JP5439887B2 - High-strength steel and manufacturing method thereof - Google Patents

Description

  The present invention relates to a high-strength steel used for ships, offshore structures, line pipes, pressure vessels, and the like, and its manufacturing method. In particular, the yield stress (YS) is 460 MPa or more, and only the strength and toughness of the base material is excellent. In particular, the present invention relates to a high-strength steel excellent in toughness (CTOD characteristics) of a welded portion and a manufacturing method thereof.

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

  Conventionally, absorbed energy by Charpy impact test has been mainly used as an evaluation standard for toughness of steel. However, in recent years, a crack opening displacement test (hereinafter referred to as “CTOD test”) is often used in order to increase reliability. In this test, a test piece that had fatigue cracks in the toughness evaluation part was bent at three points, and the amount of opening (plastic deformation) at the crack bottom just before fracture was measured to determine the resistance to brittle fracture. It is something to evaluate.

  By the way, in general, multi-layer welding is performed on steel having a large thickness as used in the above-mentioned applications. However, in such welding, since the heat-affected zone receives a complicated heat history, there is a local embrittlement region. Easily generated, especially in bond areas (boundary between weld metal and base metal) and two-phase reheat zone (coarse grains in the first cycle of welding and heated to two phases of α and γ in the second cycle) ) Has a problem of large reduction in toughness. 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 Widmanstatten 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. Further, Patent Document 1 and Patent Document 2 describe a technique for suppressing the austenite grain growth and improving the toughness of the welded portion by adding a rare earth element (REM) together 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 further to improve toughness by controlling the form of sulfide by adding Ca and REM Technology has also been proposed.

  On the other hand, the above-mentioned two-phase region reheat zone, that is, the region exposed to the high temperature immediately below the melting point in the first welding, the region that becomes the two-phase region of ferrite and austenite by the reheating at the time of the subsequent welding, This is because carbon is concentrated in the austenite region due to reheating at the time of welding after the second pass, and this generates a fragile bainite structure including island martensite during cooling and lowers 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).

  Furthermore, in Patent Document 4, as a method of suppressing the formation of an embrittled structure due to reheating during the above-described welding, the amount of Ca added for controlling the form of sulfide is controlled within an appropriate range. A technique for improving the toughness characteristic (CTOD characteristic) of the weld heat affected zone by adding Ni is disclosed.

Japanese Patent Publication No. 03-053367 JP 60-184663 A JP 05-186823 A JP 2007-231312 A

  However, although the above-described problem that the toughness of the heat-affected zone is lowered has been improved to some extent by the above-described 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 on the contrary, 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 ships, offshore structures and the like increase in size, steel materials used for them are required to have higher strength and thickness. In order to meet these requirements, 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 formation of a brittle structure is promoted by reheating during welding, and the toughness of the weld heat affected zone is reduced. In addition, the technique disclosed in Patent Document 4 has a problem in that the raw material cost increases because it is essential to add Ni effective for increasing the toughness of the matrix in order to increase the strength and increase the thickness.

  Therefore, the object of the present invention is to solve the above-mentioned problems of the prior art, and even in a thick high-strength steel plate that has to increase the amount of addition of alloy elements, it is excellent in the strength and toughness of the base material, and is welded. The object is to propose a high-strength steel excellent in the toughness of the heat-affected zone and a suitable production method thereof.

  The inventors diligently studied a method capable of improving the base metal strength and toughness of the thick high-strength steel and also improving the toughness of the weld heat affected zone. As a result, the decrease in toughness of the weld heat affected zone is due to the formation of an embrittled structure, so in order to improve the toughness of this weld heat affected zone, the austenite grains in the region heated at high temperature during welding It has been found that it is effective to disperse the transformation nuclei uniformly and finely in order to suppress the coarsening and to further promote the 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 Mn 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 has been found that if controlled, the steel sheet structure becomes a structure mainly composed of acicular ferrite, and a high-strength steel excellent in the strength and toughness of the base material can be produced. Furthermore, it has been found that 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 an unrecrystallized region in the low temperature region of austenite. The present invention has been completed only by properly combining these techniques.

