KR20250020620A - High-strength steel plate for hydrogen transport pipe and its manufacturing method and steel pipe for hydrogen transport - Google Patents

Abstract

A high-strength steel plate for hydrogen transport pipes having excellent HISc resistance and fatigue crack propagation resistance in a high-pressure hydrogen environment is provided. The high-strength steel plate for hydrogen transport pipe of the present invention contains, in mass%, C: 0.030 to 0.060%, Si: 0.01 to 0.50%, Mn: 0.80 to 1.80%, P: 0.015% or less, S: 0.0015% or less, Al: 0.010 to 0.080%, Cr: 0.05 to 0.50%, Nb: 0.005 to 0.080%, Ti: 0.005 to 0.020%, N: 0.0020 to 0.0080%, and Ca: 0.0005 to 0.0050%, and the average value + 3σ of Vickers hardness at 0.25 mm below the surface of the steel plate is 225 HV or less, and the upper 20% grain size in the central structure of the plate thickness is 30 µm. Below, when the stress intensity factor range ΔK is 45 (MPa·m ^1/2 ), the fatigue crack propagation rate is less than 2.0× ^10-2 (mm/cycle), and the tensile strength is 535 MPa or more.

Description

High-strength steel plate for hydrogen transport pipe and its manufacturing method and steel pipe for hydrogen transport

The present invention relates to a high-strength steel plate for hydrogen transport pipes, and more particularly, to a high-strength steel plate for hydrogen transport pipes suitable for use in line pipes used for transporting high-pressure hydrogen gas, and a method for manufacturing the same. In addition, the present invention relates to a hydrogen transport pipe using the above-described high-strength steel plate for hydrogen transport.

Typically, line pipes are manufactured by forming steel plates manufactured by plate mills or hot rolling mills into steel pipes by UOE forming, press bend forming, and roll forming.

Here, the line pipe used for transporting high-pressure hydrogen gas requires hydrogen embrittlement resistance in addition to strength, toughness, and weldability. In particular, since stress is repeatedly applied to the line pipe due to pressure fluctuations during operation, fatigue crack propagation resistance in a high-pressure hydrogen gas environment is required in order to extend the service life. In addition, hydrogen-induced stress cracking resistance (HISC (Hydrogen Induced Stress Cracking) resistance) in a high-pressure hydrogen gas environment is required. When the hydrogen pressure is about 15 MPa, a low-alloy steel having sufficient thickness is used. However, at pressures higher than that, the risk of hydrogen embrittlement destruction during use increases, so the low-alloy steel is not used, and an austenitic stainless steel such as SUS316L, which is less susceptible to hydrogen embrittlement than the low-alloy steel, is used.

Austenitic stainless steels are expensive and have low strength, so if they are designed to withstand high hydrogen pressure, the thickness increases, and the price of the hydrogen transport line pipe itself also increases. Therefore, for hydrogen transport line pipes, there has been a demand for steels that are lower cost and can also withstand high-pressure hydrogen gas environments.

To solve the above problem, for example, patent document 1 proposes an austenitic steel having a high manganese content.

Japanese Patent No. 6703608

The technology described in Patent Document 1 enables the provision of a low-cost steel material compared to austenitic stainless steels such as SUS316L. However, since the steel material described in Patent Document 1 is an austenitic alloy, it is expensive compared to general low-alloy steels. In addition, in the steel material described in Patent Document 1, HIS resistance and fatigue crack propagation resistance in a high-pressure hydrogen gas environment are not taken into consideration.

Therefore, the present invention, in consideration of the above problems, aims to provide a high-strength steel plate for hydrogen transport pipe having excellent HISc resistance and fatigue crack propagation resistance in a high-pressure hydrogen environment, together with an advantageous manufacturing method thereof.

In addition, the present invention aims to provide a hydrogen transport steel pipe using the high-strength steel plate for the hydrogen transport steel pipe.

The inventors of the present invention have repeatedly conducted numerous experiments and studies on the component composition, microstructure, and manufacturing conditions of steel in order to secure HIS/Crack growth resistance and fatigue crack propagation resistance in a high-pressure hydrogen gas environment. As a result, the following points were recognized. That is, when the standard deviation of the Vickers hardness at a depth of 0.25 mm below the surface of the steel sheet is σ, the average value + 3σ of the Vickers hardness at a depth of 0.25 mm below the surface of the steel sheet is controlled to 225 HV or less, and the top 20% grain size in the structure at the center of the plate thickness is set to 30 µm or less. Accordingly, the HIS/Crack growth resistance and fatigue crack propagation resistance are improved. Furthermore, in order to realize such a steel structure, it is necessary to strictly control the rolling conditions and cooling conditions, and the inventors have succeeded in discovering those conditions. The present invention has been made based on these recognitions.

That is, the gist of the present invention is as follows.

[1] In mass%,

C: 0.030∼0.060%,

Si: 0.01∼0.50%,

Mn: 0.80∼1.80%,

P: 0.015% or less,

S: 0.0015% or less,

Al: 0.010∼0.080%,

Cr: 0.05∼0.50%,

Nb: 0.005∼0.080%,

Ti: 0.005∼0.020%,

N: 0.0020∼0.0080% and,

Ca: Contains 0.0005∼0.0050%, and the remainder is Fe and unavoidable impurities.

When the standard deviation of Vickers hardness at 0.25 mm below the surface of the steel plate is σ, the average value of Vickers hardness at 0.25 mm below the surface of the steel plate + 3σ is 225 HV or less, and the upper 20% grain size at the center of the plate thickness is 30 ㎛ or less, and the structure has

A high-strength steel plate for hydrogen transport pipes, having a fatigue crack propagation rate of less than 2.0× ^10-2 (mm/cycle) when the stress intensity factor range ΔK is 45 (MPa·m ^1/2 ) and a tensile strength of 535 MPa or more.

