increasing the C content. However, on the other hand, there is a problem that the toughness of the steel material is remarkably deteriorated and delayed fracture due to hydrogen in the steel material is easily caused.

[0004] Therefore, it has been premised to select steel components so as to improve the hardness of the steel material by improving the hardenability, that is, to improve the wear resistance and to secure the toughness and the delayed fracture resistance. .

[0005] Further, in order to improve both toughness and delayed fracture resistance, the steel is hot-rolled, air-cooled, re-heat-quenched, and if necessary, tempered to provide abrasion resistance. The production of steel has been done.

Specifically, Japanese Patent Application Laid-Open No. 10-102185 discloses that old austenite grains are refined by rolling in an austenite recrystallization temperature range, and Cr, M
There is disclosed a manufacturing method in which each element of o and V is dissolved in a matrix, followed by quenching and tempering. During the use of the steel material obtained by this manufacturing method, Cr, Mo, V
Is formed, and the generated composite precipitate improves wear resistance. Therefore, the content of C, Mn, etc., which contributes to the improvement of the wear resistance but is harmful to the toughness, can be reduced.

Japanese Patent Application Laid-Open No. 11-71631 discloses that the C content is reduced, the decrease in hardenability accompanying the reduction in the C content is compensated for by increasing the Si content, and the pinning effect of Nb is used. Thus, a production method has been disclosed which suppresses austenite grains from being coarsened during a reheating operation, thereby improving toughness.

Further, Japanese Unexamined Patent Publication No. 60-59019 discloses that the delayed fracture resistance is improved by reducing the Mn content, and the decrease in hardenability due to the reduction in the Mn content is confirmed by the addition of Cr and Mo. Discloses a manufacturing method that compensates for this.

However, these conventional manufacturing methods are:
Since all of them include a reheating and quenching step, there is a disadvantage that the manufacturing process is complicated and an increase in product cost cannot be avoided. Further, an increase in the amount of additional elements such as Cr, Mo, V, and Nb further increases the manufacturing cost.

On the other hand, in order to reduce the production cost, it is desired to omit the air cooling operation and the reheating quenching operation after the above-described hot rolling, and to establish a production method in which a direct quenching operation is performed instead. I have.

Specifically, JP-A-8-41535 discloses a production method in which Si and Nb are added in combination as steel components, and the steel is directly quenched and then tempered. According to this manufacturing method, both elements of Si and Nb suppress temper embrittlement and temper softening.
b acts on the refinement of the prior austenite grains to achieve both abrasion resistance and toughness, and the addition of Mo can further improve the hardenability and toughness.

Japanese Patent Application Laid-Open No. 1-255622 discloses that, after direct quenching, high-temperature tempering is performed for the purpose of reducing tensile residual stress in a steel sheet and improving delayed fracture resistance. There is disclosed a manufacturing method in which a decrease in hardness caused by tempering is compensated by adding Nb.

Further, Japanese Patent Application Laid-Open No. 63-317623 discloses that the content of Mn, which enhances delayed fracture susceptibility, is reduced,
By compensating for the decrease in hardness of the steel material due to the reduction of the Mn content by adding Nb, and performing low-temperature tempering after direct quenching by adding Ti to precipitate Ti nitrides and Ti carbonitrides, It is disclosed that the interface with the matrix acts as a hydrogen trap site to improve delayed fracture resistance.

The above-mentioned JP-A-8-41535, JP-A-1-255622 and JP-A-63-3176 are disclosed.
The technology described in Japanese Patent Publication No. 23 is to select a steel composition and to soften the matrix by tempering to improve toughness and delayed fracture resistance, and to improve the toughness and delayed fracture resistance. It is intended to compensate for the decrease in hardness by including a specific component element and to secure wear resistance.

[0015]

However, a problem common to the above-mentioned prior arts is that the addition of a specific component element necessitates an increase in the cost, thereby offsetting the cost reduction by saving the process. Could be In addition, since the tempering step is indispensable in any of the prior arts, there is a concern that an increase in manufacturing cost and a decrease in production efficiency may be caused.

The present invention has been made in view of such circumstances, and an object of the present invention is to provide a wear-resistant steel material excellent in toughness and delayed fracture resistance without selecting a special steel. Is to do.

Another object of the present invention is to provide a manufacturing method capable of reducing the manufacturing cost in manufacturing such a wear-resistant steel material.

[0018]

The present inventors have conducted intensive studies on wear-resistant steel materials having excellent toughness and wear resistance. As a result, the austenitic grains were subjected to strong pressure reduction in the austenite non-recrystallization temperature range. Immediately after the morphology control, direct quenching is performed, and the quenching is stopped halfway in a specific temperature range to make the surface layer a martensitic structure.
By making the internal part a mixed structure of martensite and lower bainite or a lower bainite single phase structure, a wear-resistant steel material having excellent toughness and delayed fracture resistance has been developed. The present invention has been completed based on such findings.

