CN102124133A - High-tensile strength steel and manufacturing method thereof - Google Patents

Abstract

The present invention provides a high-strength steel and an advantageous manufacturing method thereof. Even when the high-strength steel is a thick-walled high-strength steel plate, the strength and toughness of the base material are excellent, and the toughness of the welded hot spot is also excellent. Specifically, the high-strength steel has the following composition: by mass %, C: 0.03 to 0.10%, Si: 0.30% or less, Mn: 1.60 to 2.30%, P: 0.015% or less, S: 0.005% or less, Al: 0.005 to 0.06%, Nb: 0.004 to 0.05%, Ti: 0.005 to 0.02%, N: 0.001 to 0.005%, Ca: 0.0005 to 0.003%, and the content of Ca, S and O satisfies the following formula (1) , and the balance consists of Fe and unavoidable impurities. 0<(Ca-(0.18+130×Ca)×O)/1.25/S<1 (1) wherein, Ca, S and O represent the contents (% by mass) of the respective elements.

Description

High-strength steel and its manufacturing method

technical field

The present invention relates to high-tensile strength steels used in ships or marine structures, line pipes, pressure vessels, etc. The stress (YS (yield stress)) is 460MPa or more, and not only the strength (strength) and toughness (toughness) of the base material (base material) are excellent, but also the toughness (CTOD (crack tip opening displacement) characteristic) of the weld zone (weld zone) is also good. Excellent high-strength steel and method for its manufacture.

Background technique

Steel used for ships, offshore structures, and the like is generally processed into desired-shaped structures and the like by welding joints. For this reason, in these steels, from the viewpoint of ensuring the safety of structures and the like, it is essential that the base metal itself has excellent strength and toughness, and it is also required that the welded part of the weld joint (weld joint) (weld metal) and heat-affected zone (heat-affected zone) are also excellent in toughness.

As evaluation standards for toughness of steel, absorbed energy measured by Charpy impact test (Charpy impact test) has been mainly used in the past. However, in recent years, in order to further improve the reliability (reliability), the crack tip opening displacement test (Crack Tip Opening Displacement Test, hereinafter abbreviated as "CTOD test") is often used. This test measures the value of opening displacement of the crack bottom (bottom of crack) before fracture by subjecting a test piece having a fatigue precrack to three-point bending in the toughness evaluation portion. (value of plastic deformation) to evaluate the occurrence resistance of brittle fracture.

However, in general, multi-pass welding can be performed on thick steel used in the above-mentioned applications. In such welding, since the heat-affected zone is subjected to a complicated thermal history, it is easy to Produce local embrittlement zones, especially the fusion line (bond) (the boundary between the weld metal and the base metal) and the inter-critically reheated zone (inter-critically reheated zone) in the duplex zone (becoming coarse in the first cycle of welding, and The problem that the toughness of the area heated into the two-phase area of α and γ in the two cycles) is greatly reduced. This is because the fusion line portion is exposed to a high temperature close to the melting point, thereby coarsening austenite grains, and then easily transforming into a fragile upper bainitic structure by cooling. In addition, since an embrittlement structure (embrittlement structure) such as Widmannstatten structure and island martensite (M-A constituent) is formed at the weld line portion, the toughness is further lowered.

As a countermeasure against the above-mentioned problems, for example, a technique of suppressing the coarsening of austenite grains by finely dispersing TiN in steel or a technique of utilizing a nucleus of ferrite transformation has been put into practical use. In addition, the following technology is disclosed in Japanese Patent Publication No. 03-053367 or Japanese Patent Application Publication No. 60-184663: adding rare earth elements (REM (rare-earth metal) together with Ti and dispersing fine particles in steel Among them, the growth of austenite grains can be suppressed and the toughness of welded joints can be improved. In addition, the technology of dispersing Ti oxides and the combination of ferrite formation ability of BN and oxide dispersion have been proposed. technology, and the technology of adding Ca or REM to control the shape of sulfide (sulfide) to improve toughness.

On the other hand, the above-mentioned two-phase region reheating part, that is, the region exposed to high temperature close to the melting point in the initial welding, also becomes a two-phase ferrite and austenite by reheating in the subsequent welding. The reason for the brittleness of the region is that carbon is enriched in the austenite region by reheating during welding after the second pass, which generates brittle bainite containing island martensite during cooling. body tissue, reducing toughness. Therefore, as a countermeasure, a technique has been disclosed in which the 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, refer to Japanese Patent Application Laid-Open No. 05-186823). .