[2] The above composition of ingredients is, additionally, in mass%,

Cu: 0.50% or less,

Ni: 0.50% or less,

Mo: 0.50% or less,

V: 0.1% or less,

Zr: 0.02% or less,

Mg: 0.02% or less and,

REM: 0.02% or less

A high-strength steel plate for hydrogen transport pipe, containing at least one type selected from among [1].

[3] After heating the steel sheet having the composition described in [1] or [2] above to a temperature of 1000 to 1250°C,

Total pressure reduction in the recrystallization temperature range: 35% or more and 55% or less,

Reduction ratio of the final rolling pass in the recrystallization temperature range: 10% or more,

(Lower limit temperature of the recrystallization temperature range -80℃) or higher, and the reduction ratio of the final rolling pass in the temperature range below the lower limit temperature of the recrystallization temperature range: 15% or higher

It is made into a steel plate by hot rolling,

After that, for the above steel plate,

Steel plate surface temperature at the start of cooling: Ar ₃ transformation point (℃) or higher,

Cooling start time difference across the entire plate: within 50 seconds;

Average cooling rate from 750℃ to 550℃ at 0.25㎜ below the steel plate surface: 15∼50℃/s

Average cooling rate from 750℃ to 550℃ at the center of thickness: 15∼50℃/s,

Cooling stop temperature: 250∼550℃ at the steel plate temperature at 0.25㎜ below the steel plate surface and at the center of the plate thickness

A method for manufacturing a high-strength steel plate for hydrogen transport pipe, which performs human cooling.

[4] A hydrogen transport pipe using the high-strength steel plate for hydrogen transport pipe described in [1] or [2] above.

The high-strength steel plate for hydrogen transport pipes of the present invention and the hydrogen transport pipes using the high-strength steel plate for hydrogen transport pipes have excellent HISc resistance and fatigue crack propagation resistance in a high-pressure hydrogen environment. The steel pipe has excellent HISc resistance in a high-pressure hydrogen environment even in a region including a welded portion in the steel pipe. Furthermore, according to the method for manufacturing a high-strength steel plate for hydrogen transport pipes of the present invention, it is possible to manufacture a high-strength steel plate for hydrogen transport pipes having excellent HISc resistance and fatigue crack propagation resistance in a high-pressure hydrogen environment.

Figure 1 is a schematic diagram explaining a method for collecting test pieces for evaluating HISc resistance in an embodiment.

(Form for carrying out the invention)

Hereinafter, the high-strength steel plate for hydrogen transport pipe of the present invention will be specifically described. In addition, the high-strength steel plate for hydrogen transport pipe of the present invention will also be simply referred to as a high-strength steel plate.

[Ingredients]

First, the composition of the high-strength steel plate of the present invention and the reasons for its limitation will be explained. In the following description, all units expressed as % are mass% unless otherwise specified.

C: 0.030∼0.060%

C effectively contributes to the improvement of strength, but sufficient strength cannot be secured when the C content is less than 0.030%, so the C content is set to 0.030% or more. The C content is preferably set to 0.035% or more. On the other hand, when the C content exceeds 0.060%, the hardness increases during accelerated cooling, so the HISc resistance deteriorates. Therefore, the C content is set to 0.060% or less. The C content is preferably set to 0.050% or less.

Si: 0.01∼0.50%

Si is added for deoxidation, but since the deoxidation effect is not sufficient when the Si content is less than 0.01%, the Si content is set to 0.01% or more. The Si content is preferably set to 0.05% or more. On the other hand, since the weldability deteriorates when the Si content exceeds 0.50%, the Si content is set to 0.50% or less. The Si content is preferably set to 0.45% or less.

Mn: 0.80∼1.80%

Mn effectively contributes to the improvement of strength, but its effect is not sufficiently expressed when the Mn content is less than 0.80%. Therefore, the Mn content is set to 0.80% or more. The Mn content is preferably set to 1.00% or more. The Mn content is more preferably set to 1.20% or more. On the other hand, when the Mn content exceeds 1.80%, the hardness increases during accelerated cooling, so that the HISc resistance deteriorates. Therefore, the Mn content is set to 1.80% or less. The Mn content is preferably set to 1.70% or less. The Mn content is more preferably set to 1.60% or less.

P: 0.015% or less

P, as an unavoidable impurity element, increases hardness and deteriorates HISc resistance. Since this tendency becomes remarkable when the P content exceeds 0.015%, the upper limit of the P content is set to 0.015%. The P content is preferably 0.008% or less. In addition, although the lower the P content, the better, excessive P removal causes an increase in refining costs, so from the viewpoint of refining costs, the P content is preferably 0.001% or more.

S: 0.0015% or less

S is an inevitable impurity element, and since it forms MnS inclusions in steel and deteriorates low-temperature toughness, the S content is preferably low, but up to 0.0015% is permissible. Therefore, the S content is set to 0.015% or less. The S content is preferably 0.0010% or less. In addition, although the lower the S content, the better, excessive deS leads to an increase in refining costs, so from the viewpoint of refining costs, the S content is preferably set to 0.0002% or more.

Al: 0.010∼0.080%

Although Al is added as a deoxidizer, its effect is not sufficiently expressed when the Al content is less than 0.010%. Therefore, the Al content is set to 0.010% or more. The Al content is preferably set to 0.015% or more. The Al content is more preferably set to 0.025% or more. On the other hand, when the Al content exceeds 0.080%, alumina clogging of the immersion nozzle occurs during continuous casting, so the Al content is set to 0.080% or less. The Al content is preferably set to 0.070% or less. The Al content is more preferably set to 0.040% or less.

