JP2003239036A - Thick steel plate excellent in fatigue strength and method of manufacturing the same - Google Patents

Abstract

(57) [Summary] [PROBLEMS] To add a large amount of special or expensive alloying elements, or to lower productivity or complicated production of thick steel plates for welded structures having excellent fatigue crack propagation resistance of a base metal. It can be obtained regardless of the method and without being greatly limited in tensile strength or steel sheet thickness. SOLUTION: An appropriate amount of C, Si, Mn, Al, N is contained, and if necessary, Ni, Cu, Cr, Mo,
W, Ti, V, Nb, Zr, Ta, B, Mg, Ca, R
In a thick steel sheet containing EM, and having a structure composed of ferrite and a hard second phase composed of bainite or martensite or a mixed structure of both, a hard second phase in a cross section parallel to the steel sheet surface is formed. , Fraction: 20-80%, average Vickers hardness: 250-80
0, the average equivalent circle diameter: 10 to 200 μm, and the maximum distance between the hard second phases: 500 μm or less. By satisfying all the conditions, the fatigue crack growth of the base metal is significantly delayed, and the fatigue strength of the joint is improved. .

Description

Detailed Description of the Invention

[0001]

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to a thick steel plate used for welded structural members requiring fatigue strength and a method for manufacturing the same. INDUSTRIAL APPLICABILITY The steel sheet of the present invention can be generally used for welded steel structures such as marine structures, pressure vessels, ships, bridges, buildings, line pipes, etc., but particularly marine structures, ships and bridges that require fatigue strength. It is useful as a steel plate for structures such as building structures. In addition, it is also applicable to steel pipes or shaped steels made of thick steel plates.

[0002]

2. Description of the Related Art With the increase in size of welded structures and the demand for environmental protection, structural members are required to have higher reliability than ever before. Welded structures are generally used in current structures, and possible fracture modes for welded structures include fatigue fracture, brittle fracture, and ductile fracture.Of these, the most frequent fracture mode is It is a brittle fracture or fatigue fracture from an initial defect, and further a brittle fracture subsequent to the fatigue fracture. In addition, it is difficult to prevent these destruction modes only by considering the design of the structure,
Moreover, it is often the cause of sudden collapse of the structure, and from the viewpoint of ensuring the safety of the structure, it is the most important form of destruction to prevent it.

Regarding brittle fracture, there are means of improvement such as addition of Ni in terms of chemical composition and optimization of transformation structure.
The manufacturing method can also be improved by controlling the structure and refining the structure by rolling or thermomechanical treatment. On the other hand, in the case of fatigue characteristics, it is possible to improve the smooth member by improving the strength, etc., but since the fatigue strength is governed by the toe shape of the welded part in the welded structure, it is possible to improve the strength and structure. It was considered impossible to improve fatigue strength (joint fatigue strength) by metallurgical means. That is, in the structure where fatigue strength is a problem, even if high-strength steel is used, the design strength cannot be increased, and the advantage of using high-strength steel cannot be obtained. Therefore,
Conventionally, in such a welded structure, joint fatigue strength has been improved by so-called toe treatment for improving the shape of the weld toe portion, which is a stress concentration portion.
For example, a method of increasing the toe radius by grinding the toe with a grinder and a method of remelting the toe portion by TIG welding to smooth the toe shape (for example, Japanese Patent Publication No.
-30386), a method in which a compressive stress is generated at the toe by shot peening.

However, since these toe treatments are extremely time-consuming, a means for improving the joint fatigue strength of the steel itself, which does not rely on the toe treatment, is awaited for cost reduction and productivity improvement. It was

In response to such demands, several steel materials having good joint fatigue strength have recently been proposed. For example, the structure of the weld heat affected zone (HAZ) is ferrite (α)
There is disclosed a technique (Japanese Patent Laid-Open No. 8-73983) capable of improving the fatigue strength of the HAZ. However, since the present technology requires the HAZ structure to have a ferrite structure, there is a limit to the strength level of the steel material that can be manufactured, and it is not possible to manufacture a high-strength steel material having a tensile strength exceeding 780 MPa.

Several means have been proposed to improve the joint fatigue strength of high-strength steels having a tensile strength of 590 MPa or more, and a high Si content is used to improve the fatigue crack initiation / propagation characteristics of the bainite structure of HAZ (Japanese Patent Laid-Open No. Hei 8 (1998)). -209295),
There is a report that high Nb (Japanese Patent Laid-Open No. 10-1743) is effective. However, if a large amount of Si and Nb are added,
It is an element that significantly deteriorates toughness and may cause manufacturing problems such as cracking of a steel slab.

In addition, all of the above-mentioned conventional techniques are means for improving fatigue crack initiation in the HAZ structure and fatigue crack propagation in the HAZ, but the HAZ is greatly affected by stress concentration at the toe. , The effect may not be produced or may be small depending on the shape of the toe.

In order to improve the joint fatigue strength regardless of the shape of the toe, it is effective to delay the propagation of the fatigue crack generated from the toe in the base material. Based on such an idea, a technique for improving fatigue crack growth characteristics of a base metal by forming a base metal structure in which coarse ferrite is dispersed in a fine grain structure having an average ferrite grain size of 20 μm or less (special feature (Kaihei 7-90481) is disclosed. However, also in this case, since it is necessary to make the structure mainly composed of ferrite, it is possible to manufacture only steel materials having a tensile strength of about 580 MPa class.

Further, as a technique for increasing the fatigue strength by suppressing the fatigue crack propagation of the base material, in the structure composed of ferrite and the hard second phase, the hardness is constant between the hardness of the ferrite and the hardness of the hard second phase. After defining the relationship of
A technique that defines the form of the second phase (aspect ratio, spacing) and / or the texture is disclosed in JP-A-11-1742.
It is disclosed in the publication. This technology is one of the best means for suppressing fatigue crack propagation among the technologies currently shown, but it is the two-phase region due to the formation of texture and the development of texture.
Since it is necessary to increase the cumulative rolling reduction in the ferrite region, there are problems such as deterioration of productivity and deterioration of steel plate shape.

[0010]

DISCLOSURE OF THE INVENTION The present invention provides a thick steel plate for a welded structure, which is excellent in fatigue crack propagation resistance of a base metal, with a large amount of a special or expensive alloying element added and poor productivity.
Alternatively, it is an object of the present invention to provide it without using a complicated manufacturing method and without being greatly restricted in tensile strength and steel plate thickness.

[0011]

Means for Solving the Problems The inventors of the present invention have found means for improving joint fatigue strength by improving the fatigue crack propagation resistance of a base material without depending on the toe shape of the joint.
It was found from the detailed experimental results of the relationship between fatigue crack growth behavior and steel microstructure. That is, although the fatigue crack generated from the stress concentration part of the joint toe propagates in the plate thickness direction,
The type, morphology, and characteristics of the structure of the crack front have a great influence on fatigue crack growth.

First, as the type of structure, it is preferable to use a mixed structure of a soft phase and a hard second phase rather than a uniform structure on the order of an optical microscope. This is a blunting of the crack when the crack progresses from the hard second phase to the soft phase, while delaying the crack growth when progressing from the soft phase to the hard second phase,
This is because crack detours and branches will occur. In order to generate such crack growth behavior, the structure must contain ferrite as the soft phase, bainite or martensite as the hard second phase, or both. Then, the inventors of the present invention effectively suppress the fatigue crack propagation rate in a steel having a mixed structure containing ferrite and at least bainite or martensite, or both based on the detailed experiment described below. In order to achieve this, it was found that the distribution state is important in addition to the hardness and fraction of the hard second phase. Especially regarding the distribution of the hard phase, it is important to strictly define the distribution of the hard second phase in the front surface of the fatigue crack, that is, in the cross section parallel to the steel plate surface (hereinafter referred to as the Z plane). I found it for the first time.

