EP2889391A1

EP2889391A1 - Thick steel plate having good ultralow-temperature toughness

Info

Publication number: EP2889391A1
Application number: EP13831371.3A
Authority: EP
Inventors: Hidenori Nako; Akira Ibano
Original assignee: Kobe Steel Ltd
Current assignee: Kobe Steel Ltd
Priority date: 2012-08-23
Filing date: 2013-08-19
Publication date: 2015-07-01
Anticipated expiration: 2033-08-19
Also published as: EP2889391A4; EP2889391B1; JP5833991B2; CN104583439A; CN104583439B; KR20150029754A; JP2014040648A; WO2014030618A1; KR101632159B1

Abstract

A thick steel plate that contains definite components in the steel, said thick steel plate satisfying the following requirements: the average roundness (A) of inclusions present in the steel, the equivalent circle diameter of said inclusions exceeding 1.0 µm, is not greater than 1.8; the volume fraction (V) of the retained austenite phase at -196°C is 2.0-12.0%; and value (B) represented by formula (1) is 1.3 or greater.

Description

Technical Field

The present invention relates to a thick steel plate having good ultralow-temperature toughness, more specifically relates to a thick steel plate having good toughness (in particular, toughness in a plate width direction (C direction)) at an ultralow temperature of -196°C or lower even if the Ni content is decreased to about 5.0 to 7.5%. Although description is made mainly on a thick steel plate for liquefied natural gas (LNG) (typically, a storage tank, a transport ship, and the like), which is to be exposed to the ultralow temperature, the thick steel plate of the invention is not limited thereto, and may be applied to general thick steel plates to be used for applications exposed to the ultralow temperature of -196°C or lower.

Background Art

A thick steel plate for a LNG storage tank used as a storage tank of liquefied natural gas (LNG) is required to have high strength and toughness high enough to withstand the ultralow temperature of -196°C. A thick steel plate containing about 9% of Ni (9% Ni steel) has been used as the thick steel plate to be used for such applications. A thick steel plate, which has good ultralow-temperature toughness despite a low Ni content of less than 9%, is now increasingly developed along with recent increase in Ni cost.
For example, NPTL 1 describes an effect of heat treatment in a α-γ coexistence region on low-temperature toughness of 6% Ni steel. In detail, NPTL 1 describes that the 6% Ni steel is subjected to heat treatment (L treatment) in the α-γ coexistence region (between Ac1 and Ac3) before tempering, thereby the 6% Ni steel has the ultralow-temperature toughness at -196°C higher than or equal to that of the 9% Ni steel subjected to normal quenching and tempering; such heat treatment also increases toughness in a C-direction (plate width direction) specimen; and such effects are caused by existence of a large amount of fine retained-austenite that is stable to an impact load at ultralow temperature. Such a technique provides good ultralow-temperature toughness in a rolling direction (L direction). However, ultralow-temperature toughness is likely to be worse in the plate width direction (C direction) than in the L direction. In addition, no description is made on percent brittle fracture.
PTL 1 and PTL 2 each describe a technique similar to that in NPTL 1. PTL 1 describes the following technique. That is, steel containing 4.0 to 10% of Ni, in which austenite grain size is controlled to be within a predetermined range, is hot-rolled and then heated to a temperature between Ac1 and Ac3 and is then cooled. Such treatment, which corresponds to the L treatment described in NPTL 1, is repeated one or more time. Subsequently, the steel is tempered at a temperature lower than or equal to the Ac1 transformation temperature. PTL 2 describes the following technique: steel containing 4.0 to 10% of Ni, in which size of AlN is controlled to be 1 µm or less before hot rolling, is subjected to heat treatment similar to that described in PTL 1 (L treatment followed by tempering). The impact value (vE_-196) at -196°C in description of such techniques is estimated to be a value in the L direction, and the impact value in the C direction is not clear. Each of such techniques makes no consideration on strength. In addition, no description is made therein on percent brittle fracture.
NPTL 2 describes development of 6% Ni steel for a LNG storage tank through a combination of the above-described L treatment (two-phase region quenching) and TMCP. NPTL 2 describes a high value of toughness in the rolling direction (L direction), but does not describe a value of toughness in the plate width direction (C direction).
PTL 3 describes high-tensile steel of 570 MPa class or higher with high toughness, which contains 0.3 to 10% of Ni and a predetermined amount of Mg, contains appropriately dispersed Mg-containing oxide particles having a predetermined particle size, and has good weld toughness. PTL 3 describes the following. That is, the Mg-containing oxide is controlled to refine grain size of austenite, and thereby base metal and a heat affected zone (HAZ) each have improved toughness. To achieve this, important are the amount of O (oxygen) before adding deoxidizing elements, and adding order of Mg and other deoxidizing elements. Specifically, Mg, Ti, and Al are added together into molten steel having an amount of dissolved oxygen of 0.001 to 0.02%, and the molten steel is then cast into billets. Alternatively, Mg, Ti, and Al are added in such a manner that Al is added at the end, and then the molten metal is cast into billets. In an embodiment of PTL 3, toughness values (fracture transition temperature vTrs) in the C direction are described, showing that while 9% Ni steel is good in the above-described properties (fracture transition temperature vTrs of -196°C or less), around 5% Ni steel has the fracture transition temperature vTrs of -140°C and is required to be further improved.
Furthermore, PTL 4 describes the following technique: 5.0 to 7.5% of Ni is added to uniformly distribute austenite, which allows production of a thick steel plate that is good in toughness (CTOD properties), arrestability, and unstable-fracture inhibiting properties for each of base metal and a weld joint. However, evaluation temperature in the CTOD test is slightly high, -165°C, showing that the disclosed technique is not suitable for ultralow temperature of -196°C or lower. As a result of scrutiny of PTL 4, no description has been found on percent brittle fracture in a Charpy impact absorption test. In addition, the steel must be heated for long time at high temperature, i.e., for 8 to 50 hr at 1250 to 1380°C in order to fabricate the thick steel plate described in PTL 4, which is disadvantageous in light of manufacturing cost.

Citation List

Patent Literature

PTL 1: Japanese Unexamined Patent Application Publication No. Sho49 (1974)-135813.
PTL 2: Japanese Unexamined Patent Application Publication No. Sho51 (1976)-13308.
PTL 3: Japanese Unexamined Patent Application Publication No. 2001-123245 .
PTL 4: Japanese Patent No. 4975888 .

Non-patent Literature

NPTL 1: Yano et al.: Effect of Heat Treatment in the Ferrite-Austenite Region on Notch Toughness of 6% Nickel Steel, Tetsu-to-Hagane, 59(1973), 6, pp.752-763.
NPTL 2: Furuya et al.: Development of 6% Ni Steel for LNG Storage Tanks, CAMP-ISIJ, Vol.23(2010), p1322.

Summary of Invention

Technical Problem

As described above, although there have been provided techniques for improving ultralow-temperature toughness at -196°C of Ni steel having the Ni content of about 5.0 to 7.5%, ultralow-temperature toughness in the C direction is not sufficiently investigated. In particular, it is strongly required to further improve ultralow-temperature toughness (improve ultralow-temperature toughness in the C direction) under high strength (in detail, tensile strength of more than 690 MPa and yield strength of more than 590 MPa) of base metal.
Any of the above-described literatures makes no description on percent brittle fracture. The percent brittle fracture indicates the percentage of brittle fracture occurring at load application in a Charpy impact test. In a region where brittle fracture occurs, energy absorbed by steel before start of fracture becomes extremely small, and thus fracture easily propagates. To reduce the fracture particularly at ultralow temperature, therefore, it is an extremely important requirement that the percent brittle fracture shown in a general Charpy impact test is controlled to a low level (10% or less). However, since brittle fracture more easily occurs along with higher strength, it is in general difficult to achieve percent brittle fracture of 10% or less under such high strength of base metal. Hence, there has not been proposed a technique that allows a high-strength thick steel plate with high strength of base metal to have high strength and good ultralow-temperature toughness.
An object of the invention, which has been made in light of the above-described circumstances, is to provide a high-strength thick steel plate of Ni steel having a Ni content of about 5.0 to 7.5%, which is good in ultralow-temperature toughness (particularly the ultralow-temperature toughness in the C direction) at -196°C is good and has percent brittle fracture of 10% or less.