Cr: 0.05∼0.50%

Cr, like Mn, is an effective element for obtaining sufficient strength even in low-C content steels, but its effect is not fully expressed when the Cr content is less than 0.05%. Therefore, the Cr content is set to 0.05% or more. The Cr content is preferably set to 0.10% or more. The Cr content is more preferably set to 0.15% or more. However, when the Cr content exceeds 0.50%, the hardenability becomes excessive, so the hardness increases during accelerated cooling and the HISc resistance deteriorates. Therefore, the Cr content is set to 0.50% or less. The Cr content is preferably set to 0.45% or less. The Cr content is more preferably set to 0.35% or less.

Nb: 0.005∼0.080%

Nb, when present as dissolved Nb, expands the non-recrystallization temperature range during hot rolling and contributes to grain refinement, but its effect is not sufficiently expressed when the Nb content is less than 0.005%. Therefore, the Nb content is set to 0.005% or more. The Nb content is preferably set to 0.010% or more. The Nb content is more preferably set to 0.025% or more. On the other hand, when the Nb content exceeds 0.080%, coarse carbides are created during solidification, so that the hydrogen-induced cracking resistance deteriorates. Therefore, the Nb content is set to 0.080% or less. The Nb content is preferably set to 0.060% or less. The Nb content is more preferably set to 0.055% or less.

Ti: 0.005∼0.020%

Ti has the effect of pinning austenite grains when heated as TiN, thereby suppressing grain growth. When the Ti content is less than 0.005%, TiN is not sufficiently generated, so the Ti content is set to 0.005% or more. The Ti content is preferably set to 0.008% or more. Furthermore, when the Ti content exceeds 0.020%, the generated TiN coarsens and sufficient toughness of the weld heat-affected zone is not obtained, so the Ti content is set to 0.020% or less. The Ti content is preferably set to 0.017% or less. The Ti content is more preferably set to 0.015% or less.

N: 0.0020∼0.0080%

N effectively contributes to the improvement of strength, but sufficient strength cannot be secured when the N content is less than 0.0020%. Therefore, the N content is set to 0.0020% or more. The N content is preferably set to 0.0025% or more. The N content is more preferably set to 0.0030% or more. On the other hand, when the N content exceeds 0.0080%, the hardness increases during accelerated cooling, so the HISc resistance deteriorates. Therefore, the N content is set to 0.0080% or less. The N content is preferably set to 0.0070% or less. The N content is more preferably set to 0.0050% or less.

Ca: 0.0005∼0.0050％

Ca is an effective element for improving hydrogen-induced cracking resistance by controlling the shape of sulfide inclusions, but when the Ca content is less than 0.0005%, the effect of its addition is not sufficient. Therefore, the Ca content is set to 0.0005% or more. The Ca content is preferably set to 0.0008% or more. The Ca content is more preferably set to 0.0015% or more. On the other hand, when the Ca content exceeds 0.0050%, not only the above-mentioned effect is saturated, but also the hydrogen-induced cracking resistance deteriorates due to the decrease in the cleanliness of the steel, so the Ca content is set to 0.0050% or less. The Ca content is preferably set to 0.0045% or less. The Ca content is more preferably set to 0.0035% or less.

Above, the basic components (essential components) of the high-strength steel plate of the present invention have been described. Among the component compositions of the high-strength steel plate of the present invention, components (remainder) other than the above may be Fe and unavoidable impurities.

The composition of the high-strength steel plate of the present invention may optionally contain, in addition to the above components, at least one selected from Cu, Ni, Mo, V, Zr, Mg, and REM within the following range.

Cu: 0.50% or less

Cu is an element effective for improving low-temperature toughness and increasing strength, and to obtain this effect, it is preferable that the Cu content be 0.05% or more. The Cu content is more preferably 0.10% or more. However, if the Cu content exceeds 0.50%, the surface of the steel sheet is likely to be scratched, so when Cu is contained, the Cu content is set to 0.50% or less. The Cu content is preferably set to 0.45% or less.

Ni: 0.50% or less

Ni is an element effective for improving low-temperature toughness and increasing strength, and to obtain this effect, it is preferable that the Ni content be 0.05% or more. The Ni content is more preferably 0.10% or more. On the other hand, since Ni is an expensive element, when containing Ni, the Ni content is set to 0.50% or less. The Ni content is preferably set to 0.45% or less.

Mo: 0.50% or less

Mo is an element effective in improving low-temperature toughness and increasing strength, and to obtain this effect, it is desirable to make the Mo content 0.05% or more. On the other hand, since Mo is an expensive element, when containing Mo, the Mo content is 0.50% or less. The Mo content is preferably 0.45% or less.

V: 0.1% or less

V is an element that can be optionally added to increase the strength and low-temperature toughness of the steel plate, but its effect is not fully expressed when the V content is less than 0.005%. Therefore, when V is contained, it is preferable that the V content be 0.005% or more. On the other hand, when the V content exceeds 0.1%, the toughness of the weld deteriorates, so when V is contained, it is preferable that the V content be 0.1% or less. The V content is more preferably 0.050% or less, and even more preferably 0.010% or less.

Zr: 0.02% or less, Mg: 0.02% or less, REM: 0.02% or less

Zr, Mg and REM (rare earth metals) are elements that can be optionally added to increase fatigue crack propagation resistance through grain refinement or to increase crack resistance through control of inclusion properties. For each element, its effect is not sufficiently expressed when the content is less than 0.0005%. Therefore, when these elements are contained, it is preferable that the content of each element be 0.0005% or more. On the other hand, when the content of each element exceeds 0.02%, its effect is saturated, so when Zr, Mg and REM are contained, it is preferable that the content of each element be 0.02% or less. The content of each of the above elements is more preferably 0.0050% or less, and even more preferably 0.0030% or less. Also, REM is a general term for 15 elements from Sc, Y, and lanthanum (La) with atomic number 57 to lutetium (Lu) with atomic number 71, and the REM content referred to here is the total content of these elements.