In addition, as a method of suppressing the formation of brittle structures by reheating during welding, Japanese Patent Application Laid-Open No. 2007-231312 discloses a technique in which the amount of Ca added to control the form of sulfide is appropriately controlled. On this basis, the toughness (CTOD characteristics) of the welded heat-affected zone can be improved by adding Ni.

However, although the above-mentioned problem of lowering the toughness of the heat-affected zone has been improved to some extent compared with the above-mentioned prior art, there are still some problems to be solved. For example, in the technology using TiN, this effect disappears in the fusion line portion heated to the temperature range where TiN melts, and more than that, due to the embrittlement of the matrix structure caused by solid-solution Ti and solid-solution N, toughness may sometimes occur significantly lowered. In addition, in the technology using oxides of Ti, there is a problem that the oxides cannot be dispersed finely and uniformly enough. Furthermore, along with the increase in size of ships, offshore structures, etc. in recent years, the steel materials used for them are required to be further increased in strength and thickness. In order to satisfy these requirements, it is effective to add a large amount of alloy elements contrary to the technique of JP-A-05-186823. However, the addition of a large amount of alloying elements has the problem of promoting the formation of an embrittlement structure caused by reheating during welding, resulting in a decrease in the toughness of the welded heat-affected zone. In addition, in the technology disclosed in Japanese Patent Application Laid-Open No. 2007-231312, as a countermeasure for high strength and thick wall, it is necessary to add Ni (effect of solid solution Ni) effective for high toughness of the matrix, which is costly. Getting expensive becomes the problem.

Therefore, in order to solve the above-mentioned problems of the prior art, the object of the present invention is to provide a high-strength steel and a preferred manufacturing method thereof, which can be used even for thick-walled steels that have to increase the addition of alloy elements. In the case of a high-strength steel plate, the strength and toughness of the base metal are also excellent, and the toughness of the welded heat-affected zone is also excellent.

Contents of the invention

The present invention is a high-strength steel characterized by having the following composition: C: 0.03-0.10% by mass, Si: 0.30% by mass or less, Mn: 1.60-2.30% by mass, P: 0.015% by 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 to 0.02 mass %, N: 0.001 to 0.005 mass %, Ca: 0.0005 to 0.003 mass %, and Ca, The content of S and O satisfies the following formula (1), and the balance is composed of Fe and unavoidable impurities,

0＜(Ca-(0.18+130×Ca)×O)/1.25/S＜1 …(1)

Here, Ca, S, and O represent the content (mass %) of each element.

In addition, the high-strength steel of the present invention is characterized in that, in addition to the composition of the above components, it further contains: B: 0.0003 to 0.0025 mass%, V: 0.2 mass% or less, Cu: 1 mass% or less, Ni: 2 One or more of Cr: 0.7 mass % or less and Mo: 0.7 mass % or less.

In addition, the present invention also proposes a method for producing high-strength steel, which is characterized in that after heating the steel billet to 1050-1200°C, the cumulative rolling rate applied in the temperature range of 950°C or higher is 30% or more. Hot rolling with a cumulative rolling ratio of 30 to 70% in a temperature range below 950°C, and then performing front cooling and rear cooling, the front cooling is cooled from the hot rolling end temperature at 5 to 45 °C/sec to 600 to The cooling stop temperature between 450°C, the cooling stop temperature in the subsequent stage is cooled from the cooling stop temperature in the previous stage at a rate of more than 1°C/s and less than 5°C/s to a cooling stop temperature below 450°C, wherein the billet has the following composition: Contains C: 0.03-0.10 mass%, Si: 0.30 mass% or less, Mn: 1.60-2.30 mass%, P: 0.015 mass% or less, S: 0.005 mass% or less, Al: 0.005-0.06 mass%, Nb: 0.004- 0.05% by mass, Ti: 0.005 to 0.02% by mass, N: 0.001 to 0.005% by mass, Ca: 0.0005 to 0.003% by mass, and the content of Ca, S and O satisfies the following formula (1), and the balance is composed of Fe and The inevitable impurity composition,

0＜(Ca-(0.18+130×Ca)×O)/1.25/S＜1 …(1)

Here, Ca, S, and O represent the content (mass %) of each element.

In addition, the production method of the present invention is characterized in that the high-strength steel contains, in addition to the above-mentioned composition, components selected from the group consisting of B: 0.0003 to 0.0025% by mass, V: 0.2% by mass or less, and Cu: 1% by mass or less. , Ni: 2% by mass or less, Cr: 0.7% by mass or less, and Mo: 0.7% by mass or less, or two or more.