The wear-resistant steel material according to the present invention, which has excellent toughness and delayed fracture resistance, has a C content of 0.05 to 0.4 in mass%.
%, Si: 0.1 to 0.8%, Mn: 0.5 to 2.0
%, Cr: 0.05 to 2.0%, Ti: 0.005 to
0.5%, B: 0.0005 to 0.005%, Al:
0.005 to 0.10%, N: 0.005% or less, the balance substantially consisting of iron and inevitable impurities,
The surface layer has a martensite structure, the internal portion has a mixed structure of a martensite structure and a lower bainite structure or a lower bainite single phase structure, and the former austenite in the rolling direction relative to the former austenite grain size (dZ) in the thickness direction. It is characterized in that the prior austenite grain elongation expressed by the ratio of particle size (dL) (dL / dZ) is 2 or more.

In this case, in mass%, Cu: 0.
1 to 1.0%, Ni: 0.1 to 1.0%, Mo: 0.1
１．０1.0% and V: preferably one or more selected from the group consisting of 0.01 to 0.2%.

In mass%, Nb: 0.005 to 0.5%.
You may make it contain 1% further.

The method for producing a wear-resistant steel material having excellent toughness and delayed fracture resistance according to the present invention is as follows.
-0.4%, Si: 0.1-0.8%, Mn: 0.5-
2.0%, Cr: 0.05 to 2.0%, Ti: 0.00
5 to 0.5%, B: 0.0005 to 0.005%, A
l: a preparation step of preparing a steel material containing 0.005 to 0.10% and N: 0.005% or less, with the balance substantially consisting of iron and unavoidable impurities, and after heating the prepared steel material, A rolling step of hot rolling to a cumulative draft of 50% or more in a temperature range of 900 ° C. or less;
Start quenching from a temperature range of 3 points or more, Ms point -250
And a quenching step of stopping quenching in a temperature range of not less than ° C and an Ms point + 100 ° C or less.

In this case, the steel material in the preparation step is as follows: Cu: 0.1-1.0%, Ni: 0.1-1.
0%, Mo: 0.1-1.0% and V: 0.01-
It is preferable to further contain one or more selected from the group consisting of 0.2%.

Further, the steel material in the preparation step is expressed as
b: 0.005 to 0.1% may be further contained.

Hereinafter, a wear-resistant steel material having excellent toughness and delayed fracture resistance according to the present invention will be described.

First, the function of each of the above components and the reasons for limiting the component ranges will be described. In addition, “%” in the following component ranges means “% by mass”.

(1) C: 0.05 to 0.40% C not only increases the hardness of the steel material but also generates fine carbides in the lath of the martensite structure described later, thereby improving toughness and delayed fracture resistance. Has a function. C content is 0.
When the content exceeds 0.4%, the weldability is deteriorated, quenching cracks are easily generated, and it is difficult to improve the toughness and the delayed fracture resistance while securing the wear resistance. C content is preferably 0.05 to 0.3%.

(2) Si: 0.1-0.8% Si has a function as a deoxidizing agent in steel making. In order to function effectively as a deoxidizing agent, the addition amount is required to be 0.1% or more. However, if the addition amount exceeds 0.8%, the toughness of the weld may be impaired. The Si content is preferably 0.25 to 0.55%.

(3) Mn: 0.5 to 2.0% Mn has a function of improving hardenability and a function of improving toughness at low cost, and its content is required to be 0.5% or more. If it exceeds 2.0%, the weldability may be impaired, and delayed fracture tends to occur. The Mn content is preferably 1.0 to 1.6%.

(4) Cr: 0.05 to 2.0% Cr has a function of improving hardenability at low cost.
The effect is small at a Cr content of less than 0.05%,
If it exceeds 2.0%, weldability and toughness may be impaired. The Cr content is preferably 0.05 to 1.5%.

(5) Ti: 0.005 to 0.5% Ti combines with N in the steel, and fixes this N to B
Has the function of securing the hardenability by
It is dispersed and precipitated as C and contributes to improvement of wear resistance. 0.
When the Ti content is less than 005%, such an effect is hardly obtained, while when it exceeds 0.5%, the cost tends to increase.

(6) B: 0.0005 to 0.005% B has a function of improving the hardenability by adding a small amount of B. If the B content is less than 0.0005%, it is difficult to exert its effect. On the other hand, if the B content exceeds 0.005%, the weldability may be detrimental and the hardenability may be reduced.

(7) Al: 0.005 to 0.10% Al has a function as a deoxidizing agent at the time of steel making, and its content needs to be 0.005% or more. If it exceeds, the toughness may be reduced. The Al content is preferably 0.015 to 0.035%.