In addition, the remainder other than the above-mentioned elements is composed of Fe and inevitable impurities. However, as long as it does not impair the effect of the present invention, it does not interfere with the inclusion of other trace elements. For example, O is an element inevitably included in steel, but if its content is 0.0050% or less, preferably 0.0040% or less, it is allowed in the present invention.

[Hardness of 0.25mm below the steel plate surface]

It is important that the high-strength steel plate of the present invention has an average value + 3σ of the Vickers hardness (HV 0.5) at a depth of 0.25 mm below the surface of the steel plate, when the standard deviation of the Vickers hardness (HV 0.5) at a depth of 0.25 mm below the surface of the steel plate is σ, of 225 HV or less. By satisfying this condition, excellent HISc resistance can be obtained in a high-pressure hydrogen environment. When the average value + 3σ of the Vickers hardness (HV 0.5) at a depth of 0.25 mm below the surface of the steel plate exceeds 225 HV, since the deviation in hardness within the steel plate is large, hydrogen accumulates locally, resulting in deterioration of the HISc resistance in the portion where the hydrogen has accumulated locally. Here, the "Vickers hardness (HV 0.5) at 0.25 mm below the steel plate surface" is the Vickers hardness (HV 0.5) at 100 points at equal intervals along the plate width direction at positions 0.25 mm below the steel plate surface (positions at a depth of 0.25 mm from the surface of the steel plate toward the center of the plate thickness) from the tip and tail ends in the rolling direction of the steel plate, respectively. In addition, the tip of the steel plate in the rolling direction is a position 1 m downstream in the rolling direction from the tip of the steel plate. The tail end of the steel plate in the rolling direction is a position 1 m upstream in the rolling direction from the tip of the steel plate. The measurement was performed on an area excluding an abnormal portion near the ends in the plate width direction. Here, the hardness of the steel plate is measured with 0.5 kgf instead of the commonly used 10 kgf because by measuring with 0.5 kgf, the indentation becomes smaller, making it possible to obtain hardness information closer to the surface or hardness information more sensitive to the microstructure. If the Vickers hardness is measured with a test force smaller than 0.5 kgf, the indentation size becomes excessively small, which increases the measurement deviation, making it undesirable. The average value of the Vickers hardness at 0.25 mm below the surface of the steel plate + 3σ is preferably 220 HV or less. Furthermore, as an example, the average value of the Vickers hardness at 0.25 mm below the surface of the steel plate + 3σ is 200 HV or more.

[Top 20% particle size in the center of the plate thickness]

By refining the average crystal grain size, the fatigue crack propagation resistance improves, but there is a limit to refining the average crystal grain size when cooling is initiated above the Ar ₃ transformation point temperature. In the present invention, it is important to suppress the formation of coarse crystal grains. That is, if the top 20% grain size is large, the fatigue crack propagation resistance deteriorates. In particular, in a structure in which the top 20% of the crystal grain size distribution in the center of the plate thickness exceeds 30 µm, the fatigue crack propagation resistance deteriorates significantly because cracks easily propagate. Therefore, it is necessary to have a structure in which the top 20% grain size in the center of the plate thickness (at the position of 1/2 the plate thickness) is 30 µm or less. The top 20% grain size is preferably 25 µm or less. Further, as an example, the top 20% grain size is 15 µm or more. In addition, the top 20% grain size refers to the grain size corresponding to the 20% position from the larger grain size when the grain sizes are arranged in order of size in the distribution of grain sizes. The measurement range of the grain size was 1 mm x 1 mm at the center of the plate thickness. More specifically, the grain size was determined as a boundary having an orientation difference of 15° or more as a result of analyzing the structure at the center of the plate thickness by the EBSD (Electron Backscatter Diffraction) method, and the diameter of the circle equivalent diameter was calculated as the grain size from the area of each grain. In addition, in the present invention, a grain size corresponding to 20% of the accumulated relative frequency from the larger grain size is referred to as the "top 20% grain size" in a frequency distribution table created for all the grains to be measured.

[Fatigue crack propagation rate]

The high-strength steel plate of the present invention, in a fatigue crack propagation test in a 21 MPa high-pressure hydrogen gas, has a fatigue crack propagation speed of less than 2.0× ^10-2 (mm/cycle) when the stress intensity factor range ΔK is 45 (MPa·m ^1/2 ). Preferably, the fatigue crack propagation speed is 1.5× ^10-2 (mm/cycle) or less. The lower the fatigue crack propagation speed, the more preferable it is. In addition, as an example, the fatigue crack propagation speed is 1.0× ^10-2 (mm/cycle) or more.

[tensile strength]

The high-strength steel plate of the present invention is mainly intended for use as a steel plate for steel pipes having a strength of API 5L X65 grade or higher, and therefore has a tensile strength of 535 MPa or higher. In addition, the upper limit of the tensile strength of the high-strength steel plate of the present invention is not particularly limited, but as an example, the tensile strength of the high-strength steel plate of the present invention is 760 MPa or lower. In addition, the tensile strength of the high-strength steel plate of the present invention may be 600 MPa or lower.

[Thickness of high-strength steel plate]

The thickness of the high-strength steel plate of the present invention is not particularly limited, but is preferably 12 mm or more. In addition, the thickness of the high-strength steel plate of the present invention is not particularly limited, but is preferably 39 mm or less.

[Manufacturing method]

Hereinafter, the manufacturing method and manufacturing conditions for manufacturing the above high-strength steel plate will be specifically described.

The method for manufacturing a high-strength steel plate of the present invention comprises heating a steel plate (slab) having the above-described composition, hot rolling the steel plate to form a steel plate (hot rolling process), and then cooling the steel plate under predetermined conditions (cooling process).

