WO2014092129A1

WO2014092129A1 - Thick steel plate having excellent cryogenic toughness

Info

Publication number: WO2014092129A1
Application number: PCT/JP2013/083239
Authority: WO
Inventors: 朗伊庭野; 秀徳名古
Original assignee: 株式会社神戸製鋼所
Priority date: 2012-12-13
Filing date: 2013-12-11
Publication date: 2014-06-19
Also published as: CN104854252B; EP2933347A1; CN104854252A; EP3190201A1; KR20150082611A; KR101711774B1; EP2933347A4

Abstract

This thick steel sheet contains prescribed steel components, the Di value composed of the steel components is 2.5 or greater, the residual austenite phase (residual (γ)) remaining at -196°C is 2.0 to 12.0% by volume fraction, and the residual (γ) stability parameter composed of the components contained in the residual austenite at -196°C satisfies the condition of 3.1 or greater.

Description

Thick steel plate with excellent cryogenic toughness

The present invention relates to a thick steel plate having excellent cryogenic toughness. Specifically, according to the present invention, even when the Ni content is reduced to about 5.0 to 7.5%, toughness at an extremely low temperature of −196 ° C. or less [particularly, toughness in the plate width direction (C direction)] Relates to a good thick steel plate. In the following, the explanation will focus on thick steel plates (typically storage tanks, transport ships, etc.) for liquefied natural gas (LNG) exposed to the above-mentioned cryogenic temperatures. The present invention is not intended to be limited, and is applied to all thick steel plates used for applications exposed to extremely low temperatures of −196 ° C. or less.

LNG tank steel plates used in LNG storage tanks are required to have high strength and high toughness that can withstand extremely low temperatures of -196 ° C. In general, it is known that the addition of Ni improves the hardness-toughness balance particularly at low temperatures. So far, thick steel plates containing about 9% Ni (9% Ni steel) have been used as the thick steel plates used in the above applications. However, since the cost of Ni has increased in recent years, the development of a thick steel plate excellent in cryogenic toughness even with a low Ni content of less than 9% has been underway.

For example, Non-Patent Document 1 describes the effect of α-γ2 phase coexistence heat treatment on the low temperature toughness of 6% Ni steel. More specifically, by applying heat treatment (L treatment) in the α-γ2 phase coexistence region (between A _c1 and A _c3 ) before tempering treatment, a large amount of fine and extremely low temperature impact loads are applied. Further, it is described that stable retained austenite is generated, and cryogenic toughness at −196 ° C. which is equal to or higher than that of 9% Ni steel subjected to normal quenching and tempering treatment can be secured. However, although the cryogenic toughness in the rolling direction (L direction) is excellent, the cryogenic toughness in the sheet width direction (C direction) generally tends to be inferior to that in the L direction. There is no description of the brittle fracture surface ratio.

A technique similar to that of Non-Patent Document 1 is described in Patent Document 1 and Patent Document 2. Of these, Patent Document 1 discloses that steel containing 4.0 to 10% Ni and whose austenite grain size is controlled within a predetermined range is hot-rolled and then heated between A _c1 and A _c3. A method is described in which a cooling treatment (corresponding to the L treatment described in Non-Patent Document 1 above) is repeated once or twice or more and then _{tempered at} a temperature not higher than the _Ac1 transformation point. Patent Document 2 discloses that the steel containing 4.0 to 10% Ni and the size of AlN before hot rolling is 1 μm or less is similar to the heat treatment (L treatment → firing) described above. A method for performing the return processing) is described. The impact value at −196 ° C. (vE ₋₁₉₆ ) described in these methods is presumed to be probably in the L direction, and the toughness value in the C direction is unknown. In these methods, strength is not taken into consideration, and there is no description of the brittle fracture surface ratio.

Further, Non-Patent Document 2 describes the development of 6% Ni steel for LNG tanks combining the above-mentioned L treatment (two-phase quenching treatment) and TMCP. According to this document, although it is described that the toughness in the rolling direction (L direction) shows a high value, the toughness value in the sheet width direction (C direction) is not described.

On the other hand, in Patent Document 3, a steel sheet having a yield strength at room temperature of 590 MPa or more in a Ni steel of more than 5.0% and less than 8.0% is assumed to be as resistant as 9% Ni steel even in a use environment. It describes a Ni-reduced steel plate for low temperature with excellent fracture safety and a method for producing the same. In patent document 3, if the yield point in the low temperature environment which is use temperature can be raised reliably, fracture safety will be improved (that is, high toughness can be obtained in a low temperature environment). In the heating process, the steel ingot is heated at a low temperature in a short time, and in the rolling process, the thickness of the steel ingot at the end of the rough rolling is the product thickness (thickness after finish rolling). The steel sheet is reduced to 3 to 8 times the thickness of the steel sheet. Further, in the examples, rolling is performed from a slab thickness of 300 mm to a finishing thickness of 50 mm or less (mostly to a finishing thickness of less than 50 mm). Thus, by securing a relatively high reduction ratio, the residual γ fraction and the fineness are reduced. Combined with a matrix structure, it achieves low temperature toughness comparable to 9% Ni steel. However, the TS at room temperature of the thick steel plate of Patent Document 3 is 741 MPa at the maximum.

In Patent Document 3, although the absorbed energy in the C direction is described, there is no description of the brittle fracture surface ratio. Moreover, TS at room temperature in Patent Document 3 is about 741 MPa at the maximum.

Japanese Unexamined Patent Publication No. 49-13581 Japanese Laid-Open Patent Publication No. 51-13308 Japanese Unexamined Patent Publication No. 2011-241419

As described above, until now, a technology excellent in cryogenic toughness at −196 ° C. has been proposed for Ni steel having a Ni content of about 5.0 to 7.5%, but the cryogenic temperature in the C direction has been proposed. Toughness has not been fully studied. In addition, if it is possible to increase the strength, it is useful in that the design margin can be increased, but a technology that has high strength and excellent cryogenic toughness has not been provided.

In addition, none of the above-mentioned documents has been studied on the brittle fracture surface ratio. The brittle fracture surface ratio indicates the ratio of brittle fracture that occurs when a load is applied in the Charpy impact test. At the site where brittle fracture occurs, the energy absorbed by the steel material until the fracture is significantly reduced, and the fracture proceeds easily. Therefore, in the cryogenic toughness improvement technology, the general-purpose Charpy impact value (vE _{In addition to} the improvement of _-196 ), the brittle fracture surface ratio should be 10% or less, which is an extremely important requirement. However, a technique that satisfies the above requirement for the brittle fracture surface ratio in a high-strength thick steel plate having a high base metal strength as described above has not yet been proposed.

The present invention has been made in view of the above circumstances, and its purpose is to achieve cryogenic toughness at −196 ° C. (particularly in the C direction) in Ni steel having a Ni content of about 5.0 to 7.5%. An object of the present invention is to provide a high-strength thick steel plate that is excellent in low-temperature toughness and can realize a brittle fracture surface ratio ≦ 10%, and a method for producing the same.

The thick steel plate having excellent cryogenic toughness according to the present invention that can solve the above problems is mass%, C: 0.02 to 0.10%, Si: 0.40% or less (excluding 0%) Mn: 0.50 to 2.0%, P: 0.007% or less (not including 0%), S: 0.007% or less (not including 0%), Al: 0.005 to 0. 050%, Ni: 5.0 to 7.5%, N: 0.010% or less (not including 0%), Cr: 1.20% or less (not including 0%), and Mo : 1.0% or less (excluding 0%), containing at least one element selected from the group consisting of steel and inevitable impurities, and the remainder is a thick steel plate composed of steel components (1) The Di value determined based on the formula is 2.5 or more,
Di value = ([C] / 10) ^0.5 × (1 + 0.7 × [Si]) × (1 + 3.33 × [Mn]) × (1 + 0.35 × [Cu]) × (1 + 0.36 × [ Ni]) × (1 + 2.16 × [Cr]) × (1 + 3 × [Mo]) × (1 + 1.75 × [V]) × 1.115 (1)
(In the formula, [] means the content (% by mass) of each component in the steel)
The residual austenite phase (residual γ) present at −196 ° C. is 2.0 to 12.0% in volume fraction, and
The gist is that the residual γ stabilization parameter determined based on the following formula (2) composed of the components contained in the retained austenite is 3.1 or more.
Residual γ stabilization parameter =
(365 × <C> + 39 × <Mn> + 30 × <Al> + 10 × <Cu> + 17 × <Ni> + 20 × <Cr> + 5 × <Mo> + 35 × <V>) / 100 (2)
(In the formula, <> means the content (% by mass) of each component contained in the retained austenite existing at -196 ° C.)

In a preferred embodiment of the present invention, the volume fraction of the residual γ phase and the residual γ stability calculated based on the following equation (3) composed of the volume fraction of the residual γ phase and the residual γ stabilization parameter: The activation parameter is 40 or less.
Volume fraction of residual γ / residual γ stabilization parameter = 10 / (volume fraction of residual γ phase × residual γ stabilization parameter) ^1/2 (3)

In addition, the Mn concentration in the residual γ is controlled in place of the residual γ stabilization parameter after defining the element content, Di value, and residual γ volume fraction in the above-mentioned thick steel plate within a more limited range. This is also a preferred embodiment of the present invention because it is possible to exhibit extremely low temperature toughness after obtaining a higher base metal strength.

Specifically, in terms of mass%, C: 0.02 to 0.10%, Si: 0.40% or less (not including 0%), Mn: 0.6 to 2.0%, P: 0.0. 007% or less (excluding 0%), S: 0.007% or less (not including 0%), Al: 0.005 to 0.050%, Ni: 5.0 to 7.5%, N: Contains 0.010% or less (excluding 0%), Mo: 0.30 to 1.0%, Cr: 1.20% or less (excluding 0%), the balance being iron and inevitable impurities A thick steel plate,
The Di value determined on the basis of the formula (1) composed of components in steel is more than 5.0,
The residual austenite phase (residual γ) present at −196 ° C. is 2.0 to 5.0% in volume fraction,
The Mn concentration in the residual austenite phase (residual γ) existing at −196 ° C. is 1.05% or more, and the Mn and Ni contents (mass%) in the steel satisfy the following formula (4). Is a thick steel plate with excellent cryogenic toughness.
[Mn] ≧ 0.31 × (7.20− [Ni]) + 0.50 (4)
(In the formula, [] means the content (mass%) of each component in the steel.)

In a preferred embodiment of the present invention, the steel sheet further contains Cu: 1.0% or less (excluding 0%).

In a preferred embodiment of the present invention, the steel sheet further includes Ti: 0.025% or less (excluding 0%), Nb: 0.100% or less (not including 0%), and V: 0.50. % Or less (not including 0%).

