JP4946092B2 - High-strength steel and manufacturing method thereof - Google Patents

Description

  The present invention relates to a high-strength steel used for offshore structures, line pipes, pressure vessels, and the like, and its manufacturing method. In particular, the yield stress is 355 MPa or more, and the strength and toughness of the base metal are not only excellent, but also the toughness of the welded portion. The present invention relates to a high-strength steel excellent in (CTOD characteristics) and a method for producing the same.

  Steel used for offshore structures and the like is usually finished by welding to a desired shape. Therefore, from the viewpoint of ensuring the safety of the structure, these steels are not only excellent in the strength and toughness of the base metal itself, but also in the toughness of the welded joint (welded metal or heat affected zone) of the welded joint. It is required to be excellent.

  As an evaluation standard for the toughness of steel, conventionally, absorbed energy mainly by Charpy test has been used. However, in recent years, a crack opening displacement test (hereinafter abbreviated as “CTOD test”) is often used in order to increase reliability. In this test, a test piece in which a fatigue precrack was generated in the toughness evaluation part was bent at three points, and the amount of opening (plastic deformation) at the crack bottom just before the fracture was measured to evaluate the resistance to brittle fracture. To do.

  By the way, the steel having a large thickness as used in the above-mentioned applications is generally subjected to multi-layer welding. However, in such welding, the heat-affected zone receives a complicated heat history, so that a local embrittlement region occurs. Easy, especially in the bond part (boundary between the weld metal and the base metal) and the two-phase region reheat part (the region that becomes coarse in the first cycle of welding and heated to the two-phase region of α and γ in the second cycle) There is a problem that the reduction in toughness is large. This is because the bond portion is exposed to a high temperature just below the melting point, so that austenite grains are coarsened and are easily transformed into a fragile upper bainite structure by subsequent cooling. In addition, since a brittle structure such as a woodman-stetten structure or island martensite is generated in the bond portion, the toughness is further reduced.

  As a countermeasure against the above problem, for example, a technique of finely dispersing TiN in steel to suppress coarsening of austenite grains or to use as a ferrite transformation nucleus has been put into practical use. Furthermore, Patent Document 1 and Patent Document 2 disclose a technique for suppressing the austenite grain growth and improving the toughness of the welded portion by adding a rare earth element (REM) in combination with Ti and dispersing fine particles in the steel. Is disclosed. In addition, technology to disperse Ti oxide, technology to combine ferrite nucleation ability of BN and oxide dispersion, and addition of Ca and REM to control the form of sulfides to achieve high toughness Obtaining technology has also been proposed.

  On the other hand, the two-phase region reheat zone, that is, the region exposed to a high temperature immediately below the melting point in the first welding, the region that becomes the two-phase region of ferrite and austenite by reheating during the subsequent lap welding is most brittle. The reason for this is that carbon is concentrated in the austenite region by reheating after the second pass, and this generates a fragile bainite structure including island martensite during cooling, thereby reducing toughness. Therefore, as a countermeasure, a technique is disclosed in which generation of island martensite is suppressed by reducing C and Si, and the strength of the base material is ensured by adding Cu (for example, Patent Document 3). reference).

Japanese Patent Publication No. 03-053367 JP 60-184663 A Japanese Patent Laid-Open No. 05-186823

  However, although the above-described problem of poor toughness of the heat affected zone has been improved to some extent by the above prior art, some problems to be solved still remain. For example, in the technology using TiN, the action is lost in the bond portion heated to a temperature range where TiN dissolves, and further, the toughness is significantly lowered due to the embrittlement of the base structure due to solute Ti and solute N. Sometimes. Further, the technology using Ti oxide has a problem that fine dispersion of the oxide cannot be sufficiently homogeneous. Furthermore, in recent years, as marine structures, ships, and the like are increased in size, steel materials used therefor have been increased in strength and thickness. In order to achieve these problems, it is effective to add a large amount of alloy elements, contrary to the technique of Patent Document 3. However, the addition of a large amount of alloy elements has a problem in that the toughness of the heat-affected zone is reduced.

  Therefore, the object of the present invention is to solve the above-mentioned problems of the prior art, and to increase the tensile strength of the base metal and the toughness of the weld heat affected zone without increasing the amount of alloy elements added, It is to propose a steel and its advantageous production method.