Furthermore, the production method of the present invention is characterized in that the tempering treatment at 450 to 650° C. is performed on the steel cooled in the subsequent stage.

According to the present invention, a high-strength steel having a high strength and excellent toughness with a yield stress of 460 MPa or more of the base material and excellent toughness (CTOD characteristics) of the heat-affected zone after welding can be produced cheaply. Therefore, it is suitable for ships and offshore structures. The upscaling of etc. contributes a lot.

Description of drawings

Fig. 1 is a graph showing the effect of the front cooling rate (the cooling rate from the rolling end temperature to the cooling stop temperature between 600°C and 450°C) after hot rolling on the properties of the base material.

Detailed ways

The inventors conducted intensive research on a method capable of increasing the base material strength and toughness of a thick-walled high-strength steel and improving the toughness of a welded heat-affected zone. As a result, it was found that the reduction of the toughness of the welded heat-affected zone leads to the formation of an embrittlement structure, and therefore, in order to improve the toughness of the welded heat-affected zone, the coarsening of austenite grains in the region heated at a high temperature during welding has been suppressed. Fundamentally, it is also effective to uniformly and finely disperse transformation nuclei in order to promote ferrite transformation during cooling after welding.

Therefore, the inventors of the present invention further studied a method for suppressing the formation of the above-mentioned brittle structure, and found that it is effective to control the amount of Ca added to control the form of sulfide in an appropriate range, and in order to increase the welding heat It is effective to add Mn to affect the toughness (CTOD characteristic) of the part.

In addition, the effect of rolling conditions on the strength and toughness of the base metal was studied. As a result, it was found that if the cooling after rolling is set to consist of front-stage cooling with a high cooling rate and post-stage cooling with a low cooling rate The two-stage cooling and appropriate control of the respective cooling rates will change the structure of the steel plate into acicular ferrite as the main structure, and high-strength steel with excellent strength and toughness of the base metal 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 nonrecrystallization zone (nonrecrystallization zone) in the low temperature range of austenite. Also, by appropriately combining these technologies, the present invention has been accomplished for the first time.

The basic technical idea of the present invention will be described below.

The first feature of the present invention is to effectively utilize the crystallization of the Ca compound (CaS) added for the purpose of shape control of sulfide in order to improve the toughness of the welded heat affected zone. (crystallization). Since CaS crystallizes at a lower temperature than oxides, it can be finely dispersed uniformly. Furthermore, by controlling the amount of CaS added and the dissolved oxygen amount in molten steel (dissolved oxygen amount) at the time of addition (dissolved oxygen amount) in an appropriate range, solid solution S can be ensured even after the crystallization of CaS, so the surface of CaS MnS is precipitated to form complex sulfides. It is known that this MnS has the potential for ferrite nucleus, and further, a Mn depleted zone (Mn depleted zone) can be formed around the precipitated MnS, so it can further promote the ferrite transformation (ferrite transformation). ). By increasing the amount of Mn added to the steel, the effect of the Mn-depleted zone can be more effectively exhibited. In addition, since ferrite nuclei such as TiN, BN, and AlN are deposited on the precipitated MnS, the ferrite transformation can be further promoted.

In addition, by increasing the amount of Mn added, it is possible to effectively increase the strength of the base metal without generating a large amount of insular martensite as an embrittlement structure in the weld heat-affected zone. This is because the insular martensite generated in the cooling process after welding is easily decomposed into cementite by increasing the amount of Mn added, thereby reducing the insular martensite in the heat-affected zone structure. As a result of these effects, the toughness of the welded heat-affected zone can be secured without adding Ni.

By the above-mentioned technology, it is possible to finely disperse ferrite transformation nuclei that do not melt at high temperatures, and to refine the microstructure of the welded heat-affected zone, and by suppressing island martensite (island martensite, The generation of M-A constituent) can obtain high toughness. In addition, due to the heat cycle (heat cycle) during multilayer welding, even in the region that is reheated into the dual-phase region, the structure of the heat-affected zone of the initial welding is refined and the phase transformation is not achieved. The toughness of the region is improved, and the re-transformed austenite grains are further refined, so that the decrease in toughness can be suppressed to a small degree.

The second feature of the present invention is that the cooling after rolling of the steel is divided into two stages, the front cooling and the rear cooling, and the cooling rate of the front cooling is controlled to be higher than that of the rear cooling. Hereinafter, this feature will be described based on experimental results.

