JP2005139517A - Method for producing high strength and high toughness thick steel plate - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a method for producing a high strength and high toughness thick steel plate which has the high strength of ≥590 MPa tensile strength and excellent balance to the strength and the toughness and uniform characteristic in the thickness direction of the plate. <P>SOLUTION: To a steel blank composed by mass% of 0.01-0.20% C, 0.01-0.60% Si, 0.50-2.50% Mn, ≤0.020% P, ≤0.0070% S and 0.001-0.100% solAl, a hot-rolling is applied at ≥Ar<SB>3</SB>point of rolling finish temperature after heating, and successively, the cooling is started at CR<SB>M</SB>martensite generating critical cooling speed from the temperature zone of ≥Ar<SB>3</SB>point, and after cooling to Ms point to 300°C quenching-cooling stop temperature, this steel is isothermally held to this temperature for a prescribed time, or cooled at a rate of <10°C/s from this temperature to a room temperature. Further, this steel blank can contain one or more groups among one or more kinds of Cu,Ni,Cr,Mo,B, one or more kinds of Ti,V,Nb and one or more kinds of Ca,REM. <P>COPYRIGHT: (C)2005,JPO&NCIPI

Description

  The present invention relates to a method for producing a high-strength, high-tough steel plate suitable for use in welded steel structures such as shipbuilding, offshore structures, building machines, buildings, bridges, tanks, pipelines, and penstocks. -Regarding improvement of toughness balance and improvement of homogeneity of sheet thickness direction characteristics. In the present invention, “high strength” refers to strength having a tensile strength of 590 MPa or more. The “thick steel plate” in the present invention refers to a steel plate having a thickness of 2 mm or more.

  Generally, as the strength of the steel sheet increases, the low temperature toughness tends to decrease. In particular, it is not easy to provide good low temperature toughness in a high strength steel sheet having a tensile strength of 590 MPa or more.

  Conventionally, high-strength thick steel plates having a tensile strength of 590 MPa or more have been manufactured by a reheating quenching process in which the steel is reheated to an austenite temperature range and then quenched. Recently, for example, Patent Document 1 proposes a method for producing a high-strength thick steel plate by direct quenching that is quenched immediately after hot rolling. However, in order to obtain good toughness, both in the case of reheating quenching treatment and in the case of direct quenching treatment, tempering treatment is performed to reduce the hardness of martensite generated during quenching and restore toughness. It was necessary and left problems in terms of productivity, manufacturing cost, construction period, and the like.

As a method for solving these problems, for example, in Patent Document 2, a steel material in which the amounts of C, Si, Mn, Cr, Ti, B, Al, and N are adjusted to appropriate amounts is set to a temperature range of 900 ° C. or lower. After performing hot rolling with a cumulative reduction ratio of 50% or more at, immediately start quenching at a temperature of Ar ₃ or higher, and quench in the temperature range from (Ms point + 100 ° C) to (Ms point -250 ° C). A method for producing a wear-resistant steel material that is stopped and has a martensite structure at the surface layer and excellent in toughness and delayed fracture resistance is described. The technique described in Patent Document 2 stops quenching at a temperature in the temperature range sandwiching the Ms point, and the surface layer part is a martensite single phase structure and the inner layer part is a mixed structure of lower bainite and martensite. is there. According to the technique described in Patent Document 2, a thick steel plate having high strength and high toughness can be manufactured without performing a tempering treatment. However, this technique has a problem in that the hardness distribution in the thickness direction varies and the homogeneity of the material in the thickness direction is inferior, and the toughness of the surface layer portion that becomes a martensite single phase structure decreases.

