JP5833991B2 - Thick steel plate with excellent cryogenic toughness - Google Patents

Description

  The present invention relates to a thick steel plate excellent in cryogenic toughness. Specifically, even if the Ni content is reduced to about 5.0 to 7.5%, the toughness at a cryogenic temperature of −196 ° C. or less [in particular, This relates to a thick steel plate having good toughness in the plate width direction (C direction). 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 lower.

  In addition to high strength, a thick steel plate for an LNG tank used for a storage tank for liquefied natural gas (LNG) is required to have high toughness that can withstand an extremely low temperature of -196 ° C. 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, but since the cost of Ni has increased in recent years, it is less than 9% However, even with a low Ni content, development of a thick steel plate excellent in cryogenic toughness has been underway.

For example, Non-Patent Document 1 describes the effect of α-γ two-phase coexistence heat treatment on the low temperature toughness of 6% Ni steel. Specifically, before the tempering treatment, a heat treatment (L treatment) in the α-γ two-phase coexistence region (between A _{c1 and} A _c3 ) is added to obtain a 9% Ni steel that has undergone a normal quenching and tempering treatment. It can provide a cryogenic toughness at −196 ° C. that is equal to or better; this heat treatment also improves the toughness of the specimen in the C direction (plate width direction); It is described that it is due to the presence of retained austenite that is stable against impact loads. However, according to the above method, although the cryogenic toughness in the rolling direction (L direction) is excellent, the cryogenic toughness in the sheet width direction (C direction) 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. Among these, Patent Document 1 discloses that a 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 below the A _c1 transformation point. Further, Patent Document 2 contains 4.0 to 10% of Ni and heat treatment similar to Patent Document 1 (L treatment → tempering) for steel in which the size of AlN before hot rolling is 1 μm or less. The method of performing the processing) is described. The impact value (vE ₋₁₉₆ ) at −196 ° C. 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.

  Non-Patent Document 2 describes the development of 6% Ni steel for LNG tanks that combines the above-described 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) exhibits a high value, the toughness value in the sheet width direction (C direction) is not described.

  Patent Document 3 includes 0.3 to 10% Ni and a predetermined amount of Mg, and Mg-containing oxide particles having a predetermined particle diameter are appropriately dispersed. A tough high tensile steel is described. Patent Document 3 discloses that the grain size of the heated austenite is refined by controlling the Mg-containing oxide, and the toughness of the base material and the weld heat affected zone (HAZ) is improved; The order of addition of O (oxygen) and Mg and other deoxidizing elements is important. After adding Mg, Ti and Al simultaneously to molten steel with dissolved oxygen content of 0.001 to 0.02%, casting In the addition of Mg, Ti, or Al, it is described that Al is added lastly and then cast into a steel piece. In the examples of Patent Document 3, the toughness value in the C direction (fracture surface transition temperature vTrs) is described, and the above properties of 9% Ni steel are good (fracture surface transition temperature vTrs ≦ −196 ° C.). However, the above-mentioned property of Ni steel in the vicinity of 5% is −140 ° C., and further improvement is required.

  Furthermore, in Patent Document 4, 5.0 to 7.5% Ni is added, and the austenite distribution is made uniform, so that the toughness (CTOD characteristics), arrestability, and unstable fracture suppression of the base metal and the welded joint A technique for obtaining a thick steel plate having excellent characteristics is described. However, the evaluation temperature in the CTOD test is slightly high at −165 ° C., and is not a technique disclosed for extremely low temperatures of −196 ° C. or lower. Further, even if Patent Document 4 is scrutinized, there is no description of the brittle fracture surface ratio in the Charpy impact absorption test. Moreover, in order to manufacture the thick steel plate described in Patent Document 4, a long-time heat treatment at a high temperature of 1250 to 1380 ° C. for 8 to 50 hours is necessary before hot rolling. Is disadvantageous.

JP-A-49-13581 Japanese Patent Laid-Open No. 51-13308 JP 2001-123245 A Japanese Patent No. 4975888

Yano et al., “Effect of heat treatment of α-γ two-phase coexistence on low temperature toughness of 6% Ni steel”, Iron and Steel, 59th (1973) No. 6, p752-763 Furuya et al., “Development of 6% Ni steel for LNG tanks”, CAMP-ISIJ, Vol. 23 (2010), p1322

  As described above, until now, a technique excellent in cryogenic toughness at −196 ° C. in Ni steel having a Ni content of about 5.0 to 7.5% has been proposed, but the cryogenic temperature in the C direction has been proposed. Toughness has not been fully studied. In particular, there is a strong demand for further improvement in cryogenic toughness under high strength (in particular, tensile strength TS> 690 MPa, yield strength YS> 590 MPa) (high improvement in cryogenic toughness in the C direction) with high base metal strength. ing.

  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. In areas where brittle fracture occurs, the energy absorbed by the steel material by the time it reaches the fracture is significantly reduced, and the fracture progresses easily. It is a very important requirement to suppress the brittle fracture surface rate appearing in the Charpy impact test to a low level (10% or less). However, since the brittle fracture is more likely to occur as the strength is higher, it is generally difficult to realize the brittle fracture surface ratio ≦ 10% under the high base metal strength as described above. For this reason, in a high-strength thick steel plate having a high base material strength, a technique having both of these has not been proposed yet.

