JP2009179840A - Method for producing high tensile steel excellent in low-temperature toughness and crack arrest property - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a method for producing a high tensile steel excellent in low-temperature toughness and crack arrest properties, the high tensile steel being excellent in crack arrest properties as well as in strength and toughness. <P>SOLUTION: A steel blank containing specified contents of C, Si, Mn, P, S, Ni, Ti and N satisfying Pcm=C+Si/30+Mn/20+Cu/20+Ni/60+Cr/20+Mo/15+V/10+5B ≤0.25% and the balance Fe with inevitable impurities, is heated to 1,000-1,200°C, and the hot-rolling in ≥8% cumulative rolling-reduction ratio at Ac<SB>3</SB>point+20°C to Ar<SB>3</SB>point+100°C is applied to the heated steel blank, then an accelerated cooling starting from Ar<SB>3</SB>point-30°C to Ar<SB>3</SB>point+80°C and completing at ≤500°C, is applied to the hot-rolled steel at 5 to 100°C/sec cooling speed, and the thickness of prior γ grain on the surface layer, is controlled to ≤12 μm, and the crystal having ≤5 μm diameter on the surface layer after transforming, is present at an area ratio of 10% or more. <P>COPYRIGHT: (C)2009,JPO&INPIT

Description

  The present invention relates to a method for producing steel sheets for low-temperature storage tanks, low-temperature pressure vessels, and the like, and in particular, by controlling the rolling conditions, the thickness of old γ grains (former austenite grains) on the steel sheet surface is controlled, The present invention relates to a method for producing high-strength steel having excellent crack propagation stopping characteristics.

  In steel sheets used for low-temperature storage tanks, low-temperature pressure vessels, line pipes, etc., when tensile strength is improved, toughness tends to deteriorate, and it is not easy to satisfy both characteristics. However, considering the steel sheet for the LPG tank, the low-temperature characteristics of the steel sheet are particularly important because the storage of liquefied gas or the like is stored or transported in the cold state. Also, weldability is important.

Conventionally, as an example of a method for suppressing toughness deterioration while improving strength, C: 0.02% to 0.22% by weight%, Si: 0.05 to 0.6%, Mn: 0.5 to 2 .5%, S: 0.025% or less, Al: 0.01 to 0.08%, N: 0.008% or less, steel made of the remaining iron and unavoidable impurities by hot working After rolling for 30 mm or more, directly quenching treatment as a primary treatment, or after rolling, air cooling, reheating, normalizing treatment, or quenching treatment, and then as secondary treatment (Ac ₃ transformation point) Heating at a temperature in the range of −30) ° C. or more and (AC ₃ transformation point + 150) ° C. or less, and assuming that the plate thickness is tmm, the holding time T (min) is 0.1 t−0.05 t ≦ T In the range of ≦ 0.1t + 0.05t, extraction and reheating in the secondary treatment, A mixed structure of ferrite and bainite that have not been stenitized is generated at 10 to 60% per unit area in the range of 12.5% of the plate thickness on the front surface side and the back surface side, centering on 1/2 of the cross section of the steel plate. Then, the cooling rate of the temperature in the range of 800 to 500 ° C. is set to 2 to 10 ° C./second, forcibly cooled to normal temperature and quenched, and then heated to a temperature equal to or lower than the Ac ₁ transformation point as the third treatment. The technology to return is known. (See Patent Document 1)

Moreover, as an example of a technique for ensuring high toughness while obtaining high strength, C: 0.01% to 0.2%, Si: 0.01 to 1%, Mn: 0.1 to 3% by weight %, P: 0.02% or less, S: 0.01% or less, Al: 0.01 to 0.1%, Ni: 0.3 to 10%, Ti: 0.03 to 0.1%, N : A steel slab containing 0.002 to 0.1% and comprising the balance iron and inevitable impurities is heated to Ac3 transformation point to 1200 ° C, and the average austenite grain size is set to 20 to 100 µm, and the starting temperature is 900 ° C. Thereafter, hot rolling is performed at an end temperature of 650 ° C. or more and a cumulative reduction ratio of 30 to 95%, and subsequently, accelerated cooling is started at 600 ° C. or more and finished at 500 ° C. or less at a cooling rate of 1 to 100 ° C./s. Techniques for performing are known. (See Patent Document 2)
JP-A-1-195242 JP 2001-123322 A

