WO2015088040A1 - 鋼板およびその製造方法 - Google Patents
鋼板およびその製造方法 Download PDFInfo
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- WO2015088040A1 WO2015088040A1 PCT/JP2014/083321 JP2014083321W WO2015088040A1 WO 2015088040 A1 WO2015088040 A1 WO 2015088040A1 JP 2014083321 W JP2014083321 W JP 2014083321W WO 2015088040 A1 WO2015088040 A1 WO 2015088040A1
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Definitions
- the present invention relates to a high-tensile steel plate used for steel structures such as marine structures, ships, pressure vessels, and penstocks, and a method for producing the same.
- the yield stress (YS) is 460 MPa or more, and the strength and toughness of the steel plate is increased.
- the present invention relates to a high-thickness and high-tensile steel sheet that is not only excellent, but also excellent in low-temperature toughness (CTOD characteristics) of multi-layer welds and strength and toughness (PWHT characteristics) after heat treatment during welding (PWHT characteristics), and a manufacturing method thereof .
- Steel plates used in ships, offshore structures, pressure vessels, etc. are welded and finished as structures of a desired shape. Therefore, these steel sheets have high strength and excellent toughness from the viewpoint of the safety of the structure, as well as welded joints (welded metals and heat-affected zones) when welding is performed. It is also required to have excellent toughness.
- the CTOD test since a fatigue precrack is used, a very small region becomes the toughness evaluation part. Therefore, if a local embrittlement region exists in the steel sheet, even if good toughness is obtained in the Charpy impact test, the CTOD test may show a low value.
- This local embrittlement region is likely to occur in a welding heat-affected zone (hereinafter also referred to as HAZ) that is subjected to a complex thermal history by multi-layer welding in a thick steel plate or the like.
- HAZ welding heat-affected zone
- the boundary of the base material) and the part where the bond part is reheated in the two-phase region (coarse grains in the first cycle welding, the region heated to the two-phase region of ferrite and austenite by the subsequent welding pass, (Referred to as a two-phase region reheating part) tends to be a local embrittlement region.
- the bond portion since the bond portion is exposed to a high temperature just below the melting point during welding, the austenite grains are coarsened and are easily transformed into an upper bainite structure having low toughness by subsequent cooling, so the toughness of the matrix itself tends to be low. Further, the bond portion is liable to generate an embrittled structure such as a woodman-stetten structure or island martensite (hereinafter also referred to as MA), and when this embrittled structure is generated, the toughness of the steel sheet is likely to further decrease. Become.
- TiN is finely dispersed in a steel plate to suppress coarsening of austenite grains or use it as a ferrite transformation nucleus.
- the bonding part may be heated to a temperature range where TiN dissolves, and the heating temperature is higher as the requirement for low temperature toughness of the welded part becomes more severe. Therefore, the effect of finely dispersing TiN described above appears. It becomes difficult.
- Patent Document 1 and Patent Document 2 suppress the austenite grain growth by adding rare earth elements (REM) together with Ti and dispersing fine particles in the steel sheet.
- REM rare earth elements
- Patent Document 1 and Patent Document 2 include a technique for dispersing an oxide of Ti, a technique for combining ferrite nucleation ability of BN and oxide dispersion, and further adding Ca and REM to form a sulfide. Techniques for increasing toughness by controlling the thickness have been proposed.
- Patent Document 3 discloses a technique for improving the HAZ toughness by dispersing Ti oxide in steel.
- the two-phase region reheating part carbon is concentrated in a region reversely transformed into austenite by the two-phase region reheating, and a brittle bainite structure including island martensite is generated during cooling.
- the toughness is reduced, in order to prevent this toughness reduction, the steel plate component is made low C and low Si, the formation of island martensite is suppressed to improve the toughness, and the base material strength is ensured by adding Cu. Techniques are disclosed (for example, Patent Documents 4 and 5).
- Patent Document 4 employs a method in which the cooling rate after rolling is set to 0.1 ° C./s or less and Cu particles are precipitated in this process, but there is a problem in manufacturing stability.
- the N / Al ratio is set to 0.3 to 3.0 to suppress the deterioration of toughness due to the adverse effect of AlN coarsening or solid solution N. Is more easily controlled by Ti.
- post-heat treatment may be performed at the time of welding. At this time, since the base material is also heated at the same time, it is necessary to maintain the base material characteristics even when subjected to the PWHT treatment. It has been common to add elements to form.
- Japanese Patent Publication No. 03-053367 JP 60-184663 A Japanese Patent No. 3697202 Japanese Patent No. 3045856 Japanese Patent No. 4432905
- Patent Documents 1 and 2 are intended for steel materials having relatively low strength and a small amount of alloy elements.
- the HAZ structure is used in the case of thick materials having higher strength and a large amount of alloy elements.
- tissue which does not contain a ferrite there exists a problem that it cannot apply.
- Patent Document 3 has a problem that it is difficult to stably finely disperse Ti oxide in steel.
- steel structures such as ships, offshore structures, pressure vessels, and penstocks have been required to have higher strength as their size has increased.
- Steel materials used for these steel structures are, for example, many thick materials having a plate thickness of 35 mm or more. Therefore, in order to ensure a yield stress of 460 MPa class or higher, steel containing a large amount of alloying elements to be added.
- the component system is advantageous.
- the present invention advantageously solves the above-mentioned problems, and has a yield stress (YS) of 460 MPa or more suitable for use in steel structures such as offshore structures, ships, pressure vessels, and penstocks.
- YS yield stress
- PWHT characteristics strength and toughness
- COD characteristics low temperature toughness
- An object is to provide a manufacturing method.
- (B) In order to improve the toughness of the weld heat affected zone, it is effective to effectively use TiN to suppress the austenite grain coarsening in the vicinity of the weld bond zone.
- TiN can be uniformly and finely dispersed in the steel.
- the present invention has been completed based on the above-described knowledge, and the gist of the present invention is as follows. 1. % By mass C: 0.020 to 0.090% Si: 0.01 to 0.35% Mn: 1.40 to 2.00% P: 0.008% or less S: 0.0035% or less Al: 0.010 to 0.060% Ni: 0.40 to 2.00% Mo: 0.05 to 0.50% Nb: 0.005 to 0.040% Ti: 0.005 to 0.025% N: 0.0020 to 0.0050% Ca: 0.0005 to 0.0050% O: 0.0035% or less, Ceq defined by the following formula (1) is in the range of 0.420 to 0.520%, and the following formulas (2), (3) and (4) While satisfying the formula, B is suppressed to less than 0.0003%, the balance is a steel plate component consisting of Fe and inevitable impurities, Ti, Nb, and Mo are changed into Ti amount ([Ti]), Nb amount ([Nb]), and Mo amount ([Mo]), [Nb] /
- the component composition (steel component) of a steel plate (hereinafter also referred to as a thick material) is limited to the above range will be described in detail for each component.
