JP3462943B2 - Steel sheet having high fatigue strength at welded portion and method for producing the same - Google Patents

Description

Detailed Description of the Invention

[0001]

TECHNICAL FIELD The present invention relates to construction machinery, shipbuilding,
The present invention relates to a steel plate which is used for a structural member that requires a long fatigue life in a welded structure such as an offshore structure / bridge, and an automobile, and has a long repeated life of fatigue fracture generated from a weld, and a manufacturing method thereof. In particular, as steel plates, thick steel plates, middle plates, and thin plates (hot rolled, cold rolled steel plates) are targeted.

[0002]

2. Description of the Related Art Due to increasing demands for environmental protection and respect for human life, structures are required to have higher reliability than ever before. In a large welded structure used under severe conditions, fatigue fracture, brittle fracture, ductile fracture, and other fractures may occur, but among these, fatigue fracture is a fracture caused by the action of low repetitive stress, It is the one most likely to occur. As a measure to lengthen the fatigue life, at present, it is largely due to design consideration such as increasing the plate thickness so that the load stress generated in the member does not become high, and as a result, the weight reduction of the structure does not progress etc. Has been pointed out.

Up to now, many techniques for improving fatigue strength have been disclosed, but most of them are related to the base metal, and those for the purpose of improving the fatigue strength of a welded portion, which is the object of the present invention, are not. Few. Further, spot welding, which is widely used for thin steel sheets, has a stress state such as stress concentration and residual stress that is very different from that of butt welding, and therefore cannot be applied to the improvement of the fatigue strength of steel sheet butt-welded portions targeted by the present invention.

In JP-A-57-108241, it is possible to improve stretch flangeability and fatigue strength by setting the area ratio of bainite to 5 to 70% and the area ratio of martensite to 1 to 30%. Is listed. Further, Japanese Patent Publication No. 1-45833 describes that the area ratio of bainite can be set to 5 to 60% and fatigue strength can be improved by limiting the cooling rate and the winding temperature of the hot rolled steel sheet.

Further, in Japanese Examined Patent Publication (Kokoku) No. 4-24418, the area ratio of bainite is 5 to 60% and the area ratio of martensite is 1 to 15% in a three-phase mixed structure of ferrite, bainite and martensite. It is described that by doing so, stretch flangeability and fatigue strength can be improved. Further, as a method for improving the fatigue strength of a thick steel plate, a ferrite having an aspect ratio of 4 or more and a minor axis of 10 μm or less can be obtained by rolling in an austenite / ferrite two-phase region as disclosed in JP-A-5-148540. Has been described, and a technique of causing the generation of a crack parallel to the plate surface as the fatigue crack grows and suppressing the propagation of the fatigue crack.

[0006]

Among these, JP-A-57-108241 discloses that the fatigue strength is improved by limiting the area ratio of bainite and martensite to a specific range. The present invention relates to the improvement of fatigue strength of a thin steel sheet base material, and cannot be applied because the stress conditions are completely different such as residual stress occurring in the welded portion targeted by the present invention.

Further, Japanese Patent Publication No. 1-46583 is intended to improve the fatigue strength by controlling the structure of the base material, and the effect is limited to the improvement of the fatigue strength of the welded portion of the steel sheet. Further, JP-B-4-24418 is intended mainly for the purpose of improving stretch-flangeability, and with regard to fatigue strength, suppression of decrease in hardness of flash butt welded portion where stress concentration is low is suppressed. By doing so, the fatigue strength is improved, and the effect cannot be expected to improve the fatigue strength of the welded portion where the stress concentration is high.

Further, Japanese Patent Laid-Open No. 148540/1993 attempts to suppress the propagation of fatigue cracks by generating a separation parallel to the plate surface. There is no effect of suppressing the generation, only the effect is exhibited after the crack generated from the welded portion penetrates into the base metal portion, and there is a limit to the improvement of the fatigue strength of the welded portion.

An object of the present invention is to improve the fatigue strength of welded portions of steel plates. Fatigue cracks are the most stress-concentrated welds, especially the weld heat-affected zone (hereinafter referred to as HAZ).
Arises from. The generated crack propagates in the HAZ, penetrates into the base metal, and further propagates. Therefore, in order to improve the fatigue strength of the welded portion, it is necessary to control both crack initiation / propagation in the HAZ and crack propagation in the base metal portion. The present invention suppresses the generation and propagation of cracks by controlling the structure of the HAZ and the base metal part, and improves the fatigue strength of the welded member.

[0010]

DISCLOSURE OF THE INVENTION As a result of microscopically observing the morphology of fatigue crack initiation and propagation in welds, the present inventors have found that in order to improve the fatigue strength of welds, HAZ
It was found that the fatigue strength of the welded part can be remarkably improved by the synergistic effect of suppressing fatigue crack initiation and propagation in, and suppressing crack propagation after the crack penetrates into the base metal part. That is, HA is used to suppress crack initiation and propagation in HAZ.
It is effective to increase the ferrite area ratio of Z,
It was newly found that texture control is effective in suppressing crack propagation in the base metal part.

The present invention was made on the basis of the above findings, and the gist thereof is as follows. (1) In weight%, 0.015 ≦ C ≦ 0.10, 0.05 ≦ Si ≦ 2.0, 0.1 ≦ Mn ≦ 1.5, P ≦ 0.05, S ≦ 0.02 The balance of Fe and inevitable impurities, the value of Ceq (f) shown in the following formula satisfies Ceq (f) ≦ 0.11, and the (200) diffraction intensity ratio in the plate thickness direction measured by X-ray. Is 2.0 to 15.0, and the area ratio of the recovered or recrystallized ferrite is 15 to 80%.