  The inventors diligently studied a method capable of improving the base material strength and toughness of high-tensile steel and also improving the toughness of the weld heat affected zone. As a result, the deterioration of the toughness of the weld heat affected zone is caused by the formation of an embrittled structure. Therefore, in order to improve the toughness of the weld heat affected zone, the coarseness of austenite grains in the region heated at high temperature during welding Further, it was found that it is effective to uniformly disperse transformation nuclei for promoting ferrite transformation during cooling after welding.

  Thus, as a result of further study on the method for suppressing the formation of the embrittlement structure, the inventors are effective to control the addition amount of Ca added for the morphology control of the sulfide within an appropriate range. In addition, it has been found that the addition of Ni is effective for improving the toughness (CTOD characteristics) of the weld heat affected zone.

  In addition, when the influence of rolling conditions on the strength and toughness of the base metal was examined, the cooling after rolling was changed to two-stage cooling consisting of a pre-stage cooling with a large cooling rate and a small post-stage cooling, and each cooling rate was appropriately set. It was found that if controlled, the steel sheet structure becomes a structure mainly composed of acicular ferrite, and high strength steel excellent in the strength and toughness of the base material can be produced. Furthermore, in order to further increase the strength and toughness of the base material, it is important to effectively use Nb, which has a large effect of forming a non-recrystallized region in the low temperature region of austenite. The present inventors have found that it is necessary to regulate the upper limit more strictly than in the past, and completed the present invention.

That is, the present invention is C: 0.05 to 0.1 mass%, Si: 0.05 to 0.5 mass%, Mn: 1 to 2 mass%, P: 0.015 mass% or less, S: 0.005 mass% or less. , Al: 0.005 to 0.06 mass%, Ni: 0.3 to 2 mass%, Nb: 0.004 to 0.05 mass%, Ti: 0.005 to 0.02 mass%, N: less than 0.003 mass% Ca: 0.0005 to 0.003 mass%,
Ca, S and O are represented by the following formula (1):
0 <(Ca− (0.18 + 130 × Ca) × O) /1.25/S <1 (1)
Here, Ca, S and O are the contents of each element (mass%).
Is a high-strength steel characterized in that it has a component composition consisting of Fe and inevitable impurities .

In addition to the above component composition, the high-tensile steel of the present invention further includes B: 0.0003 to 0.0025 mass%, Mo: 0.7 mass% or less, V: 0.022 mass% or less, Cu: 1 mass% or less, and Cr: One or more selected from 0.7 mass% or less are contained.

In the present invention, C: 0.05 to 0.1 mass%, Si: 0.05 to 0.5 mass%, Mn: 1 to 2 mass%, P: 0.015 mass% or less, S: 0.005 mass% or less , Al: 0.005 to 0.06 mass%, Ni: 0.3 to 2 mass%, Nb: 0.004 to 0.05 mass%, Ti: 0.005 to 0.02 mass%, N: less than 0.003 mass% Ca: 0.0005 to 0.003 mass%,
Ca, S and O are represented by the following formula (1):
0 <(Ca− (0.18 + 130 × Ca) × O) /1.25/S <1 (1)
Here, Ca, S and O are the contents of each element (mass%).
After heating a steel slab having a composition composed of Fe and unavoidable impurities to 1050 to 1200 ° C., the cumulative rolling reduction in a temperature range of 950 ° C. or higher is 30% or more and less than 950 ° C. Pre-stage cooling in which the cumulative rolling reduction in the region is 30 to 70%, and thereafter cooling from 5 to 20 ° C./s from the hot rolling end temperature to the cooling stop temperature between 600 to 450 ° C., We propose a method for producing high-strength steel, characterized in that post-stage cooling is performed in which cooling is carried out at a temperature of less than 1 to 5 ° C./s from the temperature of the former stage of cooling to a cooling stop temperature of less than 450 ° C. to 200 ° C.

The steel slab in the method for producing high-tensile steel according to the present invention is further provided with B: 0.0003 to 0.0025 mass%, Mo: 0.7 mass% or less, V: 0.022 mass% or less in addition to the above component composition. Cu: 1% by mass or less and Cr: 0.7% by mass or less selected from 1 type or 2 types or more.

Moreover, the manufacturing method of the high-tensile steel of this invention is characterized by performing a tempering process at 450-650 degreeC to the steel after the said back | latter stage cooling.

