JP6018453B2 - High strength 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 in 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, the development of a thick steel plate excellent in cryogenic toughness is being promoted.

For example, Non-Patent Document 1 describes the effect of heat treatment in the coexistence region of α-γ2 phase on the low temperature toughness of 6% Ni steel. Specifically, by adding heat treatment (L treatment) in the α-γ two phase coexistence region (between A _{c1 and} A _c3 ) before tempering treatment, it is equivalent to 9% Ni steel subjected to normal quenching and tempering treatment. The above-mentioned ability to impart cryogenic toughness at −196 ° C .; this heat treatment also improves the toughness of the specimen in the C direction (sheet width direction); It is described that it is due to the presence of retained austenite that is stable against the load. 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) shows 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. 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.

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

Yano et al., “Effect of heat treatment in coexistence of α-γ2 phase 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. At the site where brittle fracture occurs, the energy absorbed by the steel material until the fracture is significantly reduced, and the fracture proceeds easily. Therefore, in the cryogenic toughness improvement technology, the general-purpose Charpy impact value (vE _-196 ) as well as improving the brittle fracture surface ratio to 10% or less are extremely important requirements. However, a technique that satisfies the above requirement for the brittle fracture surface ratio in a high-strength thick steel plate having a high base metal strength as described above has not yet been proposed.

  The present invention has been made in view of the above circumstances, and its purpose is to achieve cryogenic toughness (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 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 a steel plate with iron and inevitable impurities, at -196 ° C The summary is that the residual austenite phase present is 2.0% to 12.0% in volume fraction, and the average equivalent circle diameter of inclusions having an equivalent circle diameter of more than 2.0 μm is 3.5 μm or less. It is what has.

  In a preferred embodiment of the present invention, the steel sheet is such that the residual 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: 1.0% or less (excluding 0%).

  In a preferred embodiment of the present invention, the steel sheet is further selected from the group consisting of Cr: 1.20% or less (not including 0%) and Mo: 1.0% 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 (excluding 0%), Nb: 0.100% or less (not including 0%), and V: 0.50. % Or less (not including 0%).

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

  In a preferred embodiment of the present invention, the steel sheet further includes Ca: 0.0030% or less (excluding 0%), REM: 0.0050% or less (excluding 0%), and Zr: 0.005. % Or less (not including 0%).

  According to the present invention, in Ni steel having a Ni content of about 5.0 to 7.5%, even if the base metal strength is high (specifically, tensile strength TS> 690 MPa, yield strength YS> 590 MPa), − Excellent low temperature toughness at 196 ° C. or less (particularly C direction cryogenic toughness), brittle fracture surface ratio at −196 ° C. ≦ 10% (preferably, brittle fracture surface ratio at −233 ° C. ≦ 50% ) Was able to be provided.

The characteristic part of the thick steel plate according to the present invention is (Ni) at 196 ° C. in order to further improve the C-direction cryogenic toughness in Ni steel having a Ni content of about 5.0 to 7.5%. The residual austenite phase (residual γ phase) is controlled to 2.0% to 12.0% (volume fraction) [preferably controlled to 4.0% to 12.0% (volume fraction)],
(A) The average equivalent circle diameter of coarse inclusions with an equivalent circle diameter of more than 2.0 μm (hereinafter sometimes simply referred to as coarse inclusions, sometimes abbreviated as N1) is 3.5 μm or less. It is in place. 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 of −196 ° C. or lower in a Ni steel having a Ni content of 7.5% or lower. Specifically, in the present invention, the brittle fracture surface ratio at −196 ° C. in the C direction ≦ 10%, the tensile strength TS> 690 MPa, and the yield strength YS> 590 MPa satisfying all the characteristics that are excellent in cryogenic toughness. From the viewpoint of providing a thick steel plate, first, the method taught in the literature described in the prior art was 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. In addition, considering the manufacturing method comprehensively, in the molten steel stage, the amount of dissolved oxygen before deoxidation element addition is controlled, and in this molten steel, Al is added last, and the α-γ2 phase is added. After heat treatment (L treatment) in the coexistence zone (between A _{c1 and} A _c3 ), a method of tempering at a temperature below the A _c1 transformation point is recommended, which teaches that cryogenic toughness is improved. Has been. 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 was proved that the brittle fracture surface ratio at −196 ° C. in FIG.