However, Ceq (f) = C-Si / 57 + Mn / 13 + (Cu +
Ni) / 26 + Cr / 5 + Mo / 6 + V / 5 + Nb /
1.5 (2)% by Weight, 0.1 ≦ Cu ≦ 2.0, 0.1 ≦ Ni ≦ 2.0, 0.05 ≦ Cr ≦ 0.5, 0.05 ≦ Mo ≦ 0.5, 0.005 ≦ Nb ≦ 0.10, 0.005 ≦ V ≦ 0.10 One kind or two or more kinds of welded parts according to the preceding paragraph (1) are contained. Steel plate with high fatigue strength.

(3) 0.005≤Ti≤0.05, 0.002≤N≤0.015 by weight%, and Ti / N is 2.0 to 3.4. The steel sheet having high fatigue strength of the welded part according to (1) or (2) above.

(4) One or two of 0.0005 ≦ REM ≦ 0.0050 and 0.0005 ≦ Ca ≦ 0.0050 are contained in a weight percentage, and the above-mentioned (1) to (3) are included. The steel plate with high fatigue strength of the welded part of any one of 1).

(5) A steel ingot having the chemical composition described in any one of the above items (1) to (4) is converted into an Ac ₃ transformation point to 130.
It is heated to 0 ° C., rolled in a recrystallization temperature range with a cumulative reduction of 20 to 90%, and subsequently rolled with a cumulative reduction of 10 to 80% in an unrecrystallized temperature range of the Ar ₃ transformation point or higher. ₃
A method for producing a steel sheet having a high fatigue strength of a welded portion, which comprises finish rolling at a cumulative reduction of 40 to 90% at a transformation point or lower and 600 ° C. or higher, and then allowing to cool to room temperature in the atmosphere after rolling.

(6) A steel ingot having the chemical composition described in any one of the above items (1) to (4) is converted into an Ac ₃ transformation point to 130.
It is heated to 0 ° C., rolled in a recrystallization temperature range with a cumulative reduction of 20 to 90%, and subsequently rolled with a cumulative reduction of 10 to 80% in an unrecrystallized temperature range of the Ar ₃ transformation point or higher. ₃
Finish rolling is performed at a cumulative rolling reduction of 40 to 90% below the transformation point at 600 ° C. or more, and after the rolling is finished, it is left to cool in the atmosphere for 30 to 300 seconds, and then at room temperature at a cooling rate of 5 to 100 ° C./sec. A method for producing a steel sheet having a high fatigue strength of a weld, which comprises controlled cooling to 600 ° C.

(7) A steel ingot having the chemical composition described in any one of the above items (1) to (4) is converted into an Ac ₃ transformation point to 130.
It is heated to 0 ° C., rolled in a recrystallization temperature range with a cumulative reduction of 20 to 90%, and subsequently rolled with a cumulative reduction of 10 to 80% in an unrecrystallized temperature range of the Ar ₃ transformation point or higher. ₃
After finish rolling at a temperature below the transformation point and at a cumulative rolling reduction of 40 to 90% at a temperature of 600 ° C. or higher, immediately control-cool to room temperature to 600 ° C. at a cooling rate of 5 to 100 ° C./second, and subsequently cool to room temperature in the atmosphere. Then, after further heating to 500 ° C. or higher and the Ac ₁ transformation point or lower, it is allowed to cool in the atmosphere, and a method for producing a steel sheet having high fatigue strength of a welded part.

Here, the term "recovered or recrystallized ferrite grains" is defined as a ferrite grain in which no slip band is observed in the grain by observing an optical microscope structure at a magnification of 500, and a ferrite in which no slip band is observed in the grain. The area ratio occupied by is measured and is defined as the area ratio of recovered or recrystallized grains.

[0019]

[Function] Fatigue fracture consists of crack initiation and propagation. The sum of the crack initiation life and the crack propagation life is the total life leading to fatigue failure. In the welded portion, the crack initiation often occurs from the HAZ that coincides with the weld toe where the stress concentration is most severe. The generated crack propagates through the HAZ and then rushes into the base metal portion, and further propagates to finally break the member. In order to improve the fatigue fracture life of the weld, it is necessary to suppress crack initiation / propagation in the HAZ and propagation in the base metal. Obviously, it is more effective to suppress both of them at the same time than to suppress only one of them.

The present inventors first conducted a systematic experiment on the effect of the structure on the fatigue strength of HAZ, and obtained extremely useful knowledge. That is, the HAZ structure was reproduced with a thermal cycler and a test piece was subjected to a fatigue test.
When the influence of the structure was investigated, it was found that the ratio of the fatigue limit stress to the tensile strength (hereinafter, referred to as the fatigue limit ratio) was improved as the high temperature transformation structure was obtained.

FIG. 1 shows the test results. Various steels with different alloy element contents were subjected to a simulated welding heat cycle with a maximum heating temperature of 1400 ° C, and three-point bending test pieces with a notch with a stress concentration factor of 2.6 were machined from this reproduced HAZ material. Then, it was subjected to a fatigue test. The ferrite structure fraction of the reproduced HAZ material is plotted on the horizontal axis, and the fatigue limit ratio is plotted on the vertical axis. It has been clarified that the fatigue crack initiation and propagation can be suppressed by increasing the ferrite fraction in the HAZ structure. In general welded structural mild steel and high-strength steel, when low heat input welding with a heat input of about 1.0 to 2.0 kJ / mm is performed, HAZ becomes a structure mainly composed of bainite and martensite. Therefore, the structure becomes unfavorable from the viewpoint of fatigue crack initiation and propagation in the HAZ. On the other hand, H
When the area ratio of ferrite structure of AZ is 60% or more, H
The fatigue limit ratio of AZ increases, and the fatigue life also increases accordingly.