Solution to Problem

To solve the above-described problem, according to the present invention, there is provided a thick steel plate having good ultralow-temperature toughness, which is summarized by containing, by mass percent, 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%, and N: 0.010% or less (not including 0%), with the remainder consisting of iron and inevitable impurities, in which average roundness (A) of inclusions, each having an equivalent circle diameter of more than 1.0 µm, is 1.8 or less, a volume fraction (V) of a retained austenite phase existing at -196°C satisfies 2.0 to 12.0%, and a value B represented by Formula (1) is 1.3 or more. $B = V^{2 / 3} / A$
In a preferred embodiment of the invention, a retained austenite phase existing at -196°C of the steel plate satisfies 4.0 to 12.0% in a volume fraction.
In a preferred embodiment of the invention, the steel plate further contains at least one selected from a group consisting of Cu: 1.00% or less (not including 0%), Cr: 1.2% or less (not including 0%), Mo: 1.00% or less (not including 0%), Ti: 0.025% or less (not including 0%), Nb: 0.10% or less (not including 0%), V: 0.50% or less (not including 0%), B: 0.0050% or less (not including 0%), Ca: 0.0030% or less (not including 0%), REM: 0.0050% or less (not including 0%), and Zr: 0.0050% or less (not including 0%).

Advantageous Effects of Invention

According to the invention, there can be provided a high-strength thick steel plate of low Ni steel having a low Ni content of about 5.0 to 7.5%, which has good ultralow-temperature toughness at -196°C or less (particularly ultralow-temperature toughness in the C direction) despite high strength of base metal (in detail, tensile strength TS of more than 690 MPa and yield strength YS of more than 590 MPa), and specifically satisfies percent brittle fracture at -196°C of 10% or less (preferably, percent brittle fracture at -233°C of 50% or less) in a Charpy impact absorption test in the C direction.

Description of Embodiments

The inventors have conducted investigations to provide a technique for improving ultralow-temperature toughness of a high-strength thick steel plate having a low Ni content of 7.5% or less such that percent brittle fracture at -196°C of 10% or less is satisfied in a Charpy impact test in the C direction while tensile strength TS of more than 690 MPa and yield strength YS of more than 590 MPa are satisfied. As a result, they have found that the desired object is attained by (a) controlling a retained austenite (retained γ) phase at -196°C to have a volume fraction V of 2.0 to 12.0% (preferably controlled to be 4.0 to 12.0% (in volume fraction)), and (b) decreasing the average roundness A of inclusions each having an equivalent circle diameter of more than 1.0 µm, to 1.8 or less, the inclusions promoting propagation of brittle fracture (which may be simply referred to as inclusion hereinafter), and controlling a value B represented by Formula (1) to be 1.3 or more. $B = V^{2 / 3} / A$
The latter (b) is a distinctive characteristic in terms of the above-described existing techniques. Details of achievement of the invention are now described.
The inventors have conducted various investigations to provide a thick steel plate of Ni steel having a Ni content of 7.5% or less and having good ultralow-temperature toughness at -196°C. Specifically, the inventors have first investigated a method taught in a literature on the existing techniques from the viewpoint of providing the high-strength thick steel plate having good ultralow-temperature toughness, the steel plate satisfying all properties of percent brittle fracture of 10% or less in the C direction, tensile strength TS of more than 690 MPa, and yield strength YS of more than 590 MPa.
The literature teaches that the retained austenite (retained γ) existing at -196°C is importantly stabilized to improve ultralow-temperature toughness of 5% Ni steel. In addition, a technique is recommended from comprehensive consideration of a manufacturing method, in which the amount of dissolved oxygen is controlled before adding deoxidizing elements in a stage of molten steel, and the molten steel is cast into a slab in such a manner that Al is added therein at the end, and the slab is subjected to heat treatment (L treatment) in the α-γ coexistence region (between Ac1 and Ac3) and then tempered at a temperature of lower than or equal to the Ac1 transformation temperature, and it is taught that the ultralow-temperature toughness is increased by the technique. Through the investigation results, however, the inventors have found that such a technique increases the ultralow-temperature toughness in the L direction, but does not sufficiently increase the ultralow-temperature toughness in the C direction, and consequently the above-described target level (the percent brittle fracture of 10% or less in the C direction) cannot be achieved by the technique.
As a result of further investigations, the inventors have found that further requirements must be added for a thick steel plate and a method of manufacturing the thick steel plate while the above-described technique is basically adopted in order to produce a desired thick steel plate having good ultralow-temperature toughness. In detail, it has been found that (1) a thick steel plate is effectively designed such that a retained-γ phase at -196°C is allowed to exist in a range of the volume fraction V of 2.0 to 12.0%, and while attention is focused on the inclusions each having the equivalent circle diameter of more than 1.0 µm, the inclusions being found to promote propagation of brittle fracture, the average roundness A of the inclusions is decreased to 1.8 or less, and a value B, which is represented by the relational expression (1) of the average roundness A of the inclusions and the volume fraction V (%) of the retained-γ phase existing at -196°C, is controlled to be 1.3 or more. In addition, it has been found that (2), to fabricate such a thick steel plate, it is effective to perform control of the amount of dissolved oxygen (the amount of free oxygen) before adding Al in a stage of molten steel, and perform heat treatment (L treatment) between Ac1 and Ac3 followed by tempering in a predetermined temperature range after hot rolling, and it is effective to perform further control in the stage of molten steel in such a manner that cooling time (t1) from 1450 to 1500°C in casting is controlled to be 300 sec or less (a value at a half position of the slab thickness t), and cooling time (t2) from 1300 to 1200°C in casting is controlled to be 680 sec or less (a value at a quarter position of the slab thickness t).
Furthermore, the inventors have found that (c), in the control described in (a), the retained-γ phase existing at -196°C is controlled to be 4.0 to 12.0% (in volume fraction), thereby the percent brittle fracture can be maintained to a good level of 50% or less even at a lower temperature of -233°C, and (d), to fabricate such a thick steel plate, the steel is effectively held for a predetermined time in the heat treatment (L treatment) between Ac1 and Ac3 after hot rolling, and eventually they have completed the invention.
In this specification, "ultralow-temperature toughness is good" means that when percent brittle fracture in a Charpy impact absorption test in the C direction (plate width direction) is measured by a method described in the section of Examples described later, the percent brittle fracture at -196°C of 10% or less is satisfied. In Examples described later, percent brittle fracture in the L direction (rolling direction) is not measured. This is based on the following empirical rule: if percent brittle fracture in the C direction is 10% or less, percent brittle fracture in the L direction is necessarily 10% or less.
In this specification, "thick steel plate" refers to a steel plate having a thickness of roughly 6 to 50 mm.
The invention covers the high-strength thick steel plate satisfying tensile strength TS of more than 690 MPa and yield strength YS of more than 590 MPa.
The thick steel plate of the invention is now described in detail.
As described above, the thick steel plate of the invention is characterized by containing , by mass percent, 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%, and N: 0.010% or less (not including 0%), with the remainder consisting of iron and inevitable impurities, in which the average roundness (A) of inclusions, each having an equivalent circle diameter of more than 1.0 µm, is 1.8 or less, a retained austenite phase existing at -196°C has a volume fraction (V) of 2.0 to 12.0%, and a value B represented by Formula (1) is 1.3 or more. $B = V^{2 / 3} / A$
First, components in the steel are described.

