JP2001288512A - Method for producing high strength steel with excellent toughness and ductility - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide sufficient toughness and low ductility such as uniform elongation, as well as excellent low-temperature toughness, and excellent toughness and ductility at low temperatures. To provide a method for producing high strength steel. A while optimizing the steel composition and D _I value,
In order to appropriately disperse the hard second layer capable of lowering the yield ratio and improving uniform elongation without deteriorating the toughness without deteriorating the toughness after working heat treatment or reheating treatment, the Ac ₁
After reheating to the transformation point + 20 ° C. or higher and the Ac ₃ transformation point + 150 ° C. or less, the accelerated cooling stop temperature is 300 to 600 ° C.
In addition, a heat treatment characterized by performing accelerated cooling at a cooling rate of 1 to 100 ° C./s is performed.

Description

DETAILED DESCRIPTION OF THE INVENTION

[0001]

BACKGROUND OF THE INVENTION The present invention relates to a steel having sufficient strength as a welding structural steel, a low yield ratio, excellent ductility such as uniform elongation, and excellent toughness and ductility at low temperature toughness. The present invention relates to a method for producing an excellent high-tensile steel.
The steel produced by this method can be used, for example, in general in welded steel structures such as offshore structures, pressure vessels, shipbuilding, bridges, buildings, line pipes, etc., while achieving a low yield ratio, high ductility and toughness at the same time. Because it can be used, it is particularly useful as a steel material for structures such as buildings and bridges that require earthquake resistance. Also,
Although the form of the steel material is not particularly limited, it is particularly useful for a steel plate used as a structural member and requiring low-temperature toughness, particularly a thick plate, a steel pipe material, or a shaped steel.

[0002]

2. Description of the Related Art Reduction in yield ratio (yield stress / tensile strength),
It is known that it is effective to disperse an appropriate amount of a hard phase such as a martensite phase in a soft phase ferrite (α) to improve ductility properties, particularly uniform elongation. This soft α
Various methods for producing a duplex stainless steel comprising a hard phase and a hard phase have heretofore been proposed, but a method of performing an intermediate heat treatment of heating to a ferrite (α) + austenite (γ) dual phase region between quenching and tempering heat treatment (hereinafter, referred to as “hardening and tempering”) , QLT treatment) is basically intended to mix α as a soft phase and bainite or martensite as a hard phase.

[0003] The overall strength level, yield ratio, and ductility properties have been controlled by changing the mixture ratio of these phases. Various production methods for obtaining a mixed structure of the soft phase and the hard phase have been conventionally proposed, for example,
JP-A-53-23817 discloses a method of reheating and quenching a steel sheet, then reheating between the Ac ₁ transformation point and the Ac ₃ transformation point to form two phases of γ and α and then air cooling. ,
Japanese Patent Application Laid-Open No. Hei 4-314824 discloses a method of similarly reheating to a two-phase region and then quenching. In addition, as a method for producing online without performing reheating treatment, for example, JP-A-63-286517 discloses a method in which hot rolling is performed from the γ region to the two-phase region, and then the temperature is increased from the Ar ₃ transformation point by 20%. A method is disclosed in which air is cooled to a temperature lower by １００100 ° C. to generate an α phase, and then quenched.

After reheating and quenching, Ac₁Transformation point
And Ac_ThreeReheating during the transformation point and as two phases of γ and α?
Process including two-phase zone heat treatment of air-cooling or water-cooling
The structure is relatively easy to control, but the two-phase region heat treatment remains
In this case, the toughness extremely deteriorates.₁Transformation point not yet
It is indispensable to perform a tempering process when it is full. For this reason, Q
LT processing is a complicated process, and there is a great
It becomes a title. Also, Ac ₁Perform tempering below the transformation point
And, to reduce the strength of the hard phase and strengthen precipitation in the α matrix,
The degree of lowering the yield ratio is limited, and the
High uniform elongation rather deteriorates.

[0005]

In the heat treatment step of dispersing a soft phase and a hard phase, which corresponds to a two-phase heat treatment, deterioration of toughness can be suppressed, and if tempering below the Ac ₁ transformation point can be omitted, productivity can be improved. It is clear that the improvement of properties (yield ratio, uniform elongation) can be achieved at the same time, and in the present invention, the low yield ratio and the high ductility manufactured by the reheating treatment method including the two-phase heat treatment are included. It is an object to provide means for securing toughness in steel.

[0006]

Means for Solving the Problems The present inventors have proposed a conventional O.D.
In the LT treatment, the mechanism by which the toughness is degraded by heat treatment in the two-phase region is examined in detail, and it is necessary to make the structure of the hard phase into a low-temperature transformed bainite-based structure that has a small adverse effect on toughness and a small strength decrease. I found it important. However, in a chemical composition that can achieve tensile strength as a structural steel, in order to ensure a low yield ratio and ductility characteristics, it is necessary to relatively increase the proportion of the α phase, which is a soft phase compared to a hard phase. Under the heat treatment conditions in the two-phase region for obtaining such a structure ratio, since the C concentration in the inversely transformed γ phase in the heating stage is remarkable, the hardenability of the γ phase is more than considered from the average chemical composition. Is also very high, it is not easy to uniformly low-temperature transformation bainite phase, the present inventors studied a new manufacturing method based on the reheating treatment method, in a wide chemical composition range As a bainite phase having a good toughness as the second phase, a low yield ratio and a high uniform elongation can be achieved without deteriorating the toughness.The following new means have been found, and the present invention has been completed. The summary is as follows That.