Further, for example, in Patent Document 3, hot rolling with a cumulative rolling reduction of 50% or more is performed in a temperature range of 1000 to 900 ° C., and then in an unrecrystallized γ region of less than 900 ° C. and 810 ° C. or more per pass. Production of ultra-high-strength steel sheets for welding with excellent homogeneity in the thickness direction, with a cumulative reduction of 10 to 30% by light reduction with a reduction of less than 10%, followed by quenching and subsequent tempering A method has been proposed. However, in the technique described in Patent Document 3, the homogeneity in the plate thickness direction is improved, but tempering after quenching is essential, and the process is complicated and prolonged, which is problematic from the viewpoint of productivity and manufacturing cost. Was leaving.
JP-A-2-27407 JP 2002-80930 A Japanese Patent Publication No. 6-70248

  The present invention solves the above-described problems of the prior art, and has a high strength and high toughness steel plate having a high strength of tensile strength of 590 MPa or more, an excellent strength-toughness balance, and a uniform thickness direction property. It is an object of the present invention to propose a method for producing a high-strength, high-toughness thick steel plate that can be produced at high efficiency and at low cost without performing tempering treatment.

  In order to achieve the above-mentioned problems, the present inventors diligently studied various factors affecting low-temperature toughness at a tensile strength level of 590 MPa or more. As a result, by quenching and cooling at a specific cooling rate from the austenite region temperature, quenching and cooling is stopped in a temperature region below the Ms point, and held for a specific time in the temperature region, depending on the plate thickness position of the surface layer portion, center portion, etc. It was found that a mixed structure of tempered martensite and lower bainite was obtained, the strength-toughness balance was improved, and the thickness direction characteristics were homogenized.

  In low carbon (C: 0.01 to 0.20 mass%) steel, austenite is transformed during quenching and cooling to form a lath structure. At this time, C is concentrated between the laths, and untransformed austenite (γ) film tends to remain, forming acicular island martensite (hereinafter also referred to as MA) that is harmful to toughness. Cheap.

  The present inventors stopped quenching cooling at a temperature below the Ms point and kept isothermal at that temperature, in addition to tempering the martensite generated during quenching cooling, and forming a fine and large amount of lower bainite. In addition, the MA between the laths of the lower bainite was found to be reduced and the toughness was improved. In addition, by performing the isothermal holding for an appropriate time, the present inventors proceeded with the decomposition of γ between the laths of the lower bainite, reducing the amount of untransformed γ, reducing the amount of MA produced, and reducing the deterioration of toughness. I found out.

  On the other hand, when quenching and cooling is stopped at a temperature slightly higher than the Ms point (immediately above the Ms point) and kept isothermally at that temperature, lower bainite transformation occurs, and needle-like MA or lump that extends long between laths of the lower bainite. MA is easy to form, and toughness decreases.

  In this way, the inventors stopped quenching cooling at a temperature below the Ms point, and kept isothermal at that temperature, thereby making it possible to reduce the generation of island-like martensite phases, and further finer It has been found that the lower bainite phase can be obtained at a high fraction, and that it is possible to produce a high strength and high tough steel plate.

  The present invention has been completed based on such findings and further studies. That is, the gist of the present invention is as follows.

(1) By mass%, C: 0.01 to 0.20%, Si: 0.01 to 0.60%, Mn: 0.50 to 2.50%, P: 0.020% or less, S: 0.0070% or less, solAl: 0.001 to 0.100%, After heating the steel material having the composition of the balance consisting of Fe and inevitable impurities, the steel material is subjected to hot rolling at a temperature in the temperature range higher than the Ar ₃ transformation point, and then higher than the Ar ₃ transformation point. From the temperature range, the following equation (1)
logCR _M = 2.94−0.75β (1)
Here, β = 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + 2Mo (when B ≧ 0.0005 mass%)
= 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + Mo (when B <0.0005 mass%)
CR _M: martensite critical cooling rate (° C. / s)
C, Si, Mn, Ni, Cr, Mo: Content of each element (mass%)
In in defined to start cooling with martensite critical cooling rate CR _M or more cooling rate is, after cooling to quench cooling stop temperature of martensitic transformation start temperature or lower 300 ° C. or higher temperature range,該焼purse cooling stop temperature A method for producing a high-strength, high-tough steel plate, characterized in that it is kept isothermal for a predetermined time, or the temperature range from the quenching cooling stop temperature to room temperature is cooled at a cooling rate of less than 10 ° C / s.