- the% display which shows the component composition of the steel plate described below means the mass%.
- C 0.020 to 0.090% C is an element necessary for securing strength as a high-tensile steel plate. If the addition is less than 0.020%, the hardenability decreases, and a large amount of hardenability improving elements such as Cu, Ni, Cr, and Mo are required to ensure strength, resulting in high costs. On the other hand, addition exceeding 0.090% lowers the weld zone toughness. Therefore, the C content is in the range of 0.020 to 0.090%. Preferably, it is in the range of 0.020 to 0.080%.
- Si 0.01 to 0.35%
- Si is a component added as a deoxidizing element and for obtaining steel plate strength. To obtain these effects, addition of 0.01% or more is necessary. On the other hand, a large amount of addition exceeding 0.35% causes a decrease in weldability and a decrease in weld joint toughness. Accordingly, the Si amount needs to be in the range of 0.01 to 0.35%. Preferably, it is 0.01 to 0.23%.
- Mn 1.40 to 2.00% Mn needs to be added in an amount of 1.40% or more in order to ensure the strength of the steel plate and the welded joint. On the other hand, addition exceeding 2.00% reduces weldability, the hardenability becomes excessive, and the steel plate toughness and weld joint toughness are reduced. Therefore, the amount of Mn is set to a range of 1.40 to 2.00%. More preferably, it is 1.40 to 1.95%.
- P 0.008% or less
- P which is an impurity element, lowers the steel sheet toughness and weld zone toughness, and in particular, when the content in the weld zone exceeds 0.008%, the CTOD characteristics are significantly lowered.
- the following. Preferably, it is 0.006% or less.
- the P content is preferably as small as possible, the lower limit is about 0.002% from the viewpoint of refining costs and the like.
- S 0.0035% or less
- S is an impurity element, and if contained in excess of 0.0035%, the toughness of the steel sheet and welded portion is lowered, so the content is made 0.0035% or less. Preferably, it is 0.0030% or less.
- the S content is preferably as small as possible, but the lower limit is about 0.0004% from the viewpoint of refining costs and the like.
- Al 0.010 to 0.060%
- Al is an element added for deoxidizing molten steel, and it is necessary to contain 0.010% or more.
- the toughness of the steel sheet and welded portion is lowered and mixed into the welded metal portion by dilution by welding to lower the toughness, so the content is limited to 0.060% or less.
- it is 0.017 to 0.055%.
- the amount of Al is defined by acid-soluble Al (also referred to as Sol.Al or the like).
- Ni 0.40 to 2.00%
- Ni is an element effective for improving the strength and toughness of the steel sheet, and is also effective for improving the welded portion CTOD characteristics. To obtain this effect, 0.40% or more must be added.
- Ni is an expensive element, and excessive addition tends to cause scratches on the surface of the slab during casting, so the upper limit of content is 2.00%.
- Mo 0.05 to 0.50% Mo plays an important role in the present invention and is an effective element for increasing the strength of a steel sheet by adding an appropriate amount. This is the effect of improving hardenability and softening resistance during tempering. Moreover, the composite precipitate formed with Ti and Nb is maintained finely, and there is an effect of strengthening the thick material and suppressing toughness reduction. In order to acquire these effects, it is necessary to contain 0.05% or more of Mo. On the other hand, if contained excessively, the toughness of the thick material is adversely affected, so the upper limit of the Mo amount is 0.50%.
- the Mo amount is more preferably in the range of 0.08 to 0.40%. Further, it is more preferably in the range of 0.16 to 0.30%.
- Nb 0.005 to 0.040% Since Nb forms a non-recrystallized area of austenite at a low temperature range, the structure of the steel sheet can be refined and toughened by rolling in that temperature range. Further, Nb has an effect of improving hardenability and has an effect of increasing softening resistance at the time of tempering by being added in combination with Mo and Ti, and is also an element effective for improving the strength of the steel sheet. In order to acquire these effects, it is necessary to contain Nb 0.005% or more. On the other hand, if the content exceeds 0.040%, the toughness is deteriorated, so the upper limit of the Nb amount is 0.040%, preferably 0.035%.
- Ti 0.005 to 0.025%
- Ti precipitates as TiN when the molten steel is solidified, and suppresses austenite coarsening in the welded portion, thereby contributing to improvement in the toughness of the welded portion.
- Mo and Nb has the effect of increasing the softening resistance during tempering.
- the content is less than 0.005%, the effect is small.
- TiN becomes coarse, and the effect of improving the toughness of the steel sheet or the welded portion cannot be obtained.
- the range is 0.005 to 0.025%.
- N 0.0020 to 0.0050% N reacts with Ti and Al to form precipitates, thereby refining crystal grains and improving steel sheet toughness. Moreover, it is an element required in order to form TiN which suppresses the coarsening of the structure
- Ca 0.0005 to 0.0050%
- Ca is an element that improves toughness by fixing S. In order to obtain this effect, addition of at least 0.0005% is necessary. On the other hand, even if the content exceeds 0.0050%, the effect is saturated, so Ca is added in the range of 0.0005 to 0.0050%.
- O 0.0035% or less O exceeds 0.0035%, so that the toughness of the steel sheet deteriorates, so 0.0035% or less, preferably 0.0028% or less.
- the O content is preferably as low as possible, but the lower limit is about 0.0010% from the viewpoint of refining costs and the like.
- Ceq 0.420 to 0.520%
- a thick material strength of 460 MPa class cannot be obtained.
- the content is made 0.520% or less.
- it is in the range of 0.440 to 0.520%.
- [M] represents the content (mass%) of the element M in steel.
- [Ti] / [N] 1.5 to 4.0 If the value of [Ti] / [N] is less than 1.5, the amount of TiN produced decreases, and solid solution N that does not become TiN reduces the toughness of the weld. On the other hand, when the value of [Ti] / [N] exceeds 4.0, TiN becomes coarse and the weld zone toughness is lowered. Therefore, the range of [Ti] / [N] is 1.5 to 4.0, preferably 1.8 to 3.5.
- 0 ⁇ [Ca] ⁇ (0.18 + 130 ⁇ [Ca]) ⁇ [O] ⁇ / 1.25 / [S] ⁇ 1.5 ⁇ [Ca] ⁇ (0.18 + 130 ⁇ [Ca]) ⁇ [O] ⁇ / 1.25 / [S] is a value indicating the ratio of atomic concentrations of Ca and S effective for sulfide morphology control. It can be adjusted by controlling the amount of oxygen added and the amount of dissolved oxygen in the molten steel at the time of addition to an appropriate range, and is also referred to as ACR (Atomic Concentration Ratio).