(1) C: 0.01 to 0.2% by weight
5%, Si: 0.01-1%, Mn: 0.1-3%, P
: 0.02% or less, S: 0.01% or less, Al:
0.001-0.1%, N: 0.001-0.01%
And the balance consists of Fe and inevitable impurities, and
(1) Desired Hardening critical diameter (D _I value) 0.5 indicated by the formula
After heating the slab which is 30 to the Ac ₃ transformation point to 1300 ° C., after hot rolling including rolling with a starting temperature of 950 ° C. or less, an end temperature of 700 ° C. or more, and a cumulative draft of 30 to 95%, Ac ₁ transformation point + 20 ° C or higher, Ac ₃ transformation point +150
After reheating to below ℃, accelerated cooling stop temperature is 300-6
A method for producing a high-strength steel excellent in toughness and ductility, characterized by performing accelerated cooling at a temperature of 00 ° C and a cooling rate of 1 to 100 ° C / s. D _I = 0.5 · (C%) ^1/2 · (1 + 0.64 · Si%) · (1 + 4.10 · Mn%) · (1 + 0.27 · Cu%) · (1 + 0.52 · Ni%)・ (1 + 2.33 ・ Cr%) ・ (1 + 3.14 ・ Mo%) ・ (1 + 1.50 ・ W%) (1) (2) After hot rolling, start from 650 ° C. or higher, 600 ° C
The above-mentioned (1), wherein the reheating is performed after performing the accelerated cooling at a cooling rate of 1 to 100 ° C./s which is completed below.
3. The method for producing a high-tensile steel excellent in toughness and ductility according to item 1.

(3) C: 0.01-0.2% by weight
5%, Si: 0.01-1%, Mn: 0.1-3%, P
: 0.02% or less, S: 0.01% or less, Al:
0.001-0.1%, N: 0.002-0.01%
And the balance consists of Fe and inevitable impurities, and
(1) Desired Hardening critical diameter (D _I value) 0.5 indicated by the formula
After hot rolling the slab of No. 30, normalizing the heating temperature to the Ac ₃ transformation point or more and the Ac ₃ transformation point +250 or less,
Then, Ac ₁ transformation point + 20 ° C or higher, Ac ₃ transformation point +1
After reheating to 50 ° C or less, accelerated cooling stop temperature is 300
A method for producing a high-strength steel excellent in toughness and ductility, characterized by performing accelerated cooling at a temperature of up to 600 ° C and a cooling rate of 1 to 100 ° C / s. D _I = 0.5 · (C%) ^1/2 · (1 + 0.64 · Si%) · (1 + 4.10 · Mn%) · (1 + 0.27 · Cu%) · (1 + 0.52 · Ni%)・ (1 + 2.33 ・ Cr%) ・ (1 + 3.14 ・ Mo%) ・ (1 + 1.50 ・ W%) ・ ・ ・ ・ ・ (1)

(4) C: 0.01-0.2% by weight
5%, Si: 0.01-1%, Mn: 0.1-3%, P
: 0.02% or less, S: 0.01% or less, Al:
0.001-0.1%, N: 0.002-0.01%
And the balance consists of Fe and inevitable impurities, and
(1) Desired Hardening critical diameter (D _I value) 0.5 indicated by the formula
After hot rolling a slab of No. ₃₀ , the heating temperature is higher than the Ac ₃ transformation point, lower than the Ac ₃ transformation point +250, and the cooling rate is 1 to 1.
After quenching at 00 ° C / s, the Ac ₁ transformation point +2
After reheating to 0 ° C. or higher and the Ac ₃ transformation point + 150 ° C. or lower, accelerated cooling at an accelerated cooling stop temperature of 300 to 600 ° C. and a cooling rate of 1 to 100 ° C./s is performed. A method for producing high strength steel with excellent toughness and ductility. D _I = 0.5 · (C%) ^1/2 · (1 + 0.64 · Si%) · (1 + 4.10 · Mn%) · (1 + 0.27 · Cu%) · (1 + 0.52 · Ni%)・ (1 + 2.33 ・ Cr%) ・ (1 + 3.14 ・ Mo%) ・ (1 + 1.50 ・ W%) ・ ・ ・ ・ ・ (1)

(5) The steel slab is in weight%, Ni: 0.1 to 6
%, Cu: 0.05 to 1.5%, Cr: 0.05 to 2
%, Mo: 0.1 to 2%, W: 0.2 to 4%, V:
0.005 to 0.5%, Ti: 0.003 to 0.1%,
Nb: 0.005 to 0.5%, Ta: 0.01 to 0.5
%, Zr: 0.005 to 0.1%, B: 0.0002
The production of a high-tensile steel excellent in toughness and ductility according to any one of the above (1) to (4), further comprising one or more of 0.005% to 0.005%. Method. (6) The billet is weight%, Y: 0.001-0.1%,
Ca: 0.0005 to 0.01%, Mg: 0.0001
-0.01%, REM: 0.005-0.1%, one or more of which are further contained, any one of the above (1) to (5), A method for producing high strength steel with excellent toughness and ductility. (7) Prior to hot rolling, 2 at 1150-1300 ° C.
The method for producing a high-strength steel excellent in toughness and ductility according to any one of the above (1) to (6), wherein a solution treatment for holding for h to 48 h is performed.