In a preferred embodiment of the present invention, the steel sheet further contains B: 0.0050% or less (excluding 0%).

In a preferred embodiment of the present invention, the steel sheet is further selected from the group consisting of Ca: 0.0030% or less (excluding 0%) and REM: 0.0050% or less (excluding 0%). Containing at least one kind.

In a preferred embodiment of the present invention, the steel plate further contains Zr: 0.005% or less (excluding 0%).

In addition, the method for producing a thick steel plate according to the present invention according to claim 1 or 2 that has solved the above-described problem is a heat treatment (L treatment) in an α-γ2 phase coexistence region (between A _c1 and A _c3 ). The L parameter calculated based on the following formula (5) composed of the temperature (L treatment temperature) and A _c1 and A _c3 in the steel is 0.25 or more and 0.45 or less, and Adjusting the L treatment temperature and the steel component so that the λ _L parameter calculated based on the following formula (6) composed of the L parameter and the steel component is 7 or less: , L treatment, water cooling to room temperature, and tempering treatment (T treatment) are carried out for 10 to 60 minutes at a temperature of _Ac1 or lower.
L parameter = (L treatment temperature−A _c1 ) / (A _c3 −A _c1 ) +0.25 (5)
λ _L parameter = 9.05 × (0.90 × [L parameter] +0.14) × [Mn] + 1.46 × (0.37 × [L parameter] +0.67) × [Cr] −41.5 × (0.26 × [L parameter] +0.79) × [Mo] (6)
(In formula, [] means content (mass%) of each component in steel.)

Furthermore, in the method for producing a thick steel plate according to the present invention according to claim 3, which has solved the above problem, the L parameter calculated based on the formula (5) is 0.6 or more and 1.1 or less. In addition, there is a feature in that the L treatment temperature and the steel components are adjusted so that the λ _L parameter calculated based on the equation (6) is 0 or less.

According to the present invention, in Ni steel having a Ni content of about 5.0 to 7.5%, even if the base metal strength is high (specifically, tensile strength TS> 741 MPa, yield strength YS> 590 MPa, preferably TS ≧ 830 MPa, YS ≧ 690 MPa), excellent in cryogenic toughness at −196 ° C. or less (particularly, cryogenic toughness in the C direction), brittle fracture surface ratio at −196 ° C. ≦ 10% (preferably − It was possible to provide a high-strength thick steel plate that satisfies the brittle fracture surface ratio ≦ 50% at 233 ° C.

The inventors of the present invention have a Ni content of 7.5% or less, and when performing a Charpy impact absorption test in the C direction, the brittle fracture surface ratio at −196 ° C. is 10% or less, the tensile strength TS> 741 MPa, In order to provide a thick steel plate that satisfies the yield strength YS> 590 MPa, investigations were made.

In particular, in the present invention, the following points were examined.

First, regarding the manufacturing method, in the present invention, as in Patent Documents 1 and 3, cryogenic toughness similar to that of 9% Ni steel is achieved without strict management such as cooling after rolling and T treatment. It was assumed that. Specifically, component design is performed in consideration of the case where the rolling reduction as high as that of Patent Document 3 cannot be ensured. For rolling, a rolling reduction of 830 ° C. or higher is approximately 50% or lower, and a rolling reduction of 700 ° C. or higher is approximately 85. It was presupposed that the water cooling after tempering (T treatment) after hot rolling was not performed (that is, air cooling was performed after T treatment). The reduction ratio (%) was calculated by 100 × (thickness before rolling−thickness after rolling) / (thickness before rolling).

Also, for cryogenic toughness, evaluation in the C direction, which tends to be more difficult to secure toughness than in the L direction, was adopted, and from the viewpoint of ensuring toughness, the evaluation was based on the fracture surface ratio instead of absorbed energy. Further, regarding the tensile strength (TS), in the design of a cryogenic pressure vessel, in consideration of safety, from the viewpoint that TS should be higher if it is within the standard range, in the present invention, TS> 741 MPa is assumed. did.

Specifically, based on the above manufacturing conditions, in the Charpy impact absorption test in the C direction, the brittle fracture surface ratio at −196 ° C. ≦ 10%, the tensile strength TS> 741 MPa, and the yield strength YS> 590 MPa are satisfied. Consideration has been repeated to provide a thick steel plate.

As a result, the residual γ form the following (A), (b), and lambda _L Parameters following (c)], by the following control, in Charpy impact absorption test, not transformed into martensite It was found that a stable residual γ that is plastically deformed can be ensured, and that excellent cryogenic toughness can be obtained.
(A) From the viewpoint of plastic deformation without being transformed into martensite during impact at extremely low temperature, and securing stable residual γ useful for improving toughness (enhancing the stability of residual γ), The Di value [see the above formula (1)] is controlled by an appropriate balance of the above.
(B) Balance between the steel components and the temperature (L treatment temperature) in the heat treatment (L treatment) in the α-γ2 phase coexistence region (between A _c1 and A _c3 ) by the L parameter [see the above formula (5)] After the L treatment, water cooling to room temperature, tempering treatment (T treatment) under predetermined conditions, followed by air cooling, the volume fraction of residual austenite (residual γ) existing at −196 ° C. is 2 The residual γ stabilization parameter [see the above formula (2)], which is controlled within the range of 0 to 12.0% and is determined by the components in the residual γ existing at −196 ° C., is set to 3.1 or more. Controlling (preferably, the volume fraction of residual γ and the residual γ stabilization parameter [refer to the above equation (3)] composed of the volume fraction of residual γ and the above-mentioned residual γ stabilization parameter are set to 40 or less. To control).
(C) Control the Λ _l parameter determined by the component (Mn, Cr, Mo) and the L processing temperature as shown in the above equation (5).

That is, the thick steel plate of the present invention is, by mass%, C: 0.02 to 0.10%, Si: 0.40% or less (excluding 0%), Mn: 0.50 to 2.0%, P: 0.007% or less (not including 0%), S: 0.007% or less (not including 0%), Al: 0.005 to 0.050%, Ni: 5.0 to 7.5 %, N: 0.010% or less (not including 0%), Cr: 1.20% or less (not including 0%), and Mo: 1.0% or less (not including 0%) ) Is a thick steel plate containing at least one element selected from the group consisting of iron and inevitable impurities, and the Di value determined on the basis of the following formula (1) composed of steel components Is a residual austenite phase (residual γ) present at −196 ° C. in a volume fraction of 2.0 to 12 0% and, is characterized in residual γ stabilization parameter determined based on the following equation (2) consists of components contained in the residual austenite is 3.1 or more.
Di value = ([C] / 10) ^0.5 × (1 + 0.7 × [Si]) × (1 + 3.33 × [Mn]) × (1 + 0.35 × [Cu]) × (1 + 0.36 × [ Ni]) × (1 + 2.16 × [Cr]) × (1 + 3 × [Mo]) × (1 + 1.75 × [V]) × 1.115 (1)
(In the formula, [] means the content (% by mass) of each component in the steel)
Residual γ stabilization parameter =
(365 × <C> + 39 × <Mn> + 30 × <Al> + 10 × <Cu> + 17 × <Ni> + 20 × <Cr> + 5 × <Mo> + 35 × <V>) / 100 (2)
(In the formula, <> means the content (% by mass) of each component contained in the retained austenite existing at -196 ° C.)

1. Components in steel First, components in steel will be described.

C: 0.02 to 0.10%
C is an element essential for securing strength and retained austenite. In order to effectively exhibit such an action, the lower limit of the C amount is set to 0.02% or more. The minimum with the preferable amount of C is 0.03% or more, More preferably, it is 0.04% or more. However, if added excessively, the cryogenic toughness decreases due to an excessive increase in strength, so the upper limit is made 0.10%. The upper limit with preferable C amount is 0.08% or less, More preferably, it is 0.06% or less.

Si: 0.40% or less (excluding 0%)
Si is an element useful as a deoxidizer. However, if added in excess, the formation of a hard island-like martensite phase is promoted and the cryogenic toughness decreases, so the upper limit is made 0.40% or less. The upper limit with the preferable amount of Si is 0.35% or less, More preferably, it is 0.20% or less.

Mn: 0.50 to 2.0%
Mn is an austenite (γ) stabilizing element and is an element contributing to an increase in the amount of residual γ. In order to effectively exhibit such an action, the lower limit of the amount of Mn is set to 0.50%. The minimum with the preferable amount of Mn is 0.6% or more, More preferably, it is 0.7% or more. However, if added excessively, temper embrittlement occurs and the desired cryogenic toughness cannot be secured, so the upper limit is made 2.0% or less. The upper limit with the preferable amount of Mn is 1.5% or less, More preferably, it is 1.3% or less.

P: 0.007% or less (excluding 0%)
P is an impurity element causing grain boundary fracture, and its upper limit is made 0.007% or less in order to secure the desired cryogenic toughness. The upper limit with preferable P amount is 0.005% or less. The smaller the amount of P, the better. However, it is difficult to make the amount of P 0% industrially.

S: 0.007% or less (excluding 0%)
S, like P, is an impurity element causing grain boundary fracture, and its upper limit is made 0.007% or less in order to ensure the desired cryogenic toughness. As shown in the examples described later, when the amount of S increases, the brittle fracture surface ratio increases and the desired cryogenic toughness (the brittle fracture surface ratio at −196 ° C. ≦ 10%) cannot be realized. The upper limit with the preferable amount of S is 0.005% or less. The smaller the amount of S, the better. However, it is difficult to make the amount of S 0% industrially.

Al: 0.005 to 0.050%
Al is an element that promotes desulfurization and fixes nitrogen. If the Al content is insufficient, the concentration of solute sulfur, solute nitrogen, etc. in the steel increases and the cryogenic toughness decreases, so the lower limit is made 0.005% or more. The minimum with the preferable amount of Al is 0.010% or more, More preferably, it is 0.015% or more. However, if added excessively, oxides, nitrides, and the like are coarsened and the cryogenic toughness is also lowered, so the upper limit is made 0.050% or less. The upper limit with the preferable amount of Al is 0.045% or less, More preferably, it is 0.04% or less.

Ni: 5.0 to 7.5%
Ni is an essential element for securing retained austenite (residual γ) useful for improving cryogenic toughness. In order to effectively exhibit such an action, the lower limit of the Ni amount is set to 5.0% or more. A preferable lower limit of the Ni amount is 5.2% or more, and more preferably 5.4% or more. However, if added excessively, the cost of the raw material is increased, so the upper limit is made 7.5% or less. The upper limit of the Ni content is preferably 7.0% or less, more preferably 6.5% or less, still more preferably 6.2% or less, and even more preferably 6.0% or less.