The experiment was conducted by a three-point bending test of a test piece with a surface mechanical notch shown in FIG. 4 in order to evaluate only the crack propagation characteristic of the base material. Fatigue condition is stress amplitude 378MP
a, the stress ratio was 0.1. The chemical composition of the sample steel
C: 0.05 to 0.2%, Si: 0.15 to 0.3%,
Mn: 0.5-2%, P ≦ 0.01%, S: about 0.00
5%, Al: 0.01 to 0.05%, Nb: 0 to 0.0
5%, Ti: 0 to 0.02%, Ni: 0 to 3%, and various hot rolling conditions and heat treatment conditions (including diffusion heat treatment before hot rolling). Among the microstructures, a small vacuum melting steel (steel plate thickness: 25 mm) was used, in which the type, fraction and distribution of the hard second phase were mainly changed.
The fatigue test piece was taken so that the longitudinal direction of the test piece was parallel to the rolling direction. Investigation of microstructure, hardness measurement is 1 of plate thickness
It was performed on the Z plane of the / 4 portion. Quantification of tissue is 1 of plate thickness
It was performed by using an image analysis device by using a structure photograph of 5 to 10 visual fields in an optical microscope structure of a cross section (Z plane) parallel to the steel plate surface in / 4. The hardness of the hard second phase was also measured at 10 or more points of micro Vickers hardness under a load of 5 to 10 g in the same cross section and evaluated by the average value.

In the test apparatus shown in FIG. 4, mechanical notches N are provided on the surface of the convex test piece A so that fatigue cracks can be easily generated, and rolls are positioned on both sides and the center. By the three-point bending in which a force is applied from this roll in the arrow direction, the fatigue life when an alternating stress is applied is measured, and the fatigue crack propagation characteristics can be evaluated.

FIG. 1 shows a structure in which the hard second phase (hereinafter sometimes simply referred to as the second phase) is adjusted to have a Vickers hardness in the range of 200 to 250 (a pearlite-based phase) and a Vickers. In the second phase fraction and the fatigue test, when the hardness is adjusted to the range of 550 to 600, the stratification is performed for each structure (bainite-martensite main phase) mainly composed of bainite or martensite or a mixed structure of both. The relationship with the breaking life is shown. In the pearlite-based phase, the second phase other than pearlite is less than 5%, while in the bainite to martensite-based phase, the fraction of the pearlite phase other than bainite and martensite is less than 5%. When at least the second phase contains bainite, martensite, or both, and the hardness is high, the fatigue properties are obviously better than when the second phase is mainly pearlite and the hardness is low.

Further, when the second phase is a pearlite-based phase, the fatigue properties do not largely depend on the fraction thereof, whereas when the second phase is a bainite-martensite-based phase, high fatigue properties are ensured. It is necessary to limit the ratio to In particular, if the fraction of the second phase is less than 20%, a large improvement in fatigue properties cannot be expected even if the second phase is a bainite-martensite-based phase. Further, even if the hard phase exceeds 80% and increases, the fatigue properties tend to deteriorate. This is not preferable because the hard phase is too large, and the fatigue crack propagates while causing micro-brittle fracture.

It can be seen from FIG. 1 that it is necessary to make the bainite-martensite main phase having a high hardness the second phase be present in the structure in an amount of 20 to 80% in order to improve the fatigue characteristics. Among them, the fatigue characteristics fluctuate greatly, and it is suggested that there are other factors that strongly control the fatigue characteristics. From the fatigue crack growth mechanism, the present inventors presumed that the variation in the fatigue characteristics was due to the difference in the morphology and distribution of the second phase, and conducted a more detailed study to determine the extension of the second phase. Although there is a slight influence of the elongation, for example, the ratio (aspect ratio) between the length in the rolling direction and the length in the sheet thickness direction of the second phase observed in the sheet thickness sectional structure, the second phase in the Z plane is more than that. We came to discover that the distribution of is important. That is, it is effective for the fatigue crack growth suppression that the second phase present in the fatigue crack front during development is dense and uniform, and its distribution is non-uniform even with the same second phase fraction,
If there is a place where the second phase does not exist depending on the place, the fatigue crack grows preferentially in that place, so the effect of suppressing the fatigue crack growth by the second phase is not sufficiently exerted.

FIG. 2 is a result of clarifying the above-mentioned points. In FIG. 1, the second phase has a high hardness of bainite to martensite and the fraction thereof is in the range of 20 to 80%. 3 shows the relationship between the maximum value of the interval between the second phases observed on the Z plane, based on the definition shown in FIG. 3, and the fracture life of the fatigue test. The second phase interval is measured in the direction perpendicular to the rolling direction on the Z plane at the 1/4 position of the plate thickness in order to correspond to the second phase interval existing on the fatigue crack front surface.

From FIG. 2, it is apparent that the fatigue characteristics are improved as the Z-plane maximum second phase interval is reduced within a range in which the type, hardness, and fraction of the second phase are limited. In particular, when the Z-plane maximum second phase interval exceeds 500 μm, the fatigue characteristics are significantly deteriorated. When the thickness is 500 μm or less, the adverse effect of non-uniformity of the second phase distribution is slight.

The present invention is based on detailed experiments including the above findings to find a structure morphology that is favorable for fatigue crack propagation resistance of a base metal, and to achieve the structure morphology industrially most suitable for achieving the structure morphology. I also invented the preferred means,
The summary is as follows.

(1) C: 0.04 to 0.
3%, Si: 0.01 to 2%, Mn: 0.1 to 3%, A
1: 0.001-0.1%, N: 0.001-0.0
1%, P: 0.02% or less, S
: 0.01% or less, the balance consisting of iron and unavoidable impurities, and having a structure containing at least ferrite and a hard second phase, and in the cross-sectional structure parallel to the surface, the hard second phase is In a thick steel sheet satisfying all the conditions of to, the hard second phase structure is composed of either bainite, martensite, or a mixed structure of both, and is a steel plate excellent in fatigue strength. Fraction of hard second phase: 20-80% Average Vickers hardness of hard second phase: 250-800 Average circle equivalent diameter of hard second phase: 10-200 m Maximum spacing between hard second phases: 500 m or less

(2) Further, in mass%, Ni: 0.0
1-6%, Cu: 0.01-1.5%, Cr: 0.01
~ 2%, Mo: 0.01 to 2%, W: 0.01 to 2
%, Ti: 0.003-0.1%, V: 0.005-
0.5%, Nb: 0.003 to 0.2%, Zr: 0.0
03-0.1%, Ta: 0.005-0.2%, B:
The thick steel sheet having excellent fatigue strength according to (1) above, which contains 0.0002 to 0.005% of one type or two or more types.

(3) Further, in mass%, Mg: 0.0
005-0.01%, Ca: 0.0005-0.01
%, REM: 0.005 to 0.1%, and 1 or 2 or more of the above (1) or (2).
A thick steel plate excellent in fatigue strength according to any one of 1.