C: 0.02 to 0.10%

C is an essential element to obtain strength and retained austenite. The lower limit of the C content is defined to be 0.02% or more to allow such a function to be effectively exhibited. The lower limit of the C content is preferably 0.03% or more, and more preferably 0.04% or more. However, excessively large amount of C causes excessive increase in strength, leading to reduction in ultralow-temperature toughness; hence, the upper limit of the C content is 0.10% or less. The upper limit of the C content is preferably 0.08% or less, and more preferably 0.06% or less.

Si: 0.40% or less (not including 0%)

Si is a useful element as a deoxidizer. However, excessively large amount of Si promotes formation of a hard martensite island, leading to reduction in ultralow-temperature toughness; hence, the upper limit of the Si content is 0.40% or less. The upper limit of the Si content is preferably 0.35% or less, and more preferably 0.20% or less.

Mn: 0.50 to 2.0%

Mn functions as a deoxidizer, and is an austenite (γ) stabilizing element and thus contributes to increasing the amount of retained γ. The lower limit of the Mn content is defined to be 0.50% to allow such a function to be effectively exhibited. The lower limit of the Mn content is preferably 0.6% or more, and more preferably 0.7% or more. However, excessively large amount of Mn causes temper brittleness, which prevents desired ultralow-temperature toughness from being obtained; hence, the upper limit of the Mn content is 2.0% or less. The upper limit of the Mn content is preferably 1.5% or less, and more preferably 1.3% or less.

P: 0.007% or less (not including 0%)

P is an impurity element causing grain boundary fracture. The upper limit of the P content is therefore defined to be 0.007% or less to obtain desired ultralow-temperature toughness. The upper limit of the P content is preferably 0.005% or less. Although the P content is preferably as small as possible, it is industrially difficult to decrease the P content to 0%.

S: 0.007% or less (not including 0%)

S is an impurity element causing grain boundary fracture as with P. The upper limit of the S content is therefore defined to be 0.007% or less to obtain desired ultralow-temperature toughness. As shown in Examples described later, a larger amount of S increases percent brittle fracture, which prevents desired ultralow-temperature toughness (percent brittle fracture at -196°C of 10% or less) from being obtained. The upper limit of the S content is preferably 0.005% or less. Although the S content is preferably as small as possible, it is industrially difficult to decrease the S content to 0%.

Al: 0.005 to 0.050%

Al is a deoxidizing element. When the Al content is insufficient, free oxygen concentration in molten steel increases, and secondary inclusions such as oxides or sulfides are compositely formed on surfaces of inclusions originally existing in molten steel, which makes a shape of each inclusion to be distorted, and increases the average roundness of the inclusions each having the equivalent circle diameter of more than 1.0 µm; hence, the lower limit of the Al content is defined to be 0.005% or more. The lower limit of the Al content is preferably 0.010% or more, and more preferably 0.015% or more. However, excessively large amount of Al promotes aggregation or coalescence of the inclusions, which also increases the average roundness of that inclusions; hence, the upper limit of the Al content is defined to be 0.050% or less. The upper limit of the Al content is preferably 0.045% or less, and more preferably 0.04% or less.

Ni: 5.0 to 7.5%

Ni is an essential element to provide retained austenite (retained γ) useful for improving ultralow-temperature toughness. The lower limit of the Ni content is defined to be 5.0% or more to allow such a function to be effectively exhibited. The lower limit of the Ni content is preferably 5.2% or more, and more preferably 5.4% or more. However, excessively large amount of Ni causes increase in cost of material; hence, the upper limit of the Ni content is defined to be 7.5% or less. The upper limit of the Ni content is preferably 7.0% or less, more preferably 6.5% or less, and most preferably 6.0% or less.

N: 0.010% or less (not including 0%)

N causes strain aging and thereby reduces the ultralow-temperature toughness; hence, the upper limit of the N content is defined to be 0.010% or less. The upper limit of the N content is preferably 0.006% or less, and more preferably 0.004% or less.
The thick steel plate of the invention contains the above-described components as the essential components with the remainder consisting of iron and inevitable impurities.
The thick steel plate of the invention may contain the following optional components in order to add further properties.

Cu: 1.00% or less (not including 0%)

Cu is a γ stabilizing element, i.e., an element that contributes to increasing the amount of retained γ. Cu is preferably contained 0.05% or more to allow such a function to be effectively exhibited. However, excessively large amount of Cu causes excessive increase in strength, which prevents the desired effect on the ultralow-temperature toughness from being exhibited; hence, the upper limit of the Cu content is preferably 1.00% or less. The upper limit of the Cu content is more preferably 0.8% or less, and most preferably 0.7% or less.

At Least One Selected from Group Consisting of Cr: 1.2% or less (not including 0%) and Mo: 1.00% or less (not including 0%)

Cr and Mo are each an element that increases strength. Such elements may each be contained singly or may be contained in combination. The Cr content and the Mo content are preferably 0.05% or more and 0.01% or more, respectively, to allow such a function to be effectively exhibited. However, excessively large amount of each of the elements causes excessive increase in strength, which prevents the desired ultralow-temperature toughness from being obtained. Hence, the upper limit of the Cr content is preferably 1.2% or less (more preferably 1.1% or less, further preferably 0.9% or less, and most preferably 0.5% or less), and the upper limit of the Mo content is preferably 1.00% or less (more preferably 0.8% or less, and most preferably 0.6% or less).

At Least One Selected from 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%)

Ti, Nb, and V are each an element that precipitates as carbonitride and increases strength. Such elements may each be contained singly or may be contained in combination. To allow such a function to be effectively exhibited, it is preferred that the Ti content is 0.005% or more, the Nb content is 0.005% or more, and the V content is 0.005% or more. However, excessively large amount of each of the elements causes excessive increase in strength, which prevents the desired ultralow-temperature toughness from being obtained. Hence, the upper limit of the Ti content is preferably 0.025% or less (more preferably 0.018% or less, and most preferably 0.015% or less), the upper limit of the Nb content is preferably 0.10% or less (more preferably 0.05% or less, and most preferably 0.02% or less), and the upper limit of the V content is preferably 0.50% or less (more preferably 0.3% or less, and most preferably 0.2% or less).

B: 0.0050% or less (not including 0%)

B is an element that improves hardenability and thereby contributes to increasing strength. The B content is preferably 0.0005% or more to allow such a function to be effectively exhibited. However, excessively large amount of B causes excessive increase in strength, which prevents the desired ultralow-temperature toughness from being obtained; hence, the upper limit of the B content is preferably 0.0050% or less, (more preferably 0.0030% or less, and most preferably 0.0020% or less).