N: 0.010% or less (excluding 0%)
N lowers the cryogenic toughness by strain aging, so the upper limit is made 0.010% or less. The upper limit with preferable N amount is 0.006% or less, More preferably, it is 0.004% or less.

Cr: at least one selected from the group consisting of 1.20% or less (not including 0%) and Mo: 1.0% or less (not including 0%) Cr and Mo are both strength-enhancing elements. is there. These elements may be added alone or in combination of two kinds. In order to effectively exhibit the above action, the Cr content is 0.05% or more and the Mo content is 0.01% or more. However, if excessively added, the strength is excessively increased and the desired cryogenic toughness cannot be ensured, so the upper limit of the Cr content is 1.20% or less (preferably 1.1% or less, more preferably 0). 0.9% or less, more preferably 0.5% or less) and the upper limit of the Mo amount is 1.0% or less (preferably 0.8% or less, more preferably 0.6% or less).

The thick steel plate of the present invention contains the above components as basic components, the balance: iron and unavoidable impurities.

In the present invention, the following selective components can be contained for the purpose of imparting further properties.

Cu: 1.0% or less (excluding 0%)
Cu is a γ-stabilizing element and is an element that contributes to an increase in the amount of residual γ. In order to exhibit such an action effectively, it is preferable to contain 0.05% or more of Cu. However, if added excessively, the strength is excessively improved and the desired cryogenic toughness effect cannot be obtained. Therefore, the upper limit is preferably made 1.0% or less. A more preferable upper limit of the amount of Cu is 0.8% or less, and even more preferably 0.7% or less.

Selected from the group consisting of Ti: 0.025% or less (not including 0%), Nb: 0.100% or less (not including 0%), and V: 0.50% or less (not including 0%) At least one of Ti, Nb, and V is an element that precipitates as carbonitride and increases strength. These elements may be added alone or in combination of two or more. In order to effectively exhibit the above action, it is preferable that the Ti amount is 0.005% or more, the Nb amount is 0.005% or more, and the V amount is 0.005% or more. However, if excessively added, the strength is excessively improved, and the desired cryogenic toughness cannot be ensured. Therefore, the preferable upper limit of Ti amount is 0.025% or less (more preferably 0.018% or less, More preferably 0.015% or less), a preferable upper limit of Nb amount is 0.100% or less (more preferably 0.05% or less, further preferably 0.02% or less), and a preferable upper limit of V amount is 0. .50% or less (more preferably 0.3% or less, and still more preferably 0.2% or less).

B: 0.0050% or less (excluding 0%)
B is an element that contributes to improving strength by improving hardenability. In order to effectively exhibit the above action, the B content is preferably 0.0005% or more. However, if added excessively, the strength is excessively improved and the desired cryogenic toughness cannot be ensured, so the preferable upper limit of the B amount is 0.0050% or less (more preferably 0.0030% or less, more preferably Is 0.0020% or less).

Ca: 0.0030% or less (excluding 0%) and REM (rare earth element): at least one selected from the group consisting of 0.0050% or less (excluding 0%) Ca and REM are solid It is an element that fixes dissolved sulfur and renders sulfides harmless. These elements may be added alone or in combination of two or more. If these contents are insufficient, the solid solution sulfur concentration in the steel increases and the toughness decreases, so the Ca content is 0.0005% or more, the REM content (when the REM described below is contained alone) It is a single content, and when it contains two or more kinds, it is the total amount thereof. However, if excessively added, sulfides, oxides, nitrides, etc. become coarse and the toughness is also lowered, so the preferable upper limit of Ca content is 0.0030% or less (more preferably 0.0025% or less), REM A preferable upper limit of the amount is 0.0050% or less (more preferably 0.0040% or less).

In this specification, REM (rare earth element) means addition of Sc (scandium) and Y (yttrium) to a lanthanoid element (15 elements from La with atomic number 57 to Lu with atomic number 71 in the periodic table). These elements can be used alone or in combination of two or more. Preferred rare earth elements are Ce and La. The addition form of REM is not particularly limited, and may be added in the form of a misch metal mainly containing Ce and La (for example, Ce: about 70%, La: about 20-30%), or Ce, La alone may be added.

Zr: 0.005% or less (excluding 0%)
Zr is an element that fixes nitrogen. If the Zr content is insufficient, the solid solution N concentration in the steel increases and the toughness decreases, so the Zr content is preferably 0.0005% or more. However, if added excessively, oxides, nitrides, etc. become coarse and the toughness also decreases, so the preferable upper limit of the amount of Zr is made 0.005% or less (more preferably 0.0040% or less).

As above, the components in the steel of the present invention have been described.

2. Volume fraction of residual austenite phase (residual γ) Further, in the thick steel plate of the present invention, the residual γ phase present at −196 ° C. is 2.0 to 12.0% (preferably 4.0 to 12%). 0.0%).

In order to improve the cryogenic toughness, it is effective to secure a residual γ that is easily plastically deformed during an impact test at a low temperature. In order to obtain the desired cryogenic toughness, the volume fraction of the residual γ phase in the entire structure existing at −196 ° C. is set to 2.0% or more. However, the residual γ is relatively soft compared to the matrix phase, and if the residual γ amount becomes excessive, YS cannot secure a predetermined value, so the upper limit is made 12.0%. Regarding the volume fraction of the residual γ phase, the preferable lower limit is 4.0% or more, the more preferable lower limit is 6.0% or more, the preferable upper limit is 11.5% or less, and the more preferable upper limit is 11.0% or less. .

In the thick steel sheet of the present invention, it is important to control the volume fraction of the residual γ phase among the structures existing at −196 ° C., and the structure other than the residual γ is not limited at all. Any material that normally exists in thick steel plates may be used. Examples of the structure other than the residual γ include carbides such as bainite, martensite, and cementite.

3. About Di Value Furthermore, in the present invention, the Di value determined based on the following formula (1) composed of steel components satisfies 2.5 or more.
Di value = ([C] / 10) ^0.5 × (1 + 0.7 × [Si]) × (1 + 3.33 × [Mn]) × (1 + 0.35 × [Cu]) × (1 + 0.36 × [ Ni]) × (1 + 2.16 × [Cr]) × (1 + 3 × [Mo]) × (1 + 1.75 × [V]) × 1.115 (1)
(In formula, [] means content (mass%) of each component in steel.)

The above formula (1) relating to the hardenability Di value is described as the Grossmann formula (Trans. Metal. Soc. AIME, 150 (1942), p. 227). The larger the amount of the alloy element that constitutes the Di value, the easier the firing (the Di value increases) and the finer the structure becomes. Also, the greater the Di value, the higher the strength, and it becomes easier to ensure the desired strength. According to the examination results of the present inventors, there is a correlation between the Di value and the structure size after rolling, and in order to make the structure after rolling fine and secure a desired high strength, the Di value is 2.5. It turned out that it should have done above. In detail, the Di value is such that a fine rolled structure can be obtained even when the rolling reduction of the amorphous region is small, and a sufficient volume fraction of residual γ useful for improving the cryogenic toughness is ensured by the subsequent heat treatment, and the stable residual This parameter is useful as a guideline for securing γ. In addition, the manufacturing conditions described in Patent Document 3 [reduction of reduction rate at low temperature (non-recrystallized region), time limit until the start of cooling, etc.] are relaxed to ensure good characteristics even if the process load is reduced. This is an effective parameter.

In order to effectively exhibit such an action, the Di value is set to 2.5 or more. When the Di value is less than 2.5, a fine structure cannot be sufficiently obtained after rolling, and thus a predetermined amount of residual γ cannot be obtained. Furthermore, since the residual γ stabilization parameters, volume fraction of residual γ and residual γ stabilization parameters described later cannot be controlled to a predetermined level, a stable residual γ structure cannot be obtained, and the desired cryogenic toughness is ensured. Can not. A preferable range of the Di value is 3.0 or more. On the other hand, the upper limit of the Di value is not particularly limited from the above viewpoint, but considering the viewpoint of cost and the like, and the strength standard range of the current LNG tank steel is 830 MPa or less, it is generally 5.0 or less. Preferably there is.

4). Residual γ stabilization parameter Further, in the present invention, the residual γ stabilization parameter determined based on the following equation (2) is controlled to 3.1 or more. Thereby, desired cryogenic toughness is obtained.
Residual γ stabilization parameter =
(365 × <C> + 39 × <Mn> + 30 × <Al> + 10 × <Cu> + 17 × <Ni> + 20 × <Cr> + 5 × <Mo> + 35 × <V>) / 100 (2)
(In the formula, <> means the content (% by mass) of each component contained in the retained austenite existing at -196 ° C.)

As described above, to improve the cryogenic toughness, it is effective to secure a stable residual γ that undergoes plastic deformation without being transformed into martensite during the impact test. For this purpose, it is conceivable to secure the residual γ fraction before the impact test and to increase the stability of the residual γ so that it can be plastically deformed without being transformed into martensite even when subjected to an impact. In the present invention, from the former viewpoint, the volume fraction of residual γ is defined in the above range. Furthermore, when the experiment was conducted also from the latter viewpoint, the stability of the residual γ existing at −196 ° C. was determined by the component in the residual γ existing at −196 ° C., and the parameter represented by the above equation (2) It was found effective to control. When the amount of Ni is reduced to 7.5% or less as in the present invention, the hardenability generally decreases, so the structure after rolling becomes coarse and the volume fraction of residual γ obtained after heat treatment or the above-mentioned Di value is ensured. In the present invention, these requirements are also appropriately controlled by appropriately controlling the residual γ stabilization parameter determined in consideration of the component balance in the residual γ. This residual γ stabilization parameter is derived with reference to the Ms point equation.

In order to ensure the desired cryogenic toughness, the lower limit of the residual γ stabilization parameter is set to 3.1 or more. Preferably it is 3.3 or more, More preferably, it is 3.5 or more, More preferably, it is 3.7 or more. The upper limit of the residual γ stabilization parameter is not particularly limited from the viewpoint of improving cryogenic toughness.

5. Residual γ volume fraction / residual γ stabilization parameter Further, in the present invention, for the purpose of securing a further excellent cryogenic toughness, the volume fraction of residual γ and the residual γ calculated based on the following formula (3): It is preferable to control the stabilization parameter to 40 or less.
Volume fraction of residual γ / residual γ stabilization parameter = 10 / (volume fraction of residual γ phase × residual γ stabilization parameter) ^1/2 (3)

As shown in the above equation (3), the above parameters are composed of a residual γ volume fraction and a residual γ stabilization parameter. The inventors of the present invention have determined the above parameters, considering that the improvement of the cryogenic toughness is due to the plastic deformation during the cryogenic impact test, and that the distribution of residual γ effective for the improvement of the toughness greatly affects. That is, those having a high volume fraction of residual γ and a large residual γ stabilization parameter are those in which the distance between each residual γ is short and finely dispersed, and they are martensite even at low temperatures. Since it is not transformed into a material and is responsible for plastic deformation, it exhibits good cryogenic toughness.