[Heating temperature of the steel sheet]

Heating temperature of steel sheet: 1000∼1250℃

When the heating temperature of the billet (slab) is less than 1000°C, the solid solution of carbides becomes insufficient, and the amount of solid solution strengthening by solid solution C, etc. decreases, so the required strength is not obtained. On the other hand, when the heating temperature of the billet exceeds 1250°C, the crystal grains become extremely coarsened and the fatigue crack propagation resistance deteriorates, so the heating temperature of the billet is set to 1000 to 1250°C. The heating temperature of the billet is preferably 1030°C or higher. In addition, the heating temperature of the billet is preferably 1200°C or lower. In addition, the billet (slab) is heated to the center at the above-mentioned heating temperature.

[Total pressure reduction in the recrystallization temperature range: 35% or more and 55% or less]

In order to refine the upper 20% grain size in the central structure of the plate thickness, it is necessary to promote recrystallization of the grains and suppress the formation of coarse grains in hot rolling in the recrystallization temperature range. If the total reduction ratio in the recrystallization temperature range is less than 35%, coarse grains remain because recrystallization is insufficient. Therefore, the total reduction ratio in the recrystallization temperature range is set to 35% or more, preferably 38% or more. On the other hand, if the total reduction ratio in the recrystallization temperature range exceeds 55%, coarsening of the grains can be suppressed, but grain refinement cannot be achieved because the reduction in the non-recrystallization range is insufficient. Therefore, the total reduction ratio in the recrystallization temperature range is set to 55% or less, preferably 52% or less. Here, the lower limit temperature Tnr (℃) of the recrystallization temperature range can be obtained, for example, from the component of the steel by the following equation. In addition, the temperature in hot rolling is the surface temperature of the rolled material (steel billet or steel plate), and the surface temperature can be measured using a radiation thermometer, etc.

Tnr(℃)=174×log[％Nb][％C+(12/14)%N]＋1444

However, [%X] in the above formula represents the content of element X in the steel (mass%).

[Reduction ratio of the final rolling pass in the recrystallization temperature range: 10% or more]

In addition to setting the total reduction ratio in the recrystallization temperature range to 35% or more and 55% or less, it is necessary to sufficiently secure the reduction ratio of the final rolling pass in the recrystallization temperature range and sufficiently promote recrystallization so as to initiate partial recrystallization rolling in a state of uniform grains without coarse grains. If the reduction ratio of the final rolling pass in the recrystallization temperature range is less than 10%, recrystallization is insufficient, so that coarse grains grow during the holding time from rough rolling to the start of finish rolling. Therefore, the reduction ratio of the final rolling pass in the recrystallization temperature range is set to 10% or more, preferably 11% or more. The upper limit of the reduction ratio of the final rolling pass in the recrystallization temperature range is not particularly limited, and a higher value is preferable. As an example, the reduction ratio of the final rolling pass in the recrystallization temperature range is 20% or less.

[(Lower limit temperature of recrystallization temperature range -80℃) or higher, Final rolling pass reduction ratio in temperature range below the lower limit temperature of recrystallization temperature range: 15% or higher]

Even after rolling is completed in the recrystallization temperature range, partial recrystallization occurs, so it is possible to promote recrystallization by further increasing the reduction ratio in the temperature range from above (the lower limit temperature of the recrystallization temperature range -80°C) to below the lower limit temperature of the recrystallization temperature range. Accordingly, the upper 20% grain size in the structure at the center of the plate thickness can be effectively refined. Therefore, the reduction ratio of the final rolling pass in the temperature range from above (the lower limit temperature of the recrystallization temperature range -80°C) to below the lower limit temperature of the recrystallization temperature range is set to 15% or more, preferably 16% or more. The upper limit of the reduction ratio of the final rolling pass in the above temperature range is not particularly limited, and a higher temperature is preferable. As an example, the reduction ratio of the final rolling pass in the above temperature range is 25% or less.

Rolling at a temperature lower than (the lower limit temperature of the recrystallization temperature range -80℃) is effective for grain refinement because rolling at a low temperature introduces a lot of strain. Therefore, it is preferable to roll at a low temperature lower than (the lower limit temperature of the recrystallization temperature range -80℃) within a range where the cooling start temperature of cooling can be observed.

[Rolling end temperature]

In the hot rolling process, in order to make the crystal grains fine, the lower the rolling end temperature is, the better. On the other hand, from the viewpoint of securing HISc resistance in a high-pressure hydrogen environment, it is necessary to set the rolling end temperature based on the fact that the cooling start temperature in the cooling process after the hot rolling process must be higher than the Ar ₃ transformation point in terms of the steel sheet surface temperature. Here, the Ar ₃ transformation point means the ferrite transformation start temperature during cooling, and can be obtained, for example, from the following equation from the steel composition. In addition, the surface temperature of the steel sheet can be measured with a radiation thermometer, etc.

Ar ₃ transformation point (℃) = 910-310 [％C] - 80 [％Mn] - 20 [％Cu] - 15 [％Cr] - 55 [％Ni] - 80 [％Mo]

However, in the above formula, [%X] represents the content (mass%) of element X in the steel, and elements not contained are set to 0.

[Cooling start temperature of cooling]

Cooling start temperature: Steel plate surface temperature above Ar ₃ transformation point (℃)

After the hot rolling process, cooling (controlled cooling) is performed on the steel plate. If the steel plate surface temperature at the start of cooling is lower than the Ar ₃ transformation point (℃), ferrite is generated before cooling, which significantly reduces the strength. Therefore, the steel plate surface temperature at the start of cooling is set to be higher than the Ar ₃ transformation point (℃). In addition, the steel plate surface temperature at the start of cooling is the temperature of the steel plate surface region where the cooling start temperature is the lowest. Specifically, the steel plate surface temperature at the start of cooling is, for example, the steel plate surface temperature at the end of the steel plate when cooling is performed while moving the steel plate in one direction relative to the cooling device. In addition, for example, if cooling is performed for each region of the entire steel plate and the times at which cooling is started are different between the regions, it is the steel plate surface temperature of the region that was cooled last. In addition, as an example, the upper limit of the steel plate surface temperature at the start of cooling is the rolling end temperature.