[0011]

BEST MODE FOR CARRYING OUT THE INVENTION For the purpose of improving high ductility, particularly uniform elongation, factors that cause deterioration in toughness when a structure in which a soft phase and a hard phase having significantly different hardnesses are dispersed are manufactured by heat treatment. As a result of experimentally examining the improvement guidelines in detail, the reverse transformed austenite and other untransformed structures formed in the heating stage of the two-phase treatment are both fine, and
When the structure morphology generated during cooling in the two-phase heat treatment is specific, lower yield ratio, higher ductility (uniform elongation, elongation at break),
It was found that both toughness was improved. That is, as described below, after the chemical composition is optimized, the steel structure before the heat treatment in the two-phase region is refined in the working heat treatment or the reheating treatment, and then the heat treatment for dispersing the hard phase in the soft phase is further performed. In the heat treatment, the thermal history of the cooling stage is controlled to make the structure of the hard phase a low-temperature transformed bainite-based structure, thereby reducing the yield ratio and increasing the ductility (uniform elongation, elongation at break). ), And improvement in toughness is achieved at the same time.

Specifically, after heating the slab to the Ac ₃ transformation point to 1300 ° C., the heat including rolling at a starting temperature of 950 ° C. or less, an ending temperature of 700 ° C. or more, and a cumulative rolling reduction of 30 to 95%. Cold rolling is performed, and if necessary, cooling after hot rolling is started at 650 ° C. or higher, and ended at 600 ° C. or lower.
whether to s accelerated cooling, or, after hot rolling a slab, heating temperature Ac ₃ transformation point or higher, Ac ₃ transformation point + 250 ° C. or less, or subjected to normalizing, or the cooling rate is 1 After quenching at 100 ° C / s to properly refine the prestructure,
"Ac ₁ transformation point + 20 ° C or higher, Ac ₃ transformation point + 150 ° C
After reheating below, accelerated cooling stop temperature is 300-60
Perform accelerated cooling at 0 ° C. and a cooling rate of 1 to 100 ° C./s ”is a requirement for the production method of the present invention.

The reasons for limiting the above-mentioned manufacturing method will be described first, and then the reasons for limiting the chemical composition will be described. First, it is necessary to refine the initial structure by the above method or the above method. Is a method based on thermomechanical treatment, and the method is a method based on reheating treatment.

[0014] The initial heat treatment of
In this case, the heating temperature of the slab prior to hot rolling is Ac_ThreeStrange
The temperature is higher than the state point and equal to or lower than 1300 ° C. This is because the heating temperature
Ac _ThreeIf the temperature is below the transformation point, the structure becomes uneven, and
Insufficient solution of precipitation strengthening elements, resulting in strength and toughness
May be deteriorated. On the other hand, when the heating temperature is 1300
If it exceeds ℃, the heated austenite particle size becomes coarse
Even after hot rolling, the microstructure of the pre-structure becomes insufficient,
This is not preferable because the deteriorated state of the billet surface occurs.

After heating the slab to the Ac ₃ transformation point to 1300 ° C., it is necessary to perform hot rolling including rolling at a starting temperature of 950 ° C. or less, an ending temperature of 700 ° C. or more, and a cumulative rolling reduction of 30 to 95%. There is. This is because the austenite grain size before the transformation is refined and the transformed structure is refined by introducing processing strain into the austenite. If the rolling start temperature is higher than 950 ° C., austenite is recrystallized and the refining becomes insufficient. On the other hand, if the rolling end temperature is lower than 700 ° C., ferrite is formed during processing, resulting in a coarse non-uniform structure. Is not preferred. Further, the cumulative rolling reduction of the rolling in the temperature range is 30% because the effect of the working becomes clear.
It is necessary. The higher the rolling reduction, the more effective for the refinement of the structure. However, if the rolling reduction exceeds 90%, the effect of rolling is saturated and the shape of the steel material is adversely affected. And Note that the starting temperature is 95
0 ° C or less, end temperature is 700 ° C or more, and cumulative rolling reduction is 3
If rolling of 0 to 95% is included, there is no problem in performing rolling in a recrystallization region at a temperature exceeding 950 ° C. for the purpose of adjusting the sheet thickness before the rolling.

The cooling after the hot rolling may be either standing cooling or accelerated cooling, but accelerated cooling is more preferable in order to further refine the structure and improve strength and toughness. In the case of performing accelerated cooling, the cooling rate must be in the range of 1 to 100 ° C./s, the starting temperature of the accelerated cooling must be 650 ° C. or more, and the ending temperature must be 600 ° C. or less.

The reason why the cooling rate is limited to the range of 1 to 100 ° C./s is that if the cooling rate is less than 1 ° C./s, the structure refining effect by accelerated cooling is not sufficiently exhibited. This is because while the cooling effect is saturated, there is a concern that adverse effects may occur such as an increase in residual stress and deterioration of the steel sheet shape. Further, the reason why the start temperature of the accelerated cooling is set to 650 ° C. or higher is that if the start temperature is lower than 650 ° C., the transformation starts before the accelerated cooling, and the effect of the accelerated cooling becomes insufficient. On the other hand, the end temperature needs to be 600 ° C. or less because if the end temperature is higher than 600 ° C., the rate of untransformed at the stage where accelerated cooling is completed is excessive, and the effect of accelerated cooling is not sufficiently exhibited. That's why. Any means may be used as long as the accelerated cooling conditions satisfy the present invention. That is, the effect is not different even if cooling by oil cooling or other refrigerant other than water cooling.

