JP4926406B2 - Steel sheet with excellent fatigue crack propagation characteristics - Google Patents

Description

The present invention is, architecture fatigue properties are required, shipbuilding, bridges, construction machinery, it relates to a steel plate having excellent fatigue crack propagation properties for use in welded structural members such as marine structures.

  Generally, welded joints using a wide variety of welding methods such as arc welding, plasma welding, laser welding, and electron beam welding are applied to welded structures such as buildings, shipbuilding, bridges, construction machinery, and offshore structures. Has been.

Since these welded joints are subjected to repeated loads due to wind, waves, mechanical vibrations, etc., it is extremely important to improve fatigue strength. Generally, this is a post-welding process as a method for improving fatigue strength. (1) Grinding, (2) TIG dressing, (3) shot peening, and (4) hammer peening are used, but have the following problems.
Here, the grinding and the TIG dressing improve the shape of the weld bead, but the working efficiency is remarkably poor.
Shot peening and hammer peening have an effect of improving fatigue strength, but shot peening requires a huge machine and various utilities.

In addition, hammer peening has a large reaction, the processing result is not stable, and sometimes press formability and fatigue strength are lowered. Hammer peening also has a drawback that it is difficult to use for thin plates because it gives too much plastic deformation.
Furthermore, since grinding and hammer peening are performed on the joints with low frequency machining of several Hz, the unevenness of the processed surface is severe, stress concentrates on the peaks, and repeated stress is applied to the joints. There is a problem that the fatigue strength of the entire joint is lowered because cracks are generated from the concentrated portion.

Further, residual stress is generally introduced into the weld by heat input by welding. The residual stress is one major factor that reduces the fatigue strength at the weld.
Therefore, as another means for improving the fatigue strength, a method is known in which a compressive residual stress is generated in the welded joint portion or a tensile residual stress generated in the welded joint portion is reduced to increase the fatigue strength.
For example, compressive residual stress can be applied by performing shot peening near the weld toe. Here, the shot peening treatment is a technique in which a compression residual stress is applied by hitting a large number of steel balls of less than 1 mm at the site where fatigue cracks start.
It is also known that the weld toe shape can be improved or the tensile residual stress can be reduced by remelting the weld metal by heating.
However, this shot peening process requires a steel ball, and the post-treatment or cost of this steel ball may be a problem. Furthermore, there is a problem that the amount of improvement in fatigue strength varies.

  As described above, it is difficult to employ a technique for improving fatigue strength by post-welding treatment in a welded joint, and even if it can be adopted, the allowance for improving fatigue strength remains at a low level. Therefore, there is a need for a technique that does not require post-weld processing and that can achieve improved fatigue strength of the welded joint as it is.

From this point of view, several steel sheets that suppress the propagation of fatigue cracks have been proposed in order to improve the fatigue strength of the welded joint as it is.
For example, Non-Patent Document 1 discloses that so-called SUF steel in which a superfine grain structure is formed on the surface layer by processing ferrite in the temperature rising process of general shipbuilding steel has an effect of reducing the propagation speed of fatigue cracks. Is disclosed. However, it is difficult to significantly reduce the propagation speed only by reducing the ferrite grain size, and the superfine grain structure formed on the surface layer is largely lost due to the effect of welding heat, improving the fatigue strength of welded joints. Is not fully achieved.

  Further, in Patent Documents 1 to 7, when a mixed structure having hard pearlite, bainite, and martensite as a second phase in a soft ferrite matrix is used, the hard second phase becomes an obstacle to crack propagation, and fatigue cracks are caused. A steel sheet capable of reducing the propagation speed is disclosed. However, for these technologies, the martensite area ratio, aspect ratio (major axis / minor axis), hardness, and ferrite hardness, which are important factors for delaying crack growth, are appropriately adjusted. Therefore, the fatigue crack propagation characteristics are not improved at all, the improvement is insufficient, or the toughness of the steel material is significantly deteriorated.