At Least One Selected from Group Consisting of Ca: 0.0030% or less (not including 0%), REM: 0.0050% or less (not including 0%), and Zr: 0.0050% or less (not including 0%)

Ca, REM, and Zr are each a deoxidizing element. That is, when each of such elements is contained, oxygen concentration in steel decreases and thus the amount of oxides decreases, and thereby favorable influence is exerted on toughness. Such elements may each be contained singly or may be contained in combination. To allow such a function to be effectively exhibited, the Ca content of 0.0005% or more, the REM content, which refers to the amount of one REM element when each of the following REM elements is singly contained, or refers to the total amount of the REM elements when at least two of them are contained together (the same applies to the following for the REM content), of 0.0005% or more, and the Zr content of 0.0005% or more are preferred. However, excessively large amount of each of such elements increases oxide size, and thus reduces the ultralow-temperature toughness; hence, the upper limit of the Ca content is preferably 0.0030% or less (more preferably 0.0025% or less), the upper limit of the REM content is preferably 0.0050% or less (more preferably 0.0040% or less), and the upper limit of the Zr content is preferably 0.0050% or less (more preferably 0.0040% or less).
In this specification, REM (rare earth elements) refers to an element group including lanthanoid elements (15 elements from atomic number 57 (La) to atomic number 71 (Lu) in a periodic table), Sc (scandium), and Y (yttrium). Such elements may each be contained singly or may be contained in combination. Ce and La are preferred among the rare earth elements. REM may be contained in any form without limitation, i.e., may be contained in a form of misch metal mainly containing Ce and La (for example, Ce: about 70% and La: about 20 to 30%), or may be contained in a form of a simple substance of Ce or La.
The components in the steel of the invention have been described.
Furthermore, the thick steel plate of the invention is designed such that the volume fraction V of the retained-γ phase existing at -196°C satisfies 2.0 to 12.0% (preferably 4.0 to 12.0%).
The retained-γ phase existing at -196°C is known to contribute to improvement in ultralow-temperature toughness. To allow such a function to be effectively exhibited, the volume fraction V of the retained-γ phase in the entire microstructures existing at -196°C is defined to be 2.0% or more. However, the retained γ is relatively soft compared with a matrix phase, and if the amount of the retained γ is excessive, the predetermined value of YS cannot be obtained; hence, the upper limit of the retained γ is defined to be 12.0% (see No. 43 in Table 2B described later). The lower limit of the volume fraction V of the retained-γ phase is preferably 4.0% or more, and more preferably 6.0% or more. The upper limit thereof is preferably 11.5% or less, and more preferably 11.0% or less.
Furthermore, the volume fraction V of the retained-γ phase in the entire microstructures existing at -196°C is controlled to be 4.0% or more, thereby the percent brittle fracture can be maintained to a good level of 50% or less even at -233°C that is lower than -196°C described above. When such effects are intended to be further exhibited, the lower limit of the volume fraction V of the retained-γ phase is more preferably 6.0% or more while the preferred upper limit thereof is the same as that described above.
In the thick steel plate of the invention, it is important to control the volume fraction V of the retained-γ phase in the microstructures existing at -196°C, and any of microstructures other than the retained γ may exist without limitation as long as the microstructure normally exists in the thick steel plate. Examples of the microstructures other than the retained γ include bainite, martensite, and carbide such as cementite.
Furthermore, in the thick steel plate of the invention, inclusions, each having the equivalent circle diameter of more than 1.0 µm, are controlled such that the average roundness A of the inclusions satisfies A ≤ 1.8, and the value B represented by Formula (1) satisfies 1.3 or more. $B = V^{2 / 3} / A$
Here, "equivalent circle diameter" means a diameter that is obtained in such a manner that size of each inclusion is focused, a circle having area equal to area of the inclusion is assumed, and diameter of the circle is determined.
In the invention, inclusions each having the equivalent circle diameter of more than 1.0 µm are focused because it has been found that such inclusions promote propagation of brittle fracture. Specifically, the inclusions that promote brittle fracture must be decreased in order to improve the percent brittle fracture at ultralow temperature while predetermined high-strength is maintained. From the results of investigation of the inventors, however, it has been found that if the average roundness A of the inclusions increases, the desired ultralow-temperature toughness cannot be obtained even if the volume fraction V of the retained-γ phase at -196°C is controlled to be within the above-described range (see Nos. 33, 35, and 36 in Table 2B described later). The average roundness A of the inclusions is better as it is smaller, and is preferably 1.7 or less, and more preferably 1.5 or less. The average roundness A is most preferably 1. In the invention, average size (average equivalent circle diameter) of the inclusions each having the equivalent circle diameter of more than 1.0 µm is roughly 2.0 µm or less.
The inclusions can be determined by the method mentioned in Examples described later. The invention does not limit an inclusion type of each inclusion having the equivalent circle diameter of more than 1.0 µm. This is because occurrence of brittle fracture is greatly affected by size (equivalent circle diameter) of the inclusion rather than a type of the inclusion. Examples of the type of the inclusion include particles of a single substance of oxide, sulfide, nitride, or oxynitride, particles of a compound of at least two of such single particle substances, and composite particles including such single substance particles combined with another element.
As shown in the invention, the average roundness A of the coarse inclusions each having the equivalent circle diameter of more than 1.0 µm is controlled to be 1.8 or less, thereby ultralow-temperature toughness is improved while the predetermined strength is maintained. While the mechanism of this is not clear in detail, it is estimated as follows. Each of inclusions typically has higher hardness than a matrix, and therefore stress concentration tends to occur thereon. As a result, the inclusion often serves as an origin of brittle fracture. It is thus considered that as a shape of such an inclusion is more distorted, local stress concentration around the inclusion is further promoted, and therefore brittle fracture is further easily induced. It is therefore estimated that when the distorted inclusions are reduced (i.e., when the average roundness A of the inclusions is controlled to be 1.8 or less and controlled to be close to a round (A=1) as much as possible), occurrence of stress concentration is avoided, and the ultralow-temperature toughness is improved.
Furthermore, in the invention, it is necessary that the average roundness A of the inclusions is controlled as above, and that the value B represented by Formula (1) satisfies value B ≥ 1.3.
The value B is a parameter to decrease the percent brittle fracture at ultralow temperature. As shown in Formula (1), the value B is calculated in terms of a relationship between the average roundness A of the inclusions and the volume fraction V of the retained austenite (retained γ) phase existing at -196°C. The details of introduction of the value B are now described.
It is known that the percent brittle fracture increases with increase in number of origins of brittle fracture or with decrease in resistance against propagation of brittle fracture. A coarse inclusion in general tends to become an origin of brittle fracture, and the inventors have found the following. That is, as the roundness of the coarse inclusion increases, i.e., as the coarse inclusion has a shape more distorted from the round (A=1), the coarse inclusion more easily serves as an origin of brittle fracture. In addition, as the amount of the retained γ increases, the retained γ further acts as a resistance against propagation of brittle fracture. According to such findings, the inventors have experimentally obtained a contribution rate of each of the average roundness A and the volume fraction V to the percent brittle fracture in an ultralow temperature range based on many basic experiments. As a result, they have found that the value B represented by Formula (1) is a useful parameter to evaluate the ultralow-temperature toughness. As shown in Examples described later, the value B is controlled to be 1.3 or more while the volume fraction V of the retained-γ phase and a form (average roundness) of each coarse inclusion having the equivalent circle diameter of more than 1.0 µm are maintained. This exclusively allows both strength and percent brittle fracture at each of -196°C and -233°C to be obtained at a high level.
The value B is preferably 1.6 or more, and more preferably 1.8 or more. The value B is preferably larger in light of ultralow-temperature toughness, and the upper limit thereof is not specifically limited. However, as described above, if the volume fraction V of the retained γ excessively increases, the predetermined value of YS cannot be obtained. Hence, the upper limit of the volume fraction V of the retained γ is limited to 12. 0%. In consideration of this, the upper limit of the value B is substantially limited to 5.2 (= 12.0^2/3/1) (the volume fraction V of the retained γ = 12.0% and the average roundness A = 1 are substituted into the computational expression of the value B). In consideration of a balance of strength and toughness, the value B is more preferably 3.0 or less.
A method of manufacturing the thick steel plate of the invention is now described.
Distinctive characteristics of the manufacturing method according to the invention are the following (A) and (B).

