EP2860276A1

EP2860276A1 - Steel plate

Info

Publication number: EP2860276A1
Application number: EP13883843.8A
Authority: EP
Inventors: Naoki Saitoh; Katsumi Kurebayashi; Yasunori Takahashi
Original assignee: Nippon Steel and Sumitomo Metal Corp
Current assignee: Nippon Steel Corp
Priority date: 2013-08-13
Filing date: 2013-08-13
Publication date: 2015-04-15
Anticipated expiration: 2033-08-13
Also published as: EP2860276B1; KR101542709B1; EP2860276A4; WO2015022729A1; JP5510620B1; KR20150029615A; CN104520463B; JPWO2015022729A1; CN104520463A

Abstract

In a steel sheet, an A value is 4.5% or less, a Pcm value is 0.25% or less, a yield strength is 460 N/mm² to 580 N/mm², a tensile strength is 550 N/mm² to 670 N/mm2, a difference between a hardness at 1/8t, which is a 1/8 thickness position from a surface of the steel plate in a through-thickness direction, and a hardness at 1/2t, which is a 1/2 thickness position from the surface of the steel plate in the through-thickness direction, is 20 or less by Vickers hardness, and an average grain size at 1/8t is 35 µm or less.

Description

[Technical Field of the Invention]

The present invention relates to a thick high tensile strength steel plate having a desired thickness of 80 mm or more and having a yield strength of 460 N/mm² to 580 N/mm² and a tensile strength of 550 N/mm² to 670 N/mm² which are desired strengths for use in welded structures such as buildings, construction machines, offshore structures, large cranes for ships, or civil engineering structures, in which properties in a through-thickness direction thereof are uniform, and weldability, base metal toughness, and HAZ (heat affected zone) toughness are superior.

[Related Art]

Recently, along with an increase in the size of structures, a thick high tensile strength steel plate has been widely used in welded structures such as buildings, construction machines, offshore structures, large cranes for ships, or civil engineering structures.
When a complex welded structure is designed and constructed by using a thick high tensile strength steel plate in a large structure, differences in strength and toughness in a through-thickness direction are not preferable from the viewpoint of estimating deformation behavior and fracture behavior thereof in order to secure a reasonably high level of safety. Therefore, a thick high tensile strength steel plate having uniform properties in a through-thickness direction thereof is desired.
In most cases, a thick high tensile strength steel plate is used in a portion of a large offshore structure, a large crane, or the like, where a high degree of safety is required. The most concerning problem regarding fractures of structures is that brittle fractures are initiated from a welded joint having a weld defect or the like. Accordingly, superior weldability is required to prevent the occurrence of defects in a weld zone, and high heat affected zone toughness (hereinafter, referred to as "HAZ toughness") is required to prevent brittle fractures in most cases.
In particular, a thick high tensile strength steel plate having a thickness of 80 mm or more is typically produced by adding an appropriate amount of an alloy element for improving hardenability such as C, Mn, Cr, Mo, or V to impart a predetermined strength up to a mid-thickness of the steel plate and quenching and tempering it. It is well known that, during quenching, due to a difference in cooling rate in a through-thickness direction, strength and toughness change depending on the thickness in the through-thickness direction in a region from a surface to a mid-thickness of the steel plate. In addition, as the thickness of the steel plate increases, not only a difference in the cooling rate of quenching but also a difference in the heating rate of quenching increases between the surface and the mid-thickness. The time of the surface of the steel plate being held at a high temperature is longer than that of the mid-thickness and grains are likely to be coarse as compared to those in the mid-thickness. When there is a difference in grain size between the vicinity of the surface and the mid-thickness, properties including strength may be different.
Typically, in most steel standards, properties of a 1/4 thickness position from a surface of a steel plate in a through-thickness direction thereof, that is, a position which is at a distance of the thickness of 1/4 from the surface of the steel plate to the center in the through-thickness direction (hereinafter, referred to as "1/4t") are specified. However, when the thickness of a steel plate increases and high safety is required to prevent fracture in an offshore structure or the like, stable and high properties are also required in a 1/2 thickness position from a surface to a mid-thickness of the steel plate (hereinafter, referred to as "1/2t").
From the above-described viewpoints, for a thick high tensile strength steel plate which will be used in a large structure, superior weldability, high base metal toughness and HAZ toughness, and elimination of non-uniformity in the through-thickness direction unique to the thick high tensile strength steel plate are important. Many studies clearly show that weldability is determined by the alloy composition. For example, weldability can be evaluated based on an index such as a Pcm value. In most cases, superior weldability with which preheating is unnecessary can be achieved by limiting the amount of an alloy element having high hardenability such as Cr or Mo and controlling the Pcm value to be, for example, 0.25% or less. Accordingly, in order to secure superior weldability, as described above, it is important to secure strength without adding the element for improving hardenability if possible. As one such technique of the related art, a high tensile strength steel plate which contains a large amount of Cu is disclosed.
For example, Patent Documents 1 and 2 disclose a method of producing a high tensile strength steel plate which contains 0.6% to 1.5% of Cu and a method of producing a high tensile strength steel plate which contains 0.5% to 2.0% of Cu. In these methods, controlled rolling is performed during hot-rolling, in principle, on the assumption of the use of a thermo-mechanical control process along with accelerated cooling after rolling. Therefore, the production methods disclosed in Patent Documents 1 and 2 are not suitable for the production of a thick high tensile strength steel plate having a thickness of 80 mm or more. Further, when these methods are used, the size of microstructures or the like in the vicinity of the surface and the center in the through-thickness direction is increased due to the effect of controlled rolling and the like. Consequently, properties thereof in the through-thickness direction may be significantly changed.
Patent Document 3 discloses a method of producing steel having high toughness and high strength (high tensile strength steel plate) which contains 0.5% to 4.0% of Cu in which elongation properties are superior and the tensile strength is 686 MPa or higher. The target of Patent Document 3 is a high strength steel having a tensile strength of 686 MPa or higher which exceeds the assumption of the present invention and is a high strength steel having high hardenability to which an alloy element such as Cr, Mo, or V can be added. Therefore, the production method disclosed in Patent Document 3 has a problem in material uniformity in a through-thickness direction and thus cannot be adopted as a method for solving the target problem of the present invention.
Patent Document 4 discloses a high tensile strength steel plate having superior weld zone toughness which contains 0.8% to 1.5% of Cu. Although Cu and Ni are added thereto, this high tensile strength steel plate is assumed to have a thickness of 77 mm as shown in Examples of Patent Document 4, and the target thereof is different from that of the present invention in which the desired thickness is 80 mm or more. In addition, Patent Document 4 clearly describes a technique in which, in order to produce the high tensile strength steel plate, rolling is performed while restricting a total rolling reduction at 900°C or lower and direct water cooling is performed after rolling. Therefore, there is a serious concern regarding material uniformity in a through-thickness direction. In addition, a N/Al ratio is limited to be within a range of 0.3 to 3.0, but the Al content is 0.013% or less as disclosed in Examples. As a result, ordinary deoxidation by Al cannot be performed. Therefore, the production method disclosed in Patent Document 4 slightly deviates from an ordinary production method of the related art and has the potential problems of low stability for production and high cost.
Patent Documents 5, 6, and 7 discloses methods of producing steel for high heat input welding having superior low-temperature toughness which contains 0.2% to 2.0% of Cu. These steel plates are characterized in that the S content is controlled to be 0.003% to 0.008%. By adding S to steel and controlling the S content to be in the above-described range, fine MnS precipitates in the steel, and superior HAZ toughness at high heat input welding is obtained. With these techniques, a certain level of effect can be obtained at high heat input welding, but the target thereof is a thin steel plate having a thickness of about 32 mm, which is significantly different from that of the present invention. Further, particularly in a thick high tensile strength steel plate, the addition of S accelerates the production of an MnS inclusion which is highly likely to adversely affect toughness. Therefore, the techniques disclosed in Patent Documents 5 to 7 are not superior methods in relation to the production of a thick high tensile strength steel plate.
Patent Document 8 discloses a high tensile strength steel plate having superior CTOD properties which contains 0.70% to 1.75% of Cu. However, the strength level of this steel plate is in the order of 780 MPa (the tensile strength is 780 MPa or higher), which is significantly different from the desired strength of the present invention. Further, since this steel plate contains 0.005% to 0.0015% of B, an increase in the hardness in the vicinity of a surface of the steel plate in a through-thickness direction is extremely large. Therefore, it is presumed that the steel plate disclosed in Patent Document 8 has a large difference in strength in the through-thickness direction. Further, in this steel plate, the Al content is extremely small at 0.01% or less, and ordinary deoxidation by Al cannot be performed. Therefore, the production method disclosed in Patent Document 8 slightly deviates from an ordinary production method of the related art, is poor in stability, and requires high cost. For this reason and the like, the production method disclosed in Patent Document 8 is not suitable for solving the problem of the present invention.
As described above, Cu addition is the technique which has been widely used in the related art. However, there is no technique of the related art in which, in a thick high tensile strength steel plate having a thickness of, for example, more than 80 mm, material uniformity in a through-thickness direction can be secured substantially without adding an alloy element such as Cr, Mo, or V thereto.

[Prior Art Document]

[Patent Document]

[Patent Document 1] Japanese Examined Patent Application, Second Publication No. H07-81164
[Patent Document 2] Japanese Unexamined Patent Application, First Publication No. H05-179344
[Patent Document 3] Japanese Unexamined Patent Application, First Publication No. H05-186820
[Patent Document 4] Japanese Patent No. 4432905
[Patent Document 5] Japanese Unexamined Patent Application, First Publication No. H02-254118
[Patent Document 6] Japanese Unexamined Patent Application, First Publication No. H02-250917
[Patent Document 7] Japanese Unexamined Patent Application, First Publication No. H03-264614
[Patent Document 8] Japanese Unexamined Patent Application, First Publication No. 2001-335884

[Disclosure of the Invention]

[Problems to be Solved by the Invention]

According to the present invention, it is possible to provide a thick high tensile strength steel plate having a thickness of, for example, 80 mm or more and having a yield strength of 460 N/mm² to 580 N/mm² and a tensile strength of 550 N/mm² to 670 N/mm², in which properties in a through-thickness direction thereof are uniform, and weldability, base metal toughness, and HAZ toughness are superior, which cannot be achieved in the related art.