For example, in Patent Document 1, the martensite fraction is insufficient, and sufficient improvement in fatigue crack propagation characteristics cannot be obtained. In Patent Document 2, when the martensite fraction exceeds 30%, a significant decrease in toughness occurs, and even if the hardness of the hard second phase with respect to ferrite is ensured to be 30% or more, the hardness of the ferrite is 150 or less. When the two-phase hardness is 400 or less, a sufficient improvement effect of fatigue crack propagation characteristics cannot be obtained. Similarly, Patent Document 3 has a martensite fraction exceeding 30%, and the toughness of the steel material is significantly impaired.
In Patent Documents 4 to 7, the ferrite, the hardness of the second phase, the fraction, and the interval between them are not properly controlled, and in the case of bainite having a low hardness of 400 or less, the fraction Even if there are many, toughness deterioration is suppressed, but the propagation suppression effect is small. Further, in the case where the second phase is 400 or more martensite having a high hardness, if the fraction is 30% or more, significant toughness deterioration occurs.

  Further, in Patent Document 8, after a two-phase structure of ferrite and bainite is set, the ratio of the ferrite phase portion, the hardness of the ferrite, the number of phase boundaries between the ferrite and bainite, and the like are specified in a specific range, thereby fatigue. A steel sheet capable of reducing the crack growth rate is disclosed. However, the improvement effect of fatigue crack propagation characteristics is insufficient at the bainite hardness level, and the effect is similarly small even if the hardness of the ferrite is 150 or less.

  Further, Patent Documents 9 to 11 disclose a steel sheet capable of reducing the fatigue crack propagation rate by dispersing the hard phase as a parent phase and the soft phase as a second phase, unlike the ideas listed above. ing. These are intended to absorb the plastic deformation energy necessary for crack growth in the soft phase, thereby promoting crack closing behavior and suppressing crack growth, but there is weld tensile residual stress. In such a welded joint, the crack is easily opened, so that a sufficient fatigue crack propagation characteristic improvement effect cannot be obtained only by the crack closing effect.

  Further, Patent Documents 12 and 13 disclose steel plates that can reduce the fatigue crack propagation rate by securing a recovery or recrystallization ferrite fraction and further developing a specific texture. This is intended to suppress the plastic deformation of the crack tip during crack growth by a specific texture, but sufficient fatigue crack propagation characteristics with a ferrite texture only where the second phase structure is not specified. In addition, since the plastic deformation at the crack tip can be suppressed only in the extremely low ΔK region, the applicable range is extremely narrow.

As described above, in the prior art, appropriate structure control for remarkably suppressing crack propagation has not been achieved, and the fatigue life of a steel plate that can stably reduce the fatigue crack propagation rate and the welded joint The development of steel sheets that can contribute to improvement is eagerly desired.
Japanese Patent Laid-Open No. 06-271985 Japanese Patent Application Laid-Open No. 07-090478 Japanese Patent Application Laid-Open No. 08-073980 Japanese Patent Laid-Open No. 10-168542 JP-A-11-001742 JP 2002-047531 A JP 2003-003229 A Japanese Patent Application Laid-Open No. 08-225882 JP 07-242992 A Japanese Patent Application Laid-Open No. 08-199286 Japanese Patent Laid-Open No. 09-095754 Japanese Patent Application Laid-Open No. 08-199286 Japanese Patent Laid-Open No. 09-095754 Proceedings of the 24th Fatigue Symposium of the Japan Society of Materials Science 1998 "Fatigue properties of super-fine grain steel sheets" (p157-162)

An object of the present invention is to solve the problems of conventional techniques as described above, building, shipbuilding, bridges, construction machinery, marine structures Fatigue Crack Propagation excellent steel plate used in the welding structural members, such as Is to provide. Specifically, the fatigue crack propagation rate da / dN when the stress intensity factor range ΔK with a stress ratio of 0.1 is 20 MPa√m is 10 ⁻⁸ m / cycle or less, and the heat input with the stress ratio of 0.1. There is to provide a steel plate that fatigue life when the welded joint axial force fatigue test 10~30kJ / min satisfies more than twice the conventional steel.