  The present invention has been made in view of the above circumstances, and its purpose is to achieve cryogenic toughness (especially in the C direction) at -196 ° C 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%.

The thick steel plate excellent in the cryogenic toughness according to the present invention that can solve the above problems is in mass%, C: 0.02 to 0.10%, Si: 0.40% or less (not including 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. 0. 050%, Ni: 5.0 to 7.5%, N: 0.010% or less (excluding 0%), the balance being thick steel plate with iron and inevitable impurities, present in the steel plate The average roundness (A) of inclusions having an equivalent circle diameter exceeding 1.0 μm is 1.8 or less, and the volume fraction (V) of the retained austenite phase existing at −196 ° C. is 2.0-12. It has a gist where it satisfies 0% and the B value represented by the following formula (1) is 1.3 or more.
B = V ^2/3 / A (1)

  In a preferred embodiment of the present invention, in the steel sheet, the retained austenite phase existing at -196 ° C satisfies 4.0 to 12.0% in volume fraction.

  In a preferred embodiment of the present invention, the steel sheet further contains Cu: 1.00% or less (excluding 0%).

  In a preferred embodiment of the present invention, the steel sheet is further selected from the group consisting of Cr: 1.2% or less (not including 0%) and Mo: 1.00% or less (not including 0%). Containing at least one kind.

  In a preferred embodiment of the present invention, the steel sheet further includes: Ti: 0.025% or less (not including 0%), Nb: 0.10% 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 further includes Ca: 0.0030% or less (excluding 0%), REM: 0.0050% or less (excluding 0%), and Zr: 0.0050. % Or less (not including 0%).

  According to the present invention, even in a low Ni steel whose Ni content is reduced to about 5.0 to 7.5%, even if the base material strength is high (specifically, tensile strength TS> 690 MPa, yield strength YS> 590 MPa) and excellent in cryogenic toughness at −196 ° C. or less (particularly in the C direction, cryogenic toughness). Specifically, in the Charpy impact absorption test in the C direction, the brittle fracture surface ratio at −196 ° C. ≦ It was possible to provide a high-strength thick steel plate that satisfies 10% (preferably, the brittle fracture surface ratio ≦ 50% at −233 ° C.).

In the high strength thick steel sheet in which the Ni content is reduced to 7.5% or less and the tensile strength TS> 690 MPa and the yield strength YS> 590 MPa are satisfied, the Charpy impact value in the C direction is −196 ° C. In order to provide a cryogenic toughness improvement technology satisfying the brittle fracture surface ratio of ≦ 10%, a study was conducted. As a result, (a) the volume fraction V of the retained austenite (residual γ) phase existing at −196 ° C. is controlled to 2.0 to 12.0% [preferably 4.0 to 12.0% (volume (B) for inclusions with an equivalent circle diameter exceeding 1.0 μm that promote the progress of brittle fracture (hereinafter sometimes referred to simply as inclusions). When the circularity A was reduced to 1.8 or less, and the B value represented by the following formula (1) was controlled to 1.3 or more, it was found that the intended purpose was achieved, and the present invention was completed. .
B = V ^2/3 / A (1)

  In particular, there is a characteristic part to be noted in relation to the prior art described above in the latter (A). Hereinafter, the process of reaching the present invention will be described.

  The inventors of the present invention have made extensive studies in order to provide a thick steel plate excellent in cryogenic toughness at −196 ° C. in Ni steel having a Ni content of 7.5% or less. Specifically, the present invention provides a high-strength thick steel plate excellent in cryogenic toughness that satisfies all the properties of the brittle fracture surface ratio in the C direction ≦ 10%, tensile strength TS> 690 MPa, and yield strength YS> 590 MPa. From this point of view, first, the methods taught in the documents described in the prior art were examined.

The above document teaches that it is important to stabilize the retained austenite (residual γ) existing at −196 ° C. in order to improve the cryogenic toughness of 5% Ni steel. Further, considering the manufacturing method comprehensively, in the molten steel stage, the amount of dissolved oxygen before the deoxidizing element addition is controlled, and in this molten steel, Al is added last, and α-γ two-phase is added. After heat treatment (L treatment) in the phase coexistence region (between A _{c1 and} A _c3 ), a method of tempering at a temperature below the A _c1 transformation point is recommended, which may improve cryogenic toughness. Taught. However, according to the examination results of the present inventors, although the cryogenic toughness in the L direction is improved by the above method, the cryogenic toughness in the C direction is not sufficient, and the above target level (C direction) described in the present invention is not sufficient. It has been found that the brittle fracture surface ratio at ≦ 10%) cannot be realized.