(4) A steel slab having the component described in any one of (1) to (3) above and having a casting thickness of 100 mm or less is reheated to an AC ₃ transformation point to 1250 ° C. and a reduction ratio is set. Two
The above hot rolling is performed, and after the hot rolling, the ferrite fraction is cooled to a temperature of 10% or more at a cooling rate of 0.1 to 2 ° C./s, and further to 500 ° C. or less to 5 to 100 ° C. /
A method for producing a thick steel sheet having excellent fatigue strength, which comprises quenching at s.

(5) Before the hot rolling, the billet has a heating temperature of 1150 to 1300 ° C. and a holding time of 1 to 10
The diffusion heat treatment for 0 h is performed, and the above (4)
The method for manufacturing a thick steel sheet having excellent fatigue strength as described in 1.

(6) For a steel slab containing any of the components described in (1) to (3) above and having a casting thickness of more than 100 mm, the heating temperature is 1150 to 1300 before hot rolling.
After performing a diffusion heat treatment at a temperature of 1 ° C for 1 to 100 hours,
After reheating to an AC ₃ transformation point to 1250 ° C. and performing hot rolling with a reduction ratio of 2 or more, after the hot rolling, the ferrite fraction is 10
Of a thick steel sheet having excellent fatigue strength, which comprises cooling to a temperature of 0.1% or more at a cooling rate of 0.1 to 2 ° C / s and then rapidly cooling to 500 ° C or less at 5 to 100 ° C / s. Method.

(7) In the hot rolling, at least a starting temperature of 850 ° C. or lower, an ending temperature of Ar ₃ transformation point or higher, and a cumulative rolling reduction of 30% or higher are included. The method for manufacturing a thick steel sheet having excellent fatigue strength according to any one of (4) to (6) above.

(8) In the hot rolling, at least the starting temperature is lower than the Ar ₃ transformation point, the ending temperature is 600 ° C. or higher, and the cumulative rolling reduction is 10 to 80%. The method for producing a thick steel sheet having excellent fatigue strength according to any one of (4) to (7) above.

(9) After quenching to 500 ° C. or lower, after hot rolling is completed, further (AC ₁ transformation point + 30 ° C.)
To (AC ₃ transformation point −50 ° C.) and subjected to a two-phase region heat treatment of cooling to 500 ° C. or less at 5 to 100 ° C./s, (4) to (8) The method for manufacturing a thick steel sheet having excellent fatigue strength as described in 1.

(10) After being rapidly cooled to 500 ° C. or lower, or after being subjected to a two-phase region heat treatment, 250 to 500 ° C.
The method for producing a thick steel sheet having excellent fatigue strength according to any one of (4) to (9) above, which comprises tempering with.

[0031]

BEST MODE FOR CARRYING OUT THE INVENTION The present invention is to optimize the chemical composition and
Although the optimization of the structural requirements based on the above-mentioned new knowledge is essential, first, the reasons for limiting the structural requirements are explained, then the reasons for limiting the chemical composition are stated, and finally, the method for producing the thick steel sheet of the present invention. With respect to, an embodiment of the manufacturing method proposed in the present invention will be described.

The structural requirement for increasing the fatigue strength is that "a cross section parallel to the surface of the steel sheet has a structure containing at least ferrite and a hard second phase composed of either bainite, martensite, or a mixed structure of both. In the structure, the hard second phase is a fraction of the hard second phase: 20 to 80%,
Hard second phase average Vickers hardness: 250-800,
The average equivalent circle diameter of the hard second phase: 10 to 200 μm, the maximum interval between the hard second phases: 500 μm or less ”are all satisfied.” It was decided based on knowledge.

First, it is necessary to have a structure containing ferrite and a hard second phase composed of bainite, martensite, or a mixed structure of both,
By having a mixed structure of the soft phase and the hard phase, blunting of the crack occurs when the crack propagates from the hard phase to the soft phase,
On the other hand, when the soft phase progresses to the hard phase, the fatigue crack growth rate is significantly suppressed due to the delay of crack growth, detour of cracks, and branching. In order to generate such crack growth behavior, the structure must contain ferrite as the soft phase, bainite or martensite as the hard phase, or both. Ferrite is preferable as the soft phase because ferrite is the only sufficiently soft structure in the low alloy steel for welded structures. If it is a high alloy steel, it is also possible to make the austenite phase a soft phase, in the welded structural thick steel plate that is the subject of the present invention, it is possible to leave the austenite phase in a sufficient proportion in the transformation structure Very difficult and difficult to adopt. In the case of the ferrite phase, no particular problem occurs in the fatigue properties unless the phase is hardened by extreme working or solid solution strengthening. As a guide, the Vickers hardness of the ferrite phase is preferably 220 or less.

The type of the hard second phase is preferably bainite, martensite, or a mixed structure of both, because the structure is more uniform than pearlite and the toughness is good relative to the hardness. . Perlite is not preferable because it is not easy to set the Vickers hardness to 250 or more, and even if the hardness is increased,
This is because the pearlite itself has a layered structure of ferrite and cementite, so that the fatigue crack can selectively propagate the soft ferrite in the pearlite, and the effect of suppressing fatigue crack growth is small. Further, as the hard second phase, precipitates and inclusions are also considered, but it is not easy to contain 20 to 80% of these, which are effective in suppressing fatigue crack growth, and the precipitates and inclusions are bainite. Since it is extremely brittle compared to martensite and martensite, when it is contained in such a large amount, the toughness is remarkably deteriorated, and it cannot be practically used as a structural steel.

For the above reasons, the hard second phase must be either bainite, martensite, or a mixed structure of both, and the fraction, hardness, size, and distribution state of the hard second phase. Must be strictly specified.

The lower limit of the fraction of the hard second phase is 20 from FIG.
%. This is because if the second phase fraction is less than 20%, a clear improvement in fatigue properties cannot be expected even if other structural requirements are optimized. Further, in the present invention, the upper limit of the hard second phase is 80%. This is because it is not easy in terms of chemical composition to set the hardness of the hard second phase to 250 or more after the hard second phase fraction exceeds 80%, and the hard second phase fraction is If it exceeds 80%, the toughness of the hard second phase is inferior to that of ferrite, and there is a concern that the toughness of the steel material may deteriorate. Further, restraining the deformation of ferrite existing between the hard second phases is also disadvantageous in ensuring toughness, and brittle fracture may occur during fatigue crack growth, resulting in deterioration of fatigue properties. The hard second phase fraction in the present invention means the area fraction in the cross-sectional structure in the Z plane.

If the fraction of the hard second phase composed of either bainite or martensite, or a mixed structure of both, satisfies the present invention, the hard second phase other than these is 10
Even if the content of the second phase is less than 10%, it does not substantially affect the fatigue properties even if the content of the second phase is less than 10%. Further, in the present invention, cementite, carbonitrides, and non-metallic inclusions do not show a clear effect on fatigue properties, and thus are not included in the hard second phase.

The hardness of the hard second phase is also an essential requirement for ensuring fatigue characteristics. The type of the second phase is bainite, either or both of martensite, and a mixed structure of both, and the second phase fraction is 20 to 80%, and the hardness necessary for improving the fatigue properties is: Average Vickers hardness of 250
The range is from ~ 800. If the average Vickers hardness is less than 250, the hardness with ferrite, which is a soft phase, is small, and therefore delay in the propagation of fatigue cracks, detours, and branching occur near the soft phase / hard phase interface with sufficient frequency. Therefore, the fatigue characteristics cannot be improved. On the other hand, when the average Vickers hardness of the hard second phase is more than 800, embrittlement of the hard second phase becomes remarkable,
The hard second phase causes brittle fracture even during a fatigue test, rather accelerates fatigue crack growth and deteriorates fatigue properties.