(8) N: 0.005% or less N easily combines with the above-mentioned B and inhibits the hardenability.
If the N content exceeds 0.005%, the fixation of N by Ti in the above-specified content range may be insufficient. Therefore, the upper limit of the N content is 0.005%
And

The basic components of the steel material according to the present invention are as described above. In order to further improve its properties, Cu, Ni,
One, two or more selected from Mo, V and Nb can be contained.

(9) Cu: 0.1 to 1.0% Cu has a function of further improving the hardenability. If the Cu content is less than 0.1%, this effect is small, and if it exceeds 1.0%, hot brittleness may be caused. The Cu content is preferably 0.1-0.3%.

(10) Ni: 0.1 to 1.0% Ni has a function of further improving toughness and hardenability. The Ni content is required to be 0.1% or more, but if the Ni content is less than 0.1%, the effect is small, and if it exceeds 1.0%, the cost tends to increase. Ni content is preferably 0.1 to
0.3%.

(11) Mo: 0.1 to 1.0% Mo has a function of further improving the hardenability, and its content is required to be 0.1% or more. The weldability and toughness may be impaired. The Mo content is
Preferably, it is 0.1 to 0.5%.

(12) V: 0.01 to 0.2% V has a function of further increasing the hardness of steel due to precipitation hardening, and its content is required to be 0.01% or more.
%, The weldability may be impaired. The V content is preferably 0.01 to 0.1%.

(13) Nb: 0.005 to 0.1% Nb has an effect different from those of the above-mentioned components (1) to (12), and suppresses recrystallization at the time of rolling to reduce the recrystallization. It has the function of facilitating the expansion of austenite grains and improving the toughness. If the Nb content is less than 0.005%, such a function may not be effectively achieved.
If it exceeds 0.1%, the weldability may be impaired. Nb
The content is preferably 0.005 to 0.03%.

The steel material according to the present invention contains each component in the above specific range, wherein the surface layer has a martensite structure, and the internal portion has a mixed structure of a martensite structure and a lower bainite structure or Lower bainite single phase structure. Here, the surface portion refers to a portion having a depth of 1 mm from the surface of the steel material.

In the present invention, it is sufficient that the surface layer has a martensitic structure and at least 80% of the inner part has a mixed structure of a martensite structure and a lower bainite structure or a lower bainite single phase structure. The upper part is allowed to partially contain an upper bainite structure or a ferrite structure. Further, the martensite structure may be a tempered martensite structure in which carbides are precipitated by reheating after the intermediate quenching-intermediate stop described later.

Further, the steel material of the present invention has a prior austenite grain elongation (dL / dz) expressed by a ratio of a prior austenite grain size (dL) in the rolling direction to a prior austenite grain size (dz) in the thickness direction. Is 2 or more. In the present invention, the prior austenite grain elongation is set to 2 or more because the effect of improving toughness and delayed fracture resistance is exhibited thereby. This is due to the former austenite grain elongation, the absorbed energy (J), the delayed fracture initiation stress intensity factor (N / mm ^3/2 ), and the Brinell hardness (HB10 / 3) of the steel sheet.
000) can be understood from the characteristic diagram of FIG.

The above-mentioned prior austenite grain elongation can be determined, for example, based on a heat treatment grain size test method based on a quenching and tempering method specified in Japanese Industrial Standard JIS G 0551, based on the cross section along the thickness direction of the steel sheet and the rolling direction of the steel sheet. Can be determined by measuring the particle size of the prior austenite grains that have appeared in the cross section along the line. The value of the prior austenite grain elongation taken on the horizontal axis of FIG. 1 is 5 at the plate thickness t / 2.
The average value of the measured values obtained when observed in the visual field was used.

[0045] The steel sample used here was expressed as mass%
And C: 0.23%, Si: 0.45%, Mn: 1.5
5%, P: 0.011%, S: 0.005%, Cr:
0.31%, Ti: 0.018%, B: 0.0019
%, Al: 0.035%, N: 0.0029%, and the balance being iron, after rolling under various conditions, the surface layer portion is made to have a martensite structure, and the internal portion is made up of a martensite structure and a lower portion. It is a steel plate (plate thickness 50 mm) having a mixed structure with a bainite structure or a lower bainite single phase structure. The measurement of Brinell hardness taken on the vertical axis of FIG. 1 is based on JIS Z22.
Based on No. 43, a Brinell hardness (HB10 / 3000) test was conducted to measure the diameter of a recess formed when a 10 mm diameter indenter was pushed into the steel sheet surface. The absorbed energy taken along the vertical axis in FIG. 1 is based on JIS Z 2202 taken from the center of the sheet thickness.
A Charpy impact test was performed at −40 ° C. using a 2 mm V notch test piece of × 10 mm to apply an impact force perpendicular to the rolling direction of the sample. The value of the absorbed energy was obtained by performing the above-mentioned impact test three times and calculating the average value of the obtained measured values. Further, the stress intensity factor for delayed fracture initiation stress taken on the vertical axis in FIG. 1 is obtained by immersing a test piece in a 3.5 mass% NaCl aqueous solution and applying a predetermined load in a direction perpendicular to the rolling direction of the test piece. This is the maximum stress intensity factor that leads to fracture in a cantilever type constant load delayed fracture test. The measurement of the fracture at this time was 10
00 hours was the longest.