[Means for Solving the Problem]

The present inventors have carried out many experiments on a method of producing a thick high tensile strength steel plate. As a result, it was found that, in order to secure high weldability of base metal and HAZ toughness, it is important to control a Pcm value to be in a range of 0.25% or less and to contain substantially no Cr, Mo, V, and B which have high hardenability. In the present invention, high weldability indicates that weld cracking does not occur at 0°C during actual welding. In this case, preheating is not necessary during welding.
Further it was found that, in order to secure properties after stress relief annealing and HAZ toughness, it is efficient to contain a high concentration of Cu and a high concentration of Ni. Further, it was found that, in order to obtain a thick high tensile strength steel plate having material uniformity in a through-thickness direction, it is efficient to perform quenching and tempering, and not a TMCP (Thermo Mechanical Control Process) which is mainly used in Cu-added steel of the related art, in a state where the amounts of Cu and Ni are limited to be in a specific high concentration range.
FIG. 1 is a diagram illustrating cross-sectional hardness distributions in a through-thickness direction of two types of steel plates having a thickness of 110 mm after being quenched and tempered, the steel plates containing 1.15% of Cu and 1.81% or 3.22% of Ni. Typically, the cross-sectional hardness of a thick high tensile strength steel plate in a through-thickness direction tends to increase from the inside region to the vicinity of a surface of the steel plate. This tendency increases as the amount of the alloy element for improving hardenability increases. As can be seen from FIG. 1, when the steel containing 3.22% ofNi is compared to the steel containing 1.81% of Ni (having a Ni content of 1.81%), the range of high hardness spreads over a range from a surface to the inside of the steel plate in a through-thickness direction, and a difference (ΔHv) between a Vickers hardness of a 1/8 thickness position from the surface of the steel plate in the through-thickness direction (hereinafter, referred to as "1/8t") and a Vickers hardness at 1/2t is 38. AΔHv value of the steel containing 3.22% of Ni is significantly higher than that of the steel containing 1.81% of Ni. Here, the surface of the steel plate refers to a single surface of the steel plate, not a specific surface during rolling.
As described above, ΔHv depends on the amount of the alloy element. The experimental results of a relationship between ΔHv and the amount of the alloy element are illustrated in FIG. 2. FIG. 2 illustrates ΔHv values, which is a difference between a hardness at 1/8t and a hardness at 1/2t in a steel plate, of steel plates having a thickness of 100 mm and having different amounts of Cu and Ni. Numerical values in circles of the drawing indicate ΔHv. When the hardness of a cross-section of a steel plate in a through-thickness direction is measured, a locally high hardness range caused by center segregation may appear depending on the state of a slab in the vicinity of a mid-thickness. Since such a locally high hardness region (locally hardened portion) is an extremely small region with respect to the entire thickness of a thick high tensile strength steel plate, it is considered that this region has little effect on the strength of steel. Therefore, when a hardness distribution of a cross-section of a steel plate is measured, it is preferable that data of the above-described locally hardened portion be excluded. It can be seen from FIG. 2 that: there is a correlation between an A value (A=Cu+Ni), which is the sum of the Cu content and the Ni content, and ΔHv; and when the A value is more than 4.5%, ΔHv is more than 20. Further, it was found that, even if the Cu content is small at 1.5% or less, when the Ni content is more than 3.0%, ΔHv may also be more than 20. On the other hand, the lower limit of the A value is not particularly limited. However, from the viewpoint of securing HAZ toughness and strength described below, the lower limits of the Ni content and the Cu content are set as 1.2% and 0.7%, respectively. Accordingly, the lower limit of the A value is preferably 1.9% which is the sum of the lower limits of the Cu content and the Ni content.
Further, the present inventors have carried out an impact test in which a heat-affected zone (HAZ) at -40°C is simulated in order to investigate the effects of the Cu content and the Ni content on HAZ toughness (vE (HAZ)) which is an important factor of the present invention. The results are illustrated in FIG. 3. In an ordinary large structure, it is known that, when a Charpy absorbed energy at -40°C is 42 J or higher, the initiation of brittle fractures can be prevented. Therefore, whether or not the Charpy absorbed energy at -40°C is 42 J or higher is set as a pass/fail criterion. Numerical values in circles of FIG. 3 indicate the Charpy absorbed energy at -40°C. It can be seen from FIG. 3 that: the toughness of steel is significantly improved concurrent with an increase in the Ni content; and the Ni content is required to be 1.2% or higher in order to secure an impact test value of 42 J or more as described below. However, it was also found that, when the Cu content is more than 2.5%, even if the Ni content is 1.2% or more, toughness decreases.
As described above, HAZ toughness has a strong effect on the alloy composition (the amount of the alloy component). On the other hand, it is necessary that base metal toughness is investigated in consideration of not only the alloy composition but also microstructures, specifically, a grain size. It is necessary that a grain size is investigated at each thickness position of a thick high tensile strength steel plate having a thickness of, particularly, more than 80 mm. Typically, steel having a tensile strength of 550 N/mm² to 670 N/mm² which is assumed in the present invention has microstructures in which ferrite and bainite are mixed. Therefore, it is not easy to evaluate a grain size through the microstructure observation of the related art in which an optical microscope is used. Therefore, in the present invention, using EBSD (electron backscatter diffraction pattern) analysis which has been widely used for crystal orientation analysis, a region surrounded by grain boundaries having a crystal orientation difference of 30° or more is defined as a grain, and a circle equivalent grain size of the grain is defined as a grain size. A frequency distribution of the measured grain size is calculated, and the grain size at which a cumulative frequency from the smallest grain size is 70% is defined as an average grain size. Examples which are actually measured are illustrated in FIG. 4. FIG. 4 illustrates a cumulative frequency (%) of the grain sizes of steel containing 0.08% of C, 0.15% of Si, 1.51% of Mn, 0.008% of P, 0.0010% of S, 1.15% of Cu, 1.23% of Ni, 0.012% of Ti, 0.012% of Nb, 0.035% ofAl, and 0.0039% ofN as components. When the cumulative frequency is obtained, first, the steel in which the above-described components are melted is hot-rolled into a thickness of 140 mm and then is quenched and tempered after hot-rolling. The grain sizes of a surface (that is, a surface portion or an outermost surface) and the respective thickness positions including a 1/8t, a 2/8t (1/4t), and a 3/8t of the quenched and tempered steel plate are obtained, and a cumulative frequency (%) of the grain sizes is obtained. The grain size corresponding to a cumulative frequency of 70% is an average grain size. As can be seen from FIG. 4, in this experimental result, the average grin size at each thickness position changes depending on a machined position in a through-thickness direction of the steel plate. Approximately, the average grain sizes at the outermost surface and 1/8t are 20 µm or more, whereas the average grain sizes at 2/8t and 3/8t are 15 µm or less.
Further, the present inventors have investigated a change in toughness depending on the grain sizes defined as described above. FIG. 5 illustrates a relationship between a grain size and toughness obtained in a Charpy test, which is carried out while changing a test temperature at an interval of 20°C, in the above-described quenched and tempered steel having a thickness of 140 mm and containing 0.08% of C, 0.15% of Si, 1.51% of Mn, 0.008% of P, 0.0010% of S, 1.15% of Cu, 1.23% ofNi, 0.012% of Ti, 0.012% ofNb, 0.035% of Al, and 0.0039% ofN as components. As an index of toughness, a fracture transition temperature (vTrs) obtained in the Charpy test is used. Here, vTrs denotes a temperature at which an area ratio of brittle fractures is 50% when ductile fractures and brittle fractures are identified from characteristics of fractures of a test specimen, the area ratio of the brittle fracture to all of the fractures is measured, and a relationship between the area ratio of the brittle fractures and the test temperature is obtained. The lower the value of vTrs, the higher the toughness. The machined position of the Charpy test specimens are the same as the measurement positions of the grain sizes, and the machined direction is a direction perpendicular to a rolling direction.
In FIG 5, the vertical axis represents vTrs (toughness), and d^-1/2 of the horizontal axis represents an inverse of a square root of an average grain size. In this drawing, the higher the value of d^-1/2× 100 of the horizontal axis, the smaller the grain size.
As clearly seen from FIG. 5, there is a substantially linear correlation between vTrs and d^-1/2. This correlation corresponds to the Hall-Petch relationship of the related art. The vTrs of the horizontal axis is affected by the components of the steel plate. In particular, it is known that toughness is improved along with an increase in the Ni content. FIG. 5 illustrates a case where the Ni content is 1.23%, and this Ni content is close to 1.2% which is the lower limit of the Ni content required for improving HAZ toughness. Therefore, when the Ni content which is an alloy component having the largest effect on toughness is close to the lower limit of the range according to the present invention, the necessary grain size can be estimated from FIG. 5. Hereinafter, the details will be described.
Since fractures in an ordinary large structure are initiated from a welded joint, HAZ toughness is important in steel. However, in order to further improve the safety of the structure, not only HAZ toughness but also the toughness of the base metal (portion which is not affected by welding heat) are required to be high. Typically, it is assumed that brittle fractures are initiated from weld defects. When most of these defects are present inside a steel plate, and not on the surface where they can be easily detected, the defects have the largest effect on brittle fractures. The reason is presumed to be as follows. The defects present inside the steel plate has a low possibility of being detected and, although depending on the applied stress state, is likely to be in the most severe stress state with respect to the propagation of cracks.
Assuming that fractures are initiated from a defect of a weld zone, even if brittle cracks are initiated, in order to prevent the propagation of the brittle cracks in the base metal, it is necessary that the toughness of the base metal in the vicinity of the defect is high. It is assumed that a region from 1/8t to 7/8t which is positioned inside a steel plate is generally in a severe stress state. Accordingly, toughness required for the base metal should be defined in an inside region from the 1/8t to the center of the steel plate, rather than in the vicinity of the surface of the steel plate.
For the above-described reasons, an energy value of 42 J or higher, which is typically required as the Charpy absorbed energy at -40°C (vE-40), is necessary in the inside region of the steel plate but not in a region from the surface to 1/8t of the steel plate. Therefore, in the present invention, the grain size of the inside region from 1/8t to the center is defined.
When a transition curve obtained from a steel plate of the related art is taken into consideration, it is necessary that vTrs be -10°C or lower in order to satisfy the absorbed energy of 42 J at -40°C.
In FIG. 5, the average grain size corresponding to the vTrs of -10°C (indicated by a broken line in the drawing) is 35 µm. Accordingly, it can be seen that, when the average grain size is 35 µm or less, vTrs≤-10°C can be satisfied. Each point in FIG 5 is machined from a thickness position indicated in "( )". As described above, it is considered that a surface portion of a steel plate has little effect on the fractures of an actual structure. Therefore, in the present invention, the average grain sizes at positions excluding a region from an outermost surface to 1/8t of a steel plate are defined. Since a thick steel plate is held in a heat-treatment furnace for a long period of time, the grain size of a surface portion of the steel plate tends to be coarse as compared to that of a mid-thickness thereof. Therefore, it is particularly important to control the average grain size at 1/8t of the steel plate to be 35 µm or less. Further, by controlling the average grain size at 3/8t of the steel plate to be 35 µm or less, the average grain sizes at both 1/8t and 3/8t of the steel plate can be controlled to be 35 µm or less.
As described above, as the average grain size decreases, toughness is improved. However, it is difficult to decrease the average grain size. Therefore, the lower limit of the average grain size may be set as 5 µm, 10 µm, or 15 µm.
It is considered that, in order to improve the safety of a steel structure, higher toughness is required of the base metal in consideration of strain aging. Particularly in the case of strain aging, as a result of an investigation by the present inventors, it was found that, when a strain of about 5% is applied in a cold environment and the aging treatment is performed at 250°C (held for 2 hours), the Charpy transition temperature increases by about 15°C. Therefore, when higher toughness is required in consideration of strain aging, it is preferable that vTrs further decreases by 15°C, that is, be -25°C or lower. To that end, similarly, it can be seen from FIG. 5 that the average grain size at 1/8t of the steel plate is required to be 25 µm or less. That is, for the same reason as described above, the average grain size at 3/8t of the steel plate may be set to be 25 µm or less.
In the vicinity of the surface of the steel plate, the cooling rate of quenching is higher than that of the inside of the steel plate. Therefore, a sufficiently hardened structure is easily obtained, but the strength tends to increase. Therefore, it cannot be said that the toughness in the vicinity of the surface is always higher than that of the inside (for example, 1/4t) of the steel plate. However, as described above, when the safety of a structure relative to brittle fractures is taken into consideration, under a condition where extreme bending deformation does not occur, brittle cracks tend not to be initiated easily in the inside region of the steel plate (region from the 1/8t to the center) as opposed to the vicinity of the surface where potential cracks such as weld defects and a restraining force is low, can be easily detected. Therefore, in the present invention, it is considered that the safety of a structure is sufficiently secured by improving the toughness of the inside region from 1/8t to the center, and the average grain size of the inside region from 1/8t to the center is defined.
In a steel plate produced based on the above-described technique, the uniformity in the through-thickness direction thereof is secured, and weldability, base metal toughness, and HAZ (heat affected zone) toughness are superior. In particular, in a steel plate having a thickness of 80 mm or more, the effects are significant. However, in a steel plate having a thickness of more than 200 mm, the cooling rate of a mid-thickness significantly decreases, which causes the coarsening of microstructures. As a result, predetermined strength and toughness may not be satisfied. Accordingly, the thickness of a steel plate produced according to the present invention may be controlled to be 200 mm or less. Optionally, the upper limit of the thickness may be set as 175 mm, 150 mm, or 125 mm. The lower limit of the thickness may be set as 90 mm or 100 mm.
In this way, unlike the thick high tensile strength steel plate of the related art which contains a large amount of an alloy element such as Cr or Mo, a thick high tensile strength steel plate having a thickness of, for example, 80 mm or more according to the present invention has been basically made essentially under specific conditions: the thick high tensile strength steel plate does not substantially contain the above-described alloy element; and the Cu content and the Ni content are appropriately controlled. Under these conditions, steel having uniform properties in a through-thickness direction and having superior weldability, base metal toughness, and HAZ toughness can be produced.