As described above, a technology that is compatible with both high strength and high toughness is known. However, in recent technical development requirements, the property that crack propagation does not spread over a wide range when an impact is applied to a steel sheet. Attention has been paid.
However, in the techniques disclosed in Patent Documents 1 and 2 described above, attention has not been paid to the aspect of the crack propagation stop characteristic, and other conventional techniques have excellent strength and toughness, and cracks are also observed. There was no technology that paid attention to the propagation stop characteristics.

  The invention of the present application was made in view of the above-mentioned background. In addition to excellent strength and toughness, the method for producing high-tensile steel with excellent low temperature toughness and excellent crack propagation stopping characteristics having excellent characteristics in crack propagation stopping characteristics. The purpose is to provide.

As a result of intensive studies to solve the above problems, the present inventors have reported that the toughness reported to be improved together with the strength in these conventional techniques is the Charpy impact value, and strives to improve the crack propagation stop characteristics. As a result, the present invention has been achieved.
In this type of thick steel plate, the slab, which is a steel material, is heated to 1000 ° C. or higher, then subjected to a roughing mill and a hot straightening machine, and then subjected to quenching and tempering processes. .
The present inventor adjusts the composition range of steel in a thick steel plate manufactured in this manufacturing process, focusing on the composition system capable of balancing strength and toughness, and exists in the surface layer of the obtained steel plate. It was found that the crack propagation stopping property can be improved by controlling the diameter and distribution of the old γ grains.

The present invention has been completed based on the above findings, and the gist thereof is as follows.
(1) The present invention is mass%, C: 0.02% to 0.15%, Si: 0.01 to 1.0%, Mn: 0.3 to 3.0%, P: 0.015 % Or less, S: 0.01% or less, Ni: 0.1 to 10%, Ti: 0.003 to 0.1%, N: 0.001 to 0.01%, and Pcm = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5B is 0.25% or less, and the steel material consisting of iron and inevitable impurities is heated to 1000 to 1200 ° C., Ar ₃ points + 20 ° C. to Ar ₃ points + 100 ° C. After performing hot rolling with a cumulative rolling reduction of 8% or more, cooling is started from Ar ₃ point −30 ° C. to Ar ₃ point + 80 ° C., and is terminated at 500 ° C. or lower, and cooling rate is 5 to 100 ° C./sec. The thickness of the old γ grains on the surface layer is controlled to 12 μm or less. It allows low-temperature toughness, the method of producing a high strength steel excellent in crack propagation stop characteristics, characterized in that is present more than 10% crystalline area ratio of the surface layer after metamorphosis diameter 5μm or less to.

(2) In the present invention, the steel component is further in mass%, Cu: 0.05 to 1.5%, Cr: 0.05 to 0.2%, Mo: 0.05 to 0.5%, Nb : 0.005-0.05%, V: 0.005-0.05%, B: containing at least one of 30 ppm or less, Pcm = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5B The method for producing a high-strength steel excellent in low-temperature toughness and crack propagation stopping characteristics as described in (1), wherein a steel material satisfying 0.25% or less is used.
(3) The present invention further includes, as a steel component, in mass%, Y: 0.001 to 0.1%, Ca: 0.0001 to 0.01%, Mg: 0.0001 to 0.01%, REM. : Production of high-strength steel excellent in low temperature toughness and crack propagation stopping characteristics as described in (1) or (2), wherein a steel material containing one or more of 0.005 to 0.1% is used. Method.