  According to the present invention, since the base material can produce a high-strength steel having a yield stress of 355 MPa or more and excellent toughness and also excellent toughness (CTOD characteristics) of the weld heat affected zone, it can be manufactured at low cost. This greatly contributes to increasing the size of ships and ships.

The basic technical idea of the present invention will be described.
The first feature of the present invention is to effectively use the crystallization of a Ca compound (CaS) added for the purpose of controlling the form of sulfides in order to improve the toughness of the weld heat affected zone. Since this CaS crystallizes at a lower temperature than the oxide, it can be uniformly finely dispersed. And by controlling the amount of CaS added and the amount of dissolved oxygen in the molten steel at the time of addition to an appropriate range, solid solution S is ensured even after crystallization of CaS, and MnS precipitates on the surface of CaS to form composite sulfide. Form. This MnS is known to have a ferrite nucleation ability. Further, since a Mn dilute band is formed around the precipitated MnS, the ferrite transformation is further promoted. Moreover, since ferrite-forming nuclei such as TiN, BN, and AlN also precipitate on the precipitated MnS, the ferrite transformation is further promoted.

  According to the above technique, it becomes possible to finely disperse the ferrite transformation formation nuclei that do not dissolve even at high temperatures, and the structure of the weld heat affected zone can be made into fine ferrite pearlite, thereby achieving high toughness. Even in the region that is reheated to the two-phase region due to the thermal cycle during multi-layer welding, the structure of the heat-affected zone by the first welding is refined, so that the toughness of the untransformed region is improved, and Since the transformed austenite grains are also refined, the degree of reduction in toughness can be suppressed small.

  The second feature of the present invention is that the amount of N is limited to less than 0.003 mass% as a raw material component in order to improve the strength and toughness of the base material. Addition of Nb is effective to improve the strength and toughness of the base metal. However, Nb is easily dissolved in the nitride, so if a large amount of nitride exists, Much Nb is dissolved, and it becomes impossible to secure a solid solution Nb amount effective for improving strength and toughness. Therefore, by suppressing the amount of N to less than 0.003 mass% and suppressing the formation of nitride, the effect of Nb can be utilized to the maximum, and consequently the strength and toughness of the base material can be improved. it can. Further, by reducing the N amount, the surface defects of the continuous cast slab are reduced, and the product yield can be improved.

The third feature of the present invention lies in that the cooling after rolling the steel material is divided into two stages of the pre-stage cooling and the post-stage cooling, and the cooling rate of the pre-stage cooling is controlled to be greater than the post-stage cooling. This point will be described based on experimental results.
C: 0.08 mass%, Si: 0.2 mass%, Mn: 1.4 mass%, Ni: 0.4 mass% After heating the steel slab to 1150 ° C, the cumulative reduction ratio of 950 ° C or higher Hot rolling is performed at 40%, a cumulative reduction ratio of less than 950 ° C. is 50%, and a rolling end temperature is 850 ° C., and then the rolling end temperature to 500 ° C. is cooled at a cooling rate of 2 to 25 ° C./s. After the pre-stage cooling, further post-stage cooling was performed by cooling to 350 ° C. at a cooling rate of 3 ° C./s, and then air cooling was performed to obtain a thick steel plate having a plate thickness of 10 to 50 mmt. About this thick steel plate, the area ratio of the acicular ferrite structure, the tensile strength characteristics, and the toughness (Charpy absorbed energy) at −40 ° C. were measured.

In general, when the structure is changed to a high-strength structure composed of a ferrite-pearlite structure, a relatively coarse upper bainite structure including island martensite and the like between the laths is formed, and the toughness is greatly reduced. Therefore, in order to achieve both high strength and high toughness, it is necessary to obtain a fine acicular ferrite structure by devising rolling conditions.
FIG. 1 shows the influence of the cooling rate of the pre-cooling on the base material strength and the acicular ferrite area ratio. From this figure, it can be seen that the strength increases and the toughness tends to decrease as the cooling rate of the pre-stage cooling increases. On the other hand, the area ratio of the acicular ferrite structure increases with an increase in the cooling rate. However, when the cooling rate increases at about 10 ° C./s or more, the rising gradient becomes gentle, that is, by increasing the cooling rate of the preceding stage cooling to a certain rate or more. It was found that a steel sheet excellent in strength-toughness balance can be produced by suppressing the formation of polygonal ferrite produced at a relatively high temperature to form a structure mainly composed of acicular ferrite.
On the other hand, when the latter cooling rate is faster than the former cooling rate, island martensite is generated, and the toughness of the base material is deteriorated. However, it has also been found that if it is too slow, the strength of the base material will be reduced, so that it is necessary to control within an appropriate range.

