EP3656886A1

EP3656886A1 - Steel plate

Info

Publication number: EP3656886A1
Application number: EP19220022.8A
Authority: EP
Inventors: Yusuke SANDAIJI; Masao Kinefuchi; Haruya KAWANO
Original assignee: Kobe Steel Ltd
Current assignee: Kobe Steel Ltd
Priority date: 2014-05-22
Filing date: 2015-05-15
Publication date: 2020-05-27
Also published as: KR20160144439A; EP3147379B1; KR20180115352A; EP3147379A4; EP3147379A1; WO2015178320A1; KR102218806B1; CN106460114A; JP6472315B2; JP2016000855A; CN106460114B

Abstract

Disclosed is a steel plate. The steel plate has a chemical composition meeting predetermined conditions. A microstructure in the surface layer of the steel plate includes at least one of ferrite and upper bainite in a total fraction of 80 area percent or more, and grains of the at least one of ferrite and upper bainite have an effective grain size of 10.0 µm or less. Of the microstructure of the surface layer, grains of the remainder microstructure excluding the ferrite and the upper bainite have an average equivalent circle diameter of 3.0 µm or less. The steel plate has a dislocation density ρ of 5.0× 10¹³ m^-1 to 2.5× 10¹⁵ m^-1 as determined by X-ray diffractometry. The steel plate having the configuration has excellent fatigue properties.

Description

Technical Field

The present invention relates to steel plates. Specifically, the present invention relates to a steel plate that is mainly used as materials for structures such as ships buildings, bridges, and construction machinery, has a tensile strength of 490 MPa to less than 650 MPa, and offers excellent fatigue properties.

Background Art

With upsizing, large structures such as ships, buildings, bridges, and construction machinery require still better reliability of their structural components, because such larger structures, if suffering from breakage, may receive larger damage. It has been known that most of breakages in large structures is caused by fatigue fracture. As possible solutions to this, various techniques for resisting fatigue fracture have been developed. However, there are not a few cases where fatigue fracture causes breakage of such a large structure, even now.
In general, portions of large structures susceptible to fatigue damage have been protected from fatigue fracture by schemes such as designing of the portions to have such a shape as to less cause stress concentration, and use of high-strength steel plates. However, the resulting structures obtained according to the schemes cause higher production cost due to a higher number of steps or due to the use of more expensive steel plates. Accordingly, demands are made to provide a technique for allowing a steel plate itself to have better fatigue properties. In general in steel plates, it is known that fatigue strength, in particular fatigue limit, is in proportional to tensile strength, and a steel plate having a high fatigue limit to tensile strength ratio is considered as a steel plate having excellent fatigue properties, where the fatigue limit to tensile strength ratio is calculated by dividing the fatigue limit by the tensile strength.
Many researches have been performed so as to offer better fatigue properties. For example, Nonpatent Literature (NPL) 1 presents how various influencing factors affect fatigue strength. The literature mentions that solid-solution strengthening, precipitation strengthening, grain refinement, and second phase strengthening contribute to better fatigue properties, but dislocation hardening less contributes to better fatigue properties, because the dislocation hardening is attended with increase in moving dislocations. The fatigue fracture process can be divided into a process (1) and a process (2). In the process (1), a load is repeatedly applied, and, finally, cracks are generated. In the process (2), the generated cracks grow and lead to rupture (breakage). Of the above-mentioned factors for better fatigue properties, solid-solution strengthening, precipitation strengthening, and grain refinement are considered to be effective in the process (1), because restrainment of accumulation ofdislocations is effective in this process. In contrast, grain refinement and second phase strengthening are considered to be effective in the process (2), because elimination or minimization of crack propagation is effective in this prooess.
Patent Literature (PTL) 1 proposes a technique in which the steel is controlled to have a two-phase microstructure including fine ferrite and hard martensite, and the difference in hardness between the two phases is specified to be 150 or more in terms of Vickers hardness. The technique is intended to lower the crack propagation rate and to contribute to a longer fatigue life after crack initiation.
PTL 2 proposes a technique for lowering the crack propagation rate by allowing the steel to have a mixed microstructure including fine ferrite and bainite. This technique is expected to offer longer fatigue life after crack initiation also in fatigue fracture. However, the technique lacks consideration of fatigue properties before crack initiation at all.
PTL 3 proposes a technique of allowing carbides to precipitate in ferrite phase so as to offer higher fatigue strength. This literature, however, lacks description about fatigue properties after crack initiation In addition, the technique in PTL 3 targets thin steel sheets and gives no consideration to toughness and other properties necessary for large structures at all.

Citation List

Nonpatent literature

NPL 1: Abe et al, "Tetsu-to-hagane", Vol. 70 (1984), No. 10, pp. 1459-1466

Patent Literature

PTL 1: Japanese Unexamined Patent Application Publication ( JP-A) No. Hei10-60575
PTL 2: JP-A No. 2011-195944
PTL 3: JP-A No. 2009-84643

Summary of Invention

Technical Problem

The present invention has been made under these circumstances and has a main object to provide a steel plate having excellent fatigue properties.
Other objects of the present invention will be apparent from descriptions in "Description of Embodiments" below.

Solution to Problem

The present invention has achieved the objects and provides, according to a first embodiment, a steel plate as follows. The steel plate contains C in a content of 0.02 to 0.10 mass percent, Mn in a content of 1.0 to 2.0 mass percent, Nb in a content of greater than 0 mass percent to 0.05 mass percent, Ti in a content of greater than 0 mass percent to 0.05 mass percent, Al in a content of 0.01 to 0.06 mass percent, and at least one element selected from the group consisting of Si in a content of 0.1 to 0.6 mass percent and Cu ina content of 0.1 to 0.6 mass percent, where the total content of the at least one of Si and Cu is 0.3 mass percent or more, with the remainder consisting of iron and inevitable impurities. The microstructure of the surface layer of the steel plate includes at least one of ferrite and upper bainite in a total fraction of 80 area percent or mare. Grains of the at least one of ferrite and upper bainite have an effective grain size of 100 µm or less. Of the microstructure of the surface layer, grains of the remainder microstructure excluding the ferrite and the upper bainite have an average equivalent circle diameter of 3.0 µm or less. The steel plate has a dislocation density ρ of 2.5× 10¹⁵m^-1 or less, as determined by X-ray diffractometry.
As used herein, the term "average equivalent circle diameter" refers to an average of "equivalent circle diameters" of grains of a phase in question, where the "equivalent circle diameter" refers to the diameter of a grain of the phase in terms ofa circle having an equivalent area to the grain.
The fraction (proportion) of martensite-austenite constituent in the remainder microstructure is preferably 5 area percent or less.
The steel plate according to the present invention preferably further contains at least one selected from the group consisting of (a), (b), and (c) below. The steel plate may have further improved property or properties depending on the type of an element to be contained. Specifically, the steel plate preferably further contains at least one selected from the group consisting of:

(a) Ni in a content of greater than 0 mass percent to 0.6 mass percent, where the ratio [Ni]/[Cu] of the Ni content [Ni] to the Cu content [Cu] is less than 1.2;
(b) at least one element selected from the group consisting of V in a content of greater than 0 mass percent to 0.5 mass percent, Cr in a content of greater than 0 mass percent to 1.0 mass percent or less, and Mo in a content of greater than 0 mass percent to 0.5 mass percent; and
(c) B in a content of greater than 0 mass percent to 0.005 mass percent.