(4) The present invention is characterized by low temperature toughness and crack propagation stopping characteristics according to any one of (1) to (3), wherein the tempering is performed at a temperature of 500 ° C. or higher and lower than the Ac1 transformation point after the accelerated cooling. An excellent high-strength steel manufacturing method.

The present invention contains a specific amount of C, Si, Mn, P, S, Ni, Ti, N, and in the structure, by controlling the thickness of the former γ grains of the surface layer to 12 μm or less, the diameter of the surface layer after transformation is Since crystals of 5 μm or less were present in an area ratio of 10% or more, it was possible to ensure good low temperature toughness while maintaining a high strength, and to prevent solidification of alloy elements in the suppression of γ grain size coarsening and solid solution strengthening. It is possible to provide a high-strength steel capable of improving both strength and drop weight characteristics by heating under the conditions satisfying the above conditions.
Moreover, the old γ grain thickness in the steel sheet surface layer can be controlled to 12 μm or less by carrying out hot rolling in which the cumulative rolling reduction at Ar ₃ point + 20 ° C. to Ar ₃ point + 100 ° C. is 20% or more.
Then, by starting accelerated cooling following rolling, γ grains suppressed to a thickness of 12 μm or less are present on the surface without coarsening, and precipitated from the grain boundaries of these fine γ grains. Since bainite can be precipitated, it is possible to obtain a steel sheet with improved drop weight characteristics exhibiting low temperature toughness and crack propagation.

Hereinafter, the present invention will be described in detail.
First, the reasons for limiting the steel plate components will be described. Unless otherwise specified, “%” notation relating to ingredients means mass%.

C: 0.02-0.15%
C is a strengthening element. It contributes to high strength by increasing the strength by pearlite structure or iron-based carbide, or by forming precipitates in combination with Nb or Ti. Alternatively, fixing the dislocations in the ferrite contributes to an increase in yield stress. Since these effects are difficult to obtain with an addition of less than 0.02%, 0.02% is set as the lower limit. On the other hand, addition exceeding 0.15% makes these reinforcements too strong and causes deterioration of moldability. In addition, 0.07 to 0.12% is a more preferable range.
Si: 0.01 to 1.0%
Si acts as a deoxidizer, is necessary for steelmaking, and may increase the yield stress due to solid solution strengthening. However, the upper limit of 1.0% is the upper limit for addition exceeding 1.0% because of the decrease in weldability and the ease of forming oxides on the steel sheet surface during annealing. On the other hand, when the content is less than 0.01%, the function as a deoxidizing agent is reduced and the production cost is increased, so this value is the lower limit.

Mn: 0.3 to 3.0%
Mn is a strengthening element. However, excessive addition promotes the formation of martensite, bainite, and the like, resulting in a decrease in yield ratio. For these reasons, the upper limit is 3.0%. On the other hand, if Mn is less than 0.3%, it is difficult to obtain the required strength, so 0.3% is made the lower limit. 1.0 to 1.8% is a more preferable range.

P: 0.015% or less P is a strengthening element of steel, but promotes deterioration in terms of toughness. P also has an effect of embrittlement of the welded portion of the steel plate. In view of these, the appropriate range is limited to 0.015% or less. Although the lower limit value of P is not particularly defined, it is preferable to set this value as the lower limit value because it is economically disadvantageous to set it to less than 0.001%.
S: 0.01% or less S adversely affects weldability and manufacturability during casting and hot rolling. Therefore, the upper limit is set to 0.01% or less. Although the lower limit of S is not particularly defined, it is preferable to set this value as the lower limit because it is economically disadvantageous to make it less than 0.0001%. In addition, since S is combined with Mn to form coarse MnS, the hole expandability is lowered. For this reason, it is necessary to reduce as much as possible in order to improve hole expansibility.