The volume fraction of residual γ and the residual γ stabilization parameter are preferably 35 or less, and more preferably 30 or less. From the viewpoint of improving the cryogenic toughness, the lower the parameter, the better. The lower limit of the above parameter is not particularly limited in relation to the cryogenic toughness, but is generally 10 or more in consideration of the component system of the present invention.

Furthermore, as demonstrated in Example 2 to be described later, by controlling the volume fraction of residual γ and the stabilization parameter of residual γ to a more appropriate range, at a temperature of −233 ° C., which is lower than −196 ° C. described above. However, the brittle fracture surface ratio can be maintained at a favorable level of 50% or less. Specifically, by reducing the upper limit of the volume fraction of residual γ and the residual γ stabilization parameter as much as possible (generally, 30 or less), the brittle fracture surface ratio at −233 ° C. is reduced to 50% or less. Can do.

With the above-described configuration of the present invention, it is possible to obtain a thick steel plate satisfying the brittle fracture surface ratio ≦ −10% at −196 ° C., tensile strength TS> 741 MPa, and yield strength YS> 590 MPa in the C direction Charpy impact absorption test. However, the present inventors have further studied whether it is possible to obtain a thick steel plate that exhibits higher base metal strength and also satisfies cryogenic toughness.

Specifically, in the Charpy impact absorption test in the C direction, we examined as a further goal to provide a thick steel plate that satisfies a brittle fracture surface ratio of 10% or less at −196 ° C., TS> 830 MPa, and YS> 690 MPa. I did it. As a result, after defining the element content, the Di value, and the residual γ volume fraction in the above-described thick steel plate to a more limited range, the Mn concentration in the residual γ is set in place of the residual γ stabilization parameter. It has been found that by controlling, the cryogenic toughness can be satisfied while exhibiting higher base material strength as described above.

Here, the content of the element in the thick steel plate, the Di value, and the limitation to the more limited range of the residual γ volume fraction are:
(A) The lower limit of the Mn content is 0.6%.
(B) Both Cr and Mo elements are essential additions, and the lower limit of Mo is 0.30%.
(C) The Di value is more than 5.0.
(D) The upper limit of the residual γ volume fraction existing at −196 ° C. is 5.0%.
(E) Properly control the balance of the Mn—Ni content in the steel represented by the following formula (4): [Mn] ≧ 0.31 × (7.20− [Ni]) + 0.50・ (4)
(In the formula, [] means the content (mass%) of each component in the steel.)
Means that.

In addition to these limitations,
(F) If the Mn concentration in residual γ existing at −196 ° C. is 1.05% or more, the brittle fracture surface ratio at −196 ° C. is 10 without controlling the residual γ stabilization parameter described above. % And below, TS> 830MPa, YS> 690MPa can be provided.

Hereinafter, a steel plate according to the present invention that has already been described, that is, a steel plate that satisfies the brittle fracture surface ratio ≦ 10% at −196 ° C., TS> 741 MPa, and YS> 590 MPa in the C direction Charpy impact absorption test. The above (a) to (f) having different configurations will be described.

(A) Lower limit of Mn content: 0.6%
From the viewpoint of further improving the strength of TS> 830 MPa and YS> 690 MPa and securing the Mn concentration in the residual γ, the Mn content is set to 0.6% or more. A preferable lower limit of the amount of Mn is 0.7% or more.

(B) Essential addition of both Cr and Mo elements, and lower limit of Mo content: 0.30%
(C) Di value: more than 5.0 From the viewpoint of further improving the strength of TS> 830 MPa and YS> 690 MPa, both Cr and Mo are essential additions, and Mo is contained at 0.30% or more. To do. The Di value is also greater than 5.0.

(D) Upper limit of residual γ volume fraction present at −196 ° C .: 5.0%
From the viewpoint of improving the cryogenic toughness, the volume fraction of the residual γ phase is preferably high, but the residual γ is relatively soft compared to the matrix phase, and when the residual γ amount is excessive, a predetermined YS and Since TS may not be secured, the upper limit is set to 5.0%. A preferred upper limit for the volume fraction of the residual γ phase is 4.8%. A preferred lower limit is 3.0%, and a more preferred lower limit is 3.5%.

(E) Properly control the balance of the Mn—Ni content in the steel represented by the following formula (4): [Mn] ≧ 0.31 × (7.20− [Ni]) + 0.50・ (4)
Satisfying this condition further enhances the stability of the residual γ. Hereinafter, the requirement of the above formula (4) may be referred to as “Ni—Mn balance in steel” or simply “Ni—Mn balance”.

The outline of how the above equation (4) was reached is as follows. In order to ensure a high strength-toughness balance at extremely low temperatures while reducing the amount of Ni to 7.5% or less, the present inventors use Mn which is a γ-stabilizing element among the components in steel. In addition, from the viewpoint that the balance between Mn and Ni having a relatively high content among the components in the steel is important, a design guideline in steel for improving the stability of residual γ. It was investigated. Specifically, from the viewpoints of the effect of hardenability due to Ni reduction, concentration of alloy components during L treatment, and refinement of the MA size formed during impact, the aforementioned Di value and Ms point (start of martensite generation) (Including temperature). As a result, it has been found that the MA size formed at the time of impact has a correlation with the structure size as it is rolled and correlates with the amounts of Ni and Mn in the steel. As a result of further studies based on the above findings, the above formula (4) was specified as the Ni—Mn balance in steel that can ensure the desired strength-toughness balance at extremely low temperatures.

(F) Mn concentration in residual γ existing at −196 ° C .: 1.05% or more By satisfying this condition, the stability of residual γ is enhanced and an excellent strength-toughness balance is achieved at extremely low temperatures. Is done.

The preferable Mn concentration in the residual γ existing at −196 ° C. is 1.40% or more, more preferably 1.75% or more. The preferable upper limit of the Mn concentration in the residual γ is not particularly limited from the relationship with the above action, but considering the range of the amount of Mn in steel and the like, the upper limit is preferably 2.50% or less.

Furthermore, as demonstrated in Example 4 to be described later, at least one of (i) the volume fraction of residual γ, (ii) the Mn concentration in residual γ, and (iii) λ _L parameter is more appropriate. By controlling in such a range, the brittle fracture surface ratio can be maintained at a favorable level of 50% or less even at −233 ° C., which is lower than −196 ° C. described above. Specifically, (i) the residual γ fraction is approximately 3.5 to 4.8%, (ii) the Mn concentration in the residual γ is approximately 1.40 to 2.5%, (iii) λ _{By controlling the L} parameter within a range of approximately −10 or less, toughness at −233 ° C. can be improved. Further, when at least two of the above (i) to (iii) and / or (i) the Mn concentration in the residual γ is controlled to 1.75 to 2.50%, the toughness at −233 ° C. is further improved. Can do.

The thick steel plate of the present invention has been described above.

Next, a method for producing the thick steel plate of the present invention will be described. The method for producing a thick steel plate according to the present invention as defined in claim 1 or 2 of the present invention is the temperature (L treatment temperature) in the heat treatment (L treatment) in the α-γ2 phase coexistence region (between A _c1 and A _c3 ). And the L parameter calculated based on the following formula (5) composed of A _c1 and A _c3 in the steel is 0.25 or more and 0.45 or less, and the L parameter and the steel The step of adjusting the L treatment temperature and the steel component so that the λ _L parameter calculated based on the following formula (6) composed of the components is 7 or less, and after the L treatment, It is characterized in that after water cooling to room temperature, a tempering process (T process) is performed at a temperature of _Ac1 or lower for 10 to 60 minutes.
L parameter = (L treatment temperature−A _c1 ) / (A _c3 −A _c1 ) +0.25 (5)
λ _L parameter = 9.05 × (0.90 × [L parameter] +0.14) × [Mn] + 1.46 × (0.37 × [L parameter] +0.67) × [Cr] −41.5 × (0.26 × [L parameter] +0.79) × [Mo] (6)
(In formula, [] means content (mass%) of each component in steel.)

Hereinafter, each process will be described in detail.

The manufacturing method of the present invention described above is for manufacturing a thick steel plate that satisfies the above requirements by appropriately controlling the rolling step and the subsequent tempering treatment (T treatment), and the steel making step is not particularly limited. The method used can be employed.

Hereinafter, the steps after the rolling step characterizing the present invention will be described in detail in order.

First, it is preferable to control the heating temperature to about 900 to 1100 ° C., the FRT (finish rolling temperature) to about 700 to 900 ° C., and the SCT (cooling start temperature) to about 650 to 800 ° C. Here, the SCT is preferably controlled within the above-mentioned range within 60 seconds after finish rolling, whereby a microstructure useful for improving toughness can be obtained after rolling → cooling.

Next, the temperature range from 800 ° C. to 500 ° C. is cooled at an average cooling rate of about 10 ° C./s or more. In the present invention, the average cooling rate in the above temperature range is particularly controlled in order to obtain a fine structure after cooling. The upper limit is not particularly limited.

In the present invention, it is preferable to cool at least the above temperature range at an average cooling rate of about 10 ° C./s or more, but the stop temperature at the above average cooling rate is preferably 200 ° C. or less. Thereby, untransformed γ can be reduced, and a fine and uniform structure can be obtained.

After hot rolling, it is heated and maintained at the two-phase region [ferrite (α) -γ] temperature (L treatment temperature) at points A _c1 to A _c3 and then cooled with water (L treatment). In the present invention, in order to control the volume fraction of residual γ and the residual γ stabilization parameter (preferably, the volume fraction of residual γ and the residual γ stabilization parameter) within the scope of the present invention, the above equation (5) The L treatment temperature and the components in the steel are appropriately controlled so that the L parameter represented by the formula (1) and the λ _L parameter represented by the above formula (6) fall within a predetermined range.

First, the L treatment temperature after hot rolling is preferably controlled within the range of A _c1 to (A _c1 + A _c3 ) / 2. Thereby, alloy elements such as Ni are concentrated in the generated γ phase, and a part thereof becomes a metastable residual γ phase that exists metastable at room temperature. When the L treatment temperature is less than the A _c1 point or more than [(A _c1 + A _c3 ) / 2], the residual γ fraction at −196 ° C. or the stability of the residual γ cannot be secured sufficiently as a result (described later). (See Nos. 29 and 30 in Table 2B of Example 1). A preferable L treatment temperature is approximately 620 to 650 ° C.