Next, a description will be given of limiting conditions in the case where the initial structure is refined by the reheating treatment. When the initial structure is refined by the reheating treatment, hot rolling may be performed only for shape adjustment, and there is no need to particularly limit the rolling temperature, the rolling reduction, and the like as in the thermomechanical treatment.

The refinement of the initial structure by the reheating treatment is performed by normalizing or quenching after adjusting to a desired shape by hot rolling or the like. In the case of normalization, after the heating temperature is reheated to the Ac ₃ transformation point or more and the Ac ₃ transformation point + 250 ° C. or less, the material is allowed to cool. Set the reheating temperature of normalization to A
The reason why the temperature is limited to the c ₃ transformation point or more and the Ac ₃ transformation point + 250 ° C. or less is that if the reheating temperature is less than the Ac ₃ transformation point, an untransformed structure remains and the refining becomes insufficient. ,
On the other hand, when the Ac ₃ transformation point is higher than + 250 ° C., the grain size of the heated austenite becomes coarse, and the refinement of the transformed structure is not guaranteed.

Quenching in which accelerated cooling is performed by water cooling or the like after reheating is a more preferable means for refining the structure because the cooling rate can be increased with the same sheet thickness as compared with the case where cooling is performed after reheating. When the initial structure is refined by quenching, the reheating temperature is set to A for the same reason as normalization.
The temperature is limited to the c ₃ transformation point or more and the Ac ₃ transformation point + 250 ° C. or less. Cooling after reheating is accelerated cooling by water cooling, etc.
In the present invention, the quenching conditions are limited by the cooling rate. That is, in the present invention, the cooling rate is 1 to 10
Limit to 0 ° C / s. This is because if the cooling rate is less than 1 ° C./s, the effect of microstructure refinement by accelerated cooling is not sufficient and there is no point in performing quenching. On the other hand, the higher the cooling rate, the greater the effect of microstructure refinement. If the temperature is higher than ° C./s, the effect of making the structure finer is saturated, and it is not easy to realize industrially. In addition, as long as the cooling rate of the quenching is within the range of the present invention, any means may be used. That is, the effect is not different even if cooling by oil cooling or other refrigerant other than water cooling.

After miniaturization of the initial structure by the above method, the Ac ₁ transformation point + 20 ° C. or higher, and the Ac ₃
After reheating to the transformation point + 150 ° C. or lower, the accelerated cooling stop temperature is 300 to 600 ° C., and the cooling rate is 1 to 100.
Heat treatment is performed by performing accelerated cooling at a rate of ° C./s. This heat treatment is the most important requirement for achieving the properties desired by the present invention. In other words, this heat treatment makes it possible to disperse the fine low-temperature transformed bainite phase with less toughness degradation as a hard phase, to produce a high-strength steel with a low yield ratio, excellent ductility properties such as uniform elongation, and excellent low-temperature toughness. It can be manufactured.

It is necessary to limit the reheating temperature in this heat treatment to not less than Ac ₁ transformation point + 20 ° C. and not more than Ac ₃ transformation point + 150 ° C. This is because it is necessary to form a certain amount or more of the reverse transformed austenite phase in the heating step in order to generate the hard phase. Therefore, at the reheating temperature lower than the Ac ₁ transformation point + 20 ° C., is the reverse transformed austenite not generated? Even if it is produced, its amount is small, so that the final ratio of the hard phase is also small, and it is not possible to secure the strength and improve the strength-ductility balance.
At a reheating temperature of more than the Ac ₃ transformation point + 150 ° C., even if the initial structure before heating is refined by the above method or the above method, the heated austenite grain size becomes coarse, and the generated hard phase becomes coarse, resulting in toughness. Is not preferred because it inhibits Therefore, in the present invention, the reheating temperature in the heat treatment for forming the hard second phase is set to the Ac ₁ transformation point + 20 ° C. or more,
Ac ₃ transformation point is limited to + 150 ° C or lower.

After reheating to an Ac ₁ transformation point + 20 ° C. or more and an Ac ₃ transformation point + 150 ° C. or less, an accelerated cooling process is performed to form a low-temperature transformed bainite-based structure by transformation. Is a cooling rate of 1 to 100 ° C./s and an accelerated cooling stop temperature of 300 to 600.
° C. In order to form a hard second phase mainly composed of low-temperature transformed bainite that achieves both toughness and ductility, the purpose of the present invention, after optimizing the chemical composition as described below, and ensuring a certain level of hardenability, It is necessary to accelerate cooling. If the accelerated cooling rate is less than 1 ° C./s, ferrite and coarse upper bainite unfavorable in toughness may be formed. Also, the higher the cooling rate, the greater the effect of making the structure finer. However, if the cooling speed is higher than 100 ° C./s, the effect of making the structure finer is saturated, and it is not easy to realize industrially. In addition, as long as the cooling rate of the quenching is within the range of the present invention, any means may be used. That is, the effect is not different even if cooling by oil cooling or other refrigerant other than water cooling.