The steel plate according to the present invention also preferably has a bainite transformation start temperature Bs of 640°C or higher, where the bainite transformation start temperature is calculated from the chemical composition according to Expression (1):
where [C], [Mn], [Ni], [Cr], and [Mo] represent contents (in mass percent) respectively of C, Mn, Ni, Cr, and Mo.
The present invention has archived the objects and provides, according to a second embodiment, a steel plate as follows. This steel plate contains C in a content of 0.02 to 0.10 mass percent, Mn in a content of 1.0 to 2.0 mass percent, Nb in a content of greater than 0 mass percent to 0.05 mass percent, Ti in a content of greater than 0 mass percent to 0.05 mass percent, Al in a content of 0.01 to 0.06 mass percent, at least one element selected from the group consisting of Si in a content of0.1 to 0.6 mass percent and Cu in a content of 0.1 to 0.6 mass percent, where the at least one of Si and Cu is contained in a total content of 0.3 mass percent or more, and B in a content of greater than 0 mass percent to 0.005 mass percent, with the remainder consisting of iron and inevitable impurities. The microstructure of the surface layer of the steel plate includes at least one of ferrite and upper bainite in a fraction of 80 area percent or more. Grains of the at least one of ferrite and upper bainite have an effective grain size of 10.0 µm or less. Of the microstructure of the surface layer, grains of the remainder microstructure excluding the ferrite and the upper bainite have an average equivalent circle diameter of 3.0 µm or less. The steel plate has a dislocation density ρ of 2.5× 10¹⁵ m^-1 or less as determined by X-ray diffractometry. A microstructure at a position at a depth of one-fourth the thickness t of the steel plate from the surface along the thickness direction includes upper bainite in a fraction of 80 area percent or more in a longitudinal section in parallel with the rolling direction. Grains of the upper bainite in the microstructure at the position at a depth of one-fourth the thickness t have an effective grain size of 10.0 µm or less. Of the microstructure at the position at a depth of one-fourth the thickness t, grains of the remainder microstructure excluding the upper bainite have an average equivalent circle diameter of 3.0 µm or less.
Also in the steel plate according to the second embodiment, of the microstructure of the surface layer, the remainder microstructure excluding the ferrite and the upper bainite preferably has a fraction of martensite-austenite constituent of 5 area percent or less.
The steel plate according to the second embodiment preferably further contains, in chemical composition, at least one selected from the group consisting of (a) and (b) below. The steel plate may have better property or properties according to the type of an element to be contained. Specifically, the steel plate preferably further contains at least one selected from the group consisting of:

(a) Ni in a content of greater than 0 mass percent to 0.6 mass percent, where the ratio [Ni]/[Cu] of the Ni content [Ni] to the Cu content [Cu] is less than 1.2; and
(b) at least one element selected from the group consisting of V in a content of greater than 0 mass percent to 0.5 mass percent, Cr in a content of greater than 0 mass percent to 1.0 mass percent or less, and Mo in a content of greater than 0 mass percent to 0.5 mass percent.

The steel plate according to the second embodiment also preferably has a bainite transformation start temperature Bs of 640°C or higher, where the bainite transformation start temperature is calculated from the chemical composition according to Expression (1):
where [C], [Mn], [Ni], [Cr], and [Mo] represent contents (in mass percent) respectively of C, Mn, Ni, Cr, and Mo.
Assume that the microstructure at the position one fourth the thickness t of the steel plate from the surface along the thickness direction meets the condition(s), where the microstructure is in a longitudinal section in parallel with the rolling direction. In this case, the microstructure at the position at a depth of one-fourth the thickness t preferably includes grains each having a grain average misorientation (GAM) of 1° or more in a fraction of 20 area percent to 80 area percent, where the grain average misorientation is determined per one grain by electron backscatter pattern (EBSP) analysis.
Assume that the microstructure at the position at a depth of one-fourth the thickness t includes grains each having a GAM of 1° or more in a fraction of 20 area percent to 80 area percent, where the grain average misorientation is determined per one grain by EBSP analysis. In this case, the steel plate preferably meets conditions specified by Expression (2) and Expression (3):

where [Si], [Mn], [Ni], [Cu], [Ti], [N], [Cr], [Nb], [C], and [Mo] represent contents (in mass percent) respectively of Si, Mn, Ni, Cu, Ti, N, Cr, Nb, C, and Mo, and, when at least one of the term {[Ti] - 3.4[N]} and the term {[Nb] - 7.7[C]} is negative, the calculation according to the expression is performed as treating the at least one term as "zero (0)". Advantageous Effects of Invention
Advantageous effects of the present invention as disclosed in the description will be briefly described as follows. Specifically, the present invention can practically provide steel plates that have excellent fatigue properties.

Brief Description of Drawings

Fig. 1 depicts schematic explanatory views of a steel plate according to the present invention;
Fig. 2 depicts schematic explanatory views of a test specimen used in measurements of fatigue properties;
Fig. 3 depicts schematic explanatory views of a compact tension test specimen used in crack propagation rate measurement; and
Fig. 4 depicts conceptual diagrams illustrating grain boundaries, KAM, and GAM.

Description of Embodiments

The inventors of the present invention made investigations of the total life of steel plates leading up to fatigue fracture and particularly of proportions of the prior-stage life leading to crack initiation and of the subsequent-stage life from crack initiation leading up to rupture. As a result, the inventors found that the prior-stage life leading to crack initiation occupies about one half of the total life leading up to fatigue fracture; and that the prior-stage life leading to crack initiation occupies a larger proportion with a decreasing stress level and a resulting increasing total life. The results indicate that increase of the total life leading up to fatigue fracture requires not only better fatigue properties after crack initiation, but also better fatigue properties leading to crack initiation. In particular, at around the fatigue limit, increase in prior-stage life is considered to be effective because the prior-stage life leading to crack initiation tends to occupy a larger proportion of the total life.
The inventors made various investigations on conditions for a longer prior stage life. As a result, the inventors found that a steel plate as follows can have a longer prior-stage life, and consequently have a longer total life leading up to fatigue fracture. This steel plate is appropriately controlled in the chemical composition, and conditions typically on the fraction(s) of principal phase(s), the effective grain size of grains of the phase(s), the average equivalent circle diameter of the remainder microstructure excluding the principal phase(s), and the dislocation density ρ as determined by X-ray diffractometry. The present invention has been made based on these findings.