That is, the present invention includes C: 0.03 to 0.10 mass%, Si: 0.30 mass% or less, Mn: 1.66 to 2.30 mass%, P: 0.012 mass% or less, S: 0.005 mass. %: Al: 0.005-0.06 mass%, Nb: 0.004-0.05 mass%, Ti: 0.005-0.02 mass%, N: 0.001-0.005 mass%, Ca: 0 .0005 to 0.003 mass%, and 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-strength steel of the present invention further includes B: 0.0003 to 0.0025 mass%, V: 0.2 mass% or less, Cu: 1 mass% or less, Ni: 0.75 mass% or less, One or more selected from Cr: 0.7 mass% or less and Mo: 0.7 mass% or less are contained.

In the present invention, C: 0.03 to 0.10 mass%, Si: 0.30 mass% or less, Mn: 1.66 to 2.30 mass%, P: 0.012 mass% or less, S: 0.005 mass %: Al: 0.005-0.06 mass%, Nb: 0.004-0.05 mass%, Ti: 0.005-0.02 mass%, N: 0.001-0.005 mass%, Ca: 0 .0005 to 0.003 mass%, and 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 then the hot rolling is cooled from 5 to 45 ° C./sec from the hot rolling end temperature to the cooling stop temperature between 600 to 450 ° C., The present invention proposes a method for producing a high-strength steel, characterized in that post-stage cooling is performed in which cooling is performed at a temperature of 1 ° C./sec or more and less than 5 ° C./sec from the preceding stage cooling stop temperature to a cooling stop temperature of 450 ° C. or less.

In addition to the above component composition, the steel slab used in the method for producing high-tensile steel according to the present invention further includes B: 0.0003 to 0.0025 mass%, V: 0.2 mass% or less, Cu: 1 mass% or less, Ni: It is characterized by containing one or more selected from 0.75 mass% or less, Cr: 0.7 mass% or less, and Mo: 0.7 mass% or less.

  In addition, the production method of the present invention is characterized in that a tempering treatment at 450 to 650 ° C. is performed on the steel after the subsequent stage cooling.

  According to the present invention, a high-strength steel whose base material has a high strength with a yield stress of 460 MPa or more, is excellent in toughness, and is excellent in the toughness (CTOD characteristics) of a heat-affected zone after welding is manufactured at low cost. Can greatly contribute to the increase in size of ships and offshore structures.

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 a base material characteristic.

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 morphology 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. Then, 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, so that MnS is precipitated on the surface of CaS and combined. Forms sulfides. 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. The effect of this Mn dilute band is more effectively expressed by increasing the amount of Mn added in the steel. Moreover, since ferrite-forming nuclei such as TiN, BN, and AlN also precipitate on the precipitated MnS, the ferrite transformation is further promoted.
Further, by increasing the amount of Mn added, the base metal strength can be effectively increased without generating island martensite, which is an embrittled structure, in the weld heat affected zone as much as possible. This is because the island-like martensite generated during cooling after welding is easily decomposed into cementite due to the increase in the amount of Mn added, and the island-like martensite in the heat-affected zone structure is reduced. As a result of these effects, the toughness of the weld heat affected zone can be ensured without making the addition of Ni essential.

  The above technology makes it possible to finely disperse ferrite transformation nuclei that do not melt even at high temperatures, refine the structure of the weld heat-affected zone, and minimize the formation of island martensite to obtain high toughness. be able to. Even in the region reheated to the two-phase region by the thermal cycle during multi-layer welding, the structure of the first weld heat affected zone is refined, so that the toughness of the untransformed region is improved and the retransformation is further improved. Since the austenite grains to be refined are also made fine, the degree of toughness reduction can be kept small.

The second feature of the present invention lies in that the cooling after rolling the steel material is divided into two stages, a pre-stage cooling and a post-stage cooling, and the cooling rate of the pre-stage cooling is controlled more greatly than the post-stage cooling. This point will be described based on experimental results.
After heating a steel slab containing C: 0.08 mass%, Si: 0.2 mass%, and Mn: 1.8 mass% to 1150 ° C, the cumulative reduction ratio of 950 ° C or higher is 40%, less than 950 ° C. After hot rolling with a cumulative rolling reduction of 50% and a rolling end temperature of 850 ° C., cooling from the rolling end temperature to 500 ° C. is performed at a cooling rate of 5 to 45 ° C./sec, more preferably 5 to 20 ° C./sec. Pre-stage cooling and further post-stage cooling to 350 ° C. at a cooling rate of 3 ° C./sec were performed, followed by air cooling to obtain a thick steel plate having a thickness of 10 to 50 mm. About this thick steel plate, the tensile strength characteristic and the toughness characteristic (Charpy impact absorption energy) in -40 degreeC were measured.