Although the hard phase is formed during the cooling by the accelerated cooling, formation of coarse upper bainite or martensite which deteriorates toughness is suppressed, and the hard phase is transformed into a low-temperature transformed bainite phase main structure which is preferable for achieving both toughness and ductility. Therefore, it is necessary to stop accelerated cooling in the middle. That is,
When accelerated cooling to a low temperature of less than 300 ° C., the hard second phase is likely to become a martensite-based structure.
If the temperature exceeds 00 ° C., the proportion of retained austenite is high at the stage when the accelerated cooling is stopped, and there is a high possibility that coarse upper bainite is formed after the accelerated cooling is stopped. In order to secure the bainite phase, the accelerated cooling stop temperature is limited to 300 to 600C.

According to the present invention described above, when compared at the same strength, superior ductility, particularly uniform elongation, can be achieved as compared with steel produced by the existing method. That is,
According to the existing method, the uniform elongation is, on average, the uniform elongation (U-El) in% and the tensile strength (TS) in MPa, U-El (%) ≦ 20− 0.017 · T. S (MPa) (2) In contrast to the level represented by the above equation (2), according to the present invention, the uniform elongation represented by the equation (3) is achieved. You. That is, when viewed at the same strength, the uniform elongation is better by 4% or more on average than the existing method. U-El (%) ≧ 24−0.017 · T. S (MPa) ... (3)

According to the present invention, the uniform elongation level of the above equation (3) is achieved, and the yield ratio (yield stress / tensile strength) is obtained.
Can be reduced to 80% or less, and the toughness can be improved to -40 ° C. or less at a fracture surface transition temperature in a 2 mmV notch Charpy impact test without performing a subsequent tempering treatment.

In the present invention, the reheating temperature of the final heat treatment for forming the hard second phase is set to the Ac ₁ transformation point + 20 ° C.
As described above, if further limited to the Ac ₃ transformation point −50 ° C. or less,
The yield ratio can be improved to 75% or less and the uniform elongation level can be improved to the extent shown by the expression (4). U-El (%) ≧ 27−0.017 · T. S (MPa) ... (4)

Further, in order to improve toughness and ductility,
If necessary, prior to hot rolling, 1150-130
A solution treatment in which the solution is held at 0 ° C. for 2 hours to 48 hours can be performed. That is, although the slab has inevitably a micro-segregated portion in which the alloy component is enriched, the segregated portion has a low heating transformation point and high hardenability, so that it preferentially becomes a hard phase generation portion. Therefore, if the micro segregation part is wide and / or the component concentration of the segregation part is large, a coarse martensite phase is generated in the final heat treatment step, which leads to deterioration of toughness and ductility. Therefore, depending on the micro-segregation state of the steel slab, toughness and ductility are improved by performing a solution treatment for reducing the micro-segregated portion and reducing the concentration prior to hot rolling.

In the present invention, the solutionizing condition is set at a heating temperature of 1: 1.
150-1300 ° C, holding time: limited to 2h-48h. If the heating temperature is lower than 1150 ° C., even if the heating temperature is maintained for a long time, the diffusion of segregated elements such as Mn is not sufficient, and if the temperature is higher than 1300 ° C., oxidation of the surface becomes remarkable, which is not preferable. However, if the holding time is less than 2 hours, the solution is not enough even at a heating temperature of 1300 ° C., so the lower limit of the holding time is 2 hours. A longer holding time is more effective in reducing segregation, but holding at a high temperature for a long time is economically unfavorable, and there is a problem of surface oxidation. Therefore, even at 1150 ° C., the effect of improving toughness and ductility is clear. The upper limit of the holding time is 48 h. Since the solution treatment is sufficiently achieved in the heating and holding stage of the solution treatment and the structure adjustment is achieved in the subsequent manufacturing process of the present invention regardless of the steel slab structure, cooling after heating and holding in the solution treatment is performed. The conditions are not particularly limited.

The above is the reason for limiting the present invention relating to the production method. The production of a steel for welded structures is sufficiently performed, and a low yield ratio high tensile strength steel sheet having a low low yield ratio and excellent plastic deformability is produced. Therefore, it is necessary to define the chemical components as well. The reasons for limiting each chemical component are described below.

First, C is added as an effective component for improving the strength of steel. If it is less than 0.01%, it is difficult to secure the strength required for structural steel, and more than 0.25%. If the hard phase is a low-temperature-transformed bainite phase, excessive addition causes an excessive amount of C in the bainite to remarkably deteriorate toughness and remarkably reduce weld cracking resistance. The range was 01 to 0.25%.

Next, Si is an element effective as a deoxidizing element and for ensuring the strength of the base material. Addition of less than 0.01% results in insufficient deoxidation and is disadvantageous for securing strength.
Conversely, an excessive addition exceeding 1% forms a coarse oxide and causes deterioration in ductility and toughness. Therefore, the range of Si is 0.01
１1%.

Mn is an element necessary for securing the strength and toughness of the base material, and must be added at least 0.1% or more. However, excessive addition of more than 3% causes toughness degradation due to the hard phase as well as excessive C content, and the toughness of the welded portion,
The upper limit is set to 3% because the cracking property is also deteriorated.