For the above reasons, the fraction and hardness of the hard second phase composed of bainite, martensite, or a mixed structure of both are limited to appropriate ranges, but in order to further improve fatigue properties, There is a need to further limit the size and distribution of the hard second phase.

The size is limited as the harder the second phase becomes harder, the more brittle it becomes, leading to deterioration of toughness and fatigue characteristics of the steel material. It is necessary in order to satisfy both the requirement and the suppression of the toughness deterioration of the hard second phase. The toughness of the hard second phase is dominated by the average equivalent circle diameter, and if it exceeds 200 μm, deterioration of the toughness cannot be ignored. Therefore, in the present invention, the average equivalent circle diameter of the hard second phase is limited to 200 μm or less. To do.
From the viewpoint of ensuring toughness, the finer the size of the hard second phase, the more preferable it is, but if the size of the hard second phase is too small, the fatigue crack growth suppression effect will be insufficient, so fatigue in the present invention. The lower limit of the size of the hard second phase that can reliably exhibit the effect of suppressing crack growth is 10 μm.

Further, as the distribution of the hard second phase, as shown in the result shown in FIG. 3, it is necessary to optimize the interval between the hard second phases observed on the Z plane. In the present invention, based on the result of FIG. 3, the maximum second phase distance in the Z plane of 500 μm where the deterioration of the fatigue characteristics due to the expansion of the hard second phase distance is small is defined as the upper limit. The interval between the second phases in the Z plane is an interval in the direction corresponding to the second phase interval existing on the fatigue crack front surface, and for example, when the crack surface propagates at right angles to the rolling direction. Is also a value in the Z-plane second phase interval in the direction perpendicular to the rolling direction.

The above-mentioned fractions, hardnesses, and distribution states of the structure are all for the Z plane, but since fatigue cracks start at the weld and propagate from the surface in the plate thickness direction, It suffices that the average structure state from the center to the center of the plate thickness satisfies the present invention. If the change in the structure in the plate thickness direction is smaller than the change in the structure in the Z plane, it may be evaluated by the measured value in the Z plane at 1/4 of the plate thickness. When the change in the structure in the plate thickness direction is large, several places in the plate thickness direction, for example, the steel plate surface 1-2 m
m, 1/4 of the plate thickness, or the average value of the plate thickness central portion may be used for evaluation.

The above are the reasons for limiting the organizational requirements in the present invention. In order to secure fatigue properties and to secure the strength and toughness required for structural steel, it is necessary to further optimize the chemical composition as shown below.

That is, C is an effective component for increasing the hardness of the hard second phase. If it is less than 0.04%, it is not easy to stably allow 20% or more of the hard second phase having a Vickers hardness of 250 or more. Therefore, in the present invention, the lower limit of C is 0.04%. However, since an excessive content exceeding 0.3% lowers the toughness and weld crack resistance of the base material and the welded portion, the upper limit was made 0.3%.

Si is an element that is effective as a deoxidizing element and for securing the strength of the base material, but if it is contained in an amount less than 0.01%, deoxidation becomes insufficient and it is disadvantageous for securing the strength. On the contrary, if the content exceeds 2%, a coarse oxide is formed and ductility and toughness are deteriorated. Therefore, the range of Si is 0.01
~ 2%.

Mn is an element necessary to secure the strength and toughness of the base material, and it is necessary to contain at least 0.1%,
If it is contained excessively, the toughness of the base metal, the toughness of the welded portion, the weld cracking property, and the like are deteriorated due to the formation of a hard phase, the grain boundary embrittlement, and the like.

Al is an element effective for deoxidation, refining of the heated austenite grain size, etc., but in order to exert the effect, it is necessary to contain 0.001% or more. On the other hand, 0.1%
If it is contained in excess, the coarse oxide is formed and the ductility is extremely deteriorated. Therefore, it is necessary to limit the content to 0.001% to 0.1%.

N works effectively for austenite grain refinement in combination with Al and Ti, so that a small amount is effective for improving mechanical properties. Further, it is impossible to industrially completely remove N in steel, and it is not preferable to reduce N more than necessary because it puts an excessive load on the manufacturing process. Therefore, the lower limit is set to 0.001% as a range in which industrial control is possible and the load on the manufacturing process is allowable. If it is contained excessively, the solid solution N increases, which may adversely affect the ductility and toughness. Therefore, the upper limit of the allowable range is 0.
It is set to 01%.

Since P is an impurity element and is harmful to various properties of steel, it is preferable to reduce it as much as possible, but in the present invention, the upper limit is 0.02% as an amount that can be adversely affected in practical use. And

S is also an impurity element basically, and particularly has a great adverse effect on the ductility, toughness and fatigue properties of steel.
Reduction is preferred. In practice, the upper limit is set to 0.01% as the amount that can adversely affect. However, in the trace amount of S, fine sulfides are formed and the weld heat affected zone (HAZ)
When considering HAZ toughness, it is preferable to add it in the range of 0.0005 to 0.005% because it contributes to the improvement of toughness.

The above are the reasons for limiting the basic components of the thick steel plate of the present invention. In the present invention, however, Ni, Cu, Cr, Mo, W, T may be added as necessary to adjust the strength and toughness.
One or more of i, V, Nb, Zr, Ta and B can be contained.

Ni is a very effective element because it can improve the strength and toughness of the base material at the same time, but it is necessary to add 0.01% or more in order to exert the effect. As the amount of Ni increases, the strength and toughness of the base material are improved, but if added in excess of 6%, the effect saturates while HAZ
There is a concern that toughness and weldability may be deteriorated, and since it is an expensive element, the upper limit of Ni is set to 6% in the present invention in consideration of economical efficiency.

Cu is an element which has almost the same effect as Ni, but in order to exert the effect, it is necessary to add 0.01% or more. If it exceeds 1.5%, hot workability and H
In the present invention, AZ toughness causes a problem.
It is limited to the range of 01 to 1.5%.

Cr is an element effective for improving strength by solid solution strengthening and precipitation strengthening, and 0.01% is required for producing the effect.
The above is necessary, but if Cr is added in excess, it adversely affects the toughness of the base metal and HAZ through an increase in quenching hardness, formation of coarse precipitates, etc., so the upper limit is limited to 2% as an allowable range. To do.

Similar to Cr, Mo and W are also effective elements for increasing strength by solid solution strengthening and precipitation strengthening.
It is an element that is also effective in securing the hardness of the hard second phase,
Both Mo and W are limited to 0.01 to 2% as a range in which the effect can be exhibited and other characteristics are not adversely affected.

Since Ti forms stable TiN in austenite and contributes to the refinement of the heated austenite grain size of the HAZ as well as the base material, Ti is an element effective not only for improving the strength but also for improving the toughness. However, in order to exert its effect, it is necessary to contain 0.003% or more,
If it is contained in excess of 0.1%, coarse TiN is formed and the toughness is adversely deteriorated. Therefore, in the present invention, the content is limited to 0.003 to 0.1%.

V is an element effective in improving the strength of the base material by precipitation strengthening, but 0.005% is required to exert the effect.
The above is necessary. As the amount of addition increases, the amount of strengthening also increases, but along with that, the base material toughness and HAZ toughness deteriorate,
In addition, since the precipitate becomes coarse and the strengthening effect tends to be saturated, the upper limit is set to 0.5% as the range in which the deterioration of toughness is small with respect to the strengthening amount.