N: 0.001 to 0.01%
N is preferably added because it is effective in fixing dislocations and obtaining a large yield point elongation. On the other hand, it is necessary to suppress the content to 0.01% or less because coarse nitrides are formed or blowholes are generated during welding.
Cu: 0.05 to 1.5%
Ni: 0.1 to 10%
Cu and Ni are elements that can increase strength while maintaining high toughness, have little effect on HAZ toughness, are useful elements for increasing strength, and are selected according to desired characteristics. Can exist now. Cu is preferably contained in an amount of 0.05% or more, but when the content exceeds 1.5%, hot brittleness tends to occur. Ni is preferably contained in an amount of 0.1% or more. However, even if it is contained in a large amount, the effect is saturated, and an effect commensurate with the content cannot be expected, which is economically disadvantageous. The Cu content is preferably in the range of 0.1 to 0.5%, and the Ni content is preferably in the range of 0.3 to 0.7%.

Nb: 0.005 to 0.05%
Nb is preferably contained in an amount of 0.005% or more, but if it exceeds 0.05%, the base material toughness and the HAZ toughness are adversely affected. For this reason, when it contains, it is desirable to limit to 0.05% or less.
Ti: 0.003-0.1%
Ti contributes to improving the strength of the steel, and has a strong affinity with N and precipitates as TiN during solidification, thereby suppressing the austenite grain coarsening in HAZ and contributing to high toughness. However, if contained in a large amount, the base material toughness deteriorates, so a range of 0.003 to 0.1% is desirable. The Ti content is preferably in the range of 0.005 to 0.025%.

  In addition, B, Cr, Mo, V, Ca, REM (REM refers to elements of the La and lanthanoid series, for example, elements of the series such as La and Ce, etc.) 1) or more of Mg may be added.

B: 30 ppm or less B has an effect of improving the strength of steel through the improvement of hardenability by adding a small amount. On the other hand, if it is contained in a large amount, the hardenability is remarkably increased, and the toughness and ductility of the base metal are deteriorated and the weldability is lowered. For this reason, B content shall be 30 ppm or less.
Cr: 0.05-0.2%
Mo: 0.05-0.5%
Cr and Mo increase the steel plate strength and the yield ratio by solid solution strengthening and structure strengthening in the same manner as Mn. However, since this effect cannot be obtained unless it is 0.01% or more, the lower limit was made 0.05%. Moreover, addition of an amount exceeding the respective upper limit values is not preferable because it causes a delay of pearlite transformation in the annealing line and lowers the yield ratio. The Cr content is preferably in the range of 0.05 to 0.15%, and the Mo content is preferably in the range of 0.05 to 0.20%.

V: 0.005-0.05%
Since V is a carbide forming element, it can be added because the strength and yield ratio can be increased by precipitation strengthening or fine grain strengthening, as with Nb and Ti. Since this effect is easily obtained by addition of 0.005% or more, the lower limit is 0.005%. On the other hand, addition of a large amount not only increases the cost but also affects the structure, so 0.05% is made the upper limit. The V content is preferably in the range of 0.010 to 0.050%.
Ca: 0.0001 to 0.01%
Since Ca has an effect of improving toughness through refinement of crystal grains, Ca may be added to improve toughness. Addition of 0.0001% or more makes it easier to obtain the effect, so 0.0001% or more is preferable. On the other hand, if the content exceeds 0.01%, the effect is saturated. Become.

Mg: 0.0001 to 0.01%
REM: 0.005-0.1%
Since REM and Mg contribute to the fine dispersion of inclusions, particularly oxides, by adding an appropriate amount, 0.0001% or more of Mg and 0.005% or more of REM may be added. On the other hand, excessive addition reduces the manufacturability such as castability and hot workability, and the ductility of the steel sheet product, so the upper limit is set.
REM (for example, a lanthanoid series element such as Ce or La) is often added by misch metal, and often contains a lanthanoid series element such as La or Ce in combination. Even if these elements are included in combination, the effect of refinement of inclusions can be obtained if the REM satisfies the scope of the present invention. However, the effects of the present invention are exhibited even when metal La or Ce is added.
In addition to the above-mentioned elements, there are unavoidable impurities such as Sn and Sb. However, even if these elements are contained in a total content of 0.2% or less, the effect of the present invention is not impaired.