(1) That is, according to one aspect of the present invention, there is provided a steel plate comprising, as a chemical composition, by mass%, C: 0. 03% to 0.12%, Si: 0.05% to 0.30%, Mn: 1.20% to 1.65%, Cu: 0.7% to 2.5%, Ni: 1.2% to 3.0%, Nb: 0.005% to 0.030%, Ti: 0.005% to 0.030%, Al: 0.015% to 0.065%, N: 0.0020% to 0.0060%, Mo: 0% to 0.04%, Cr: 0% to 0.08%, V: 0% to 0.01%, B: 0% to 0.0005%, P: 0.010% or less, S: 0.002% or less, Ca: 0% to 0.0030%, Mg: 0% to 0.0030%, REM: 0% to 0.0030%, and a balance consisting of Fe and impurities, wherein an A value represented by the following expression (a) is 4.5% or less, a Pcm value represented by the following expression (b) is 0.25% or less, a yield strength is 460 N/mm² to 580 N/mm², a tensile strength is 550 N/mm² to 670 N/mm², a difference between a hardness at 1/8t, which is a 1/8 thickness position from a surface of the steel plate in a through-thickness direction, and a hardness at 1/2t, which is a 1/2 thickness position from the surface of the steel plate in the through-thickness direction, is 20 or less by a Vickers hardness, and when a region surrounded by grain boundaries having a crystal orientation difference of 30° or more in a crystal orientation analysis using an electron backscatter diffraction pattern analysis is defined as a grain, a circle equivalent grain size of the grain is defined as a grain size, and the grain size at which a cumulative frequency from the smallest grain size side is 70% in a frequency distribution of the grain size is defined as an average grain size, the average grain size at 1/8t is 35 µm or less. $A = Cu + Ni$
$Pcm = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5 \times B$

(wherein C, Si, Mn, Cu, Ni, Cr, Mo, V, and B represent the amounts of the respective elements, and the unit thereof is mass%)
(2) In the steel plate according to (1), the average grain size at 3/8t, which is a 3/8 thickness position from the surface of the steel plate in the through-thickness direction, may be 35 µm or less.
(3) In the steel plate according to (1), the average grain size at 1/8t may be 25 µm or less.
(4) In the steel plate according to (3), the average grain size at 3/8t, which is a 3/8 thickness position from the surface of the steel plate in the through-thickness direction, may be 25 µm or less.
(5) In the steel plate according to any one of (1) to (4), the thickness of the steel plate may be 80 mm or more.

[Effects of the Invention]

According to the present invention, it is possible to provide a thick high tensile strength steel plate having superior uniformity of base metal properties in a through-thickness direction and having superior weldability, base metal toughness, and HAZ toughness.

[Brief Description of the Drawing]