The reason why the fatigue limit ratio becomes higher as the ferrite structure fraction of the HAZ becomes higher is not clear. However, the softer the structure, the more prominent the crack closure becomes, and the microcrack propagation is delayed. On the contrary, the dislocation density becomes higher. It is considered that, in the bainite-martensite structure, dislocation rearrangement occurs due to repeated deformation and dislocation strengthening is invalidated, resulting in a low fatigue limit ratio.

The present inventors have further investigated the chemical composition of steel and HA.
As a result of detailed examination of the relationship of the Z structure, the results shown in FIG. 2 were obtained. That is, HAZ is represented by the following carbon equivalent formula.
The ferrite area ratio can be expressed. Ceq (f) = C-Si / 57 + Mn / 13 + (Cu +
Ni) / 26 + Cr / 5 + Mo / 6 + V / 5 + Nb /
1.5 Here, the welding heat input was set to 1.7 kJ / mm, and the structure of the coarse grain area HAZ in the vicinity of the welding fusion line (hereinafter referred to as FL) was set to 20.
The area ratio of the ferrite structure was determined by observing with a 0 × optical microscope. As is clear from FIG. 2, Ceq (f) is set to 0.
When it is 11 or less, the ferrite area ratio of HAZ can be 60% or more.

Next, as a result of detailed experiments on the relationship between the base metal structure and the fatigue crack propagation behavior, it was discovered that there is a close relationship between the texture and the fatigue crack propagation. The present inventors have focused on the plastic deformation at the propagating fatigue crack tip, and based on the idea that the crack propagation rate can be reduced by suppressing the plastic deformation at the crack tip, the plasticity at the fatigue crack tip is reduced. Various studies were made on the method of suppressing deformation. As a result, it was clarified that there is a close relationship between the crystal orientation of ferrite and plastic deformation of the crack tip.

In general, the slip deformation in the crystal of a ferrite having a bcc crystal structure is such that the slip plane is the {110} plane and the slip direction is the <111> direction. Twelve slip systems are determined by the combination of slip surface and slip direction. Among many slip systems, the slip system in which the shear stress determined by the slip plane and slip direction is the highest due to the relationship between the crystal orientation and the stress field at the crack tip (called the main slip system)
At this point, dislocations are most active and slip deformation occurs in the slip system. When the shear stress of the main slip system is much larger than that of the second or lower slip system, only the main slip system is active and there is no interference due to deformation in the other slip systems. Deformation easily occurs, and as a result, plastic deformation at the crack tip is facilitated, there are few obstacles for crack growth, and crack propagation is easy.

On the other hand, when the values of the shear stresses of the main slip system and the slip system of the second order or lower are equal to or small, the slip deformation is active in the slip system of the second order or lower together with the main slip system. become. In this case, since slip deformations tend to occur on different slip planes, interference between dislocations frequently occurs, and slip deformation becomes extremely difficult. As a result, plastic deformation at the crack tip is significantly suppressed, and crack propagation is suppressed.

From the above viewpoint, the relationship between the crystal orientation of ferrite and fatigue crack propagation was examined in more detail. As described above, the fatigue crack of the welded member occurs from the weld toe,
After propagating through the HAZ, it penetrates into the base metal, and the crack propagation direction in this case is the plate thickness direction. Therefore, when examining the relationship between the ferrite crystal orientation and the crack propagation velocity, it is assumed that the crack propagation direction is the plate thickness direction. As shown in FIG. 3 (a), when a crack propagates in the plate thickness direction in a crystal grain in which the {100} plane of ferrite has an azimuth relationship perpendicular to the plate thickness direction, the main slip system and the secondary slip system It was clarified that the difference in shear stress between slip systems below became smaller and slip interference occurred, and crack propagation was most suppressed. This is schematically shown in FIG. On the other hand, when a crack propagates in a crystal having a crystal orientation other than the above, for example, a {111} plane having an orientation perpendicular to the plate thickness direction (illustrated in FIG. 3 (b)), shear of the main slip system is used. The stress is predominantly larger than that of a slip system of second order or less, and slip deformation is active only in the main slip system. This is schematically shown in FIG. In this case, the crack propagation speed is high because plastic deformation easily occurs at the crack tip.