Pcm: 0.25% or less In the present invention, the content of each component is further adjusted so that Pcm defined by the following formula (1) is 0.25% or less.
Pcm = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5B (1)
Here, C, Si, Mn, Ni, Cu, Cr, Mo, V, B: The content (% by mass) of each element is 0.
Pcm is an index of the cold cracking property of the welded portion, and is desirably as low as possible. When Pcm exceeds 0.25, the weldability is remarkably deteriorated, so Pcm is limited to 0.25% or less.

Moreover, Ceq which shows a carbon equivalent is computed in the following formula | equation.
Ceq (%) = C + Mn / 6 + (Cr + Mo + V) / 5 + (Ni + Cu) / 15
About this carbon equivalent, the range of 0.30-0.50 is preferable.

Next, the reasons for limiting the production conditions of the steel sheet of the present invention will be described.
In the present invention, molten steel is melted by a conventional method using a converter, electric furnace, vacuum melting furnace or the like to obtain a steel material, and this steel material (slab) is reheated to a temperature in the range of 1000 ° C to 1200 ° C. . In addition, after this reheating treatment, a solution treatment that is maintained at 1000 to 1200 ° C. for 2 to 48 hours is performed, and after the steel material is strained or the solid solution of unnecessary precipitates is promoted, the subsequent rolling step You may use for.
When the reheating temperature is less than 1000 ° C., Nb and Ti carbonitrides precipitated during casting cannot be re-dissolved, deformation resistance in hot rolling is increased, and a reduction amount per pass is large. Since it cannot be removed, the number of rolling passes increases, the rolling efficiency decreases, and casting defects in the steel material (slab) cannot be crimped.
On the other hand, if the reheating temperature exceeds 1200 ° C., the γ grains existing as crystals may be coarsened, and if the γ grains are coarsened, the toughness may be reduced. If coarse γ grains remain, the steel sheet structure may not be adjusted by subsequent processing. Further, as the heating temperature becomes higher, surface flaws are more likely to occur due to the scale during heating, and the maintenance load after rolling increases. For this reason, it is preferable to make the reheating temperature of a steel raw material into the range of 1000-1200 degreeC.

Immediately after the reheating, hot rolling is performed in which the cumulative rolling reduction at Ar ₃ point + 20 ° C. to Ar ₃ point + 100 ° C. is 20% or more. If rolling is performed at a temperature lower than Ar ₃ point + 20 ° C. during rolling, there is a possibility that the rolling cannot be sufficiently reduced, and work hardening may occur and strength and toughness may be reduced. Further, when rolling at a temperature exceeding Ar ₃ point + 100 ° C. during rolling, the structure may not be refined even if a subsequent step is performed.
The cumulative rolling reduction in the rolling process performed in the temperature range is 8% or more and 50% or less, more preferably 20% or more and 50% or less. Unless the steel sheet is sufficiently reduced in the above temperature range, the austenite of the steel sheet structure cannot be sufficiently flattened. Further, if the rolling reduction rate exceeds 50%, an appropriate rolling amount per pass in normal equipment cannot be obtained, so 50% or less is preferable.
In this step, for example, a condition for terminating the hot rolling step at a temperature of 830 ° C. or lower can be adopted as an example. For example, by rolling at 830 ° C. or less, it becomes possible to give anisotropy to the γ grains of the steel sheet surface layer (for example, a region having a distance of 10 μm or less from the steel sheet surface) to form a structure that is long in the rolling direction. By controlling the thickness (thickness in the thickness direction of the steel plate), the crack propagation stop characteristic can be improved.