FIG. 1 is a diagram illustrating the results of measuring cross-sectional hardness distributions in a through-thickness direction of two types of steel plates having a thickness of 110 mm after being quenched and tempered, the steel plates containing 0.12% of C, 0.08% of Si, 1.45% of Mn, and 1.15% of Cu as basic components and containing 1.81% or 3.22% of Ni and the other components in ranges of an embodiment of the present invention.
FIG. 2 is a diagram illustrating effects of the Cu content and the Ni content by calculating difference (Hv: 98 N) between a hardness of a position, which is at a distance of 12.5 mm (1/8t) from a surface each of steel plates having a thickness of 100 mm after being hot-rolled, quenched, and tempered, and a hardness at 1/2t of each of the steel plates, the steel plates containing 0.13% of C, 0.12% of Si, and 1.55% of Mn as basic components and containing 0.3% to 3.6% of Cu, 0.57% to 3.5% ofNi, and the other components in the ranges of the embodiment.
FIG. 3 is diagram illustrating a relationship between the amounts of Cu and Ni and a Charpy absorbed energy (J) when an impact test at -40°C is carried after applying a weld thermal cycle, which corresponds to heat input (3.5 kJ/mm) during welding, to 1/2t of each of quenched and tempered steel plates having a thickness of 100 mm, the steel plates containing 0.13% of C, 0.12% of Si, 1.55% of Mn, 0.012% of TI, and 0.013% of Nb as basic components and containing 0.3% to 3.6% of Cu, 0.57% to 3.5% of Ni, and the other components in the ranges of the embodiment.
FIG. 4 is a diagram illustrating a relationship between a cumulative frequency (%) and grain sizes of the respective thickness positions including an outermost surface and 1/8t to 3/8t of each of quenched and tempered steel plates having a thickness of 140 mm, the steel plates containing 0.08% of C, 0.15% of Si, 1.51% of Mn, 0.008% of P, 0.0010% of S, 1.15% of Cu, 1.23% of Ni, 0.012% of Ti, 0.012% of Nb, 0.035% of Al, and 0.0039% of N as components.
FIG. 5 is a diagram illustrating a relationship (Hall-Petch relationship) between vTrs and inverses of square roots of average grain sizes of the respective thickness positions including an outermost surface (the center of a test specimen is at a distance of 6 mm from the surface) and 1/8t to 3/8t of each of quenched and tempered steel plates having a thickness of 140 mm, vTrs being obtained from a test using test specimens which are machined from the respective thickness positions in a direction perpendicular to a rolling direction, the steel plates containing 0.08% of C, 0.15% of Si, 1.51% of Mn, 0.008% of P, 0.0010% of S, 1.15% of Cu, 1.23% ofNi, 0.012% of Ti, 0.012% of Nb, 0.035% of Al, and 0.0039% of N as components.
FIG. 6 is a diagram illustrating a relationship between an average of critical CTOD values (δc) and the Mo content when each of steel plates having a thickness of 100 mm is subjected to multi-layer welding with heat input of 25 kJ/mm, a CTOD test specimen of full thickness is machined from the steel plate in a direction perpendicular to a weld line, and the CTOD test is carried out at a test temperature of -10°C three times at each notch location including a fusion line (FL) between weld metal and base metal and at a position (FL+3 mm) which is at a distance of 3 mm from FL, the steel plates containing 0.06% of C, 0.18% of Si, 1.35% of Mn, 1.05% of Cu, 1.35% of Ni, 0.013% of Ti, and 0.015% of Nb as basic components and having different Mo contents in a range of up to 0.12%.
FIG. 7 is a diagram illustrating a relationship between the Cr content and an average of critical CTOD values (δc) obtained in an CTOD test of each of steel plates having a thickness of 100 mm which is quenched at 900°C, tempered at 580°C, and is subjected to multi-layer welding with heat input of 25 kJ/mm when the CTOD test is carried at a test temperature of -10°C three times using a CTOD test specimen of full thickness in which a machined direction is a direction perpendicular to a weld line and a notch location is a position (FL+3 mm) which is at a distance of 3 mm from a fusion line (FL) between weld metal and base metal, the steel components containing 0.05% to 0.06% of C, 0.15% to 0.18% of Si, 1.30% to 1.35% of Mn, 1.05% to 1.10% of Cu, 1.30% to 1.35% of Ni, 0.012% to 0.013% of Ti, and 0.012% to 0.015% of Nb as basic components and containing 0.05% to 0.14% of Cr.
FIG. 8 is a diagram illustrating a relationship between the V content and an average of critical CTOD values (δc) obtained in an CTOD test of each of steel plates having a thickness of 100 mm which is quenched at 900°C, tempered at 580°C, and is subjected to multi-layer welding with heat input of 25 kJ/mm when the CTOD test is carried at a test temperature of -10°C three times using a CTOD test specimen of full thickness in which a machined direction is a direction perpendicular to a weld line and a notch location is a position (FL+3 mm) which is at a distance of 3 mm from a fusion line (FL) between weld metal and base metal, the steel components containing 0.05% to 0.06% of C, 0.15% to 0.18% of Si, 1.30% to 1.35% of Mn, 1.05% to 1.10% of Cu, 1.30% to 1.35% of Ni, 0.012% to 0.013% of Ti, and 0.012% to 0.015% ofNb as basic components and containing 0.005% to 0.05% of V
FIG. 9 is a diagram illustrating a relationship between a holding temperature of a preheating process for a heat treatment and an average grain size at 1/8t in each of steel plates having a thickness of 140 mm which are subjected to a rolling process, a preheating process in which the holding temperatures are 450°C and 550°C and the holding times are different, a quenching process in which the preheated steel plate is held at 920°C for 120 minutes and is water-cooled, and a tempering process in which the quenched steel plate is held at 590°C for 100 minutes and is air-cooled, the steel plates containing 0.08% of C, 0.15% of Si, 1.51% of Mn, 0.008% of P, 0.0010% of S, 1.15% of Cu, 1.23% ofNi, 0.012% of Ti, 0.012% ofNb, 0.035% of Al, and 0.0039% of N as components.

[Embodiments of the Invention]

Hereinafter, a steel plate according to an embodiment of the present invention (hereinafter, referred to as the steel plate according to the embodiment) will be described in detail.
First, the reason for limiting a chemical composition of the steel plate according to the embodiment will be described.

C: 0.03% to 0.12%

C is an element for improving the strength of base metal. In order to obtain this effect, the C content is necessarily 0.03% or more. In order to improve the strength, the lower limit of the C content may be set as 0.04%, 0.05%, 0.06%, or 0.07%. On the other hand, when the C content is more than 0.12%, material uniformity in a through-thickness direction decreases due to an increase in hardenability. In addition, when the hardness of a weld zone increases, HAZ toughness decreases. Therefore, the upper limit of the C content is set as 0.12%. In order to improve HAZ toughness, the upper limit of the C content may be set as 0.11%, 0.10%, 0.09%, or 0.08%.

Si: 0.05% to 0.30%

Si is an element which is effective for deoxidation and improves strength. In order to obtain these effects, it is necessary that the Si content is 0.05% or more. In order to improve the strength, the lower limit of the Si content may be set as 0.06%, 0.08%, 0.10%, or 0.13%. On the other hand, when the Si content is more than 0.30%, HAZ toughness decreases. Therefore, the upper limit of the Si content is set as 0.30%. In order to improve HAZ toughness, the upper limit of the Si content may be set as 0.25%, 0.22%, 0.20%, or 0.18%.

Mn: 1.20% to 1.65%,

Mn is an element which is effective for deoxidation and improves strength. In order to obtain these effects, it is necessary that the Mn content is 1.20% or more. In order to improve the strength, the lower limit of the Mn content may be set as 1.25%, 1.28%, 1.30%, 1.33%, 1.35%, or 1.37%. On the other hand, when the Mn content is more than 1.65%, material uniformity in a through-thickness direction decreases due to an increase in hardenability, and HAZ toughness decreases due to significant segregation in a slab. Therefore, the upper limit of the Mn content is set as 1.65%. In order to improve HAZ toughness, the upper limit of the Mn content may be set as 1.60%, 1.58%, 1.55%, 1.52%, 1.50%, or 1.47%.

Cu: 0.7% to 2.5%

Cu is a major alloy element of the steel plate according to the embodiment and is one of few elements for improving the strength of base metal without causing a decrease in weldability and HAZ toughness. By controlling the Cu content to be 0.7% or more, an effect of significantly increasing strength is obtained. Therefore, the lower limit of the Cu content is set as 0.7%. In order to improve the strength, the lower limit of the Cu content may be set as 0.75%, 0.8%, 0.85%, 0.9%, 0.95%, 1.0%, 1.05%, or 1.1%. On the other hand, when the Cu content is more than 2.5%, hardenability increases, which may decrease HAZ toughness as illustrated in FIG. 3. Therefore, the upper limit of the Cu content is set as 2.5%. In order to improve HAZ toughness, the upper limit of the Cu content may be set as 2.3%, 2.1%, 1.9%, 1.7%, 1.6%, 1.5%, or 1.4%.

Ni: 1.2% to 3.0%

Ni is also a major alloy element of the steel plate according to embodiment which is effective for improving base metal strength and toughness and for improving HAZ toughness. From the viewpoint of improving HAZ toughness, it is necessary that the Ni content is 1.2% or more as illustrated in FIG. 3. In order to improve the above-described properties, the lower limit of the Ni content may be set as 1.25%, 1.3%, 1.35%, 1.4%, 1.45%, 1.5%, 1.55%, or 1.6%. On the other hand, when the Ni content is more than 3.0%, there is a difference in properties in a through-thickness direction as illustrated in FIG. 2. Therefore, the upper limit of the Ni content is limited to 3.0%. In order to decrease the difference in properties in a through-thickness direction, the upper limit of the Ni content may be set as 2.8%, 2.6%, 2.4%, 2.2%, 2.0%, 1.9%, or 1.8%.

Nb: 0.005% to 0.030%,

Nb is an element which is effective for improving strength and refining the grains of base metal. In order to obtain these effects, it is necessary that the Nb content is 0.005% or more. In order to improve the strength and refine the grains, the lower limit of the Nb content may be set as 0.007%, 0.010%, 0.012%, 0.013%, or 0.015%. On the other hand, when the Nb content is more than 0.030%, HAZ toughness decreases. Therefore, the upper limit of the Nb content is set as 0.030%. In order to improve HAZ toughness, the upper limit of the Nb content may be set as 0.027%, 0.025%, 0.022%, or 0.020%.

Ti: 0.005% to 0.030%,

Ti is an element which forms a nitride and contributes to the refinement of grains in a heat affected zone. In order to obtain these effects, it is necessary that the Ti content is 0.005% or more. In order to improve HAZ toughness, the lower limit of the Ti content may be set as 0.007%, 0.010%, or 0.012%. On the other hand, when the Ti content is more than 0.030%, a nitride may be coarsened and HAZ toughness may decrease. Therefore, the upper limit of the Ti content is set as 0.030%. In order to prevent a decrease in HAZ toughness, the upper limit of the Ti content may be set as 0.025%, 0.020%, or 0.018%.

Al: 0.015% to 0.065%

Al is an element which is effective for deoxidation, forms a nitride, and is effective for refining the grains of base metal and HAZ. In order to obtain these effects, it is necessary that the Al content is 0.015% or more. In order to refine grains of base metal and HAZ, the lower limit of the Al content may be set as 0.020%, 0.025%, 0.028%, 0.031%, or 0.035%. On the other hand, when the Al content is more than 0.065%, a coarse nitride is formed, which tends to decrease toughness. Therefore, the upper limit of the Al content is set as 0.065%. In order to prevent a decrease in toughness, the upper limit of the Al content may be set as 0.060%, 0.055%, 0.052%, 0.050%, or 0.048%.

N: 0.0020% to 0.0060%

Ni is an element which binds to an element such as Ti or Al to form a nitride. From the viewpoint of forming a nitride, it is necessary that the N content is 0.0020% or more. In order to more reliably form a nitride, the lower limit of the N content may be set as 0.0024% or 0.0028%. On the other hand, when the N content is more than 0.0060%, HAZ toughness decreases. Therefore, the upper limit of the N content is set as 0.0060%. In order to prevent a decrease in HAZ toughness, the upper limit of the N content may be set as 0.055%, 0.050%, or 0.045%.