- the form of sulfide can be estimated from this ACR value, in the present invention, it is defined as an index for finely dispersing the ferrite transformation nuclei CaS that does not dissolve even at high temperatures.
- the ACR value when the ACR value is 0 or less, CaS does not crystallize. Therefore, since S precipitates in the form of MnS alone, it easily dissolves in the weld heat affected zone and ferrite formation nuclei cannot be obtained.
- MnS precipitated alone is elongated during rolling and causes a reduction in the toughness of the steel sheet. Therefore, in the present invention, the ACR value needs to exceed zero.
- the ACR value when the ACR value is 1.5 or more, the ratio of the oxide in the Ca-based inclusion increases, the ratio of the sulfide functioning as a transformation nucleus decreases, and the toughness improving effect cannot be obtained. Therefore, in the present invention, the ACR value needs to be less than 1.5. Therefore, when the ACR value is controlled to be more than 0 and less than 1.5, a composite sulfide mainly composed of CaS can be effectively formed and function effectively as a ferrite forming nucleus.
- the ACR value is preferably in the range of 0.15 to 1.30. More preferably, it is in the range of 0.20 to 1.00.
- the CTOD test is a test at the full thickness of the steel sheet, the test piece includes center segregation, and when the concentration of components at the center segregation is significant, a hardened zone is generated in the weld heat affected zone. Good results cannot be obtained. Therefore, in the present invention, by controlling the Ceq * value within an appropriate range, an excessive increase in hardness in the center segregation portion is suppressed, and excellent CTOD characteristics can be obtained even in a welded portion of a steel plate having a large plate thickness.
- Appropriate range of Ceq * values are those determined experimentally, Ceq * value is to 3.70 or less because CTOD properties deteriorate exceeds 3.70. Preferably it is 3.50 or less.
- limiting in particular in the minimum of Ceq * value About 2.2 is preferable from a viewpoint of productivity.
- Cu less than 0.7%
- Cr 0.1 to 1.0%
- V 0.005 to 0.05 in order to improve hardenability. 1 type or 2 types or more chosen from% can be contained.
- Cu Less than 0.7% By adding Cu, the steel plate strength can be improved. However, since addition exceeding 0.7% reduces hot ductility, it limits to 0.7% or less. Preferably, it is 0.1 to 0.6%.
- Cr 0.1 to 1.0% Cr is an element effective for increasing the strength of a steel sheet, and in order to exert this effect, it contains 0.1% or more. However, if it is excessively contained, the toughness is adversely affected. Therefore, when it is contained, the range is preferably 0.1 to 1.0%, and more preferably 0.2 to 0.8%.
- V 0.005 to 0.05%
- V is an element effective in improving the strength and toughness of the steel sheet when contained in an amount of 0.005% or more. However, if the content exceeds 0.05%, the toughness is reduced. It is preferably 0.05%.
- one or two selected from Mg: 0.0002 to 0.0050% and REM: 0.0010 to 0.0200% Seeds can be included.
- Mg and REM are elements having an effect of improving toughness due to oxide dispersion. In order to exhibit such effects, 0.0002% or more of Mg and 0.0010% or more of REM are added. On the other hand, even if Mg exceeds 0.0050% and REM exceeds 0.0200%, the effect is only saturated. Therefore, when adding these elements, it is preferable to set it as the above-mentioned range, respectively. More preferably, Mg is 0.0005 to 0.0020% and REM is 0.0020 to 0.0150%.
- the above components other than the steel plate components are Fe and inevitable impurities, but particularly B segregates at the austenite grain boundaries when the steel plate is cooled from the austenite region, and suppresses the ferrite transformation.
- B segregates at the austenite grain boundaries when the steel plate is cooled from the austenite region, and suppresses the ferrite transformation.
- Fig. 1 shows the relationship between precipitate size and precipitate composition after PWHT, and changes in strength and toughness ( ⁇ TS, ⁇ vTrs) before and after PWHT.
- Fig. 2 shows TEM replica observation of precipitates in steel. An EDX analysis result is shown.
- ⁇ TS must satisfy the ranges of 5 to ⁇ 15 MPa and ⁇ vTrs of 10 to ⁇ 5 ° C., respectively, from the viewpoint of stability. And in order to satisfy the range, while suppressing the average size of the precipitates to 20 nm or less, the Ti amount (represented as [Ti]), the Nb amount (represented as [Nb]) and the Mo amount (represented by [Nb]). It can be seen from FIG. 1 that [Mo] is required to satisfy the relationship [Nb] / ([Ti] + [Nb] + [Mo]) ⁇ 0.3.
- the above-mentioned precipitate is a precipitate of Ti, Nb and Mo, but the amount of Ti in the precipitate, Nb Since the amount and the amount of Mo only need to satisfy the relationship [Nb] / ([Ti] + [Nb] + [Mo]) ⁇ 0.3, the precipitates should be at least Nb precipitates, Ti and Mo precipitates may be included within a range satisfying this relationship.
- the PWHT characteristics are excellent when ⁇ TS is in the range of 5 to ⁇ 15 MPa and ⁇ vTrs is in the range of 10 to ⁇ 5 ° C.
- the precipitate (composite precipitate) in the present invention is a precipitate of Mo, Ti, Nb, specifically, a carbide, nitride, or carbonitride of Mo, Ti, Nb, Or a mixture of these.
- the method for obtaining the precipitate particle diameter in the present invention is based on the TEM replica method. That is, after appropriately collecting precipitates of carbides of Ti, Nb, and Mo in steel, an average equivalent circle diameter was obtained using image processing from observation with four fields of view at 100,000 times, and this was determined as the particle diameter of the precipitates. And In the present invention, the lower limit value of the measurement target of the precipitate particle size is 2 nm. This is because it is difficult to measure with precipitates having a particle size of less than this.
- the steel of the present invention is preferably produced by the production method described below.
- a desired plate thickness is obtained by hot rolling, followed by cooling and tempering as necessary.
- the slab heating temperature and the rolling reduction are specified.
- the temperature condition of the steel sheet is defined by the temperature at the center of the thickness of the steel sheet.
- the temperature at the center of the plate thickness is obtained by simulation calculation or the like from the plate thickness, surface temperature, cooling conditions, and the like. For example, the temperature at the center of the plate thickness can be obtained by calculating the temperature distribution in the plate thickness direction using the difference method.
- Slab heating temperature 950-1150 ° C
- the slab heating temperature is set to 950 ° C. or higher so that casting defects existing in the slab are steadily pressed by hot rolling.
- the upper limit of the heating temperature is set to 1150 ° C.
- Cumulative rolling reduction of hot rolling in a temperature range of 900 ° C. or higher 30% or higher Heat in a temperature range of 900 ° C. or higher in order to make the austenite grains finer by recrystallization by detoxification of casting defects.
