THICK HIGH-STRENGTH STEEL PLATE AND PROCESS FOR PRODUCING THE SAME

Abstract

 A high-strength thick steel plate having a composition of Ni-containing steel and bainite-dominated structure, wherein a pearlite fraction is not greater than 5%, fraction of coarse ferrite having a circle equivalent diameter of larger than 25 µm is not greater than 10% in front and back surface portions each having a depth of 5% of the plate thickness, mean circle equivalent diameter of cementite is not greater than 0.5µm, and mean circle equivalent diameter of iso-crack propagation resisting domains is not greater than d (µm)=(7.1 ×[Ni]+11) ×(1.2-t/300) (µm), where each of the iso-crack propagation resisting domains is defined such that inner portion excluding the surface portions in a section perpendicular to the rolling direction of the steel plate is partitioned to iso-orientation domains, the measuring line of cutting method is drawn along T direction in parallel to the plate thickness, the plural iso-orientation domains continuously aligned and adjacent to each other are identified excluding iso-orientation domains having a circle equivalent diameter of smaller than 8µm, the angle between <001> axes close to the T direction of each pair of the identified iso-orientation domains is measured, and the continuously aligned iso-orientation domains satisfying the angle of not greater than 20° are regarded to constitute one iso-crack propagation resisting domain.

Description

TECHNICAL FIELD

[0001] The present invention relates to a high-strength thick steel plate (hereafter also referred to as a high-strength high-arrest thick steel plate or a high-arrest steel plate) excellent ability to arrest propagation of brittle cracks (hereafter referred to as arrestability) and method of producing the same.
Specifically, the present invention relates to a high-strength thick steel plate having excellent brittle-crack propagation arrestability and a method of producing same, where the steel plate thickness is 50 mm or more (hereafter referred to as a heavy thick plate), and the temperature at which Kca=6000N/mm^1.5 is satisfied (hereafter referred to as an arrestability indicator) is not higher than -10°C.
A steel plate according to the present invention may be applied in weld constructions such as shipbuilding, buildings, bridges, tanks, and offshore structures. The steel plate according to the present invention may also be worked to steel pipes, steel columns or the like, and may be distributed in the form of a secondary product.
Priority is claimed on Japanese Patent Application, No. 2007-54279 filed on March 5, 2007, the content of which is incorporated herein by reference.

BACKGROUND ART

[0002] In accordance with the recent trend of increasing the size of steel constructions, use of thick and high-strength steel plates is required. In addition, in terms of ensuring safety, the steel plate is strictly required to have brittle crack propagation arrestability (arrestability). On the other hand, increasing the strength or thickness of a steel generally results in steeply increasing the difficulty in ensuring the arrestability, thereby disturbing the application of the thick high-strength steel in the steel construction. In addition, there is an increasing demand of consumers to shorten the delivery time. Therefore, enhancement of productivity in the steel production process is strongly demanded.

[0003] Known metallurgical factors for enhancing arrestability of steel include (i) refining crystal grain size, (ii) Ni addition, (iii) control of brittle secondary phase, (iv) control of texture, or the like.
For example, Patent Reference 1 (Japanese Unexamined Patent Application, First Publication, No. H02-1293189) describes a method (i) of refining the crystal grain size. In this method, steel is subjected to rolling with a reduction of 50% or more at a non-recrystallization region at a temperature of not lower than Ar₃ temperature, and is subsequently subjected to rolling with a reduction of 30 to 50% at two-phase region at 700 to 750°C. As a specific method for refining crystal grain size of a steel plate, Patent Reference 2 (Japanese Examined Patent Application, Second Publication No.H06-004903) and Patent Reference 3 (Japanese Unexamined Patent Application, First Publication No. 2003-221619) each describe a method including cooling a surface of a slab before rolling or after preliminary rolling, and making the slab to recuperate by starting rolling while maintaining the temperature difference between the inner portion and the surface of the slab, thereby generating fine ferrites in the surface of the slab.

(ii) It is reported that the addition of Ni suppresses the propagation of brittle cracks by promoting cross slips in a low temperature region (Non-Patent Reference 1: Imao Tamura, "Study on steel strength" published by Nikkan Kogyo Shinbunsha, July 5, 1969, p125), and thereby enhancing the arrestability of a matrix (Non Patent Reference 2: Hasebe and Kawaguchi, "Brittle fracture propagation arrestability of Ni-added steel plate examined by taper-shaped DCB test", Iron and Steel, Vol. 61, 1975, p 875).

(iii) Patent Reference 4 (Japanese Patent Application, first Publication, No. S59-047323) describes a method of controlling brittle (embrittling) secondary phase. In this method, martensite as a brittle phase is controlled to occur as fine dispersed phase dispersed in a ferrite matrix.

(iv) With respect to the control of texture, Patent Reference 5 (Japanese Patent Application, First Publication No. 2002-241891) describes a method including rolling of low-carbon bainite steel at low-temperature and heavy-reduction conditions, thereby accumulating (211) planes in parallel to the roll plane.

DISCLOSURE OF INVENTION

Problems to be solved by the invention

[0004] However, the method described in Patent Reference 1 is directed to steel used in low-temperature conditions, where the steel has relatively low strength because of its microstructure mainly composed of ferrite, and the plate thickness is up to about 20 mm. Therefore, if the method is applied to a thick plate with a thickness of 50 mm or more, as designated by the present invention, it is difficult to ensure a sufficient reduction (reduction ratio) in terms of slab thickness. In addition, there is another problem in that prolonged stand-by time before reaching the process temperature reduces the productivity of the steel.
In addition, it is difficult to ensure a yield strength of 390MPa or more by the method described in this reference.