Cr: 0% to 0.08%

Mo: 0% to 0.04%

V: 0% to 0.01%

Cr, Mo, and V are elements which increase hardenability and increase a difference in hardness between a surface and a mid-thickness of a thick high tensile strength steel plate. In addition, when Cr, Mo, and V are contained, HAZ toughness may decrease. Therefore, in the steel plate according to the embodiment, it is necessary that these elements be reduced in amount.
As described above, a Charpy test is used for evaluation of HAZ toughness in most cases. However, recently, a CTOD test in which a CTOD value which can be reflected in the design in consideration of fracture mechanics has also been carried out. The CTOD value refers to the value of crack tip opening displacement and refers to, when brittle fractures are initiated from a fatigue crack tip, the amount of opening at the crack tip. The CTOD test is an experimental method of obtaining this CTOD value. Typically, the CTOD test is carried out at a design temperature which is actually set for a structure. The CTOD value is affected by the microstructures of a steel plate having a fatigue crack tip, that is, by the state of hardness, grain size, or a carbide and whether or not a brittle structure is present. It is said that the CTOD test is more sensitive to these metallurgical factors than the Charpy test. In most cases, it is determined that, when the CTOD value is 0.1 mm or more, a steel plate has sufficient resistance to brittle fractures.
The present inventors have verified the effects of the amounts of Cr, Mo, and V, which are elements having particularly high hardenability, on the CTOD value. FIG. 6 is a diagram illustrating the results of evaluating the effect the Mo content by carrying out the CTOD test on welded joints of plural steel plates having different Mo contents. In this test, the steel plates having a thickness of 100 mm were produced by melting steel, which contained 0.06% of C, 0.18% of Si, 1.35% of Mn, 1.05% of Cu, 1.25% ofNi, and 0.013% of Ti as basic components and contained free (amount of Mo contained as an impurity) to 0.12% of Mo, and hot-rolling the steel. Next, each of the steel plates was quenched at 900°C and was tempered at 580°C and then was subjected to multi-layer welding with heat input of 25 kJ/mm. A CTOD test specimen of full thickness was machined from the welded steel plate in a direction perpendicular to a weld line. Notch locations of the CTOD test are a fusion line (FL) between weld metal and base metal and a position (FL+3 mm) which is at a distance of 3 mm from FL. The CTOD test was carried out on the machined test specimens three times for each at a test temperature of -10°C.
In FIG. 6, the vertical axis represents the average of three critical CTOD values δc (also referred to as δc-10°C) at -10°C, and the horizon axis represents the Mo content. It can be seen from FIG. 6 that Mo decreases CTOD properties of the welded joints, in particular, CTOD properties at the FL+3 mm position and the FL position. In addition, it can be seen that, when δc≥0.1 mm is set as a criterion of passing, the Mo content is required to be 0.04% or less.
The lower the Mo content, the better. However, it is not preferable that the steel plate is made not to contain Mo at all because the cost increases. In addition, in consideration of a case where Mo is contained as an impurity or on purpose, the upper limit of the Mo content is set as 0.04%. More preferably, the upper limit of the Mo content is set as 0.03%, 0.02%, or 0.01%.
Likewise, effects of the Cr content and the V content on HAZ toughness have been investigated. The results are illustrated in FIGS. 7 and 8. These figures show a relationship between δc and the amounts of Cr and V, δc being obtained by preparing a welded joint with the same method as that of FIG.6, forming a notch at the FL+3 mm position, and carrying out the CTOD test at a test temperature of -10°C. As the amounts of Cr and V increase, δc becomes less than 0.1 mm at certain contents. When the upper limits of the each content where δc is 0.1 mm or more are obtained from FIGS. 7 and 8, the upper limit of the Cr content is 0.08% and the upper limit of the V content is 0.01%. Therefore, irrespective of whether Cr is contained as an impurity or on purpose, the upper limit of the Cr content is 0.08%. In order to improve HAZ toughness, the upper limit of the Cr content may be set as 0.06%, 0.05%, 0.04%, or 0.03%. In addition, the upper limit of the V content is set as 0.01% irrespective of whether V is contained as an impurity or on purpose. In order to improve HAZ toughness, the upper limit of the V content may be set as 0.008%, 0.005%, 0.003%, or 0.001%.
Cr, Mo, and V may be incorporated from scrap or the like as impurities during the preparation of molten steel, and the lower limits thereof are not particularly limited and are 0%.
As in the case of Cr, Mo, and V, B is also an element which increases hardness after quenching with a small amount thereof and is effective for improving hardenability. However, in the case of a thick high tensile strength steel plate, a difference in quenching hardness between a surface and a mid-thickness thereof increases due to the addition of B. Accordingly, from the viewpoint of the uniformity of properties in a through-thickness direction, it is not preferable that B be contained. However, it is technically difficult for a steel plate not to contain these elements at all. Accordingly, in consideration of a case where B is contained as an impurity, the upper limit of the B content is set as 0.0005%. Even if B is contained on purpose, the upper limit of the B content is 0.0005%. In order to further improve the uniformity of properties in a through-thickness direction, the upper limit of the B content may be set as 0.0004%, 0.0003%, 0.0002%, or 0.0001%. B may be incorporated from scrap or the like as an impurity during the preparation of molten steel, and the lower limit thereof is not particularly limited and is 0%.
P and S are impurity elements contained in steel and decrease base metal toughness and HAZ toughness. Therefore, the lower the amounts of P and S, the better. In the present invention, the upper limit of the P content is limited to be 0.010% or less and preferably 0.007% or less, 0.005% or less, or 0.003% or less. The upper limit of the S content is preferably limited to be 0.002% or less. The upper limit of the S content may be limited to 0.001% or 0.0008%. The lower limits of the P content and the S content are not particularly limited and are 0%.
Ca has an effect of spheroidizing a sulfide of the steel plate to reduce the effect of MnS which is harmful to toughness. In order to obtain this effect, 0.0001% or more of Ca may be contained. However, when the Ca content is excessively high, weldability decreases. Therefore, the Ca content is limited to be 0.0050% or less. In order to improve weldability, the upper limit of the Ca content may be set as 0.0040%, 0.0035%, or 0.0030%. Although Ca may be incorporated from scrap, refractory, or the like as an impurity during the preparation of molten steel, the lower limit thereof is not particularly limited and is set as 0%.
Mg and REM are elements which form an oxide in the steel plate and improve HAZ toughness. In order to obtain these effects, 0.0001% or more of Mg and REM may be contained. However, when the amounts of Mg and REM are excessively high, a coarse oxide is formed, which may decrease toughness. Therefore, the Mg content and the REM content are limited to be 0.0030% or less, for each. Optionally, the upper limits of these contents may be set as 0.0025% or 0.0020%. Mg and REM may be incorporated from scrap, refractory, or the like as impurities during the preparation of molten steel, and the lower limits thereof are not particularly limited and are 0%.
Here, REM is a general term for 17 elements including 15 lanthanide elements, Y, and Sc. Among these elements, one or more elements may be contained. The REM content refers to the total amount of these elements.
Irrespective of whether alloy elements having no lower limits (for example, Mo, Cr, V, B, P, S, Ca, Mg, and REM) are added to a steel plate on purpose or are incorporated into a steel plate as impurities, as long as the contents thereof are within the ranges described in Claims, the steel plate is included in the scope of Claims of the present invention.
The steel plate according to the embodiment contains the above-described components and a balance including Fe and impurities. However, in order to improve the strength, toughness and the like of the steel or as impurities from auxiliary materials such as scrap, the steel plate according to the embodiment may further contain Sb, As, Sn, Pb, Zr, Zn, W, and Co in addition to the above-described components. However, it is preferable that the upper limits of the amounts of the above elements be set as described below.
Since Sb decreases HAZ toughness, the upper limit of the Sb content may be set as 0.02%. In order to improve HAZ toughness, the upper limit of the Sb content may be set as 0.01%, 0.005%, or 0.002%.
Since As and Sn decrease HAZ toughness, the upper limits of the As content and the Sn content may each be set as 0.02%. Optionally, the upper limits of the As content and the Sb content may be set as 0.01%, 0.005%, or 0.002%.
In addition, in order to improve strength and toughness, the amounts of Pb, Zr, Zn, and W may be set to be 0.1% or less, 0.01% or less, 0.005% or less. These lower limits are not particularly limited and are 0%.
Co may be contained in Ni as an impurity. Since Co decreases HAZ toughness, the upper limit of the Co content may be set as 0.3%, 0.1%, or 0.05%. The lower limit is not particularly limited and is 0%.

A value (=Cu+Ni): 4.5% or less

In the embodiment, regarding a through-thickness direction of base metal, it is necessary to control ΔHv which is an index mainly indicating the uniformity of strength. As illustrated in FIG 2, when the A value is more than 4.5%, ΔHv is more than 20 and properties in a through-thickness are non-uniform, the A value being Cu+Ni, that is, the sum of the Cu content and the Ni content and being represented by the following expression (1), ΔHv being a difference between a Vickers hardness at 1/8t and a Vickers hardness at 1/2t. Based on this result, the upper limit of the A value is set as 4.5% in addition to the limitation of the ranges of the above-described each element. In order to reduce the difference in hardness in the through-thickness direction, optionally, the upper limit of the A value may be set as 4.2%, 4.0%, 3.8%, 3.5%, 3.3%, or 3.0%. The lower limit of the A value is not particularly limited but is, in practice, 1.9% which is the sum of the lower limits of the Cu content and the Ni content. $A = Cu + Ni$

(wherein Cu and Ni in the expression (1) represent the amounts of the respective elements, and the unit thereof is mass%).
Further, in order to secure weldability in the steel plate according to the embodiment, in addition to the limitation of the ranges of the individual elements, the chemical composition is limited such that a Pcm value obtained from the following expression (2) is 0.25% or less. The Pcm value is widely used as an index indicating a weld cracking parameter as in the case of carbon equivalent (Ceq) and is calculated from the amount of an alloy element contained in steel. In the expression (2), elements such as Cr, Mo, V, or B which are not substantially contained in the present invention are expressed. However, these elements may be incorporated from various alloy raw materials as impurities in the industrial process. Therefore, in order to evaluate weldability, it is necessary that the amounts of alloy elements including the above impurities are evaluated. When each alloy element is not contained (is not detected), an amount thereof only needs to be set as 0 in the calculation. $Pcm = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5 \times B$