- the cumulative rolling reduction of the hot rolling is set to 30% or more. If it is less than 30%, coarse grains generated during heating remain, which adversely affects the toughness of the steel sheet.
- the upper limit of the cumulative rolling reduction of hot rolling in a temperature range of 900 ° C. or higher is not particularly limited, it is about 95% industrially.
- accelerated cooling to at least 500 ° C. with a cooling rate of 1.0 ° C./s or more. This is because if the cooling rate is less than 1.0 ° C./s, sufficient steel sheet strength cannot be obtained.
- the minimum of the stop temperature of accelerated cooling is not specifically limited, You may carry out to room temperature.
- Tempering temperature 450-650 ° C
- a sufficient tempering effect cannot be obtained at a tempering temperature of less than 450 ° C.
- tempering at a temperature exceeding 650 ° C. is not preferable because precipitates may become coarse and the toughness may be reduced or the strength may be reduced.
- the tempering treatment of the present invention is more preferable because induction heating is used to suppress the coarsening of carbides during tempering. In that case, it is desirable that the center temperature of the steel sheet calculated by a simulation such as a difference method is 450 to 650 ° C.
- the tempering process may not be performed.
- the thick material of the present invention has a thickness of 15 mm or more. Accordingly, in the present invention, the term “thick” means that the thickness of the steel is 15 mm or more, but the effect of the present invention is most obtained when the thickness of the steel is in the range of 40 to 100 mm.
- the manufacturing conditions other than the manufacturing conditions of the above-described thick high-strength steel may be in accordance with ordinary methods.
- the thick-walled high-strength steel according to the present invention refines the structure of the weld heat-affected zone by finely dispersing ferrite transformation nuclei that do not melt even at high temperatures while suppressing the coarsening of austenite grains in the weld heat-affected zone. Therefore, high toughness can be obtained. Even in the region where reheating is performed in the two-phase region by the thermal cycle during multi-layer welding, the structure of the weld heat affected zone by the first welding is refined, so the toughness of the untransformed region in the two-phase region reheating region As a result, the austenite grains that are retransformed can be made finer, and the degree of decrease in toughness can be reduced. In addition, by forming fine composite precipitates of Ti, Nb, and Mo, a thick high-tensile steel sheet having excellent CTOD characteristics and PWHT characteristics can be obtained.
- the precipitation part in steel was extract
- the PWHT heat treatment was performed at 580 ° C. for 4 hours and at a temperature increase / decrease rate of 70 ° C./h.
- Table 3 shows steel sheet characteristics, Charpy impact test results and CTOD test results, precipitate size / composition, and steel sheet characteristics changes after PWHT, along with hot rolling conditions and heat treatment conditions.
- steel symbols A to E are compatible steels of the present invention
- steel symbols F to Z are comparative steels whose steel components are outside the scope of the present invention.
- Sample No. 1, 2, 5, 6, 8, and 11 are all inventive examples, the Charpy impact test result of the weld bond part, the three-point bending CTOD test result of the weld bond part, the precipitate size and composition in the steel plate, and the PWHT A result that satisfies the target in all of the characteristics is obtained.