After heating a steel slab whose basic components are C: 0.08% by mass, Si: 0.2% by mass, and Mn: 1.8% by mass to 1150°C, the cumulative rolling ratio above 950°C is 40%, and the cumulative rolling ratio below 950°C After hot rolling at a rate of 50% and a rolling finish temperature of 850°C, cooling from the rolling finish temperature to 500°C is performed at a cooling rate of 5 to 45°C/sec, more preferably 5 to 20°C/sec The first stage of cooling and the subsequent stage of cooling to 350°C at a cooling rate of 3°C/s, and then air cooling to produce a thick steel plate with a plate thickness of 10-50mm. For this thick steel plate, the tensile strength characteristics and the toughness characteristics (Charpy impact absorbed energy) at -40°C were measured.

Figure 1 is a graph showing the influence of the front-stage cooling rate on the strength and toughness of the base metal based on the above-mentioned measurement results. In the range of 45° C./sec, a high-strength steel sheet having a yield stress of 460 MPa or more and a vE-40° C. of 200 J or more can be obtained with an excellent strength-toughness balance.

In addition, it is also known that the steel plate cooled at the above-mentioned cooling rate has a structure mainly of acicular ferrite. In general, when high-strength steel is to be obtained, toughness is greatly reduced when a relatively coarse upper bainite structure containing island martensite and the like is formed between laths. Therefore, in order to achieve both high strength and high toughness, it is necessary to obtain a fine acicular ferrite structure by considering rolling conditions and the like. However, the inventors found that by dividing the cooling after rolling into the front stage cooling and the rear stage cooling whose cooling rate is slower than that of the previous stage cooling, and appropriately controlling the respective cooling rates, it is possible to obtain acicular ferrite as the main structure, and A steel sheet with an excellent strength-toughness balance is obtained. This is because: by accelerating the cooling rate of the front stage, the nucleation density of the phase transformation can be increased, so that the structure after the phase transformation becomes a dense acicular ferrite structure instead of a coarse bainite structure. It was further found that: when the cooling rate of the rear stage is too fast compared with the cooling rate of the front stage, island-shaped martensite is formed, and the toughness of the base metal is deteriorated. On the other hand, when the cooling rate of the rear stage is too slow, The strength of the base material is reduced, so it needs to be controlled within an appropriate range.

The present invention has been accomplished based on the above findings.

The composition of components that the high-strength steel according to the present invention should have will be described below.

C: 0.03 to 0.10% by mass

C is an element most affecting the strength of steel, and it needs to be contained in an amount of 0.03% by mass or more in order to ensure the strength (YS≧460 MPa) required as structural steel. However, on the contrary, if the content of C is too large, the toughness of the base metal will decrease or cold cracking during welding will occur, so the upper limit of the content of C is 0.10% by mass.

Si: 0.30% by mass or less

Si is a component added as a deoxidizing material to increase the strength of steel. In order to obtain this effect, it is preferable to add 0.01 mass % or more of Si. However, if the content of Si exceeds 0.30% by mass, the toughness of the base metal and the weld will decrease, so it needs to be controlled to 0.30% by mass or less. Preferably, it is the range of 0.01-0.20 mass %.

Mn: 1.60 to 2.30% by mass

Mn is an effective element for ensuring the strength of the base metal. In the present invention, it is to promote the microstructure of the welded heat-affected zone and suppress the formation of embrittlement structures as much as possible, thereby improving the toughness of the welded heat-affected zone (CTOD features) to add important elements. In order to obtain this effect, it is necessary to add 1.60% by mass or more of Mn. On the other hand, when the content of Mn exceeds 2.30% by mass, the toughness of the base metal and the weld will be significantly reduced, so it is controlled to be 2.30% by mass or less. Preferably, it is the range of 1.65-2.15 mass %.

P: 0.015% by mass or less

P is an unavoidable impurity, and if the content of P exceeds 0.015% by mass, the toughness of the base metal and the weld will be reduced, so it is controlled at 0.015% by mass or less. Preferably it is 0.010 mass % or less.

S: 0.005% by mass or less

S is an unavoidable impurity, and if the content of S exceeds 0.005% by mass, the toughness of the base metal and the weld will be reduced, so it is controlled at 0.005% by mass or less. Preferably it is 0.0035% by mass or less.

Al: 0.005 to 0.06% by mass

Al is an element added for deoxidation of molten steel, and needs to be contained in an amount of 0.005% by mass or more. On the other hand, if Al is added in excess of 0.06% by mass, the toughness of the base metal will decrease, and it will be mixed into the weld metal part due to dilution by welding, thereby reducing the toughness. Therefore, it is necessary to control it to 0.06% by mass or less. . Preferably it is 0.010 to 0.055% by mass.