P is an impurity element and is preferably reduced as much as possible, but the upper limit is set to 0.02% as an allowable amount from the viewpoint of securing the toughness of the heat affected zone.

Since S forms MnS and degrades the ductility value, S is an element that needs to be particularly reduced in a steel sheet that is required to ensure ductility as in the present invention. However, the content is set to 0.01% or less as a practically allowable upper limit where ductility is not significantly deteriorated.

Al is an element effective for deoxidation and refining of γ particle size, and it is necessary to contain 0.001% or more in order to exhibit the effect. , A coarse oxide is formed and the ductility is extremely deteriorated, so it is necessary to limit the range to 0.001% to 0.1%.

N is effective in refining γ grains in combination with Al and Ti.
On the other hand, it is necessary to contain 1% or more. On the other hand, if it is added excessively, the amount of dissolved N increases, leading to an increase in yield ratio and deterioration in toughness. The upper limit is set to 0.01% as an allowable range from the viewpoint of securing the toughness of the weld heat affected zone.

The above are the basic components of the steel of the present invention. However, in order to make the hard second phase a low-temperature transformed bainite phase effective for achieving both ductility and toughness, the above-mentioned production method of the present invention is premised. It is necessary to optimize the hardenability of steel. The effect of chemical composition on the hardenability of steel is determined by the ideal hardening critical diameter (D _I
Value). The present inventors have determined the D _I value hardenability of the second phase in the final heat treatment stage can be expressed accurately experimentally as shown in the following equation (1) in the production method of the present invention, the D _I value and ductility was investigated in detail the relationship between the appearance condition of toughness preferred second phase. From the results, it was decided to limit the scope of 0.5 to 30 and D _I value represented by the equation (1). That, D _I value 0.5 to 30
Within this range, the production method of the present invention can suppress the formation of martensite and a coarse upper bainite phase, which are unfavorable for the material, and can reliably form a low-temperature transformed bainite phase that is excellent in uniform elongation and can ensure toughness. When D _I value is less than 0.5, since hardenability is excessively small, low yield ratio, strength - not preferable because the hard phase is not formed play an important role in the uniform elongation balance improved, whereas, 30 If it is more than that, the hardenability becomes excessive, and the martensite phase having poor toughness cannot be suppressed even by the method of the present invention, and it is difficult to form a second phase which is favorable in ductility and toughness, which is not preferable. D _I = 0.5 · (C%) ^1/2 · (1 + 0.64 · Si%) · (1 + 4.10 · Mn%) · (1 + 0.27 · Cu%) · (1 + 0.52 · Ni%)・ (1 + 2.33 ・ Cr%) ・ (1 + 3.14 ・ Mo%) ・ (1 + 1.50 ・ W%) ・ ・ ・ ・ ・ (1)

The above are the basic components of the steel of the present invention. Ni, Cu, Cr, Mo, W, V, Ti, and Nb may be used, if necessary, for the purpose of increasing the base metal strength according to the desired strength level. ,
One, two or more of Ta, Zr, and B can be contained.

First, Ni is a very effective element that can simultaneously improve the strength and toughness of the base material, but it is necessary to contain 0.1% or more in order to exert the effect. When the content is increased, the strength and toughness are improved, but the effect is saturated even if added in excess of 6%.
%. Next, Cu has almost the same effect as Ni, but since addition of more than 1.5% causes a problem in hot workability, it is limited to the range of 0.05 to 1.5%. In addition, Cr
Is an element effective for improving the strength of the base material, but is required to be 0.05% or more in order to produce a clear effect, while 2%
If added in excess of, the toughness tends to deteriorate, so the content is made 0.05 to 2%. Mo is also an effective element for improving the strength of the base material, but it needs to be 0.1% or more in order to produce a clear effect. On the other hand, if added over 2%, the toughness tends to deteriorate. , 0.1 to 2%. W is an element effective for improving the strength of the base material, similarly to Mo. However, in order to produce a clear effect, W is 0.2%.
On the other hand, if added in excess of 4%, the toughness tends to deteriorate, so the content is set in the range of 0.2 to 4%. Both V and Nb mainly contribute to the improvement of the strength of the base material by precipitation strengthening, but the toughness is deteriorated by excessive addition. Therefore, V is 0.005 to 0.5% and Nb is 0.0% as a range in which the effect can be exhibited without inducing the deterioration of toughness.
0.5 to 0.5%. Ti is an element effective for improving the toughness by refining γ grains by forming TiN, but it is necessary to add 0.003% or more to exhibit the effect. On the other hand, if it exceeds 0.1%, as in the case of Al, a coarse oxide is formed to deteriorate toughness and ductility, so the upper limit is made 0.1%. Ta is also effective for precipitation strengthening and grain refinement, but is required to be 0.01% or more in order to exhibit the effect. If it exceeds 0.5%, the toughness is adversely deteriorated. 0.01% to 0.5%. Zr is an element that exerts an effect on precipitation strengthening and grain refinement, but in order to exhibit the effect, 0.005% or more must be added. On the other hand, since an excessive addition of more than 0.1% causes deterioration of toughness due to coarsening of precipitates, the content is limited to the range of 0.005% to 01%. B is very effective in increasing the strength by increasing the hardenability of steel by adding a very small amount of 0.0002% or more, but when added excessively, it forms BN and conversely reduces the hardenability and decreases the toughness. The upper limit is 0.005% to greatly deteriorate
And

Further, in the present invention, for the purpose of improving ductility and toughness of a welded portion (HAZ toughness),
One, two or more of Y, Ca, Mg, and REM can be contained. In each case, the ductility is improved by fine dispersion of oxides and sulfides, and the heat affected zone (HA)
Although the structure of Z) is refined to improve the HAZ toughness,
In order to exhibit the effect, Y is 0.001% or more,
Ca must be contained 0.0005% or more, Mg is contained 0.0001% or more, and REM is contained 0.005% or more.
On the other hand, if added in excess, the oxides and sulfides coarsen and themselves become the starting point of brittle fracture, deteriorating the HAZ toughness.
%, Y and REM are limited to 0.1%.