- sample No. 3, 4, 7, 9, 10, 12 to 31 are at least one of steel plate components, production conditions, precipitate size and composition is outside the scope of the present invention, steel plate characteristics, and Charpy impact test results of welded bonds.
- One of the three-point bending CTOD test results and PWHT characteristics of the weld bond part did not satisfy the target.
- a horizontal line item means that the item could not be measured.
- the steel of the inventive example according to the present invention has a yield stress (YS) of the steel plate of 460 MPa or more and a Charpy absorbed energy (vE- 40 ° C.) of 200 J or more.
- Yield stress (YS) of the steel plate of 460 MPa or more
- Charpy absorbed energy (vE- 40 ° C.) of 200 J or more.
- vE- 40 ° C is 150 J or more
- the CTOD value is 0.5 mm or more
- the weld heat affected zone is also excellent in toughness.
- the average particle size of the precipitate is 20 ⁇ m or less and [Nb] / ([Ti] + [Nb] + [Mo]) ⁇ 0.3
- the steel sheet characteristics after PWHT are also excellent.
- the comparative example outside the scope of the present invention it can be seen that only a steel sheet inferior in any of the above characteristics is obtained.
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Abstract
Description
併せて、特許文献1や特許文献2には、Tiの酸化物を分散させる技術や、BNのフェライト核生成能と酸化物分散を組み合わせる技術、さらにはCaやREMを添加して硫化物の形態を制御することにより、靭性を高める技術などが提案されている。
なお、YSが460MPaを超える厚肉材においては、溶接施工時に後熱処理(PWHT)を施される場合がある。このとき母材も同時に加熱されるため、PWHT処理を受けても母材特性を保持しなければならないが、従来、熱を受けた場合の強度低下を抑えるためには、その温度で析出物を形成する元素を添加することが一般的であった。
これら鉄鋼構造物に用いられる鋼材は、例えば、板厚が35mm以上の厚肉材が多いので、降伏応力460MPa級やそれ以上の強度を確保するためには、添加する合金元素を多く含有する鋼成分系が有利となっている。
(a)CTOD特性は鋼板全厚の試験片で評価されるため、成分の濃化する中心偏析部が破壊の起点となる。
従って、溶接熱影響部のCTOD特性を向上するためには、鋼板の中心偏析として濃化しやすい元素を適正量に制御し、中心偏析部の硬化を抑制することが効果的である。また、溶鋼が凝固する際に最終凝固部となるスラブの中心において、C、Mn、P、NiおよびNbが他の元素に比べて濃化度が高いため、これらの元素の添加量を中心偏析部の硬さ指標により制御して、中心偏析部での硬さを抑制することが効果的である。
CaSは、酸化物に比べて低温で晶出するため、均一に微細分散することができる。そして、Caの添加量および添加時の溶鋼中の溶存酸素量を適正範囲に制御することによって、CaS晶出後でも固溶Sが確保されるので、CaSの表面上にMnSが析出して複合硫化物を形成する。このMnSの周囲には、Mnの希薄帯が形成されるので、フェライト変態がより促進される。
従来、YS:460MPa超級の鋼板では、PWHT後に強度の低下が顕著であったが、開発鋼板では、微細なMo,Ti,Nb複合析出物(炭化物、窒化物または炭窒化物)が安定して存在することで、析出強化を維持することができ、鋼板の強度低下を抑制することが可能であることが分かった。また、微細なMo,Ti,Nb複合析出物の存在により、鋼板の靭性も併せて維持できることが分かった。
1.質量%で、
C:0.020~0.090%
Si:0.01~0.35%
Mn:1.40~2.00%
P:0.008%以下
S:0.0035%以下
Al:0.010~0.060%
Ni:0.40~2.00%
Mo:0.05~0.50%
Nb:0.005~0.040%
Ti:0.005~0.025%
N:0.0020~0.0050%
Ca:0.0005~0.0050%
O:0.0035%以下
を含有し、下記(1)式で規定されるCeqが0.420~0.520%の範囲であって、下記(2)式、(3)式および(4)式を満たすと共に、Bを0.0003%未満に抑制し、残部がFeおよび不可避的不純物からなる鋼板成分と、
Ti、NbおよびMoを、Ti量(〔Ti〕)、Nb量(〔Nb〕)およびMo量(〔Mo〕)が、〔Nb〕/(〔Ti〕+〔Nb〕+〔Mo〕)≧0.3の関係を満足する範囲で含みかつ平均粒子径が20nm以下の析出物を有する鋼板。
記
Ceq=[C]+[Mn]/6+([Cu]+[Ni])/15+([Cr]+[Mo]+[V])/ 5・・・(1)
1.5≦[Ti]/[N]≦4.0 ・・・(2)
0<{[Ca]−(0.18+130×[Ca])×[O]}/1.25/[S]<1.5 ・・・(3)
5.5[C](4/3)+15[P]+0.90[Mn]+0.12[Ni]+7.9[Nb](1/2)+0.53[Mo] ≦3.70 ・・・(4)
但し、[M]は、鋼板中の元素Mの含有量(質量%)を表す。
まず、本発明において、鋼板(以下、厚肉材ともいう)の成分組成(鋼成分)を上記の範囲に限定した理由について、成分毎に詳しく説明する。なお、以下に述べる鋼板の成分組成を示す%表示は特に断らない限り質量%を意味する。
C:0.020~0.090%
Cは、高張力鋼板としての強度確保に必要な元素である。0.020%未満の添加では、焼入性が低下し、強度確保のために、Cu、Ni、CrおよびMoなどの焼入性向上元素の多量添加が必要となって、コスト高を招く。一方、0.090%を超える添加は溶接部靭性を低下させる。従って、C量は0.020~0.090%の範囲とする。好ましくは、0.020~0.080%の範囲である。
Siは、脱酸元素として、また、鋼板強度を得るために添加する成分であり、これらの効果を得るためには0.01%以上の添加が必要である。一方、0.35%を超える多量の添加は、溶接性の低下と溶接継手靭性の低下を招く。従って、Si量は0.01~0.35%の範囲とする必要がある。好ましくは、0.01~0.23%である。
Mnは、鋼板強度および溶接継手強度を確保するため、1.40%以上添加する必要がある。一方、2.00%を超える添加は、溶接性を低下させ、焼入性が過剰となって、鋼板靭性および溶接継手靭性を低下させる。従って、Mn量は1.40~2.00%の範囲とする。さらに好ましくは、1.40~1.95%である。
不純物元素であるPは、鋼板靭性および溶接部靭性を低下させ、特に溶接部における含有量が0.008%を超えるとCTOD特性が著しく低下するので、0.008%以下とする。好ましくは、0.006%以下である。なお、Pの含有量は、極力少ないほうが良いが、精錬コスト等の点から、その下限値は、0.002%程度である。
Sは、不純物元素であり、0.0035%を超えて含有すると鋼板および溶接部靭性を低下させるため、0.0035%以下とする。好ましくは、0.0030%以下である。なお、Sの含有量は、極力少ないほうが良いが、精錬コスト等の点から、その下限値は、0.0004%程度である。