  (2) The method for producing a high strength and high toughness thick steel plate according to (1), wherein the predetermined time is 1 to 60 minutes.

(3) In (1) or (2), in addition to the above-mentioned composition, in mass%, the following groups A to C group A: mass%, Cu: 0.10 to 1.00%, Ni: 0.10 to 5.00%, Cr : 0.10 to 0.80%, Mo: 0.01 to 0.80%, B: One or more selected from 0.0002 to 0.0025% Group B: mass%, Ti: 0.03% or less, V: 0.005 to 0.100% , Nb: One or more selected from 0.050% or less Group C: Mass%, Ca: 0.010% or less, REM: From one or two selected from 0.020% or less A method for producing a high-strength, high-toughness steel plate, comprising one or more selected groups.

  According to the present invention, a high-strength, high-toughness thick steel plate having a high strength with a tensile strength of 590 MPa or more, an excellent balance between strength and toughness, and uniform thickness direction characteristics can be obtained without tempering. It can be manufactured efficiently and inexpensively, and has a remarkable industrial effect. Moreover, the high-strength thick steel plate obtained by this invention also has the effect that the reliability of a heat processing precision is high.

  The reason for limiting the composition of the steel material used in the present invention will be described first. Hereinafter, the mass% in the composition is simply expressed as%.

C: 0.01-0.20%
C is an element that increases the strength of steel, and in order to obtain a desired high strength, it needs to be contained in an amount of 0.01% or more. On the other hand, if the content exceeds 0.20%, the weldability deteriorates, the weld cracking is likely to occur, and the base metal toughness and the HAZ toughness are lowered. For this reason, C was limited to the range of 0.01 to 0.20%. In addition, Preferably it is 0.02 to 0.16%.

Si: 0.01-0.60%
Si is an element that acts as a deoxidizing material and further increases the strength of the steel material by solid solution strengthening. In order to obtain such an effect, a content of 0.01% or more is required, but a content exceeding 0.60% significantly deteriorates the HAZ toughness. For this reason, Si was made into the range of 0.01 to 0.60%. In addition, Preferably, it is 0.05 to 0.20%.

Mn: 0.50-2.50%
Mn is an element that has the effect of enhancing the hardenability of steel and improving toughness. In the present invention, it is necessary to contain 0.50% or more, but the content exceeding 2.50% may deteriorate weldability. There is. For this reason, in this invention, Mn was limited to 0.50 to 2.50% of range. In addition, Preferably, it is 0.80 to 2.50%.

P: 0.020% or less P is an element that increases strength by solid solution strengthening. However, in order to deteriorate toughness and weldability, P is preferably reduced as much as possible in the present invention, but the content up to 0.020% is acceptable. For this reason, P was limited to 0.020% or less. In addition, Preferably it is 0.017% or less. In addition, since extreme reduction leads to an increase in melting cost, it is preferably 0.008% or more in the present invention.

S: 0.0070% or less S is present as a sulfide in steel and exhibits an effect of reducing ductility. For this reason, it is desirable to reduce S as much as possible, but it is acceptable up to 0.0070%. In addition, Preferably it is 0.0030% or less. Moreover, since an extreme reduction leads to an increase in melting cost, it is preferable to make it 0.0005% or more in the present invention.

solAl: 0.001 to 0.100%
Al acts as a deoxidizer during steelmaking, and in the present invention, it needs to be contained in an amount of 0.001% or more. However, if it exceeds 0.100%, the toughness is lowered. For this reason, solAl was limited to the range of 0.001 to 0.100%. In addition, Preferably, it is 0.001 to 0.060%.

In addition to the basic composition described above, in the present invention, the following groups A to C, group A: mass%, Cu: 0.10 to 1.00%, Ni: 0.10 to 5.00%, Cr: 0.10 to 0.80%, Mo: 0.01 to One or more selected from 0.80%, B: 0.0002-0.0025% Group B: mass%, Ti: 0.03% or less, V: 0.005-0.100%, Nb: selected from 0.050% or less 1 type or 2 types or more selected: C group:% by mass, Ca: 0.010% or less, REM: 0.020% or less selected from 1 type or 2 types selected from 1 type or 2 types Can be contained.