Based on the above basic new knowledge, the relationship between the steel sheet texture and the fatigue crack propagation rate in the sheet thickness direction was examined. The laboratory vacuum-melted steel ingot whose chemical composition is shown in Table 1 is 115
After heating to 0 ° C., rolling in the recrystallized region and unrecrystallized region, and further rolling in the austenite / ferrite two-phase region or the ferrite single-phase region below the Ar ₃ transformation temperature, the texture of 20 mm thick was changed. A steel plate was created. From this, a test piece having a plate thickness of 10 mm, a width of 18 mm, a length of 100 mm, and a notch depth of 5 mm was processed. Here, the plate thickness direction of the steel plate was made to coincide with the crack propagation direction, and the longitudinal direction of the test piece was made parallel to the rolling direction. A crack was propagated by repeatedly applying a load to the test piece under the condition that the ratio of the minimum load and the maximum load was 0.1. Stress intensity factor range ΔK is 70 kgf · mm
The crack propagation rate da / dN came in ^-3/2 was measured. On the other hand, the (200) diffraction intensity in the plate thickness direction was measured by X-ray, and the ratio to the (200) diffraction intensity of the comparative material having random orientation was determined. Here, the (200) diffraction intensity ratio corresponds to the existence probability of the crystal grains having the above-described {100} plane having an orientation relationship perpendicular to the plate thickness direction. FIG. 5 shows the result of plotting da / dN against the (200) diffraction intensity ratio. It was found that the crack propagation speed in the plate thickness direction decreased when the (200) diffraction intensity ratio was 2.0 or more. The crack propagation speed tends to decrease as the diffraction intensity ratio increases, but extremely strong low-temperature rolling is required to achieve the (200) diffraction intensity ratio of 15.0 or higher. It will be extremely difficult. (200) Diffraction intensity ratio of 2.0
Although the effect of lowering the crack propagation speed in the plate thickness direction can be obtained with the above, it is preferably 4.0 or more, and considering the capability of the rolling mill, 15.0 or less, that is, 4.0 to 4.0. It is desirable to set it to 15.0.

[0029]

[Table 1]

As described above, it has been clarified that the crack propagation can be controlled by controlling the texture. In addition to this, even the ferrite having the same crystal orientation has a lower dislocation density and is softer than the above. It was also revealed that the effect was more prominent. This is presumably because even if the ferrite is oriented so that interference of slip deformation is likely to occur, if dislocations are introduced in advance, the dislocations will easily move, and slip suppression will be less likely to occur. Furthermore, crack propagation is also affected by crack opening and closing behavior, and crack closure is more pronounced in soft tissues, and crack propagation slows down because the effective stress intensity factor range is reduced. In the present invention, rolling at a low temperature is necessary to develop the (200) diffraction intensity ratio, but when the low temperature rolling is performed, ferrite having a high dislocation density is generated. Even if such low temperature rolling is carried out, it is effective to utilize recovery or recrystallization in order to generate soft ferrite having a low dislocation density. If the dislocation density is reduced and the high (200) diffraction intensity ratio is maintained in both the recovered ferrite and the recrystallized ferrite, the crack propagation suppressing effect is similarly exhibited.

Controlling the crystal orientation of ferrite,
The crack propagation of the base metal is suppressed by the synergistic effect of controlling the recovery and recrystallization of ferrite. The present invention has been constructed based on the above new findings. In order to realize the above findings, the following limitations are necessary. C is effective in increasing the strength of the base material. If it is less than 0.015%, the strength as a steel sheet cannot be secured, so the lower limit is 0.015.
%. On the other hand, if the content exceeds 0.10%, Ar
_{(3) The} transformation temperature is remarkably lowered, the rolling temperature is lowered, and the rolling load is increased, so that the rolling becomes extremely difficult, and the pearlite fraction is increased and the fatigue crack propagation suppressing effect is lowered,
Further, the toughness is significantly reduced, so the upper limit of C is set to 0.1.
It was set to 0%.

Si is not only effective in increasing the strength of the base material, but is an important element as a deoxidizing element. If it is less than 0.05%, strength increase cannot be obtained, deoxidation is weak, and inclusions increase, which easily becomes the starting point of fatigue fracture. Therefore, the lower limit of Si is set to 0.05%. Further, Si raises the transformation temperature to raise the ferrite structure fraction of HAZ. If it is less than 0.05%, this effect is not remarkable.
To increase the strength of the base metal and the fraction of HAZ ferrite, S
It is desirable to increase the i content, but if the content exceeds 2.0%, the toughness is significantly reduced. Therefore, the upper limit of Si is set to 2.0%, preferably 1.0%.

Mn has the effect of increasing the strength of the base material. If it is less than 0.1%, the strength increasing effect cannot be obtained, so
The lower limit of Mn was 0.1%. On the other hand, if the content of Al exceeds 1.5%, the Ar ₃ transformation temperature becomes too low and rolling becomes difficult, and in addition, the toughness decreases significantly, so the upper limit of Mn was made 1.5%. Since P is an impurity element and easily causes grain boundary destruction, a lower P content is preferable. When the content exceeds 0.05%, the toughness is significantly reduced due to grain boundary fracture, so the upper limit of P was set to 0.05%.

Since S produces MnS to reduce ductility, particularly elongation in the sheet thickness direction, and causes a large variation in fatigue strength as a starting point of fatigue fracture, it is preferably low. This effect becomes remarkable when the content exceeds 0.02%, so the upper limit of S is set to 0.02%. Cu, Ni, Cr, Mo, Nb, V, Ti, N, R selectively contained
The content of EM and Ca is limited for the following reasons.

Cu is an element effective in increasing the strength of the base material by strengthening the solid solution and increasing the hardenability. If it is less than 0.1%, this effect is not remarkable, so the lower limit of Cu is set to 0.1%. On the other hand, if added in excess of 2.0%, the Ar ₃ transformation temperature will be too low and rolling will be difficult, and in addition, the toughness will be significantly reduced, so the upper limit of Cu was made 2.0%. Ni has the effect of improving the toughness as well as increasing the strength of the base material by increasing the hardenability. Since this effect is not remarkable when it is less than 0.1%, the lower limit of Ni is set to 0.1%. Conversely, 2.0
%, The Ar ₃ transformation temperature becomes too low and rolling becomes difficult, so the upper limit of Ni was made 2.0%.