[0005] Application of the method described in Patent References 2 and 3 to a production of a thick member having a plate thickness of 50 mm or more, as designated in the present invention, also causes problems. Even though a similar microstructure is obtained, it is difficult to ensure arrestability. The effect of refined ferrite in the surface portion is relatively reduced. Further, the production process itself includes a problem in that thermal control along the plate thickness is further made difficult, and it is impossible to avoid increasing the reduction during the recuperation process, thereby largely disturbing the productivity.

[0006] On the other hand, too much cost is required to form an alloy
to provide a desired arrestability to a steel plate only by the addition of Ni as described in the above (ii). It is considered that refinement of the microstructure could be combined with the Ni addition to ensure arrestability while reducing the amount of Ni addition. However, there is no examination to separate and quantify the influence of factors other than the Ni addition on the arrestability. In the present circumstance, there still is no clear guideline for producing a Ni-added type high-arrest steel plate.

[0007] Further, in a thick steel plate, it is difficult to disperse fine martensite in the steel as described in Patent Reference 4. In a thick high-strength steel plate, such a brittle phase possibly deteriorate the brittle-fracture initiation property.

[0008] If the invention described in Patent Reference 5 is applied to a thick plate, efficiency of rolling is extremely reduced. Thus, such a method is not applicable in industrial production.

[0009] As described above, there has been no established technology for stable and effective production of a thick high-arrest steel plate, as directed by the present invention, having a plate thickness of 50 mm or more, arrestability indicator T_Kca=6000 of - 10°C or less even when the yield strength is in a range of 390 to 460 MPa, and being applicable to large constructions.

[0010] Based on the above-described circumstance, an object of the present invention is to provide a thick high-strength steel plate having excellent brittle crack propagation arrestability that is sufficient as a steel for large construction, and to provide a production method that enables industrially stable and effective production of the same steel plate.

[Means for solving the problems]

[0011] The present invention concerns a thick high-strength steel plate having excellent brittle crack propagation arrestability and a method of producing the same that solve the above-described problems, and includes below-described aspects.

[1] A thick high-strength steel plate having excellent brittle crack propagation arrestability, having:

a composition containing in mass %, C: 0.01 to 0.14%, Si: 0.03 to 0.5%, Mn: 0.3 to 2.0%, P: 0.020% or less, S: 0.010% or less, Ni: 0.5 to 4.0%, Nb: 0.005 to 0.050%, Ti: 0.005 to 0.050%, Al: 0.002 to 0.10%, N: 0.0010 to 0.0080%, and the balance consisting of Fe and unavoidable impurities, where Ceq defined by the below described formula (1) is 0.30 to 0.50%;

a microstructure dominated by bainite of 60% or more in volume fraction, wherein

a pearlite fraction is not greater than 5%,

mean circle equivalent diameter of cementite is not greater than 0.5µm,

fraction of coarse ferrite having a circle equivalent diameter larger than 25 µm is not greater than 10% in surface portions each having a depth of 5% of the plate thickness from the front and back surfaces of the plate,

and

mean circle equivalent diameter of iso-crack propagation resisting domains is not smaller than 8 µm and not greater than d (µm) defined by the below described formula (2),

where each of the iso-crack propagation resisting domains is defined such that

a section (section plane) perpendicular to a rolling direction of the plate is defined as a T section, and a direction in parallel to the plate surface is defined as T direction in the T section, inner portion in the T section is defined as a portion excluding the surface portions,

the microstructure of the T section is partitioned to domains each having the same crystal orientation (iso-orientation domain) by orientation analysis of the inner portion using Electron Back Scattering Pattern (EBSP),

an arbitrary measurement line is drawn on the T section microstructure partitioned to the iso-orientation domains by applying a dissection method in accordance with JIS G0551,

a plurality of the iso-orientation domains having a circle equivalent diameter of 8µm continuously aligned and adjacent to each other on the measurement line are identified excluding iso-orientation domains having a circle equivalent diameter of smaller than 8µm,

<001> axis closest to the T direction is selected from three <001> axes of each of the identified iso-orientation domains,

an angle (crack-propagation deviation angle) formed by the <001> axes closest to the T direction of the identified iso-orientation domains is measured in each adjacent pair of the iso-orientation domains on the measurement line,

the continuously aligned iso-orientation domains satisfying the angle of not greater than 20° and adjacent iso-orientation domains having a circle equivalent diameter of less than 8µm are regarded to constitute one iso-crack propagation resisting domain.




where [X] denotes content (in mass %) of an element X and t denotes a plate thickness (mm).

[2] A thick high-strength steel plate having excellent brittle crack propagation arrestability as described in the above [1], further containing in mass %, one or two or more selected from Cu: 0.05 to 1.5%, Cr: 0.05 to 1.0%, Mo: 0.05 to 1.0%, V: 0.005 to 0.10%, B: 0.0002 to 0.0030%.

[3] A thick high-strength steel plate having excellent brittle crack propagation arrestability as described in the above [1] or [2], further containing in mass %, one or two or more selected from Mg: 0.0003 to 0.0050%, Ca: 0.0005 to 0.0030%, REM: 0.0005 to 0.010%.