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, (a) In a thick steel plate, in addition to the presence of a residual γ phase at −196 ° C. in the range of 2.0% to 12.0% (volume fraction), it becomes a starting point for brittle fracture. It is effective to refine the average equivalent circle diameter (N1) of coarse inclusions having an equivalent circle diameter of more than 2.0 μm to 3.5 μm or less. (A) In order to produce such a thick steel plate, In addition to the control of the dissolved oxygen amount (free O amount) before Al addition in the molten steel stage and the heat treatment (L treatment) between A _{c1 and} A _c3 after hot rolling → tempering treatment in a predetermined temperature range Further control of the molten steel stage is effective, the holding time (t1) from the addition of Al to the start of casting is set to 15 minutes or more, and the cooling time (t2) at 1450 to 1500 ° C. during casting is 300 seconds or less. I found out that it is effective to control the system.

Furthermore, (c) In the above (a), by controlling the residual γ phase existing at −196 ° C. to 4.0% to 12.0% (volume fraction), it is brittle even at a lower temperature of −233 ° C. The fracture surface ratio can be maintained at a favorable level of 50% or less, and (d) in order to produce such a thick steel plate, heat treatment between A _{c1 and} A _c3 after hot rolling (L treatment) ), It was found that holding for a predetermined time was effective, and the present invention was completed.

The "excellent cryogenic temperature toughness" as used herein, when measured vE _-196 and brittle fracture rate in a Charpy impact absorption test of C direction (sheet width direction) by the method described in the column of the examples described later The brittle fracture surface ratio at −196 ° C. ≦ 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 0.5%, N: 0.010% or less (excluding 0%), the balance being a steel plate with iron and inevitable impurities, the residual austenite phase present at -196 ° C. in the volume fraction The inclusion is characterized by the fact that the average equivalent circle diameter of inclusions having an equivalent circle diameter of more than 2.0 μm is 3.5 μm or less.

  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%. The upper limit with preferable C amount is 0.08% or less, More preferably, it is 0.06% or less.

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

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

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

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

Al: 0.005 to 0.050%
Al is a deoxidizing element. If the Al content is insufficient, the oxygen concentration in the steel increases and coarse inclusions increase, 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, the aggregation and coalescence of inclusions are promoted and the increase in inclusion size is caused, so the upper limit is made 0.050% or less. The upper limit with preferable Al amount 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. The minimum with preferable Ni amount is 5.2% or more, More preferably, it is 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 the upper limit is made 0.010% or less. The upper limit with preferable N amount is 0.006% or less, More preferably, it is 0.004% or less.

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

Cr: at least one selected from the group consisting of 1.20% or less (not including 0%) and Mo: 1.0% or less (not including 0%) Cr and Mo are both strength-enhancing elements. is there. These elements may be added alone or in combination of two kinds. In order to effectively exhibit the above action, 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, so the preferable upper limit of Cr content is 1.20% or less (more preferably 1.1% or less, even more Preferably 0.9% or less, and still more preferably 0.5% or less), and a preferable upper limit of the Mo amount is 1.0% or less (more preferably 0.8% or less, still more preferably 0.6% or less). And