(A) The amount of free oxygen [O] before adding Al is controlled to be 100 ppm or less in a stage of molten steel, cooling time (t1) from 1450 to 1500°C in casting is controlled to be 300 sec or less (a value at a half position of the slab thickness t), and cooling time (t2) from 1300 to 1200°C in casting is controlled to be 680 sec or less (a value at a quarter position of the slab thickness t). According to the technique (A), particularly the average roundness A of the inclusions can be decreased to a level within a predetermined range.
(B) After hot rolling, a slab is heated and held within a temperature range from Ac1 to Ac3, and is then water-cooled. Subsequently, the slab is tempered for 10 to 60 min within a temperature range from 520°C to Ac1, and is then air-cooled or water-cooled. According to the technique (B), particularly the volume fraction of the retained-γ phase existing at -196°C can be appropriately controlled.

The value B defined in the invention is a parameter on both the average roundness of the inclusions and the volume fraction of the retained γ; hence, appropriately controlling the techniques (A) and (B) makes it possible to control the value B to be within a predetermined range.
In terms of comparison with the above-described existing techniques, it is the most distinctive characteristic that t1 and t2 are each controlled in the technique (A).
Each step is now described in detail.

(Melting Step)

In the invention, a method of adding Al is particularly noticed. This is because of the following fact. The inclusions, each having the equivalent circle diameter of more than 1.0 µm, to be controlled in the invention are formed through composite formation of secondary inclusions such as oxides or sulfides nucleating on Al-based inclusions formed in molten metal during cooling. The Al-based inclusions are easily coarsened through aggregation or coalescence, and thus each easily have a distorted shape having a large roundness.
First, in adding Al as a deoxidizer into molten steel, the amount of free oxygen (the amount of dissolved oxygen, which may be abbreviated as [O] amount hereinafter) before adding Al is controlled to be 100 ppm or less. If the [O] amount exceeds 100 ppm, an increased number of Al-based inclusions are formed during addition of Al, and the roundness of each inclusion exceeds the predetermined range (see No. 33 in Table 2B described later). The [O] amount, which is better as it is smaller, is preferably 80 ppm or less, and more preferably 50 ppm or less. The lower limit of the [O] amount is not particularly limited in light of controlling the average roundness of the inclusions.
Examples of a method of controlling the [O] amount as described above include deoxidizing the molten steel by adding deoxidizing elements of Mn and Si into the molten steel. When deoxidizers such as Ti, Ca, REM, and Zr are contained as optional components in addition to the above-described elements, the [O] amount can be controlled through addition of such elements.
To control the Al-based inclusions, controlling the [O] amount before adding Al is important regardless of adding order of Al and other deoxidizing elements. However, if Al is added with high [O] amount, temperature of the molten steel increases due to an oxidation reaction, which is dangerous in operation; hence, Si and Mn are preferably added prior to Al. The optional components such as Ti are preferably added into the molten steel after adding Al.
Subsequently, casting is started. While the temperature range in casting is roughly 1650°C or less, the invention has revealed that it is particularly important that cooling time (t1) in a temperature range from 1450 to 1500°C is controlled to be 300 sec or less, and cooling time (t2) from 1300 to 1200°C is controlled to be 680 sec or less, and thereby the average roundness of the inclusions each having the equivalent circle diameter of more than 1.0 µm is appropriately controlled. This is described in detail below.
First, the cooling time (t1) in the temperature range from 1450 to 1500°C is controlled to be 300 sec or less. If the t1 exceeds 300 sec, composite formation of the secondary inclusions nucleating on the Al-based inclusions is promoted, and each inclusion having the equivalent circle diameter of more than 1.0 µm has a distorted shape, leading to increase in average roundness and decrease in value B. As a result, the desired ultralow-temperature toughness is not obtained (see Nos. 34 and 35 in Table 2B described later). From such a viewpoint, t1 is better as it is shorter, and is preferably 290 sec or less, and more preferably 280 sec or less. The lower limit of t1 is not specifically limited from such a viewpoint.
In the invention, attention is particularly focused on the temperature range from 1450 to 1500°C among temperature ranges in casting. This is because that temperature range is a temperature region during which the molten steel is progressively solidified and each component is increasingly segregated into the molten steel, so that growth of the inclusions is accelerated.
The temperature range from 1450 to 1500°C corresponds to temperature of the center (t/2) of the slab thickness t. The reason for this is as follows. As described before, the oxide-based secondary inclusions are compositely formed mainly in molten metal; hence, cooling time of a molten metal region must be controlled. However, since the temperature range from 1450 to 1500°C is a temperature region during which the molten metal is progressively solidified, cooling time of the molten metal region may not be accurately measured due to solidification of the molten metal depending on temperature measurement positions. In the invention, therefore, cooling time is measured at the t/2 position at which the molten metal remains even at the lowest temperature. The temperature of the center of the slab thickness can be measured by inserting a thermocouple into a mold.
Subsequently, cooling time (t2) from 1300 to 1200°C is controlled to be 680 sec or less. If the t2 exceeds 680 sec, composite formation of mainly the sulfide-based secondary inclusions on the Al-based inclusions is promoted, and the average roundness of the inclusions also increases (see No. 36 in Table 2B described later). From such a viewpoint, t2 is more preferable as it is shorter since shorter t2 leads to a shape closer to a round. Preferably, t2 is 650 sec or less. More preferably, t2 is 600 sec or less. However, excessively short t2 increases a cooling load; hence, t2 is recommended to be roughly 400 sec or more.
The temperature range from 1300 to 1200°C corresponds to the temperature of the quarter part (t/4) of the slab thickness t. The reason for this is as follows. The cooling time from 1300 to 1200°C is controlled to control the sulfide-based secondary inclusions that are compositely formed mainly in solid iron. Since solidification has been substantially completed in such a temperature region, measurement of cooling time is conducted at the t/4 position at which the percent brittle fracture is measured. The temperature of the t/4 part of the slab thickness can be measured by inserting a thermocouple into a mold.
In the invention, the cooling time (t1) in the temperature range from 1450 to 1500°C and the cooling time (t2) from 1300 to 1200°C should be controlled as described above, and any approach may be taken to achieve this without limitation. For example, for t1, the molten metal may be cooled at uniform velocity, i.e., at an average cooling rate of about 0.17 °C/sec or less, in the above-described temperature range such that cooling time in the temperature range is 300 sec or less. Alternatively, the molten metal may be cooled at nonuniform velocity such that cooling time in the temperature range is 300 sec or less. For t2, similar approach may be made.
In the invention, the molten metal may be cooled in any of temperature ranges in casting other than the above-described temperature ranges in any method without limitation, i.e., may be cooled by a typical method (air cooling or water cooling).
The molten steel is cast into a slab as described above, and then the slab is hot-rolled and subjected to heat treatment.
The hot rolling step may be performed by any of typically used methods without limitation so that a predetermined thickness is given. Specifically, the slab is heated for 1 to 4 hr at about 1100°C, and then (finish rolling) temperature and rolling reduction are appropriately adjusted.
After the hot rolling, the slab is heated into the temperature range (TL) from the Ac1 point to the Ac3 point and held therein, and is then water-cooled. This treatment corresponds to the L treatment in the above-described existing technique, and allows the retained γ to stably exist at -196°C within a range of the predetermined amount.
In detail, the slab is heated to a temperature (TL) of the two-phase region [ferrite (α)-γ] between the Ac1 point and the Ac3 point. The slab is heated into such a temperature region, thereby alloy elements such as Ni are concentrated in the formed γ phase, resulting in formation of a metastable retained-γ phase that exists metastably at room temperature. A temperature less than the Ac1 point or more than the Ac3 point eventually prevents the retained-γ phase from sufficiently existing at -196°C (see Nos. 37 and 38 in Table 2B described later). Preferred heating temperature is roughly 660 to 710°C.
The heating time at the two-phase region temperature (holding time, tL) is preferably roughly 10 to 50 min. In the heating time of less than 10 min, the alloy elements are not sufficiently concentrated into the γ phase. In the heating time of more than 50 min, the α phase is annealed, and strength is lowered. A preferred heating time is roughly 15 to 30 min.
Furthermore, the heating time is controlled to be 15 min or more, thereby at least 4.0% of the volume fraction of the retained-γ phase at -196°C is obtained. Consequently, the percent brittle fracture is 50% or less at -233°C, i.e., good toughness is obtained even under a further ultralow temperature. The upper limit of preferred heating time is the same as that described above (30 min or less).
Subsequently, the slab is water-cooled to room temperature and then tempered. The tempering is performed for 10 to 60 min (t3) in a temperature range (T3) from 520°C to A_c1. Consequently, C is concentrated in the metastable retained-γ during tempering, which increases stability of the metastable retained-γ phase, resulting in formation of the retained γ phase that stably exists even at -196°C. If the tempering temperature T3 is lower than 520°C, the metastable retained-γ phase formed during holding of the two-phase coexistence region is decomposed into the α phase and a cementite phase, and thus the retained-γ phase at -196°C cannot be sufficiently given (see No. 41 in Table 2B described later). For the tempering temperature T3 of more than the Ac1 point or the tempering time t3 of less than 10 min, C is not sufficiently concentrated in the metastable retained-γ phase, and the desired amount of retained γ at -196°C cannot be given (see No. 55 (an example of short t3) in Table 2 described later). For the tempering time t3 of more than 60 min, the retained γ at -196°C is excessively formed, and the predetermined strength cannot be obtained (see No. 43 in Table 2 described later).
A preferred tempering condition includes tempering temperature T3 of 570 to 620°C, and tempering time t3 of 15 to 45 min (more preferably 15 to 35 min, and most preferably 15 to 25 min).
The slab is tempered as described above and then cooled to room temperature. The slab may be cooled by any method without limitation, i.e., may be cooled by air cooling or water cooling.
In this specification, the Ac1 point and the Ac3 point are calculated according to the following formulas (from "Modern Metallurgy Course, Material Edition 4, Steels", The Japan Institute of Metals and Materials). $Ac 1 point = 723 - 10.7 \times [Mn] - 16.9 \times [Ni] + 29.1 \times [Si] + 16.9 \times [Cr] + 290 \times [As] + 6.38 \times [W]$
$Ac 3 point = 910 - 203 \times {[C]}^{1 / 2} - 15.2 \times [Ni] + 44.7 \times [Si] + 104 \times [V] + 31.5 \times [Mo] + 13.1 \times [W]$
In the formulas, [] means concentration (mass%) of an alloy element in steel. In the invention, As and W are not contained in the steel composition. In the Formulas, therefore, calculation is made assuming that each of [As] and [W] is 0%.