Next, the reason for limiting the component composition of the high-strength steel according to the present invention will be described.
C: 0.05-0.1 mass%
C is an element that has the greatest influence on the strength of the steel. In order to ensure the strength necessary for structural steel (YS ≧ 355 MPa), it is necessary to contain 0.05 mass% or more. However, conversely, if too much, weld cracking is caused, so the upper limit is made 0.1 mass%.

Si: 0.05-0.5 mass%
Si is a component added as a deoxidizer, and it is necessary to add 0.05 mass% or more. On the other hand, when it exceeds 0.5 mass%, it is necessary to make it 0.5 mass% or less in order to reduce the toughness of the base material.

Mn: 1 to 2 mass%
Mn needs to be added in an amount of 1 mass% or more in order to ensure the strength of the base material. On the other hand, if it exceeds 2 mass%, the toughness of the welded portion is remarkably lowered, so it is necessary to set it to 2 mass% or less. Preferably, it is in the range of 1.2 to 1.8 mass%.

P: 0.015 mass% or less P is an impurity that is inevitably mixed. If it exceeds 0.015 mass%, the toughness of the welded portion is lowered, so that it is limited to 0.015 mass% or less. Preferably, it is 0.012 mass% or less.

S: 0.005 mass% or less S is an impurity that is inevitably mixed, and if contained in excess of 0.005 mass%, the toughness of the base metal and the welded portion is lowered, so the content is made 0.005 mass% or less. Preferably, it is 0.0035 mass% or less.

Al: 0.005-0.06 mass%
Al is an element added to deoxidize molten steel, and it is necessary to contain 0.005 mass% or more. On the other hand, if added over 0.06 mass%, the toughness of the base metal is reduced and mixed into the weld metal part by dilution by welding to reduce toughness. Therefore, it is necessary to limit to 0.06 mass% or less. .

Ni: 0.3-2 mass%
Ni is an element effective for improving the strength of steel and the CTOD characteristics of the weld heat affected zone. This effect is manifested by addition of 0.3 mass% or more. However, since Ni is an expensive element, the upper limit is set to 2 mass%.

Nb: 0.004 to 0.05 mass%
Since Nb forms a non-recrystallized region in the low temperature range of austenite, the base material structure can be refined and toughened by rolling in that temperature range. In addition, precipitation strengthening can be achieved by performing a tempering treatment after rolling and cooling. Therefore, Nb is an important additive element from the viewpoint of strengthening steel. In order to acquire the said effect, Nb needs to add 0.004 mass% or more. However, when Nb is added excessively in excess of 0.05 mass%, the toughness of the welded portion is deteriorated, so the upper limit is made 0.05 mass%.

Ti: 0.005-0.02 mass%
Ti precipitates as TiN when the molten steel is solidified, and serves as a suppression of coarsening of austenite and ferrite transformation nuclei in the welded portion, thereby contributing to high toughness. If it is less than 0.005 mass%, the effect is small. On the other hand, if it exceeds 0.02 mass%, the effect expected by the coarsening of TiN particles cannot be obtained. Therefore, the amount of Ti added is in the range of 0.005 to 0.02 mass%.

N: Less than 0.003 mass% N is required to be less than 0.003 mass% in order to ensure the amount of solute Nb necessary for improving the strength and toughness of the base material. As described above, Nb is an element effective for refining the structure of the base material, increasing the toughness, and precipitation strengthening. In order to obtain these effects, Nb needs to be in a solid solution state during heating before rolling. However, when Nb and Ti are added at the same time, a (Ti, Nb) (C, N) composite carbonitride is formed, and this precipitate exists stably up to a high temperature as compared with NbC. In the heating stage, a part of the solid solution remains without being dissolved, and the solid solution Nb is reduced. Furthermore, when the amount of N increases, (Ti, Nb) (C, N) tends to be more difficult to dissolve, so the amount of solid solution Nb further decreases. Therefore, in order to effectively use the effect of the solute Nb, the upper limit of the N amount is limited after adding an amount of Ti necessary for ensuring the toughness of the welded portion. 0.003 mass%.