(wherein C, Si, Mn, Cu, Ni, Cr, Mo, V, and B represent the amounts of the respective elements, and the unit thereof is mass%)
In the steel plate according to the embodiment, when the Pcm value is more than 0.25%, low temperature cracking is likely to occur during welding at 0°C. Therefore, the upper limit of the Pcm value is set as 0.25%. The lower limit of the Pcm value is not particularly limited but may be set as 0.15% or 0.18%.
Next, the steel plate can be produced using the following production method.
First, molten iron containing steel components (chemical composition) which are controlled to be within the above-described ranges are cast into a slab by a continuous casting process or an ingot casting slabbing process (casting process: S1). Next, the obtained slab is heated (heating process: S2). In order to impart a sufficient rolling reduction effect up to a mid-thickness during the rolling of the thick high tensile strength steel plate, the lower limit of a desired heating temperature in the heating process is preferably set as 950°C. On the other hand, when the heating temperature is higher than 1250°C, a scale cannot be peeled off from a steel plate, and surface defects are generated in the steel plate. Therefore, the upper limit of the heating temperature is preferably set as 1250°C.
After the heating process, the heated slab is hot-rolled into a steel plate (hot-rolling process: S3). After the hot-rolling process, the steel plate is cooled to 350°C or lower (cooling process: S4). At this time, if there is a limit to a cooling place or the like, optionally, accelerated cooling may be carried out because the steel plate is reheated to the Ac3 transformation point or higher after the cooling process. It is not preferable that a cooling stop temperature in the cooling process is higher than 350°C because embrittlement may occur due to a coarse precipitate such as aluminum nitride.
The Ac1 transformation point described herein refers to a temperature at which ferrite starts to be locally transformed into austenite when steel is heated from room temperature. In addition, when the steel is further heated, a two-phase state of ferrite and austenite is transformed into a single-phase state of austenite. The Ac3 transformation point refers to a temperature at which the two-phase state is transformed into the single-phase state of austenite. These transformation points can be experimentally obtained, typically, by using a difference in the thermal dilatation coefficient between ferrite and austenite. That is, the transformation points can be experimentally obtained from changing points of thermal dilatation in an dilatation-temperature curve which is obtained by heating the steel at a fixed heating rate (for example, 2.5 °C/min).
After the cooling process, a quenching process of heating the steel plate to the Ac3 transformation point or higher and water cooling it and a tempering process of heating the steel plate to the Ac1 transformation point or lower and air-cooling it are carried out (quenching and tempering process: S5).
When the heating temperature during quenching is lower than the Ac3 transformation point, a sufficiently hardened structure cannot be obtained. Therefore, strength or toughness may decrease. On the other hand, from the viewpoint of preventing the coarsening of grains, the heating temperature during quenching is preferably low. Therefore, the upper limit of the heating temperature may be set as 930°C, 910°C, or 890°C. In addition, when the heating temperature during tempering is higher than the Ac1 transformation point, strength or toughness may significantly decrease.
In a recent example, unlike in the method according to the embodiment in which the steel plate cooled after rolling is reheated to be quenched and tempered, a method in which the steel plate is directly cooled after rolling and the cooled steel plate is tempered (direct quencing+tempering) is applied to the production of a high tensile strength steel plate. However, this method is not suitable for the steel plate according to the embodiment. The reason is as follows.
The grain size of a steel plate which is directly quenched after rolling depends on heating and rolling temperatures. When low-temperature heating is performed or low-temperature rolling is performed to refine grains, a rolling temperature decreases on the surface of the steel plate which is likely to be cooled. As a result, in many cases, immediately after rolling, the surface of the steel plate has an austenite structure having flat and fine grains due to hot-rolling, and the center thereof has an austenite structure having isotropic and slightly coarse grains which are produced by recrystallization and are not easily affected by rolling. When the steel plate having these austenite structures is directly quenched, a region from the surface to 1/8t thereof which is affected by rolling has microstructures in which ferrite and bainite structures having fine grains which are transformed from the processed austenite are mainly present. Conversely, an inside region from 2/8t to the center of the steel plate has microstructures in which ferrite and bainite structures having coarse grains are present. As a result, the average grain size at 3/8t of the steel plate is 35 µm or more.
That is, in the steel plate which is directly quenched as described above (directly quenched steel), the grain size of the region from the surface to 1/8t is less than that of the mid-thickness. This microstructure is completely opposite to that of the steel plate according to the embodiment. That is, in the directly quenched steel, even if the grain size of a base metal at 1/8t thereof is defined, the grain size in an inside region from 1/8t to the center is more than that at 1/8t. Therefore, the base metal toughness cannot be specified by the limitations of the ranges according to the present invention. Further, since the grain size on the surface is refined, the surface side tends to be harder in a hardness distribution in a through-thickness direction and thus cannot satisfy ΔHv≤20.
As described above, in the thick high tensile strength steel plate, the direct quenching+tempering method is not suitable as a means for securing material uniformity in a through-thickness direction and imparting superior toughness. In order to secure material uniformity in a through-thickness direction and control the average grain size at 1/8t to be 35 µm or less, it is necessary to perform a quenching process and a tempering process after cooling.
Further, in the embodiment, in order to obtain uniform grains in a through-thickness direction during quenching, it is preferable that, as an intermediate process between the hot-rolling process and the quenching and tempering process, a process (preheating process: S6) of preheating the steel plate be further provided such that the temperature of the steel plate is in a range of 550°C to the Ac1 transformation point and is held in this temperature range for 5 hours to 500 hours. By performing this preheating process, a difference in grain size in a through-thickness direction illustrated in FIG. 4 can be reduced. That is, this preheating process is the process which is performed before quenching in order to prevent the coarsening of grains caused when the heating time in the region from the surface to 1/8t is long during the heating process during the quenching of the thick high tensile strength steel plate described above. The metallurgically meaningful of this preheating process is to coarsen carbonitrides of Ti and Nb or aluminum nitride precipitates, which finely precipitate after hot-rolling, into an appropriate size by Ostwald growth so as to function as pinning particles during quenching. FIG. 9 is a diagram illustrating a change in the average grain size at 1/8t of each of steel plates which are obtained by rolling steel into a size of 140 mm, preheating the rolled steel at temperatures of 450°C and 550°C for different holding times, holding the preheated steel at 920°C for 120 minutes and water cooling it during a quenching process, and holding the quenched steel at 590°C for 100 minutes and air-cooling it during a tempering process, the steel containing 0.08% of C, 0.15% of Si, 1.51% of Mn, 0.008% of P, 0.0010% of S, 1.15% of Cu, 1.23% of Ni, 0.012% of Ti, 0.012% of Nb, 0.035% of Al, and 0.0039% of N as components.
As can be seen from FIG. 9, when the temperature of the preheating process is 450°C, the average grain size tends to gradually decrease along with an increase in the holding time. However, in order to control the average grain size to be 25 µm or less, an extremely long holding time of 100 hours or longer is required. On the other hand, when the temperature of the preheating process is 550°C, the average grain size is controlled to be 25 µm or less with a holding time of 5 hours or longer, and clear grain refinement is shown. It can be seen from the above-described results that, in order to decrease the average grain size in the vicinity of 1/8t of the steel plate where grains are likely to be coarsened, the preheating process is performed, preferably, at 550°C or higher for 5 hours or longer. By decreasing the average grain size, toughness is further improved. Further, the grain refinement effect by the above-described preheating process is greater on the grain size of the surface than those of the mid-thickness. Therefore, a difference in toughness between the surface and the mid-thickness decreases, and toughness in the through-thickness direction tends to be uniform. However, when the holding time in the preheating process is longer than 500 hours, the coarsening of precipitate particles significantly advances during the preheating process, and the number density of the particles decreases, and thus a pinning effect decreases. Therefore, the upper limit of the holding time is preferably set as 500 hours. When the preheating temperature is higher than the Ac1 transformation point, austenite transformation partially occurs in the steel plate. In this case, due to a difference in the growth speed of precipitates between ferrite and austenite, the uniform growth of precipitates in the steel plate cannot be expected. Therefore, the heating temperature (holding temperature) of the preheating process is preferably controlled to be the Ac1 transformation point or lower.
After the preheating process, the steel plate is cooled to 350°C or lower and then is quenched. The quenching process is the process in which the steel plate heated to a temperature higher than the Ac3 transformation point is water-cooled. From the viewpoint of preventing the coarsening of grains, the heating temperature during quenching is preferably low. Therefore, the upper limit of the heating temperature may be set as 930°C, 910°C, or 890°C.
After the quenching process, a tempering process is performed. The tempering process is a process which is important to control strength and toughness to be within the predetermined ranges. In the embodiment, the tempering process is performed at the Ac1 transformation point or lower in order to secure material uniformity in a through-thickness direction. The temperature range is preferably a range of 500°C to 650°C and more preferably 550°C to 610°C. In order to control the difference ΔHv in Vickers hardness between 1/8t and 1/2t from the surface of the steel plate in the hardness distribution in the through-thickness direction to be 20 or less, it is effective to perform the tempering process at the above-described temperature.