Examples

Although the invention is now described in detail with some Examples, the invention should not be limited thereto, and modifications or alterations thereof may be made within the scope without departing from the gist described before and later, all of which are included in the technical scope of the invention.

Example 1

Test steel samples having compositions shown in Table 1, in each of which the remainder consists of iron and inevitable impurities and the unit is mass percent, were each melted using a vacuum induction furnace (150 kg VIF) and cast, and then an ingot 600 mm long, 150 mm wide, and 150 mm high was fabricated through hot forging. In Example 1, mish metal containing about 50% Ce and about 25% La was used as REM. Adding order of the deoxidizing elements was as follows. That is, when the optional components were not contained, Si and Mn were added together (simultaneous adding) and then Al was added. When the optional components of Ti, REM, Zr, and Ca were contained, Si and Mn were added together (simultaneous adding), and then Al and Ti were added in this order, and then REM, Zr, and Ca were added together (simultaneous adding). In Example 1, time from addition of Al to start of casting was about 10 min in each case (not shown in Tables).
In Table 2, [O] denotes the amount of dissolved oxygen (ppm) before adding Al, t1 denotes cooling time (sec) from 1450 to 1500°C in casting, and t2 denotes cooling time (sec) from 1300 to 1200°C in casting. Each type of steel was cooled in each of such temperature regions by air cooling or water cooling while cooling time was controlled to be the above-described cooling time.
Subsequently, each of the ingots was heated for 1 to 4 hr at 1100°C, and then the ingot was rolled into a thickness of 75 mm at a temperature of 830°C or more, and was rolled at a final rolling temperature of 780°C and then water-cooled, thereby a thick steel plate having a thickness of 25 mm was produced. The steel plates produced in this way were heated at respective temperatures (TL in Table 2) shown in Table 2, and were then held for 5 to 60 min while being heated (see tL in Table 2), and were then water-cooled to room temperature. Subsequently, the steel plates were tempered as shown in Table 2 (T3 is tempering temperature, and t3 is tempering time), and were then air-cooled or water-cooled to room temperature.
According to the procedures described below, the thick steel plates produced in this way were each evaluated in average roundness A of the inclusions each having the equivalent circle diameter of more than 1.0 µm, volume fraction (%) of the retained-γ phase existing at -196°C, tensile properties (tensile strength TS and yield strength YS), and ultralow-temperature toughness (percent brittle fracture in the C direction at -196°C or -233°C).

(1) Measurement of average roundness A of Inclusions Each Having Equivalent Circle Diameter of More Than 1.0 µm

Each of the steel plates was mirror-polished at its t/4 position (t: thickness), and was subjected to photography in four viewing fields at 400 magnifications with a light microscope. Area of one viewing field was 0.04 mm², and the total area of the four viewing fields was 0.15 mm². Inclusions observed in the four viewing fields were subjected to image analysis with "Image-Pro Plus" from Media Cybernetics. In addition, roundness of each inclusion having an equivalent circle diameter of more than 1.0 µm was calculated according to the following formula, and an average of the calculated values was defined as the average roundness A of the inclusions. When an inclusion has a round shape, the roundness calculated by the formula is 1. As the inclusion has a more distorted shape, the roundness calculated by the formula has a larger value.
$Roudness = L^{2} / 4 π / S$
In the formula, L is the perimeter (µm) of the inclusion, and S is area (µm²) of the inclusion.
In Example 1, the inclusions each having an equivalent circle diameter of more than 1.0 µm were observed in number density of about 200 to 300 per square millimeter.

(2) Volume Fraction of Retained-γ phase Existing at -196°C

A test specimen 10 by 10 by 55 mm was taken from the t/4 position of each steel plate. The test specimen was held for 5 min at liquid nitrogen temperature (-196°C), and was then subjected to X-ray diffraction measurement by a two-dimensional micro-part X-ray diffractometer "RINT-RAPID II" from Rigaku Corporation. Subsequently, an integrated intensity ratio was obtained on each of peaks of the lattice planes of (110), (200), (211), and (220) of the ferrite phase, and on each of peaks of the lattice planes of (111), (200), (220), and (311) of the retained-γ phase. Based on such integrated intensity ratios, the volume fraction of each of the lattice planes of (111), (200), (220), and (311) of the retained-γ phase was calculated, and an average of such volume fractions was obtained and defined as "volume fraction (%) of retained γ".

(3) Measurement of Tensile Properties (Tensile Strength TS and Yield Strength YS)

A JIS Z 2241-4 test specimen was taken parallel to the C direction from the t/4 position of each steel plate, and was subjected to a tensile test by a method according to JIS Z 2241 to determine tensile strength TS and yield strength YS. In Example 1, steel having TS of more than 690 MPa and YS of more than 590 MPa was determined to be good in strength of base metal.