First Embodiment

The first embodiment according to the present invention will be illustrated below. The inventors made investigations on various steel plates about how elements to be added affect the fatigue strength. Consequently, the inventors found that the addition of Si and/or Cu significantly improves the fatigue strength. In general, moving dislocations, which move by repeated stress, move irreversibly due typically to cross slip, and this causes fatigue cracking. It is known that the dislocations form a cell structure in this process. The inventors found that the addition of Si and/or Cu in a total amount of 0.3 mass percent or more restrains the cell structure formation.
Detailed observations demonstrate that these elements do not form precipitates and are not noticeably dissolved as solutes typically in carbides present in steel plates, but are probably dissolved and present as solutes in the matrix. Specifically, it is considered that a steel plate can have a longer fatigue life up to (prior to) crack initiation by allowing these elements to be dissolved and present as solutes sufficiently in the matrix to thereby restrain the irreversible movements of dislocations. The inventors also found that other elements to be added, such as Mn and Cr, do not noticeably offer the effect of restraining dislocations from forming a cell structure, but rather lower the transformation temperature to cause the formation of lower bainite, which has a high dislocation density; and that such other elements to be added do not so much improve fatigue strength as compared with static strength.
In addition, the inventors prepared various steel plates under different rolling conditions and made investigations on how the rolling conditions affect the mechanical properties and fatigue strength of the resulting steel plates. As a result, the inventors found that steel plates, when caused to have a higher dislocation density typically by rapidly cooling the steel plates down to a low temperature, or by applying compression reduction (performing rolling) at a temperature equal to or lower than the Ar₃ transformation temperature, have higher static strength such as yield stress and tensile strength, but have not so higher fatigue strength as compared with the static strength, and have a lower fatigue limit to tensile strength ratio.
This is probably because moving dislocations introduced upon steel plate production serve as a dislocation hardening factor and contribute to higher static strength, but cause cyclic softening during repeated deformation, namely, during a fatigue test, and thereby less contribute to higher fatigue strength. In particular, a steel plate having a dislocation density p of greater than 2.5× 10¹⁵ m^-1 as determined by X-ray diffractometry (XRD) undergoes significant improvement in static strength due to dislocation hardening and tends to have a lower fatigue limit to tensile strength ratio. The dislocation density ρ is preferably 2.0× 10¹⁵ m^-1 or less, and more preferably 1.5× 10¹⁵ m^-1 or less. The dislocation density ρ may be about 5.0× 10¹³ m^-1 or more in terms of lower limit.
Fig. 1 depicts schematic explanatory views of a steel plate according to the present invention. Specifically, Figs. 1(a) and 1(b) are a schematic perspective view and a schematic side view, respectively, of the steel plate according to the present invention. Figs. 1(a) and 1(b) illustrates a rolling direction L, a transverse direction (width direction) W, a thickness direction D, a steel plate surface S1, and a section S2 in the thickness direction D in parallel with the rolling direction L.
The inventors made investigations on various steel plates having different chemical compositions and different microstructure morphologies to examine fatigue properties. As a result, the inventors found that a steel plate is allowed to have excellent fatigue properties by controlling the microstructure of a longitudinal section (namely, the section S2 in Figs. 1(a) and 1(b)) as follows, where the longitudinal section is in parallel with the rolling direction, and where the microstructure is in a surface layer at a position adjacent to the steel plate surface S1, for example, in a surface layer at a position about 1 to about 3 mm deep from the steel plate surface S1 in the thickness direction. The surface layer herein is defined to be at a portion about 1 to about 3 mm deep from the steel plate surface, in order to evaluate a surface layer of the steel plate itself, excluding a scale layer, because the steel plate surface immediately after production may include the scale layer of about 0.1 to about 2 mm depth (thickness) when produced under some production conditions.
The ferrite and upper bainite are phases that are relatively resistant to introduction of moving dislocations upon formation of the phases, as compared with other phases. This restrains the fatigue limit to tensile strength ratio from decreasing and contributes to longer life leading to crack initiation. To offer these advantageous effects, the microstructure in the surface layer is controlled to include at least one of ferrite and upper bainite in a total fraction of 80 area percent or more. The fraction of the at least one of ferrite and upper bainite is preferably 85 area percent or more, and more preferably 90 area percent or more. In terms of upper limit, the fraction of the at least one of ferrite and upper bainite may be 100 area percent, but is typically about 98 area percent or less.
The grains of ferrite and upper bainite in the surface layer, if coarsen, may often undergo stress concentration, which causes cracks. When cracks after crack initiation collide with grains, the cracks stop and bypass the grains, and this restrains crack propagation. However, if the grains are coarse, cracks collide with the grains in a relatively smaller frequency. Thus, it is considered that the fatigue crack propagation in the surface layer is not sufficiently restrained. Based on these findings and considerations, the effective grain size of grains is specified to be 10.0 µm or less, where the grains herein are each defined as a region surrounded by high-angle grain boundaries with a misorientation of 15° or more between adjacent grains of ferrite or upper bainite. The term "effective grain size" refers to an average length of the grains in the thickness direction.
The effective grain size of the grains of at least one of ferrite and upper bainite is preferably 6 pm or less, and more preferably 5 µm or less. The lower limit of the effective grain size of the grains of at least one of ferrite and upper bainite is not limited, but is typically greater than about 2 µm.
The upper bainite phase can have a smaller size as compared with the ferrite phase, but is attended with shear deformation upon transformation, to which moving dislocations are readily introduced. In particular, bainitic transformation, when allowed to occur at a low temperature, often gives a lower bainite phase containing a large amount of moving dislocations. To restrain the lower bainite phase from forming, the bainite transformation start temperature Bs is preferably controlled appropriately. From this viewpoint, the bainite transformation start temperature Bs calculated according to Expression (1) is preferably 640°C or higher, and more preferably 660°C or higher.
The remainder microstructure excluding the ferrite and the upper bainite in the surface layer is controlled to have an average equivalent circle diameter of 3.0 pm or less. The remainder microstructure is controlled to have an average equivalent circle diameter of 3.0 pm or less, because the remainder microstructure, if having an average equivalent circle diameter greater than 3.0 pm, may cause toughness and other properties to significantly deteriorate. The average equivalent circle diameter of the remainder microstructure is preferably 2.5 pm or less, and more preferably 2.0 µm or less in terms of upper limit; and is preferably about 0.5 pm or more in terms of lower limit. The remainder microstructure excluding the ferrite and the upper bainite in the surface layer basically includes martensite, martensite-austenite constituent (MA), pearlite, and pseudo-pearlite.
Among them, the martensite-austenite constituent, which is formed typically in cooling process after rolling, undergoes expansive transformation in its formation process, introduces moving dislocations into the matrix, and causes the steel plate to have a shorter life leading to crack initiation. To eliminate or minimize this, the proportion of the martensite-austenite constituent in the remainder microstructure in the surface layer is preferably controlled to be an area percent of 5% or less. The area percent of the martensite-austenite constituent is preferably minimized, and is more preferably 3% or less, furthermore preferably 1% or less, and most preferably 0%.
Next, such a chemical composition of the steel plate according to the present invention as to have better fatigue properties will be described. The steel plate according to the present invention, as being incorporated with C, Mn, Nb, and other alloy elements as appropriate, is allowed to surely have a fine ferrite phase and/or an upper bainite phase. Simultaneously, the steel plate, as being controlled in elements to be added such as Si and Cu as appropriate, restrains dislocations from forming a cell structure, which causes fatigue crack initiation. Thus, the steel plate can have excellent fatigue properties. From these viewpoints, the elements are controlled in the following manner.

C: 0.02 to 0.10 mass percent

Carbon (C) is important to allow the steel plate to have strength at certain level To this end, the carbon content is specified to be 0.02 mass percent or more. The carbon content is preferably 0.03 mass percent or more, and more preferably 0.04 mass percent or more. In contrast, the steel plate, if containing carbon in an excessively high content, may have excessively high strength to fail to have desired tensile strength. In addition, this steel plate, when undergoing accelerated cooling, may have excessive hardenability, have a large dislocation density p, and offer lower fatigue properties. To eliminate or minimize these, the carbon oontent is controlled to be 0.10 mass percent or less, and is preferably 0.08 mass percent or less, and more preferably 0.06 mass percent or less.

Mn: 1.0 to 2.0 mass percent

Manganese (Mn) has a significance to ensure hardenability so as to give a fine microstructure. To effectively offer this activity, the Mn content is specified to be 1.0 mass percent or more. The Mn content is preferably 1.2 mass percent or more, and more preferably 1.4 mass percent or more. However, the steel plate, if containing Mn in an excessively high content, may have excessive hardenability, have a higher dislocation density p, and fail to have sufficient fatigue properties. To eliminate or minimize these, the Mn content is controlled to be 2.0 mass percent or less, and is preferably 18 mass percent or less, and more preferably 1.6 mass percent or less.

Nb: greater than 0 mass percent to 0.05 mass percent

Niobium (Nb) is effective for better hardenability and for a finer microstructure. To effectively offer these activities, the Nb content is preferably controlled to be 0.01 mass percent or more, and more preferably 0.02 mass percent or more. However, the steel plate, if containing Nb in an excessively high content, may have excessive hardenability and fail to have desired fatigue properties. To eliminate or minimize this, the Nb content is controlled to be 0.05 mass percent or less, and is preferably 0.04 mass percent or less, and more preferably 0.03 mass percent or less.

Ti: greater than 0 mass percent to 0.05 mass percent

Titanium (Ti) effectively contributes to better hardenability and, simultaneously, forms TiN to allow the heat-affected zone upon weldingto have a finer microstructure and to restrain reduction in toughness. To offer these activities, the Ti content is preferably 0.01 mass percent or more, and more preferably 0.02 mass percent or more. However, Ti, if contained in an excessively high content, may form coarse TiN particles and may cause the steel plate to have properties such as toughness at lower levels. To eliminate or minimize this, the Ti content is controlled to be 0.05 mass percent or less, and is preferably 0.04 mass percent or less, and more preferably 0.03 mass percent or less.