Nb is an element effective for strengthening a small amount by precipitation strengthening and transformation strengthening, and has a great influence on the working and recrystallization behavior of austenite, and therefore is also effective for improving the base material toughness. Furthermore, it is also effective for improving the fatigue characteristics of HAZ. 0.003 to show effect
% Or more is necessary. However, if added in excess of 0.2%, the toughness is extremely deteriorated, so in the present invention, it is limited to the range of 0.003 to 0.2%.

Zr is also an element effective in improving strength mainly by precipitation strengthening, but 0.000 in order to exert the effect.
3% or more is required. On the other hand, if added in excess of 0.1%, coarse precipitates are formed and the toughness is adversely affected, so the upper limit is made 0.1%.

Ta also has the same effect as Nb, and contributes to the improvement of strength and toughness by adding an appropriate amount, but 0.005
If it is less than 0.1%, the effect is not clearly produced, and if it is added excessively over 0.2%, the toughness deterioration due to coarse precipitates becomes remarkable, so the range is made 0.005 to 0.2%.

B is an element that enhances the hardenability in a very small amount,
It is an element effective for strengthening. B is effective even at a very small amount because it enhances hardenability by segregating to austenite grain boundaries in a solid solution state, but if it is less than 0.0002%, a sufficient segregation amount to grain boundaries cannot be secured, so quenching It is not preferable because the effect of improving the property tends to be insufficient or the effect tends to vary. On the other hand, if added in excess of 0.005%, coarse precipitates are often formed during the production of the billet or during the reheating stage, so the effect of improving hardenability becomes insufficient, or cracks in the billet occur. The risk of deterioration of toughness due to precipitates also increases. Therefore, in the present invention, the range of B is 0.0002 to 0.005%.

Further, in the present invention, improvement of ductility,
In order to improve joint toughness, Mg, Ca,
One or two or more types of REM can be contained.

Any of Mg, Ca, and REM is effective in improving ductility by suppressing the expansion of sulfide during hot rolling.
It also works effectively to improve joint toughness by refining the oxide. The lower limit of the content for exerting the effect is 0.0005% for Mg, 0.0005% for Ca, and REM for REM of 0.005%.
It is 005%. On the other hand, if it is contained excessively, coarsening of sulfides and oxides occurs, leading to deterioration of ductility, toughness, and fatigue properties, so the upper limits are 0.01% for Mg and Ca, respectively.
REM is 0.1%.

The above are the reasons for limiting the microstructure and chemical composition, which are the basic requirements of the present invention. In addition, the present invention also presents a suitable manufacturing method for satisfying the organizational requirements of the present invention. However, for the microstructure of the present invention, the effect is exhibited regardless of the means for achieving it, and the method for producing a thick steel sheet excellent in fatigue strength according to claims 1 to 3 of the present invention is It is not limited to the method shown in 4-10.

According to the first manufacturing method, if necessary, before hot rolling, a steel piece is subjected to diffusion heat treatment at a heating temperature of 1150 to 1300 ° C. and a holding time of 1 to 100 hours, and a casting thickness of 10 is obtained.
A steel piece of 0 mm or less is reheated to an AC ₃ transformation point to 1250 ° C., hot rolling is performed with a reduction ratio of 2 or more, and after hot rolling,
0.1 to 2 up to the temperature at which the ferrite fraction becomes 10% or more
After cooling at a cooling rate of ° C / s, it is characterized by further rapidly cooling to 500 ° C or less at 5 to 100 ° C / s.

The second manufacturing method has a casting thickness of 100 mm.
For steel slab that is over, before hot rolling, the steel slab has a heating temperature of 1150 to 1300 ° C. and a holding time of 1 to 100.
After the diffusion heat treatment of h, AC ₃ transformation point to 1250 ° C
Reheating, hot rolling with a reduction ratio of 2 or more, and after hot rolling to a temperature at which the ferrite fraction is 10% or more, 0.
After cooling at a cooling rate of 1 to 2 ° C / s, further 500 ° C
It is characterized by rapid cooling up to the following at 5 to 100 ° C./s.

Further, in both the first and second methods, if necessary, rolling with a start temperature of 850 ° C. or lower, an end temperature of Ar ₃ transformation point or higher, and a cumulative rolling reduction of 30% or higher, or / And, the hot rolling including the rolling of which the starting temperature is the Ar ₃ transformation point or lower and the ending temperature is 600 ° C or higher and the cumulative rolling reduction is 10 to 80% can be performed.

Further, for the steel sheet after hot rolling, (A
C ₁ reheated to transformation point + 30 ℃) ~ (AC ₃ transformation point -50 ° C.), two-phase region a heat treatment is cooled at 5 to 100 ° C. / s up to 500 ° C. or less, or / and a heating temperature of from 250 to 50
It can be tempered at 0 ° C.

The size and distribution of the second phase are greatly influenced by the distribution and transformation behavior of microsegregated portions such as Mn and C which are generated during solidification. It is effective to finely disperse the microsegregation portion for the refining of the hard second phase and the reduction of the interval which are the requirements of the present invention, for that purpose, increasing the cooling rate before and after solidification, It is effective to reduce the secondary dendrite arm spacing by hot rolling. Further, it is also effective to reduce the degree of segregation itself of microsegregation once generated by the diffusion heat treatment.

The casting thickness is set to 100 mm or less in order to make the secondary dendrite arm spacing finer by increasing the solidification rate so that the hard second phase size and spacing of the Z plane in the final structure fall within the range of the present invention. This is because If the casting thickness exceeds 100 mm, it is difficult to reliably set the average equivalent circle diameter of the hard second phase to 200 μm or less and the maximum interval between the hard second phases to 500 μm or less.

A method for finely dispersing the hard second phase regardless of the casting thickness is diffusion heat treatment at a heating temperature of 1150 to 1300 ° C. and a holding time of 1 to 100 hours. For effective diffusion of the alloying elements mainly of Mn, 1 at 1150 ° C
It is necessary to hold at least h. However, if the diffusion heat treatment temperature exceeds 1300 ° C., the heated austenite becomes extremely coarse and adversely affects the toughness of the steel sheet, and the surface of the steel sheet may be roughened, which is not preferable. Hold time is 1h
The longer the above, the better, but the diffusion heat treatment temperature is 1
If the holding temperature is 150 ° C. or higher, the alloying elements can be sufficiently homogenized within a holding time of 100 hours.
00h. Since the effects of the diffusion heat treatment and the reduction of the casting thickness are additive, in the first method,
It is effective to subject the steel piece having a casting thickness of 100 mm or less to the main diffusion heat treatment if necessary.

In the present invention, the casting thickness is 100 mm.
The following steel pieces and / or heating temperature is 1150 to 1300
The steel slab subjected to the diffusion heat treatment at a temperature of 1 ° C for 1 to 100 hours is reheated to an AC ₃ transformation point to 1250 ° C and the reduction ratio is 2
The above hot rolling is performed, and after the hot rolling, the ferrite fraction is cooled to a temperature of 10% or more at a cooling rate of 0.1 to 2 ° C./s, and further to 500 ° C. or less to 5 to 100 ° C. /
The basic condition of the steel sheet manufacturing is to rapidly cool the steel sheet into a thick steel sheet.