The present invention has been made as a result of intensive studies in order to solve the above-described problems, and the gist thereof is as follows.
(1) In mass%,
C: 0.03-0.2%, Si: 0.01-1.6%,
Mn: 0.5-2%, P: 0.02% or less,
S: 0.02% or less, Al: 0.001-0.1%,
N: 0.001 to 0.008%
Containing the balance Fe and unavoidable impurities, and the microstructure of the base material is a ferrite having a Vickers hardness of 150 or more, the Vickers hardness is 400 to 900, the area ratio is 5 to 30%, and the aspect ratio (Major axis / minor axis) is a layered structure having flat martensite of 3 or more as a second phase, and the average layer spacing in the plate thickness direction of ferrite and martensite is 3 to 50 μm, and the stress ratio is 0. A steel plate having excellent fatigue crack propagation characteristics, wherein a fatigue crack propagation rate da / dN when a stress intensity factor range ΔK of 1 is 20 MPa√m is 10 ⁻⁸ m / cycle or less.
(2) Furthermore, in mass%,
Cu: 0.1 to 2.5%, Ni: 0.1 to 5%,
Cr: 0.01 to 1.5%, Mo: 0.01 to 1.5%,
W: 0.01-1.5%, Ti: 0.001-0.05%,
Nb: 0.005 to 0.2%, Zr: 0.005 to 0.2%,
V: 0.005-0.2%, B: 0.0002-0.005%
1 or 2 types or more, The steel plate excellent in the fatigue crack propagation characteristic as described in said (1) characterized by the above-mentioned.
(3) Furthermore, in mass%,
Mg: 0.0005 to 0.01%, Ca: 0.0005 to 0.01%,
REM: 0.005 to 0.05%
The steel plate excellent in fatigue crack propagation characteristics as described in (1) or (2) above, comprising one or more of the following.

According to the present invention, construction, shipbuilding, bridges, construction machinery, it is possible to provide a superior steel plate fatigue crack propagation properties for use in welded structural members such as marine structures.
Specifically, the fatigue crack propagation rate da / dN when the stress intensity factor range ΔK with a stress ratio of 0.1 is 20 MPa√m is 10 ⁻⁸ m / cycle or less, and the fatigue life of the welded joint is 2 It can be improved more than twice, and there are significant industrially useful effects such as improving the reliability against fatigue fracture of welded steel structures.

In general, it is known that the fatigue crack propagation rate does not depend on the structure and strength of the steel material. However, as a result of intensive investigations, the present inventors have conducted a layer dispersion with ferrite as a parent phase and martensite as a second phase, and further, hardness, area ratio, aspect ratio (major axis / minor axis), each phase It was found that the fatigue crack propagation rate is significantly reduced by controlling the layer spacing in the thickness direction of the steel.
The mechanism by which the fatigue crack propagation rate decreases is due to the change in internal stress around martensite that occurs when martensite transformation occurs during steel sheet rolling cooling, and has the effect of reducing the driving force for crack propagation. Due to this effect, the crack stagnates just above the martensite and cannot easily propagate inside the martensite, and the crack circumvents or branches along the martensite interface. Such a delay due to crack stagnation, an increase in propagation distance due to crack detouring and branching, and a significant crack closing behavior associated with crack detouring and bifurcation enabled a significant decrease in fatigue crack propagation rate. .

The reason for the limited range of the microstructure will be described below.
Among the factors affecting the fatigue crack propagation rate, the most influential factor is the martensite area ratio, and the propagation rate rapidly decreases at 5% or more. This is due to the increase in crack propagation failure due to an increase in the martensite fraction. However, if it exceeds 30%, the toughness deteriorates remarkably, so the range was made 5-30%.

In order to increase internal stress and lower the driving force for crack propagation more effectively, it is necessary to lower the martensitic transformation start temperature. This is because when martensitic transformation occurs at a low temperature, the ferrite acting as a constraint of the transformation is hard, and the internal stress is increased by the reaction force.
The martensitic transformation start temperature decreases as the amount of carbon in the austenite during hot rolling increases. And since the hardness of a martensite becomes so large that there is much quantity which carbon concentrates, in order to make a martensite transformation start temperature into 400 degrees C or less, the hardness of a martensite must be 400 or more.
The reason why the martensitic transformation start temperature must be 400 ° C. or lower is that if it exceeds 400 ° C., the internal stress is relaxed by the thermal contraction after transformation, and the fatigue crack propagation delay effect is reduced. In addition, when the hardness of martensite exceeds 900, it is difficult to secure a martensite fraction of 5% or more, and since martensite may be the starting point and cause brittle fracture, the hardness of martensite is It was set to 400 to 900 Hv.

  Furthermore, as described above, the harder the ferrite, the more constrained it is during martensitic transformation, and the reaction force increases and the internal stress increases. Therefore, the hardness of the ferrite is set to 150 Hv or more.

  The larger the martensite aspect ratio, the more frequently it hits martensite, which is an obstacle to crack growth, and the detour / branch distance increases, which is effective in reducing the fatigue crack propagation rate. When the aspect ratio is less than 3, the detour / branch distance is small even when the crack hits martensite, and the effect of improving crack propagation characteristics is small. Therefore, the martensite aspect ratio (major axis / minor axis) was set to 3 or more.