Ca: 0.0005 to 0.003 mass%
Ca has an effect of fixing S and improving toughness. In order to exhibit this effect, it is necessary to add at least 0.0005 mass%. However, even if it contains 0.003 mass% or more, since the effect is saturated, Ca is made into the range of 0.0005 to 0.003 mass%.

0 <(Ca− (0.18 + 130 × Ca) × O) /1.25/S <1
In order to finely disperse the ferrite transformation nuclei CaS that does not dissolve even at high temperatures, Ca, S and O are represented by the following formula (1):
0 <(Ca− (0.18 + 130 × Ca) × O) /1.25/S <1 (1)
Here, Ca, S, O: Content of each element (mass%)
It is necessary to contain and satisfy the relationship. In the above formula, (Ca− (0.18 + 130 × Ca) × O) / (1.25 / S) is a value indicating the ratio of the atomic concentrations of Ca and S effective for sulfide morphology control, and this The form of sulfide can be estimated from the value (Mochida et al., “Iron and Steel”, Japan Iron and Steel Institute, 66th (1980), No. 3, P354-362). In the case of (Ca− (0.18 + 130 × Ca) × O) / (1.25 / S) ≦ 0, since CaS does not crystallize, S precipitates in the form of MnS alone. It is not possible to achieve fine dispersion of ferrite nuclei at the weld heat affected zone, which is the main focus. Further, MnS precipitated alone is elongated during rolling of the steel sheet, causing a reduction in the toughness of the base material. In addition, when (Ca− (0.18 + 130 × Ca) × O) / (1.25 / S) ≧ 1, S is completely fixed by Ca, and MnS acting as a ferrite nucleation precipitates on CaS. Therefore, the composite sulfide cannot sufficiently function as a ferrite formation nucleus. On the other hand, when Ca, S, and O satisfy the above formula (1), MnS precipitates on CaS to form a composite sulfide, and can effectively function as a ferrite nuclei. . In addition, ((Ca− (0.18 + 130 × Ca) × O) / (1.25 / S) is preferably in the range of 0.2 to 0.8.

In addition to the above essential components, the high strength steel of the present invention can further contain one or more selected from B, V, Cu, Cr and Mo in order to increase the strength and toughness.
B: 0.0003-0.0025 mass%
B segregates at the austenite grain boundaries and has the effect of increasing the strength of the steel by suppressing the ferrite transformation that occurs from the grain boundaries and increasing the fraction of the bainite structure. The effect can be obtained by adding 0.0003 mass% or more. However, if added over 0.0025 mass%, the toughness is conversely reduced. When adding B, a more preferable range is 0.0005 to 0.002 mass%.

V: 0.2 mass% or less V is an element effective for improving the strength and toughness of the base material, and is also an element that precipitates as VN and also serves as a ferrite formation nucleus. However, if the addition amount exceeds 0.2 mass%, the toughness is deteriorated on the contrary. Therefore, it is preferable to add 0.2 mass% or less. More preferably, it is 0.15 mass% or less.

Cu: 1 mass% or less Cu is an element having the same effect as Ni. However, if it exceeds 1 mass%, it causes hot brittleness and deteriorates the surface properties of the steel sheet, so it is added in a range of 1 mass% or less. Is preferred. More preferably, it is 0.8 mass% or less.

Cr: 0.7 mass% or less Cr is an element effective for increasing the strength of the base material. However, if added in a large amount, it adversely affects toughness, so the upper limit should be 0.7 mass%. preferable. More preferably, it is 0.5 mass% or less.

Mo: 0.7 mass% or less Mo, like Cr, is an element effective for increasing the strength of the base material, but if added in a large amount, adversely affects toughness, so the upper limit is 0.7 mass%. Is preferable. More preferably, it is 0.5 mass% or less.

Next, the manufacturing method of the high strength steel of this invention is demonstrated.
The above-described molten steel adjusted to the component composition suitable for the present invention is melted by a normal method using a converter, an electric furnace, a vacuum melting furnace, etc., and a normal process such as continuous casting or ingot-bundling rolling is performed. After that, steel material such as slab is used. The steel material is hot-rolled to obtain a thick high-strength steel. At this time, the heating temperature of the steel material prior to the hot-rolling needs to be in the range of 1050 to 1200 ° C. The reason for heating to 1050 ° C. or higher is to ensure that the casting defects present in the steel material are crimped by hot rolling. However, when heated to a temperature exceeding 1200 ° C., TiN deposited during solidification becomes coarse and the toughness of the welded portion decreases, so the heating temperature needs to be regulated to 1200 ° C. or less.