[Cooling start time of cooling]

Cooling start time difference across the entire plate: within 50 seconds

If the cooling start time difference for the entire steel sheet exceeds 50 seconds, the temperature difference within the steel sheet becomes large, so the variation in the temperature of the steel sheet when cooling is stopped becomes large, and the variation in the Vickers hardness at 0.25 mm below the surface of the steel sheet becomes large, and the HISc resistance deteriorates. Therefore, the cooling start time difference for the entire steel sheet is set to within 50 seconds, and preferably within 45 seconds. Specifically, for example, in the case of cooling while moving the steel sheet in one direction with respect to a cooling device, the difference between the cooling start time at the tip of the steel sheet and the cooling start time at the tail of the steel sheet is set to within 50 seconds. In addition, for example, in the case of cooling being performed for each predetermined region of the entire steel sheet and the times at which cooling is started between the regions are different, the difference between the cooling start time of the first region and the cooling start time of the last region is set to within 50 seconds. In addition, in the case where the entire steel sheet can be cooled at once, the cooling start time difference for the entire steel sheet may be 0 seconds.

[Cooling rate of cooling]

In order to achieve high strength while obtaining excellent HIS resistance, it is necessary to control the cooling rate at 0.25 mm below the steel plate surface and at the center of the plate thickness.

Average cooling rate from 750℃ to 550℃ at 0.25㎜ below the steel plate surface: 15∼50℃/s

It is important to make the average cooling rate from 750°C to 550°C at a steel plate temperature of 0.25 mm below the steel plate surface as slow as possible to form granular bainite. Since the temperature range from 750°C to 550°C is an important temperature range for bainite transformation, it is important to control the cooling rate in this temperature range. If the average cooling rate in the above temperature range exceeds 50°C/s, there is a concern that hardness deviation may occur, and the HISc resistance after pipe forming deteriorates. Therefore, the average cooling rate is set to 50°C/s or less. Preferably, it is set to 45°C/s or less. On the other hand, if the cooling rate is excessively small, ferrite or pearlite is generated, resulting in insufficient strength. Therefore, from the viewpoint of preventing this, the average cooling rate in the above temperature range is set to 15°C/s or more, and preferably 17°C/s or more. In addition, for cooling in a temperature range of 550℃ or less at a steel plate temperature at 0.25mm below the steel plate surface, if the cooling rate is slow, cooling in a stable nucleate boiling state is not achieved, and there is a concern that hardness may vary at the extreme surface layer of the steel plate. Therefore, the average cooling rate from 550℃ at a steel plate temperature at 0.25mm below the steel plate surface to the cooling stop temperature is preferably 150℃/s or more. From the viewpoint of making it easier to suppress hardness variation, the average cooling rate is preferably 250℃/s or less.

Average cooling rate from 750℃ to 550℃ at the center of plate thickness: 15∼50℃/s

When the average cooling rate from 750°C to 550°C in the center of the plate thickness is less than 15°C/s, granular bainite structure is not obtained, resulting in a decrease in strength. Therefore, the average cooling rate from 750°C to 550°C in the center of the plate thickness is set to 15°C/s or more. From the viewpoint of suppressing the deviation of the structure, the average cooling rate is preferably set to 17°C/s or more. On the other hand, in order to suppress the deviation of the grain size, the average cooling rate is set to 50°C/s or less, and preferably 45°C/s or less. In addition, there is no particular limitation on cooling in a temperature range of 550°C or less from the steel plate temperature in the center of the plate thickness, but from the viewpoint of suppressing the deviation of the structure or grain size, the average cooling rate in the temperature range is preferably set to 15°C/s or more. In addition, from the above viewpoint, the average cooling rate in the temperature range is preferably set to 50°C/s or less.

In addition, the steel plate temperature at 0.25 mm below the steel plate surface and at the center of the plate thickness cannot be measured directly physically. However, based on the surface temperature at the start of cooling measured by a radiation thermometer and the surface temperature at the target stop of cooling, the temperature distribution within the plate thickness cross-section can be calculated by differential calculation using, for example, a process computer, and the temperature can be obtained in real time from the result. The temperature at 0.25 mm below the steel plate surface in the temperature distribution is referred to as "the steel plate temperature at 0.25 mm below the steel plate surface" in this specification, and the temperature at the center of the plate thickness in the temperature distribution is referred to as "the steel plate temperature at the center of the plate thickness" in this specification.

[Cooling stop temperature]

Cooling stop temperature: 250-550℃, the steel plate temperature at 0.25mm below the steel plate surface and at the center of the plate thickness

If the cooling stop temperature exceeds 550°C at a steel plate temperature of 0.25 mm below the steel plate surface and at the center of the plate thickness, bainite transformation becomes incomplete and sufficient strength cannot be obtained. Therefore, the cooling stop temperature is preferably set to 550°C or lower and 500°C or lower. In addition, if the cooling stop temperature is less than 250°C, the hardness increases and the HISC resistance deteriorates. Therefore, the cooling stop temperature is preferably set to 250°C or higher and 300°C or higher.

[Steel pipe for hydrogen transport]

By forming the high-strength steel plate of the present invention into a tubular shape by press bend forming, roll forming, UOE forming, etc. and then welding the butt joints, a hydrogen transport steel pipe (UOE steel pipe, electric resistance welded steel pipe, spiral steel pipe, etc.) suitable for transporting high-pressure hydrogen gas can be manufactured. In addition, by manufacturing a steel pipe using the high-strength steel plate of the present invention, a steel pipe having excellent HISc resistance can be manufactured even if a high-hardness region exists in the welded portion. In addition, in the present invention, high-pressure hydrogen means, for example, a hydrogen gas environment of 15 MPa or more.