[4] A method of producing a thick high-strength steel plate having excellent brittle crack propagation arrestability, including:

reheating a steel slab having a composition as described in any of the above [1] to [3] at 950 to 1150°C;

performing rough-rolling (preliminary rolling) the steel at a temperature of not lower than 900°C with a cumulative reduction of 30% or more;

performing finish-rolling the steel at a temperature of not lower than Ar₃ temperature and not higher than T(°C) defined by the below-described formula (3) with a cumulative reduction of 40%;

performing accelerated cooling of the steel with a cooling rate of 8°C/s or more averaged throughout the plate thickness from a temperature not lower than A₃ to a temperature of not higher than 500°C.


where [Ni] denotes the Ni content (in mass %) and t denotes a plate thickness (mm).

[6] A method of producing thick steel plate having excellent brittle crack propagation arrestability as described in the above [5], further including performing tempering of the steel at a temperature of 300 to 600°C after finishing the accelerated cooling.

Effect of the invention

[0012] By applying the present invention, high arrestability steel plates applicable to large constructions can be provided by a stable and effective production method, where the arrestability indicators T_KCa=6000 of the plates are -10°C or less, even when the steel plates are thick plates having a plate thickness of 50 mm or more, and the yield strength of the steel is in a range of 390 to 460 MPa. Therefore, the present invention has large industrial efficiency.

BRIEF EXPLANATION OF DRAWINGS

[0013] 

FIG. 1 shows an example of crystal orientation mapping in accordance with EBSP and analysis of boundary of iso-crack propagation resisting domains.

FIG. 2 is a graph showing a change of arrestability depending on the amount ofNi addition.

FIG. 3 is a graph showing the effects of Ni addition and effective grain size on arrestability.

FIG. 4 is a graph showing a relationship between the pearlite fraction and arrestability.

FIG. 5 is a graph showing the relationship between circle-equivalent diameter of cementite and arrestability.

FIG. 6 is a graph showing the relationship between the arrestability and the fraction of coarse ferrite having a circle equivalent diameter of 25 µm or more in surface portions having a depth of 5% of plate thickness from the front and back surfaces of the plate.

FIG. 7 is a graph showing the relationship between the Ni content and the effective grain size required to provide predetermined arrestability to a steel.

FIG. 8 is a graph showing the plate thickness-dependence of effective grain size required to provide predetermined arrestability to a steel.

FIG. 9 is a graph showing the relationship between the Ni content and finish-rolling temperature required to provide a predetermined arrestability to a steel.

FIG. 10 is a graph showing plate thickness-dependence of the finish-rolling temperature required to provide a predetermined arrestability to a steel.

BEST MODE FOR CARRYING OUT THE INVENTION

[0014] In the following, an embodiment of the present invention is explained in detail.
The inventors carried out experimental research on factors dominating the arrestability of a steel having a yield strength of 390 to 460 MPa and a bainite-dominated (60% or more in volumetric fraction) microstructure. As a result, the inventors found a method that ensures arrestability even in a thick steel plate having a plate thickness of 50 mm or more. Important points in the present invention include new discoveries described in the below (1) to (5).

[0015] 

(1) The unit of fracture during the propagation of a brittle crack does not correspond to apparent crystal grain boundaries but does closely correspond to the crystal grain boundaries obtained by crystal orientation analysis using EBSP. Specifically, a mean circle equivalent diameter (effective crystal grain size) of an iso-crack propagation resisting domain (including grains having circle equivalent diameter of less than 8 µm shows good correlation with the arrestability, where the domain boundaries are defined by identifying <001> axes closest to a T direction perpendicular to the rolling direction. That is, where angles of 20° or more are required to make the identified <001> axes of the grains (exuding grains having circle equivalent diameter of less than 8 µm) to be coincident with each other, domain boundaries are defined between the grains.

[0016] 

(2) Where Ni is added in an amount of 0.5 % or more, the apparent effect of improving the arrestability appears. The effect of the Ni and the effect of grain size refinement are independent from each other, and therefore a substantially additive effect can be obtained. That is, by the addition of Ni, similar arrestability can be ensured even in a coarse microstructure. As a result it is possible to reduce the burden in the production process such as increasing the finish-rolling temperature.

[0017] 

(3) Where pearlite fraction exceeds 5 %, coarse pearlite tends to act as a starting point for brittle fracture thereby decreasing the arrestability even when the steel has a fine effective grain size. In order to avoid such a phenomenon, it is necessary to control the cooling rate of the accelerated cooling and the target (stopping) temperature at which the cooling is stopped.

[0018] 

(4) Fine cementites having a mean circle equivalent diameter of 0.5 µm or less contribute to improve the arrestability. In order to maintain the fine grain size of cementite, it is necessary to control the accelerated cooling after the rolling and the conditions of subsequent heat treatment.

[0019] 

(5) Even though the steel have fine effective grain size averaged by the plate thickness, arrestability of the steel deteriorates when the fraction of coarse ferrite generated in the surface portions exceeds 10%. In order to avoid such a deterioration, it is necessary to control the finish-rolling temperature and the starting temperature of cooling to be not too low of temperatures.

[0020] In the following, the constitutions of the present invention are explained.
In general, the basic structural unit dominating the toughness of bainite steel is not prior-austenite grain size, but the size of a region of a so-called packet or block (With respect to a packet and block, see Non-Patent Reference: Matsuda, Inoue, Mimura, and Okamura, "Toughness and effective grain size of low-alloy heat-treated high-tensile steel", Proc. of Int. Symp. on Towered Improved Ductility and Toughness. Climax Molybdenum Co., Kyoto (1971), p47). The toughness is improved as the region has small size.
However, in accordance with the microstructural observation using an usual optical microscope, it is difficult to measure the size of packets or blocks. Further, where ferrite is mixed in the microstructure, it is very difficult to define the basic structure unit obj ectively.