Nb: 0.004 to 0.05% by mass

Since Nb forms a non-recrystallization zone (non-recrystallization zone) in the low temperature range of austenite, the microstructure of the base metal and high toughness can be achieved by performing rolling in this temperature range. In addition, precipitation strengthening (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 obtain the above effects, it is necessary to add 0.004% by mass or more of Nb. However, if Nb is added excessively exceeding 0.05% by mass, the toughness of the welded portion will deteriorate, so the upper limit is made 0.05% by mass.

Ti: 0.005 to 0.02% by mass

Ti precipitates in the form of TiN during molten steel solidification, suppresses the coarsening of austenite in the welded zone, and serves as a transformation nucleus of ferrite, thus contributing to high toughness of the welded zone. In order to obtain this effect, it is necessary to add 0.005% by mass or more of Ti. However, when the added Ti is less than 0.005% by mass, this effect is small. On the other hand, when the added Ti exceeds 0.02% by mass, the TiN particles are coarsened, and the effect of improving the toughness of the base metal and the weld cannot be obtained. Therefore, the addition amount of Ti is controlled within the range of 0.005 to 0.02% by mass.

N: 0.001 to 0.005% by mass

N is an element necessary to form TiN that suppresses the coarsening of the structure of the welded portion, and needs to be added in an amount of 0.001% by mass or more. On the other hand, when the added N exceeds 0.005% by mass, solid solution N significantly reduces the toughness of the base metal and the weld, so the upper limit is controlled to be 0.005% by mass. In addition, in order to form a sufficient amount of TiN during pinning to suppress coarsening of the structure, it is preferable to control N to a range of 0.003 to 0.005% by mass.

Ca: 0.0005 to 0.003% by mass

Ca is an element that improves toughness by fixing S. In order to exhibit this effect, it is necessary to add at least 0.0005% by mass of Ca. However, when the content of Ca exceeds 0.003% by mass, this effect is saturated, so the added Ca is limited to the range of 0.0005 to 0.003% by 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 melt even at high temperatures, the contents of Ca, S, and O need to satisfy the relationship of the following formula (1). Here, Ca, S, and O represent the content (mass %) of each element.

0＜(Ca-(0.18+130×Ca)×O)/1.25/S＜1 …(1)

(Ca-(0.18+130×Ca)×O)/(1.25/S) in the above formula is a value representing the ratio of atomic concentrations of Ca and S effective in controlling the form of sulfide, and sulfidation can be estimated from this value. The form of things (Mochida et al., "Iron and Steel", Japan Iron and Steel Association, 66th (1980), No. 3, P.354~362).

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 is precipitated in the form of MnS alone, fine dispersion of ferrite nuclei in the weld heat-affected zone, which is the main focus of the present invention, cannot be achieved. In addition, independently precipitated MnS is elongated during steel sheet rolling, 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 does not precipitate on CaS as ferrite. MnS, which acts as a nucleus, cannot fully exert the function of complex sulfides as ferrite nuclei.

On the other hand, when Ca, S, and O satisfy the above formula (1), MnS is precipitated on CaS to form complex sulfides, which can effectively function as ferrite formation nuclei. In addition, the value of ((Ca-(0.18+130×Ca)×O)/1.25/S) is preferably in the range of 0.2 to 0.8.

In order to improve strength and toughness, the high-strength steel of the present invention may further contain one or two or more selected from B, V, Cu, Ni, Cr, and Mo in addition to the above-mentioned essential components.

B: 0.0003 to 0.0025% by mass

B segregates to austenite grain boundaries, suppresses ferrite transformation caused by grain boundaries, increases the percentage of bainite structure, and has the effect of increasing the strength of steel. This effect is obtained by adding 0.0003% by mass or more of B. However, when the amount of B added exceeds 0.0025% by mass, the toughness decreases on the contrary. A more preferable range of B is 0.0005 to 0.002% by mass.

V: 0.2% by mass or less

V is an element effective in improving the strength and toughness of the base material, and is also an element that precipitates in the form of VN and functions as a ferrite formation nucleus. In order to obtain this effect, V is preferably added in an amount of 0.01% by mass or more. However, if the added amount exceeds 0.2% by mass, the toughness will conversely decrease, so it is preferable to add V at 0.2% by mass or less. More preferably, it is 0.15 mass % or less.