Selected from the group consisting of Ti: 0.025% or less (not including 0%), Nb: 0.100% or less (not including 0%), and V: 0.50% or less (not including 0%) At least one of Ti, Nb, and V is an element that precipitates as carbonitride and increases strength. These elements may be added alone or in combination of two or more. In order to effectively exhibit the above action, it is preferable that the Ti amount is 0.005% or more, the Nb amount is 0.005% or more, and the V amount is 0.005% or more. However, if excessively added, the strength is excessively improved and the desired cryogenic toughness cannot be ensured. Therefore, the preferable upper limit of the Ti amount is 0.025% or less (more preferably 0.018% or less, Even more preferably, 0.015% or less), the preferred upper limit of the Nb amount is 0.100% or less (more preferably 0.05% or less, even more preferably 0.02% or less), the preferred upper limit of the V amount Is 0.50% or less (more preferably 0.3% or less, 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 excessively added, the strength is excessively improved and the desired cryogenic toughness cannot be ensured. Therefore, the preferable upper limit of the B amount is 0.0050% or less (more preferably 0.0030% or less, even more Preferably it is 0.0020% or less.

From Ca: 0.0030% or less (excluding 0%), REM (rare earth element): 0.0050% or less (not including 0%), and Zr: 0.005% or less (not including 0%) At least one selected from the group consisting of Ca, REM, and Zr is a deoxidizing element, and the addition reduces the oxygen concentration in the steel and reduces coarse inclusions. 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 excessively added, coarse inclusions increase and cryogenic toughness decreases, so the preferable upper limit of the Ca amount is 0.0030% or less (more preferably 0.0025% or less), and the REM amount is preferable. The upper limit is made 0.0050% or less (more preferably 0.0040% or less), and the preferred upper limit of the amount of Zr is made 0.005% 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 residual γ phase existing at −196 ° C. satisfies the volume fraction of 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 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 is excessive, YS cannot secure a predetermined value, so the upper limit is set to 12.0% (Table 2 described later) No. 39). Regarding the volume fraction of the residual γ phase, the preferable lower limit is 4.0% or more, the more preferable lower limit is 6.0% or more, the preferable upper limit is 11.5% or less, and the more preferable upper limit is 11.0% or less. .

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

  Furthermore, the thick steel plate of the present invention satisfies the average equivalent circle diameter N1 of inclusions (coarse inclusions) having an equivalent circle diameter of more than 2.0 μm of 3.5 μm or less. In contrast to the above-described prior art, the thick steel plate of the present invention has the greatest feature in that the coarse inclusions are refined to 3.5 μm or less.

  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.

That is, according to the examination results of the present inventors, coarse inclusions having an equivalent circle diameter of more than 2.0 μm become the starting point of brittle fracture, and when the average size of the coarse inclusions (average equivalent circle diameter N1) increases, Even if the volume fraction of the residual γ phase at −196 ° C. was controlled within the above range, it was found that the desired cryogenic toughness could not be realized (Nos. 33 to 35, 45 to 49 in Table 2 described later). See). The average equivalent circle diameter of N1 is preferably as small as possible, preferably 3.2 μm or less, and more preferably 3.0 μm or less. In the present invention, there are about 10 to 100 / mm ² inclusions having an equivalent circle diameter of more than 2.0 μm.

  The said inclusion can be measured by the method as described in the Example mentioned later. Here, the kind of “inclusion” in inclusions having an equivalent circle diameter of more than 2.0 μm is not particularly limited in the present invention. The occurrence of brittle fracture is because the size of inclusions (average equivalent circle diameter) has the greatest influence, not the type of inclusions. 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.

From the viewpoint of inclusion control only, similar technology is disclosed in Patent Document 3 described above, but the direction of inclusion control is greatly different from the present invention. That is, in the above-mentioned Patent Document 3, focusing on Mg in particular, the coarsening of austenite grains at high temperature is suppressed by dispersing a large number of fine Mg-containing oxide particles having a size of 2 μm or less, and toughness is improved. On the other hand, in the present invention, regardless of the type, coarse inclusions that become the starting point of brittle fracture and reduce toughness are reduced, and the control method of inclusions is completely different between the two. Therefore, the present invention does not control fine inclusions at all. However, according to the preferable production method of the present invention described later, fine inclusions having a circle-equivalent diameter of 2.0 μm or less are generally 100 to 1000. About 1 piece / mm ² comes to exist. Of the fine inclusions having a circle-equivalent diameter of 2.0 μm or less, if limited to Mg-containing oxides, there is almost no present invention.