In the present specification, the _Ac1 point and the _Ac3 point are calculated based on the following formulas ("Lecture / Modern Metallurgy Materials 4 Steel Materials", Japan Institute of Metals).
A _c1 point = 723-10.7 × [Mn] −16.9 × [Ni] + 29.1 × [Si] + 16.9 × [Cr] + 290 × [As] + 6.38 × [W]
A _c3 point = 910−203 × [C] ^1/2 −15.2 × [Ni] + 44.7 × [Si] + 104 × [V] + 31.5 × [Mo] + − 30 × [Mn] + 11 × [Cr] + 20 × [Cu]
In the above formula, [] means the concentration (mass%) of the alloying element in the steel material. In the present invention, As and W are not included as components in the steel, and in the above formula, [As] and [W] are both calculated as 0%.

The heating time (holding time) at the above two-phase region temperature is preferably about 10 to 50 minutes. If it is less than 10 minutes, the alloy element concentration to the γ phase does not proceed sufficiently, whereas if it exceeds 50 minutes, the α phase is annealed and the strength decreases. The upper limit of the preferred heating time is 30 minutes.

Furthermore, in the present invention, the L parameter represented by the above formula (5) is set to 0.25 or more and 0.45 or less for each component. The L parameter is a parameter set in order to efficiently use the alloy concentration during the L treatment in order to finally have both the volume fraction of the residual γ and the stability of the residual γ. As shown in the examples described later, when the L parameter is out of the above range, the desired residual γ fraction and / or the stability of the residual γ cannot be sufficiently obtained. Preferably they are 0.28 or more and 0.42 or less, More preferably, they are 0.30 or more and 0.40 or less.

Further, in the present invention, as in the above equation (6), the λ _L parameter determined by the respective contents of Mn, Cr and Mo and the L parameter is controlled to be 7 or less. This λ _L parameter is set to suppress the adverse effect of temper embrittlement that occurs in the concentrated part when P is segregated to the old γ grain boundary during L treatment and Mn and Cr are excessively concentrated. It is. Since it is difficult to directly measure the amount of P segregated at the old γ grain boundary, the λ _L parameter can be regarded as an alternative parameter for the amount of P segregated at the old γ grain boundary. Those having a small segregation of P to the former γ grain boundary have a small λ _L parameter. Preferably it is 0.0 or less, more preferably -10.0 or less. The lower limit is not particularly limited, but it is preferable to suppress the amount of addition of Mo as much as possible from the viewpoint of cost. In addition, generally considering the content and the preferable range of the L parameter, approximately −30 or more It is preferable that

Specifically, in the extremely low temperature range of −196 ° C., adverse effects of trace impurities such as P are easily manifested, and temper embrittlement is large when segregation of P to the old γ grain boundary is large (that is, the λ _L parameter is large). Case), it is estimated to have an adverse effect on the cryogenic toughness. For example, No. 1 in Table 1 of Example 1 described later. When comparing 1, 2, and 25 (both examples of the present invention), the volume fraction of residual γ and the residual γ stabilization parameter are comparable [No. No. 1, volume fraction of residual γ = 8.0%, residual γ stabilization parameter = 3.7; No. 2, volume fraction of residual γ = 9.4%, residual γ stabilization parameter = 3.8; 25, the volume fraction of residual γ = 7.9%, the residual γ stabilization parameter = 3.7], and the λ _L parameter were −6.8 (No. 1) and −10.9 (No. 2) It is very different from 5.2 (No. 25). Therefore, in the above three examples, the No. with the lowest λ _L parameter. 2 is most excellent in cryogenic toughness.

Next, after water cooling to room temperature, tempering (T treatment) is performed.

The tempering process is performed at a temperature of _Ac1 or lower for 10 to 60 minutes. By such low temperature tempering, C is concentrated in the metastable residual γ and the stability of the metastable residual γ phase is increased, so that a residual γ phase that exists stably even at −196 ° C. is obtained. Moreover, a low Ms point can be secured by the low temperature tempering.

When the tempering temperature exceeds the _Ac1 temperature, the metastable residual γ phase generated while maintaining the two-phase coexistence region is decomposed into an α phase and a cementite phase, and a sufficient residual γ phase at −196 ° C. cannot be secured. On the other hand, when the tempering temperature is less than 540 ° C. or when the tempering time is less than 10 minutes, C concentration into the metastable residual γ phase does not proceed sufficiently, and the desired residual γ at −196 ° C. The amount cannot be secured. Further, when the tempering time exceeds 60 minutes, the dislocation density of the α phase is excessively reduced, and a predetermined strength (TS and YS) cannot be secured (No. 33 in Table 2B of Example 1 described later is set. reference).

Preferred tempering conditions are tempering temperature: 540 to 560 ° C., tempering time: 15 minutes or more and 45 minutes or less (more preferably 35 minutes or less, more preferably 25 minutes or less).

After tempering as above, cool to room temperature. The cooling method after tempering is not water cooling but air cooling. This is because carbon concentrates in the residual γ during air cooling, so that the residual γ stabilization parameter is larger in air cooling than in water cooling.

Next, the manufacturing method of the thick steel plate which concerns on this invention prescribed | regulated to this-application Claim 3 is demonstrated.
In the production method of the present invention, the L parameter calculated based on the formula (5) is 0.6 or more and 1.1 or less, and the L parameter and the component in steel ( 6) The step of adjusting the L treatment temperature and steel components so that the λ _L parameter calculated based on the equation is 0 or less, and the L treatment, followed by water cooling to room temperature and tempering treatment (T treatment) is characterized in that it is performed at a temperature of _{Ac 1} or lower for 10 to 60 minutes.

Hereinafter, although each process is explained in full detail, the part which overlaps with the method of manufacturing the thick steel plate of the present invention according to claim 1 or 2 described above (conditions of rolling process, conditions of temperature and holding time of L treatment, etc.) ) Will be omitted.

In the method for manufacturing a thick steel plate according to the present invention defined in claim 3 of the present application, the L parameter represented by the formula (5) is set to 0.6 or more and 1.1 or less. The L parameter is a parameter set to finally combine the volume fraction of residual γ and the stability of residual γ (particularly expressed by the Di value and the Mn concentration in the residual). In particular, the upper limit (1.1 or less) was specified from the viewpoint of the components of the thick steel plate and the desired structure conditions. Note that increasing the stability of the residual γ by the L treatment (that is, concentrating Mn into the residual γ) means that the Mn concentration of the parent phase (in the steel) is diluted when reversed. In this state, since the strength is adversely affected or the volume fraction of residual γ and the stability of residual γ cannot be combined, the lower limit of L parameter (0.6 or more) is set in the present invention. A preferable L parameter is 0.7 or more and 1.0 or less.

Further, in the present invention, as shown in the above equation (6), the contents of Mn, Cr and Mo in the steel and the λ _L parameter determined by the _L parameter are controlled to be 0 or less. As described above, this λ _L parameter is used to suppress the adverse effect of temper embrittlement that occurs in the concentrated portion when P is segregated to the old γ grain boundary during L treatment and Mn and Cr are excessively concentrated. It is set. Since the amount of P segregating at the old grain boundary cannot be directly measured, the λ _L parameter can be regarded as an alternative parameter for the amount of P segregating at the old γ grain boundary. Those having a small segregation of P to the former γ grain boundary have a small λ _L parameter. Preferably it is -10.0 or less. The lower limit is not particularly limited, but it is preferable to suppress the amount of addition of Mo as much as possible from the viewpoint of cost. In addition, generally considering the content and the preferable range of the L parameter, approximately −30 or more It is preferable that

The tempering process is performed at a temperature of _Ac1 or lower for 10 to 60 minutes. As described above, C is concentrated in the metastable residual γ by such low temperature tempering, and the stability of the metastable residual γ phase is increased. Therefore, a residual γ phase that exists stably even at −196 ° C. is obtained. Moreover, a low Ms point can be secured by the low temperature tempering.

When the tempering temperature exceeds A _c1 , the metastable residual γ phase generated while maintaining the two-phase coexistence region is decomposed into an α phase and a cementite phase, and a sufficient residual γ phase at −196 ° C. cannot be secured. On the other hand, if the tempering time is less than 10 minutes, the C concentration in the metastable residual γ phase does not proceed sufficiently, and the desired residual γ amount at −196 ° C. cannot be ensured. Further, when the tempering time exceeds 60 minutes, the dislocation density of the α phase is excessively reduced, and a predetermined strength (TS) cannot be secured (see No. 7 in Table 2B of Example 3 described later). . A preferable tempering time is 15 minutes or more and 45 minutes or less, more preferably 20 minutes or more and 35 minutes or less.

Further, the tempering temperature is a temperature of A _c1 or less, and the preferable tempering temperature is 510 ° C. to 520 ° C.

After tempering as above, cool to room temperature. The cooling method after tempering is not water cooling but air cooling. This is because carbon is concentrated in the residual γ during air cooling, so that the stability of the residual γ is higher in air cooling than in water cooling.

Hereinafter, the present invention will be described in more detail with reference to examples, but the present invention is not limited by the following examples, and can be implemented with modifications within a range that can meet the purpose described above and below. They are all included in the technical scope of the present invention.

Example 1: Example relating to a thick steel plate satisfying a brittle fracture surface ratio ≦ 10% at −196 ° C., a tensile strength TS> 741 MPa, and a yield strength YS> 590 MPa. A test steel having a component composition (remainder: iron and inevitable impurities, the unit is mass%) was melted and cast, and then a 150 mm × 150 mm × 600 mm ingot was produced by hot forging. In this example, misch metal containing about 50% Ce and about 25% La was used as REM.

Next, after heating the above ingot to 1100 ° C., it is rolled to a sheet thickness of 75 mm at a temperature of 830 ° C. or higher, the finish rolling temperature (FRT) is 700 ° C., SCT within 60 seconds after FRT is 650 ° C., As a result, the sheet was rolled to a thickness of 25 mm (a rolling reduction of 83%). The average cooling rate from 800 to 500 ° C. was 19 ° C./s, and cold rolling was performed to a stop temperature of 200 ° C. or lower.

The steel plate thus obtained was subjected to L treatment at the L treatment temperature shown in Table 2, and heated and held for 30 minutes, and then cooled with water. Further, T treatment (tempering) was performed at the temperature (T treatment temperature) and time (T time) shown in Table 2, and then cooled to room temperature.

With respect to the thick steel plate thus obtained, the amount of residual γ phase (volume fraction) present at −196 ° C., the residual γ stabilization parameter, tensile properties (tensile strength TS, yield strength YS) are as follows. ), And cryogenic toughness (the brittle fracture surface ratio in the C direction at −196 ° C. or −233 ° C.).