Alは、溶鋼を脱酸するために添加される元素であり、0.010%以上含有させる必要がある。一方、0.060%を超えて添加すると鋼板および溶接部靭性を低下させるとともに、溶接による希釈によって溶接金属部に混入し、靭性を低下させるので、0.060%以下に制限する。好ましくは、0.017~0.055%である。なお、本発明においてAl量は、酸可溶性Al(Sol.Alなどとも称される)で規定するものとする。
Niは、鋼板の強度と靭性の向上に有効な元素であり、溶接部CTOD特性の向上にも有効である。この効果を得るには0.40%以上の添加が必要である。一方、Niは高価な元素であること、また過度の添加は鋳造時にスラブの表面にキズを発生しやすくするので、含有する上限は2.00%とする。
Moは、本発明において重要な役割を果たし、適量添加によって鋼板を高強度化するのに有効な元素である。これは、焼入れ性と、焼き戻し時の軟化抵抗性の向上による効果である。また、TiやNbと形成する複合析出物を微細に維持し、厚肉材の強化と靭性低下抑制の効果がある。これらの効果を得るためには、Moを0.05%以上含有する必要がある。一方、過剰に含有すると、厚肉材の靭性に悪影響を与えるので、Mo量の上限は0.50%とする。なお、Mo量は0.08~0.40%の範囲であることがより好ましい。また、0.16~0.30%の範囲であることがさらに好ましい。
Nbは、低温域で、オーステナイトの未再結晶域を形成するので、その温度域で圧延を施すことにより、鋼板の組織微細化や、高靭化を図ることができる。また、Nbは、焼入れ性の向上効果を有するとともに、MoやTiと複合添加することで、焼戻し時の軟化抵抗を高める効果を有し、鋼板強度の向上に有効な元素でもある。これらの効果を得るためには、Nbを0.005%以上含有する必要がある。一方、0.040%を超えて含有すると靭性を劣化させるので、Nb量の上限は、0.040%とし、好ましくは0.035%とする。
Tiは、溶鋼が凝固する際にTiNとなって析出し、溶接部におけるオーステナイトの粗大化を抑制して溶接部の靭性向上に寄与する。さらにMo,Nbと併せて複合添加することで焼戻し時の軟化抵抗を高める効果がある。しかし、0.005%未満の含有ではその効果が小さい一方で、0.025%を超えて含有すると、TiNが粗大化し、鋼板や溶接部の靭性改善効果が得られないため、Tiは、0.005~0.025%の範囲とする。
Nは、TiやAlと反応して析出物を形成することで、結晶粒を微細化し、鋼板靭性を向上させる。また、溶接部の組織の粗大化を抑制するTiNを形成させるために必要な元素である。これらの作用を発揮するには、Nを0.0020%以上含有することが必要である。一方、Nは、0.0050%を超えて添加すると固溶Nが鋼板や溶接部の靭性を著しく低下させたり、TiおよびNbの複合析出物生成による固溶Nb減少に伴う強度低下を招いたりするので、上限を0.0050%とする。
Caは、Sを固定することによって靭性を向上させる元素である。この効果を得るためには、少なくとも0.0005%の添加が必要である。一方、0.0050%を超えて含有してもその効果は飽和するため、Caは0.0005~0.0050%の範囲で添加する。
Oは、0.0035%を超えると鋼板の靭性が劣化するため、0.0035%以下、好ましくは0.0028%以下とする。なお、Oの含有量は、極力少ないほうが良いが、精錬コスト等の点から、その下限値は、0.0010%程度である。
以下の式で規定されるCeqが0.420%未満の場合、460MPa級の厚肉材強度が得られない。一方、0.520%を超えると、厚肉材の溶接性や溶接部靭性が低下するため、0.520%以下とする。好ましくは、0.440~0.520%の範囲である。なお、以下、[M]は元素Mの鋼中含有量(質量%)を表す。また、含有しない元素は0で計算する。
Ceq=[C]+[Mn]/6+([Cu]+[Ni])/15+([Cr]+[Mo]+[V])/ 5
[Ti]/[N]の値が1.5未満では生成するTiN量が減少し、TiNとならない固溶Nが溶接部の靭性を低下させてしまう。一方、[Ti]/[N]の値が4.0を超えると、TiNが粗大化し、溶接部靭性を低下させる。従って、[Ti]/[N]の値の範囲は1.5~4.0、好ましくは、1.8~3.5とする。
{[Ca]−(0.18+130×[Ca])×[O]}/1.25/[S]は、硫化物形態制御に有効なCaとSの原子濃度の比を示す値で、Caの添加量および添加時の溶鋼中の溶存酸素量を適正範囲に制御することによって調整することができ、ACR(Atomic Concentration Ratio)とも称される。このACR値によって硫化物の形態を推定することができるが、本発明では、高温でも溶解しないフェライト変態生成核CaSを微細分散させる指標として規定する。
ここで、ACR値が0以下の場合、CaSが晶出しない。そのため、Sは、MnS単独の形態で析出するので、溶接熱影響部では容易に固溶してしまいフェライト生成核が得られない。また、単独で析出したMnSは、圧延時に伸長されて、鋼板の靭性低下を引き起こしてしまう。従って、本発明では、ACR値を0超とする必要がある。
一方、ACR値が1.5以上の場合には、Ca系介在物中の酸化物の割合が多くなって、変態核として機能する硫化物の割合が低下し、靭性向上効果が得られない。従って、本発明では、ACR値を1.5未満とする必要がある。
従って、ACR値を0超かつ1.5未満に制御すると、CaSを主体とする複合硫化物が効果的に形成し、フェライト生成核として有効に機能させることができる。なお、ACR値は、好ましくは0.15~1.30の範囲である。より好ましくは、0.20~1.00の範囲である。
上記式の左辺(5.5[C](4/3)+15[P]+0.90[Mn]+0.12[Ni]+7.9[Nb](1/2)+0.53[Mo])は、中心偏析部に濃化しやすい成分で構成された中心偏析部の硬さ指標であり、以下の説明ではCeq*値と称する。
CTOD試験は、鋼板全厚での試験のため、試験片は中心偏析を含み、中心偏析での成分濃化が顕著な場合には、溶接熱影響部に硬化域が生成するので、CTOD試験として、良好な結果が得られない。
そこで、本発明では、Ceq*値を適正範囲に制御することによって、中心偏析部における過度の硬度上昇を抑制し、板厚が厚い鋼板の溶接部においても優れたCTOD特性が得られるのである。
Ceq*値の適正範囲は、実験的に求められたものであり、Ceq*値が3.70を超えるとCTOD特性が低下するので3.70以下とする。好ましくは3.50以下である。なお、Ceq*値の下限に特に制限はないが、生産性の観点などから2.2程度が好ましい。
Cuは添加することで、鋼板強度を向上させることができる。ただし、0.7%を超えての添加は熱間延性を低下させるので、0.7%以下に制限する。好ましくは、0.1~0.6%である。
Crは、鋼板を高強度化するのに有効な元素であり、この効果を発揮するには0.1%以上を含有する。しかし、過剰に含有すると靭性に悪影響を与えるので、含有する場合は0.1~1.0%の範囲が好ましく、0.2~0.8%の範囲であることがより好ましい。
Vは、0.005%以上の含有で鋼板の強度と靭性の向上に有効な元素であるが、含有量が0.05%を超えると靭性低下を招くので、含有する場合は0.005~0.05%であることが好ましい。
本発明における析出物粒子径の求め方は、TEMレプリカ法に準拠する。すなわち、鋼中の、Ti、NbおよびMoの炭化物の析出部を適宜採取したのち、10万倍で4視野による観察から画像処理を用いて平均円相当径を求め、これを析出物の粒子径とする。なお、本発明では、析出物粒径の測定対象の下限値を2nmとする。これ未満の析出物粒径の析出物では、測定が難しくなるからである。
前記した本発明範囲内の鋼板成分に調整した溶鋼を、転炉や、電気炉、真空溶解炉などを用いた通常の方法で溶製し、次いで、連続鋳造の工程を経てスラブとした後、熱間圧延により所望の板厚とし、その後冷却して必要に応じて焼戻し処理を施す。なお、本発明における熱間圧延ではスラブ加熱温度と、圧下率を規定する。
なお、本発明において、特に記載しない限り、鋼板の温度条件は、鋼板の板厚中心部の温度で規定するものとする。板厚中心部の温度は、板厚、表面温度および冷却条件などから、シミュレーション計算などにより求められる。たとえば、差分法を用い、板厚方向の温度分布を計算することによって、板厚中心部の温度を求めることができる。
スラブ加熱温度は、スラブに存在する鋳造欠陥を熱間圧延によって着実に圧着させるため950℃以上とする。一方、スラブを、1150℃を超える温度に加熱するとオーステナイト結晶粒が粗大化して鋼板の靭性が低下するため、加熱温度の上限を1150℃とする。