As a result of further studies, in order to obtain the desired thick steel plate with excellent cryogenic toughness, additional requirements were added to the thick steel plate and its manufacturing method while basically following the above-mentioned technology. I found out that it was essential. Specifically, (i) In the thick steel plate, in addition to the presence of the volume fraction V of the residual γ phase at −196 ° C. in the range of V = 2.0 to 12.0%, the progress of brittle fracture Focusing on inclusions with an equivalent circle diameter of more than 1.0 μm that were found to promote the above, the average roundness A of the inclusions was reduced to 1.8 or less, and the average roundness A of the inclusions It is effective to control the B value represented by the relational expression (1) with the volume fraction V (%) of the residual γ phase existing at −196 ° C. so that the B value ≧ 1.3. ii) In order to produce such a thick steel plate, in the molten steel stage, control of the dissolved oxygen amount (free O amount) before Al addition and heat treatment between A _{c1 and} A _c3 after hot rolling ( L treatment) → In addition to tempering treatment in a predetermined temperature range, further control of the molten steel stage is effective, and 1450-15 at the time of casting The cooling time (t1) at 00 ° C. is controlled to 300 seconds or less (value at a half position of the slab thickness t), and the cooling time (t2) at 1300 to 1200 ° C. during casting is 680 seconds or less ( It has been found that it is effective to control the slab thickness to a value at a quarter position of the slab thickness t.

Further, (c) In the above (a), by controlling the residual γ phase existing at −196 ° C. to 4.0 to 12.0% (volume fraction), the brittle fracture can be achieved even at a lower temperature of −233 ° C. (A) In order to manufacture such a thick steel plate, heat treatment between A _{c1 and} A _c3 after hot rolling (L treatment) is possible. The present invention has been completed by finding that holding for a predetermined time is effective.

  In this specification, “excellent in cryogenic toughness” means −196 ° C. when the brittle fracture surface ratio in the C direction (plate width direction) Charpy impact absorption test is measured by the method described in the column of Examples described later. The brittle fracture surface ratio at ≦ 10% is satisfied. In the examples to be described later, the brittle fracture surface ratio in the L direction (rolling direction) is not measured. However, if the brittle fracture surface ratio in the C direction is 10% or less, the brittle fracture surface in the L direction. The rate is based on an empirical rule that the rate is necessarily 10% or less.

  In the present specification, the “thick steel plate” means a steel plate having a thickness of about 6 to 50 mm.

  In the present invention, a high-strength thick steel plate that satisfies the tensile strength TS> 690 MPa and the yield strength YS> 590 MPa is an object.

  Hereinafter, the thick steel plate of the present invention will be described in detail.

As described above, the thick steel plate of the present invention is in mass%, C: 0.02 to 0.10%, Si: 0.40% or less (not including 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 (excluding 0%), the balance being thick steel plate with iron and unavoidable impurities, with an equivalent circle diameter exceeding 1.0 μm present in the steel plate The average roundness (A) of the product is 1.8 or less, the volume fraction (V) of the retained austenite phase existing at −196 ° C. satisfies 2.0 to 12.0%, and the following (1 ) Is characterized in that the B value represented by the formula is 1.3 or more.
B = V ^2/3 / A (1)

  First, the components in steel will be described.

C: 0.02-0.10%
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% or less. 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 is lowered, 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 ensured, 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 ensure 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 a deoxidizing element. If the Al content is insufficient, the free oxygen concentration in the molten steel increases, and secondary inclusions such as oxides and sulfides are complexly formed on the surface of the inclusions originally present in the molten steel during the cooling of the casting. As a result, the shape of the inclusion becomes distorted, and the average roundness of the inclusion having an equivalent circle diameter of more than 1.0 μm increases, 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, aggregation and coalescence of the inclusions are promoted, and the average roundness of the inclusions is also increased, 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: 5.0-7.5%
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 with preferable Ni amount is 7.0% or less, More preferably, it is 6.5% or less, More preferably, it is 6.0% or less.

N: 0.010% or less (excluding 0%)
N lowers the cryogenic toughness by strain aging, so its 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.

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

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

Cu: 1.00% 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, so the upper limit is preferably made 1.00% or less. A more preferable upper limit of the amount of Cu is 0.8% or less, and further preferably 0.7% or less.

At least one selected from the group consisting of Cr: 1.2% or less (not including 0%) and Mo: 1.00% 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, it is preferable that 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 improved and the desired cryogenic toughness cannot be ensured. Therefore, the upper limit of Cr content is preferably 1.2% or less (more preferably 1.1% or less, still more preferably 0.9% or less, still more preferably 0.5% or less), and a preferable upper limit of the Mo amount is 1.00% or less (more preferably 0.8% or less, still more preferably 0.6% or less). .

Selected from the group consisting of Ti: 0.025% or less (not including 0%), Nb: 0.10% 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 0.10% 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%), REM (rare earth element): 0.0050% or less (not including 0%), and Zr: 0.0050% or less (not including 0%) At least one selected from the group consisting of Ca, REM, and Zr are all deoxidizing elements, and their addition has a positive effect on toughness by reducing the oxygen concentration in the steel and reducing the amount of oxide. Effect. These elements may be added alone or in combination of two or more. In order to effectively exhibit the above action, the Ca amount is 0.0005% or more, the REM amount (when the REM described below is contained alone, it is a single content, and when two or more are contained. Is the total amount thereof, hereinafter, the same applies to the REM amount) is preferably 0.0005% or more, and the Zr amount is preferably 0.0005% or more. However, when added in excess, the oxide size increases and the cryogenic toughness decreases, so the preferable upper limit of Ca amount is 0.0030% or less (more preferably 0.0025% or less), and the preferable upper limit of REM amount. Is 0.0050% or less (more preferably 0.0040% or less), and a preferable upper limit of the amount of Zr 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.