(1) Measurement of amount of residual γ phase (volume fraction) present at −196 ° C. From a t / 4 position of each steel plate, a 10 mm × 10 mm × 55 mm test piece was taken and liquid nitrogen temperature (−196 ° C.) Was held for 5 minutes, and then X-ray diffraction measurement was performed with a two-dimensional microscopic X-ray diffractometer (RINT-RAPIDI value I) manufactured by Rigaku Corporation. Next, the peaks of the lattice planes (110), (200), (211), and (220) of the ferrite phase and the lattices of (111), (200), (220), and (311) of the residual γ phase For the peak of the surface, based on the integrated intensity ratio of each peak, calculate the volume fraction of (111), (200), (220), (311) of the residual γ phase, and obtain the average value of these, Was defined as “volume fraction of residual γ”.

(2) Measurement of residual γ stabilization parameter In order to measure the residual γ stabilization parameter calculated based on the above equation (2), each of the residual γ existing at −196 ° C. constituting the above equation (2) Components, that is, C content <C>, Mn content <Mn>, Al content <Al>, Cu content <Cu>, Ni content <Ni>, Cr content <Cr>, Mo content <Mo>, V content <V > Was measured as follows.

(2-1) Measurement of C content <C> in residual γ existing at −196 ° C. Simultaneously with the measurement in (1) above, the standard material Si was applied to the surface of the test steel, and the angle was corrected with the Si peak. To obtain a precise lattice constant [a ₀ (Å)] of γ-Fe. From the refined lattice constant of γ-Fe and the following components other than C, the amount of C in the residual γ was calculated by back calculation.

(2-2) Measurement of Ni content <Ni> in residual γ existing at −196 ° C. From a t / 4 position of each steel plate, a 10 mm × 10 mm × 55 mm test piece was collected and liquid nitrogen temperature (−196 ° C. ), The Ni concentration was measured under the conditions of an acceleration voltage of 15 kV, an irradiation current of 50 nA, and a minimum beam diameter using an EPMA apparatus manufactured by JEOL, JXA-8500F. Measurement was performed three times for each sample, and the maximum value was defined as the amount of Ni in the residual γ.

(2-3) Measurement of Al amount <Al> in residual γ existing at −196 ° C. Al is assumed that the entire amount is consumed by oxide or nitride, and Al in residual γ is 0 (zero) It was.

(2-4) Measurement of Mn content <Mn>, Cu content <Cu>, Cr content <Cr>, Mo content <Mo>, and V content <V> in residual γ existing at −196 ° C. Then, the concentration of each alloy element <Mn>, <Cu>, <Cr>, <Mo>, <V> after the L treatment → T treatment is the measured Ni amount obtained by the method (2-2) above. Considering that it is proportional to <Ni>, it was calculated as follows.

The behavior of Ni concentration during each heat treatment of L treatment and T treatment is expressed by the following equation.
(Constant during each heat treatment) x (Driving force of γ reverse transformation) x (Diffusion coefficient of each alloy element)

Here, (driving force of γ reverse transformation) in the above formula was calculated by commercially available calculation software (thermocalc) based on the temperature during each heat treatment. The (diffusion coefficient of each alloy element) in the above formula was calculated based on the temperature and holding time during each heat treatment using the value of “DiffusionDin Solid Metals and Alloys” (H. Mehrer, 1990).

And (constant at the time of each heat treatment) in the above formula was experimentally determined as follows. According to the above formula, the measured Ni concentration after L treatment → T treatment is {(constant during L treatment) × (driving force of γ reverse transformation) × (diffusion coefficient of Ni during L treatment)} {(Constant at the time of T treatment) × (driving force for γ reverse transformation) × (diffusion coefficient of Ni at the time of L treatment)}. That is, the measured Ni concentration after L treatment → T treatment includes both (constant for L treatment) and (constant for T treatment), and (constant for T treatment) is (constant for L treatment). ) To change the value of the product so that it is closest to the measured Ni concentration after the L treatment → T treatment, the constants [(constants at the L treatment) and ( Constant during T treatment)]. Using the constants thus obtained, the alloy element concentrations of <Mn>, <Cu>, <Cr>, <Mo>, and <V> were calculated.

(3) Measurement of tensile properties (tensile strength TS, yield strength YS) From the t / 4 position of each steel plate, No. 4 test piece of JIS Z2241 was taken in parallel to the C direction, and a tensile test was performed by the method described in ZIS Z2241. The tensile strength TS and the yield strength YS were measured. In this example, TS> 740 MPa and YS> 590 MPa were evaluated as having excellent base material strength.

(4) Measurement of cryogenic toughness (brittle fracture surface ratio in the C direction) t / 4 position (t: plate thickness) and W / 4 position (W: plate width), t / 4 position and W of each steel plate / 3 position, three Charpy impact test pieces (V-notch test piece of JIS Z 2242) were taken in parallel with the C direction, and the brittle fracture surface rate at -196 ° C (%) was measured by the method described in JIS Z2242. Were measured and the average value of each was calculated. Of the two average values calculated in this way, the average value that is inferior in characteristics (that is, the brittle fracture surface ratio is large) is adopted, and this value is 10% or less. Then, it evaluated that it was excellent in cryogenic toughness.

These results are also shown in Table 2.

From Table 2, it can be considered as follows.

First, No. in Table 2A. Examples 1 to 25 are examples that satisfy all of the requirements of the present invention. Even if the base metal strength is high, the cryogenic toughness at −196 ° C. (specifically, the average value of the brittle fracture surface ratio in the C direction ≦ 10 %), An excellent thick steel plate could be provided.

In contrast, No. in Table 2B. Nos. 26 to 45 are comparative examples that do not satisfy the requirements of the present invention because they do not satisfy any of the components in the steel and the preferred production conditions of the present invention, and the desired characteristics cannot be obtained.

First, No. No. 26 is an example in which the Di value does not satisfy the requirements of the present invention, and the desired volume fraction of residual γ cannot be obtained, and the residual γ stabilization parameter also decreases. Furthermore, the volume fraction of residual γ and the residual γ stabilization parameter also exceeded the predetermined range. As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized at -196 ° C. Moreover, since Di value was low, TS also fell.

No. No. 27 in Table 1B with a large amount of C. In this example, the cryogenic toughness was lowered.

No. No. 28 is No. in Table 1B with a large amount of P. In this example, the desired volume fraction of residual γ was not obtained, and the residual γ stabilization parameter was also reduced. Furthermore, the volume fraction of residual γ and the residual γ stabilization parameter also exceeded the predetermined range. As a result, the cryogenic toughness decreased.

No. No. 29 in Table 1B in which the components in the steel satisfy the requirements of the present invention. 29 is used, but heating is performed at a temperature lower than the two-phase region temperature (L treatment temperature), and the L parameter is low. Therefore, the amount of residual γ was insufficient and the residual γ stabilization parameter was also reduced. Furthermore, the volume fraction of residual γ and the residual γ stabilization parameter also exceeded the predetermined range. As a result, the cryogenic toughness decreased.

No. No. 30 is No. in Table 1B with a large amount of Si. 30 and heating at a temperature exceeding the two-phase region temperature (L treatment temperature), and the L parameter and the λ _L parameter are high. Therefore, the amount of residual γ was insufficient and the residual γ stabilization parameter was also reduced. Furthermore, the volume fraction of residual γ and the residual γ stabilization parameter also exceeded the predetermined range. As a result, the cryogenic toughness decreased.

No. No. 31 in Table 1B in which the components in the steel satisfy the requirements of the present invention. 31 was used, but because the tempering temperature (T treatment temperature) was low, the amount of residual γ was insufficient and the residual γ stabilization parameter was also lowered. Furthermore, the volume fraction of residual γ and the residual γ stabilization parameter also exceeded the predetermined range. As a result, the cryogenic toughness decreased.

No. No. 32 is No. in Table 1B with a large amount of Mn. 32, and the λ _L parameter is high. As a result, the cryogenic toughness decreased.

No. No. 33 in Table 1B where the components in the steel satisfy the requirements of the present invention. Although 33 was used, this was an example in which the tempering time (T time) was long, and the strength (TS and YS) decreased.

No. No. 34 in Table 1B has a small amount of Mn and a small Di value. In this example, the desired volume fraction of residual γ was not obtained, and the residual γ stabilization parameter was also reduced. Furthermore, the volume fraction of residual γ and the residual γ stabilization parameter also exceeded the predetermined range. As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized at -196 ° C. Moreover, since Di value was low, TS also fell.

No. No. 35 in Table 1B with a large amount of S. This is an example using 35. For this reason, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

No. No. 36 in Table 1B in which the components in the steel satisfy the requirements of the present invention. Although 36 was used, this is an example where the L parameter is high. As a result, the amount of residual γ was insufficient, and the volume fraction of residual γ and the residual γ stabilization parameter exceeded a predetermined range. As a result, the cryogenic toughness decreased.

No. No. 37 in Table 1B has a small amount of C, a large amount of Al, and a small amount of Ni. Since 37 was used, the amount of residual γ was insufficient and the residual γ stabilization parameter was also reduced. Furthermore, the volume fraction of residual γ and the residual γ stabilization parameter also exceeded the predetermined range. As a result, the cryogenic toughness decreased. TS also decreased.

No. No. 38 in Table 1B has a small amount of Al and a large amount of N. Since 38 was used, the cryogenic toughness decreased.

No. No. 39 in Table 1B, which has a large amount of Cu and Ca as the selection components. Since 39 was used, cryogenic toughness decreased.

No. No. 40 in Table 1B, which has a large amount of Cr and Zr as selective components. Since 40 was used, the cryogenic toughness decreased.

No. No. 41 in Table 1B with a large amount of Nb and REM as the selection components. Since 41 was used, the cryogenic toughness decreased.

No. No. 42 in Table 1B has a large amount of Mo as a selection component and a large Di value. Since 42 was used, the cryogenic toughness decreased.

No. No. 43 in Table 1B with a large amount of Ti as a selected component. Since 43 was used, the cryogenic toughness decreased.

No. No. 44 of Table 1B with a large amount of V as a selected component. Since 44 was used, the cryogenic toughness decreased.

No. No. 45 of Table 1B with a large amount of B as a selected component. Since 45 was used, the cryogenic toughness decreased.

Example 2:
In this example, the brittle fracture surface rate at −233 ° C. was evaluated for some of the data used in Example 1 (all of the examples of the present invention).

Specifically, No. in Table 3 (No. in Table 3 corresponds to No. in Table 1 and Table 2 described above) Three test pieces were collected from the t / 4 position and the W / 4 position, and -233 was obtained by the method described below. A Charpy impact test at ℃ was conducted to evaluate the average brittle fracture surface ratio.