In FIG. 1, the value of Brinell hardness is almost constant at about 400 regardless of the value of the prior austenite grain elongation. Also, the former austenite grain extension elongation dL / dz
Where dL / dz is less than 2, the absorbed energy and the delayed fracture initiating stress intensity factor are extremely low in the region where dL / dz is less than 2, while the absorbed energy and the delayed fracture stress are in the region where dL / dz is 2 or more. A remarkable transition in which the expansion coefficient sharply increased was observed, and it was found that excellent toughness and delayed fracture resistance were exhibited. That is, by setting the austenite grain extension and elongation to 2 or more, the surface layer portion to have a martensite structure, and the internal portion to have a mixed structure of martensite and lower bainite or a lower bainite single phase structure, thereby achieving a desired resistance. It has been found that it has excellent toughness and delayed fracture resistance while maintaining wearability.

The remarkable improvement in toughness and delayed fracture resistance is due to the fact that the lath length is reduced in the mixed structure of the martensite structure and the lower bainite structure or in the lower bainite single phase structure, and in the lower bainite lath. It is considered that fine carbides are preferentially precipitated in the above. That is, regarding toughness, (a) the crack length is likely to be bent or branched due to the shortened lath length, and (b) the fine carbide preferentially precipitated in the lath structure is a barrier that suppresses crack propagation. The above, (a),
It is considered that the toughness is improved by both actions (b). On the other hand, the delayed fracture resistance is considered to be because the interface between the fine carbide and the matrix in the lath structure serves as a hydrogen trap site, and the occurrence of delayed fracture is suppressed.

Next, a method for producing a wear-resistant steel material having excellent toughness and delayed fracture resistance according to the present invention will be described.

The method for producing steel according to the present invention comprises:
C, Si, Mn, Cr, Ti, B, Al and N are contained so as to satisfy the specific ranges shown in the above (1) to (8),
A steel material whose balance substantially consists of iron and unavoidable impurities is prepared.

In this preparation step, one or more of Cu, Ni, Mo, and V may be further contained as component elements in the component ranges described in (9) to (12) above. Further, Nb may be further contained as a component element in the component range shown in the above (13).

Next, after heating the prepared steel material,
Rolling is performed at a cumulative draft of 50% or more in a temperature range of 0 ° C. or less. Here, the upper limit temperature of 900 ° C. means the average temperature from the surface of the steel sheet to the center of the steel sheet, and the same applies to the temperature in the following description. In actual production, the temperature is substantially controlled by the steel sheet surface temperature. However, it is necessary to calculate the average temperature in real time and control the temperature based on the average temperature.

The heating temperature of the steel before rolling is 950
The temperature is preferably 1250C. This heating temperature is 9
If the temperature is lower than 50 ° C., the deformation resistance of the steel increases, so that it becomes difficult to perform rolling. If the heating temperature exceeds 1250 ° C., the crystal grains of the steel become coarse, and it becomes difficult to obtain desired strength and toughness.

The reason for setting the temperature range at the time of rolling to 900 ° C. or lower is that the temperature range of 900 ° C. or lower corresponds to a temperature range lower than the austenite recrystallization temperature, and the austenite grains expanded by rolling disappear. This is to maintain the form without causing it to occur.

In the production method of the present invention, when the rolling conditions are such that the cumulative rolling reduction is less than 50%, the austenite grain extension / extension dL / dz becomes less than 2. On the other hand, dL / dz can be made 2 or more by making the cumulative draft be 50% or more. In fact, we have 90
This is made clear from the characteristic diagram of FIG. 2 showing the correlation between the cumulative draft (%) in the temperature range of 0 ° C. or lower and the prior austenite grain elongation dL / dz. The steel sample used here is a steel having the same composition as that used in FIG. 1 rolled to various cumulative reductions in a temperature range of 900 ° C. or less, the surface layer has a martensite structure, and the inner portion has a martensite structure. It has a mixed structure of a site structure and a lower bainite structure or a lower bainite single phase structure.