Cr has the effect of increasing the strength of the base material by increasing the hardenability. If less than 0.05%, this effect is not remarkable, so the lower limit of Cr was made 0.05%. vice versa,
If added in excess of 0.5%, the Ar ₃ transformation temperature will be too low and rolling will be difficult, and in addition, the toughness will be significantly reduced, so the upper limit of Cr was made 0.5%. Mo has the effect of increasing the hardenability and increasing the strength of the base material. 0.05
If it is less than%, this effect is not remarkable, so the lower limit of Mo is set to 0.05%. On the other hand, if it is added in excess of 0.5%, the deformation resistance at high temperature increases and rolling becomes difficult, and in addition, the toughness deteriorates significantly. Therefore, the upper limit of Mo is 0.5.
%.

Nb has the effect of increasing the strength of the base material by increasing hardenability and precipitation hardening. If less than 0.005%, this effect is not remarkable, so the lower limit of Nb is set to 0.005%.
%. On the contrary, if the content of Ni exceeds 0.10%, a large amount of precipitates are formed and the toughness is remarkably deteriorated.
Was set to 0.10%. V shows the effect of increasing the strength of the base material by increasing the hardenability and precipitation hardening. 0.0
If it is less than 05%, this effect is not remarkable, so the lower limit of V is made 0.005%. On the other hand, if the content exceeds 0.10%, a large amount of precipitates are formed and the toughness is significantly reduced, so the upper limit of V was made 0.10%.

Ti produces TiN, which suppresses the growth of austenite grains during slab heating prior to rolling, and is effective in refining ferrite grains after rolling. A smaller ferrite grain size is less likely to cause slip deformation at the crack tip and is effective in suppressing crack propagation. If less than 0.005%, this effect is not significant, so the lower limit of Ti is 0.00
It was set to 5%. On the other hand, if the content exceeds 0.05%, a large amount of precipitates are formed and the toughness is remarkably reduced.
The upper limit of i was set to 0.05%.

N is TiN when added in combination with Ti.
To produce the above effect. N content is 0.002%
If it is less than 1.0, this effect is not remarkable, so the lower limit is set to 0.00
It was set to 2%. On the contrary, when the content exceeds 0.015%,
Since it forms a solid solution in ferrite and causes deterioration in toughness, the upper limit of N is set to 0.015%. The purpose of N addition is to produce TiN. Therefore, the Ti / N ratio is 2.0 to
It must be in the range of 3.4. When the Ti / N ratio is less than 2.0, the amount of N in ferrite increases due to excess N. On the contrary, when the Ti / N ratio exceeds 3.4, the amount of Ti carbide produced increases due to excess Ti. Therefore, outside this range, the toughness is significantly reduced.

REM fixes S to suppress the formation of MnS, and is effective in improving ductility and reducing variations in fatigue strength. As REM, both lanthanoid series and actinoid series show similar effects, but typical ones are lanthanoid series La and Ce. If less than 0.0005%, this effect is not significant, so the lower limit of REM is 0.0005%.
And On the contrary, if it exceeds 0.0050%, the REM becomes coarse.
Ductility is reduced due to the formation of oxides and sulfides, and this also serves as the starting point for fatigue cracks, increasing the variation in fatigue strength. Therefore, the upper limit of REM is set to 0.0050%.

Ca is MnS by fixing S like REM.
It suppresses the generation and shows the effect of improving the ductility and reducing the variation in fatigue strength. If less than 0.0005%, this effect is not remarkable, so the lower limit of Ca was made 0.0005%. On the other hand, if it exceeds 0.0050%, coarse Ca oxides and sulfides are formed, the ductility is lowered, and it becomes the starting point of fatigue cracks and increases the variation of fatigue strength. Therefore, the upper limit of Ca is set to 0.0050%.

After limiting each of the above components, the value of Ceq (f) expressed by the following equation must be 0.11 or less.
This is because the HAZ ferrite area ratio is 60% in this range.
This is because the effect of suppressing the fatigue crack initiation / propagation of HAZ becomes remarkable. Ceq (f) = C-Si / 57 + Mn / 13 + (Cu +
Ni) / 26 + Cr / 5 + Mo / 6 + V / 5 + Nb /
1.5 Next, the limitation necessary for suppressing the fatigue crack propagation of the base material will be described.

As described above, the (200) diffraction intensity ratio in the plate thickness direction measured by X-ray must be 2.0 to 15.0. When the (200) diffraction intensity ratio in the plate thickness direction is less than 2.0, the development of texture is insufficient and the effect of suppressing propagation is insufficient. Conversely, the (200) diffraction intensity ratio in the plate thickness direction is 1
If it exceeds 5.0, it is necessary to carry out strong rolling at a low temperature,
Since it becomes substantially impossible to perform plate rolling, the upper limit was set to 15.0.

Further, the area ratio of the recovered or recrystallized ferrite must be in the range of 15-80%. If the area ratio of recovered / recrystallized ferrite is less than 15%, the density of mobile dislocations in the ferrite is high, and the effect of suppressing slip deformation becomes weak even in the crystal orientation that is advantageous for suppressing crack propagation. Conversely, 80
%, Ferrite grains in an undesired direction grow, and the plate thickness direction (200) diffraction intensity ratio inevitably decreases.
The crack propagation suppression effect is reduced. Therefore, the upper limit of the area ratio of recovered / recrystallized ferrite is set to 80%.