  It is necessary to disperse the ferrite phase and the martensite phase in layers, and if the layer spacing is smaller than 3 μm, the internal stress introduced at the time of martensite transformation does not work effectively, and it becomes difficult to delay the crack growth. In addition, when the layer interval exceeds 50 μm, the frequency with which the crack hits martensite, that is, the stagnation of the crack and the detour / branch effect becomes small, so the range of the layer interval is set to 3 to 50 μm.

Next, the reason for limiting the range of each alloy element will be described below. In the following,% means mass%.
C is one of the main elements as a component of the present invention, and is contained as an effective component for controlling the martensite fraction and improving the strength of the steel. When it is less than 0.03%, the martensite component is contained. Although it is difficult to secure a rate of 5% or more , the lower limit was set to 0.05% based on the amount of C in Invention Examples 3 and 13 shown in Table 1 of the Examples . In 0.2 percent since lowering the toughness and resistance to weld cracking resistance of the base metal and weld portion was 0.0 5 to 0.2 percent.

In addition to ensuring strength, Si is an essential element as a deoxidizing element, and in order to obtain the effect, addition of 0.01% or more is necessary, and excessive content exceeding 1.6% is a coarse oxide. The upper limit is 0.8% based on the Si amount of Invention Example 11 shown in Table 1 of the Examples , and the amount is 0.01 to 0.8 %. did.

  Mn is an essential element for increasing the strength, but if it is less than 0.5%, the strength of the base material cannot be secured. On the other hand, an excessive content exceeding 2% deteriorates the base metal toughness, the toughness of the welded portion, the weld cracking property, etc. due to grain boundary embrittlement or the like.

  P is an element that affects the toughness of the steel. If it exceeds 0.02%, not only the base metal but also the toughness of HAZ is significantly inhibited.

  S is preferably as low as P, and when it exceeds 0.02%, MnS precipitation becomes prominent, and the HAZ toughness of the base metal is inhibited to reduce the ductility in the thickness direction. did.

  Al is an element effective for deoxidation, austenite grain size reduction, etc., and in order to exhibit the effect, it is necessary to contain 0.001% or more. On the other hand, if it exceeds 0.1% and excessively contained, a coarse oxide is formed and the ductility is extremely deteriorated, so the amount was made 0.001 to 0.1%.

  Since N combines with Al and Ti and effectively works to refine the austenite grains, it contributes to the improvement of the mechanical properties if the amount is small. Moreover, it is impossible to remove N in steel completely industrially, and reducing it more than necessary is not preferable because it places an excessive burden on the manufacturing process. Therefore, it is industrially controllable and the lower limit is set to 0.001% as a range in which the load on the manufacturing process can be tolerated. If excessively contained, the solid solution N increases and the strain aging characteristics deteriorate, so the upper limit was made 0.008%.

  The above is the reason for limiting the basic components of the present invention. In the present invention, Cu, Ni, Cr, Mo, W, Ti, Nb, Zr, V, One or more of B can be contained. The reasons for limiting the components of each element will be described below.

  Cu is an element effective for increasing the strength without reducing toughness, but if it is less than 0.1%, there is no effect, and if it exceeds 2.5%, it tends to cause cracks during heating of the steel slab or during welding. Therefore, the amount is set to 0.1 to 2.5%.

  Ni is an element effective for improving toughness and strength. To obtain the effect, Ni is required to be added in an amount of 0.1% or more. However, excessive addition exceeding 5% saturates the effect. There is a concern that the HAZ toughness and weldability may be deteriorated, and since it is an expensive element, its amount is set to 0.1 to 5% in consideration of economy.

  Cr is required to be 0.01% or more for enhancing the hardenability and ensuring the strength. On the other hand, if it exceeds 1.5%, it is not preferable for the same reason as Ni. Therefore, the amount is set to 0.01 to 1.5%.

  Mo is an element effective for improving hardenability, improving strength, resistance to temper embrittlement, and suppressing recrystallization. To obtain the effect, addition of 0.01% or more is necessary. If exceeded, toughness and weldability deteriorate. Therefore, the amount is set to 0.01 to 1.5%.

  W is an element necessary for enhancing the hardenability and securing the strength, but the amount is set to 0.01 to 1.5% as a range that can exert the effect and does not adversely affect other characteristics.