  FIG. 1 shows the influence of the preceding cooling rate on the base material strength and toughness with respect to the above measurement results. The cooling rate of the former cooling from the rolling end temperature to 500 ° C. is 5 to 45 ° C./sec. It can be seen that by controlling to the range, a steel sheet having a high strength with a yield stress of 460 MPa or more and an excellent strength-toughness balance with a vE-40 ° C. of 200 J or more can be obtained.

Furthermore, it has been found that the steel sheet cooled at the cooling rate has a structure mainly composed of acicular ferrite. Generally, when trying to obtain high-strength steel, the toughness is greatly reduced when a relatively coarse upper bainite structure including island-like martensite and the like between the laths is obtained. 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. However, the inventors divided the cooling after rolling into pre-stage cooling and post-stage cooling with a slower cooling rate than that, and by appropriately controlling each cooling rate, the structure was mainly composed of acicular ferrite, and was excellent. It has been found that a steel sheet having a strength-toughness balance can be obtained. This is because by increasing the cooling rate in the previous stage, the transformation nucleation density can be increased, and the transformed structure can be made into a dense acicular ferrite structure instead of a coarse bainite structure. Furthermore, with regard to the cooling rate of the subsequent stage, if it is too faster than the cooling rate of the previous stage, island-shaped martensite is generated and the toughness of the base material is deteriorated. It has also been found that it is necessary to control within an appropriate range since the strength decreases.
The present invention has been completed based on the above findings.

Next, the component composition that the high-strength steel according to the present invention should have will be described.
C: 0.03-0.10 mass%
C is an element that has the greatest influence on the strength of the steel, and in order to ensure the strength necessary for structural steel (YS ≧ 460 MPa), it is necessary to contain 0.03 mass% or more. However, conversely, if too much, lowering of the toughness of the base metal and cold cracking during welding are caused, so the upper limit is made 0.10 mass%.

Si: 0.30 mass% or less Si is a component to be added as a deoxidizer and to increase the strength of steel, and in order to obtain the effect, it is preferable to add 0.01 mass% or more. . However, if it exceeds 0.30 mass%, the toughness of the base metal and the welded portion is lowered, so that it is necessary to be 0.30 mass% or less. Preferably, it is the range of 0.01-0.20 mass%.

Mn: 1.60 to 2.30 mass%
Mn is an element effective for securing the strength of the base material, but in the present invention, the refinement of the structure of the heat affected zone of the weld is promoted and the formation of the brittle structure is suppressed as much as possible, thereby affecting the heat of welding. It is an important element to be added in order to improve the toughness of the part (CTOD characteristics). In order to acquire this effect, it is necessary to add 1.60 mass% or more. On the other hand, if it exceeds 2.30 mass%, the toughness of the base metal and the welded portion is remarkably lowered, so that it is 2.30 mass% or less. Preferably, it is the range of 1.65 to 2.15 mass%.

P: 0.015 mass% or less P is an impurity that is inevitably mixed, and if it exceeds 0.015 mass%, the toughness of the base metal and the welded portion is lowered, so that it is limited to 0.015 mass% or less. Preferably, it is 0.010 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 the toughness. Therefore, it is necessary to limit to 0.06 mass% or less. Preferably, it is 0.010-0.055 mass%.

Nb: 0.004 to 0.05 mass%
Since Nb forms a non-recrystallized region in the low temperature range of austenite, the microstructure of the base material can be refined and the toughness can be increased 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, it is necessary to add Nb 0.004 mass% or more. However, if it is added excessively exceeding 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 solidifies, suppresses coarsening of austenite in the welded portion, and becomes a transformation nucleus of ferrite, thereby contributing to the increase in toughness of the welded portion. In order to obtain these effects, it is necessary to add 0.005 mass% or more. On the other hand, if added over 0.02 mass%, the TiN particles become coarse, and the effect of improving the toughness of the base metal and the welded portion cannot be obtained. Therefore, the amount of Ti added is in the range of 0.005 to 0.02 mass%.