Next, the reasons for limiting the rolling, cooling and heat treatment conditions of the steel sheet will be described below. Prior to hot rolling, the steel ingot needs to be austenitized to 100%, and for this purpose, the temperature of the steel ingot needs to be heated to the Ac ₃ transformation temperature or higher. However, if heated above 1300 ° C., austenite grains are significantly coarsened, and fine-grained ferrite cannot be obtained after rolling. Therefore, the upper limit of the heating temperature is 130.
Set to 0 ° C.

Since the austenite grains are coarsened by heating the steel ingot, it is necessary to reduce the austenite grain size before rolling in the non-recrystallization temperature region by rolling in the recrystallization temperature region. Cumulative rolling reduction of recrystallization temperature range rolling is 20%
If it is less than the above, recrystallization does not proceed sufficiently and the recrystallized grains are not sufficiently small. Therefore, the lower limit of the cumulative rolling reduction in this temperature range is set to 20%. Further, in order to make the recrystallized grains finer, it is desirable to increase the cumulative reduction ratio, but if it is too large, the reduction in the subsequent rolling at a low temperature cannot be secured. Therefore, the upper limit of the cumulative rolling reduction in this temperature range is set to 90%.

In the non-recrystallization temperature range rolling subsequent to the recrystallization temperature range rolling, a deformed body is introduced into austenite to refine ferrite grains after transformation. This effect is not significant when the cumulative rolling reduction in the non-recrystallization temperature range rolling is less than 10%, so the lower limit was made 10%. It is preferable that the cumulative rolling reduction itself in the non-recrystallization temperature region rolling is high in order to refine the crystal grains, but if it is too high, it is not possible to secure the subsequent rolling below the Ar ₃ transformation temperature. Therefore, the upper limit of the cumulative rolling reduction in this temperature range is set to 80%.

In the present invention, it is necessary to increase the (200) diffraction intensity ratio in the plate thickness direction. For this reason, finish rolling at an Ar ₃ transformation temperature or lower plays an extremely important role and is an essential step in the present invention. Is. If the temperature is below the Ar ₃ transformation temperature, even in the austenite-ferrite two-phase region,
Alternatively, it may be in the ferrite single phase region. A lower rolling temperature is desirable only from the viewpoint of increasing the (200) diffraction intensity ratio in the plate thickness direction, but the lower the temperature, the higher the deformation resistance, and the higher the rolling load and the more difficult the rolling.
Further, in the case where recovery or recrystallization is caused by cooling the steel sheet in the air following the rolling (claims 5 and 6), if the finishing temperature is less than 600 ° C., recovery / recrystallization is less likely to occur. Therefore, the lower limit of the rolling finishing temperature is set to 600 ° C.

The higher the cumulative rolling reduction below the Ar ₃ transformation temperature, the higher the (200) diffraction intensity ratio in the sheet thickness direction, which is desirable from the viewpoint of suppressing fatigue crack propagation. If the cumulative rolling reduction in this temperature range is less than 40%, the (200) diffraction intensity ratio in the plate thickness direction does not become sufficiently high, so the lower limit was made 40%. On the contrary, in order to make the cumulative reduction rate in this temperature range to be more than 90%, strong rolling at a low temperature is necessary, and substantially thick plate rolling becomes impossible. Therefore, the upper limit of the cumulative rolling reduction in this temperature range is set to 90%.

After finish rolling below the Ar ₃ transformation temperature, the dislocation density in the ferrite must be reduced by allowing the steel sheet to cool in the air to recover or recrystallize. When cooling to room temperature in the air, recovery / recrystallization proceeds after rolling, and therefore no particular control is required. However, since the cooling rate is low, recovery / recrystallization may proceed too much, so it is desirable to carry out controlled rolling. The process of allowing to cool to room temperature in the atmosphere is particularly effective when manufacturing a steel plate having a thin plate thickness.

When the controlled cooling is performed after the finish rolling to increase the strength of the base material, it is necessary to secure the cooling time after the finish rolling until the start of the controlled cooling and to promote the recovery / recrystallization. If the cooling time is less than 30 seconds, the fraction of ferrite with a low dislocation density does not exceed 15%, so the lower limit is set to 30
Seconds The longer the cooling time, the more the recovery and recrystallization proceed, and the fraction of ferrite with a low dislocation density increases. However, if the cooling time exceeds 300 seconds, recrystallization proceeds too much, the (200) diffraction intensity ratio in the plate thickness direction decreases, and the fatigue crack propagation suppressing effect decreases. Therefore, the upper limit of the cooling time is set to 300 seconds.

Controlled cooling is necessary in order to freeze the ferrite structure after cooling in the atmosphere to cause recovery and recrystallization and to reduce the grain size of ferrite formed by transformation from untransformed austenite. is there. If the cooling rate of controlled cooling is less than 5 ° C / sec, this effect cannot be obtained.
The lower limit was 5 ° C / sec. Conversely, this cooling rate is 100 ℃
If it exceeds / sec, untransformed austenite transforms into bainite or martensite, and the effect of suppressing fatigue crack propagation is weakened, so the upper limit was made 100 ° C / sec.

Further, it is necessary to perform controlled cooling to a temperature at which the transformation is completely completed and recovery / recrystallization does not occur. Therefore, the upper limit of the cooling stop temperature is set to 600 ° C. Even if the cooling stop temperature is kept below room temperature by cooling with a special cooling medium, transformation, recovery and recrystallization do not occur after the cooling is stopped, but in normal controlled cooling the temperature is almost the same as room temperature. Is generally used, and in this case, the lower limit of the cooling stop temperature is room temperature. Therefore, the lower limit of the cooling stop temperature is set to room temperature.