  Ti is an element that contributes to improving the strength of the base metal by precipitation strengthening, and is also an effective element for refining the heated austenite grain size by forming TiN that is stable even at high temperatures. There is a need to. On the other hand, if it exceeds 0.05%, a coarse oxide is formed and the ductility is extremely deteriorated. Therefore, the amount is set to 0.001 to 0.05%.

  Nb, Zr, and V contribute to improving the strength of the base metal by precipitation strengthening, but if less than 0.005%, there is no effect, and if it exceeds 0.2%, ductility and toughness deteriorate. Therefore, the amount of Nb, Zr, and V is set to 0.005 to 0.2%.

  B is an element capable of improving the hardenability in a small amount by segregating at the austenite grain boundary in a solid solution state. However, B is also effective in suppressing recrystallization of austenite in the state segregated at the grain boundary. Addition of 0.0002% or more is necessary to exert effects on hardenability and recrystallization suppression. On the other hand, excessive addition exceeding 0.005% produces coarse precipitates and deteriorates toughness. Therefore, the amount was made 0.0002 to 0.005%.

Furthermore, in this invention, 1 type, or 2 or more types of Mg, Ca, and REM can be added as needed for the improvement of ductility and the improvement of joint toughness.
Mg, Ca, and REM are all effective in improving ductility by suppressing extension during hot rolling of sulfides. It effectively works to improve joint toughness by refining oxides. The lower limit content for exhibiting the effect is 0.0005% for Mg, 0.0005% for Ca, and 0.005% for REM. On the other hand, excessive content causes coarsening of sulfides and oxides, leading to deterioration of ductility and toughness. Therefore, the upper content is 0.01% for Mg, 0.01% for Ca, and 0 for REM. .05%.

The above is the reason for limiting the microstructure and chemical components, which are the basic requirements of the present invention. In addition, a suitable manufacturing method for satisfying the organizational requirements of the present invention is also presented. In this production method, a steel slab containing an appropriate chemical component is heated to a temperature not lower than the Ac3 transformation point and not higher than 1350 ° C, and then the cumulative reduction ratio is 53 to 80% in the austenite single phase region of Ar3 transformation point to 1250 ° C. After rolling, finish rolling is performed in an austenite-ferrite two-phase region in which the rolling start temperature is equal to or lower than the Ar3 transformation point and the rolling end temperature is 600 ° C. or higher, and the cumulative rolling reduction is 40 to 90%. After finish rolling, accelerated cooling to 20 to 400 ° C. may be performed at a cooling rate of 5 to 80 ° C./s. Furthermore, you may temper in the temperature range of 300-500 degreeC. However, the microstructure of the present invention, Ru der which effective regardless of its achievement means.

  Prior to hot rolling, it is necessary to make the steel ingot 100% austenitic. For this purpose, the temperature of the steel ingot needs to be heated to the Ac3 transformation point or higher. However, when heated above 1350 ° C., the austenite grains become extremely coarse and fine-grained ferrite cannot be obtained after rolling, so the upper limit of the heating temperature is 1350 ° C.

The subsequent hot rolling was limited to the temperature range of Ar3 transformation point to 1250 ° C. By rolling in the austenite single phase region, the transformation temperature was increased and the transformation structure was refined. This is because fine-grained ferrite can be obtained by rolling. This effect is cumulative rolling reduction below 10 percent less Iga, lower limit, it is appropriate to 53% based on the cumulative rolling reduction of the Invention Example A2 shown in Table 2 of Example. On the other hand, if it exceeds 80%, it becomes impossible to secure the reduction in the subsequent two-phase rolling . In this case, it is preferable to perform controlled rolling in the austenite region and further refine the austenite grains before the two-phase region rolling.

  In the present invention, it is necessary to disperse flat and hard martensite in a layer form in hard ferrite, and for this reason, finish rolling below the Ar3 transformation point plays an extremely important role and is an essential step in the present invention. . In order to improve the hardness of ferrite, the hardness of martensite, flattening, and lowering the transformation start temperature, finishing rolling at Ar3 point or lower is necessary, and the lower the rolling temperature is desirable, the lower the deformation resistance. As a result, the rolling load increases and rolling is difficult. On the other hand, when the temperature is lower than 600 ° C., it is impossible to secure a martensite fraction of 5% or more. Therefore, the rolling end temperature is set to 600 ° C. or higher.