As shown in FIG. 2, the old austenite grain elongation dL / dz increases almost proportionally with an increase in the cumulative rolling reduction. According to this, the cumulative rolling reduction is 50%.
DL / dz is less than 2 in a region of less than 2, and dL / dz is 2 or more in a region of a cumulative reduction ratio of 50% or more. Therefore, it can be seen that by performing rolling at a cumulative draft of 50% or more in a temperature range of 900 ° C. or less, dL / dz can be controlled in form to 2 or more.

In the production method of the present invention, the steel material rolled to a cumulative draft of 50% or more in the above-mentioned temperature range of 900 ° C. or less is provided.
Immediately quenching directly from the temperature range of Ar3 or more
The quenching is stopped in a temperature range of s point -250 ° C or more and Ms point + 100 ° C or less.

Here, the Ar3 point is, for example, Ar3
(° C) = 910-310C% -80Mn% -20Cu%
−15Cr% −55Ni% −80Mo% (“%” shown here is “mass%” of each component element in the steel material, and the same applies to the following Ms point). The relational expression can be derived based on the composition of the steel material. Similarly, the above Ms point is, for example, Ms
(° C) = 517-300C% -33Mn% -22Cr%
It can be derived based on the component composition of the steel material by a relational expression expressed by -17Ni% -11Mo% -11Si%.

As described above, the direct quenching after rolling is because the above-mentioned rolling effect is weakened in the case of reheating quenching.
This is because there is a possibility that the effect of improving the toughness and delayed fracture resistance of the steel material obtained after quenching may not be obtained, and the manufacturing process is increased because the process is complicated.

The above-mentioned quenching start temperature is set to a temperature range of Ar3 point or higher because a desired surface hardness (for example, Brinell hardness HB (10
/ 3000) cannot be obtained.

Further, the quenching stop temperature is set at Ms point-250 ° C.
The reason why the temperature range is defined as the Ms point + 100 ° C. or less is that the surface layer of the steel material has a martensite structure and the internal portion has a mixed structure of martensite and lower bainite or a lower bainite single phase structure. is there.

As described above, by direct quenching and stopping halfway, only the surface portion of the steel sheet having a high cooling rate can have a martensite structure, and the internal portion can be cooled completely so as not to have a martensite single phase structure. Since the quenching is stopped halfway, a mixed structure of martensite and lower bainite or a lower bainite single phase structure can be obtained.
This makes it possible to obtain a wear-resistant steel material like a composite material that can improve the toughness and the delayed fracture resistance as a whole while securing the surface hardness.

The cooling rate during direct quenching is 10 to
The cooling rate is preferably 50 ° C./second, and such a cooling rate can be easily achieved by, for example, water quenching.

In the present invention, tempering by reheating may be performed after quenching is stopped halfway.

FIG. 3 shows that the vertical axis represents the Brinell hardness HB (10 /
3000), absorbed energy (J), and stress intensity factor of delayed fracture initiation stress (N / mm ^3/2 ), and quenching stop temperature (° C.) is plotted on the horizontal axis, and Brinell hardness, absorbed energy, and delayed fracture initiation stress It is a characteristic diagram which shows the result of having investigated the relationship between three characteristics of an expansion coefficient and quenching stop temperature, respectively. When the sample used here is directly quenched after rolling a steel having the same composition as that described with reference to FIG. 1 in a temperature range of 900 ° C. or less, the cumulative rolling reduction during rolling is 65% and 35%. It was obtained by changing the quenching stop temperature variously while changing the temperature. Note that FIG.
The Brinell hardness, absorbed energy and delayed fracture initiation stress intensity factor measured on the vertical axis were measured in the same manner as described with reference to FIG.

In FIG. 3, curves A1 and A connecting white circles
2, A3 are characteristic lines showing the results when rolling was performed to a cumulative draft of 65% in a temperature range of 900 ° C. or lower, and curves B1, B2, and B3 connecting black circles were 35% in a temperature range of 900 ° C. or lower.
6 is a characteristic line showing a result when rolling is performed to a cumulative reduction ratio of%.

Focusing on the characteristic lines B1, B2, and B3 in FIG. 3, in the region where the quenching stop temperature is 485 ° C. (Ms point + 100 ° C.) or less, although the Brinell hardness is as high as 300 or more, delayed fracture occurs. It was found that the stress intensity factor and the absorbed energy were extremely low, and the toughness and delayed fracture resistance were extremely poor. The quenching stop temperature is 48
In the region exceeding 5 ° C. (Ms point + 100 ° C.), both the delayed fracture initiation stress intensity factor and the absorbed energy tend to increase, but the Brinell hardness tends to decrease, and the Brinell hardness decreases to less than 300, resulting in poor wear resistance. It turned out to be inferior.