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

The characteristic parts of the manufacturing method according to the present invention are the following (A) and (B).
(A) In the molten steel stage, the amount of free oxygen [O] before the addition of Al is 100 ppm or less, the holding time (t1) from the addition of Al to the start of casting is 15 minutes or more, and the cooling time at 1450 to 1500 ° C. during casting ( t2) is controlled to 300 seconds or less. By the method (A), the average equivalent circular diameter of the coarse inclusions described above is refined to 3.5 μm or less.
(B) After hot rolling, after heating and holding in a temperature range of A _{c1 to} A _c3 points, tempering is performed in a temperature range of 520 ° C. to A _c1 points for 10 to 60 minutes. In particular, the volume fraction of the residual γ phase existing at −196 ° C. is appropriately controlled by the method (B).

  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, based on the viewpoint that the Al inclusions are coarsened by aggregation and coalescence, and it is easy to form coarse inclusions that are the starting points of brittle fracture, the generation of such coarse Al inclusions is prevented. Special attention is paid to the method of adding Al.

  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. If the amount of [O] exceeds 100 ppm, the size of inclusions produced when Al is added increases, N1 cannot be controlled properly, and the desired cryogenic toughness cannot be achieved (No in Table 2 described later). .33). The smaller the amount of [O], the better, preferably 80 ppm or less, and more preferably 50 ppm or less. In addition, the lower limit of the [O] amount is not particularly limited from the viewpoint of refining coarse 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, after adding Al to the molten steel, the holding time (t1) from the addition of Al to the start of casting is set to 15 minutes or more. As a result, coarse inclusions float and are separated. Conventionally, casting was started within 13 minutes at the same time as Al addition or after Al addition, but when t1 is less than 15 minutes, the effect of removing coarse inclusions is not effectively exhibited, It was found that the desired cryogenic toughness was not exhibited because the coarse inclusions were not refined (see No. 34 and No. 55 in Table 2 below). From the above viewpoint, t1 is preferably as long as possible, preferably 18 minutes or more, and more preferably 20 minutes or more. Although the upper limit of t1 is not specifically limited from the said viewpoint, Since holding for a long time causes the increase in manufacturing cost, it is preferable that it is 180 minutes or less, More preferably, it is 150 minutes or less.

  Next, casting starts. The temperature range at the time of casting is generally 1650 ° C. or less. However, in the present invention, it is particularly important to control the cooling time (t2) in the temperature range of 1450 to 1500 ° C. to 300 seconds or less. It was found that coarse inclusions were appropriately refined. When t2 exceeds 300 seconds, secondary inclusions are complexly formed with inclusions as nuclei, the size of coarse inclusions is increased, and the desired cryogenic toughness is not exhibited (table to be described later). 2 No. 35 and No. 56). From the above viewpoint, t2 is preferably as short as possible, preferably 290 seconds or less, and more preferably 280 seconds or less. The lower limit of t2 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 at the center of the slab thickness. The slab thickness is generally 150 to 250 mm, and the surface temperature tends to be about 200 to 1000 ° C. lower than the temperature at the center. Since the surface temperature has a large variation in temperature difference, the surface temperature is the temperature at the central portion (near the thickness t × 1/2) where the variation is small. The temperature at the center of the slab thickness can be measured by inserting a thermocouple into the mold.

  Further, in the present invention, it is only necessary to control the cooling time (t1) in the temperature range of 1450 to 1500 ° C. to 300 seconds or less, and the means is not limited. For example, the temperature range may be cooled at 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 You may cool with a different cooling rate so that the cooling time of a range may be 300 seconds or less.