  If the cumulative rolling reduction of finish rolling is less than 40%, the effect of improving the hardness of the ferrite, improving the hardness of the martensite, and flattening is small, and further, the layer spacing in the thickness direction of the ferrite and martensite increases. The larger the cumulative rolling reduction, the better. Therefore, the finish rolling cumulative reduction ratio is set to 40 to 90%.

As a cooling method after the two-phase region rolling, in order to perform martensite transformation, it is necessary to accelerate cooling to 20 to 400 ° C. at a cooling rate of 5 to 80 ° C./s up to the martensite transformation start temperature or lower.
The reason why the cooling rate in the case of accelerated cooling is limited to 5 to 80 ° C./s is that if it is less than 5 ° C./s, martensitic transformation is difficult for accelerated cooling, and improvement in strength and toughness of the base material cannot be expected. This is because if the temperature exceeds 80 ° C./s, a difference in structure or characteristics between the surface layer and the inside greatly occurs, which is not preferable.

  The accelerated cooling stops at 20 to 400 ° C. depending on the desired strength and toughness level of the steel sheet. Setting the stop temperature of accelerated cooling to less than 20 ° C. has no effect in controlling the material, and merely causes an increase in manufacturing cost and is meaningless. Conversely, if accelerated cooling is stopped at over 400 ° C., martensitic transformation is difficult, internal stress is relaxed, and fatigue crack propagation characteristics cannot be expected.

  The tempering process that is subsequently carried out as necessary after rolling and cooling is intended to improve the toughness of the base metal structure by recovery. Therefore, the heating temperature must be less than Ac1, which is a temperature range in which reverse transformation does not occur. Don't be. Further, if the temperature exceeds 500 ° C., the fatigue crack propagation characteristics deteriorate due to relaxation of internal stress, so the upper limit was set to 500 ° C. Further, the recovery is to reduce the lattice defect density by the disappearance and coalescence of dislocations, and in order to realize this, it is necessary to heat to 300 ° C. or higher, so the lower limit was set to 300 ° C. In addition, the tempered martensite produced | generated by this tempering heat processing is also defined as a martensite which is the structure | tissue requirements of this invention.

Hereinafter, the effects of the present invention will be described more specifically with reference to examples.
Table 1 shows the chemical composition of the test steel used in the examples. Each test steel is made into a steel slab by ingot rolling, by ingot rolling, or by continuous casting. Steel numbers 1 to 20 in Table 1 satisfy the chemical composition range of the present invention, and steel numbers 21 to 25 do not satisfy the chemical composition range of the present invention.

Steel strips having the chemical components shown in Table 1 were produced into steel plates under the conditions shown in Table 2. Test No. A1 to A23 were produced by the method according to claims 4-6. In addition, Test No. B1 to B12 do not satisfy the production conditions of the present invention. Table 2 shows the mechanical properties of each at room temperature.
Table 3 shows steel Nos. 1 to 25, test Nos. The microstructure check result of the steel plate which consists of A1-A23 and B1-B12 and the fatigue test result are shown.

  The microstructure is obtained by mirror-polishing the thickness cross section in the rolling direction of the steel sheet, revealing it by nital corrosion and repeller corrosion, observing it with an optical microscope, and identifying the generated phase in combination with the hardness test results described below. did. The hardness was measured using a micro Vickers hardness tester with a load of 10 g. The fraction, aspect ratio, and layer spacing of each phase were determined by image analysis of optical micrographs.

FIG. 1 is a view showing a test piece used in a fatigue crack propagation test. The fatigue crack propagation test conditions were as follows.
・ Loading method: 3-point bending,
-Stress ratio: 0.1,
・ Environment: In room temperature atmosphere,
・ Crack length measurement: DC potential difference method

FIG. 2 is a view showing a test piece used in a welded joint fatigue test. For welding, carbon dioxide arc welding was performed at a heat input of 18 kJ / min. The fatigue test conditions were as follows.
・ Loading method: axial force,
-Stress ratio: 0.1,
・ Environment: In room temperature atmosphere,
Test stress range: 150 MPa

Test No. Each of A1 to A20 is a steel material produced by producing a steel slab having the chemical composition of the present invention in accordance with the requirements of the present invention, which also satisfies the structural requirements, and a fatigue crack when the stress intensity factor range ΔK is 20 MPa√m. Propagation velocity da / dN is 10 ⁻⁸ m / cycle or less, and the weld joint fatigue life is Test No. As compared with the comparative example of B1, the fatigue characteristics were excellent, twice or more.