鋳造欠陥の圧着による無害化と、オーステナイト粒を再結晶により微細なミクロ組織とするために、900℃以上の温度域における熱間圧延の累積圧下率を30%以上とする。30%未満では、加熱時に生成した粗大粒が残存して、鋼板の靭性に悪影響を及ぼすからである。なお、900℃以上の温度域における熱間圧延の累積圧下率の上限は特に限定されないが、工業的には95%程度である。
この温度域で圧延されたオーステナイト粒は十分に再結晶しないため、圧延後のオーステナイト粒は偏平に変形したままで、内部に変形帯などの欠陥を多量に含む内部歪の高い状態となる。これらは、フェライト変態の駆動力として働き、相変態を促進する。
しかし、累積圧下率が30%未満では、内部歪による内部エネルギーの蓄積が十分でないためフェライト変態が起こりにくく鋼板靭性が低下する一方で、累積圧下率が70%を超えると、逆にポリゴナルフェライトの生成が促進されて、高強度と高靭性が両立しない。従って、本発明では、900℃未満の温度域における熱間圧延の累積圧下率を30~70%の範囲とする。
熱間圧延後、冷却速度を1.0℃/s以上として少なくとも500℃まで加速冷却する。冷却速度が1.0℃/s未満では十分な鋼板の強度が得られないからである。また、500℃より高い温度で冷却を停止するとフェライト+パーライト組織の分率が高くなって、厚肉材の高強度と高靭性とが両立しない。なお、加速冷却の停止温度の下限は特に限定されるものではなく、室温まで行っても良い。
本発明で焼戻し処理を行う場合、450℃未満の焼戻し温度では十分な焼戻しの効果が得られない。一方、650℃を超える温度で焼戻しを行うと、析出物が粗大になって靭性が低下したり、強度が低下したりすることもあるため好ましくない。
また、本発明の焼戻し処理は、誘導加熱を用いることにより、焼戻し時の炭化物の粗大化が抑制されるためより好ましい。その場合は、差分法などのシミュレーションによって計算される鋼板の中心温度が450~650℃となるようにすることが望ましい。
なお、本発明において、TMCP鋼板等、鋼板の所望の性能が得られている場合には、上記焼戻し処理を行わなくても良い。
表2に示す成分組成を有する鋼記号A~Zの連続鋳造スラブを素材とした後、表3に示す熱間圧延と熱処理とを行い、厚さが50~150mmの厚鋼板を製造した。鋼板の評価方法として、引張試験は鋼板の板厚の1/2位置より試験片の長手方向が鋼板の圧延方向と垂直になるようにJIS4号試験片を採取し、降伏応力(YS)および引張強さ(TS)を測定した。
また、シャルピー衝撃試験は、鋼板の板厚の1/2位置より試験片の長手方向が鋼板の圧延方向と垂直になるようにJIS4号2mmVノッチ試験片を採取し、−40℃における吸収エネルギーvE−40℃を測定した。なお、本実施例では、YS≧460MPa、TS≧570MPaおよびvE−40℃≧200Jの全てを満たすものを鋼板特性が良好であると評価した。
溶接部靭性の評価は、レ型開先を用いて、溶接入熱35kJ/cmのサブマージアーク溶接による多層盛溶接継手を作製し、鋼板の板厚の1/2位置のストレート側の溶接ボンド部をシャルピー衝撃試験のノッチ位置として、−40℃の温度における吸収エネルギーvE−40℃を測定した。そして、3本の平均がvE−40℃≧150Jを満足するものを溶接部靭性が良好と判断した。
また、ストレート側の溶接ボンド部をCTOD試験片のノッチ位置として、−10℃におけるCTOD値であるδ−10℃を測定し、試験数量3本のうちCTOD値(δ−10℃)の最小値が0.5mm以上である場合を、溶接継手のCTOD特性が良好と判断した。
さらに、鋼中の析出部をTEMレプリカ法により採取し、10万倍4視野による観察から画像処理により平均円相当径を求め、これを析出物サイズとした。また、EDXにより粒子径がほぼ平均に近い析出物を選び、その析出物組成を求め、3個の平均として〔Nb〕/(〔Ti〕+〔Nb〕+〔Mo〕)を求めた。
PWHT後の鋼板特性変化については、ΔTS(=TS(PWHT後)−TS(PWHT前))、ΔvTrs(=vTrs(PWHT後)−vTrs(PWHT前))を求めた。PWHT熱処理は、580℃で4h保持とし、昇温、降温速度を70℃/hとして行った。
表3に、熱間圧延条件、熱処理条件とともに、鋼板特性および上記溶接部のシャルピー衝撃試験結果とCTOD試験結果、析出物サイズ・組成、PWHT後の鋼板特性変化を併記する。
一方、試料No.3、4、7、9、10、12~31は、鋼板成分、製造条件、析出物サイズ・組成の少なくとも一つが本発明の範囲外であり、鋼板特性や、溶接ボンド部のシャルピー衝撃試験結果、溶接ボンド部の三点曲げCTOD試験結果、PWHT特性のいずれかが目標を満足しなかった。なお、表3中、ヨコ線の項目は、当該項目の測定ができなかったことを意味する。
Claims (5)
- 質量%で、
C:0.020~0.090%
Si:0.01~0.35%
Mn:1.40~2.00%
P:0.008%以下
S:0.0035%以下
Al:0.010~0.060%
Ni:0.40~2.00%
Mo:0.05~0.50%
Nb:0.005~0.040%
Ti:0.005~0.025%
N:0.0020~0.0050%
Ca:0.0005~0.0050%および
O:0.0035%以下
を含有し、下記(1)式で規定されるCeqが0.420~0.520%の範囲であって、下記(2)式、(3)式および(4)式を満たすと共に、Bを0.0003%未満に抑制し、残部がFeおよび不可避的不純物からなる鋼板成分と、
Ti、NbおよびMoを、Ti量(〔Ti〕)、Nb量(〔Nb〕)およびMo量(〔Mo〕)が、〔Nb〕/(〔Ti〕+〔Nb〕+〔Mo〕)≧0.3の関係を満足する範囲で含みかつ平均粒子径が20nm以下の析出物を有する鋼板。
記
Ceq=[C]+[Mn]/6+([Cu]+[Ni])/15+([Cr]+[Mo]+[V])/ 5・・・(1)
1.5≦[Ti]/[N]≦4.0 ・・・(2)
0<{[Ca]−(0.18+130×[Ca])×[O]}/1.25/[S]<1.5 ・・・(3)
5.5[C](4/3)+15[P]+0.90[Mn]+0.12[Ni]+7.9[Nb](1/2)+0.53[Mo] ≦3.70 ・・・(4)
但し、[M]は、鋼板中の元素Mの含有量(質量%)を表す。 - 前記鋼板成分に、さらに質量%で、Cu:0.7%未満、Cr:0.1~1.0%およびV:0.005~0.05%のうちから選ばれる1種または2種以上を含有する請求項1に記載の鋼板。
- 前記鋼板成分に、さらに質量%で、Mg:0.0002~0.0050%およびREM:0.0010~0.0200%のうちから選ばれる1種または2種を含有する請求項1または2に記載の鋼板。
- 請求項1~3のいずれか1項に記載の鋼板成分を有する鋼に、950~1150℃に加熱後、900℃以上の温度域における累積圧下率が30%以上、900℃未満の温度域における累積圧下率が30~70%となる熱間圧延を施し、その後、少なくとも500℃までを冷却速度1.0℃/s以上で冷却する鋼板の製造方法。
- 前記冷却後、さらに450~650℃で焼戻し処理を施す請求項4に記載の鋼板の製造方法。
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Citations (11)
Publication number | Priority date | Publication date | Assignee | Title |
---|---|---|---|---|
JPS60184663A (ja) | 1984-02-29 | 1985-09-20 | Kawasaki Steel Corp | 大入熱溶接用低温用高張力鋼 |
JPH0345856B2 (ja) | 1988-02-10 | 1991-07-12 | Hitachi Ltd | |
JPH0353367B2 (ja) | 1984-01-20 | 1991-08-14 | Kawasaki Steel Co | |
JP2001247932A (ja) * | 2000-03-08 | 2001-09-14 | Nippon Steel Corp | 高ctod保証低温用鋼 |
JP3697202B2 (ja) | 2001-11-12 | 2005-09-21 | 新日本製鐵株式会社 | 溶接熱影響部の靭性が優れた鋼及びその製造方法 |
JP2006291349A (ja) * | 2005-03-17 | 2006-10-26 | Jfe Steel Kk | 高変形性能を有するラインパイプ用鋼板およびその製造方法。 |