[Examples]

Using slabs obtained by melting steel A1 to A10 and B1 to B29 having chemical compositions shown in Tables 1 and 2, steel plates having a thickness of 80 mm to 200 mm were produced under production conditions shown in Tables 3 and 4.
During production, the heating temperature was 950°C to 1250°C. Subsequently, the steel plates were hot-rolled and air-cooled or water-cooled. Next, the steel plates of Test Nos. 5, 10, 15, and 26 were preheated before being quenched. The steel plates of Test Nos. 1 to 51 were quenched and tempered except for the steel plate of Test No. 18. The steel plate of Test No. 18 was water-cooled to 100°C immediately after rolling and was simply tempered without being quenched. Next, in order to evaluate strength properties of base metal, a tensile test specimen defined in JIS Z 2201 (No. 14) was machined, and a tensile test defined in JIS Z 2241 was performed. As a result of the test, a case where the yield strength was 460 N/mm² to 580 N/mm² and the tensile strength was 550 N/mm² to 670 N/mm² was determined as "Pass". Further, an impact test specimen was machined according to JIS Z 2242, and a test was performed. In the impact test which was performed to evaluate base metal toughness, the average value of three absorbed energy values at -40°C is shown as vE-40 (base metal), and a case where the average value was 42 J or higher was determined as "Pass". The tensile test specimen was machined from 1/4t of each of the steel plates which was generally defined in ordinary steel standards. The impact test specimen was machined from three positions including 1/8t, 1/4t, and 1/2t of each of the steel plates. In Tables 3 and 4, only the test result of 1/2t (mid-thickness) where toughness was lowest is shown. The machined direction of all the test specimens was a direction perpendicular to a rolling direction. The Ac1 and Ac3 transformation points were read from a change in thermal dilatation in a longitudinal direction of a cylindrical test specimen having a diameter of 3 mmϕ and a length of 10 mm when the test specimen was machined from 1/4t of each of the steel plates by machining, a thermocouple was mounted on a tip of the test specimen, and the test specimen was heated from room temperature to 950°C at a heating rate of 2.5°C/min by high frequency induction heating.
In addition, microstructure test specimens were machined from 1/8t and 3/8t of each of the steel plates in a direction perpendicular to a rolling direction and were mirror-polished. In these microstructure test specimens, a region surrounded by grain boundaries having a crystal orientation of 30° or more obtained by EBSD method was defined as a grain, and a circle equivalent diameter of the grain was defined as a grain size. A frequency distribution of the grain size of each sample was measured, and the grain size at which a cumulative frequency from the smallest grain size was 70% was defined as an average grain size.
Further, a Vickers hardness distribution (load: 98 N) of a cross-section of each of the steel plates in a through-thickness direction was measured, and a difference in hardness between 1/8t and 1/2t of the steel plate is shown as ΔHv which is an index of material uniformity. In addition, a case where ΔHv was 20 or less was determined as "Pass". Here, two 1/8t locations were present in the steel plate (that is, positions including 1/8t and 7/8t from one surface), and ΔHv refers to a larger value among the differences between the two 1/8t locations and 1/2t.
In order to evaluate weldability, a y-groove weld cracking test defined in JIS Z 3158 was performed. CO₂ welding was performed with heat input of 1.5 kJ/mm. As a steel plate provided for the test, a steel plate whose front and back surfaces were cut such that the thickness was 50 mm centering on a mid-thickness was used. As a result of the test, a preheat temperature at which a root crack ratio was 0% was obtained, and a case where the test temperature was 0°C was determined as "Pass".
On the other hand, in order to evaluate HAZ toughness, a K-groove butt joint was prepared by submerged arc welding with heat input of 3.5 kJ/mm to 4.5 kJ/mm. Three impact test specimens defined in JIS Z 3128 were machined from the butt joint with a fusion line as a notch location, and an impact test was performed at a test temperature of -40°C. The average values of the three test specimens are shown in Table 3 and 4 as vE-40 (HAZ).
In addition, a full-thickness CTOD test specimen (B×B type) defined in BS 7448 was machined from the above butt joint with a fusion line called CGHAZ (Coarse grain HAZ) as a notch location, and CTOD tests defined in API (American Petroleum Institute) RP 2Z and BS (British Standards) 7448 were performed three times at a test temperature of -10°C, for each. Minimum values of the test results are shown in Tables 3 and 4 as δc-10°C. In the impact test, a case where the value was 42 J or higher was determined as "Pass". In the CTOD test (δc), a case where the value was 0.1 mm or more was determined as "Pass".
It is said that there is a rough correlation between the result of the impact test and the result of the CTOD test; however, even if one of the results is satisfactory, the other one may be unsatisfactory. Therefore, in a structure where requirements for fractures are strict, it is necessary to satisfy the results of both the tests as HAZ toughness.

In Tables 1 and 2, the steel components, the A values (Cu+Ni), and the Pcm values which are underlined indicate that the values thereof are out of the range according to the present invention. In Tables 3 and 4, the underlined numerical values indicate that the properties thereof are insufficient. In addition, in Tables 1 and 2, the balance includes Fe and impurities.

[Table 2]

(mass%)
STEEL	C	Si	Mn	P	S	Cu	Ni	Nb	Ti	Al	N	Cr	Mo	V	B	others	A VALUE (%)	Ac1 (°C)	Ac3 (°C)	Pcm (%)
B1	0.02	0.07	1.21	0.005	0.002	0.91	1.20	0.021	0.017	0.033	0.0035	0.01	0.01	0.002	0.0002	-	2.11	694	861	0.15
B2	0.14	0.07	1.21	0.005	0.002	0.91	1.20	0.021	0.017	0.033	0.0035	0.01	0.01	0.002	0.0002	-	2.11	691	804	0.27
B3	0.06	0.01	1.20	0.005	0.002	1.02	1.52	0.023	0.015	0.051	0.0042	0.02	0.03	0.002	0.0001	-	2.54	684	829	0.20
B4	0.06	0.37	1.22	0.004	0.002	1.04	1.55	0.018	0.012	0.053	0.0042	0.02	0.03	0.002	0.0001	-	2.54	687	848	0.21
B5	0.06	0.07	0.58	0.005	0.002	0.96	1.75	0.019	0.012	0.042	0.0048	0.02	0.01	0.001	0.0002	-	2.71	687	839	0.17
B6	0.05	0.06	1.89	0.003	0.002	0.96	1.77	0.017	0.011	0.047	0.0035	0.01	0.005	0.001	0.0002	-	2.71	670	817	0.24
B7	0.07	0.06	1.58	0.012	0.002	0.99	1.42	0.021	0.015	0.048	0.0032	0.03	0.02	0.005	0.0003	-	2.41	682	823	0.23
B8	0.08	0.18	1.35	0.003	0.004	1.13	1.65	0.022	0.019	0.036	0.0038	0.02	0.02	0.002	0.0001	-	2.78	678	821	0.24
B9	0.08	0.05	1.21	0.003	0.001	0.65	2.12	0.021	0.014	0.025	0.0029	0.02	0.01	0.001	0.0002	Ca:0.0011	2.77	677	811	0.21
B10	0.05	0.06	1.32	0.003	0.001	2.54	1.41	0.015	0.013	0.046	0.0051	0.02	0.01	0.002	0.0001	-	3.95	649	812	0.27
B11	0.08	0.08	1.64	0.003	0.001	1.08	0.92	0.021	0.015	0.047	0.0055	0.01	0.02	0.002	0.0002	-	2.00	691	830	0.24
B12	0.07	0.06	1.22	0.005	0.002	1.15	3.48	0.015	0.012	0.038	0.0042	0.01	0.01	0.001	0.0001	-	4.63	633	772	0.25
B13	0.04	0.08	1.35	0.004	0.002	2.62	1.05	0.018	0.014	0.045	0.0035	0.01	0.02	. 0.002	0.0002	-	3.67	656	826	0.26
B14	0.06	0.12	1.21	0.005	0.002	2.85	3.15	0.024	0.012	0.033	0.0042	0.01	0.02	0.001	0.0001	-	6.00	606	762	0.32
B15	0.05	0.07	1.32	0.004	0.002	0.61	3.29	0.012	0.011	0.031	0.0039	0.02	0.01	0.002	0.0002	-	3.90	649	794	0.21
B16	0.07	0.06	1.33	0.005	0.002	0.57	1.11	0.014	0.015	0.032	0.0048	0.02	0.03	0.002	0.0002	-	1.68	700	843	0.19
B17	0.04	0.06	1.54	0.002	0.002	0.93	1.35	0.004	0.014	0.034	0.0039	0.04	0.01	0.004	0.0003	Mg:0.0021	2.28	683	841	0.19
B18	0.08	0.18	1.33	0.004	0.002	1.22	1.41	0.038	0.015	0.049	0.0031	0.02	0.02	0.001	0.0001	-	2.63	686	826	0.24
B19	0.07	0.05	1.58	0.002	0.001	1.14	1.54	0.014	0.003	0.040	0.0033	0.02	0.01	0.002	0.0001	-	2.68	674	816	0.24
B20	0.08	0.06	1.35	0.002	0.002	1.03	1.47	0.019	0.036	0.044	0.0041	0.03	0.02	0.001	0.0002	-	2.50	682	820	0.23
B21	0.07	0.08	1.32	0.002	0.001	1.05	1.54	0.022	0.012	0.014	0.0047	0.01	0.02	0.001	0.0001	Ca:0.0008	2.50	681	825	0.22
B22	0.08	0.08	1.21	0.002	0.001	0.98	1.52	0.023	0.011	0.077	0.0055	0.02	0.02	0.002	0.0002	-	2.50	685	824	0.22
B23	0.07	0.07	1.54	0.005	0.002	1.19	1.51	0.019	0.015	0.046	0.0075	0.02	0.03	0.005	0.0001	-	2.70	676	819	0.24
B24	0.08	0.06	1.64	0.005	0.002	0.95	1.44	0.021	0.012	0.042	0.0044	0.11	0.01	0.002	0.0002	-	2.39	684	816	0.24
B25	0.05	0.18	1.61	0.004	0.001	1.02	2.27	0.022	0.015	0.048	0.0041	0.02	0.05	0.002	0.0001	-	3.29	665	816	0.24
B26	0.04	0.05	1.35	0.002	0.001	0.99	1.88	0.021	0.019	0.036	0.0038	0.02	0.01	0.012	0.0002	-	2.87	674	830	0.19
B27	0.10	0.06	1.33	0.002	0.002	0.91	2.31	0.015	0.014	0.035	0.0036	0.02	0.03	0.002	0.0006	-	3.22	664	791	0.26
B28	0.05	0.07	1.32	0.003	0.001	1.68	2.92	0.018	0.012	0.049	0.0048	0.01	0.02	0.002	0.0001	Ca:0.0009	4.60	635	787	0.25
B29	0.07	0.15	1.47	0.004	0.002	1.44	2.48	0.021	0.016	0.059	0.0044	0.02	0.01	0.002	0.0003	REM:0.0012	3.92	650	794	0.27

In Test Nos. 1 to 17 of Table 3, the steel components and the production conditions were all in the ranges according to the present invention. Each steel satisfied the desired values regarding the tensile properties and toughness (impact properties) of base metal and ΔHv which is an index of uniformity in a through-thickness direction. Further, regarding weldability, cracking was not found in all of the steel at 0°C, and HAZ toughness, the absorbed energy (vE-40) and the CTOD value (δc-10°C) satisfied the desired values.
Among these, in Test Nos. 5, 10, and 15 in which the preheating process was performed in the range according to the present invention, the average grain sizes at 1/8t and 3/8t of each of the steel plates were 25 µm or less as compared to those of the other steel plates. In addition, as a result, the base metal toughness of Test Nos. 5, 10, and 15 were higher than those of the other steel.
On the other hand, in Test Nos 18 to 22 in Table 4, the components thereof were within the range according to the present invention, but the production conditions were not preferable. Therefore, base metal properties and/or uniformity in a through-thickness direction did not satisfy the desired values. In addition, Test Nos. 23 to 51 were steel sheets produced using respective steel having chemical compositions which deviate from the ranges according to the present invention. As shown in Table 4, Test Nos. 23 to 51 did not satisfy the desired values regarding at least one of the strength and toughness of base metal, ΔHv, the critical preheat temperature, vE-40 (HAZ), and δc-10°C.
Test No. 18 is the steel plate which was water-cooled (directly quenched) immediately after rolling and was simply tempered without being quenched. In this steel sheet, the base metal toughness was low at 29 J, and ΔHv was high at 29. In the example of Test No. 19, the quenching temperature was that in the two-phase region. As a result, the tensile properties of base metal did not satisfy the desired values. In the example of Test No. 20, the tempering temperature was 705°C higher than the Ac1 transformation point. As a result, the yield strength is low and ΔHv does not satisfy the desired value. In the example of Test No. 21, the cooling stop temperature after rolling was high at 395°C, and heating for quenching was performed at this high heat. In this example, since the cooling stop temperature was high, the coarsening of precipitates was caused in the heating step of the quenching process which was the next step, and base metal toughness decreased.
Further, in the example of Test No. 22, the quenching temperature was 950°C which deviated from the preferable range. In Test No. 22, the grain size was large, and base metal toughness was low.
In the examples of Test Nos. 23, 25, and 27, the amounts of C, Si, and Mn were low and deviated from the ranges according to the present invention. In these examples, the tensile strength did not satisfy the desired values. In Test Nos. 23 and 25, the yield strength was also low.
Conversely, in the example of Test No. 24, the C content was 0.14% which was higher than and deviated from the range according to the present invention, and the Pcm value is 0.27% and from the range according to the present invention. As a result, base metal toughness was low, ΔHv was 32, uniformity in a through-thickness direction was low, the critical preheat temperature was high at 25°C, and the absorbed energy δc of a welded zone was low. Similarly, Test No. 28 contained 1.89 % of Mn, Test No. 46 contained 0.11% of Cr, and Test No. 49 contained 0.0006% of B. All the contents were higher than and deviated from the ranges according to the present invention. These elements are elements for improving the hardenability of base metal. Therefore, in Test Nos. 28,46, and 49, all the values of ΔHv were more than 20, and the yield strength and base metal toughness did not satisfy the ranges according to the present invention in some cases.
On the other hand, Test No. 26 contained 0.37% of Si, Test No. 29 contained 0.012% of P, Test No. 30 contained 0.004% of S, Test No. 40 contained 0.038% ofNb, Test No. 42 contained 0.036% of Ti, Test No. 44 contained 0.077% ofAl, Test No. 45 contained 0.0075% of N, Test No. 47 contained 0.05% of Mo, and Test No. 48 contained 0.012% of V All the contents were higher than and deviated from the ranges according to the present invention. When the amounts of these elements are higher than the ranges according to the present invention, HAZ toughness decreases. Accordingly, either or both values of vE-40 (HAZ) and δc-10°C do not satisfy the desired values.
Next, the effects of Cu and Ni which are major elements in the steel according to the present invention will be described. In Test Nos. 31,37, and 38, the Cu content was lower than and deviated from the ranges according to the present invention. Therefore, in Test No. 31, the tensile strength was low, and in Test Nos. 37 and 38, the tensile strength and the yield strength were low. Further, in Test No. 37, the Ni content was 3.29% which was higher than and deviated from the range according to the present invention. As a result, ΔHv was higher at 55. In Test No. 38, conversely, the Ni content was 1.11 which was lower than and deviated from the range according to the present invention. As a result, the toughness of a weld zone was low.
Further, in the examples of Test Nos. 32, 35, and 36, the Cu content was higher than and deviated from the range according to the present invention, and the Pcm value was higher than 0.25%. As a result, in all the examples, the critical preheat temperature was 25°C and did not satisfy the desired value, and HAZ toughness was also low. Among these examples, in Test No. 35, the Ni content was 1.05% which was lower than and deviated from the range according to the present invention, and vE-40 (HAZ) and δc-10°C were low. In addition, in the example of Test No. 36, conversely, the Ni content was higher than and deviated from the range according to the present invention, and Cu+Ni was 6.00% and deviated above 4.5% which was the upper limit of the range according to the present invention. As a result, ΔHv was 59 and did not satisfy the desired value.
In the examples of Test Nos. 33 and 34, the Cu content was within the range according to the present invention, and the Ni content deviated from the range according to the present invention. That is, in Test No. 33, the Ni content was 0.92% which was lower than and deviated from the range according to the present invention. As a result, the toughness of base metal and a welded zone did not satisfy the desired values. On the other hand, in the example of Test No. 34, the Ni content was 3.15% which was higher than and deviated from the range according to the present invention, and Cu+Ni was 4.63 and deviated above 4.5% which was the upper limit of the range according to the present invention. As a result, ΔHv was high at 45.
In Test No. 39, the Nb content was lower than and deviated from the range according to the present invention. As a result, the yield strength and the tensile strength of base metal were low. In Test No. 41, the Ti content was 0.003 which was lower than and deviated from the range according to the present invention, and vE-40 (HAZ) was low. In Test No. 43, the Al content was 0.014% which was lower than and deviated from the range according to the present invention, the refinement of grains of base metal was insufficient, and base metal toughness was low. In the examples of Test Nos. 50 and 51, the individual element contents were within the ranges according to the present invention, but the A value and the Pcm value deviated from the ranges according to the present invention. In the example of Test No. 50, the A value was 4.60% and deviated above 4.5% which was the upper limit of the range according to the present invention. In this case, ΔHv was 31 and did not satisfy the range according to the present invention. In Test No. 51, the Pcm value was 0.27% and deviated from the range according to the present invention. As a result, the critical preheat temperature was high at 25°C and did not satisfy the desired value.

[Industrial Applicability]

According to the present invention, it is possible to provide a thick high tensile strength steel plate having superior uniformity of base metal properties in a through-thickness direction and having superior base metal toughness, weldability, and HAZ toughness.

Claims

A steel plate comprising, as a chemical composition, by mass%,
C: 0.03% to 0.12%,

Si: 0.05% to 0.30%,

Mn: 1.20% to 1.65%,

Cu: 0.7% to 2.5%,

Ni: 1.2% to 3.0%,

Nb: 0.005% to 0.030%,

Ti: 0.005% to 0.030%,

Al: 0.015% to 0.065%,

N: 0.0020% to 0.0060%,

Mo: 0% to 0.04%,

Cr: 0% to 0.08%,

V: 0% to 0.01 %,

B: 0% to 0.0005%,

P: 0.010% or less,

S: 0.002% or less,

Ca: 0% to 0.0030%,

Mg: 0% to 0.0030%,

REM: 0% to 0.0030%, and

a balance consisting of Fe and impurities,
wherein an A value represented by the following expression (1) is 4.5% or less, a Pcm value represented by the following expression (2) is 0.25% or less,
a yield strength is 460 N/mm² to 580 N/mm²,
a tensile strength is 550 N/mm² to 670 N/mm²,
a difference between a hardness at 1/8t, which is a 1/8 thickness position from a surface of the steel plate in a through-thickness direction, and a hardness at 1/2t, which is a 1/2 thickness position from the surface of the steel plate in the through-thickness direction, is 20 or less by a Vickers hardness, and
when a region surrounded by grain boundaries having a crystal orientation difference of 30° or more in a crystal orientation analysis using an electron backscatter diffraction pattern analysis is defined as a grain, a circle equivalent grain size of the grain is defined as a grain size, and the grain size at which a cumulative frequency from the smallest grain size side is 70% in a frequency distribution of the grain size is defined as an average grain size, the average grain size at 1/8t is 35 µm or less. $A = Cu + Ni$
$Pcm = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5 \times B$

(wherein C, Si, Mn, Cu, Ni, Cr, Mo, V, and B represent the amounts of the respective elements, and the unit thereof is mass%)
The steel plate according to Claim 1,
wherein the average grain size at 3/8t, which is a 3/8 thickness position from the surface of the steel plate in the through-thickness direction, is 35 µm or less.
The steel plate according to Claim 1,
wherein the average grain size at 1/8t is 25 µm or less.
The steel plate according to Claim 3,
wherein the average grain size at 3/8t, which is a 3/8 thickness position from the surface of the steel plate in the through-thickness direction, is 25 µm or less.
The steel plate according to any one of Claims 1 to 4,
wherein the thickness of the steel plate is 80 mm or more.