N: 0.001 to 0.005 mass%
N is an element necessary for forming TiN that suppresses the coarsening of the structure of the welded portion, and is added in an amount of 0.001 mass% or more. On the other hand, if added over 0.005 mass%, the solid solution N significantly reduces the toughness of the base metal and the welded portion, so the upper limit is made 0.005 mass%. In order to form a sufficient amount of TiN for pinning that suppresses the coarsening of the structure, a range of 0.003 to 0.005 mass% is preferable.

Ca: 0.0005 to 0.003 mass%
Ca is an element that improves toughness by fixing S. In order to develop this effect, it is necessary to add at least 0.0005 mass%. However, since the effect is saturated even if it contains exceeding 0.003 mass%, Ca is added in 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 From the value, the form of sulfide can be estimated (Mochida et al., “Iron and Steel”, Japan Iron and Steel Institute, 66th (1980), No. 3, pages 354 to 362).

That is, when the value of ((Ca− (0.18 + 130 × Ca) × O) /1.25/S) is 0 or less, CaS does not crystallize. Therefore, since S precipitates in the form of MnS alone, it is impossible to realize fine dispersion of ferrite-forming nuclei in the weld heat affected zone, which is the main point of the present invention. Further, MnS precipitated alone is elongated during rolling of the steel sheet, causing a reduction in the toughness of the base material.
On the other hand, when the value of ((Ca− (0.18 + 130 × Ca) × O) /1.25/S) is 1 or more, S is completely fixed by Ca, and MnS acting as ferrite nuclei is CaS. Since the composite sulfide does not precipitate on the surface, the composite sulfide cannot sufficiently function as the ferrite nuclei.
On the other hand, when Ca, S, and O satisfy the above formula (1), MnS precipitates on CaS to form a composite sulfide, which effectively functions as a ferrite nuclei. it can. The value of ((Ca− (0.18 + 130 × Ca) × O) /1.25/S) is preferably in the range of 0.2 to 0.8.

The high-tensile steel of the present invention contains one or more selected from B, V, Cu, Ni, Cr and Mo in order to further increase the strength and toughness in addition to the above essential components. Can do.
B: 0.0003 to 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 is acquired by addition of 0.0003 mass% or more. However, if added over 0.0025 mass%, the toughness is conversely reduced. A more preferable range of B 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. In order to acquire the effect, it is preferable to add 0.01 mass% or more. However, if the added amount exceeds 0.2 mass%, the toughness is reduced instead. 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 an effect of improving the strength of steel. In order to acquire the effect, it is preferable to add 0.05 mass% or more. However, if it exceeds 1 mass%, it causes hot brittleness and deteriorates the surface properties of the steel sheet. Therefore, it is preferably added in the range of 1 mass% or less. More preferably, it is 0.8 mass% or less.

Ni: 2 mass% or less Ni is an element effective for improving the strength of steel and the CTOD characteristics of the weld heat affected zone. In order to acquire the effect, it is preferable to add 0.05 mass% or more. However, since Ni is an expensive element, the upper limit is preferably set to 2.0 mass%. When Mn is added in an amount of 1.6 mass% or more as in the present invention, Ni is more preferably less than 0.3 mass% from the viewpoint of reducing raw material costs.

Cr: 0.7 mass% or less Cr is an element effective for increasing the strength of the base material. In order to acquire the effect, it is preferable to add 0.05 mass% or more. However, if added in a large amount, the toughness is adversely affected, so the upper limit is preferably 0.7 mass%. 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. In order to acquire the effect, it is preferable to add 0.05 mass% or more. However, if added in a large amount, the toughness is adversely affected, so the upper limit is preferably 0.7 mass%. More preferably, it is 0.5 mass% or less.

Next, the manufacturing method of the high strength steel of this invention is demonstrated.
The high-strength steel of the present invention is prepared by melting a molten steel adjusted to the above-described component composition by a usual method using a converter, an electric furnace, a vacuum melting furnace, etc. It is preferable that a steel material such as a slab is obtained through a normal process such as lump-slab rolling, and then this steel material is hot-rolled to produce a thick high-strength steel. Under the present circumstances, it is necessary to make the heating temperature of the steel raw material performed prior to hot rolling into the range of 1050-1200 degreeC. The reason why the heating temperature is set to 1050 ° C. or higher is to reliably press-bond casting defects existing in the steel material by hot rolling. However, when heated to a temperature exceeding 1200 ° C., TiN deposited during solidification becomes coarse and the toughness of the base material and the welded portion decreases, so the heating temperature needs to be regulated to 1200 ° C. or lower.