In this example, the brittle fracture surface ratio ≦ 50% was evaluated as being excellent in the brittle fracture surface ratio at −233 ° C.
"High pressure gas", Vol. 24, page 181, "Cryogenic impact test of austenitic cast stainless steel"

These results are shown in Table 3.

No. in Table 3 1 to 3, 6, 8, 9, 14, 18, and 20 all have not only −196 ° C. but also a good brittle fracture surface ratio at a lower temperature of −233 ° C., and extremely excellent cryogenic temperature. Toughness could be achieved. The reason for this may be that the volume fraction of residual γ and the residual γ stabilization parameter are both small (21 or less).

On the other hand, No. Compared with the above example, 10 and 16 had a large volume fraction of residual γ and a residual γ stabilization parameter of about 35, so the brittle fracture surface ratio at −233 ° C. was large.

From the above experimental results, in order to obtain a thick steel sheet having not only −196 ° C. but also a brittle fracture surface ratio at a lower temperature of −233 ° C., in the above requirements of the present invention, in particular, the volume fraction of residual γ It can be seen that it is effective to make the residual γ stabilization parameter as small as possible.

Example 3: Example relating to a thick steel plate satisfying a brittle fracture surface ratio ≦ 10% at −196 ° C., a tensile strength TS> 830 MPa, and a yield strength YS> 690 MPa. A test steel having a component composition (remainder: iron and inevitable impurities, the unit is mass%) was melted and cast, and then a 150 mm × 150 mm × 600 mm ingot was produced by hot forging. In this example, misch metal containing about 50% Ce and about 25% La was used as REM.
Next, after heating the above ingot to 1100 ° C., it is rolled to a sheet thickness of 75 mm at a temperature of 830 ° C. or higher, the finish rolling temperature (FRT) is 700 ° C., SCT within 60 seconds after FRT is 650 ° C., By doing so, it rolled to plate | board thickness 25mm (rolling rate 85%). The average cooling rate from 800 to 500 ° C. was 19 ° C./s, and cold rolling was performed to a stop temperature of 200 ° C. or lower.

The steel plate thus obtained was subjected to L treatment at the L treatment temperature shown in Table 5, heated and held for 30 minutes, and then cooled with water. Further, T treatment (tempering) was performed at the temperature (T treatment temperature) and time (T time) shown in Table 2, and then cooled to room temperature.

For the thick steel plate thus obtained, the amount of residual γ phase (volume fraction) present at −196 ° C., the amount of Mn in the residual γ phase, tensile properties (tensile strength TS, yield strength YS), cryogenic temperature Toughness (the brittle fracture surface ratio in the C direction at -196 ° C or -233 ° C) was evaluated.

Regarding the measurement of the amount of residual γ phase (volume fraction) present at 196 ° C., the measurement of tensile properties (tensile strength TS, yield strength YS), and the measurement of cryogenic toughness (brittle fracture surface ratio in the C direction) Since this is the same as in Example 1, measurement of the amount of Mn in the residual γ phase existing at −196 ° C. will be described.

The average amount of Mn in the residual γ phase was measured with TEM-EDX and calculated according to the following procedure. In the calculation, it was assumed that the component in the residual γ phase was Fe—Mn—Ni. Actual components may include, for example, C, Si and the like in addition to Fe, Mn, and Ni, but these elements are in a small amount and can be substantially ignored in this embodiment.

First, a 10 mm × 10 mm × 55 mm test piece was taken from the t / 4 position of each steel plate, held at a liquid nitrogen temperature (−196 ° C.) for 5 minutes, and then the test piece was made into a size of 10 mm × 10 mm × 2 mm. After cutting and mechanically polishing the thickness t from 2 mm to 0.1 mm, it was punched out to a size of 3 mmΦ to prepare a thin film sample by electrolytic polishing. The thin film sample thus obtained was identified with a transmission image and a reciprocal lattice using a transmission electron microscope H-800 manufactured by Hitachi, Ltd., and then the above-mentioned γ was detected using an EDX analyzer EMAX7000 manufactured by Horiba. The Mn concentration in the phase was measured. Measurement by EDX was performed under the conditions of an acceleration voltage of 200 kV and an observation magnification of 75000 times, and each sample was measured at five points, and the average value was taken as the amount of Mn in the residual γ.

In Example 3, unlike Example 1, TS> 830 MPa and YS> 690 MPa were evaluated as having excellent base material strength.

These results are also shown in Table 5.

From Table 5, it can be considered as follows.

First, No. in Table 5A. Nos. 1 to 21 are No. 1 in Table 4A in which the components in the steel satisfy the requirements of the present invention. 1 to 21 is an example prepared under the production conditions of the present invention, and the cryogenic toughness at −196 ° C. (specifically, the average value of the brittle fracture surface ratio in the C direction) even if the base metal strength is high It was possible to provide a thick steel plate excellent in ≦ 10%.

In contrast, No. in Table 5B. Nos. 1 to 21 are comparative examples not satisfying any of the components in the steel and the production conditions of the present invention, and the desired characteristics were not obtained.

First, No. in Table 5B. No. 1 in Table 4B in which the components in the steel satisfy the requirements of the present invention. Although 1 was used, the Di value does not satisfy the requirements of the present invention, and the desired volume fraction of residual γ was not obtained. As a result, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized at -196 ° C.

No. in Table 5B No. 2 in Table 4B has a large amount of C and a small amount of Mo. In this example, the cryogenic toughness decreased.

No. in Table 5B 3 is No. in Table 4B with a large amount of P. In this example, the cryogenic toughness was lowered.

No. in Table 5B No. 4 in Table 4B in which the components in the steel satisfy the requirements of the present invention. 4 is used, however, heating is performed at a temperature lower than the two-phase region temperature (L treatment temperature), and the L parameter is low. Therefore, the amount of residual γ was insufficient. As a result, the cryogenic toughness decreased.

No. in Table 5B. No. 5 in Table 4B with a large amount of Si and Mo. 5 is heated at a temperature exceeding the two-phase region temperature (L treatment temperature), and the L parameter and the λ _L parameter are high. Therefore, the amount of residual γ was insufficient. As a result, the cryogenic toughness decreased.

No. in Table 5B. No. 6 in Table 4B has a large amount of Mn and a small amount of Mo. 6 was used, but the tempering temperature (T treatment temperature) was high, the λ _L parameter was high, and the desired volume fraction of residual γ could not be obtained. As a result, the cryogenic toughness decreased.

No. in Table 5B No. 7 in Table 4B in which the components in the steel satisfy the requirements of the present invention. 7 was used, but the tempering time (T time) was long, and the Ni—Mn balance of the above formula (2) was below the preferred range. As a result, the low temperature toughness decreased. Furthermore, the strength (TS) also decreased.

No. in Table 5B No. 8 in Table 4B with a small amount of Mn. In this example, the Ni—Mn balance in the above formula (2) was below the preferred range, the Mn concentration in the residual γ was lowered, and the residual γ amount was insufficient. As a result, the cryogenic toughness decreased.

No. in Table 5B No. 9 in Table 4B with a large amount of S. 9 is an example. For this reason, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

No. in Table 5B No. 10 in Table 4B has a small amount of C, a large amount of Al, a small amount of Ni, and the Ni—Mn balance of the above formula (2) is below the preferred range. 10 is an example. Since the amount of C and Ni useful for securing the amount of residual γ is small, the volume ratio of residual γ is small. As a result, the cryogenic toughness decreased and YS was good. However, since the amount of C and Ni effective for strength improvement are small, TS decreased.

Table 5BNo. 11 has less amount of Al and Mo content, the number is N quantity, lambda _L parameter is high Table 4B No. Since 11 was used, the cryogenic toughness decreased.

No. in Table 5B No. 12 in Table 4B, which has a large amount of Cu and Ca as the selection components. Since 12 was used, the cryogenic toughness decreased.

No. in Table 5B. No. 13 in Table 4B has a small amount of Mo as a selection component, a large amount of Cr and Zr, and a high λ _L parameter. Since 13 was used, the cryogenic toughness decreased.

No. in Table 5B No. 14 in Table 4B, which has a large amount of Nb and REM as the selection components. Since 14 was used, the cryogenic toughness decreased.

No. in Table 5B No. 15 in Table 4B with a large amount of Mo as a selection component. Since 15 was used, the cryogenic toughness decreased.

No. in Table 5B No. 16 in Table 4B with a large amount of Ti as a selected component. Since 16 was used, the cryogenic toughness decreased.

No. in Table 5B No. 17 in Table 4B with a large amount of V as a selected component. Since No. 17 was used, the cryogenic toughness decreased.

No. in Table 5B No. 18 in Table 4B with a large amount of B as a selected component. Since 18 was used, the cryogenic toughness decreased.

No. in Table 5B No. 19 in Table 4B in which the components in the steel satisfy the requirements of the present invention. 19 is used, but the L parameter is high and the L processing temperature is also high. Therefore, the Mn concentration in the residual γ amount was low, the residual γ amount was insufficient, and the cryogenic toughness was lowered.

No. in Table 5B No. 20 in Table 5B in which the components in the steel satisfy the requirements of the present invention. 20 is an example in which the tempering temperature (T treatment temperature) is high, and the Ni—Mn balance defined by the above formula (2) is below the preferred range, and the desired volume fraction of residual γ can be obtained. In addition, the Mn concentration in the residual γ also decreased. As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized at -196 ° C. Furthermore, YS and TS also decreased.

No. in Table 5B. No. 21 in Table 4B has a small Mo amount and a high _L parameter and a high λ _L parameter. 21 is an example. As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized at -196 ° C.

Example 4
In this example, the brittle fracture surface rate at −233 ° C. was evaluated for the inventive examples of Table 5A used in Example 3 above.

Specifically, No. 1 described in Table 6 was obtained. (No. in Table 6 corresponds to Nos. In Table 4A and Table 5A described above) Three test pieces were collected from the t / 4 position and the W / 4 position, and -233 was obtained by the method described below. A Charpy impact test at ℃ was conducted to evaluate the average brittle fracture surface ratio. In this example, the brittle fracture surface ratio ≦ 50% was evaluated as being excellent in the brittle fracture surface ratio at −233 ° C.
"High pressure gas", Vol. 24, page 181, "Cryogenic impact test of austenitic cast stainless steel"

These results are listed in Table 6. In Table 6, (i) volume fraction of residual γ, (ii) Mn concentration in residual γ, and (iii) λ _L parameter values are extracted from Table 5A for reference. Details of each are as follows.