On the other hand, focusing on the characteristic lines A1, A2, and A3, in the region where the quenching stop temperature is less than 135 ° C. (Ms-250 ° C.), although the Brinell hardness is maintained at a high value of about 500, delayed fracture The generated stress intensity factor and absorbed energy are significantly reduced, and the quenching stop temperature is 4
In the region exceeding 85 ° C. (Ms point + 100 ° C.), although the delayed fracture initiation stress intensity factor and the absorbed energy show high values, the Brinell hardness becomes low at less than 300, and the wear resistance is secured in any region. It was found that toughness and delayed fracture resistance could not be improved. On the other hand, the quenching stop temperature is 135 ° C. (Ms
C.) to 485.degree. C. (Ms point + 100.degree. C.), not only the Brinell hardness is as high as 300 or more, but also the delayed fracture initiation stress intensity factor and the absorbed energy increase to show high values. It has been found to have excellent toughness and delayed fracture resistance while maintaining hardness.

The quenching stop temperature is set to Ms point -2 even when the austenite grain elongation is set to 2 or more by rolling at a cumulative draft of 50% or more in a temperature range of 900 ° C. or less.
When the temperature range is lower than 50 ° C., the effect of improving toughness and delayed fracture resistance is not recognized because a martensitic single-phase structure is formed from the surface layer portion to the central portion of the plate thickness.

On the other hand, if the quenching is stopped in a temperature range exceeding the Ms point + 100 ° C., the Brinell hardness decreases because the surface layer cannot have a martensitic structure.

FIG. 4 shows that the vertical axis represents the Brinell hardness HB (10 /
3000), the absorbed energy (J) and the stress intensity factor for delayed fracture ^initiation (N / mm ^3/2 ), the quenching stop temperature (° C.) and the tempering temperature (° C.) on the horizontal axis, and the direct quenching-tempering process. About the steel sheet obtained by
The relationship between the tempering temperature, the Brinell hardness, absorbed energy, and the stress intensity factor for delayed fracture initiation was determined for the steel sheet obtained by the direct quenching-intermediate stop process, with the quenching termination temperature, Brinell hardness, absorbed energy, and delayed fracture initiation. FIG. 8 is a characteristic diagram showing the results of examining the relationship with the stress intensity factor.

Curves A4, A5, A connecting white circles in FIG.
6 are characteristic lines showing the results in the case of the direct quenching-intermediate stop process, respectively, and the curves B4, B5,
B6 is a characteristic line showing the result in the case of the direct quenching-tempering process, respectively. The composition of the sample used here is the same steel type as that described with reference to FIG. 900 ° C. in the direct quenching-intermediate stop process
The cumulative draft in the following temperature range was 65%. On the other hand, in the direct quenching-tempering process, the cumulative draft in a temperature range of 900 ° C. or less was set to 35%, and the steel was directly quenched in a temperature range of 100 ° C. or less and then tempered at various tempering temperatures.

In FIG. 4, focusing on the characteristic lines A6 and B6, if both the quenching stop temperature and the tempering temperature are the same value, the Brinell hardness shows almost the same value, and almost the same wear resistance is obtained. It turns out that it has. Next, focusing on the characteristic lines A4 and B4, 135 ° C.
(Ms point-250 ° C)-485 ° C (Ms point + 100 ° C)
In the above temperature range, the characteristic line A4 has a significantly higher value of the delayed fracture initiation stress intensity factor over the above temperature range than the characteristic line B4, and it has been found that the characteristic line A4 has excellent delayed fracture resistance. Focusing on the characteristic lines A5 and B5, 135
° C (Ms point -250 ° C) to 485 ° C (Ms point +100
In the temperature range of (° C.), the characteristic line A5 has a significantly higher absorbed energy value over the above temperature range than the characteristic line B5, and it has been found that the characteristic line A5 has excellent toughness. That is, even if they have the same Brinell hardness, the steel sheet obtained under the specified manufacturing conditions as in the present invention has higher absorption energy and delayed fracture initiation stress intensity factor, and secures wear resistance. It was also found that they also had excellent toughness and delayed fracture resistance.

In the direct quenching-tempering process described with reference to FIG. 4, the cumulative rolling reduction in the temperature range of 900 ° C. or less is 50%.
However, if the cumulative rolling reduction is to be rolled to 50% or more, tempering may be performed after quenching without stopping halfway after rolling. In this case, the upper limit of the tempering temperature is a temperature at which a surface hardness of 300 or more in Brinell hardness HB (10/3000) can be secured.

[0074]

Embodiments of the present invention will be described below.

(Examples) Steels adjusted to the component compositions of steel types A to H shown in Table 1 were smelted. In Table 1, the Ar 3 point (° C.) and the Ms point (° C.) determined based on the composition of each steel type are also shown.

Next, using these steel types A to H, Table 2
A steel sheet is manufactured according to the manufacturing conditions shown in
A 100 mm steel plate was obtained.