After completion of the rolling process, the cooling is performed by means such as water cooling without taking time (time: t ≦ 120 seconds). The temperature range of the cooling start point is Ar ₃ points−30 ° C. to Ar ₃ points + 80 ° C. Cooling is started from this temperature, and an accelerated cooling process of cooling to a temperature range of 500 ° C. or lower is performed at a cooling rate of 5 to 100 ° C./sec.
In this step, for example, water cooling is started at a cooling rate of 10 ° C./sec or higher from a temperature of 680 ° C. or higher, and after cooling at 150 to 350 ° C., the water is cooled and tempered. can do.
In the accelerated cooling step, it is important to adjust the temperature range when cooling is started by applying water, and for this purpose, the range of Ar ₃ points −30 ° C. to Ar ₃ points + 80 ° C. is used. The upper limit of the temperature range was defined as Ar ₃ point + 80 ° C., which was defined as a temperature slightly lower than the maximum temperature Ar ₃ point + 100 ° C. in the rolling process.
The upper limit of the stop temperature of this accelerated cooling is 500 ° C., and it is necessary to cool to 500 ° C. or less. This is because the steel sheet structure does not transform unless accelerated cooling to this temperature or less, and the transformation is satisfactory. Otherwise, the required strength cannot be obtained.
The transformation of the steel sheet structure is to impart anisotropy that lengthens in the rolling direction by rolling to the γ grains having the anisotropy of the surface layer, and the thickness of the old γ grains is 12 μm or less, more preferably 10 μm or less. By controlling it, a fine bainite crystal having a diameter of 5 μm or less can be deposited on the surface layer, which means a transformation when this bainite is deposited.

As described above, if it is a structure containing a large amount of refined old γ grains that suppresses the generation of coarse γ grains, a fine bainite structure precipitates from the grain boundaries of the γ grains, and the target steel sheet structure Can be obtained.
If the above accelerated cooling is applied to the rolled steel sheet having the composition and structure described above, the ratio of precipitation of the martensite structure can be reduced, and the structure can be tempered to a structure mainly composed of a ferrite structure and a bainite structure. A steel sheet structure having a balance between high strength and high toughness can be obtained.
In the steel sheet structure obtained by the method described above, the diameter of the old γ grains existing in the surface layer portion in the above-mentioned γ grains can be adjusted to a grain size of 12 μm or less, preferably 10 μm or less, from which a bainite structure is precipitated. Therefore, the bainite crystals having a diameter of 5 μm or less can be precipitated in an area ratio of 10% or more in the surface layer portion of the steel sheet.

The steel plate manufactured as described above exhibits a high strength of the 610 MPa class and exhibits excellent characteristics even in the NRL drop weight test, which is an index of crack propagation stop characteristics, and is -40 ° C or higher. Even in the low temperature range, excellent crack propagation stop characteristics can be obtained. Therefore, if it is the thick steel plate obtained by this invention, it has the characteristic outstanding as low temperature tanks, such as a LPG tank, or various low temperature container uses.
In addition, you may use, after performing surface treatments, such as pickling, with respect to the thick steel plate obtained in this way as needed. Pickling can remove oxides on the surface of the steel sheet. Pickling may be performed inline or offline. Moreover, pickling may be performed once, or pickling may be performed in a plurality of times.

Next, the present invention will be described in detail with reference to examples.
A steel material (slab) having the components shown in Table 1 is melted, and the steel material is heated to 1000 to 1200 ° C. to obtain the sheet thickness shown in Table 2 at the rolling temperature and rolling reduction shown in Table 2, and then Table 2 Accelerated cooling to the cooling end temperature was performed at the cooling coarse rate shown in FIG.
Γ grain thickness (μm), area ratio (%) of old γ grain crystal (former austenite grain crystal), tensile strength (TS), non-ductile transition temperature (NDT), mother The results of measuring the material ductility / brittle transition temperature (vTrs) are shown in Table 2.
Non-ductile transition temperature (NDT) is a value measured by a method called NRL (Naval Research Laboratory) drop weight test, and is an index for evaluating crack propagation stop characteristics.