[0021] Based on the above-describe circumstances, the inventors firstly produced steel plate of 50 mm in plate thickness under various conditions using a steel slab not containing Ni, worked the steel plates to test pieces of 500mm square each notched to a depth of 29 mm, and subjected the test specimens to a temperature-gradient type ESSO test so as to evaluate the arrestability in accordance with the method described in WES 3003. After that, by observing the fracture surface of the test specimens using a scanning electron microscope, and measured the unit (fracture facet size) of cleavage surface surrounded by the fracture portion called a tear-ridge, and confirmed its good correlation with arrestability.

[0022]  Next, EBSP measurement was carried out on a sectional plane perpendicular to the above-described fracture surface. The result of analysis of crystal orientation of grains subjacent to the fracture surface was compared with a photograph. Thus, boundary conditions of the fracture unit were examined carefully. FIG. 1 shows an example of the examination. Analysis was carried out on representative points in iso-orientation domains were analyzed based on the orientation map obtained by EBSP. Cubes constituted of {100} planes (that are regarded as cleavage plane) and the direction of crack propagation assuming that the crack propagates along the {100} planes were shown in FIG. 1.

[0023] The numbers in FIG. 1 denote the angles (crack propagation deviation angles) required to make closest <001> axes to coincident with each other, that is, the angles required to make {100} planes perpendicular to the T-direction coincident with each other by allowing their rotation. By this analysis it was confirmed that orientation of crack propagation appears to change where the crack-propagation deviation angle was 20° or more as shown in (a), (b), (c), and (f). As a result of observation of fracture surface, it was confirmed that the boundary of the above-described critical conditions actually constituted the boundaries of the fracture facet size. On the other hand, in some cases as shown in (d) to (e), crack-propagation does not change its direction in small-sized domains even when the deviation angle is more than 20°C. Such phenomenon is considered to correspond to wraparound of the crack and ductile fracture surface. Such phenomenon is observed in the domains having a circle equivalent diameter smaller than 8 µm. As a result of observation of the fracture surface, it was also confirmed that apparent boundaries were not constituted by such small domains. In the determination of effective grain size, the existent domains smaller than 8 µm may be united to any of the adjacent grains on both sides, and the domain boundaries of iso-crack propagation resisting domain may be determined by examining the deviation angles between the two adjacent domains. As described above, the effective grain size can be estimated by excluding the grains smaller than 8µm from the result of EBSP analysis, determining the boundaries of domains that show crack-propagation deviation angle of not smaller than 20°, and calculating the mean circle equivalent diameter of domains surrounded by the boundaries.

[0024] As a result of careful examination of the relationship between the thus measured effective grain size and the arrestability of steel, it was found that finish-rolling had to be performed at a low temperature of not higher than 800°C so as to provide an arrestability of sufficient level to be applied in large construction steel, and that it was required to have the cooling to be started from high temperature so as to control the yield strength to be 390MPa or more, making it very difficult to produce the steel effectively and stably.

[0025] Therefore, the effect of Ni addition was carefully examined in detail as a solution for the above-described problem. Slabs were cast by variously changing the balances of Ni and Mn contents such that substantially the same microstructure and strength were obtained. Under the same conditions, steel plates of 50 mm in thickness, and steel plates of 80 mm in thickness were produced from the slabs, and their arrestability was examined based on a ESSO test. As a result, it was confirmed that arrestability tended to increase with increasing Ni content even though the effective grain size was almost unchanged. This tendency is shown in FIG. 2.

[0026] Hear, arrestability was evaluated at a temperature (T_Kca=6000) at which brittle-crack propagation arrestability K_ca satisfied K_ca=6000N/mm^1.5. From FIG. 2, it was confirmed that arrestability was improved apparently where Ni content was 0.5% or more. By the observation of fracture surfaces of ESSO test specimens, the appearance of three-dimensional irregularity apparently increased as the Ni content increased. This phenomenon can be explained by soluble Ni enhancing cross slips thereby making the propagation direction of cracks random.

[0027]  Next, in order to individually distinguish and quantify the effects of Ni addition and refining of effective grain size, steel plates obtained by rolling the above-described Ni-containing steel slab under various conditions were subjected to examination of the arrestabilities thereof. As a result, it was found that arrestability-improving effect by grain-refining did not depend on Ni content, and could substantially be summarized by the effect of the Ni content. This behaviour is shown in FIG. 3. That is, by adding an appropriate amount of Ni, it is possible to ensure arrestability without refining the effective grain size. Therefore, where a high production efficiency of the steel is required irrespective of the cost of Ni-alloying, it is possible to increase the finish-rolling temperature by adding Ni, thereby shortening the waiting time for the process temperature. As a result, it is possible to remarkably enhance the productivity of a thick steel.

[0028] The inventors also examined the influence of microstructural factors other than effective grain size on the arrestability, since it was confirmed that steel plates of fine effective grain size occasionally showed insufficient arrestability.

[0029] One of the factors is the occurrence of pearlite mixed in the bainite-dominated structure. Where fraction of pearlite structure increases, arrestability tends to be deteriorated because of an increasing amount of coarse pearlite that acts as crack initiation point of brittle fracture. Therefore, as shown in FIG. 4, it is necessary to control the pearlite fraction to be 5% or less.