Al: 0.01 to 0.06 mass percent

Aluminum (Al) is useful for deoxidation and, if contained in a content less than 0.01 mass percent, may fail to offer effective deoxidation. The Al content is preferably 0.02 mass percent or more, and more preferably 0.03 mass percent or more. However, the steel plate, if containing Al in an excessively high content, may have excessive hardenability, have a higher dislocation density p, and fail to offer desired fatigue properties. To eliminate or minimize these, the content is controlled to be 0.06 mass percent or less, and is preferably 0.05 mass percent or less, and more preferably 0.04 mass percent or less.
At least one of Si in a content of 0.1 to 0.6 mass percent and Cu in a content of 0.1 to 0.6 mass percent
Silicon (Si) contributes to solid-solution strengthening to a large extent and is necessary for ensuring the strength of the base metal. Simultaneously, this element restrains the movements of dislocations and restrains the cell structure formation. To effectively offer these activities, the Si content is specified to be 0.1 mass percent or more. The Si content is preferably 0.2 mass percent or more, and more preferably 0.3 mass percent or more. However, the steel plate, if containing Si in an excessively high content, may include the remainder microstructure formed in excess and coarsely and may suffer from reduction in other properties such as toughness. To eliminate or minimize this, the Si content is controlled to be 0.6 mass percent or less, and is preferably 0.55 mass percent or less, and more preferably 0.5 mass percent or less.
Copper (Cu) restrains the cross slip of dislocations and effectively restrains the cell structure formation. To effectively offer the activities, the Cu content is specified to be 0.1 mass percent or more. The Cu content is preferably 0.2 mass percent or more, and more preferably 0.3 mass percent or more. However, the steel plate, if containing Cu in an excessively high content, may not only have excessive hardenability, but also become susceptible typically to cracking upon hot working. To eliminate or minimize these, the Cu content is controlled to be 0.6 mass percent or less, and is preferably 0.55 mass percent or less, and more preferably 0.5 mass percent or less.
Si and Cu can offer the common activity of restraining cell structure formation of dislocations. From this viewpoint, the steel plate may contain each of these elements alone or in combination. The effect of restraining cell structure formation of dislocations by Si and Cu is effectively offered when the total content of Si and Cu ([Si] + [CuD is 0.3 mass percent or more. The total content is preferably 0.4 mass percent or more. A preferred upper limit of the total content ([Si] + [Cu]) is the total of the preferred upper limits of the two elements.
The steel plate according to the present invention includes the elements as mentioned above as a basic composition, with the remainder consisting of approximately iron. However, it is naturally accepted that inevitable impurities such as P, S, and N are contained in the steel. The inevitable impurities are brought into the steel in circumstances of raw materials, facility materials, and production equipment. It is also effective that the steel plate according to the present invention positively contains one or more of elements below. The steel plate can have a still better property or properties according to the type(s) of element(s) to be contained.

Ni: greater than 0 mass percent to 0.6 mass percent

Nickel (Ni) effectively contributes to better hardenability and contributes to a finer microstructure. Simultaneously, this element effectively restrains cracking upon hot working, where the cracking may more readily occur by the addition of Cu. To effectively offer these effects, Ni is preferably contained in a content of 0.1 mass percent or more, and more preferably 0.2 mass percent or more. However, the steel plate, if containing Ni in an excessively high content, may have excessive hardenability, have an excessively high dislocation density p, and thereby fail to have desired fatigue properties. To eliminate or minimize these, the Ni content is preferably controlled to be 0.6 mass percent or less, more preferably 0.5 mass percent or less, and furthermore preferably 0.4 mass percent or less.
The steel plate, if having an excessively high Ni content [Ni] with respect to the Cu content [Cu], may hardly enjoy the effect of restraining cell structure formation of dislocations by Cu. To eliminate or minimize this, the ratio ([Ni]/[Cu]) of the Ni content [Ni] to the Cu content [Cu] is preferably controlled to be less than 1.2, and more preferably 1.1 or less. The ratio ([Ni]/[Cu]) may be about 0.5 or more in terms of lower limit
At least one element selected from the group consisting of V in a content of greater than 0 mass percent to 0.5 mass percent, Cr in a content of greater than 0 mass percent to 1.0 mass percent or less, and Mo in a content of greater than 0 mass percent to 0.5 mass percent
Vanadium (V), chromium (Cr), and molybdenum (Mo) effectively allow the steel plate to have better hardenability and to have a finer nacrostructure. To offer these activities, the steel plate preferably contains each of V in a content of 0.01 mass percent or more, Cr in a content of 0.1 mass percent or more, and Mo in a content of 0.01 mass percent or more alone or in combination. However, the steel plate, if containing at least one of these elements in an excessively high content, may have excessive hardenability, have an excessively high dislocation density p, and fail to have desired fatigue properties. To eliminate or minimize these, the contents of V, Cr, and Mo are preferably controlled to be respectively 0.5 mass percent or less, 1.0 mass percent or less, and 0.5 mass percent or less. The contents of V, Cr, and Mo are more preferably controlled to be respectively 0.4 mass percent or less, 0.8 mass percent or less, and 0.4 mass percent or less.

B: greater than 0 mass percent to 0.005 mass percent

Boron (B) contributes to better hardenability and, in particular, restrains a coarse ferrite phase from forming, and thereby allows a fine upper bainite phase to form more readily. To offer these effects, the boron content is preferably controlled to be 0.0005 mass percent or more, and more preferably 0.001 mass percent or more. However, the steel plate, if containing boron in an excessively high content, may have excessive hardenability, have an excessively high dislocation density p, and fail to have desired fatigue properties. To eliminate or minimize these, the boron content is preferably controlled to be 0.005 mass percent or less, and more preferably 0.004 mass percent or less.
The steel plate according to the present invention not limited, but having an excessively small thickness, may less offer longer crack propagation life. From this viewpoint, the steel plate has a thickness of preferably 6 mm or more, and more preferably 10 mm or more.
The steel plate according to the present invention meets the conditions (requirements) and is not limited in production method. However, it is preferred to control production conditions as mentioned below, so as to give the microstructure morphology for better fatigue properties. The production conditions are conditions in a series of production process for the steel plate using a slab, such as a slab, having a chemical composition within the ranges. In the production process, a steel is made via ingot making and casting, and is subjected to hot rolling. Specifically, the production conditions include the heating temperature before hot rolling; the cumulative compression reduction in the entire hot rolling process; the finish-rolling temperature; the average cooling rate from the finish-rolling temperature or 800°C, whichever is lower, down to 600°C; and the cooling stop temperature.

Temperature of heating before hot rolling: 1000°C to 1200°C
Cumulative compression reduction in entire hot rolling process: 70% or more
Finish-rolling temperature: from the Ar₃ transformation temperature to (the Ar₃ transformation temperature + 150°C)
Average cooling rate from the finish-rolling temperature or 800°C, whichever is lower, down to 600°C: 15°C/second or less
Cooling stop temperature: 500°C or higher