[0042]

Next, the effects of the present invention will be described more specifically with reference to examples. Table 1 shows the chemical compositions of the test steels used in the examples. Slab numbers 1 to 10 having the chemical components of the present invention
And slab number 11 deviating from the chemical composition range of the present invention.
A steel sheet was manufactured using the slabs Nos. To 15 under the manufacturing conditions shown in Tables 2 and 3, and the mechanical properties were investigated. In each case, a test piece was collected from the center of the plate thickness. The sampling directions were all parallel to the hot rolling direction (L direction). Tables 2 and 3 also show the results of the mechanical tests. The steel sheets shown in Table 2 are based on the method by the thermomechanical treatment of the present invention and the comparative method. The steel sheets shown in Table 3 are based on the method based on the reheating treatment of the present invention and the comparative method. It is.

In Tables 2 and 3, the steel material No. A1 to A
No. 20 was manufactured according to the present invention, B
1 to B11 are comparative examples that do not satisfy any of the requirements of the present invention. As shown in Tables 2 and 3, the steels manufactured according to the present invention achieved a low yield ratio and good toughness at any strength level, and had the same level of ductility, especially uniform elongation, at the same strength level. , It is clear that it is significantly improved as compared with the comparative method. On the other hand, steel material No. Since B1 to B11 do not satisfy the requirements of the present invention, any or all of the yield ratio, uniform elongation, and toughness are clearly inferior to those of the present invention.

Among the comparative examples, steel material No. B1 to B5 are
In the requirements of the present invention, the chemical composition does not satisfy the present invention. Steel No. B1 is inferior in toughness because C is excessively contained. Steel No. B2 is also inferior in toughness because Mn is excessive. Steel No. B3, since D _I value which is an index of hardenability is too small, even by the production method of the present invention, a hardenability lack of hard phase, yield ratio and uniform elongation is not sufficient. On the other hand, when the steel material No. B4 and B5, since D _I value which is an index of hardenability is excessively high, even by the production method of the present invention, since the hard phase inevitably becomes martensite phase undesirable hard toughness, toughness Is significantly deteriorated.

Steel material No. B6 to B11 are examples in which the manufacturing method does not satisfy the requirements of the present invention. Steel No. B6
Nos. To B8 are comparative examples in the manufacturing method by thermomechanical treatment. B6 is inferior in uniform elongation and yield ratio because the heating temperature in the final heat treatment is too high. B7
In the case of the steel material No., the uniform elongation and the toughness were poor because the stop temperature of the accelerated cooling in the final heat treatment was too low. In B8, since the stop temperature of the accelerated cooling in the final heat treatment is too high, formation of a hard phase is not sufficient, uniform elongation and yield ratio are inferior, and toughness is slightly inferior. On the other hand, when the steel material No. B9 to B11 are comparative examples of the manufacturing method by the reheating treatment. Steel No. B9
In steel No. 2, since the normalizing temperature was too high, the deterioration of toughness was large. In B10, since the heating temperature in the final heat treatment was too low, no reverse transformed austenite phase was formed during heating, so that the strength was low, the yield ratio and the uniform elongation were poor, and In B11, since the stop temperature of the accelerated cooling in the final heat treatment is too low, the toughness is significantly deteriorated.

As can be seen from the above examples, it is clear that the production method according to the present invention can achieve a low yield ratio and excellent uniform elongation characteristics without deteriorating toughness.

[0047]

[Table 1]

[0048]

[Table 2]

[0049]

[Table 3]

[0050]

[Table 4]

[0051]

[Table 5]

[0052]

According to the present invention, the steel has sufficient strength as a welded structural steel, has a low yield ratio, has excellent ductility properties such as uniform elongation, and has excellent toughness and ductility excellent in low-temperature toughness. It is possible to provide high-strength steel. Steel produced according to the present invention, for example, offshore structures, pressure vessels,
It can be used in general for welded steel structures such as ships, bridges, buildings, and line pipes.However, because it can achieve both low yield ratio, high ductility and toughness, structures such as buildings and bridges that require earthquake resistance in particular It is particularly useful as a steel material for use, and the industrial value of the present invention is extremely high.