Cu: 1% by mass or less

Cu is an element having an effect of increasing the strength of steel. In order to obtain this effect, it is preferable to add 0.05% by mass or more of Cu. However, when the content of Cu exceeds 1% by mass, hot embrittlement occurs to deteriorate the surface properties of the steel sheet, so it is preferable to add Cu in a range of 1% by mass or less. More preferably, it is 0.8 mass % or less.

Ni: 2% by mass or less

Ni is an element effective in improving the strength of steel and improving the CTOD characteristics of the welded heat-affected zone. In order to obtain this effect, it is preferable to add 0.05% by mass or more of Ni. However, since Ni is an expensive element, the upper limit is preferably 2.0% by mass. In the present application, when 1.6% or more of Mn is added, the Ni content is more preferably less than 0.3% from the viewpoint of cost reduction.

Cr: 0.7% by mass or less

Cr is an element effective in increasing the strength of the base material. In order to obtain this effect, it is preferable to add 0.05% by mass or more of Cr. However, adding a large amount of Cr adversely affects the toughness, so the upper limit is preferably 0.7% by mass. More preferably, it is 0.5 mass % or less.

Mo: 0.7% by mass or less

Like Cr, Mo is an element effective in increasing the strength of the base material. In order to obtain this effect, it is preferable to add 0.05% by mass or more of Mo. However, adding a large amount of Mo adversely affects the toughness, so the upper limit is preferably 0.7% by mass. More preferably, it is 0.5 mass % or less.

Next, the structure of the high-strength steel of the present invention will be described.

The structure of the high-strength steel of the present invention is mainly composed of acicular ferrite, and its area ratio is preferably 60% or more, more preferably 70% or more. When the area ratio of acicular ferrite is less than 60%, the coarse upper bainite structure increases and the toughness decreases. In addition, the upper limit of the area ratio is not particularly limited. In addition, the acicular ferrite structure of the high-strength steel of the present invention is bainetic ferrite (bainetic ferrite) having a thin needle-like or lath-like form and a high dislocation density in the crystal grains. Polygonal ferrite is different from coarse upper bainite.

Next, a method for producing the high-strength steel of the present invention will be described.

The high-strength steel of the present invention is preferably produced by melting molten steel adjusted to the above-mentioned composition suitable for the present invention by a common method using a converter, an electric furnace, a vacuum melting furnace, etc., followed by continuous casting or ingot-casting Thick-walled high-strength steel is produced by hot-rolling the steel raw materials such as billets after ordinary processes such as rolling. In this case, the heating temperature of the steel material before hot rolling needs to be in the range of 1050 to 1200°C. The reason why the heating temperature is 1050° C. or higher is that casting defects existing in the steel raw material can be reliably press-welded by hot rolling. However, when heated to a temperature exceeding 1200°C, the TiN precipitated during solidification will coarsen and the toughness of the base metal and the weld will decrease, so the heating temperature needs to be limited to 1200°C or lower.

Thereafter, the steel material heated to the above temperature is subjected to hot rolling with a cumulative rolling ratio of 30% or more in a temperature range of 950°C or higher and a cumulative rolling ratio of 30 to 70% in a temperature range of lower than 950°C. , made of high-strength steel with a predetermined plate thickness. The reason why hot rolling is performed with a cumulative rolling ratio of 30% or more in a temperature range of 950°C or higher is that by controlling the cumulative rolling ratio in this temperature range to 30% or higher, recrystallization of austenite grains occurs , so that the structure can be refined. When the cumulative rolling ratio is lower than 30%, the abnormally coarse grains generated during heating will remain, which will adversely affect the toughness of the base metal.

In addition, the reason for implementing hot rolling with a cumulative draft of 30 to 70% in the temperature range below 950°C is that the austenite grains rolled in this temperature range will not be sufficiently formed. After recrystallization, the austenite grains after rolling are deformed into a flat state, and it becomes a steel with high internal strain that contains a large number of defects such as deformation bands. And, the accumulated internal energy acts as a driving force for the subsequent ferrite transformation, thereby promoting the ferrite transformation. However, when the rolling reduction is less than 30%, the above-mentioned accumulated internal energy is insufficient, so ferrite transformation is difficult to occur, and the toughness of the base metal decreases. On the other hand, when the rolling ratio exceeds 70%, the formation of polygonal ferrite will be promoted and the formation of acicular ferrite will be inhibited, making it impossible to achieve both high strength and high toughness.