  Cu, Ni, Cr, Mo, and B in group A are all elements that further improve the hardenability of the steel, and can be selected as necessary to contain one or more.

  Cu has the effect of further improving the hardenability. If the content is less than 0.10%, such an effect cannot be expected. On the other hand, if the content exceeds 1.00%, the risk of causing hot brittleness increases. For this reason, it is preferable to limit Cu to 0.10 to 1.00%. In addition, More preferably, it is 0.10 to 0.30%.

  Ni has the effect of improving the hardenability of the steel and improving the toughness. Such an effect is recognized at a content of 0.10% or more, but a content exceeding 5.00% tends to increase the manufacturing cost. For this reason, it is preferable to limit Ni to the range of 0.10 to 5.00%. In addition, More preferably, it is 0.20 to 2.00%.

  Cr is an inexpensive element that improves the hardenability of steel, and such an effect is recognized with a content of 0.10% or more. However, a content exceeding 0.80% deteriorates weldability and toughness. For this reason, it is preferable to limit Cr to the range of 0.10 to 0.80%. In addition, More preferably, it is 0.20 to 0.80%.

  Mo is an element that has the effect of further improving the hardenability of steel, and such an effect is recognized with a content of 0.01% or more. However, a content exceeding 0.80% deteriorates weldability and toughness. For this reason, it is preferable to limit Mo to the range of 0.01 to 0.80%. In addition, More preferably, it is 0.10 to 0.60%.

  B is an element that improves the hardenability of the steel in a small amount, and such an effect is recognized with a content of 0.0002% or more. However, when it exceeds 0.0025%, the hardenability is lowered. For this reason, it is preferable to limit B to 0.0002 to 0.0025% of range. In addition, More preferably, it is 0.0005 to 0.0020%.

  In the group B, Ti, V, and Nb are all elements that influence the increase in strength through the formation of carbides and / or nitrides, and can be selected as needed to contain one or more.

  Ti combines with N in the steel to form TiN, suppresses the coarsening of crystal grains and contributes to the improvement of HAZ toughness, and reduces the solid solution N and ensures the effect of improving the hardenability of B. Have. Such an effect is recognized when the content is 0.005% or more. However, when the content exceeds 0.03%, TiN becomes coarse, the effect of refining γ grains disappears, and toughness deteriorates. For this reason, it is preferable to limit Ti to 0.03% or less.

  V precipitates as carbides and / or nitrides and has the effect of increasing the strength of the steel by precipitation hardening. Such an effect is recognized at a content of 0.005% or more, but if it exceeds 0.100%, the weldability deteriorates. For this reason, it is preferable to limit V to 0.005 to 0.100% of range. In addition, More preferably, it is 0.010 to 0.060%.

  Nb has the effect of suppressing recrystallization of austenite grains during hot rolling, facilitating the expansion of austenite grains by hot rolling, and refining ferrite to improve strength and toughness. Such an effect becomes remarkable when the content is 0.003% or more, but the content exceeding 0.050% deteriorates the weldability and the HAZ part toughness. For this reason, Nb is preferably limited to a range of 0.050% or less. In addition, More preferably, it is 0.010 to 0.030%.

  Ca and REM of group C are both sulfide-forming elements, elements that spheroidize sulfides and improve ductility, and can be selected as necessary to contain one or two kinds. Such effects become remarkable when the content of Ca is 0.001% or more and REM: 0.001% or more. However, when the content exceeds Ca: 0.010% and REM: 0.020%, sulfide is excessively formed and the toughness deteriorates. For this reason, it is preferable to limit to Ca: 0.010% or less and REM: 0.020% or less.

  The balance other than the components described above consists of Fe and inevitable impurities.

  The molten steel having the above composition is melted by a normal melting means such as a converter or an electric furnace, and is made into a steel material such as a slab by a normal casting method such as a continuous casting method or an ingot-bundling method. Is preferred. The melting method and the casting method are not limited to the methods described above.

  The steel material is heated to a temperature at which it becomes an austenite single phase.