  In the above, the component in steel of this invention was demonstrated.

  Further, in the thick steel plate of the present invention, the volume fraction V of the residual γ phase existing at −196 ° C. satisfies 2.0 to 12.0% (preferably 4.0 to 12.0%). .

  It is known that residual γ existing at −196 ° C. contributes to improvement of cryogenic toughness. In order to effectively exhibit such an action, the volume fraction V of the residual γ phase in all tissues 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 set to 12.0% (Table 2B described later) No. 43). For the volume fraction V of the residual γ phase, a preferred lower limit is 4.0% or more, a more preferred lower limit is 6.0% or more, a preferred upper limit is 11.5% or less, and a more preferred upper limit is 11.0% or less. is there.

  Furthermore, by controlling the volume fraction of residual γ in all tissues existing at −196 ° C. to 4.0% or more, the brittle fracture surface ratio can be reduced even at −233 ° C., which is lower than −196 ° C. described above. It can be kept at a good level of 50% or less. When it is desired to further exert such effects, the more preferable lower limit is 6.0% or more, and the preferable upper limit is the same as described above.

  In the thick steel plate of the present invention, control of the volume fraction V of the residual γ phase is important among the structures existing at −196 ° C., and other structures other than the residual γ are not limited in any way. There is no limitation as long as it is normally present in a thick steel plate. Examples of the structure other than the residual γ include carbides such as bainite, martensite, and cementite.

Further, the thick steel plate of the present invention has an average roundness A of the inclusion satisfying A ≦ 1.8 for inclusions having an equivalent circle diameter exceeding 1.0 μm present in the steel plate, and the following (1 B value represented by the formula satisfies 1.3 or more.
B = V ^2/3 / A (1)

  Here, the “equivalent circle diameter” refers to the diameter of a circle that is assumed to have the same area by paying attention to the size of the inclusions.

  The reason for focusing on inclusions having an equivalent circle diameter of more than 1.0 μm in the present invention is that it has been found that the inclusions promote the progress of brittle fracture. That is, in order to improve the brittle fracture surface ratio at extremely low temperatures while ensuring a predetermined high strength, it is necessary to reduce inclusions that promote brittle fracture. For example, when the average roundness A of the inclusion is increased, the desired cryogenic toughness cannot be achieved even if the volume fraction V of the residual γ phase at −196 ° C. is controlled within the above range. (See Nos. 33, 35, and 36 in Table 2B described later). The average roundness A of the inclusions is preferably as small as possible, preferably 1.7 or less, and more preferably 1.5 or less. Most preferably 1. In the present invention, the average size of inclusions having an equivalent circle diameter of more than 1.0 μm (average equivalent circle diameter) is approximately 2.0 μm or less.

  The said inclusion can be measured by the method as described in the Example mentioned later. Here, the type of “inclusion” in inclusions having an equivalent circle diameter of more than 1.0 μm is not particularly limited in the present invention. This is because the occurrence of brittle fracture is influenced not by the type of inclusion but by the size of the inclusion (equivalent circle diameter). Examples of the inclusions include single particles such as oxides, sulfides, nitrides, and oxynitrides, composites in which two or more of these single particles are combined, or single particles and others. Composite particles in which these elements are bonded.

  As in the present invention, by setting the average roundness A of coarse inclusions having an equivalent circle diameter of more than 1.0 μm to 1.8 or less, the mechanism for improving the cryogenic toughness while ensuring a predetermined strength is as follows. The details are unknown, but it is presumed as follows. In general, inclusions have higher hardness than matrix, and therefore stress concentration tends to occur, and as a result, often acts as a starting point of brittle fracture. At this time, as the shape of the inclusion becomes distorted, stress concentration around the inclusion is locally promoted, so that it is considered that brittle fracture is more easily induced. Therefore, if the irregular inclusions are reduced [that is, if the average roundness A of the inclusions is set to 1.8 or less and is controlled to be as close to the perfect circle (A = 1) as possible), the stress concentration is reduced. It is presumed that the occurrence of this is avoided and the cryogenic toughness is improved.

  Further, in the present invention, it is necessary not only to control the average roundness A of the inclusions but also to satisfy the B value ≧ 1.3 in the B value represented by the above formula (1).

  The B value is a parameter for reducing the brittle fracture surface ratio at an extremely low temperature. As shown in the equation (1), the average roundness A of the inclusions and residual austenite existing at −196 ° C. It is calculated in relation to the volume fraction V of the (residual γ) phase. Hereinafter, the process of deriving the B value will be described.

  It is known that the brittle fracture surface ratio increases as the starting point of brittle fracture increases and the resistance to the progress of brittle fracture decreases. In general, coarse inclusions are likely to be the starting point of brittle fracture, but the present inventors, as the roundness of coarse inclusions increases, in other words, a more distorted shape that deviates from a perfect circle (A = 1). It has been found that it becomes easier to act as a starting point of brittle fracture; and the more residual γ, the more resistant to the progress of brittle fracture. Based on these findings, the contribution ratio of both to the brittle fracture surface ratio in the cryogenic temperature range was experimentally determined based on a number of basic experiments. As a result, the B value represented by the above equation (1) is the cryogenic toughness. It was found that this was a useful parameter for evaluation. As shown in the examples described later, the B value is set to 1 after ensuring the volume fraction V of the residual γ phase and the form of coarse inclusions (average roundness) exceeding the equivalent circle diameter of 1.0 μm. For the first time, the strength and the brittle fracture surface ratio at −196 ° C. and −233 ° C. can be compatible at a high level by controlling to 3 or more.