  The steel material heated to the above temperature is then subjected to hot rolling in which the cumulative reduction ratio in the temperature range of 950 ° C. or higher is 30% or more, and heat in which the cumulative reduction ratio in the temperature range of less than 950 ° C. is 30 to 70%. Hot rolling is performed to obtain a high-strength steel having a predetermined thickness. The reason why hot rolling with a cumulative reduction ratio of 30% or more is performed in a temperature range of 950 ° C. or higher is that when a reduction with a cumulative reduction ratio of 30% or more is applied in this temperature range, the austenite grains recrystallize. This is because, when the cumulative rolling reduction is less than 30%, abnormal coarse particles generated during heating remain, which adversely affects the toughness of the base material.

  Also, the reason for performing hot rolling with a cumulative reduction ratio of 30 to 70% in a temperature range of less than 950 ° C. is that austenite grains rolled in this temperature range are not sufficiently recrystallized. It remains flat and has a high internal strain that is contained in a large amount in defects such as a deformation zone. This is because the accumulated internal energy acts as a driving force for the subsequent ferrite transformation and promotes the ferrite transformation. However, if the rolling reduction is less than 30%, the accumulated internal energy is not sufficient, so that ferrite transformation hardly occurs and a bainite structure is generated. On the other hand, if the rolling reduction exceeds 70%, the formation of polygonal ferrite is conversely accelerated and the formation of acicular ferrite is suppressed.

  Cooling after the end of the subsequent hot rolling is divided into pre-stage cooling and post-stage cooling, and the cooling rate of the former is made relatively larger than that of the latter, that is, in the pre-stage cooling, 600 to 450 ° C. from the hot rolling end temperature. Until the cooling stop temperature between, preferably from the hot rolling end temperature to the cooling stop temperature between 580 to 480 ° C. is cooled at a cooling rate of 5 to 20 ° C./s, preferably 6 to 16 ° C./s, and thereafter In the latter-stage cooling, from the first-stage cooling stop temperature to the second-stage cooling stop temperature between less than 450 ° C. and 200 ° C., preferably from the first-stage cooling stop temperature to the cooling stop temperature between 400 to 300 ° C., less than 1 to 5 ° C./s. It is necessary to cool at a cooling rate of preferably 2 to 4 ° C./s.

  When the stop temperature in the pre-cooling is higher than the above temperature range, there is almost no increase in strength. Further, if the cooling rate at the front stage is less than the lower limit of the above range, the structure is mainly composed of polygonal ferrite, and the strength cannot be improved. Conversely, if the upper limit of the above range is exceeded, the toughness is lowered. Furthermore, when the cooling stop temperature in the latter stage cooling is higher than the upper limit of the temperature range, the increase in strength is insufficient. Moreover, if the latter stage cooling rate is less than the lower limit of the above range, the base material strength is insufficient, and conversely if the upper limit of the above range is exceeded, the toughness of the base material is lowered.

  In the present invention, it is preferable to temper the steel material after cooling at a temperature of 450 to 650 ° C. for the purpose of reducing the residual internal stress. This is because if the tempering temperature is less than 450 ° C., the effect of removing residual stress is small, while if it exceeds 650 ° C., various carbonitrides precipitate to cause precipitation strengthening and lower toughness.

  As described above, in the method for producing high-strength steel according to the present invention, it is important to appropriately control the rolling reduction according to the rolling temperature in hot rolling and to properly control the two-stage cooling conditions after the end of rolling. In particular, when the cooling rate of the pre-stage cooling is made larger than that of the post-stage cooling, the base material becomes a structure mainly composed of acicular ferrite, and a steel material excellent in strength and toughness can be obtained.

  Steel slabs adjusted to various component compositions shown in Table 1 were used as raw materials, and thick steel plates having a thickness of 55 mm or 65 mm were manufactured under the manufacturing conditions shown in Tables 2-1 and 2-2. A sample was taken from each thick steel plate thus obtained and subjected to a tensile test and a Charpy test. In the tensile test, a JIS No. 4 tensile test piece was taken in the rolling width direction from the center of the thickness of each thick steel plate, and yield stress (YS) and tensile strength (TS) were determined. In the Charpy impact test, a JIS No. 4 impact test piece was taken in the rolling width direction from the center of the thickness of each steel plate, and the absorbed energy at -40 ° C (vE-40 ° C) was determined.