  The heating temperature of the steel material is a temperature at which the steel material becomes an austenite single phase structure, preferably 1050 to 1250 ° C. If the heating temperature of the steel material is less than 1050 ° C., the hot deformation resistance is too high, so that the rolling reduction per time cannot be taken high, and the productivity is lowered. Further, when a precipitate-forming element such as V or Nb is contained, these elements are not sufficiently dissolved in austenite, and it is difficult to sufficiently exhibit the effects of these elements. On the other hand, when the heating temperature exceeds 1250 ° C., the crystal grains become coarse and the amount of scale loss and the frequency of furnace repairs increase. For this reason, it is preferable to make the heating temperature of a steel raw material into the range of 1050-1250 degreeC.

The heated steel material is subjected to hot rolling with the rolling end temperature set to a temperature in the temperature range equal to or higher than the Ar ₃ transformation point to obtain a thick steel plate. If the rolling end temperature is lower than the Ar ₃ transformation point, ferrite precipitates during rolling, and a desired structure cannot be obtained even if quenching is performed thereafter, and a desired strength cannot be ensured. The cumulative rolling reduction in the temperature range above the Ar ₃ transformation point is preferably 30% or more. If the cumulative rolling reduction is less than 30%, sufficient austenite grain refinement cannot be achieved.

After the hot rolling is finished, the thick steel plate is quenched and cooled from a temperature range not lower than the Ar ₃ transformation point. When the quenching cooling start temperature is lower than the Ar ₃ transformation point, the structure at the quenching cooling start is not an austenite single phase, and the transformation to a part of ferrite has started. The amount of sites is small and the desired strength cannot be ensured.

The cooling rate of the quenching cooling, the martensite critical cooling rate CR _M or more cooling rate. Note that martensite critical cooling rate CR _M in the present invention, the following (1)
logCR _M = 2.94−0.75β (1)
Here, β = 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + 2Mo (when B ≧ 0.0005 mass%)
= 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + Mo (when B <0.0005 mass%)
CR _M: martensite critical cooling rate (° C. / s)
C, Si, Mn, Ni, Cr, Mo: Content of each element (mass%)
Refers to the cooling rate defined by.

In the present invention, in martensite critical cooling rate CR _M or more cooling rate, cooling to quench cooling stop temperature of martensitic transformation starting temperature (Ms point) or less 300 ° C. or higher temperature range is subjected to quenching treatment. Thereby, martensite is first partially generated at each position in the plate thickness direction.

  The partial generation of martensite here aims to generate strain using expansion during martensite transformation at the interface between the generated martensite and untransformed austenite. This strain energy facilitates transformation of untransformed austenite to lower bainite, and it is possible to produce a lower bainite phase that is finer and more abundant than conventional ones.

The cooling rate is martensite subcritical cooling rate CR _M quench cooling, martensite coarse bainite is generated before transformation, generation of distortion due to martensitic transformation described above is insufficient, can be expected desired to effect Disappear.

  Further, when the quenching cooling stop temperature exceeds the Ms point, the strain generation effect due to the formation of martensite cannot be expected, the transformation to the lower bainite phase is insufficiently promoted, and the subsequent isothermal holding or cooling ( (Slow cooling) increases the amount of island martensite harmful to toughness. On the other hand, if the quenching and cooling stop temperature is less than 300 ° C., the diffusion of C becomes insufficient, and carbide effective for crack propagation resistance does not precipitate in bainitic ferrite. For this reason, the quenching and cooling stop temperature was set to a temperature in the temperature range of 300 ° C. or higher below the Ms point. In addition, Preferably, it is the temperature range below 350 degreeC above Ms point.

  Next, the thick steel plate is kept isothermal for a predetermined time at the quenching cooling stop temperature in the above-described range, or is cooled at a cooling rate of less than 10 ° C./s in the temperature range from the quenching cooling stop temperature to room temperature.