[0077]

[Table 1]

With respect to each of the obtained steel sheets, the structure at a position 1 mm below the surface layer and the structure at the center of the sheet thickness were observed with an optical microscope and a transmission electron microscope. In observation of these structures, for example, martensite and lower bainite can be discriminated by a difference in the precipitation form of carbide by observing a thin film sample roughly with an optical microscope, specifically, a transmission electron microscope.

Further, the austenite grains were revealed based on the heat treatment grain size test method according to the quenching and tempering method of JIS G 0551, and the austenite grain elongation (dL / dZ) at the center of the sheet thickness was determined. Further, a hardness measurement, a Charpy impact test, and a delayed fracture test were performed in the same manner as described with reference to FIG. In this example, the Brinell hardness value obtained by the above hardness measurement was 3
The condition was that the absorption energy value in the Charpy impact test was 17 J or more, and the stress intensity factor for delayed fracture initiation stress in the delayed fracture test was 980 N / mm ^3/2 or more.

Table 2 shows the evaluation results examined above.

[0081]

[Table 2]

As shown in Table 2, each of the steel materials of Examples 1 to 7 satisfying the specific component composition and the specific manufacturing conditions has a prior austenite grain elongation dL / dz of 2 or more. The surface layer had a martensite structure, and the internal portion had a mixed structure of a martensite structure and a lower bainite structure or a lower bainite single phase structure. Each of the steel materials of Examples 1 to 7 had a Brinell hardness of 300 or more, an absorbed energy of significantly more than 17 J, and a delayed fracture initiation stress intensity factor of 980 N / mm.
It was found that the steel material greatly exceeded ^3/2 and had not only excellent wear resistance but also excellent toughness and delayed fracture resistance.

On the other hand, in each of the steel sheets of Comparative Examples 1 to 4 in which the cumulative rolling reduction in the temperature range of 900 ° C. or less was less than 50%, the surface layer had a martensitic structure and the inner part had a martensitic structure. Although it was a mixed structure of a martensite structure and a lower bainite structure or a lower bainite single phase structure, the former austenite grain extension elongation dL / dz was less than 2. Each of the steel sheets of Comparative Examples 1 to 4 had a Brinell hardness of 30.
0 or more, but the absorbed energy and the delayed fracture initiation stress intensity factor are less than 17 J and 980 N / mm, respectively.
^{It was} less than ^3/2, which proved to be inferior in toughness and delayed fracture resistance.

Further, the steel sheet of Comparative Example 5 in which the quenching stop temperature was higher than the Ms point Ms point + 100 ° C. based on the composition of steel type C had a prior austenite grain elongation dL / dz of 2
As described above, the surface layer had a lower bainite structure, and the internal portion had an upper bainite structure. This steel sheet was found to have a Brinell hardness of less than 300 and poor wear resistance.

The steel sheet of Comparative Example 6 in which the quenching stop temperature was lower than the Ms point of −250 ° C. had a martensitic single-phase structure over the entire sheet thickness, although the austenite grain extension / extension dL / dz was 2 or more. there were. Although this steel sheet had a remarkably high Brinell hardness and excellent wear resistance, it was found that the absorbed energy and the stress intensity factor for delayed fracture initiation were extremely low, respectively, and both toughness and delayed fracture resistance were poor.

[0086] The steel sheet of Comparative Example 7 in which the quenching start temperature was lower than the Ar3 point was the former austenite grain extension elongation dL.
/ Dz was 2 or more, and although the internal part had a mixed structure of martensite and lower bainite, the surface layer had a martensite structure in which a ferrite structure was significantly mixed. Although this steel sheet satisfies the above conditions with the absorbed energy and the delayed fracture initiation stress intensity factor, it was found that the Brinell hardness did not satisfy the above conditions and was inferior in wear resistance.

The steel sheet of Comparative Example 8 manufactured using steel type H having a composition in which the C content deviated from the specific range was reduced from the prior austenite grain elongation dL / dz of 2 or more to the surface layer structure and the surface layer. Although the structure from the part to the central part of the sheet thickness was a specified structure, it was found that the Brinell hardness was extremely low and the wear resistance was poor.

In Comparative Example 9, since the cumulative draft in the temperature range of 900 ° C. was 65%, which is 50% or more, immediately after rolling, the austenite grain elongation dL / dz was expanded to 2 or more. However, since it was reheated to the austenite region and quenched, dL / dz was 1.2, which was less than 2. Further, since the quenching stop temperature was set to a temperature lower than the Ms point of −250 ° C., a martensite single phase structure was formed over the entire plate thickness. Although this steel sheet satisfied the above conditions with the Brinell hardness and the delayed fracture initiation stress intensity factor, it was found that the absorbed energy was less than 17 J and the toughness was poor.