The reheating temperature of the steel slab is set to the AC ₃ transformation point to 125.
When the reheating temperature is lower than the AC ₃ transformation point, 0 ° C. does not result in an austenite single phase in the heating stage, and the solid solution of precipitates is not sufficient. On the other hand, if the reheating temperature is higher than 1250 ° C., the heated austenite grain size becomes extremely coarse, and the fine grain is not sufficiently refined by the subsequent hot rolling. This is because the toughness may deteriorate.

For hot rolling, the reduction ratio (casting thickness / steel plate thickness)
Is 2 or more. This is because if the reduction ratio is 2 or more, Z
It is also advantageous for fine dispersion of the hard second phase on the surface, and by reducing the spacing of the hard second phase in the plate thickness direction,
This is because it contributes to the improvement of fatigue characteristics. Further, if the reduction ratio is less than 2, it is difficult to press-bond the porosity existing in the slab due to solidification shrinkage, which is another reason why the reduction ratio is 2 or more.

After hot rolling, the ferrite fraction is cooled to a temperature of 10% or more at a cooling rate of 0.1 to 2 ° C./s, and then rapidly cooled to 500 ° C. or less at 5 to 100 ° C / s. Has an average Vickers hardness of 2 to improve fatigue properties.
The fraction in the structure of the hard second phase which is 50 to 800 is 20
This is to secure ~ 80%. In the low C steel having a C content of 0.3% or less as in the present invention, in order to generate the hard second phase, the transformation temperature range is rapidly cooled and the austenite phase before transformation is enriched with C. Therefore, it is necessary to cause ferrite transformation before quenching. If the ferrite fraction before quenching is less than 10%, the concentration of C in untransformed austenite may be insufficient, so in the present invention, the ferrite fraction before quenching is set to 10% or more. Since the main purpose of the formation of the ferrite is the concentration of C in austenite, it is preferable that the ferrite transformation temperature is high. Therefore, the cooling rate at the time of ferrite formation is 2 ° C./s or less. If the cooling rate is too high, the ferrite transformation temperature is lowered, and the diffusion rate of C is not sufficient, so that C in austenite is reduced.
Is not preferable for thickening. From the viewpoint of the concentration of C in austenite, a smaller cooling rate in the ferrite formation process is preferable, but 0.1 ° C./s or more is sufficient, and the effect is saturated even if gradually cooled. In the present invention, the lower limit cooling rate is 0.1 ° C./s.

The untransformed austenite in which C is sufficiently concentrated is rapidly cooled at a temperature of 5 to 100 ° C./s to 500 ° C. or lower to be transformed at a low temperature to form a hard second phase. When the average cooling rate before and after the transformation region is less than 5 ° C./s, the average Vickers hardness is 250 even within the chemical composition range of the present invention.
It becomes difficult to stably form the above hard second phase. The higher the cooling rate, the more advantageous it is for the formation of the hard second phase.
In addition, cooling at such an excessive cooling rate leads to an increase in manufacturing cost and deterioration of the steel plate shape. For the above reason, in the present invention, the cooling rate in the rapid cooling when forming the hard second phase from untransformed austenite is set in the range of 5 to 100 ° C / s. Quenching at the cooling rate is necessary until the transformation is almost completed. In the chemical composition range of the present invention, if the quenching stop temperature is 500 ° C. or less, a desired hard second phase can be obtained.

Note that the effect of the present invention is impaired even if slabs such as ingots or slabs are subjected to slabbing for the purpose of shape adjustment or the like before hot rolling to form a steel sheet. is not. Further, as long as the conditions of the diffusion heat treatment of the present invention are satisfied, the diffusion heat treatment and the slab rolling are combined, that is, the steel slab is 1 to 1150 to 1300 ° C. which is the diffusion heat treatment condition of the present invention. There is no problem in performing slabbing in the cooling stage after holding for 100 hours.

The above are the reasons for limiting the basic requirements of the manufacturing method in the present invention. In the manufacturing method of the present invention, further, in order to obtain the structural requirements of the present invention and for the purpose of improving mechanical properties, etc. If necessary, in addition to satisfying the basic requirements of the production method of the present invention, the following (a) ~
The process of (d) can be performed. (A) Hot rolling is performed with a start temperature of 850 ° C. or lower, an end temperature of Ar ₃ transformation point or higher, and a cumulative rolling reduction of 30% or higher. (B) Start temperature is below the Ar ₃ transformation point and end temperature is 600
Hot rolling with a cumulative reduction of 10 to 80% is performed at a temperature of ℃ or more. (C) After being rapidly cooled to 500 ° C. or lower, it is reheated to (AC ₁ transformation point + 30 ° C.) to (AC ₃ transformation point −50 ° C.),
A two-phase zone heat treatment of cooling to 500 ° C. or less at 5 to 100 ° C./s is performed. (D) After quenching to 500 ° C. or lower, or after performing a two-phase region heat treatment, tempering at 250 to 500 ° C.

(A) is a step of refining the austenite before transformation and / or accumulating strain in the unrecrystallized austenite to refine the transformation structure. As a result of refining the transformation structure, the hard second phase is finely dispersed and stably,
The average circle equivalent diameter of the hard second phase: 10 to 200 μm and the maximum interval between the hard second phases: 500 μm or less can be satisfied, and the fatigue characteristics are improved. In addition, since the ferrite grain size is also reduced, this means is effective for achieving high toughness as well as fatigue characteristics.

In order to exert the above effects, hot rolling with a cumulative rolling reduction of 30% or more is started at a temperature of 850 ° C. or less,
It is necessary to carry out the finishing temperature at the Ar ₃ transformation point or higher. If the cumulative rolling reduction is less than 30%, the recrystallized austenite grains are not sufficiently refined in the recrystallization region rolling, and the strain is not sufficiently accumulated in the austenite grains in the non-recrystallization region rolling, so that the rolling Regardless of temperature, the refinement of the transformation structure is not sufficient. On the other hand, even if the cumulative rolling reduction is 30% or more, if the rolling temperature is not within the proper range, the rolling effect does not effectively contribute to the refinement of the structure, which is not preferable. That is, when the rolling start temperature is higher than 850 ° C., the refinement of the recrystallized austenite grains is not sufficient, or the speed of recovery of the introduced dislocations is high, and strain is not effectively accumulated. In the present invention, the rolling in the means (a) has an upper limit of the rolling start temperature of 850 from the viewpoint of all contributing to the refinement of the structure.
℃. As long as the rolling is completed in the austenite region, it is effective for the refinement of the structure, so the rolling end temperature in the additional treatment of (a) may be the Ar ₃ transformation point or higher. If the rolling temperature is in the range of 850 ° C. to Ar ₃ transformation point, the effect of rolling is almost accumulated. Therefore, the rolling reduction may be defined by the cumulative rolling reduction, and it is not necessary to specify the rolling reduction condition for each pass. Absent.

(B) is an effective means for promoting the ferrite transformation by the two-phase region rolling and promoting the C concentration in the untransformed austenite to secure the hardness of the hard second phase. In order to enrich C in austenite, it is more reliable to diffuse C from ferrite to austenite at a higher temperature so that the enrichment of C in austenite will increase, and the amount of C in austenite will increase. The hardness of the phase increases.