(4) Measurement of Ultralow-Temperature Toughness (Percent Brittle Fracture in C Direction)

Three Charpy impact test specimens (JIS Z 2242 V-notch test specimens) were taken parallel to the C direction from the position of t/4 (t: thickness) at W/4 (W: width) and from the position of t/4 at W/2, and were subjected to measurement of percent brittle fracture (%) at -196°C by the method according to JIS Z 2242, and an average of the three measured values was calculated for each position. One average having worse properties (i.e., larger percent brittle fracture) was selected from the two averages calculated in this way. In Example 1, when the selected average has a value of 10% or less, the relevant steel plate was determined to have good ultralow-temperature toughness.

Table 2 collectively shows results of such measurements. Tables 1 and 2 also show the Ac1 point and the Ac3 point for reference.

Table 2A

No.	Al → Casting → Hot rolling → Tempering										Average round-ness (A)	Retained γ (%) V	Ultralow-temperature toughness	Tensile properties
No.	[O] (ppm)	t1 (sec)	t2 (sec)	TL (°C)	tL (min)	A_c1	A_c3	T3 (°C)	t3 (min)	Cooling method after tempering	Average round-ness (A)	Retained γ (%) V	Percent brittle fracture at -196°C (%)	YS	TS
1	45	261	500	660	10	622	775	600	20	Air cooling	1.5	3.8	5	641	721
2	34	268	530	660	15	598	762	580	20	Water coding	1.4	5.4	0	633	722
3	45	258	500	660	15	620	784	600	20	Air cooling	1.5	6.5	0	639	731
4	42	255	530	660	15	633	784	600	30	Water cooling	1.4	7.8	0	628	728
5	45	260	530	660	10	632	779	600	30	Water cooling	1.5	3.8	7	625	742
6	40	271	500	660	15	617	775	600	20	Air cooling	1.5	5.5	0	630	725
7	42	255	500	660	10	613	777	600	20	Air cooling	1.5	3.6	5	621	715
8	48	246	600	660	15	618	778	600	20	Water cooling	1.6	6.1	0	638	721
9	38	195	530	700	15	632	780	620	35	Air cooling	1.3	7.2	0	645	750
10	42	268	530	700	15	611	782	600	20	Air cooling	1.4	6.6	0	638	746
11	41	236	650	660	15	633	796	600	20	Air cooling	1.5	6.3	0	650	778
12	37	165	500	700	15	599	757	580	40	Air cooling	1.3	10.1	0	622	749
13	32	213	530	660	15	622	771	600	20	Water coding	1.3	5.7	2	673	796
14	83	244	530	660	10	618	817	600	40	Air cooling	1.8	3.9	8	619	695
15	29	241	500	660	50	614	802	600	20	Air cooling	1.4	11.6	0	608	716
16	61	255	530	660	15	622	796	600	15	Air cooling	1.7	3.5	7	681	750
17	31	220	530	660	15	618	792	570	15	Air cooling	1.4	5.3	3	659	751
18	45	286	530	660	15	610	804	600	15	Air cooling	1.7	4.1	5	736	791
19	36	245	530	660	15	621	787	580	40	Water cooling	1.5	8.9	2	634	715
20	41	190	530	660	10	638	809	600	25	Air cooling	1.8	3.8	3	681	758
21	35	269	530	710	10	630	805	600	20	Air cooling	1.4	2.7	8	655	747
22	41	257	600	710	10	627	808	620	20	Air cooling	1.5	3.9	0	652	752
23	39	287	530	660	15	594	754	570	50	Water cooling	1.8	11.8	2	606	793
24	30	275	660	710	15	615	825	600	20	Water cooling	1.7	7.0	3	664	783
25	35	249	530	660	15	631	792	600	401	Water cooling	1.5	8.4	2	661	762
26	40	216	480	660	15	630	804	600	55	Air cooling	1.4	11.6	0	603	770
27	36	254	530	660	15	644	795	600	20	Air cooling	1.4	4.6	0	684	774
28	38	255	530	700	15	636	798	600	20	Water cooling	1.5	5.3	0	738	801
29	42	273	480	660	15	611	781	600	30	Water cooling	1.3	7.2	0	744	805
30	46	258	530	660	15	636	826	600	20	Water cooling	1.5	5.9	0	759	819
31	34	210	530	660	15	608	760	600	25	Air cooling	1.4	6.8	2	699	790
32	42	260	500	660	15	593	775	570	20	Water cooling	1.5	7.3	0	648	736

Table 2B

No.	Al → Casting → Hot rolling → Tempering										Average round-ness (A)	Retained γ (%) V	Ultralow-temperature toughness	Tensile properties
No.	[O] (ppm)	t1 (sec)	t2 (sec)	TL (°C)	tL (min)	A_c1	A_c3	T3 (°C)	t3 (min)	Cooling method after tempering	Average round-ness (A)	Retained γ (%) V	Percent brittle fracture at -196°C (%)	YS	TS
33	103	259	500	660	15	623	795	600	20	Water cooling	2.0	4.9	15	631	721
34	39	303	530	660	10	629	795	570	40	Air cooling	1.8	3.0	12	635	722
35	36	310	530	660	15	626	796	570	20	Air cooling	1.9	6.3	18	649	720
36	34	266	720	660	15	612	751	600	15	Water cooling	1.9	8.2	17	671	803
37	40	268	600	590	15	611	765	600	20	Water cooling	1.5	1.8	17	692	762
38	31	265	530	760	15	598	755	570	25	Air cooling	1.4	1.8	20	731	780
39	42	244	660	660	5	620	778	600	25	Air cooling	1.4	1.9	13	646	715
40	45	258	500	660	60	621	771	600	20	Air cooling	1.5	12.3	0	545	672
41	36	266	500	660	15	607	774	500	25	Air cooling	1.5	1.4	22	670	718
42	29	239	500	660	15	581	762	580	20	Air cooling	1.3	1.8	18	732	776
43	41	262	500	660	15	622	791	600	65	Water coding	1.5	12.9	0	585	735
44	78	271	530	660	15	635	783	620	20	Water coding	1.8	1.6	22	633	692
45	42	240	450	660	15	604	787	600	25	Air cooling	1.5	5.6	12	671	763
46	45	212	530	660	15	636	859	620	25	Air cooling	2.0	1.4	28	635	678
47	43	238	530	660	15	610	762	600	40	Water coding	1.9	11.3	25	702	821
48	30	245	530	660	15	608	770	600	20	Water coding	1.4	8.2	23	672	795
49	47	197	530	660	15	645	785	620	20	Air cooling	1.4	4.7	25	715	790
50	39	259	500	660	15	592	760	570	20	Water cooling	1.5	10.0	28	710	804
51	38	270	600	700	15	621	821	600	20	Water coding	1.5	6.1	12	702	818
52	40	251	450	710	15	599	771	570	15	Air cooling	1.6	4.2	13	733	789
53	42	263	530	660	15	633	848	600	35	Water coding	1.5	8.8	12	728	812
54	45	247	530	660	15	619	778	600	20	Water cooling	1.5	7.6	13	712	800
55	34	259	530	660	15	613	767	610	5	water cooling	1.4	1.5	12	642	741