In the method of claim 7, it is necessary to cool to room temperature after the controlled cooling. In particular, even if the control cooling stop temperature is 600 ° C. or lower, if the temperature is high, recovery may proceed before the subsequent heat treatment. Therefore, it is necessary to once cool to room temperature. As described above, in order to suppress fatigue crack propagation, it is necessary to set the fraction of ferrite having a low dislocation density to 15 to 80% by recovery / recrystallization. For this purpose, it is possible to control the cooling immediately after the completion of rolling and then heat the steel sheet to set the fraction of the recovered / recrystallized ferrite to 15 to 80%. Therefore, the heating temperature needs to be 500 ° C. or higher. Further, since no transformation should occur during heating, the heating temperature must be below the Ac ₁ transformation temperature.

The effect of improving the fatigue life of the welded portion by the synergistic effect of the HAZ structure control and the base metal structure control described above is as follows.
It exhibits a particularly remarkable effect in thick steel plate welded portions. INDUSTRIAL APPLICABILITY The present invention exerts an effect irrespective of the deoxidizing method of molten steel, and steel deoxidized with elements such as Si, Al and Ti can be used.

[0056]

DETAILED DESCRIPTION OF THE INVENTION

[0057]

EXAMPLES Examples of the present invention will be described below. Steel was melted by a converter in a factory and cast into a 240 mm thick slab by continuous casting. Tables 2 and 3 (continued from Table 2) show the chemical components and Ceq (f) values of the steels of the present invention and comparative steels.
Steels numbered 17 and 18 have chemical compositions within the scope of the invention, but manufacturing conditions are outside the scope of the invention.

Tables 4 and 5 (continued from Table 4) show the manufacturing conditions for steel materials. A steel plate having a plate thickness of 15 to 25 mm was manufactured by the manufacturing method corresponding to claims 5, 6, and 7. Numbers 17 and 18
The cumulative rolling reduction below the Ar ₃ transformation temperature is outside the range of the present invention. Tables 6 and 7 (continued from Table 6) show the tensile properties, Charpy impact properties, texture strength, and recovered / recrystallized ferrite fraction of the base material. Furthermore, a T-shaped fillet welded joint was created by carbon dioxide welding with a heat input of 1.7 kJ / mm, and HA near FL was used.
The ferrite fraction of Z was measured. The results are also shown in Table 7
Shown in.

Here, the steel plates were set to 15 mm. Further, a fatigue test piece shown in FIG. 6 was prepared from this welded joint,
It was subjected to a fatigue test. The fracture life of the test piece was measured when the bending stress range at the weld toe position was 25 kgf / mm ² . A strain gauge is affixed at a position 5 mm away from the toe, and the point at which the strain gauge output drops by 5% from the initial value is defined as the crack initiation life. It was the life. The crack initiation life according to this definition roughly corresponds to the crack initiation and propagation in the HAZ, and the crack propagation life corresponds to the propagation after the crack rushes into the base metal.

[0060]

[Table 2]

[0061]

[Table 3]

[0062]

[Table 4]

[0063]

[Table 5]

[0064]

[Table 6]

[0065]

[Table 7]

The steels 1 , 2, 4 to 16 of the present invention were produced according to the present invention, and the (200) diffraction intensity ratio in the plate thickness direction was 2.0 or more. Comparative steels 17 and 1 whose chemical composition is within the scope of the present invention but whose manufacturing method is outside the scope of the present invention
8. Further, in Comparative Steels 19 to 21 whose chemical composition and manufacturing method are out of the range of the present invention, the (200) diffraction intensity ratio in the plate thickness direction is lower than that of the present invention steel. Further, the HAZ ferrite fraction is the present invention steels 1 , 2, 4 to 1 having chemical components within the present invention range.
6 and comparative steels 17 and 18 are more than 60%.

The crack initiation life of a T-shaped fillet welded joint is HA
Inventive steels 1 , 2, 4 to 16 having a high Z ferrite fraction and comparative steels 17 and 18 are long. Further, since the fatigue limit mainly depends on the fatigue characteristics of HAZ and not on the fatigue characteristics of the base material,
Inventive Steels 1 to 16 and Comparative Steels 17 and 18 having high HAZ ferrite fractions are high. On the other hand, the crack propagation life is as high as (200) diffraction intensity ratio of 2.0 or more in the plate thickness direction, and recovery /
The recrystallized ferrite fraction of the present invention steels 1 , 2, 4 to 16 having a range of 15 to 80% is long. The comparative steels 17 and 18 have a relatively high fatigue limit, but the fatigue life is not significantly improved because the propagation control of the base metal is not performed. It was confirmed that the steels 1 , 2, 4 to 16 of the present invention had a significantly long fracture life due to the long life of both crack initiation and crack propagation, and also had a high fatigue limit.

Next, the effect of improving the fatigue strength of the welded portion according to the present invention was confirmed for a thin steel plate having a plate thickness of 3 mm. Table 8,
Table 9 (continued from Table 8) shows the chemical components. Numbers 1 to 3 are steels of the present invention, and number 4 is a comparative steel. Table 10 shows the strength of the base material and the X-ray (200) diffraction intensity ratio in the plate thickness direction. A weld bead was placed on the surface of the steel plate to prepare a welded specimen having stress concentration at the toe. The welding heat input was 3 kJ / cm. Fatigue test piece is processed from the welded part and subjected to repeated bending load of full swing, the stress range is 25 kgf / mm ²
Fatigue life and fatigue limit were measured. The results are also shown in Table 10. Table 10 also shows the ferrite fraction of HAZ. It is clear that the steel of the present invention has improved life and fatigue limit as compared with the comparative steel.