FIG. 1 is a graph illustrating the correlation between the surface layer old γ grain thickness (μm) of each sample shown in Tables 1 and 2 and the value of the non-ductile transition temperature (NDT) measured for each sample.
From the correlation shown in FIG. 1, in order to bring the NDT value to a temperature lower than −40 ° C., in other words, to bring the NDT value far below −40 ° C. in the sample obtained in this example, the surface layer It has been found that it is important to make the thickness of the old γ grains 12 μm or less, and it is more preferable to make the thickness of the surface old γ grains 10 μm or less.

FIG. 2 shows the crystal grain size distribution of the surface layer portion of the thick steel plate of sample M having an NDT value of −40 ° C., and FIG. 3 shows the surface layer of the thick steel plate of sample C having a NDT value of + 5 ° C. The crystal grain size distribution of the part is shown. By measuring the area ratio of the left part from the line of the crystal grain diameter of 5 μm on the horizontal axis in FIGS. 2 and 3, the area ratio of the crystal having a diameter of 1 μm to 5 μm is calculated. This value can be the crystal area ratio of 5 μm or less. 2 and 3 indicate the area ratio of each particle size.
The crystal area ratio of the sample shown in FIG. 2 was 10%, the crystal area ratio of the sample shown in FIG. 3 was 3%, and the crystal area ratio of each sample shown in Table 2 was obtained for each sample. It is obtained from the crystal grain distribution.

  The present invention is a thick steel plate suitable for use in low temperature tanks, low temperature containers, pipelines, etc., with a balance of toughness that is excellent in strength of 610 MPa class and excellent crack propagation stopping characteristics, and has an industrial effect extremely high.

FIG. 1 is a graph illustrating the correlation between the surface layer old γ grain thickness (μm) of each sample and the value of the non-ductile transition temperature (NDT) measured for each sample. FIG. 2 is a view showing a crystal grain size distribution of a surface layer portion in a thick steel plate of Sample M in which an NDT value is −40 ° C. FIG. 3 is a graph showing the crystal grain size distribution of the surface layer portion of the thick steel plate of Sample C in which the NDT value is + 5 ° C.

Claims (4)

% By mass
C: 0.02% to 0.15%,
Si: 0.01 to 1.0%,
Mn: 0.3-3.0%
P: 0.015% or less,
S: 0.01% or less,
Ni: 0.1 to 10%,
Ti: 0.003 to 0.1%,
N: 0.001 to 0.01%, and
Pcm = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5B is heated to 1000-1200 ° C. with a balance of 0.25% or less, the balance being iron and inevitable impurities, Ar ₃ points + 20 ° C.- After hot rolling with a cumulative rolling reduction of 8% or more at Ar ₃ points + 100 ° C., cooling starts from Ar ₃ points −30 ° C. to Ar ₃ points + 80 ° C. and ends at 500 ° C. or less. Performing at 5 to 100 ° C./sec, and controlling the thickness of the old γ grains on the surface layer to 12 μm or less, the crystals having a diameter of 5 μm or less are present in the surface layer after transformation at 10% or more by area ratio. A high-strength steel manufacturing method with excellent toughness and crack propagation stopping properties. Furthermore, as a steel component, in mass%,
Cu: 0.05-1.5%, Cr: 0.05-0.2%, Mo: 0.05-0.5%, Nb: 0.005-0.05%, V: 0.005- A steel material containing at least one of 0.05% and B: 30 ppm or less and satisfying Pcm = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5B is 0.25% or less is used. The method for producing a high-strength steel excellent in low-temperature toughness and crack propagation stopping characteristics according to claim 1.   Furthermore, as a steel component, in mass%, Y: 0.001-0.1%, Ca: 0.0001-0.01%, Mg: 0.0001-0.01%, REM: 0.005-0. The method for producing a high-strength steel excellent in low-temperature toughness and crack propagation stopping characteristics according to claim 1 or 2, wherein a steel material containing 1% or more of 1% is used.   The method for producing high-tensile steel excellent in low-temperature toughness and crack propagation stopping characteristics according to any one of claims 1 to 3, wherein the tempering is performed at a temperature of 500 ° C or higher and lower than the Ac1 transformation point after the accelerated cooling.

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