Ferrite transformation is promoted by performing the two-phase region rolling. Therefore, in the chemical composition range of the present invention, the starting temperature is not higher than the Ar ₃ transformation point and the ending temperature is 60.
It is effective to additionally perform hot rolling at a cumulative rolling reduction of 10 to 80% at 0 ° C or higher. The reason why the starting temperature is set to the Ar ₃ transformation point or lower is that if the transformation point is higher than the Ar ₃ transformation point, the working amount in the two-phase region will be insufficient even if the transformation point is increased by working. The end temperature of rolling is set to 600 ° C. or higher because when the temperature is lower than 600 ° C., transformation from austenite occurs during or after working and before start of quenching,
This is because the pearlite transformation, which has a lower hardness than the hard second phase that should be generated by quenching, may occur. In order to ensure the effect of promoting ferrite transformation by the two-phase region rolling, the cumulative rolling reduction of the two-phase region rolling must be 10% or more. If it is less than 10%, the promotion of ferrite transformation is not sufficient and it is meaningless to additionally perform two-phase region rolling. The upper limit of the cumulative rolling reduction of two-phase rolling is 80%. This is 80%
If an excessively large two-phase region rolling is performed, the transformation from untransformed austenite may be promoted, and the desired hard second phase may not be formed.In addition, industrially, the lower limit of the rolling end temperature is not reached. This is because it becomes difficult to secure the temperature of 600 ° C. By additionally using this means, a texture develops, and it is possible to expect that the fatigue characteristics and toughness of the texture are improved.

(C) secures the hardness and the fraction of the hard second phase more reliably by the two-phase zone heat treatment. When heated to the two-phase region, austenitization in the reverse transformation occurs from the microsegregated portion or pearlite portion where the components are concentrated, so if the hot rolling conditions before heat treatment satisfy the requirements of the present invention, dispersion Regarding the condition, it falls within the scope of the present invention.
The fraction of reverse transformed austenite and the degree of enrichment of C in the austenite are determined according to the heating conditions of the two-phase region heat treatment. That is, the fraction and hardness of the hard second phase after transformation from austenite are determined. If the heating temperature in the two-phase heat treatment is less than (AC ₁ transformation point + 30 ° C.), the reverse transformation austenite fraction, and thus the resulting hard second phase fraction, becomes too small as compared with the present invention, and the fatigue properties are reduced. to degrade. On the other hand, the heating temperature in the two-phase region heat treatment is
If it exceeds (AC ₃ transformation point −50 ° C.), the austenite fraction increases, but instead, the concentration of C in the austenite is insufficient and the hardness of the hard second phase deviates from the present invention. As a result, the fatigue property is also inferior, which is not preferable. Therefore, in the present invention, the heat treatment temperature in the two-phase region is limited to (AC ₁ transformation point + 30 ° C) to (AC ₃ transformation point -50 ° C). (AC ₁ transformation point + 30 ° C) ~ (AC ₃ transformation point-
After reheating to 50 ° C), cooling is 5 to 500 ° C or less.
Cooling is performed at 100 ° C./s. This is similar to the quenching after the hot rolling described above. This is to form a phase.

(D) is a tempering treatment performed after being rapidly cooled to 500 ° C. or lower, or after being subjected to a two-phase region heat treatment, and is performed to adjust the strength and toughness as necessary. However, it is necessary to consider so as not to excessively reduce the hardness of the hard second phase. That is, in the present invention, the upper limit of the tempering temperature is 500 ° C. This has a tempering temperature of 50.
If it is higher than 0 ° C, the average Vickers hardness of the hard phase is 250 depending on the chemical composition in the hot rolling or the two-phase heat treatment stage.
The above is 250 by tempering
This is because there is a concern that it will fall below this level. Further, in the present invention, the lower limit of the tempering temperature is defined as 250 ° C. This is because if the tempering temperature is less than 250 ° C., the material adjusting effect by tempering is not clear.

Next, the effects of the present invention will be described more specifically by way of examples.

[0086]

Examples Table 1 shows the chemical composition of the test steel used in the examples. Each of the test steels is made into a slab by ingot casting, slab rolling, or continuous casting. In Table 1, steel slab numbers 1 to 10 satisfy the chemical composition range of the present invention, and steel slab numbers 11 to 15 do not satisfy the chemical composition range of the present invention. The heating transformation points (AC ₁ , AC ₃ ) are also shown in Table 1, which shows that the heating rate is 5 ° C./min. Although it is the actual measured value at that time, it substantially agrees with the transformation point shown in Table 2 under the actual temperature rising condition at the time of billet heating or heat treatment of the steel sheet.

Steel pieces having the chemical composition shown in Table 1 were subjected to diffusion heat treatment, hot rolling, heat treatment and tempering under the conditions shown in Table 2 to produce a steel plate having a plate thickness of 25 mm or 50 mm, and a tensile property at room temperature of 2 mmV. Notch Charpy impact properties and fatigue properties of welded joints were investigated. The tensile test piece and the Charpy impact test piece were sampled at right angles to the rolling direction (C direction) from the center of the plate thickness. The tensile properties were measured at room temperature, and the Charpy impact properties were evaluated at the 50% fracture surface transition temperature (vTrs). The fatigue test is performed in order to evaluate the fatigue characteristics when a fatigue crack is generated from the weld toe of a structure and propagates through the base metal portion.
The rotating welded joint shown in Fig. 3 was performed. The test piece S is made of a steel plate and has a length in the steel plate longitudinal direction of 300 mm and a length in the width direction of 80.
mm, plate thickness: 25 mm (total thickness for 25 mm thick material,
50 mm thick material is sampled from the surface), and a test plate is sampled in the size of: width: 10 mm, length: 30 mm, height:
A 30 mm rib plate B was turned to the center of the test plate by carbon dioxide gas welding (CO ₂ welding) and welded by welding C. The carbon dioxide welding at this time has a chemical composition of C: 0.06 mass%,
Si: 0.5 mass%, Mn: 1.4 Mass%, using a 1.4 mm diameter welding wire, current: 270
A, voltage: 30 V, welding speed: 20 cm / min. I went there. In the fatigue test, the span of the load fulcrum F and the lower span:
70 mm, upper span: 220 mm, maximum load (P
max): 5,500 kgf, and stress ratio (R): 0.1 was applied repeatedly, and fatigue life was measured.

Table 3 shows the structural morphology (type, fraction, Vickers hardness, average circle equivalent diameter, maximum interval) and mechanical properties of the hard second phase of the steel sheet. The quantification of the structure is 1/4 of the plate thickness.
The optical microscope structure of the cross section (Z plane) parallel to the steel plate surface in FIG. Using a tissue photograph of 5 to 10 fields of view,
It was quantified by an image analyzer. The hardness of the hard second phase was also measured at 10 or more points of micro Vickers hardness under a load of 5 to 10 g in the same cross section and evaluated by the average value.

Steel plate numbers A1 to A13 in Tables 2 and 3 are
It is a steel sheet that satisfies all the requirements relating to the chemical composition and structure of the present invention, and not only has the strength and toughness (2 mm V notch Charpy impact characteristics) required for structural steel, but also an extremely good joint. It is clear that it also has fatigue properties.

On the other hand, steel plate numbers B1 to B11 are comparative steel plates that do not satisfy any of the requirements of the present invention.
It is obvious that the joint fatigue properties and toughness are inferior to those of the steel sheet of the present invention having the same composition and strength level.

Steel sheets Nos. B1 to B5 were good in spite of not satisfying the structural requirements of the present invention or satisfying the structural requirements of the present invention because their chemical compositions did not satisfy the present invention. This is an example in which various characteristics could not be achieved.

In other words, the steel plate No. B1 has an excessively large amount of C, and thus has poor toughness, and also has extremely poor toughness. Therefore, even in a fatigue test, the hard phase is brittle and fractures. In comparison, the joint fatigue properties are inferior.