The cumulative rolling reduction in the temperature range of 900 ° C. or less is 5
35%, which is less than 0%, and the quenching stop temperature is Ms point−
Temperature below 250 ° C, 450 ° C after quenching
The steel sheet of Comparative Example 10 which had been tempered had a prior austenite grain extension elongation dL / dz of less than 2, and had a martensite single phase structure over the entire sheet thickness. This steel sheet was found to have a Brinell hardness of less than 300 and inferior wear resistance, although the absorbed energy and the delayed fracture initiation stress intensity factor satisfied the above conditions.

[0090]

As described above, according to the present invention, a wear-resistant steel material having excellent toughness and delayed fracture resistance while maintaining desired hardness is provided. Further, according to the present invention, since such a wear-resistant steel material can be manufactured online, the manufacturing cost can be significantly reduced. Therefore, according to the present invention, both toughness and delayed fracture resistance can be improved in the wear-resistant steel material, and the effect of contributing to the industry such as improvement of reliability and workability of industrial machines such as civil engineering machines is extremely large.

[Brief description of the drawings]

FIG. 1: Old austenite grain elongation, Brinell hardness,
FIG. 7 is a characteristic diagram showing the results of examining the relationship between the absorbed energy and the three characteristics of the delayed fracture initiation stress intensity factor.

FIG. 2 is a characteristic diagram showing a result of examining a relationship between a cumulative rolling reduction rolled in a temperature range of 900 ° C. or less and prior austenite grain elongation.

FIG. 3 is a characteristic diagram showing the results of examining the relationship between the quenching stop temperature and three characteristics of Brinell hardness, absorbed energy, and stress intensity factor for delayed fracture initiation.

FIG. 4 shows the relationship between the tempering temperature and the three properties of Brinell hardness, absorbed energy, and the stress intensity factor for delayed fracture initiation in the steel sheet obtained by the direct quenching-tempering process. FIG. 9 is a characteristic diagram showing the results of examining the relationship between the quenching stop temperature and three properties of Brinell hardness, absorbed energy, and delayed fracture initiation stress intensity factor for each of the obtained steel sheets.

 ────────────────────────────────────────────────── ─── Continued on the front page F term (reference) 4K032 AA01 AA02 AA04 AA05 AA11 AA12 AA14 AA16 AA19 AA21 AA23 AA31 AA35 AA36 BA01 CA02 CB02 CC03 CD06

Claims (6)

[Claims] 1. C .: 0.05 to 0.4% by mass, S
i: 0.1 to 0.8%, Mn: 0.5 to 2.0%, C
r: 0.05 to 2.0%, Ti: 0.005 to 0.5
%, B: 0.0005 to 0.005%, Al: 0.00
5 to 0.10%, N: 0.005% or less, the balance substantially consists of iron and unavoidable impurities, the surface layer has a martensite structure, and the internal portion has a martensite structure and a lower bainite structure. And a lower bainite single phase structure, and the ratio (dL / d) of the former austenite grain size (dL) in the rolling direction to the former austenite grain size (dZ) in the thickness direction.
A wear-resistant steel material excellent in toughness and delayed fracture resistance, characterized in that the austenite grain extension and elongation represented by Z) is 2 or more. 2. In mass%, Cu: 0.1-1.0%, N
i: 0.1 to 1.0%, Mo: 0.1 to 1.0%, and V: One or more selected from the group consisting of 0.01 to 0.2%. The wear-resistant steel material according to claim 1, wherein: 3. Nb: 0.005 to 0.1% by mass%
The wear-resistant steel material according to claim 1, further comprising: 4. In mass%, C: 0.05-0.4%, S
i: 0.1 to 0.8%, Mn: 0.5 to 2.0%, C
r: 0.05 to 2.0%, Ti: 0.005 to 0.5
%, B: 0.0005 to 0.005%, Al: 0.00
A preparation step of preparing a steel material containing 5 to 0.10%, N: 0.005% or less, and the balance substantially consisting of iron and unavoidable impurities; and after heating the steel material, a temperature range of 900 ° C or less. Rolling step of hot rolling to a cumulative draft of 50% or more at, and quenching of the rolled steel material immediately from the temperature range of Ar3 point or more and quenching at the temperature range of Ms point -250 ° C or more and Ms point + 100 ° C or less And a quenching step of stopping the abrasion-resistant steel material having excellent toughness and delayed fracture resistance. 5. The steel material in the preparation step is represented by mass%
u: 0.1 to 1.0%, Ni: 0.1 to 1.0%, M
5. The production according to claim 4, further comprising one or more selected from the group consisting of o: 0.1 to 1.0% and V: 0.01 to 0.2%. Method. 6. The steel material in the preparation step is represented by mass%
The method according to claim 4, further comprising b: 0.005 to 0.1%.