On the other hand, test no. In A21 to 23, the production requirements of the present invention are satisfied, but the chemical composition is not within the limited range. Test No. A21 and 23 have a ferrite-martensite structure, but since the martensite fraction is small or the layer spacing is large, the propagation speed when ΔK = 20 MPa√m is 10 ⁻⁸ m / cycle or more. Therefore, the weld joint fatigue life is Test No. The fatigue property was inferior to that of the comparative steel of B1, and the fatigue properties were inferior to the steel of the present invention. In addition, Test No. Since A22 has an excessive martensite fraction, the toughness was greatly deteriorated, and the fatigue life of the welded joint was significantly inferior to that of the steel of the present invention because brittle fracture occurred during the fatigue test. In addition, the propagation characteristics were inferior to the newly developed steel due to the too small layer spacing.

In addition, Test No. Although B1-B10 is satisfying the limited range of the chemical composition of the present invention, the production requirements are deviated. Test No. In B1, B6, B7, B8, and B10, the second phase is not martensite, and other than martensite is less likely to be an effective obstacle to crack growth. Therefore, fatigue crack propagation characteristics deteriorate compared to the steel of the present invention. The fatigue life of welded joints was not improved.
Test No. In B2 and B3, the second phase is martensite, but the ferrite hardness is small and the internal stress cannot be increased. In addition, the aspect ratio is small or the layer spacing is large. Since the frequency of hitting martensite was low and could not be an effective obstacle, the fatigue crack propagation characteristics were deteriorated compared to the steel of the present invention, and the fatigue life of the welded joint was not improved.

Test No. Since B4 had a high tempering temperature and the internal stress was relaxed, B4 did not become an obstacle to crack growth, and the fatigue properties were inferior to the steel of the present invention. Test No. B5 had a high finish rolling start temperature and an accelerated cooling start temperature, and most of the second phase became bainite, so it did not become an obstacle to crack growth and the fatigue properties were inferior to the steel of the present invention. Test No. In B9, the second phase is martensite, but since the finish cumulative reduction ratio was small and the aspect ratio was very small, the frequency of hitting martensite during crack growth was low and could not be an effective failure. It was inferior to the inventive steel.
Furthermore, test no. Regarding B11 to B12, both the chemical composition and the production method did not satisfy the limited range of the present invention, so the fatigue characteristics were significantly deteriorated compared to the steel of the present invention.

It is a figure which shows the test piece used for the fatigue crack propagation test. It is a figure which shows the test piece used for the weld joint fatigue test.

Claims (3)

% By mass
C: 0.0 5 ~0.2%,
Si: 0.01~ 0.8%,
Mn: 0.5-2%
P: 0.02% or less,
S: 0.02% or less,
Al: 0.001 to 0.1%,
N: 0.001 to 0.008%
Containing the balance Fe and unavoidable impurities, and the microstructure of the base material is a ferrite having a Vickers hardness of 150 or more, the Vickers hardness is 400 to 900, the area ratio is 5 to 30%, and the aspect ratio (Major axis / minor axis) is a layered structure having flat martensite of 3 or more as a second phase, and the average layer spacing in the plate thickness direction of ferrite and martensite is 3 to 50 μm, and the stress ratio is 0. A steel plate having excellent fatigue crack propagation characteristics, wherein a fatigue crack propagation rate da / dN when a stress intensity factor range ΔK of 1 is 20 MPa√m is 10 ⁻⁸ m / cycle or less. In addition,
Cu: 0.1 to 2.5%,
Ni: 0.1 to 5%,
Cr: 0.01 to 1.5%
Mo: 0.01 to 1.5%,
W: 0.01 to 1.5%,
Ti: 0.001 to 0.05%,
Nb: 0.005 to 0.2%,
Zr: 0.005 to 0.2%
V: 0.005 to 0.2%,
B: 0.0002 to 0.005%
The steel plate excellent in fatigue crack propagation characteristics according to claim 1, comprising one or more of the following. In addition,
Mg: 0.0005 to 0.01%,
Ca: 0.0005 to 0.01%,
REM: 0.005 to 0.05%
The steel plate excellent in fatigue crack propagation characteristics according to claim 1 or 2, characterized by containing one or more of the following.

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