Before the hot rolling, the slab is preferably heated up to a temperature range of 1000°C to 1200°C, and more preferably up to 1050°C or higher. The heating is preferably performed up to a temperature range of 1000°C or higher so as to eliminate or minimize coarsening of grains and to still ensure a cumulative compression reduction in hot rolling of 70% or more, as mentioned below. However, the heating, if performed up to an excessively high temperature of higher than 1200°C, may fail to contribute to refinement (size reduction) of the microstructure even when sufficient compression reduction is applied. To eliminate or minimize this, the heating temperature is preferably controlled to be 1200°C or lower, and more preferably 1150°C or lower.
The cumulative compression reduction in the entire hot rolling process is preferably 70% or more, and more preferably 75% or more. To reduce the microstructure size, in particular to reduce the effective grain size, sufficient compression reduction is to be applied in the non-recrystallization temperature range.
In addition, the finish-rolling temperature is preferably controlled within the range of the Ar₃ transformation temperature to the (Ar₃ transformation temperature + 150°C) so as to ensure desired fine microstructure and to still restrain excessive dislocations from being introduced into the microstructure after rolling (as-rolled microstructure). The finish-rolling temperature is more preferably controlled within the range of (the Ar₃ transformation temperature + 20°C) to (the Ar₃ transformation temperature + 100°C).
The "cumulative compression reduction" is a value calculated according to Expression (4):
where to represents the slab rolling start thickness (in millimeter (mm)) when the temperature at a position 3 mm deep from the surface fells within the rolling temperature range; t₁ represents the rolling finish thickness of the slab (in mm) when the temperature at a position 3 mm deep from the surface fells within the rolling temperature range; and t₂ represents the thickness of the slab, such as a slab, before the rolling.
The Ar₃ transformation temperature employed herein is a value determined according to Expression (5):
where [C], [Si], [Mn], [Cu], [Ni], [Cr], [Mo], and [Nb] represent contents (in mass percent) respectively of C, Si, Mn, Cu, Ni, Or, Mo, and Nb.
After hot rolling finish, cooling is preferably performed at an average cooling rate of 15°C/second or less from the finish-rolling temperature or 800°C, whichever is lower, down to600°C or lower. The cooling, if performed at an average cooling rate greater than 15°C/second, may cause the microstructure transformation to complete at an approximately low temperature unless a process such as isothermal holding is performed This causes excessive introduction of dislocations and fails to give desired fatigue properties. The average cooling rate is more preferably 10°C/second or less.
The cooling at the average cooling rate may be stopped at a temperature (namely, cooling stop temperature) of 500°C or higher. This restrains coarse ferrite phase formation and ensures a fine ferrite or upper bainite phase. The cooling, if stopped at a temperature lower than 500°C, may cause the transformation to complete at a low temperature, cause excessive dislocations to be introduced, and fail to give desired fatigue properties.
The temperature range within which cooling is performed at the average cooling rate is from 800°C down to 600°C when the finish-rolling temperature is higher than 800°C; and is from the finish-rolling temperature down to 600°C when the finish-rolling temperature is lower than 800°C. The average cooling rate in terms of lower limit is preferably 3.0°C/second or more, from the viewpoint of microstructure control in the steel plate, as mentioned below.

Second Embodiment

Next, the second embodiment according to the present invention will be illustrated. Steel plates for use in large structures also require lower crack propagation rate, namely, better crack propagation properties (better crack propagation resistance). This is because, when the crack propagation rate is low, even in case of fatigue cracking generation, the damaged portion can be found and repaired before the crack leads to fracture.
The inventors performed crack propagation tests and microstructure observations on various steel plates. As a result, the inventors found that control of a microstructure morphology at a specific position, in addition to the microstructure control in the first embodiment, allows a steel plate to have not only excellent fatigue properties, but also excellent crack propagation properties. The microstructure to be controlled herein is a microstructure at a position at a depth of one-fourth the thickness t of the steel plate from the surface along the thickness direction and in a longitudinal section in parallel with the rolling direction, as illustrated in Fig. 1(b). The position at a depth of one-fourth the thickness t is selected herein for evaluations at an average position in the interior of the steel plate in the thickness direction. The longitudinal section that is in parallel with the rolling direction and at the position at a depth of one-fourth the thickness t is basically a region on a line, but the actual microstructure observation is performed in a region with a certain spread around the position (see after-mentioned experimental examples).
Specifically, the microstructure at the position at a depth of one-fourth the thickness t of the steel plate is preferably controlled to include upper bainite in a fraction of 80 area percent or more, to have an effective grains size of grains of the upper bainite of 10.0 µm or less, and to have an average equivalent circle diameter of the remainder microstructure excluding the upper bainite of 3.0 µm or less. To mainly include upper bainite in the microstructure in the interior thereof, the steel plate has to contain boron (B).
The upper bainite phase is a phase that allows fine grain boundaries to be uniformly dispersed in the microstructure. This can restrain crack propagation. To offer this effect, the microstructure at the position at a depth of one-fourth the thickness t of the steel plate preferably includes upper bainite in a fraction of 80 area percent or more. The upper bainite fraction in the interior of the steel plate is more preferably 85 area percent or more, and furthermore preferably 90 area percent or more. The upper bainite fraction in the interior of the steel plate in terms of upper limit may be 100 area percent, but is typically about 98 area percent or less.
The grain size of upper bainite at the position at a depth of one-fourth the thickness t of the steel plate affects fatigue crack propagation properties and, if the steel plate has a larger grain size of upper bainite (includes coarse upper bainite grains), may fail to sufficiently restrain crack propagation. To eliminate or minimize this, grains preferably have an average length in the thickness direction, namely, an effective grain size of 10.0 µm or less, where the grains are each defined as a region surrounded by high-angle grain boundaries having a misorientation between adjacent upper bainite grains of 15° or more. The lower limit of the grain size is not specified, because the smaller the grain size is, the better crack propagation properties are. The effective grain size is more preferably 8 µm or less, and furthermore preferably 7 µm or less.
The remainder microstructure excluding the upper bainite is preferably controlled to have an average equivalent circle diameter of 3.0 µm or less. This is because the remainder microstructure, if having an average equivalent circle diameter greater than 3.0 pm, may cause the steel plate to suffer from significant reduction in other properties such as toughness. The remainder microstructure basically includes martensite and MA, as with the surface layer. These hard phases as the remainder microstructure can contribute to lower crack propagation rate. The average equivalent circle diameter in terms of lower limit is about 0.5 µm or more.
In addition, the inventors focused attention on kernel average misorientation (KAM) and grain average misorientation (GAM) in grains at the position at a depth of one-fourth the thickness t of the steel plate. The KAM is a misorientation between each measurement point and an adjacent point in one grain. The GAM is an average of KAMs in the one grain.
Figs. 4(a), 4(b), and 4(c) are conceptual diagrams respectively of grain boundaries, KAM, and GAM The hexagons in figs. 4(a), 4(b), and 4(c) represent EBSP measurement points. The periphery of a region indicated with a thick line in Fig. 4(a) is a high-angle grain boundary having a misorientation of 15° or more, and a region surrounded by the periphery is defined as a "grain".
The KAM is an average of misorientations in the one grain. Fig. 4(b) schematically illustrates how to determine the KAM. For example, the measurement point 1 is in contact with three measurement points in one grain, i.e., n=3. The KAM as the average of misorientations (numerical values in squares) between the measurement points is calculated as 0.5. The KAM may be calculated according to the expression: $KAM = \frac{\sum_{i = 1}^{n} a_{i}}{n}$
The GAM is an average of RAMs in one grain. Fig. 4(c) schematically illustrates how to determine the GAM. In an example in Fig. 4(c), there are nine measurement points, ie., m = 9, and the GAM as an average of KAMs in one grain is calculated as 0.64. The GAM may be calculated according to the expression: $GAM = \frac{\sum_{i = 1}^{m} {KAM}_{i}}{m}$
The crack propagation may be restrained when both grains each having a large misorientation in the grain, and grains each having a small misorientation in the grain are appropriately dispersed in the microstructure. To effectively restrain the crack propagation, grains having a GAM greater than 1° are preferably present in a proportion of 20% or more, more preferably 30% or more, and furthermore preferably 40% or more, of the entire microstructure. In contrast, grains having a large grain misorientation, if present in an excessively large proportion, may weaken the effect of restraining the crack propagation by the presence of grains having different misorientations as a mixture. To eliminate or minimize this, the proportion in area percent in terms of upper limit is preferably 80% or less, more preferably 70% or less, and furthermore preferably 60% or less.
The control of the proportion of grains having a GAM greater than 1° in the interior of the steel plate within the range specified in the present invention may be performed typically, but not limitatively, by allowing the upper bainite to be a mixed microstructure of bainitic ferrite and granular bainitic ferrite, where the bainitic ferrite has a large misorientation, and the granular bainitic ferrite has a small misorientation.
To give such a mixed microstructure of bainitic ferrite and granular bainitic ferrite, it is preferred to control the steel plate to meet conditions specified by Expression (2) and Expression (3), in addition to allow the steel plate to have a chemical composition within the range specified in the present invention and to control the production method of the steel plate in a manner mentioned later. Expressions (2) and (3) are expressed as follows:

where [Si], [Mn], [Ni], [Cu], [Ti], [N], [Cr], [Nb], [C], and [Mo] represent contents (in mass percent) respectively of Si, Mn, Ni, Cu, Ti, N, Cr, Nb, C, and Mo, and, when at least one of the term {[Ti] - 3.4[N]} and the term {[Nb] - 7.7[C]} is negative, the calculation according to the expression is performed as treating the at least one term as "zero (0)".
The elements indicated in Expression (2) are elements having low carbide formation ability. With an increasing left-side value of Expression (2), all the transformation curves of ferrite, bainitic ferrite, and granular bainitic ferrite shift to a longer time side. Specifically, bainitic ferrite and granular bainitic ferrite are more easily form with an increasing left-side value of Expression (2), assuming that the cooling rates are identical.
In contrast, the elements indicated in Expression (3) are elements having high carbide forming ability. With an increasing left-side value of Expression (3), only the transformation curves of ferrite and granular bainitic ferrite shift toward a longer time side, but the transformation curve of bainitic ferrite changes (shifts) little. Specifically, with an increasing left-side value of Expression (3), bainitic ferrite, which has a large misorientation, is more readily formed, as compared with granular bainitic ferrite.
The control of the left-side values of Expression (2) and Expression (3) allows the microstructure to be a mixed microstructure of bainitic ferrite and granular bainitic ferrite and to include these phases in appropriate proportions.
The left-sade value of Expression (2) and the left-side value of Expression (3) may be adjusted as appropriate in consideration of the proportion of grains having a GAM greater than 1°, are not limited, but are preferably 40 or more and 2 or less, respectively. The Left-side value of Expression (2) is more preferably 45 or more, and furthermore preferably 50 or more. The left-side value of Expression (3) is mare preferably 1.5 or less, and furthermore preferably 1.0 or less. The upper limits of the left-side values of Expression (2) and Expression (3) are inevitably determined by the ranges of contents of the elements.
To ensure the microstructure morphology as mentioned above, the steel plate is controlled to meet the conditions for the control of the microstructure morpbology in the surface layer, and, in addition, the cumulative compression reduction and the compression reduction in a non-recrystallization temperature range upon hot rolling are preferably controlled as fellows.

Cumulative compression reduction in the entire hot rolling process: 80% or more
Compression reduction in the non-recrystallization temperature range: less than 85%
Average cooling rate from the finish-rolling temperature or 800°C, whichever is lower, down to 600°C or lower: 3.0°C/second or more

To control the upper bainite in the interior of the steel plate to have an effective grain size of 10.0 µm or less, the cumulative compression reduction during the hot rolling process may be increased and is preferably 80% or more. The hot rolling, if performed at an insufficient cumulative compression reduction, may allow the surface layer microstructure to be fine, but may fail to allow the microstructure in the interior of the steel plate to be sufficiently fine and fail to sufficiently reduce the crack propagation rate. The cumulative compression reduction is more preferably 85% or more.
To surely give an upper bainite fraction of 80 area percent or more at the position at a depth of one-fourth the thickness t, the cooling after the hot rolling is preferably performed at an average cooling rate of 3.0°C/second or more, and more preferably 5°C/second or more, where the cooling is from the finish-rolling temperature or 800°C, whichever is lower, down to 600°C or lower.
The hot rolling, if performed at an excessively high compression reduction in the non-recrystallization temperature range, may cause ferrite nucleation sites to increase in number, cause ferritic transformation to readily occur, and cause the upper bainite fraction to decrease. This fail to give sufficient effects of lowering the crack propagation rate. According, excessive compression reduction in the non-recrystallization temperature range is to be avoided so as to surely give an upper bainite fraction in the interior of the steel plate of 80 area percent or more. From this viewpoint, the cumulative compression reduction in the non-recrystallization temperature range may be controlled to be preferably less than 85%, and more preferably 80% or less.
This application claims priority to Japanese Patent Application No. 2014-106307 filed on 22 May 2014 and to Japanese Patent Application No. 2015-086047 filed on 20 April 2015 . The entire contents of these applications are incorporated herein by reference. Examples
The present invention will be illustrated in further detail with reference to several examples (experimental examples) below. It should be noted, however, that the examples are by no means intended to limit the scope of the invention; that various changes and modifications can naturally be made therein without deviating from the spirit and scope of the invention as described herein; and all such changes and modifications should be considered to be within the scope of the invention.

Experimental Example 1

Ingots of steels having chemical compositions corresponding to Steels A to W as given in Table 1 were made via melting and casting according to a common ingot-making technique, subjected to hot rolling under conditions of rolling condition types "a" to "l" given in Table 2, and yielded steel plates having a thickness of 20 mm. In Table 1, an element indicated with "-" was not added; and the symbol "[Si]+[Cu]" refers to the total content of Si and Cu. The Ar₃ transformation temperatures given in Table 1 are values determined according to Expression (5). In Table 2, the term "entire hot rolling process cumulative compression reduction" refers to the cumulative compression reduction in the entire hot rolling process.

The steel plates were each subjected to measurements of the microstructure and effective grain size of the steel plate, the size of the remainder microstructure as a second phase, the tensile strength, the fatigue properties, and the dislocation density p, according to procedures as fellows. Test specimens in all the measurements were sampled so that the measurement position be a position 3 mm deep from the steel plate surface.

Steel Plate Surface Layer Microstructure

A sample was cut out at a position 3 mm deep from the steel plate surface so as to expose a plane in parallel with the rolling direction of the steel plate and in perpendicular to the steel plate surface. This was polished using wet emery papers of#150 to #1000 and was then polished to a mirror-smooth state using a diamond abrasive as an abrasive. The mirror-smooth test specimen was etched with 2% nitric acid-ethanol solution, ie., Nital solution, the etched test specimen was observed in three view fields in an observation area of 3.71× 10^-2 mm² at 400-fold magnification, images of which were taken and analyzed using an image analyzing software Image Pro Plus ver. 4.0 supplied by Media Cybernetics so as to fractionate phases in the microstructure. Values in the three view fields were averaged, and the average was defined as the area percentages of the individual phases. The observation area was such that one view field in a size of 166 µm in the thickness direction and 222.74 µm in the rolling direction was defined around the position 3 mm deep from the steel plate surface.

Ferrite and Upper Bainite Effective Grain Sizes in Surface Layer

The effective grain size of ferrite and/or upper bainite was analyzed at a position 3 mm deep from the steel plate surface in a longitudinal section in parallel with the rolling direction of the steel plate. The measurement was performed by scanning electron microscope (SEM)-electron backscatter pattern analysis (EBSP). Specifically, a grain size was measured, where the "grain" is defined as a region surrounded by a high-angle grain boundary having a misorientation between adjacent grains of 15° or more. The measurement was performed using an EBSP system (trade name OIM) supplied by TEX SEM Laboratories in combination with a SEM, in a measurement area of 200 µm by 200 µm at a measurement step (interval) of 0.5 µm. A measurement point having a confidence index of less than 0.1 was excluded from the analysis object, where the confidence index indicates the reliability of a measurement orientation The cut lengths of the grain boundaries thus determined were measured at 100 points in the thickness direction, and an average of the cut lengths was defined as the effective grain size. However, a measurement with an effective grain size of2.0 µm or less was determined as a measurement noise and excluded. The observation area was determined as a region around the position 3 mm deep from the steel plate surface with a spread of 100 µm on both sides in the thickness direction.