WO2008078917A1 (en) * | 2006-12-26 | 2008-07-03 | Posco | High strength api-x80 grade steels for spiral pipes with less strength changes and method for manufacturing the same |
JP2009263777A (ja) * | 2008-03-31 | 2009-11-12 | Jfe Steel Corp | 高張力鋼およびその製造方法 |
JP4432905B2 (ja) | 2003-11-27 | 2010-03-17 | 住友金属工業株式会社 | 溶接部靱性に優れた高張力鋼および海洋構造物 |
CN102691015A (zh) * | 2011-03-25 | 2012-09-26 | 宝山钢铁股份有限公司 | 一种低温韧性优良的TMCP型YP500MPa级厚板及其制造方法 |
WO2013118313A1 (ja) * | 2011-02-15 | 2013-08-15 | Jfeスチール株式会社 | 溶接熱影響部の低温靭性に優れた高張力鋼板およびその製造方法 |
Family Cites Families (8)
Publication number | Priority date | Publication date | Assignee | Title |
---|---|---|---|---|
JPH0353367A (ja) | 1989-07-20 | 1991-03-07 | Toshiba Corp | 分散型情報処理システム |
JPH0792360B2 (ja) | 1989-07-10 | 1995-10-09 | 防衛庁技術研究本部長 | 金属燃焼器の製造方法 |
JPH06184663A (ja) | 1992-05-15 | 1994-07-05 | Kobe Steel Ltd | セラミックス強化アルミニウム合金複合材料 |
JP4507669B2 (ja) * | 2004-03-31 | 2010-07-21 | Jfeスチール株式会社 | 溶接部靭性に優れた低温用低降伏比鋼材の製造方法 |
KR100851189B1 (ko) * | 2006-11-02 | 2008-08-08 | 주식회사 포스코 | 저온인성이 우수한 초고강도 라인파이프용 강판 및 그제조방법 |
WO2011096456A1 (ja) * | 2010-02-08 | 2011-08-11 | 新日本製鐵株式会社 | 厚鋼板の製造方法 |
JP5304925B2 (ja) * | 2011-12-27 | 2013-10-02 | Jfeスチール株式会社 | 脆性亀裂伝播停止特性に優れた構造用高強度厚鋼板およびその製造方法 |
JP5516784B2 (ja) * | 2012-03-29 | 2014-06-11 | Jfeスチール株式会社 | 低降伏比高強度鋼板およびその製造方法並びにそれを用いた高強度溶接鋼管 |
-
2014
- 2014-12-10 KR KR1020167016202A patent/KR101846759B1/ko active IP Right Grant
- 2014-12-10 EP EP14869973.9A patent/EP3081662B1/en active Active
- 2014-12-10 US US15/103,093 patent/US20160312327A1/en not_active Abandoned
- 2014-12-10 JP JP2015524556A patent/JP5950045B2/ja active Active
- 2014-12-10 WO PCT/JP2014/083321 patent/WO2015088040A1/ja active Application Filing
- 2014-12-10 CN CN201480067195.XA patent/CN105980588B/zh active Active
Patent Citations (11)
Publication number | Priority date | Publication date | Assignee | Title |
---|---|---|---|---|
JPH0353367B2 (ja) | 1984-01-20 | 1991-08-14 | Kawasaki Steel Co | |
JPS60184663A (ja) | 1984-02-29 | 1985-09-20 | Kawasaki Steel Corp | 大入熱溶接用低温用高張力鋼 |
JPH0345856B2 (ja) | 1988-02-10 | 1991-07-12 | Hitachi Ltd | |
JP2001247932A (ja) * | 2000-03-08 | 2001-09-14 | Nippon Steel Corp | 高ctod保証低温用鋼 |
JP3697202B2 (ja) | 2001-11-12 | 2005-09-21 | 新日本製鐵株式会社 | 溶接熱影響部の靭性が優れた鋼及びその製造方法 |
JP4432905B2 (ja) | 2003-11-27 | 2010-03-17 | 住友金属工業株式会社 | 溶接部靱性に優れた高張力鋼および海洋構造物 |
JP2006291349A (ja) * | 2005-03-17 | 2006-10-26 | Jfe Steel Kk | 高変形性能を有するラインパイプ用鋼板およびその製造方法。 |
WO2008078917A1 (en) * | 2006-12-26 | 2008-07-03 | Posco | High strength api-x80 grade steels for spiral pipes with less strength changes and method for manufacturing the same |
JP2009263777A (ja) * | 2008-03-31 | 2009-11-12 | Jfe Steel Corp | 高張力鋼およびその製造方法 |
WO2013118313A1 (ja) * | 2011-02-15 | 2013-08-15 | Jfeスチール株式会社 | 溶接熱影響部の低温靭性に優れた高張力鋼板およびその製造方法 |
CN102691015A (zh) * | 2011-03-25 | 2012-09-26 | 宝山钢铁股份有限公司 | 一种低温韧性优良的TMCP型YP500MPa级厚板及其制造方法 |
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US20170137905A1 (en) * | 2014-03-31 | 2017-05-18 | Jfe Steel Corporation | High-tensile-strength steel plate and process for producing same |
US10316385B2 (en) * | 2014-03-31 | 2019-06-11 | Jfe Steel Corporation | High-tensile-strength steel plate and process for producing same |
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JP2019183205A (ja) * | 2018-04-05 | 2019-10-24 | Jfeスチール株式会社 | 鋼板およびその製造方法 |
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JP5950045B2 (ja) | 2016-07-13 |
US20160312327A1 (en) | 2016-10-27 |
EP3081662A4 (en) | 2016-12-07 |
EP3081662A1 (en) | 2016-10-19 |
WO2015088040A8 (ja) | 2016-05-06 |
KR101846759B1 (ko) | 2018-04-06 |
JPWO2015088040A1 (ja) | 2017-03-16 |
KR20160088375A (ko) | 2016-07-25 |
EP3081662B1 (en) | 2019-11-13 |
CN105980588A (zh) | 2016-09-28 |
CN105980588B (zh) | 2018-04-27 |
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