The cooling after the subsequent hot rolling is divided into front-stage cooling and rear-stage cooling, and the cooling rate of the former is relatively greater than the cooling rate of the latter, that is, in the front-stage cooling, it is necessary to cool at 5 to 45° C./second, preferably 5 to 40° C. A cooling rate of 20°C/sec, more preferably 6-16°C/sec, a cooling stop temperature between 600°C and 450°C from the end temperature of hot rolling, preferably a cooling stop temperature between 580°C and 480°C from the end temperature of hot rolling Temperature; in the subsequent subsequent cooling, it is necessary to cool from the stop temperature of the previous cooling to below 450 °C at a cooling rate of 1 °C/s or more and less than 5 °C/s, more preferably 2 to 4.5 °C/s. The stage cooling stop temperature is preferably a cooling stop temperature between 400°C and 250°C from the stop temperature of the previous stage cooling.

When the stop temperature of the pre-cooling is higher than the above temperature range, there is almost no increase in strength, and on the contrary, when it is lower than the above temperature range, the toughness is deteriorated. In addition, when the front-stage cooling rate is lower than the lower limit of the above range, the structure mainly consists of polygonal ferrite, and the strength cannot be improved. On the contrary, when the cooling rate exceeds the upper limit of the above range, the toughness is lowered. Furthermore, when the cooling stop temperature of the post-cooling is higher than the upper limit of the temperature range, the increase in strength becomes insufficient. Also, when the subsequent cooling rate is lower than the lower limit of the above range, the strength of the base material is insufficient, and conversely, when it exceeds the upper limit of the above range, the toughness of the base material decreases. In addition, when the cooling rate of the subsequent stage is too fast compared with the cooling rate of the preceding stage, insular martensite is formed and the toughness of the base material is deteriorated.

In addition, in the present invention, in order to reduce residual internal stress, tempering treatment (tempering) in a temperature range of 450 to 650° C. may be performed on the steel material after cooling. When the tempering temperature is lower than 450°C, the removal effect of residual stress is small. On the other hand, when the tempering temperature becomes higher than 650°C, various carbonitrides (carbonitride) precipitate and occur. Precipitation strengthens and toughness decreases, so it is not preferable.

As explained above, in the manufacturing method of the high-strength steel of the present invention, it is important to control an appropriate rolling ratio corresponding to the rolling temperature of hot rolling and to properly control the two-stage cooling conditions after the completion of rolling, In particular, by making the cooling rate of the front stage cooling faster than that of the rear stage cooling, the base metal has a structure mainly of acicular ferrite, and a steel material having an excellent strength-toughness balance can be obtained.

In addition, in the present invention, by making the content of N in the chemical composition exceed 0.0030%, the cooling rate of the front stage cooling is set to be more than 20°C/sec and not more than 45°C/sec, and the stop temperature of the front stage cooling is set to be 450°C or higher. And below 500°C, it is possible to inexpensively manufacture a high-strength steel that has high strength with a yield stress of 550 MPa or more of the base material, excellent toughness, and excellent toughness (CTOD characteristics) of the heat-affected zone after welding.

Example

Using steel billets No. 1 to 31 having the compositions shown in Table 1-1 and Table 1-2 as raw materials, hot rolling and pre-cooling were carried out under the conditions shown in Table 2-1 and Table 2-1. After cooling, a thick steel plate with a thickness of 25-80 mm is produced. In addition, the temperature described in Table 2-1 and Table 2-2 is the temperature at 1/4 of the plate thickness calculated from the surface layer temperature of the steel plate measured by a radiation thermometer. Samples were cut from the thick steel plates thus obtained, and subjected to tensile tests and Charpy impact tests. In the tensile test, the JIS No. 4 tensile test piece is cut from 1/4 of the plate thickness of the thick steel plate in such a way that the longitudinal axis of the test piece is parallel to the rolling direction, and the yield stress (YS) and tensile strength (TS) are measured. . In addition, in the Charpy impact test, JIS No. 4 impact test pieces were cut out from the thickness 1/4 of each thick steel plate in the rolling width direction, and the absorbed energy (vE-40°C) at a temperature of -40°C was measured. Subsequently, the test pieces satisfying all the conditions of YS≥460MPa, TS≥570MPa and vE-40°C≥200J were evaluated as having good properties of the base material.

In addition, in principle, a single bevel groove (single bevel groove) (single bevel groove) (groove angle ( bevel angle)30°), conduct carbon dioxide gas arc welding (CO ₂ arc welding) with an input heat of 25kJ/cm, make a welded joint, and cut a straight line of fusion (straight The CTOD test piece with a notch on the bond) part was subjected to the CTOD test at a temperature of -10°C. In addition, the production and test conditions of the CTOD test piece were performed in accordance with British Standard BS7448. In addition, a JIS No. 4 impact test piece with the notch position as the weld line portion was cut out, and a Charpy impact test was performed at a temperature of -40°C to measure absorbed energy (vE-40°C).