The following consideration can be made from Table 2.
Nos. 1 to 32 in Table 2A are examples satisfying all the requirements of the invention, in each of which the resultant steel plate had good ultralow-temperature toughness (in detail, the average of the percent brittle fracture in the C direction of 10% or less) even if base metal had high strength.
In contrast, Nos. 33 to 41, 43, and 55 in Table 2B each dissatisfy at least one of the preferred manufacturing conditions of the invention, each of which is therefore a comparative example that did not satisfy a requirement of the invention, and failed to have the desired properties.
Specifically, No. 33 is an example in which No. 33 steel in Table 1B was used, the steel having a composition satisfying the requirements of the invention, but since a large amount of dissolved oxygen [O] existed before adding Al, the average roundness A of the inclusions was also large. As a result, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
No. 34 is an example in which No. 34 steel in Table 1B was used, the steel having a composition satisfying the requirements of the invention, but since the cooling time (t1) from 1500 to 1450°C in casting was long, the value B was below the predetermined range. As a result, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
No. 35 is an example in which No. 35 steel in Table 1B was used, the steel containing a large amount of P, and the cooling time (t1) from 1500 to 1450°C in casting was long, and therefore the average roundness A of the inclusions was large. As a result, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
No. 36 is an example in which No. 36 steel in Table 1B was used, the steel containing a large amount of C, and the cooling time (t2) from 1300 to 1200°C in casting was long, and therefore the average roundness A of the inclusions was large. As a result, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
No. 37 is an example in which No. 37 steel in Table 1B was used, the steel having a composition satisfying the requirements of the invention, but since the steel plate was heated at a temperature below the two-phase region temperature (TL), the amount of retained γ was insufficient. As a result, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
No. 38 is an example in which No. 38 steel in Table 1B was used, the steel containing a large amount of Si, and the steel plate was heated at a temperature above the two-phase region temperature (TL), and therefore the amount of retained γ was insufficient. As a result, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
No. 39 is an example in which No. 39 steel in Table 1B was used, the steel having a composition satisfying the requirements of the invention, but since the heating and holding time (tL) at the two-phase region temperature (TL) was short, the amount of retained γ was insufficient. As a result, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
No. 40 is an example in which No. 40 steel in Table 1B was used, the steel having a composition satisfying the requirements of the invention, but since the heating and holding time (tL) at the two-phase region temperature (TL) was long, the amount of retained γ was too large. As a result, yield strength YS and tensile strength TS were lowered, and the desired base metal strength was not obtained.
No. 41 is an example in which No. 41 steel in Table 1B was used, the steel having a composition satisfying the requirements of the invention, but since tempering temperature (T3) was low, the amount of retained γ was insufficient. As a result, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
No. 43 is an example in which No. 43 steel in Table 1B was used, the steel having a composition satisfying the requirements of the invention, but since tempering time (t3) was long, the amount of retained γ was too large. As a result, yield strength YS was lowered, and the desired base metal strength was not obtained.
No. 55 is an example in which No. 55 steel in Table 1B was used, the steel having a composition satisfying the requirements of the invention, but since tempering time (t3) was short, the amount of retained γ was insufficient. As a result, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
Nos. 42 and 44 to 54 are each a comparative example in which the steel plate was fabricated according to the method of the invention except that used steel did not satisfy the requirements of steel composition.
In detail, No. 42 is an example in which No. 42 steel in Table 1B was used, the steel containing a large amount of Mn, and therefore the amount of retained γ was insufficient. As a result, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
No. 44 is an example in which No. 44 steel in Table 1B was used, the steel containing a small amount of Mn, and therefore the amount of retained γ was insufficient. As a result, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
No. 45 is an example in which No. 45 steel in Table 1B was used, the steel containing a large amount of S. Hence, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
No. 46 is an example in which No. 46 steel in Table 1B was used, the steel containing a small amount of C, a large amount of Al, and a small amount of Ni, and therefore the average roundness A of the inclusions was large, and the amount of retained γ was insufficient. As a result, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained. Furthermore, TS was low.
No. 47 is an example in which No. 47 steel in Table 1B was used, the steel containing a small amount of Al and a large amount of N, and therefore the average roundness A of the inclusions was large. As a result, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
No. 48 is an example in which No. 48 steel in Table 1B was used, the steel containing a large amount of each of Cu and Ca as the optional components. As a result, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
No. 49 is an example in which No. 49 steel in Table 1B was used, the steel containing a large amount of each of Cr and Zr as the optional components. As a result, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
No. 50 is an example in which No. 50 steel in Table 1B was used, the steel containing a large amount of each of Nb and REM as the optional components. As a result, the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
No. 51 is an example in which No. 51 steel in Table 1B was used, the steel containing a large amount of Mo as the optional component, and therefore the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
In No. 52, No. 52 steel in Table 1B was used, the steel containing a large amount of Ti as the optional component, and therefore the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
In No. 53, No. 53 steel in Table 1B was used, the steel containing a large amount of V as the optional component, and therefore the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.
In No. 54, No. 54 steel in Table 1B was used, the steel containing a large amount of B as the optional component, and therefore the percent brittle fracture was high, and the desired ultralow-temperature toughness was not obtained.

Example 2

In Example 2, part of steel plate (each being an example of the invention) used in Example 1 were evaluated in percent brittle fracture at -233°C.
Specifically, for each number of steel plate in Table 3 (No. in Table 3 corresponds to No. in each of Tables 1 and 2), three specimens were taken from the position of t/4 at W/4, and were each subjected to a Charpy impact test at -233°C by a method described below, and an average of the percent brittle fracture was determined. In Example 2, a steel plate showing a percent brittle fracture of 50% or less was determined to have good percent brittle fracture at -233°C.
"High Pressure Gas", vol.24, p.181, "Ultralow-Temperature Impact Test of Austenitic Stainless Cast Steel".

Table 3 shows results of such determination.

Table 3

No.	Retained γ (%)	Ultralow-temperature toughness
No.	Retained γ (%)	Percent brittle fracture at -196°C (%)	Percent brittle fracture at -233°C (%)
1	3.8	5	53
3	6.5	0	43
4	7.8	0	28
5	3.8	7	55
6	5.5	0	38
7	3.6	5	53
14	3.9	8	70
15	11.6	0	12
19	8.9	2	40
20	3.8	3	52
21	2.7	8	67
22	3.9	0	60
24	7.0	3	40

Each of Nos. 3, 4, 6, 15, 19, and 24 is an example in which the heating and holding time (tL) at the two-phase region temperature was controlled to be 15 min or more (see Table 2A), and at least 4.0% of retained-γ phase was given. As a result, good percent brittle fracture was shown not only at -196°C but also at a further low temperature, -233°C, i.e., extremely good ultralow-temperature toughness was obtained.
Although the invention has been described in detail with reference to specific embodiments, it should be understood by those skilled in the art that various alterations and modifications thereof may be made without departing from the spirit and the scope of the invention.
The present application is based on Japanese patent application ( JP-2012-184593 ) filed on August 23, 2012, the content of which is hereby incorporated by reference.

Industrial Applicability

The thick steel plate of the invention is good in ultralow-temperature toughness, and is particularly useful as steel for a storage tank, a transport ship, and the like for liquefied natural gas (LNG).

Claims

A thick steel plate having good ultralow-temperature toughness, comprising:
by mass percent,

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%; and

N: 0.010% or less (not including 0%),

with the remainder consisting of iron and inevitable impurities,

wherein average roundness (A) of inclusions, each having an equivalent circle diameter of more than 1.0 µm, is 1.8 or less,

a volume fraction (V) of a retained austenite phase existing at -196°C satisfies 2.0 to 12.0%, and

a value B represented by Formula (1) is 1.3 or more. $B = V^{2 / 3} / A$
The thick steel plate according to claim 1, wherein the retained austenite phase existing at -196°C is 4.0 to 12.0% in a volume fraction.
The thick steel plate according to claim 1 or 2, further comprising at least one selected from a group consisting of
Cu: 1.00% or less (not including 0%),
Cr: 1.2% or less (not including 0%),
Mo: 1.00% or less (not including 0%),
Ti: 0.025% or less (not including 0%),
Nb: 0.10% or less (not including 0%),
V: 0.50% or less (not including 0%),
B: 0.0050% or less (not including 0%),
Ca: 0.0030% or less (not including 0%),
REM: 0.0050% or less (not including 0%), and
Zr: 0.0050% or less (not including 0%).