About the said B value, Preferably it is 1.6 or more, More preferably, it is 1.8 or more. The B value is better from the viewpoint of cryogenic toughness, and the upper limit is not particularly limited. However, as described above, if the volume fraction V of residual γ becomes excessive, YS cannot secure a predetermined value, so the upper limit of the volume fraction V of residual γ is limited to 12.0%. , The upper limit of the B value is substantially limited to 5.2 (= 12.0 ^2/3 / 1) (in the B value calculation formula, the volume fraction V of residual γ = 12.0. %, Average roundness A = 1 is substituted). Considering the balance between strength and toughness, a more preferable B value is 3.0 or less.

  Next, a method for producing the thick steel plate of the present invention will be described.

The characteristic part of the manufacturing method which concerns on this invention exists in following (A)-(B).
(A) In the molten steel stage, the amount of free oxygen [O] before addition of Al is 100 ppm or less, and the cooling time (t1) at 1450-1500 ° C. during casting is 300 seconds or less (value at the 1/2 position of the slab thickness t). ) And the cooling time (t2) at 1300 to 1200 ° C. during casting is controlled to 680 seconds or less (value at ¼ position of the slab thickness t). By the method (A), the average roundness A of the inclusions described above can be reduced to a predetermined range.
(B) After hot rolling, after heating and holding in the temperature range of A _{c1 to} A _c3 points, water cooling, and subsequently tempering in the temperature range of 520 ° C. to A _c1 points for 10 to 60 minutes, then air cooling Or cool with water. In particular, the volume fraction of the residual γ phase existing at −196 ° C. is appropriately controlled by the method (B).

  Note that the B value defined in the present invention is a parameter related to both the average roundness of inclusions and the volume fraction of residual γ, and therefore by appropriately controlling the above (A) to (B) The B value can be controlled within a predetermined range.

  Speaking of the above-mentioned related art, the method (A) has the greatest feature in that t1 and t2 are particularly controlled.

  Hereinafter, each process is explained in full detail.

(About melting process)
In the present invention, special attention is paid to the method of adding Al. This is because inclusions with an equivalent circle diameter of more than 1.0 μm to be controlled in the present invention mainly include secondary inclusions such as oxides and sulfides starting from Al-based inclusions generated in the molten metal. This is because the Al inclusions are likely to be coarsened by agglomeration and coalescence, and to have an irregular shape with a high roundness, although they are produced in a complex manner upon cooling.

  First, when adding Al, which is a deoxidizing material, to a molten steel, the amount of free oxygen before adding Al (sometimes abbreviated as the amount of dissolved oxygen or [O]) is controlled to 100 ppm or less. When the amount of [O] exceeds 100 ppm, Al-based inclusions generated when Al is added increase and the roundness exceeds a predetermined range (see No. 33 in Table 2B described later). The smaller the amount of [O], the better, preferably 80 ppm or less, and more preferably 50 ppm or less. The lower limit of the amount [O] is not particularly limited from the viewpoint of controlling the average roundness of the inclusions.

  As a method for controlling the amount of [O] as described above, for example, a method of deoxidizing by adding deoxidizing elements of Mn and Si into molten steel can be mentioned. In addition to the above elements, when a deoxidizing material such as Ti, Ca, REM, or Zr is added as a selective component, the amount of [O] can be controlled also by adding these.

  In order to control Al inclusions, it is important to control the amount of [O] before addition of Al, and the order of addition of Al and other deoxidizing elements is not limited. However, if Al is added in a state where the amount of [O] is high, the temperature of the molten steel rises due to an oxidation reaction, which makes operation dangerous. Therefore, it is preferable to add Si and Mn prior to Al. Moreover, it is preferable to add the said selective components, such as Ti, in molten steel after addition of Al.

  Next, casting starts. The temperature range during casting is generally 1650 ° C. or lower. In the present invention, the cooling time (t1) in the temperature range of 1450 to 1500 ° C. is particularly controlled to 300 seconds or less, and cooling at 1300 to 1200 ° C. It was important to control the time (t2) to 680 seconds or less, and it was found that the average roundness of inclusions having an equivalent circle diameter exceeding 1.0 μm was appropriately controlled. Details will be described below.

  First, the cooling time (t1) in the temperature range of 1450 to 1500 ° C. is controlled to 300 seconds or less. When t1 exceeds 300 seconds, the complex formation of secondary inclusions with Al inclusions as the core is promoted, and the shape of inclusions with an equivalent circle diameter exceeding 1.0 μm becomes distorted and the average roundness is increased. Or the B value decreases, and the desired cryogenic toughness is not exhibited (see Nos. 34 and 35 in Table 2B described later). From the above viewpoint, t1 is preferably as short as possible, preferably 290 seconds or less, and more preferably 280 seconds or less. The lower limit of t1 is not particularly limited from the above viewpoint.