Steel plate No. B2 has an excessive amount of Mn, so C
For the same reason as when the amount is too large, both the toughness and fatigue properties are significantly inferior to those of the present invention.

Steel sheets Nos. B3 and B4 have an excessive amount of P and N, respectively, and make the steel brittle, so that the toughness and fatigue properties are also significantly inferior to those of the present invention.

Steel plate number B5 has an excessive amount of S,
Since the fatigue characteristics are greatly deteriorated through the deterioration of ductility,
Fatigue properties are inferior to those of the present invention.

Steel sheets Nos. B6 to B11 are examples in which the chemical composition satisfies the present invention, but the joint fatigue characteristics are inferior because the structural requirements do not satisfy the present invention.

That is, Steel Sheet Nos. B6 and B9 are examples in which the fatigue characteristics are inferior to those of the steels of the present invention having the same composition because the intervals between the hard second phases are excessive.

Steel plate numbers B7, B8 and B10, B11
Since the second phase is pearlite which is not preferable for the fatigue property, the fatigue property is significantly inferior to that of the present invention even though the dispersed state of the second phase satisfies the present invention.

From the above examples, it is clear that according to the present invention, it is possible to obtain excellent joint fatigue properties while ensuring sufficiently high toughness as a structural steel.

[0100]

[Table 1]

[0101]

[Table 2]

[0102]

[Table 3]

[0103]

INDUSTRIAL APPLICABILITY The present invention relates to a thick steel plate used for a welded structural member requiring fatigue strength, which has been conventionally difficult to improve in a welded portion. Without relying on a simple manufacturing process,
The industrial usefulness is extremely large in that it can be manufactured without being greatly limited in tensile strength and steel plate thickness.

[Brief description of drawings]

FIG. 1 is a diagram showing a relationship between a rupture life in a surface mechanical notch three-point bending fatigue test, a type of hard second phase, and a fraction.

FIG. 2 is a diagram showing the relationship between the rupture life in the above fatigue test and the maximum distance in the Z plane of the hard second phase.

FIG. 3 is a diagram showing a definition of a maximum interval in a Z plane of a hard second phase.

FIG. 4 is a schematic diagram of a surface mechanical notch three-point bending test piece and a test apparatus for examining a base material fatigue crack propagation characteristic.

FIG. 5 is a schematic diagram of a four-point bending test piece and a test apparatus for measuring a fatigue life when a fatigue crack propagates in a base steel plate.

[Explanation of symbols]

A test piece N mechanical notch S test piece B rib plate C turning welding F load fulcrum

   ─────────────────────────────────────────────────── ─── Continued front page    (72) Inventor Hiroyuki Shirahata             No. 1 Nishinoshu, Oita-shi, Nippon Steel Corporation             Company Oita Works F-term (reference) 4K032 AA01 AA02 AA04 AA05 AA08                       AA11 AA12 AA14 AA15 AA16                       AA17 AA19 AA20 AA21 AA22                       AA23 AA24 AA27 AA29 AA31                       AA32 AA33 AA35 AA36 AA37                       AA39 AA40 BA01 CA01 CA02                       CA03 CB01 CB02 CC02 CC03                       CC04 CD02 CF03

Claims (10)

[Claims] 1. In mass%, C: 0.04 to 0.3%,
Si: 0.01-2%, Mn: 0.1-3%, Al:
0.001-0.1%, N: 0.001-0.01%
As an impurity, P: 0.02% or less, S:
0.01% or less is contained, the balance is composed of iron and unavoidable impurities, and has a structure containing at least ferrite and a hard second phase, and in the cross-sectional structure parallel to the surface, the hard second phase is In the thick steel sheet satisfying all of the above conditions, the hard second phase structure is composed of either bainite, martensite, or a mixed structure of both, and is a steel plate excellent in fatigue strength. Fraction of hard second phase: 20-80% Average Vickers hardness of hard second phase: 250-800 Average circle equivalent diameter of hard second phase: 10-200 m Maximum spacing between hard second phases: 500 m or less 2. Further, in mass%, Ni: 0.01 to 6
%, Cu: 0.01 to 1.5%, Cr: 0.01 to 2
%, Mo: 0.01 to 2%, W: 0.01 to 2%, T
i: 0.003 to 0.1%, V: 0.005 to 0.5
%, Nb: 0.003 to 0.2%, Zr: 0.003 to
0.1%, Ta: 0.005-0.2%, B: 0.0
The thick steel sheet excellent in fatigue strength according to claim 1, containing 002 to 0.005% of one kind or two or more kinds. 3. Further, in mass%, Mg: 0.0005
~ 0.01%, Ca: 0.0005-0.01%, RE
M: 0.005-0.1% of 1 type (s) or 2 or more types are contained, The thick steel plate excellent in the fatigue strength in any one of Claim 1 or 2 characterized by the above-mentioned. 4. A steel slab containing the component according to any one of claims 1 to 3 and having a casting thickness of 100 mm or less is converted into AC ₃
Reheating to the transformation point to 1250 ° C., hot rolling with a reduction ratio of 2 or more, and after hot rolling at a cooling rate of 0.1 to 2 ° C./s until the temperature at which the ferrite fraction becomes 10% or more. After cooling, it is further rapidly cooled to 500 ° C. or less at 5 to 100 ° C./s, and a method for manufacturing a thick steel sheet having excellent fatigue strength. 5. The excellent fatigue strength according to claim 4, wherein before the hot rolling, the steel piece is subjected to diffusion heat treatment at a heating temperature of 1150 to 1300 ° C. and a holding time of 1 to 100 hours. Method for manufacturing thick steel plate. 6. A steel slab containing the component according to claim 1, which has a casting thickness of more than 100 mm, has a heating temperature of 1150 to 1300 ° C. and a holding time before hot rolling. After performing the diffusion heat treatment for 1 to 100 hours, the AC ₃ transformation point
After reheating to 1250 ° C., hot rolling with a reduction ratio of 2 or more, and after hot rolling, after cooling at a cooling rate of 0.1 to 2 ° C./s to a temperature at which the ferrite fraction is 10% or more. And a method for producing a thick steel sheet having excellent fatigue strength, further comprising rapidly cooling to 500 ° C. or lower at 5 to 100 ° C./s. 7. In the hot rolling, at least a starting temperature is 850 ° C. or lower and an ending temperature is Ar ₃ transformation point or higher,
The method for manufacturing a thick steel sheet having excellent fatigue strength according to any one of claims 4 to 6, wherein hot rolling including rolling with a cumulative reduction of 30% or more is performed. 8. In the hot rolling, at least the starting temperature is below the Ar ₃ transformation point and the ending temperature is above 600 ° C.,
The method for producing a thick steel sheet having excellent fatigue strength according to any one of claims 4 to 7, wherein hot rolling including rolling with a cumulative reduction of 10 to 80% is performed. 9. After being rapidly cooled to 500 ° C. or lower, after hot rolling is completed, (AC ₁ transformation point + 30 ° C.) to (AC ₃
Reheat to a transformation point of -50 ° C, and keep it at 5-1 below 500 ° C.
The method for producing a thick steel sheet having excellent fatigue strength according to any one of claims 4 to 8, wherein a two-phase heat treatment of cooling at 00 ° C / s is performed. 10. The fatigue strength according to any one of claims 4 to 9, which is characterized by quenching to 500 ° C. or lower, or two-phase region heat treatment, and then tempering at 250 to 500 ° C. Excellent manufacturing method for thick steel plate.