  The steel material heated to the above temperature is then subjected to hot rolling in which the cumulative reduction rate in the temperature range of 950 ° C. or higher is 30% or more, and the cumulative reduction rate in the temperature range of less than 950 ° C. is 30 to 70%. A high-strength steel having a predetermined thickness is used. The reason for performing hot rolling with a cumulative reduction ratio of 30% or more in a temperature range of 950 ° C. or higher is that the austenite grains are recrystallized and the structure becomes fine by setting the cumulative reduction ratio in this temperature range to 30% or more. However, if the cumulative rolling reduction is less than 30%, abnormal coarse grains generated during heating remain, which adversely affects the toughness of the base material.

  In addition, the reason for performing hot rolling to set the cumulative reduction ratio in the temperature range below 950 ° C. to 30 to 70% is that the austenite grains rolled in this temperature range are not sufficiently recrystallized. It remains deformed flat and has a high internal strain that is contained in a large amount in defects such as deformation bands. The accumulated internal energy works 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 the base material toughness deteriorates. On the other hand, when the rolling reduction exceeds 70%, the formation of polygonal ferrite is promoted, the formation of acicular ferrite is suppressed, and high strength and high toughness are not compatible.

  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. Between 5 to 45 ° C / sec, preferably 5 to 20 ° C / sec, more preferably 6 to 16 ° C. Cooling at a cooling rate of / sec, and in the subsequent rear cooling, from the previous cooling stop temperature to the lower cooling stop temperature of 450 ° C. or less, preferably from the previous cooling stop temperature to the cooling stop temperature between 400 to 250 ° C. It is necessary to cool at a cooling rate of 1 ° C./sec or more and less than 5 ° C./sec, preferably 2 to 4.5 ° C./sec.

  When the stop temperature in the pre-cooling is higher than the above temperature range, there is almost no increase in strength, and conversely, when it is lower than the above temperature range, the toughness deteriorates. 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 decreases. 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. On the other hand, if the subsequent cooling rate is too higher than the preceding cooling rate, island martensite is generated and the toughness of the base material is deteriorated.

  In the present invention, the steel material after cooling may be tempered in a temperature range of 450 to 650 ° C. for the purpose of reducing residual internal stress. When the tempering temperature is less than 450 ° C., the effect of removing residual stress is small. On the other hand, when the temperature exceeds 650 ° C., various carbonitrides precipitate to cause precipitation strengthening and reduce 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-toughness balance can be obtained.

  Further, in the present invention, N of the steel components is more than 0.0030 mass%, the cooling rate of the pre-stage cooling after hot rolling is more than 20 ° C./sec. 45 ° C./sec., And the stop temperature of the pre-stage cooling is 450 ° C. or more. By making the temperature lower than 500 ° C., high-strength steel with high yield strength of the base metal of 550 MPa or more, excellent toughness, and excellent heat-affected zone toughness (CTOD characteristics) after welding is inexpensive. Can be manufactured.

  No. having the component composition shown in Table 1-1 and Table 1-2. Using steel slabs 1 to 31 as raw materials, hot rolling, pre-cooling and post-cooling were performed under the conditions shown in Table 2-1 and Table 2-2 to produce a thick steel plate having a thickness of 50 to 80 mm. In addition, the temperature described in Table 2-1 and Table 2-2 is the temperature of 1/4 part thickness calculated by calculating from the steel plate surface temperature measured with the radiation thermometer. A sample was taken from the thick steel plate thus obtained and subjected to a tensile test and a Charpy impact test. In the tensile test, a JIS No. 4 tensile test piece was sampled from 1/4 thickness of a thick steel plate so that the direction of the longitudinal axis of the test piece was parallel to the rolling direction, and yield stress (YS), tensile strength ( TS) was measured. In the Charpy impact test, JIS No. 4 impact test specimens were collected in the rolling width direction from 1/4 thickness of each thick steel plate, and the absorbed energy (vE-40 ° C) at a temperature of -40 ° C was measured. And the thing which satisfy | fills all of YS> = 460MPa, TS> = 570MPa, and vE-40 degreeC> = 200J was evaluated that the base material characteristic was favorable.