Furthermore, a groove (groove angle 30 °) is processed on a test plate taken from each steel plate, and a welded joint is produced by performing submerged arc welding with a heat input of 45 kJ / cm. A CTOD test piece in which the previous straight bond part was notched was collected and subjected to a CTOD test at −10 ° C. In addition, preparation of CTOD test pieces and test conditions were performed in accordance with British Standard BS7448.
Moreover, the JIS4 impact test piece which makes a notch position a bond part was extract | collected, the Charpy impact test was implemented at test temperature -40 degreeC, and the absorbed energy (vE-40 degreeC) was calculated | required.

  The test results are shown together in Table 2-1 and Table 2-2. From this result, the steel sheet of the example of the present invention has a yield stress (YS) of the base material of 355 MPa or more and a Charpy absorbed energy (vE-40 ° C.) of 200 J or more, and the strength and toughness of the base material are both. It can be seen that the submerged arc welded joint bond portion is excellent, and vE-40 ° C. is 200 J or more and the CTOD value is 0.50 mm or more, and the weld heat affected zone has excellent toughness. On the other hand, in the comparative example outside the scope of the present invention, only the steel sheet having any one or more of the above-described characteristics is obtained.

It is a graph which shows the influence which the pre-stage cooling rate after a hot rolling (cooling rate from the rolling completion temperature to the cooling stop temperature between 600-450 degreeC) has on the base material characteristic and the area rate of acicular ferrite.

Claims (5)

C: 0.05 to 0.1 mass%, Si: 0.05 to 0.5 mass%, Mn: 1 to 2 mass%, P: 0.015 mass% or less, S: 0.005 mass% or less, Al: 0.005 ~0.06mass%, Ni: 0.3 ~2mass% , Nb: 0.004~0.05mass%, Ti: 0.005~0.02mass%, N: less than 0.003mass%, Ca: 0.0005 Contains 0.003 mass%,
A high-strength steel characterized in that Ca, S and O satisfy the following formula (1) , and the balance has a composition composed of Fe and inevitable impurities .
0 <(Ca− (0.18 + 130 × Ca) × O) /1.25/S <1 (1)
Here, Ca, S and O are the contents of each element (mass%). In addition to the above component composition, B: 0.0003 to 0.0025 mass%, Mo: 0.7 mass% or less, V: 0.022 mass% or less, Cu: 1 mass% or less, and Cr: 0.7 mass% or less The high-tensile steel according to claim 1, comprising one or more selected from among them. C: 0.05 to 0.1 mass%, Si: 0.05 to 0.5 mass%, Mn: 1 to 2 mass%, P: 0.015 mass% or less, S: 0.005 mass% or less, Al: 0.005 ~0.06mass%, Ni: 0.3 ~2mass% , Nb: 0.004~0.05mass%, Ti: 0.005~0.02mass%, N: less than 0.003mass%, Ca: 0.0005 Contains 0.003 mass%,
Ca, S and O satisfy the following formula (1) and the steel slab having a composition composed of Fe and unavoidable impurities in the balance is heated to 1050 to 1200 ° C., and then cumulative pressure in a temperature range of 950 ° C. or higher. The hot rolling is performed so that the cumulative reduction in the temperature range of 30% or more and less than 950 ° C. is 30 to 70%, and thereafter, from the hot rolling end temperature to the cooling stop temperature of 600 to 450 ° C. High tension characterized by performing pre-stage cooling that cools at 20 ° C./s, and post-stage cooling that cools from the pre-stage cooling stop temperature to a cooling stop temperature between 450 ° C. and 200 ° C. at less than 1 to 5 ° C./s. Steel manufacturing method.
0 <(Ca− (0.18 + 130 × Ca) × O) /1.25/S <1 (1)
Here, Ca, S and O are the contents of each element (mass%). In addition to the above component composition , the steel slab further comprises B: 0.0003 to 0.0025 mass%, Mo: 0.7 mass% or less, V: 0.022 mass% or less, Cu: 1 mass% or less, and Cr: 0 The method for producing high-tensile steel according to claim 3, comprising one or more selected from 0.7 mass% or less. The method for producing high-tensile steel according to claim 3 or 4, wherein the steel after the latter stage cooling is subjected to tempering treatment at 450 to 650 ° C.

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