For example, a UOE steel pipe is manufactured by groove-cutting the end of a steel plate, forming it into a pipe shape by a C press, a U press, or an O press, and then seam-welding the butt joint by internal and external welding, and further performing an expansion process as necessary. In addition, any welding method may be used as long as sufficient joint strength and joint toughness are obtained, but from the viewpoints of excellent weld quality and manufacturing efficiency, submerge arc welding is preferably used. In addition, expansion can also be performed on a steel pipe that has been formed into a pipe shape by press bend forming and then seam-welded at the butt joint.

(Example)

Steel (grades A to W) having the component composition shown in Table 1 was made into a slab by a continuous casting method, heated to the heating temperature shown in Table 2, and then hot rolled and cooled under the conditions shown in Table 2 to obtain a steel plate having the final plate thickness shown in Table 2. In the cooling process, controlled cooling was performed using a water-cooled controlled cooling device while the steel plate was run in one direction. Thereafter, the end portions of the steel plate were subjected to improvement processing, and formed into a steel pipe shape by a C press, a U press, and an O press, and then the buttress portions of the inner and outer surfaces were deep-welded by submerged arc welding, and a steel pipe was obtained through a pipe expansion process. In addition, the Ar ₃ transformation point and the lower limit temperature Tnr of the recrystallization temperature range in Table 1 were each obtained from the above-mentioned equations.

[Measurement of Vickers hardness]

In accordance with JIS Z 2244 (2009), the Vickers hardness (HV 0.5) of 100 points was measured at equal intervals along the sheet width direction at a position 0.25 mm below the surface of the steel plate, respectively, from the tip and tail in the rolling direction of the steel plate, for a cross-section perpendicular to the rolling direction. Then, the average value and the standard deviation σ of the Vickers hardness (HV 0.5) of a total of 200 points were obtained. In addition, the tip in the rolling direction of the steel plate is a position 1 m downstream in the rolling direction from the tip of the steel plate. The tail in the rolling direction of the steel plate is a position 1 m upstream in the rolling direction from the tip of the steel plate. In addition, the measurement was performed on an area excluding an abnormal portion near the ends in the sheet width direction. The average value + 3σ of the Vickers hardness at 0.25 mm below the surface of the steel plate is shown in Table 3.

[Calculation of the top 20% of admissions]

A sample for metal structure observation was collected from the central portion of the plate width obtained as described above. For this sample, a cross-section perpendicular to the plate width direction was polished to a mirror surface, and then etched with colloidal silica. Thereafter, crystal data was collected from a 1 mm × 1 mm field of view at the central position of the plate thickness by the EBSD (Electron Backscatter Diffraction) method (measurement step: 0.8 μm). After data collection, using OIM-Analysis (EDAX, OIM Analysis software), a boundary having an orientation difference of 15° or more was judged to be a crystal grain boundary, and the diameter of the circle equivalent diameter was calculated as the crystal grain size from the area of each crystal grain. In addition, a frequency distribution table was created for all the crystal grains of the measurement target, and the crystal grain size for which the cumulative relative frequency from the larger calculated crystal grain size was 20% was referred to as the "top 20% grain size." The results of the measurement are shown in Table 3.

[Derivation of fatigue crack propagation rate]

From the steel plate obtained as described above, a CT test specimen conforming to ASTM E 647 was collected so that the load application direction was parallel to the rolling direction. The CT test specimen was a 10 mm thick test specimen collected from a position 1/2 of the plate thickness. Then, the length of a fatigue crack was measured by the compliance method using a clip gage, and the fatigue crack propagation speed in 21 MPa high-pressure hydrogen gas was obtained. Then, the fatigue crack propagation speed (mm/cycle) in the stress intensity factor range ΔK of 45 (MPa·m ^1/2 ) was evaluated. The results are shown in Table 3.

[Measurement of tensile strength]

A tensile test was performed in accordance with the provisions of JIS Z2241 (2011) using a full thickness test piece perpendicular to the rolling direction as a tensile test piece, and the tensile strength and yield strength were measured. The results are shown in Table 3.

[Evaluation of my HISc]

As shown in Fig. 1, the HIS resistance was measured by flattening a test piece (coupon) cut from the obtained steel pipe, and then collecting a 3 mm x 10 mm x 50 mm test piece from the inner surface of the steel pipe. At this time, in addition to a test piece of only the base metal not including the weld, a test piece including both the weld and the base metal was collected. The inner surface, which is the surface to be inspected, was left with the black skin attached in order to leave the state of the outermost layer. In other words, 0.25 mm below the surface of the steel plate is included in the test piece. A stress of 90% of the actual yield strength (0.5% YS) of each steel pipe was applied to the test pieces collected in this manner, and a four-point bending test was performed in 21 MPa high-pressure hydrogen gas. After 720 hours of exposure, if no cracks were observed in both the test piece of only the base metal not including the weld and the test piece including both the weld and the base metal, it was judged that the HIS resistance was excellent (good) and was given an ○. Additionally, if cracks occurred in at least one test piece, it was judged as defective and marked as ×. The results are shown in Table 3.

The target scope of the present invention is as follows. A high-strength steel plate for hydrogen transport pipe, wherein the average value + 3σ of Vickers hardness at 0.25 mm below the surface of the steel plate is 225 HV or less. The top 20% grain size in the structure at the center of the plate thickness is 30 ㎛ or less. The fatigue crack propagation speed when the stress intensity factor range ΔK is 45 (MPa·m ^1/2 ) is less than 2.0×10 ^-2 (mm/cycle). The tensile strength is 535 MPa or more. In addition, no cracking was confirmed in the evaluation of the HISc resistance (4-point bending test).