Remainder Microstructure Size

The size of the remainder microstructure excluding the ferrite and the upper bainite was determined in the following manner. A sample was cut out at a position 3 mm deep from the steel plate surface so as to expose a plane in parallel with the rolling direction of the steel plate and in perpendicular to the steel plate surface. This was polished using wet emery papers of#150 to#1000 and was then polished to a mirror-smooth state using a diamond abrasive as an abrasive. The mirror-smooth test specimen was etched with 2% nitric acid-ethanol solution, i.e., Nital solution, the etched test specimen was observed in an observation area of 3.71× 10² mm² at 400-fold magnification. The observation area was determined as a region around the position 3 mm deep from the steel plate surface with a spread of 100 µm on both sides in the thickness direction. Images of the observed test specimen were taken and analyzed using the image analyzing software, an area per one grain of the remainder microstructure was calculated, and the equivalent circle diameter of grains of the remainder microstructure was determined from the calculated area. In this experimental example, measurements in three view fields were averaged, and the average was defined as the equivalent circle diameter.
Of the remainder microstructure, the area percentage of MA was determined in the following manner. The mirror-smooth test specimen after polishing to a mirror-smooth state was etched with a LePera etchant and observed in an observation area of 3.71× 10^-2 mm² at 400-fold magnification. A phase corroded to white was defined as the MA, images of which were taken and analyzed using the image analyzing software to fractionate phases. Measurements in five view fields were averaged, and the average was defined as the area percentage of MA. The LePera etchant was a 5:6:1 mixture of a solution A, a solution B, and ethanol. The solution A was a solution of 3 g of picric acid in 100 ml of ethanol. The solution B was a solution of 1 g of sodium disulfite in 100 ml of distilled water.

Tensile Strength

The tensile strength TS was measured by sampling a tensile test specimen having a thickness of4 mm and a gauge length of 35 mm from each steel plate at a position 2 to 6 mm deep from the steel plate surface, and subjecting the test specimen to a tensile test according to JIS Z 2241:2011.

Fatigue Properties

The fatigue properties were determined in the following manner. A steel plate sample having a thickness of 4 mm was cut out from each steel plate at a position 2 to 6 mm deep from the steel plate surface, from which a test specimen as illustrated in Fig. 2 was prepared. The test specimen surface was polished with emery papers up to#1200 to eliminate or minimize influence of surface conditions. The resulting test specimen was subjected to a fatigue test using a servo-electric hydraulic fatigue tester supplied by INSTRON Co., Ltd. under conditions as follows.

Testing environment:	room temperature, in the air
Control type:	load control
Control waveform:	sinusoidal wave
Stress ratio R:	-1
Testing frequency:	20 Hz
Number of cycles tocomplete testing:	5000000

The fatigue properties are affected by the tensile strength TS. To eliminate or minimize the influence of the tensile strength, a 5000000-cycle fatigue limit to tensile strength ratio was determined, and a sample, when having a 5000000-cycle fatigue limit to tensile strength ratio of greater than 0.51, was accepted herein. The 5000000-cycle fatigue limit to tensile strength ratio is a value determined by dividing a 5000000-cycle fatigue limit by the tensile strength TS. The 5000000-cycle fatigue limit was determined in the following manner. Each test specimen was subjected to a fatigue test at such a stress amplitude as to give a stress amplitude σa to tensile strength TS ratio (σa/TS) of 0.51, and whether the test specimen underwent rupture upon the 5000000th cycles was examined.

Dislocation Density ρ

The dislocation density ρ was determined by subjecting each sample to X-ray diffractometry to determine a half peak width (full-width at half maximum) of α-Fe, and calculating the dislocation density from the half peak width. The measurement conditions and principle will be illustrated below. An analyzer used herein was an X-ray diffractometer RAD-RU300 (trade name, supplied by Rigaku Corporation), with a cobalt tube as a target The half peak width was calculated from the results of X-ray difiractometry via peak fitting, based on which the dislocation density ρ was calculated according to Expression (6):
where ε represents the strain; and b represents the Burgers vector (=0.25× 10^-9 m).
The strain ε is a value calculated according to the Hall method based on Expression (7) and Expression (8):

where β represents the true half peak width (in radian (rad)); θ represents the Bragg angle (in degree (°)), λ represents the wavelength (in nanometer (nm)) of incident X ray; D represents the crystal size (in nm); β_m represents the measured half peak width; and β_s represents the half peak width (apparatus constant) of a sample with no strain. The true half peak width β was calculated from β_m and β_s according to Expression (8), and this was substituted into Expression (7), based on which βcos θ/λ - sin θ/λ was plotted. Three points, ie., (110), (211), and (220) points were fitted by the method of least squares. The strain ε was calculated from the slope (2ε) of the fitting line and was substituted into Expression (6) to calculate the dislocation density ρ.
Table 3 presents, of each steel plate, the microstructure, the effective grain size, the remainder microstructure size, the tensile strength TS, the fatigue properties, and the dislocation density ρ.
These results indicate as follows. Specifically, Test Nos. 1 to 17 were produced under appropriately controlled conditions using steels having appropriately controlled chemical compositions, met the conditions in the surface layer specified in the present invention, and offered excellent fatigue properties.
In contrast, Test Nos.18 to 34 were samples failing to meet at least one of the conditions specified in the present invention, and each had poor fatigue properties. Among them, Test No. 18 employed a steel plate derived from Steel K having a low carbon content and failed to have a tensile strength TS at the predetermined leveL Accordingly, other properties than the microstructure were not evaluated in this sample. Test No. 19 employed a steel plate derived from Steel L having an excessively high carbon content and had an excessively high tensile strength TS. Accordingly, other properties than the microstructure were not evaluated in this sample.
Test No. 20 employed a steel plate derived from Steel M not meeting the condition that "the total content of Si and Cu is 0.3% or more", felled to restrain cell structure formation of dislocations, and had inferior fatigue properties. Test No. 21 employed a steel plate derived from Steel N having an excessively high Si content, had an excessively large size of the remainder microstructure, and had inferior fatigue properties.
Test No. 22 employed a steel plate derived from Steel O having an excessively high Mn content, had a high tensile strength TS and a high dislocation density p, and offered inferior fatigue properties. Test No. 23 employed a steel plate derived from Steel P having an excessively low Mn content, failed to have a tensile strength TS at the predetermined level, had an excessively large effective grain size, and offered inferior fatigue properties.
Test No. 24 employed a steel plate derived from Steel Q having an excessively high Cu content, had an excessively high dislocation density ρ, and offered inferior fatigue properties. Test No. 25 employed a steel plate derived from Steel R having an excessively high Ni content and not meeting the condition: [Ni]/[Cu] < 1.2. This sample had an excessively high dislocation density ρ, and offered inferior fatigue properties.
Test No. 26 employed a steel plate derived from Steel S having an excessively high Cr content, had an excessively high dislocation density ρ, and offered inferior fatigue properties. Test No. 27 employed a steel plate derived from Steel T having an excessively high Mo content, had an excessively high dislocation density ρ, and offered inferior fatigue properties.
Test No. 28 employed a steel plate derived from Steel U having an excessively high V content, had an excessively high dislocation density ρ, and offered inferior fatigue properties. Test No. 29 employed a steel plate derived from Steel V having a bainite transformation start temperature Bs lower than 640°C, had an excessively high dislocation density ρ, and offered inferior fatigue properties.
Test No. 30 was a sample produced via rolling under conditions of the type g with an excessively high heating temperature in hot rolling, had an excessively large effective grain size, and offered inferior fatigue properties. Test No. 31 was a sample produced via rolling under conditions of the type h with an excessively low cumulative compression reduction in hot rolling, had an excessively large effective grain size, and offered inferior fatigue properties.
Test No. 32 was a sample produced via rolling under conditions of the type i with an excessively low finish-rolling temperature, had an excessively high dislocation density ρ, and offered inferior fatigue properties. Test No. 33 was a sample produced via rolling under conditions of the type j with an excessively high average cooling rate down to 600°C, had an excessively high dislocation density ρ, and offered inferior fatigue properties. Test No. 34 was a sample produced via rolling under conditions of the type k with an excessively low cooling step temperature, had an excessively high dislocation density ρ, and offered inferior fatigue properties.

Experimental Example 2

The steel plates of Test Nos.1 to 17 given in Table 3 were each subjected to evaluations of the fraction and effective grain size of upper bainite in the interior of the steel plate, namely, at the position at a depth of one-fourth the thickness t of the steel plate; and the size of the remainder microstructure as a second phase, according to procedures similar to those in Experimental Example 1. Test specimens were sampled by procedures similar to the above procedures, except for sampling them at a position at a depth of one-fourth the thickness t of the steel plate. These steel plates were also subjected to measurements of the portion of grains having a GAM of 1° or more, and the crack propagation rate by methods as follows.