 ────────────────────────────────────────────────── ─── Continued on the front page F term (reference) 4K032 AA01 AA02 AA04 AA05 AA08 AA11 AA12 AA14 AA15 AA16 AA17 AA19 AA20 AA21 AA22 AA23 AA24 AA27 AA29 AA31 AA33 AA35 AA36 AA37 CAA03 CD02 CA03 CD02

Claims (7)

[Claims] 1. C: 0.01 to 0.25% by weight
%, Si: 0.01-1%, Mn: 0.1-3%, P
: 0.02% or less, S: 0.01% or less, Al:
0.001-0.1%, N: 0.001-0.01%
And the balance consists of Fe and inevitable impurities, and
(1) Desired Hardening critical diameter (D _I value) 0.5 indicated by the formula
After heating the slab which is 30 to the Ac ₃ transformation point to 1300 ° C., after hot rolling including rolling with a starting temperature of 950 ° C. or less, an end temperature of 700 ° C. or more, and a cumulative draft of 30 to 95%, Ac ₁ transformation point + 20 ° C or higher, Ac ₃ transformation point +150
After reheating to below ℃, accelerated cooling stop temperature is 300-6
A method for producing a high-strength steel excellent in toughness and ductility, characterized by performing accelerated cooling at a temperature of 00 ° C and a cooling rate of 1 to 100 ° C / s. D _I = 0.5 · (C%) ^1/2 · (1 + 0.64 · Si%) · (1 + 4.10 · Mn%) · (1 + 0.27 · Cu%) · (1 + 0.52 · Ni%)・ (1 + 2.33 ・ Cr%) ・ (1 + 3.14 ・ Mo%) ・ (1 + 1.50 ・ W%) ・ ・ ・ ・ ・ (1) 2. Starting from 650 ° C. or higher after hot rolling,
The method for producing a high-strength steel excellent in toughness and ductility according to claim 1, wherein the reheating is performed after performing accelerated cooling at a cooling rate of 1 to 100 ° C / s, which ends at 600 ° C or lower. . 3. C: 0.01 to 0.25% by weight
%, Si: 0.01-1%, Mn: 0.1-3%, P
: 0.02% or less, S: 0.01% or less, Al:
0.001-0.1%, N: 0.002-0.01%
And the balance consists of Fe and inevitable impurities, and
(1) Desired Hardening critical diameter (D _I value) 0.5 indicated by the formula
After hot rolling the slab of No. 30, normalizing the heating temperature to the Ac ₃ transformation point or more and the Ac ₃ transformation point +250 or less,
Then, Ac ₁ transformation point + 20 ° C or higher, Ac ₃ transformation point +1
After reheating to 50 ° C or less, accelerated cooling stop temperature is 300
A method for producing a high-strength steel excellent in toughness and ductility, characterized by performing accelerated cooling at a temperature of up to 600 ° C and a cooling rate of 1 to 100 ° C / s. D _I = 0.5 · (C%) ^1/2 · (1 + 0.64 · Si%) · (1 + 4.10 · Mn%) · (1 + 0.27 · Cu%) · (1 + 0.52 · Ni%)・ (1 + 2.33 ・ Cr%) ・ (1 + 3.14 ・ Mo%) ・ (1 + 1.50 ・ W%) ・ ・ ・ ・ ・ (1) 4. C: 0.01 to 0.25 by weight%
%, Si: 0.01-1%, Mn: 0.1-3%, P
: 0.02% or less, S: 0.01% or less, Al:
0.001-0.1%, N: 0.002-0.01%
And the balance consists of Fe and inevitable impurities, and
(1) Desired Hardening critical diameter (D _I value) 0.5 indicated by the formula
After hot rolling a slab of No. ₃₀ , the heating temperature is higher than the Ac ₃ transformation point, lower than the Ac ₃ transformation point +250, and the cooling rate is 1 to 1.
After quenching at 00 ° C / s, the Ac ₁ transformation point +2
After reheating to 0 ° C. or higher and the Ac ₃ transformation point + 150 ° C. or lower, accelerated cooling at an accelerated cooling stop temperature of 300 to 600 ° C. and a cooling rate of 1 to 100 ° C./s is performed. A method for producing high strength steel with excellent toughness and ductility. D _I = 0.5 · (C%) ^1/2 · (1 + 0.64 · Si%) · (1 + 4.10 · Mn%) · (1 + 0.27 · Cu%) · (1 + 0.52 · Ni%)・ (1 + 2.33 ・ Cr%) ・ (1 + 3.14 ・ Mo%) ・ (1 + 1.50 ・ W%) ・ ・ ・ ・ ・ (1) 5. A steel slab in weight%, Ni: 0.1-6%,
Cu: 0.05 to 1.5%, Cr: 0.05 to 2%, M
o: 0.1 to 2%, W: 0.2 to 4%, V: 0.0
05-0.5%, Ti: 0.003-0.1%, Nb:
0.005 to 0.5%, Ta: 0.01 to 0.5%, Z
r: 0.005-0.1%, B: 0.0002-0.
The method for producing a high-strength steel excellent in toughness and ductility according to any one of claims 1 to 4, further comprising at least one of 005% and 005%. 6. The steel slab in weight%, Y: 0.001 to 0.001
0.1%, Ca: 0.0005 to 0.01%, Mg:
0.0001-0.01%, REM: 0.005-0.
The method for producing a high-tensile steel having excellent toughness and ductility according to any one of claims 1 to 5, further comprising 1% or more of 1%. 7. Prior to hot rolling, 1150 to 130
The method for producing a high-strength steel excellent in toughness and ductility according to any one of claims 1 to 6, wherein a solution treatment is performed at 0 ° C for 2 hours to 48 hours.