  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. 36 and 37 in Table 2 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. The upper limit of the preferred heating time is 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. The tempering treatment is performed for 10 to 60 minutes (t3) in a temperature range of 520 ° C. to A _c1 point (T3). 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 during 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. 40 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 at −196 ° C. The amount of γ cannot be secured [No. 41 (example where T3 is high), No. 54 (example in which 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 ensured (see No. 42 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 Table 2, [O] is the amount of dissolved oxygen (ppm) before addition of Al, t1 is the time (minutes) from the addition of Al to the start of casting, t2 is the cooling time (seconds) of 1500 to 1450 ° C. during casting. ). Cooling at 1500 to 1450 ° C. was controlled by air cooling or water cooling so that the cooling time was reached.

  Next, after heating the above ingot to 1100 ° C., the steel sheet is rolled to a thickness of 75 mm at a temperature of 830 ° C. or higher, rolled at a final rolling temperature of 780 ° C., and then water-cooled. Got. 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.

  For the thick steel plate thus obtained, the amount of residual γ phase (volume fraction) present at an average equivalent circle diameter N1 of inclusions with an equivalent circle diameter of more than 2.0 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 the average equivalent circle diameter N1 of inclusions having an equivalent circle diameter of more than 2.0 μm The t / 4 position (t: plate thickness) of the above steel sheet is mirror-polished and photographed at a magnification of 400 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 subjected to image analysis using “Image-Pro Plus” manufactured by Media Cybernetics, and the equivalent circle diameter (diameter) of inclusions with an equivalent circle diameter (diameter) exceeding 2.0 μm was calculated. The average value was calculated.

(2) Measurement of amount of residual γ phase (volume fraction) present at −196 ° C. From a t / 4 position of each steel plate, a 10 mm × 10 mm × 55 mm test piece was sampled and liquid nitrogen temperature (−196 ° C.) Was held for 5 minutes, and then 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, No. 2 in Table 2A. 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 42 and 54 to 56 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, this is an example in which the coarse inclusions were not refined because the amount of dissolved oxygen [O] before addition of Al was large. As a result, the brittle fracture surface ratio also increased, and the desired cryogenic toughness could not be realized at -196 ° C.

  No. No. 34 is No. in Table 1B with a large amount of C. 34, and the time (t1) from the addition of Al to the start of casting is short; No. 55 of Table 1B in which the components in the steel satisfy the requirements of the present invention. 55 is used, but t1 is a short example. In any case, since t1 was short, coarse inclusions were not refined. As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized.

  No. No. 35 is No. in Table 1B with a large amount of P. No. 35 and a cooling time (t2) of 1500 to 1450 ° C. at the time of casting is a long example; No. 56, No. in Table 1B in which the components in the steel satisfy the requirements of the present invention. 56 is used, but the above t2 is an example. In either case, since t2 was long, coarse inclusions were not refined. As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized.

  No. No. 36 in Table 1B in which the components in the steel satisfy the requirements of the present invention. Although 36 was used, the amount of residual γ was insufficient because it was heated at a temperature lower than the two-phase region temperature (TL). As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized.

  No. No. 37 is No. in Table 1B with a large amount of Si. This is an example in which the amount of residual γ is insufficient because it is heated at a temperature exceeding the two-phase region temperature (TL). As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized.

  No. No. 38 in Table 1B where the components in the steel satisfy the requirements of the present invention. In this example, the amount of residual γ is insufficient because the heating and holding time (tL) at the two-phase region temperature (TL) is short. As a result, the brittle fracture surface ratio also 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 increased because the heat holding time (tL) at the two-phase region temperature (TL) was long. As a result, the yield strength YS decreased, and the desired base material strength could not be ensured.

  No. No. 40 in Table 1B in which the components in the steel satisfy the requirements of the present invention. Although 40 was used, the amount of residual γ was insufficient because the tempering temperature (T3) was low. As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized.