Measurement of Proportion of Grains Having a GAM of 1° or more

The proportion of grains each having a GAM of 1° or more at the position at a depth of one-fourth the thickness t of the steel plate was measured by SEM-EBSP. Specifically, the grain size was measured while defining the "grain" as a region surrounded by a high-angle grain boundary having a misorientation between adjacent grains of 15° or more. The measurement was performed using the EBSP equipment (trade name OIM) supplied by TEX SEM Laboratories in combination with a SEM. The measurement was performed in a measurement area of 200 µm by 200 µm at a measurement step (interval) of 0.5 µm. The measurement area was a region around the position at a depth of one-fourth the thickness t of the steel plate with a 100 µm spread on both sides in the thickness direction. A measurement point having a confidence index CI of less than 0.1 was excluded from the analysis object, where the confidence index indicates the reliability of a measurement orientation. The GAM as an average of KAMs in one grain was determined, where each KAM represents the misorientation between a measurement point and an adjacent point Thus, grains each having a GAM of 1° or more were identified The term "grain" herein refers to a grain into which high strain is introduced. The measurement was performed in three view fields per one steel type, and an average of area fractions of grains having a GAM of 1° or more was calculated.

Crack Propagation Rate Measurement

The crack propagation rate was measured by preparing a compact tension test specimen, subjecting the compact tension test specimen to a fatigue crack propagation test in accordance with American Society for Testing Materials (ASTM) standard E647 using a servo-electric hydraulic fatigue tester under conditions as follows. The compact tension test specimen was sampled at the position at a depth of one-fourth the thickness t of the steel plate and had dimensions illustrated in Fig. 3. The crack length was determined using the compliance method.

Testing environment: room temperature, in the air

Control type: load control

Control waveform: sinusoidal wave

Stress ratio R: -1

Testing frequency: 5 to 20 Hz
In this test, a value in a stable growth region at ΔK of 20 (MPa m^1/2), in which Paris' law holds, was used as a representative value for evaluation, where ΔK is specified by Expression (9). A sample having a crack propagation rate of 5.0× 10^-5 mm/cycle or less at ΔK of 20 (MPa m^1/2) was evaluated as having excellent crack propagation properties. Expression (9) is expressed as follows:
where "a" represents the crack length (in millimeter (mm)); n represents the number of cycles (in cycle); and C and p are constants independently determined by conditions such as materials and load.
Results of these are presented in Table 4. In the "evaluation" in Table 4, a sample having a crack propagation rate of 5.0×10^-5 mm/cycle or less was evaluated as having excellent fatigue crack propagation properties and is indicated with "A"; and a sample having a crack propagation rate of 40×10^-5 mm/cycle or less was evaluated as having still further excellent fatigue crack propagation properties and is indicated with "AA".
These results indicate and demonstrate as follows. Specifically, Test Nos.1 to 6, 10, 11, 13, and 15 employed steels having appropriately controlled chemical compositions and were produced under appropriately controlled production conditions, met the preferred conditions in the interior of the steel plate, had a crack propagation rate of 40× 10^-5 mm/cycle or less, and had still further excellent fatigue crack propagation properties.

Claims

A steel plate comprising:
C in a content of 0.02 to 0.10 mass percent;

Mn in a content of 1.0 to 2.0 mass percent;

Nb in a content of greater than 0 mass percent to 0.05 mass percent;

Ti in a content of greater than 0 mass percent to 0.05 mass percent;

Al in a content of 0.01 to 0.06 mass percent;

Cu in a content of 0.1 to 0.6 mass percent;

Ni in a content of greater than 0 mass percent to 0.6 mass percent, where a ratio [Ni]/[Cu] of the Ni content [Ni] to the Cu content [Cu] is 0.5 to less than 1.2; and

B in a content of greater than 0 mass percent to 0.005 mass percent,

optionally comprising Si in a content of 0.1 to 0.6 mass percent, where a total content of Si and Cu is 0.3 mass percent or more;

optionally comprising, in chemical composition, at least one element selected from the group consisting of:
V in a content of greater than 0 mass percent to 0.5 mass percent;

Cr in a content of greater than 0 mass percent to 1.0 mass percent or less; and

Mo in a content of greater than 0 mass percent to 0.5 mass percent, with the remainder consisting of iron and inevitable impurities,

wherein a microstructure of a surface layer of the steel plate comprises at least one of ferrite and upper bainite in a total fraction of 80 area percent to 100 area percent as determined with the measurement method defined in the description under the heading "Steel Plate Surface Layer Microstructure", wherein, of the microstructure of the surface layer, the remainder microstructure excluding the ferrite and the upper bainite optionally has a fraction of martensite-austenite constituent of 0 area percent to 5 area percent as determined with the measurement method defined in the description under the heading "Remainder Microstructure Size",

wherein grains of the at least one of ferrite and upper bainite have an effective grain size of 2 µm to 10.0 µm, wherein the effective grain size refers to an average length of the grains in the thickness direction as determined with the measurement method defined in the description under the heading "Ferrite and Upper Bainite Effective Grain Sizes in Surface Layer",

wherein, of the microstructure of the surface layer, grains of a remainder microstructure excluding the ferrite and the upper bainite have an average equivalent circle diameter of 0.5 µm to 3.0 µm as determined with the measurement method defined in the description under the heading "Remainder Microstructure Size",

wherein the steel plate has a dislocation density ρ of 5.0× 10¹³ m^-1 to 2.5× 10¹⁵ m^-1 as determined by X-ray diffractometry with the measurement method defined in the description under the heading "Dislocation Density ρ",

wherein a microstructure at a position at a depth of one-fourth the thickness t of the steel plate from the surface along a thickness direction comprises upper bainite in a fraction of 80 area percent to 100 area percent in a longitudinal section in parallel with a rolling direction,

wherein grains of the upper bainite in the microstructure at the position at a depth of one-fourth the thickness t have an effective grain size of 10.0 µm or less, and wherein, of the microstructure at the position at a depth of one-fourth the thickness t, grains of the remainder microstructure excluding the upper bainite have an average equivalent circle diameter of 0.5 µm to 3.0 µm.
The steel plate according to claim 1,
wherein the microstructure at the position at a depth of one-fourth the thickness t comprises grains each having a grain average misorientation (GAM) of 1° or more in a fraction of 20 area percent to 80 area percent, where the grain average misorientation is determined per one grain by electron backscatter pattern observation with the measurement method defined in the description under the heading "Measurement of Proportion of Grains Having a GAM of 1° or more".
The steel plate according to claim 2,
wherein the steel plate further meets conditions specified by Expression (2) and Expression (3): $35 [Si] + 18 [Mn] + 17 [Ni] + 16 [Cu] \geq 40$
$21 \times \{[Ti] - 3.4 [N]\} + 19 [Cr] + 11 \times \{[Nb] - 7.7 [C]\} + 10 [Mo] \leq 2$
wherein [Si], [Mn], [Ni], [Cu], [Ti], [N], [Cr], [Nb], [C], and [Mo] represent contents (in mass percent) respectively of Si, Mn, Ni, Cu, Ti, N, Cr, Nb, C, and Mo, and, when at least one of the term {[Ti] - 3.4[N]} and the term {[Nb] - 7.7[C]} is negative, calculation according to the expression is performed as treating the at least one term as "zero (0)".
The steel plate according to any one of claims 1 to 3,
wherein the steel plate has a bainite transformation start temperature Bs of 640°C or higher, where the bainite transformation start temperature is calculated from the chemical composition according to Expression (1): $Bs (°C) = 830 - 270 [C] - 90 [Mn] - 37 [Ni] - 70 [Cr] - 83 [Mo]$
wherein [C], [Mn], [Ni], [Cr], and [Mo] represent contents (in mass percent) respectively of C, Mn, Ni, Cr, and Mo.