  In the present invention, among the temperature ranges at the time of casting, particularly the temperature range of 1450 to 1500 ° C. was focused on that the solidification at the time of casting proceeds and the concentration of components in the molten steel proceeds. In this temperature range, the growth of inclusions is promoted.

  The temperature range of 1450 to 1500 ° C. means the temperature of the central portion (t / 2) of the slab thickness t. The reason is as follows. As described above, since oxide-based secondary inclusions are mainly formed in the molten metal, it is necessary to control the cooling time of the molten metal part. However, since 1450-1500 degreeC is a temperature range which solidification is advancing, it solidifies during a measurement depending on a temperature measurement position, and there exists a possibility that the cooling time of a molten metal part cannot be measured correctly. Therefore, in the present invention, the cooling time at the t / 2 position where the molten metal exists up to the lowest temperature was measured. The temperature at the center of the slab thickness can be measured by inserting a thermocouple into the mold.

  Next, the cooling time (t2) at 1300 to 1200 ° C. is controlled to 680 seconds or less. When t2 exceeds 680 seconds, composite formation of mainly sulfide-based secondary inclusions in the Al-based inclusions is promoted, and the average roundness of the inclusions is also increased (see the table described later). 2B No. 36). From the above point of view, the shorter t2, the closer to a perfect circle, the more useful. Preferred t2 is 650 seconds or less, more preferably 600 seconds or less. However, if t2 is too short, the cooling load increases, so in practice, it is recommended that the time be approximately 400 seconds or more.

  The temperature range of 1300 to 1200 ° C. means a temperature of ¼ part (t / 4) of the slab thickness t. The reason for this is that the cooling time of 1300 to 1200 ° C. is mainly for controlling sulfide-based secondary inclusions that are complexly formed in solid iron, but in the above temperature range, solidification is almost completed. Therefore, it was decided to measure the cooling time at the t / 4 position where the brittle fracture surface ratio was measured. The temperature at t / 4 part of the slab thickness can be measured by inserting a thermocouple into the mold.

  In the present invention, the cooling time (t1) in the temperature range of 1450 to 1500 ° C. and the cooling time (t2) at 1300 to 1200 ° C. need only be controlled as described above, and the means is not limited. . For example, with respect to t1, the temperature range may be cooled at a constant speed and an average cooling rate of about 0.17 ° C./second or less so that the cooling time in the temperature range is 300 seconds or less, or The cooling may be performed at different cooling rates so that the cooling time in the above temperature range is 300 seconds or less. The same applies to t2.

  Moreover, in this invention, the cooling method about the temperature range at the time of casting other than the said temperature range is not limited at all, A normal method (air cooling or water cooling) can be employ | adopted.

  After casting as described above, it is hot-rolled and subjected to heat treatment.

  Here, the hot rolling step is not particularly limited, and a commonly used method can be adopted so that a predetermined plate thickness is obtained. Specifically, the slab is heated at about 1100 ° C. for 1 to 4 hours. After that, the (finish rolling) temperature, reduction amount, etc. may be adjusted.

After hot rolling, it is heated to the temperature range (TL) of points A _{c1 to} A _c3 , held, and then cooled with water. This process corresponds to the L process described in the above-mentioned prior art, and by this, it is possible to secure the residual γ that exists stably at −196 ° C. within a predetermined range.

In particular, heating to the two-phase region of the A _c1 to A _c3 point [ferrite (α) -γ] Temperature (TL). By heating to this temperature range, alloy elements such as Ni are concentrated in the produced γ phase, and a metastable residual γ phase existing metastable at room temperature is obtained. If it is less than A _c1 point or more than A _c3 point, as a result, a sufficient residual γ phase at −196 ° C. cannot be secured (see Nos. 37 and 38 in Table 2B described later). A preferable heating temperature is approximately 660 to 710 ° C.

  The heating time (holding time, tL) 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. A preferred heating time is approximately 15 to 30 minutes.

  Further, by setting the heating time to 15 minutes or more, the volume fraction of the residual γ phase at −196 ° C. is ensured to be 4.0% or more, and thereby the brittle fracture surface ratio at −233 ° C. is increased. Good toughness is ensured even at a very low temperature of 50% or less. A more preferable lower limit in the case where such an effect is desired to be exhibited effectively is 5.0% or more. In addition, the upper limit of a preferable heating time is the same as the above (30 minutes or less).

Subsequently, after water-cooling to room temperature, it tempers. A tempering process is performed for 10 to 60 minutes (t3) in the temperature range (T3) of 520 degreeC- _Ac1 point. As a result, C is concentrated in the metastable residual γ during tempering, and the stability of the metastable residual γ phase increases, so that a residual γ phase that exists stably even at −196 ° C. is obtained. When the tempering temperature T3 is lower than 520 ° C., 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 (described later). (See No. 41 in Table 2). On the other hand, when the tempering temperature T3 exceeds the A _c1 point or the tempering time t3 is less than 10 minutes, the C concentration in the metastable residual γ phase does not proceed sufficiently, and the desired residual γ at −196 ° C. The amount cannot be ensured [No. 55 (see example where t3 is short)]. Moreover, when the tempering time t3 exceeds 60 minutes, a residual γ phase at −196 ° C. is excessively generated, and a predetermined strength cannot be secured (see No. 43 in Table 2 described later).