  By maintaining isothermal for a predetermined time at the quenching and cooling stop temperature, martensite is self-annealed, while transformation of untransformed austenite to lower bainite is promoted, and a mixed structure of tempered martensite and lower bainite can be obtained. . The predetermined holding time in the isothermal holding is preferably 1 to 60 minutes. If the isothermal holding time is less than 1 min, the transformation of untransformed austenite to lower bainite is not completed, and the untransformed austenite enriched with C becomes island-like martensite by subsequent cooling, so that the toughness deteriorates. On the other hand, even if the isothermal holding time is longer than 60 min, no significant change is observed in the internal tissue. For this reason, the isothermal holding time is preferably 1 to 60 minutes.

  In the present invention, instead of isothermal holding, the temperature range from the quenching cooling stop temperature to room temperature may be cooled (slowly cooled) at a cooling rate of less than 10 ° C./s. By slow cooling, untransformed austenite is easily transformed into lower bainite. When the cooling rate from the quenching cooling stop temperature to room temperature is 10 ° C./s or more, martensite is further generated during cooling and the toughness is lowered. For this reason, it is preferable to cool the temperature range from the quenching cooling stop temperature to room temperature at a cooling rate of less than 10 ° C./s.

  The thick steel plate obtained under the manufacturing conditions described above has the above-described composition and has a mixed structure of tempered martensite and lower bainite regardless of the position in the plate thickness direction. The structural fraction of tempered martensite and lower bainite is 80% or more by volume. In addition, it is preferable that the structure fraction of the lower bainite is 50% or more of the whole by volume. In addition, as phases other than tempered martensite and lower bainite, mixing of upper bainite and ferrite having a volume ratio of 20% or less is acceptable. Here, “tempered martensite” refers to martensite in which carbides are precipitated or spheroidized. The “lower bainite” here also includes tempered lower bainite in which carbides are spheroidized.

  In actual operation, the temperature of the steel sheet is generally controlled by the surface temperature of the steel sheet, and the average temperature of the entire steel sheet is calculated in real time, and temperature control and speed control are generally performed based on this average temperature. Therefore, “temperature” in the present invention means the average temperature of the entire steel plate, and “cooling rate” means the average cooling rate of the whole steel plate.

In the present invention, the transformation points of Ar ₃ and Ms are expressed by the following formulas (2) and (3) based on the content of each element in each steel material (thick steel plate).
Ar ₃ = 910-273C-74Mn-56Ni-16Cr-9Mo-5Cu (2)
Ms = 517-300C-11Si-33Mn-22Cr-17Ni-11Mo (3)
(Here, C, Mn, Si, Cr, Ni, Mo, Cu: content of each element (%))
The value obtained by calculating using is used.

  A steel ingot obtained by melting molten steel having the composition shown in Table 1 in a vacuum melting furnace was hot rolled to form a slab (steel material) having a thickness of 100 mm. Subsequently, these steel slabs were subjected to hot rolling under the conditions shown in Table 2, followed by quenching to obtain thick steel plates (plate thickness: 12 to 50 mm).

  With respect to the obtained thick steel plate, a No. 4 tensile test piece was collected from the position in the thickness direction 1/2 in accordance with the provisions of JIS Z 2201, and a tensile test was performed in accordance with the provisions of JIS Z 2241. The tensile strength TS was determined. The surface layer portion means a position 1 mm from the surface.

In addition, a Charpy impact test piece with a V-notch standard dimension was collected from the position of the obtained thick steel plate in the thickness direction 1/2 in accordance with JIS Z 2202, and in accordance with JIS Z 2242. An impact test was performed to determine Charpy impact absorption energy vE- ₄₀ (J) at -40 ° C.

  In addition, from the obtained thick steel plate, a hardness measurement test piece is collected, and the hardness is continuously measured using a Vickers hardness tester in the plate thickness direction from the front surface to the back surface in a cross section (C cross section) orthogonal to the rolling direction. did. The difference between the maximum value and the minimum value, ΔHv, was obtained from the hardness distribution in the plate thickness direction, and the homogeneity in the plate thickness direction of the thick steel plate was evaluated.