  No. No. 41 in Table 1B with a large amount of Mn. This is an example in which the amount of residual γ is insufficient because No. 41 is used and the tempering temperature (T3) is high. As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized.

  No. No. 42 in Table 1B in which the components in the steel satisfy the requirements of the present invention. Although 42 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. 54 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 time (t3) was short. As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized.

  No. Nos. 43 to 53 are comparative examples produced by the method of the present invention using those from which only the components in steel are removed.

  Specifically, no. No. 43 of Table 1B with a small amount of Mn. This is an example in which the amount of residual γ is insufficient because 43 is used. As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized.

  No. No. 44 of Table 1B with a large amount of S. This is an example using 44. For this reason, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 45 of Table 1B has a small amount of C, a large amount of Al, and a small amount of Ni. In this example, coarse inclusions were not refined and the amount of residual γ was insufficient. As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized. TS also decreased.

  No. No. 46 in Table 1B has a small amount of Al and a large amount of N. In this example, coarse inclusions were not refined because 46 was used. As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized.

  No. No. 47 of Table 1B with a large amount of Cu and Ca as the selection components. In this example, coarse inclusions were not refined because 47 was used. As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized.

  No. No. 48 in Table 1B, which has a large amount of Cr and Zr as selective components. This is an example where coarse inclusions were not refined because 48 was used. As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized.

  No. No. 49 in Table 1B, which has a large amount of Nb and REM as selective components. In this example, coarse inclusions were not refined because 49 was used. As a result, the brittle fracture surface ratio also increased and the desired cryogenic toughness could not be realized.

  No. No. 50 of Table 1B with a large amount of Mo as a selection component. Since 50 was used, the brittle fracture surface ratio increased and the desired cryogenic toughness could not be realized.

  No. No. 51 in Table 1B with a large amount of Ti as a selected 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 V as a selected 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 B as a selected component. Since 53 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, 13, 15, 19, and 23 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 Of 4.0% or more. 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 (8)

% 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 residual austenite phase present at −196 ° C. is 2.0% to 12.0% in volume fraction, and
Wherein an average circle equivalent diameter of a circle equivalent diameter 2.0μm greater inclusions is 3.5μm or less, excellent high cryogenically toughness at -196 ° C. in the plate width direction of the steel sheet (C-direction) Strength thick steel plate. The high-strength thick steel plate according to claim 1, wherein a residual austenite phase existing at -196 ° C is 4.0% to 12.0% in volume fraction. Furthermore,
The high-strength thick steel plate according to claim 1 or 2, containing Cu: 1.0% or less (not including 0%). Furthermore,
Cr: 1.20% or less (not including 0%), and Mo: 1.0% or less (not including 0%)
The high-strength 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: 0.100% or less (not including 0%), and V: 0.50% or less (not including 0%)
The high-strength 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 high-strength thick steel plate according to any one of claims 1 to 5, containing 0.0050% or less (excluding 0%). Furthermore,
Ca: 0.0030% or less (excluding 0%),
REM: 0.0050% or less (not including 0%), and Zr: 0.005% or less (not including 0%)
The high-strength thick steel plate according to claim 1, comprising at least one selected from the group consisting of: A method for producing the high-strength thick steel plate according to any one of claims 1 to 7,
(A) In the molten steel stage, the amount of free oxygen [O] before the addition of Al is 100 ppm or less, the holding time (t1) from the addition of Al to the start of casting is 15 minutes or more, and the cooling time at 1450 to 1500 ° C. during casting ( controlling t2) to 300 seconds or less;
(B) After hot rolling, heating and holding for 10 to 50 minutes in the temperature range of A _{c1 to} A _c3 points, and then tempering for 10 to 60 minutes in the temperature range of 520 ° C. to A _c1 points;
The manufacturing method of the high strength thick steel plate characterized by including.