The above-mentioned test results are described together and shown in Table 2-1 and Table 2-2. From these results, it can be seen that the steel plate 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 excellent. The vE-40°C of the fusion line of the arc welded joint is also more than 200J, the CTOD value is more than 0.10mm, and the toughness of the welded heat-affected zone is also excellent. On the other hand, in the steels of the comparative examples outside the scope of the present invention, only steel sheets with one or more of the above-mentioned properties deteriorated were obtained.

In addition, among the steel sheets of the examples of the present invention, steel sheet Nos. 11 to 17 in which N exceeded 0.0030% by mass were all excellent in CTOD of the welded portion due to the pinning effect of TiN, and all were 0.45 or more.

Furthermore, in the steel sheet of the example of the present invention, when N exceeds 0.0030% by mass, the cooling rate of the pre-stage cooling after hot rolling exceeds 20°C/sec and is 45°C/sec or less, and the stop temperature of the pre-stage cooling is 450°C or more and low. Steel sheets No. 15 and No. 16 produced under the condition of 500° C. both had high strengths with a yield stress of the base metal being 550 MPa or more.

Industrial Utilization Possibility

The high-strength steel of the present invention is not only suitable for ships or offshore structures, pipeline steel pipes, and pressure vessels, but also for welded and assembled steel structures in the fields of construction/civil engineering and the like.

Claims (6)

1. A high-strength steel, which has the following components: Contains C: 0.03-0.10 mass%, Si: 0.30 mass% or less, Mn: 1.60-2.30 mass%, P: 0.015 mass% or less, S: 0.005 mass% or less, Al: 0.005-0.06 mass%, Nb: 0.004- 0.05% by mass, Ti: 0.005 to 0.02% by mass, N: 0.001 to 0.005% by mass, Ca: 0.0005 to 0.003% by mass, And, the content of Ca, S and O satisfies the following formula (1), and the balance consists of Fe and unavoidable impurities, 0＜(Ca-(0.18+130×Ca)×O)/1.25/S＜1 …(1) Wherein, Ca, S and O represent the mass % content of each element. 2. The high-strength steel according to claim 1, which further contains, in addition to the composition of the above-mentioned components, selected from the group consisting of: B: 0.0003-0.0025 mass%, V: 0.2 mass% or less, Cu: 1 mass% or less, Ni: One or more of 2% by mass or less, Cr: 0.7% by mass or less, and Mo: 0.7% by mass or less. 3. A method for producing high-strength steel, which comprises heating a billet to 1050 to 1200°C, applying a cumulative rolling ratio of 30% or more in a temperature range of 950°C or higher, and in a temperature range lower than 950°C Hot rolling with a cumulative rolling ratio of 30-70%, and then performing front-stage cooling and rear-stage cooling. The front-stage cooling is cooled from the end temperature of hot rolling at 5-45°C/second to the cooling stop temperature between 600-450°C, and the rear-stage cooling The stage cooling is cooled from the preceding stage cooling stop temperature to the cooling stop temperature below 450°C at a rate of 1°C/sec or more and less than 5°C/sec, Wherein, the billet has the following composition: Contains C: 0.03-0.10 mass%, Si: 0.30 mass% or less, Mn: 1.60-2.30 mass%, P: 0.015 mass% or less, S: 0.005 mass% or less, Al: 0.005-0.06 mass%, Nb: 0.004- 0.05% by mass, Ti: 0.005 to 0.02% by mass, N: 0.001 to 0.005% by mass, Ca: 0.0005 to 0.003% by mass, And, the content of Ca, S and O satisfies the following formula (1), and the balance consists of Fe and unavoidable impurities, 0＜(Ca-(0.18+130×Ca)×O)/1.25/S＜1 …(1) Wherein, Ca, S and O represent the mass % content of each element. 4. The method for producing high-strength steel according to claim 3, wherein the high-strength steel further contains B: 0.0003-0.0025% by mass, V: 0.2% by mass or less, One or more of Cu: 1% by mass or less, Ni: 2% by mass or less, Cr: 0.7% by mass or less, and Mo: 0.7% by mass or less. 5. The method for producing high-strength steel according to claim 3 or 4, wherein the tempering treatment at 450 to 650° C. is performed on the steel after subsequent cooling. 6. The method for producing high-strength steel according to any one of claims 3 to 5, wherein the pre-stage cooling is performed at a rate of 5 to 20° C./second.

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