No. in Table 6 Nos. 1 to 3, 5 to 14, and 17 to 20 are all No. 1 in Table 5A that satisfy at least one of the above (i) to (iii). Examples 1 to 3, 5 to 14, and 17 to 20 were used, and the brittle fracture surface ratio at −233 ° C. was good at 50% or less. On the other hand, no. Nos. 4, 15, 16, and 21 satisfy No. 1 in Table 5A that do not satisfy any of the requirements (i) to (iii) above. In this example, 4, 15, 16, and 21 were used, and the desired toughness could not be obtained at -233 ° C.

First, No. in Table 6 1 to 3 are Nos. 1 to 5 in Table 5A that satisfy the requirement (ii). Since 1 to 3 were used, the brittle fracture surface ratio at −233 ° C. was as good as 50%.

In contrast, No. in Table 6 No. 4 in Table 5A, which does not have any of the above requirements (i) to (iii). 4 was used, the desired toughness could not be obtained at -233 ° C.

Next, no. No. 5 in Table 5A, which has all the above requirements (i) to (iii), and (ii) controls the Mn concentration in the residual γ to a more preferable range of 1.75 to 2.50%. 5 was used, the toughness at −233 ° C. could be further increased to 15%.

In addition, No. in Table 6 No. 6 in Table 5A that satisfies the requirements (i) and (iii) above. 6 was used, the toughness at −233 ° C. could be further increased to 40%.

No. in Table 6 No. 7 in Table 5A that satisfies the requirement (iii) above. 7 was used, the brittle fracture surface rate at −233 ° C. was as good as 50%.

No. in Table 6 No. 8 in Table 5A, which combines the above requirements (i) and (ii), and (ii) controlled the Mn concentration in the residual γ to a more preferable range of 1.75 to 2.50%. Since 8 was used, the toughness at −233 ° C. could be further increased to 25%.

No. in Table 6 No. 9 in Table 5A that satisfies the requirements (i) and (iii) above. 9 was used, the toughness at −233 ° C. could be further increased to 40%.

No. in Table 6 No. 10 in Table 5A that satisfies the requirements (ii) and (iii) above. Since 10 was used, the toughness at −233 ° C. could be further increased to 40%.

No. in Table 6 No. 11 in Table 5A that satisfies the requirement (ii) above. 11 was used, the brittle fracture surface rate at −233 ° C. was as good as 50%.

No. in Table 6 No. 12 in Table 5A that satisfies the requirements (ii) and (iii) above. Since 12 was used, the toughness at −233 ° C. could be further increased to 40%.

No. in Table 6 No. 13 in Table 5A, which has all the requirements (i) to (iii) described above, and (ii) the Mn concentration in the residual γ is controlled to a more preferable range of 1.75 to 2.50%. Since 13 was used, the toughness at −233 ° C. could be further increased to 15%.

No. in Table 6 No. 14 in Table 5A that satisfies the requirement (ii) above. 14 was used, the brittle fracture surface rate at −233 ° C. was as good as 50%.

In contrast, No. in Table 6 Nos. 15 and 16 are No. 5 in Table 5A that do not have any of the requirements (i) to (iii). Since 15 and 16 were used, the desired toughness could not be obtained at -233 ° C.

On the other hand, no. No. 17 in Table 5A that satisfies the requirement (iii) above. Since No. 17 was used, the brittle fracture surface ratio at -233 ° C. was good at 50%.

No. in Table 6 No. 18 in Table 5A that satisfies the requirement (i) above. 18 was used, the brittle fracture surface rate at -233 ° C. was as good as 50%.

No. in Table 6 No. 19 in Table 5A that satisfies the requirement (ii) above. 19 was used, the brittle fracture surface rate at -233 ° C. was as good as 50%.

No. in Table 6 No. 20 in Table 5A that satisfies the requirement (i) above. 20 was used, the brittle fracture surface rate at -233 ° C. was as good as 50%.

In contrast, No. in Table 6 No. 21 in Table 5A does not have any of the requirements (i) to (iii) above. Since No. 21 was used, the desired toughness could not be obtained at -233 ° C.

Although the present invention has been described in detail and with reference to specific embodiments, it will be apparent to those skilled in the art that various changes and modifications can be made without departing from the spirit and scope of the invention.

This application is based on a Japanese patent application filed on December 13, 2012 (Japanese Patent Application No. 2012-272184) and a Japanese patent application filed on December 27, 2012 (Japanese Patent Application No. 2012-285916). Incorporated herein by reference.

The steel plate of the present invention is useful as a steel plate that comes into contact with a cryogenic substance, such as a storage tank for liquefied natural gas.

Claims

% By mass
C: 0.02 to 0.10%,
Si: 0.40% or less (excluding 0%),
Mn: 0.50 to 2.0%,
P: 0.007% or less (excluding 0%),
S: 0.007% or less (excluding 0%),
Al: 0.005 to 0.050%,
Ni: 5.0 to 7.5%,
N: 0.010% or less (excluding 0%)
And containing
Containing at least one element selected from the group consisting of Cr: 1.20% or less (not including 0%) and Mo: 1.0% or less (not including 0%);
The balance is a steel plate with iron and inevitable impurities,
Di value determined on the basis of the following formula (1) composed of components in steel is 2.5 or more,
Di value = ([C] / 10) 0.5 × (1 + 0.7 × [Si]) × (1 + 3.33 × [Mn]) × (1 + 0.35 × [Cu]) × (1 + 0.36 × [ Ni]) × (1 + 2.16 × [Cr]) × (1 + 3 × [Mo]) × (1 + 1.75 × [V]) × 1.115 (1)
(In the formula, [] means the content (% by mass) of each component in the steel)
The residual austenite phase (residual γ) present at −196 ° C. is 2.0 to 12.0% in volume fraction, and
A thick steel plate excellent in cryogenic toughness, characterized in that a residual γ stabilization parameter determined based on the following formula (2) composed of components contained in residual austenite is 3.1 or more.
Residual γ stabilization parameter =
(365 × <C> + 39 × <Mn> + 30 × <Al> + 10 × <Cu> + 17 × <Ni> + 20 × <Cr> + 5 × <Mo> + 35 × <V>) / 100 (2)
(In the formula, <> means the content (% by mass) of each component contained in the retained austenite existing at -196 ° C.)
The volume fraction of residual γ phase and the residual γ stabilization parameter calculated based on the following equation (3) composed of the volume fraction of the residual γ phase and the residual γ stabilization parameter are 40 or less: Item 2. The thick steel plate according to Item 1.
Residual γ volume fraction / residual γ stabilization parameter = 10 / (volume fraction of residual γ phase × residual γ stabilization parameter) 1/2 (3)
% By mass
C: 0.02 to 0.10%,
Si: 0.40% or less (excluding 0%),
Mn: 0.6 to 2.0%,
P: 0.007% or less (excluding 0%),
S: 0.007% or less (excluding 0%),
Al: 0.005 to 0.050%,
Ni: 5.0 to 7.5%,
N: 0.010% or less (excluding 0%)
Mo: 0.30 to 1.0%,
Cr: 1.20% or less (excluding 0%),
Is a thick steel plate with the balance being iron and inevitable impurities,
Di value determined on the basis of the following formula (1) composed of steel components is more than 5.0,
Di = ([C] / 10) 0.5 × (1 + 0.7 × [Si]) × (1 + 3.33 × [Mn]) × (1 + 0.35 × [Cu]) × (1 + 0.36 × [Ni ]) × (1 + 2.16 × [Cr]) × (1 + 3 × [Mo]) × (1 + 1.75 × [V]) × 1.115 (1)
(In the formula, [] means the content (% by mass) of each component in the steel)
The residual austenite phase (residual γ) present at −196 ° C. is 2.0 to 5.0% in volume fraction,
The Mn concentration in the residual austenite phase (residual γ) existing at −196 ° C. is 1.05% or more, and the Mn and Ni contents (mass%) in the steel satisfy the following formula (4). A steel plate with excellent cryogenic toughness characterized by
[Mn] ≧ 0.31 × (7.20− [Ni]) + 0.50 (4)
(In the formula, [] means the content (mass%) of each component in the steel.)
The thick steel plate according to any one of claims 1 to 3, further comprising at least one group of the following groups (a) to (e):
(A) Cu: 1.0% or less (excluding 0%),
(B) Ti: 0.025% or less (not including 0%), Nb: 0.100% or less (not including 0%), and V: 0.50% or less (not including 0%) At least one selected from the group,
(C) B: 0.0050% or less (excluding 0%),
(D) Ca: 0.0030% or less (excluding 0%) and REM: at least one selected from the group consisting of 0.0050% or less (excluding 0%),
(E) Zr: 0.005% or less (excluding 0%).
It is a manufacturing method of the thick steel plate of Claim 1 or 2,
Based on the following formula (5) composed of the temperature (L treatment temperature) in the heat treatment (L treatment) in the α-γ2 phase coexistence region (between A c1 and A c3 ), and A c1 and A c3 in the steel L parameter calculated in the above is 0.25 or more and 0.45 or less, and
Adjusting the L treatment temperature and the steel component so that the λ L parameter calculated based on the following formula (6) composed of the L parameter and the steel component is 7 or less. When,
A step of water cooling to room temperature after L treatment and tempering treatment (T treatment) at a temperature of Ac1 or lower for 10 to 60 minutes;
The manufacturing method of the thick steel plate characterized by performing.
L parameter = (L treatment temperature−A c1 ) / (A c3 −A c1 ) +0.25 (5)
λ L parameter = 9.05 × (0.90 × [L parameter] +0.14) × [Mn] + 1.46 × (0.37 × [L parameter] +0.67) × [Cr] −41.5 × (0.26 × [L parameter] +0.79) × [Mo] (6)
(In formula, [] means content (mass%) of each component in steel.)
It is a manufacturing method of the thick steel plate according to claim 3,
Based on the following formula (5) composed of the temperature (L treatment temperature) in the heat treatment (L treatment) in the α-γ2 phase coexistence region (between A c1 and A c3 ), and A c1 and A c3 in the steel. L parameter calculated in the above is 0.6 or more and 1.1 or less, and
Adjusting the L treatment temperature and the steel component so that the λ L parameter calculated based on the following formula (6) composed of the L parameter and the steel component is 0 or less. A method for producing a thick steel plate.
L parameter = (L treatment temperature−A c1 ) / (A c3 −A c1 ) +0.25 (5)
λ L parameter = 9.05 × (0.90 × [L parameter] +0.14) × [Mn] + 1.46 × (0.37 × [L parameter] +0.67) × [Cr] −41.5 × (0.26 × [L parameter] +0.79) × [Mo] (6)
(In the formula, [] means the content (mass%) of each component in the steel.)