  Furthermore, in principle, the groove (30 ° groove angle) is processed into a test plate taken from a thick steel plate in which all of the base material characteristics YS, TS and vE-40 ° C. satisfy the above criteria, and the heat input is A carbon dioxide arc welding of 25 kJ / cm is performed to produce a welded joint, and a CTOD test piece in which a straight bond portion of a groove is notched is taken from this welded joint, and a CTOD test is performed at a temperature of −10 ° C. It was. In addition, preparation of CTOD test pieces and test conditions were performed in accordance with British Standard BS7448. Moreover, the JIS No. 4 impact test piece which makes a notch position a bond part was extract | collected, the Charpy impact test was done at the temperature of -40 degreeC, and the absorbed energy (vE-40 degreeC) was measured.

  The test results are shown together in Table 2-1 and Table 2-2. From these results, the steel sheet of the example of the present invention has a yield stress (YS) of the base material of 460 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 high. It is understood that both are excellent, and the carbon dioxide arc welded joint has a vE-40 ° C. of 200 J or more, a CTOD value of 0.10 mm or more, and excellent toughness of the weld heat affected zone. . On the other hand, in the steel of the comparative example which is out of the scope of the present invention, only a steel plate having any one or more of the above characteristics is obtained.

In addition, steel plate No. with N content exceeding 0.0030 mass%. In 11-17, the CTOD characteristic of the weld is excellent due to the pinning effect of TiN.
Also, the N content is over 0.0030 mass%, the cooling rate of the pre-cooling after hot rolling is 20 ° C./sec to 45 ° C./sec or less, and the pre-cooling stop temperature is 450 ° C. or more and less than 500 ° C. Steel plate No. manufactured in 15 and 16 both have high strength in which the yield stress of the base material is 550 MPa or more.

  The high-tensile steel of the present invention can be suitably used not only for ships and offshore structures, line pipes, pressure vessels, but also for steel structures that are assembled by welding in the fields of construction and civil engineering.

Claims (6)

C: 0.03-0.10 mass%,
Si: 0.30 mass% or less,
Mn: 1.66 ~2.30mass%,
P: 0.012 mass% or less,
S: 0.005 mass% or less,
Al: 0.005 to 0.06 mass%,
Nb: 0.004 to 0.05 mass%,
Ti: 0.005-0.02 mass%,
N: 0.001 to 0.005 mass%,
Ca: 0.0005 to 0.003 mass%,
And high strength steel characterized by Ca, S, and O satisfying | filling following (1) Formula, and having the component composition which remainder consists of Fe and an unavoidable impurity.
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%, V: 0.2 mass% or less, Cu: 1 mass% or less, Ni: 0.75 mass% or less, Cr: 0.7 mass% or less, and The high-tensile steel according to claim 1, comprising one or more selected from Mo: 0.7 mass% or less. C: 0.03-0.10 mass%,
Si: 0.30 mass% or less,
Mn: 1.66 ~2.30mass%,
P: 0.012 mass% or less,
S: 0.005 mass% or less,
Al: 0.005 to 0.06 mass%,
Nb: 0.004 to 0.05 mass%,
Ti: 0.005-0.02 mass%,
N: 0.001 to 0.005 mass%,
Ca: 0.0005 to 0.003 mass%,
And after heating the steel slab which has the component composition which Ca, S, and O satisfy | fill following (1) Formula, and remainder consists of Fe and an unavoidable impurity to 1050-1200 degreeC, in the temperature range of 950 degreeC or more Hot rolling is performed so that the cumulative rolling reduction is 30% or more and the cumulative rolling reduction is 30 to 70% in a temperature range of less than 950 ° C., and then the hot rolling end temperature to the cooling stop temperature between 600 to 450 ° C. Pre-stage cooling is performed at 5 to 45 ° C./sec, and post-stage cooling is performed from 1 ° C./sec to less than 5 ° C./sec from the pre-stage cooling stop temperature to a cooling stop temperature of 450 ° C. or lower. Manufacturing method of high-strength steel.
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%, V: 0.2 mass% or less, Cu: 1 mass% or less, Ni: 0.75 mass% or less, Cr: 0.7 mass% or less, and The method for producing a high-strength steel according to claim 3, comprising one or more selected from Mo: 0.7 mass% or less. The high-strength steel manufacturing method according to claim 3 or 4, wherein the pre-cooling is performed at 5 to 20 ° C / sec. The method for producing a high-strength steel according to any one of claims 3 to 5, wherein the steel after post-stage cooling is subjected to a tempering treatment at 450 to 650 ° C.

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