[0069]

[Table 8]

[0070]

[Table 9]

[0071]

[Table 10]

[0072]

As described above, in the steel of the present invention, the HAZ of welded joints has a long crack initiation life, and at the same time, the crack propagation life of the base metal is long. It has become possible to significantly increase the fatigue strength of welded joints for an extremely long time. By using the steel of the present invention, not only can the reliability of the welded structure against fatigue fracture be improved, but it is also possible to reduce the plate thickness and increase the design stress without shortening the fatigue life. Weight reduction is also possible. Therefore, the present invention is extremely effective industrially.

[Brief description of drawings]

FIG. 1 is a diagram showing an influence of a ferrite structure fraction of HAZ on a fatigue limit ratio.

FIG. 2 is a diagram showing a relationship between a ferrite fraction of HAZ and a carbon equivalent value.

FIG. 3 is a schematic diagram showing a crystal orientation of ferrite.

FIG. 4 is a schematic view showing slip deformation of a fatigue crack tip.

FIG. 5 is a diagram showing the influence of the texture on the fatigue crack propagation rate.

FIG. 6 is a view showing the shape of a T-shaped fillet welded joint.

─────────────────────────────────────────────────── ─── Continuation of front page (58) Fields surveyed (Int.Cl. ⁷ , DB name) C22C 38/00-38/60 C21D 8/02

Claims (7)

(57) [Claims] 1. In weight%, 0.015 ≦ C ≦ 0.10, 0.05 ≦ Si ≦ 2.0, 0.1 ≦ Mn ≦ 1.5, P ≦ 0.05, S ≦ 0.02 Containing the balance Fe and unavoidable impurities, the value of Ceq (f) shown in the following formula satisfies Ceq (f) ≦ 0.11, and (200) diffraction in the plate thickness direction measured by X-ray. A steel sheet having a high fatigue strength in a welded portion, which has a strength ratio of 2.0 to 15.0 and an area ratio of recovered or recrystallized ferrite of 15 to 80%. However, Ceq (f) = C-Si / 57 + Mn / 13 + (Cu +
Ni) / 26 + Cr / 5 + Mo / 6 + V / 5 + Nb /
1.5 2. In% by weight, 0.1 ≦ Cu ≦ 2.0, 0.1 ≦ Ni ≦ 2.0, 0.05 ≦ Cr ≦ 0.5, 0.05 ≦ of the base metal strength increasing element group. Fatigue strength of the welded part according to claim 1, characterized in that it contains one or more of Mo ≦ 0.5, 0.005 ≦ Nb ≦ 0.10, 0.005 ≦ V ≦ 0.10. High steel plate. 3. The composition contains 0.005 ≦ Ti ≦ 0.05 and 0.002 ≦ N ≦ 0.015 by weight%, and Ti / N is 2.0 to 3.4. The steel sheet having high fatigue strength of the welded portion according to claim 1 or 2. 4. One or two kinds of 0.0005 ≦ REM ≦ 0.0050 and 0.0005 ≦ Ca ≦ 0.0050 are contained by weight%.
The steel plate with high fatigue strength of the welded part as described in any one of 1-5. 5. A steel ingot having the chemical composition according to claim 1 is heated to an Ac ₃ transformation point to 1300 ° C. and a cumulative rolling reduction of 20 to 90% in a recrystallization temperature range. Rolling, and then 10 at the unrecrystallized temperature range above the Ar ₃ transformation point.
Rolling a cumulative reduction rate of 80%, further Ar ₃ or less transformation point, and finish rolling at 40% to 90% of the cumulative reduction rate at 600 ° C. or higher, characterized by cooling in air to room temperature after rolling welding Method for manufacturing steel sheet with high fatigue strength of the part. 6. A steel ingot having the chemical composition according to claim 1 is heated to an Ac ₃ transformation point to 1300 ° C., and a cumulative rolling reduction of 20 to 90% in a recrystallization temperature range. Rolling, and then 10 at the unrecrystallized temperature range above the Ar ₃ transformation point.
Rolling at a cumulative reduction of -80%, finish rolling at a cumulative rolling reduction of 40 to 90% at 600 ° C or higher at an Ar ₃ transformation point or lower, and after standing to cool in the atmosphere for 30 to 300 seconds,
Then, at room temperature to 600 at a cooling rate of 5 to 100 ° C / sec.
A method for manufacturing a steel sheet having high fatigue strength of a weld, which is characterized by controlled cooling to ℃. 7. A steel ingot having the chemical composition according to claim 1 is heated to an Ac ₃ transformation point to 1300 ° C. and a cumulative rolling reduction of 20 to 90% in a recrystallization temperature range. Rolling, and then 10 at the unrecrystallized temperature range above the Ar ₃ transformation point.
After rolling at a cumulative rolling reduction of -80% and further finish rolling at a cumulative rolling reduction of 40 to 90% below the Ar ₃ transformation point and at 600 ° C or higher, immediately at room temperature at a cooling rate of 5 to 100 ° C / sec.
After controlled cooling to 600 ° C., cooling to room temperature in the air, and heating to 500 ° C. or higher and Ac ₁ transformation point or lower,
A method for producing a steel sheet having high fatigue strength of a weld, which is characterized by allowing to cool in the atmosphere.