As shown in Table 2, No. 1 to No. 9 and No. 33 to No. 35 are invention examples whose component compositions and manufacturing conditions satisfy the appropriate ranges of the present invention. As shown in Table 3, No. 1 to No. 9 and No. 33 to No. 35 were all high-strength steel plates, and the average value + 3σ of the Vickers hardness at 0.25 mm below the surface of the steel plate was 225 HV or less. The top 20% grain size in the structure at the center of the plate thickness was 30 ㎛ or less. When the stress intensity factor range ΔK was 45 (MPa·m ^1/2 ), the fatigue crack propagation rate was less than 2.0×10 ^-2 (mm/cycle). The tensile strength was 535 MPa or more. In addition, the HIS resistance was also good.

In this regard, Nos. 10 to 20 have steel plate component compositions outside the scope of the present invention. Nos. 10, No. 12, No. 15, and No. 19 had insufficient solid solution strengthening, and thus had insufficient strength. Nos. 11, No. 13, No. 14, No. 16, and No. 20 had poor HISc resistance because the Vickers hardness at 0.25 mm below the steel plate surface increased. Nos. 17 and No. 18 had insufficient suppression of grain growth by precipitates, and thus had poor fatigue crack propagation resistance.

No. 21 to No. 32 are comparative examples whose component compositions are within the scope of the present invention, but whose manufacturing conditions are outside the scope of the present invention. In No. 21, since the heating temperature of the slab was low, the solid solution of carbides was insufficient, resulting in low strength. In No. 22, since the heating temperature of the slab was high, the grains became coarser, and the fatigue crack propagation resistance deteriorated. In No. 23, since the total reduction ratio in the recrystallization temperature range was insufficient, coarse grains remained, and the fatigue crack propagation resistance deteriorated. In No. 24, since the total reduction ratio in the recrystallization temperature range was excessive, the top 20% grain size in the structure in the center of the plate thickness was large, and the fatigue crack propagation resistance deteriorated. In No. 25, since the reduction ratio in the final rolling pass in the recrystallization temperature range was insufficient, coarse grains remained, and the fatigue crack propagation resistance deteriorated. No. 26 had a large top 20% grain size in the central structure of the plate thickness because the reduction ratio in the final rolling pass in the temperature range from (the lower limit temperature of the recrystallization temperature range -80°C) to below the lower limit temperature of the recrystallization temperature range was insufficient, and the fatigue crack propagation resistance deteriorated. No. 27 had a low cooling initiation temperature and some ferrite was generated, and thus had low strength. No. 28 had a large cooling initiation time difference over the entire steel plate, and thus had a large variation in Vickers hardness at 0.25 mm below the surface of the steel plate, and thus had deteriorated HISc resistance. No. 29 had a low strength because the average cooling rate from 750°C to 550°C was low and some ferrite was generated. No. 30 had poor HIS resistance because the average cooling rate from 750℃ to 550℃ was high and the variation in Vickers hardness at 0.25mm below the surface of the steel plate was large. No. 31 had poor HIS resistance because the cooling stop temperature was low and the variation in Vickers hardness at 0.25mm below the surface of the steel plate was large. No. 32 had low strength because the cooling stop temperature was high and some ferrite was formed.

(Industrial applicability)

According to the present invention, it is possible to supply a high-strength steel plate for hydrogen transport pipe having excellent HISc resistance and fatigue crack propagation resistance in a high-pressure hydrogen environment.

Claims (4)

In mass %,
C: 0.030∼0.060%,
Si: 0.01∼0.50%,
Mn: 0.80∼1.80%,
P: 0.015% or less,
S: 0.0015% or less,
Al: 0.010∼0.080%,
Cr: 0.05∼0.50%,
Nb: 0.005∼0.080%,
Ti: 0.005∼0.020%,
N: 0.0020∼0.0080% and,
Ca: Contains 0.0005∼0.0050%, and the remainder is Fe and unavoidable impurities.
When the standard deviation of Vickers hardness at 0.25 mm below the surface of the steel plate is σ, the average value of Vickers hardness at 0.25 mm below the surface of the steel plate + 3σ is 225 HV or less, and the upper 20% grain size at the center of the plate thickness is 30 ㎛ or less, and the structure has
A high-strength steel plate for hydrogen transport pipes, having a fatigue crack propagation rate of less than 2.0× ^10-2 (mm/cycle) when the stress intensity factor range ΔK is 45 (MPa·m ^1/2 ) and a tensile strength of 535 MPa or more. In the first paragraph,
The above composition of ingredients, additionally, in mass%,
Cu: 0.50% or less,
Ni: 0.50% or less,
Mo: 0.50% or less,
V: 0.1% or less,
Zr: 0.02% or less,
Mg: 0.02% or less and,
REM: 0.02% or less
High-strength steel plate for hydrogen transport pipe, containing at least one type selected from the following. After heating the steel sheet having the composition described in clause 1 or clause 2 to a temperature of 1000 to 1250°C,
Total pressure reduction in the recrystallization temperature range: 35% or more and 55% or less,
Reduction ratio of the final rolling pass in the recrystallization temperature range: 10% or more,
(Lower limit temperature of the recrystallization temperature range -80℃) or higher, and the reduction ratio of the final rolling pass in the temperature range below the lower limit temperature of the recrystallization temperature range: 15% or higher
It is made into a steel plate by hot rolling,
After that, for the above steel plate,
Steel plate surface temperature at the start of cooling: Ar ₃ transformation point (℃) or higher,
Cooling start time difference across the entire plate: within 50 seconds;
Average cooling rate from 750℃ to 550℃ at 0.25㎜ below the steel plate surface: 15∼50℃/s
Average cooling rate from 750℃ to 550℃ at the center of plate thickness: 15∼50℃/s
Cooling stop temperature: 250∼550℃ at the steel plate temperature at 0.25㎜ below the steel plate surface and at the center of the plate thickness
A method for manufacturing a high-strength steel plate for hydrogen transport pipe, which performs human cooling. A hydrogen transport pipe using the high-strength steel plate for hydrogen transport pipes described in claim 1 or 2.