  Preferred tempering conditions are tempering temperature T3: 570 to 620 ° C., and tempering time t3: 15 minutes or more and 45 minutes or less (more preferably 35 minutes or less, more preferably 25 minutes or less).

  After tempering as described above, it is cooled to room temperature. The cooling method is not particularly limited, and may be either air cooling or water cooling.

In the present specification, the A _c1 point and the A _c3 point are calculated based on the following formulas (“Lecture / Modern Metallographic 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] + 13.1 × [W]
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%.

  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
Using a vacuum melting furnace (150 kgVIF), after melting and casting a test steel having the component composition shown in Table 1 (remainder: iron and inevitable impurities, unit is mass%) under the melting conditions shown in Table 2 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. The order of addition of the deoxidizing elements is Si, Mn (simultaneous addition) → Al when no selective component is included; on the other hand, when the selective component of Ti, REM, Zr, Ca is included, Si, Mn (Simultaneous addition) → Al → Ti → REM, Zr, Ca (simultaneous addition). In this example, the time from the addition of Al to the start of casting was all about 10 minutes (not shown in the table).
In Table 2, [O] is the amount of dissolved oxygen (ppm) before Al addition, t1 is the cooling time (seconds) of 1450 to 1500 ° C. during casting, and t2 is the cooling time of 1300 to 1200 ° C. during casting. (Seconds). The cooling in each temperature range was controlled by air cooling or water cooling so that the cooling time was reached.

  Next, after heating the above ingot at 1100 ° C. for 1 to 4 hours, rolling to a plate thickness of 75 mm at a temperature of 830 ° C. or higher, rolling at a final rolling temperature of 780 ° C., and then cooling with water, A 25 mm thick steel plate was obtained. The steel plate thus obtained was heated to the temperature shown in Table 2 (TL in Table 2), heated and held for 5 to 60 minutes (see tL in Table 2), and then cooled to room temperature. Next, as shown in Table 2, after tempering (T3 = tempering temperature, t3 = tempering time), air cooling or water cooling to room temperature was performed.

  With respect to the thick steel plate thus obtained, the volume fraction of residual γ phase (%) existing at an average circularity A of inclusions with an equivalent circle diameter exceeding 1.0 μm at −196 ° C. is as follows. , Tensile properties (tensile strength TS, yield strength YS), and cryogenic toughness (brittle fracture surface ratio in the C direction at -196 ° C or -233 ° C) were evaluated.

(1) Measurement of average roundness A of inclusions having an equivalent circle diameter of more than 1.0 μm The t / 4 position (t: plate thickness) of the steel sheet is mirror-polished, and a four-view photograph at 400 times using an optical microscope. I took a picture. The area per field of view is 0.04 mm ² , and the total area of the four fields of view is 0.15 mm ² . The inclusions observed in these four fields of view were analyzed by “Image-Pro Plus” manufactured by Media Cybernetics, and the roundness of inclusions with an equivalent circle diameter (diameter) exceeding 1.0 μm was calculated based on the following formula: The average value was defined as the average vacuum degree A of the inclusions. When the shape of the inclusion is a perfect circle, the roundness calculated by the following formula is 1. The roundness value calculated by the following formula increases as the shape of the inclusion becomes distorted.

Roundness = L ² / 4π / S
In the formula, L is the perimeter of the inclusion (unit: μm),
S is the area of the inclusion (unit: μm ² ).

In this example, about 200 to 300 / mm ² of inclusions having an equivalent circle diameter of more than 1.0 μm were observed.

(2) Measurement of volume fraction of residual γ phase existing at −196 ° C. A test piece of 10 mm × 10 mm × 55 mm was taken from the t / 4 position of each steel plate, and 5 at a liquid nitrogen temperature (−196 ° C.). After holding for a minute, X-ray diffraction measurement was performed with a two-dimensional micro part X-ray diffractometer (RINT-RAPIDII) 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 (%)”.

(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> 690 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. (%) 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. For reference, Tables 1 and 2 also show points A _c1 and A _c3 .

  From Table 2, it can be considered as follows.

  First, in Table 2A No. 1-32 are examples that satisfy all of the requirements of the present invention. Even if the base material 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.

  On the other hand, No. in Table 2B. Since 33 to 41, 43, and 55 do not satisfy at least one of the preferable production conditions of the present invention, they are comparative examples that do not satisfy the requirements of the present invention, and the desired characteristics were not obtained.

  Specifically, no. No. 33 in Table 1B in which the components in the steel satisfy the requirements of the present invention. Although 33 was used, since the amount of dissolved oxygen [O] before adding Al is large, the average roundness A of the inclusions is increased. As a result, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 34 in Table 1B in which the components in the steel satisfy the requirements of the present invention. 34 was used, but because the cooling time (t1) at 1500 to 1450 ° C. during casting was long, the B value was below the predetermined range. As a result, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 35 is No. in Table 1B with a large amount of P. This is an example in which the average roundness A of the inclusions is increased because the cooling time (t1) at 1500 to 1450 ° C. during casting is long. As a result, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 36 is No. in Table 1B with a large amount of C. This is an example in which the average roundness A of the inclusions is increased because the cooling time (t2) at 1300 to 1200 ° C. during casting is long. As a result, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 37 in Table 1B where the components in the steel satisfy the requirements of the present invention. Although 37 was used, it was heated at a temperature lower than the two-phase region temperature (TL), so that the amount of residual γ was insufficient. As a result, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 38 in Table 1B with a large amount of Si. This is an example in which the amount of residual γ is insufficient because of heating at a temperature exceeding the two-phase region temperature (TL). As a result, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 39 in Table 1B in which the components in the steel satisfy the requirements of the present invention. 39 was used, but the amount of residual γ was insufficient because the heating and holding time (tL) at the two-phase region temperature (TL) was short. As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized.