[0030] In addition, it was confirmed that the arrestability was also affected by the size of cementite included in the bainite. As shown in FIG. 5, arrestability is reduced where the mean circle equivalent diameter of cementite exceeds 0.5 µm. It can be estimated that fine cementites generate microcracks in their grain boundaries in the matrix before propagation of the main crack, thereby decreasing the stress state in the tip of the crack. On the other hand, where the cementite is coarsened, like as pearlite, the cementite acts as an origin for causing the brittle fracture to occur, thereby decreasing the arrestability.

[0031] Further, it was confirmed that coarse ferrite generated in the surface portions also reduced the arrestability. This surface coarse ferrite is generated where a steel of relatively low hardenability (quenchability) is rolled at a temperature lower than Ar₃ or where the rolling is finished at a temperature not lower than Ar₃ but where the starting temperature of cooling is below Ar₃. As shown in FIG. 6, remarkable reduction of arrestability can be avoided where the fraction of ferrite having a circle equivalent diameter larger than 25 µm is not greater than 10% in the surface portions each having a depth of 5% of the plate thickness from the front and back surfaces of the plate.

[0032] In order to make clear the production principle of a thick high-strength high-arrest steel plate while taking the above-described metallurgical factors into consideration, the influence of the effective grain size, Ni content, and plate thickness on the arrestability were examined in more detail using steel plates that satisfied the above-described conditions regarding pearlite, cementite and surface ferrite. As a result, it was found that the effective grain size needed to be equal to or lower than the below described d.


where [Ni] denotes Ni content ( in mass%) and t denotes a plate thickness (mm).

[0033] In the determination of the above-described d, the linear equation derived from FIG. 7 based on the influence of the effective grain size and Ni on the arrestability of plate having a thickness of 50 mm was combined to the equation of thickness effect derived from FIG. 8 based on the result of testing in which plate thickness was varied by grinding the front and back surfaces of plate of 80 mm in thickness that contained Ni: 2%. Where the effective grain size is greater than the above-described d, tear-ridge is not formed sufficiently frequently when a brittle crack propagates from one grain to another grain. Therefore, the effect of suppressing the propagation of a crack is reduced, and the arrestability is reduced.

[0034] Next, the reason for the limitation of production conditions is explained.
In the present invention, the reheating temperature of a slab was controlled to be 950 to 1150°C. Where the reheating temperature is lower than 950°C, homogenization of the alloying elements is insufficient, thereby causing inhomogeneous properties. Where the reheating temperature exceeds 1150°C, the grain sizes of austenite are coarsened. Therefore, there is a possibility that it is difficult to obtain a fine microstructure in the final state.

[0035] The subsequent rough rolling must be performed at a temperature of 900°C or more with a cumulative reduction of 30% or more. If the above-described conditions are not satisfied, recrystallization of austenite grains does not proceed sufficiently, resulting in mixed grain microstructure which may cause inhomogeneous properties.

[0036] The following finish-rolling is the most important process for refining the effective grain size that dominates the arrestability. The finish-rolling is performed at a temperature of not lower than Ar₃(a temperature at which formation of ferrite from austenite starts during cooling of steel) and not higher than the below described T(°C) with a cumulative reduction of 40% or more.


where [Ni] denotes Ni content (in mass %) and t denotes a plate thickness (mm).

[0037] In the above-described T, a linear equation is combined with an equation of thickness effect, where the linear equation is derived from FIG. 9 that shows the relationship between the Ni-content and the finish-rolling temperature required for satisfy T_Kca=6000 ≤ -10°C based on the above-described experimental result, and the equation of thickness effect is derived from FIG. 10 based on the experimental results obtained while variously changing the plate thickness and finish-rolling temperature using slabs that contained 2% of Ni. Where the temperature is lower than Ar₃, coarse ferrites having circle equivalent diameters larger than 25µm are generated, thereby decreasing the arrestability, the strength, the toughness, and the ductility of the steel plate. On the other hand, where the temperature exceeds the above-described T, or where the cumulative reduction is less than 40%, arrestability is reduced since the effective grain size is not refined sufficiently. By selecting a temperature slightly lower than the above-described T in accordance with the amount of added Ni, it is possible to shorten the waiting time for process temperature before the finish-rolling, thereby making it possible to produce thick high-strength steel plates effectively.

[0038] After the completion of finish-rolling, the steel plate is subjected to accelerated cooling from the temperature of not lower than Ar₃ to a temperature of 500°C or less with a cooling rate of 8°C/s or more. Where the starting temperature of cooling is lower than Ar₃, the fraction of coarse ferrite in the surface portions exceeds 10%, thereby deteriorating the arrestability. Where the cooling rate is less than 8°C/s, or where the finish temperature of cooling is higher than 500°C, sufficient strength is not obtained. In addition, arrestability is reduced by insufficient refinement of effective grain size, coarsening of cementite that could contribute the improvement of arrestability, or by generation of pearlite exceeding 5%.

[0039] After the accelerated cooling, a tempering treatment may be performed at a temperature of 300 to 600°C so as to control the strength and toughness of the steel plate. Where the tempering temperature is less than 300°C, ductility and toughness are not improved sufficiently. Where the tempering temperature exceeds 600°C, arrestability is reduced by coarsening of cementite.

[0040] Next, the reasons for limiting the composition of the present invention are explained.
C (carbon) is an element that contributes to generation of cementite and preventing coarsening of microstructure. In addition, carbon is an inevitable element for enhancing the strength of steel at low cost. Therefore, carbon is added in an amount of 0.01 % or more. On the other hand, too much addition of carbon makes it difficult to assure a HAZ (Heat Affected Zone) toughness in the time of large heat input welding, and easily coarsens cementite. Therefore, the upper limit of carbon content is controlled to be 0.14%.

[0041] Si (silicon) is an inexpensive deoxidizing element and is added in an amount of 0.03% or more for solid solution-strengthening of matrix. On the other hand, a silicon content exceeding 0.5% deteriorates weldability and HAZ toughness. Therefore, the upper limit of silicon content is controlled to be 0.5%.

[0042] Mn (manganese) is added in an amount of 0.3% or more since manganese is an effective element for improving strength and toughness of the steel. On the other hand, excessive Mn deteriorates HAZ toughness and weld-crack property. Therefore, the upper limit of manganese content is controlled to be 2.0%.

[0043] Although the content of P (phosphorus) and S (sulfur) are preferably controlled to be as low as possible, large cost is required to reduce the content of P and S industrially. Therefore, the upper limits are controlled to be 0.02% for P and 0.01 % for S.

[0044] Ni (nickel) is added in an amount of 0.5% or more since nickel is effective for assuring the strength and for improving the arrestability and HAZ toughness. The content of Ni is controlled to be not more than 4.0% since increasing amounts of Ni result in increasing costs of a slab.

[0045] Where added in a small amount, Nb (niobium) is an element that contributes to refining of microstructure, transformation strengthening, and precipitation strengthening, and is effective for ensuring the strength of the matrix. Therefore, Nb is added in an amount of 0.005% or more. On the other hand, excessive addition of Nb hardens the HAZ and deteriorates the toughness remarkably. Therefore, the upper limit ofNb is controlled to be 0.050%.

[0046] Where added in a small amount, Ti (titanium) is effective in refining the structure, precipitation strengthening that improve strength and toughness of the base metal, and generation of TiN that improves HAZ toughness of the welded joint. Therefore, Ti is added in an amount of 0.005% or more. On the other hand, excessive addition of Ti remarkably deteriorates HAZ toughness. Therefore, the upper limit of Ti is controlled to be 0.050%.

[0047] Al (aluminum) is an important deoxidizing element and is added in an amount of 0.002% or more. On the other hand, excessive addition of aluminum deteriorates the surface quality of slab and forms inclusions that disturb the toughness of the steel. Therefore, the upper limit of aluminum is controlled to be 0.10%.

[0048] N (nitrogen) is combined with Ti and forms nitrides that improves HAZ toughness. Therefore, N is added in an amount of 0.0010% or more. On the other hand, excessive addition ofN generates brittleness by soluble N. Therefore, the amount ofNi is controlled to be 0.0080% or less.
Optional additional elements are limited for the below described reason.

[0049] Each of Cu (copper), Cr (chromium), and Mo (molybdenum) enhances hardenability and is effective at strengthening the steel. Therefore, they are added in an amount of 0.05% or more. On the other hand, Cu is limited to be 1.5% or less, and Cr and Mo are limited to be 1.0% or less since their excessive addition deteriorates HAZ toughness.

[0050] V (vanadium) contributes to enhancement of strength by the effect of precipitation strengthening. Therefore, V is added in an amount of 0.005% or more. On the other hand, its upper limit is controlled to be 0.10% since an addition of V exceeding 0.10% reduces HAZ toughness.

[0051] B (boron) is an element for improving hardenability and is effective for enhancing the strength of steel by addition thereof in an appropriate amount. On the other hand, excessive addition of B deteriorates weldability. Therefore, boron content is controlled to be 0.0002 to 0.0030%.

[0052] Mg (magnesium), Ca (calcium), and REM form fine oxides or sulfides and contribute to improvement of HAZ toughness. On the other hand, excessive addition coarsens the inclusions and reduces toughness. Therefore, Mg is controlled to be in a range of 0.0003 to 0.0050%, Ca is controlled to be in a range of 0.0005 to 0.0030%, REM is controlled to be in a range of 0.0005 to 0.010%. REM denotes a rare earth element (rare earth metal) such as La, Ce, or the like.
Further, in order to ensure base-metal strength and joint-strength consistently, Ceq shown by the below described formula must be controlled to be in a range of 0.30 to 0.50%. Where Ceq is less than 0.30%, it is difficult to ensure the yield strength to be 390MPa or more in a base metal made of thick steel having a plate thickness of 50 mm or more. Where Ceq exceeds 0.50%, it is difficult to ensure weldability and joint toughness. In addition, there is a possibility that arrestability is reduced by too high of a strength.


where [symbol of element] denotes a content (in mass %) of the element. That is, where X denotes the symbol of an element, [X] denotes the content (in mass %) of the element X.

Examples

[0053] In the following, the effects of the present invention is shown apparently in accordance with examples. The present invention is not limited to the below described examples, and can be embodied in modified manner within a range of a scope of the invention.

[0054] Steel plates each having a thickness of 50 to 80 mm were produced in accordance with the production method shown in Tables 2 and 3 using slabs each having a composition shown in Table 1. The microstructure, base-metal strength, and arrestability of the steel plates are shown in Tables 4 and 5.
Fraction of surface coarse ferrite (surface coarse a fraction) was measured based on image analysis of an optical micrograph of a T-section of an outermost surface portion of a steel plate.
Fraction of pearlite was measured based on an optical micrograph of a T-sections obtained from a portion of 5 mm depth from the plates surface, a portion of 1/4 plate thickness, and central portion of the plate thickness.
Replica samples were obtained from the above-described three portions along the plate thickness. Grain size of cementite (0 diameter) was determined as the mean circle equivalent diameter calculated from photographs taken by a transmission electron microscope from the replica samples.
EBSP samples were obtained from the above-described three portions along the plate thickness such that T sections were subjected to the measurement. In each sample, after measuring 500 × 500 µm region with every 1 µm pitch, orientation analysis was carried out with a 3 to 5 µm pitch for a running length of 2 mm based on a crystal orientation map. Thus grain boundaries were determined. Then, an effective grain size was calculated by a dissection method in accordance with JIS G0551.
Yield strength (YP) and tensile strength (TS) were evaluated using a tensile strength test specimen having a dimension of JIS Z 2201 No. 4 sampled along T direction from the center portion of the plate thickness.
Arrestability was evaluated by temperature-gradient type ESSO test based on a temperature at which Kca=6000N / mm^1.5 was satisfied.

[0055] 

Table 1

(mass%)
Steeel	C	Si	Mn	P	S	Ni	Nb	Ti	Al	N	Cu	Cr	Mo	V	B	Mg	Ca	REM	Ceq
A	0.033	0.22	1.42	0.013	0.003	0.56	0.036	0.018	0.026	0.0037	-	-	-	-	-	-	0.0016	-	0.31
B	0.090	0.15	0.72	0.005	0.003	1.80	0.007	0.015	0.025	0.0041	-	-	-	-	0.0011	-	-	-	0.33
C	0.072	0.36	1.03	0.006	0.004	1.22	0.011	0.007	0.011	0.0021	0.28	0.18	-	-	-	-	-	-	0.38
D	0.061	0.24	0.56	0.008	0.002	3.54	0.006	0.009	0.022	0.0031	-	-	-	-	-	-	-	-	0.39
E	0.080	0.45	1.00	0.007	0.003	1.41	0.018	0.022	0.075	0.0054	0.45	-	-	0.044	-	-	-	-	0.38
F	0.110	0.33	0.65	0.009	0.002	1.98	0.009	0.013	0.030	0.0038	-	-	0.56	-	-	-	-	-	0.46
G	0.077	0.20	1.50	0.004	0.002	1.16	0.013	0.009	0.039	0.0018	-	-	-	-	0.0015	-	-	0.0041	0.40
H	0.080	0.18	1.03	0.006	0.004	1.56	0.011	0.010	0.006	0.0041	-	0.72	-	-	-	0.0023	-	-	0.50
I	0.014	0.14	1.71	0.010	0.003	0.75	0.014	0.042	0.019	0.0035	1.39	-	-	0.066	0.0025	-	-	-	0.45
J	0.132	0.23	0.79	0.008	0.002	0.97	0.009	0.008	0.028	0.0036	-	-	-	-	-	0.0018	0.0012	-	0.33
K	0.083	0.27	1.44	0.006	0.002	0.63	0.005	0.009	0.020	0.0027	-	-	-	-	-	-	-	-	0.37
L	0.093	0.08	1.37	0.007	0.005	1.11	0.012	0.012	0.015	0.0030	0.84	-	-	-	-	-	-	-	0.45
M	0.058	0.04	1.05	0.005	0.002	1.46	0.017	0.018	0.024	0.0031	-	0.68	-	-	-	-	-	-	0.47
N	0.083	0.16	1.16	0.004	0.003	1.37	0.025	0.020	0.038	0.0071	-	-	-	0.089	-	-	-	-	0.39
O	0.052	0.25	1.20	0.006	0.004	2.88	0.011	0.011	0.029	0.0046	-	-	-	-	-	0.0030	-	-	0.44
P	0.030	0.30	0.97	0.012	0.008	2.35	0.020	0.007	0.012	0.0022	-	-	-	-	-	-	-	0.0028	0.35
Q	0.060	0.12	1.02	0.008	0.001	0.52	0.008	0.014	0.032	0.0025	-	0.50	0.34	-	-	-	0.0009	-	0.43
R	0.146	0.34	0.92	0.003	0.003	1.20	0.010	0.008	0.018	0.0026	-	-	-	-	-	-	-	-	0.38
S	0.099	0.45	0.98	0.010	0.009	0.30	0.015	0.020	0.072	0.0050	0.44	0.14	0.15	0.050	-	-	-	-	0.38
T	0.071	0.30	1.76	0.017	0.008	2.38	0.020	0.008	0.014	0.0025	-	-	-	-	-	-	-	0.0070	0.52

Underline denotes values outside the present invention.

[0056] 

[0057] 

[0058] 

[0059] 

[0060] Steel plates No. 1 to 22 as Examples according to the present invention had chemical compositions within a predetermined range and were produced under predetermined conditions. Therefore, each of the plates had sufficient strength as a steel of YP: 390 to 460MPa class, and had satisfactory arrestability.
On the other hand, in steel plates No. 23 to 45 as Comparative Examples, the chemical composition or production condition was outside the predetermined range of the present invention. As a result, arrestability was reduced in each case.
In Nos. 23 and 41, the finish-rolling temperatures were lower than Ar₃, and coarse ferrite was generated in large amounts in the surface portions, resulting in reduction of arrestability.
In Nos.28 and 42, the finish-rolling temperature was not lower than Ar₃, but accelerated cooling was started at a temperature lower than Ar₃. Therefore, fractions of surface coarse ferrite were increased resulting in reduction of arrestability.
In Nos. 24 and 37, accelerated cooling was performed at a low cooling rate.
In Nos. 33 and 40, the stopping temperature of cooling was higher than 500°C.
In Nos. 26 and 38, the heat treatment temperatures were higher than 600°C. Therefore, cementite had a large size in each of the cases, and sufficient arrestability could not be obtained.
No. 34 was subjected to air cooling without performing accelerated cooling. Therefore, effective grain size was not refined, resulting in a reduction of arrestability.
Nos. 27 and 35 were rolled with a small cumulative reduction of the finish-rolling.
In Nos. 25, 30, and 36, the effective grain size was coarsened because of high finish-rolling temperatures, resulting in a reduction of arrestability.
No. 29 was reheated at high temperature.
Nos. 31 and 39 were rough-rolled with small cumulative reductions.
No. 32 was heated at high temperature and was rough-rolled with a small cumulative reduction. Therefore, effective grain size was coarsened, resulting in reduction of arrestability.
In No. 43, cementite had large size because of excessive C content. As a result, the arrestability was reduced, and the HAZ toughness was reduced.
In No.44, arrestability was insufficient because of the small Ni content.
In No. 45, the strength was increased too much because of high Ceq, resulting in reduction of arrestability.

INDUSTRIAL APPLICABILITY

[0061] By applying the present invention, high arrestability steel plates applicable to large construction can be provided by a stable and effective production method, where the arrestability indicators T_Kca=6000 of the plates are -10°C or less, even when the steel plates are thick plates having a plate thickness of 50 mm or more, and yield strength of the steel is in a range of 390 to 460 MPa. Therefore, the present invention has large industrial applicability.

Claims

1. A thick high-strength steel plate, comprising:

a composition containing in mass %, C: 0.01 to 0.14%, Si: 0.03 to 0.5%, Mn: 0.3 to 2.0%, P: 0.020% or less, S: 0.010% or less, Ni: 0.5 to 4.0%, Nb: 0.005 to 0.050%, Ti: 0.005 to 0.050%, Al: 0.002 to 0.10%, N: 0.0010 to 0.0080%, and the balance consisting of Fe and unavoidable impurities, where Ceq defined by the below described formula (1) is 0.30 to 0.50%; and

a microstructure dominated by bainite of 60% or more in volume fraction, wherein

a pearlite fraction is not greater than 5%,

a fraction of coarse ferrite having a circle equivalent diameter larger than 25 µm is not greater than 10% in surface portions each having a depth of 5% of the plate thickness from the front and back surfaces of the plate,

a mean circle equivalent diameter of cementite is not greater than 0.5µm,

and

a mean circle equivalent diameter of iso-crack propagation resisting domains is not smaller than 8 µm and not greater than d (µm) defined by the below described formula (2),

where each of the iso-crack propagation resisting domains is defined such that

a section perpendicular to a rolling direction of the plate is defined as a T section,

a direction in parallel to the plate surface is defined as a T direction in the T section,

inner portion in the T section is defined as a portion excluding the surface portions,

microstructure of the T section is partitioned to domains each having the same crystal orientation (iso-orientation domain) by orientation analysis of the inner portion using Electron Back Scattering Pattern,

an arbitrary measurement line is drawn on the T section microstructure partitioned to the iso-orientation domains by applying a dissection method in accordance with JIS G0551,

a plurality of the iso-orientation domains having a circle equivalent diameter of 8 µm continuously aligned and adjacent to each other on the measurement line are identified excluding iso-orientation domains having a circle equivalent diameter of smaller than 8µ m,

<001> axis closest to the T direction is selected from three <001> axes of each of the identified iso-orientation domains,

an angle (crack-propagation deviation angle) formed by the <001> axes closest to the T direction of the identified iso-orientation domains is measured in each adjacent pair of the iso-orientation domains on the measurement line,

the continuously aligned iso-orientation domains satisfying the angle of not greater than 20° and adjacent iso-orientation domains having a circle equivalent diameter of less than 8µm are regarded to constitute one iso-crack propagation resisting domain.




where [X] denotes content (in mass %) of an element X and t denotes a plate thickness (mm).

2. A thick high-strength steel plate according to claim 1, further containing in mass %, one or two or more selected from Cu: 0.05 to 1.5%, Cr: 0.05 to 1.0%, Mo: 0.05 to 1.0%, V: 0.005 to 0.10%, and B: 0.002 to 0.0030%.

3. A thick high-strength steel plate according to claim 1, further containing in mass %, one or two or more selected from Mg: 0.0003 to 0.0050%, Ca: 0.0005 to 0.0030%, and REM: 0.0005 to 0.010%.

4. A thick high-strength steel plate according to claim 2, further containing in mass %, one or two or more selected from Mg: 0.0003 to 0.0050%, Ca: 0.0005 to 0.0030%, and REM: 0.0005 to 0.010%.

5. A method of producing a thick high-strength steel plate, comprising:

reheating a steel slab having a composition according to any one of claims 1 to 4 at 950 to 1150°C;

performing rough-rolling the steel at a temperature of not lower than 900°C with cumulative reduction of 30% or more;

performing finish-rolling the steel at a temperature of not lower than Ar₃ temperature and not higher than T(°C) defined by the below described formula (3) with a cumulative reduction of 40%;

performing accelerated cooling of the steel with a cooling rate of 8°C/s averaged throughout the plate thickness from a temperature not lower than A₃ point to a temperature of not higher than 500°C.


where [Ni] denotes Ni content (in mass %) and t denotes a plate thickness (mm).

6. A method of producing thick steel plate according to claim 5, further comprising performing tempering of the steel at a temperature of 300 to 600°C after finishing the accelerated cooling.

Drawing