  In addition, a specimen for tissue observation is collected from the obtained thick steel plate, and the structure is observed at a position in the plate thickness direction 1/2 with a scanning electron microscope and a transmission electron microscope. The tissue fraction was determined. In addition, tempered martensite and lower bainite were discriminated by the precipitation form of carbides. The structure fraction of each structure is determined by measuring the average austenite (γ) particle size by a line segment method using a scanning electron microscope, and randomly selecting 10 particles having an average γ particle size. The area of each structure was determined as the cross-sectional area ratio, and the average value of the 10 cross-sectional area ratios was taken as the structure fraction at each position of the steel sheet. In addition, the steel plate temperature and the cooling rate in Table 2 are displayed using the average temperature and the average cooling rate.

  The obtained results are shown in Table 3.

  Each of the examples of the present invention is a thick steel plate having a high strength with a tensile strength of 590 MPa or more and a high toughness according to each strength, and also has a small ΔHv in the thickness direction and excellent uniformity in the thickness direction characteristics. ing. On the other hand, all the comparative examples outside the scope of the present invention have a lower amount of lower bainite, and any one of strength, toughness, and homogeneity is lowered. In particular, it can be seen that the increase in the amount of lower bainite greatly contributes to the improvement of toughness.

  For example, steel plates No. 29, No. 30, and No. 31 with excellent quenching cooling rate outside the scope of the present invention have excellent toughness, but a desired strength level that matches the alloy element content is not obtained. Further, the toughness is deteriorated in the steel plates No. 32 and No. 34 whose quenching and cooling stop temperature is higher than the Ms point and out of the scope of the present invention. The same applies to steel plate No. 33, where the quenching and cooling stop temperature is set to room temperature (50 ° C.) and is outside the scope of the present invention.

  In addition, the steel plate No. 35 and No. 36, which have a cooling rate of 10 ° C./s or more by cooling from the quenching cooling stop temperature and are outside the scope of the present invention, have toughness compared to the present invention example of the same strength level. Deteriorated greatly.

  In addition, the steel plates Nos. 37 to 41 whose chemical components deviate from the scope of the present invention are greatly deteriorated in toughness as compared with the present invention example having the same strength level.

Claims (3)

% By mass
C: 0.01-0.20%, Si: 0.01-0.60%,
Mn: 0.50-2.50%, P: 0.020% or less,
S: 0.0070% or less, solAl: 0.001 to 0.100%
After heating a steel material having a composition consisting of Fe and the inevitable impurities in the balance, the steel is subjected to hot rolling in which the rolling end temperature is equal to or higher than the Ar ₃ transformation point, and then the Ar ₃ transformation point. from the above temperature range, the following (1) martensite critical cooling rate CR _M or more cooling rate defined by start cooling by the formula, the martensitic transformation starting temperature below 300 ° C. over a temperature range of quench cooling stop temperature After cooling to the quenching cooling stop temperature, the isothermal holding is performed for a predetermined time, or the temperature range from the quenching cooling stop temperature to room temperature is cooled at a cooling rate of less than 10 ° C / s. A method for producing a tough steel plate.
Record
logCR _M = 2.94−0.75β (1)
Here, β = 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + 2Mo (when B ≧ 0.0005 mass%)
= 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + Mo (when B <0.0005 mass%)
CR _M: martensite critical cooling rate (° C. / s)
C, Si, Mn, Ni, Cr, Mo: Content of each element (mass%)   The method for producing a high-strength, high-toughness thick steel plate according to claim 1, wherein the predetermined time is 1 to 60 minutes. The high-strength and high-toughness thick steel plate according to claim 1 or 2, further comprising one group or two or more groups selected from the following groups A to C in mass% in addition to the composition. Manufacturing method.
Group A: Mass%, Cu: 0.10 to 1.00%, Ni: 0.10 to 5.00%, Cr: 0.10 to 0.80%, Mo: 0.01 to 0.80%, B: 0.0002 to 0.0025% Alternatively, two or more types B group: mass%, Ti: 0.03% or less, V: 0.005 to 0.100%, Nb: 0.050% or less selected from one or more types C group: mass%, Ca: One or two selected from 0.010% or less, REM: 0.020% or less