  No. No. 40 in Table 1B in which the components in the steel satisfy the requirements of the present invention. Although 40 was used, this is an example in which the amount of residual γ increased because the heating and holding time (tL) at the two-phase region temperature (TL) was long. As a result, the yield strength YS and the tensile strength TS decreased, and the desired base material strength could not be ensured.

  No. No. 41 in Table 1B in which the components in the steel satisfy the requirements of the present invention. In this example, the amount of residual γ was insufficient because the tempering temperature (T3) was low. As a result, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 43 in Table 1B in which the components in the steel satisfy the requirements of the present invention. Although 43 was used, the amount of residual γ increased because the tempering time (t3) was long. As a result, the yield strength YS decreased, and the desired base material strength could not be ensured.

  No. No. 55 of Table 1B in which the components in the steel satisfy the requirements of the present invention. In this example, the amount of residual γ was insufficient because the tempering time (t3) was short. As a result, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. Nos. 42 and 44 to 54 are comparative examples manufactured by the method of the present invention using those from which only the components in steel are removed.

  Specifically, no. No. 42 of Table 1B with a large amount of Mn. This is an example in which the amount of residual γ is insufficient because 42 is used. As a result, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 44 of No. 1 in Table 1B with a small amount of Mn. This is an example in which the amount of residual γ is insufficient because 44 is used. As a result, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 45 in Table 1B with a large amount of S. This is an example using 45. For this reason, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 46 of Table 1B has a small amount of C, a large amount of Al, and a small amount of Ni. 46 is used, the average roundness A of the inclusions is increased, and the amount of residual γ is insufficient. As a result, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized. TS also decreased.

  No. No. 47 in Table 1B has a small amount of Al and a large amount of N. This is an example in which the average roundness A of the inclusion is increased because 47 is used. As a result, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 48 of Table 1B with a large amount of Cu and Ca as the selection components. This is an example using 48. As a result, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 49 in Table 1B, which has a large amount of Cr and Zr as selective components. 49 is an example. As a result, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 50 in Table 1B with a large amount of Nb and REM as the selection components. This is an example using 50. As a result, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 51 of Table 1B with a large amount of Mo as a selection component. Since 51 was used, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 52 of Table 1B with a large amount of Ti as a selection component. Since 52 was used, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 53 of Table 1B with a large amount of V as a selected component. Since 53 was used, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 54 of Table 1B with a large amount of B as a selected component. 54 was used, the brittle fracture surface ratio increased, and the desired cryogenic toughness could not be realized.

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

Specifically, No. 1 described in Table 3 is used. (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 according to 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 3.

  No. in Table 3 3, 4, 6, 15, 19, and 24 are examples in which the heating time (tL) in the two-phase region temperature is controlled to 15 minutes or more (see Table 2A), and the residual γ phase is 4 0.0% or more could be secured. As a result, not only −196 ° C. but also a brittle fracture surface rate at a lower temperature of −233 ° C. was good, and a very excellent cryogenic toughness could be achieved.

Claims (7)

% By mass
C: 0.02-0.10%,
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%)
Is a thick steel plate with the balance being iron and inevitable impurities,
The average roundness (A) of inclusions having an equivalent circle diameter of more than 1.0 μm present in the steel sheet is 1.8 or less,
The volume fraction (V) of the residual austenite phase present at −196 ° C. satisfies 2.0 to 12.0%, and
A thick steel plate having excellent cryogenic toughness, wherein the B value represented by the following formula (1) is 1.3 or more.
B = V ^2/3 / A (1)   The thick steel plate according to claim 1, wherein the residual austenite phase existing at -196 ° C is 4.0 to 12.0% in terms of volume fraction.   Furthermore, the thick steel plate of Claim 1 or 2 containing Cu: 1.00% or less (0% is not included). Furthermore,
Cr: 1.2% or less (not including 0%), and Mo: 1.00% or less (not including 0%)
The thick steel plate according to any one of claims 1 to 3, comprising at least one selected from the group consisting of: Furthermore,
Ti: 0.025% or less (excluding 0%),
Nb: 0.10% or less (not including 0%), and V: 0.50% or less (not including 0%)
The thick steel plate according to any one of claims 1 to 4, comprising at least one selected from the group consisting of:   Furthermore, B: The thick steel plate in any one of Claims 1-5 containing 0.0050% or less (0% is not included). Furthermore,
Ca: 0.0030% or less (excluding 0%),
REM: 0.0050% or less (not including 0%), and Zr: 0.0050% or less (not including 0%)
The thick steel plate according to any one of claims 1 to 6, comprising at least one selected from the group consisting of: