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

Line pipes are typically produced by forming steel plates, which are produced by plate mills or hot rolling mills, into steel pipes using, for example, UOE forming, press bend forming, or roll forming.

Line pipes used for transporting high-pressure hydrogen gas are required to have hydrogen embrittlement resistance in addition to strength, toughness, weldability, and so forth. In particular, since line pipes are subjected to repeated stress due to pressure fluctuations during operation, fatigue crack growth resistance in a high-pressure hydrogen gas environment is required to extend their service life. In addition, hydrogen-induced stress cracking resistance (HISC resistance) in a high-pressure hydrogen gas environment is required. When the hydrogen pressure is about 15 MPa, low-alloy steel having a sufficient wall thickness is used. However, at a pressure higher than or equal to this, the risk of hydrogen embrittlement fracture during use increases. Thus, low-alloy steel is not used. Instead, austenitic stainless steel such as SUS316L, which is less susceptible to hydrogen embrittlement than low-alloy steel, is used.

Austenitic stainless steel has low strength in addition to the high cost of the steel material. Therefore, if austenitic stainless steel is designed to withstand high hydrogen pressure, the wall thickness of the austenitic stainless steel will be increased, and the hydrogen transport line pipe itself will also be more expensive. For this reason, there has been a demand for a steel material for hydrogen transport line pipes that is lower in cost and can withstand a high-pressure hydrogen gas environment.

To solve the above problems, for example, Patent Literature 1 discloses an austenitic steel material having a high Mn content.

PTL 1: Japanese Patent No. 6703608

According to the technique described in Patent Literature 1, it is possible to provide a steel material that is lower in cost than austenitic stainless steel such as SUS316L. However, the steel material described in Patent Literature 1 is an austenitic alloy and thus is more expensive than a typical low-alloy steel. Furthermore, in the steel material described in Patent Literature 1, HISC resistance and fatigue crack growth resistance in a high-pressure hydrogen gas environment are not taken into consideration.

In view of the above problems, it is an object of the present invention to provide a high strength steel plate for a hydrogen transport steel pipe, the high strength steel plate having excellent HISC resistance and fatigue crack growth resistance in a high-pressure hydrogen environment, together with an advantageous method for manufacturing the high strength steel plate.

It is another object of the present invention to provide a steel pipe for transporting hydrogen using the above-mentioned high strength steel plate for hydrogen transport steel pipe.

The inventors have repeatedly conducted numerous experiments and studies on the chemical compositions, microstructures, and production conditions of steel materials in order to ensure HISC resistance and fatigue crack growth resistance in a high-pressure hydrogen gas environment, and have found the following. That is, letting the standard deviation of a Vickers hardness at 0.25 mm below a surface of a steel plate be denoted as σ, the average value of the Vickers hardness at 0.25 mm below the surface of the steel plate + 3σ is controlled to 225 HV or less, and the top 20% grain size in the microstructure at the center of the plate thickness is 30 µm or less. This improves HISC resistance and fatigue crack growth resistance. To achieve such a steel microstructure, it is necessary to strictly control the rolling and cooling conditions. The inventors have succeeded in finding those conditions. The present invention has been made based on these findings.

The gist and configuration of the present invention are as described below.

1. [1] A high strength steel plate for a hydrogen transport steel pipe, including:
   a chemical composition containing, by 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%, the balance being Fe and incidental impurities; and
   • a microstructure in which letting the standard deviation of a Vickers hardness at 0.25 mm below a surface of the steel plate be denoted as σ, the average value of the Vickers hardness at 0.25 mm below the surface of the steel plate + 3σ is 225 HV or less, and the top 20% grain size at the center of the thickness of the plate is 30 µm or less,
   • in which a fatigue crack growth rate is less than 2.0 × 10^-2 (mm/cycle) when a stress intensity factor range ΔK is 45 (MPa·m^1/2), and a tensile strength is 535 MPa or more.
2. [2] In the high strength steel plate for a hydrogen transport steel pipe described in [1], the chemical composition further contains, by mass%, one or more selected from:
   • 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.
3. [3] A method for manufacturing a high strength steel plate for a hydrogen transport steel pipe, the method including:
   • heating a steel slab having the chemical composition described in [1] or [2] to a temperature of 1,000°C to 1,250°C;
   • then subjecting the steel slab to hot rolling at:
     • a total rolling reduction in a recrystallization temperature range: 35% or more and 55% or less,
     • a rolling reduction in a final rolling pass in the recrystallization temperature range: 10% or more, and
     • a rolling reduction in the final rolling pass in a temperature range of (the lower limit temperature in the recrystallization temperature range - 80°C) or higher and lower than the lower limit temperature in the recrystallization temperature range: 15% or more, to form a steel plate; and
   • thereafter subjecting the steel plate to cooling at:
     • the temperature of a surface of the steel plate at the start of cooling: an Ar₃ transformation point (°C) or higher,
     • a difference in cooling start time for the entirety of the steel plate: within 50 seconds,
     • an average cooling rate from 750°C to 550°C in terms of a steel plate temperature at 0.25 mm below the surface of the steel plate: 15 to 50 °C/s,
     • an average cooling rate from 750°C to 550°C in terms of the steel plate temperature at the center of the thickness of the plate: 15 to 50 °C/s, and
     • a cooling stop temperature in terms of the steel plate temperature at 0.25 mm below the surface of the steel plate and at the center of the thickness of the plate: 250°C to 550°C.
4. [4] A steel pipe for transporting hydrogen using the high strength steel plate for a hydrogen transport steel pipe described in [1] or [2].

The high strength steel plate of the present invention for the hydrogen transport steel pipe and the steel pipe for transporting hydrogen using the high strength steel plate for the hydrogen transport steel pipe have excellent HISC resistance and fatigue crack growth 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 the weld of the steel pipe. According to the method of the present invention for manufacturing a high strength steel plate for a hydrogen transport steel pipe, it is possible to manufacture a high strength steel plate for a hydrogen transport steel pipe, the high strength steel plate having excellent HISC resistance and fatigue crack growth resistance in a high-pressure hydrogen environment.

[Fig. 1] Fig. 1 is a schematic view illustrating a method for collecting a test specimen for evaluating HISC resistance in examples.

A high strength steel plate of the present invention for a hydrogen transport steel pipe will be specifically described below. Hereinafter, the high strength steel plate of the present invention for a hydrogen transport steel pipe is also referred to simply as a "high strength steel plate".

The chemical composition of the high strength steel plate of the present invention and the reasons for limitation thereof will be described. In the following description, units indicated by % are all % by mass unless otherwise specified.

C effectively contributes to improving strength. However, when the C content is less than 0.030%, sufficient strength cannot be ensured. Thus, the C content is 0.030% or more. The C content is preferably 0.035% or more. When the C content is more than 0.060%, the hardness increases during accelerated cooling, thereby deteriorating the HISC resistance. For this reason, the C content is 0.060% or less. The C content is preferably 0.050% or less.

Si is added for deoxidation. A Si content of less than 0.01% results in an insufficient deoxidation effect. For this reason, the Si content is 0.01% or more. The Si content is preferably 0.05% or more. A Si content of more than 0.50% results in a deterioration in weldability. For this reason, the Si content is 0.50% or less. The Si content is preferably 0.45% or less.

Mn effectively contributes to improving strength. However, when the Mn content is less than 0.80%, this effect is not sufficiently provided. For this reason, the Mn content is 0.80% or more. The Mn content is preferably 1.00% or more. The Mn content is more preferably 1.20% or more. A Mn content of more than 1.80% results in an increase in hardness during accelerated cooling, thereby deteriorating the HISC resistance. For this reason, the Mn content is 1.80% or less. The Mn content is preferably 1.70% or less. The Mn content is more preferably 1.60% or less.

P is an incidental impurity element that increases the hardness and thereby deteriorates the HISC resistance. When the P content is more than 0.015%, this tendency is marked. For this reason, the upper limit of the P content is 0.015%. The P content is preferably 0.008% or less. A lower P content is more preferred. However, excessive dephosphorization leads to an increase in refining cost. From the viewpoint of the refining cost, the P content is preferably 0.001% or more.

S is an incidental impurity element that forms MnS inclusions in steel to deteriorate low-temperature toughness. For this reason, a lower S content is more preferred. However, a S content of up to 0.0015% is acceptable. For this reason, the S content is 0.015% or less. The S content is preferably 0.0010% or less. A lower S content is more preferred. However, excessive desulfurization leads to an increase in refining cost. From the viewpoint of the refining cost, the S content is preferably 0.0002% or more.

Al is added as a deoxidizing agent. When the Al content is less than 0.010%, the effect is not sufficiently provided. For this reason, the Al content is 0.010% or more. The Al content is preferably 0.015% or more. The Al content is more preferably 0.025% or more. An Al content of more than 0.080% results in clogging of a submerged nozzle with alumina during continuous casting. For this reason, the Al content is 0.080% or less. The Al content is preferably 0.070% or less. The Al content is more preferably 0.040% or less.

Like Mn, Cr is an effective element for achieving sufficient strength even in steel having a low C content. When the Cr content is less than 0.05%, this effect is not sufficiently provided. For this reason, the Cr content is 0.05% or more. The Cr content is preferably 0.10% or more. The Cr content is more preferably 0.15% or more. A Cr content of more than 0.50% results in excessive hardenability to increase the hardness during accelerated cooling, thereby deteriorating the HISC resistance. For this reason, the Cr content is 0.50% or less. The Cr content is preferably 0.45% or less. The Cr content is more preferably 0.35% or less.

When Nb is present as solute Nb, Nb extends the non-recrystallization temperature range during hot rolling and contributes to refinement of crystal grain size. When the Nb content is less than 0.005%, this effect is not sufficiently provided. For this reason, the Nb content is 0.005% or more. The Nb content is preferably 0.010% or more. The Nb content is more preferably 0.025% or more. A Nb content of more than 0.080% results in crystallization of coarse carbides during solidification, deteriorating hydrogen induced cracking resistance. For this reason, the Nb content is 0.080% or less. The Nb content is preferably 0.060% or less. The Nb content is more preferably 0.055% or less.

Ti has the effect of pinning austenite grains in the form of TiN during heating, thereby inhibiting the growth of the grains. When the Ti content is less than 0.005%, TiN is not sufficiently formed. Thus, the Ti content is 0.005% or more. The Ti content is preferably 0.008% or more. When the Ti content is more than 0.020%, the formed TiN coarsens, resulting in a failure to obtain sufficient toughness of the weld heat affected zone. For this reason, the Ti content is 0.020% or less. The Ti content is preferably 0.017% or less. The Ti content is more preferably 0.015% or less.

N effectively contributes to improving strength. However, when the N content is less than 0.0020%, sufficient strength cannot be ensured. For this reason, the N content is 0.0020% or more. The N content is preferably 0.0025% or more. The N content is more preferably 0.0030% or more. A N content of more than 0.0080% results in an increase in hardness during accelerated cooling, thereby deteriorating the HISC resistance. For this reason, the N content is 0.0080% or less. The N content is preferably 0.0070% or less. The N content is more preferably 0.0050% or less.

Ca is an element effective in improving hydrogen induced cracking resistance by shape control of sulfide-based inclusions. However, when the Ca content is less than 0.0005%, this addition effect is insufficient. For this reason, the Ca content is 0.0005% or more. The Ca content is preferably 0.0008% or more. The Ca content is more preferably 0.0015% or more. A Ca content of more than 0.0050% results in saturation of the above-mentioned effect and a decrease in the cleanliness of the steel, deteriorating the hydrogen induced cracking resistance. For this reason, the Ca content is 0.0050% or less. The Ca content is preferably 0.0045% or less. The Ca content is more preferably 0.0035% or less.

The basic components (essential components) of the high strength steel plate of the present invention have been described above. In the chemical composition of the high strength steel plate of the present invention, components other than those mentioned above (the balance) can be Fe and incidental impurities.

The chemical composition of the high strength steel plate of the present invention can further contain, in addition to the above-mentioned components, one or more selected from Cu, Ni, Mo, V, Zr, Mg, and REM within the following ranges.

Cu is an element effective in improving low-temperature toughness and increasing strength. To provide these effects, the Cu content is preferably 0.05% or more. The Cu content is more preferably 0.10% or more. When the Cu content is more than 0.50%, defects on the steel plate surface are likely to occur. For this reason, when Cu is contained, the Cu content is 0.50% or less. The Cu content is preferably 0.45% or less.

Ni is an element effective in improving low-temperature toughness and increasing strength. To provide these effects, the Ni content is preferably 0.05% or more. The Ni content is more preferably 0.10% or more. Ni is an expensive element. For this reason, when Ni is contained, the Ni content is 0.50% or less. The Ni content is preferably 0.45% or less.

Mo is an element in improving low-temperature toughness and increasing strength. To provide these effects, the Mo content is preferably 0.05% or more. Mo is an expensive element. For this reason, when Mo is contained, the Mo content is 0.50% or less. The Mo content is preferably 0.45% or less.

V is an element that can be added to the steel plate to enhance its strength and low-temperature toughness. When the V content is less than 0.005%, the effect is not sufficiently provided. For this reason, when V is contained, the V content is preferably 0.005% or more. A V content of more than 0.1% results in a deterioration in the toughness of the weld. For this reason, when V is contained, the V content is preferably 0.1% or less. The V content is more preferably 0.050% or less, still more preferably 0.010% or less.

Zr, Mg, and a rare earth metal (REM) are elements that can be optionally added to enhance fatigue crack growth resistance through the refinement of crystal grains and to enhance crack resistance through control of inclusion properties. When the amount of each element contained is less than 0.0005%, the effects are not sufficiently provided. For this reason, when these elements are contained, the amount of each element contained is preferably 0.0005% or more. When the amount of each element contained is more than 0.02%, the effects are saturated. For this reason, when Zr, Mg, and REM are contained, the amount of each element contained is preferably 0.02% or less. The amount of each of the above elements contained is more preferably 0.0050% or less, still more preferably 0.0030% or less. REM is a general term for Sc, Y, and 15 elements from lanthanum (La) of atomic number 57 to lutetium (Lu) of atomic number 71. The amount of REM contained refers to the total amount of these elements.

The balance other than the above elements is Fe and incidental impurities. However, other trace elements may be contained as long as they do not impair the effects of the present invention. For example, O is an element that is incidentally contained in steel, and is acceptable in the present invention as long as the O content is 0.0050% or less, preferably 0.0040% or less.

In the high strength steel plate of the present invention, letting the standard deviation of a Vickers hardness (HV 0.5) at 0.25 mm below a surface of the steel plate be denoted as σ, it is important that the average value of the Vickers hardness (HV 0.5) at 0.25 mm below the surface of the steel plate + 3σ be 225 HV or less. Meeting this condition can provide excellent HISC resistance in a high-pressure hydrogen environment. When the average value of the Vickers hardness (HV 0.5) at 0.25 mm below the surface of the steel plate + 3σ is more than 225 HV, the hardness varies greatly within the steel plate, causing hydrogen to accumulate locally. The HISC resistance deteriorates in a region where hydrogen has accumulated locally. Here, the "Vickers hardness (HV 0.5) at 0.25 mm below the surface of the steel plate" refers to the Vickers hardness (HV 0.5) measured at 100 equally spaced points in the plate width direction from each of the leading end and the trailing end in the rolling direction of the steel plate, the 100 equally spaced points being located at 0.25 mm below the surface of the steel plate (at a depth of 0.25 mm from the surface of the steel plate toward the center of the plate thickness). The leading end of the steel plate in the rolling direction is a position 1 m downstream of the leading edge of the steel plate in the rolling direction. The trailing end of the steel plate in the rolling direction is a position 1 m upstream of the trailing edge of the steel plate in the rolling direction. The measurement was performed in a region excluding non-steady portions near the ends in the plate width direction. Here, the reason for measuring the hardness of the steel plate at 0.5 kgf instead of the commonly used 10 kgf is that measurement at 0.5 kgf results in a smaller indentation, making it possible to obtain hardness information at a position closer to the surface and hardness information that is more sensitive to the microstructure. Measuring the Vickers hardness with a test load of less than 0.5 kgf results in a too small indentation size and thus an increase in measurement variability, which is not preferred. 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. 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.

Refinement of the average crystal grain size improves fatigue crack growth resistance. However, when cooling is started at a temperature equal to or higher than the Ar₃ transformation point, there is a limit to the refinement of the average crystal grain size. In the present invention, it is important to inhibit the formation of coarse crystal grains. That is, when the top 20% grain size is large, the fatigue crack growth resistance deteriorates. In particular, in a microstructure in which the top 20% of a crystal grain size distribution at the center of the plate thickness is more than 30 µm, cracks propagate easily, thus significantly deteriorating the fatigue crack growth resistance. Therefore, it is necessary to have a microstructure in which the top 20% grain size at the center of the plate thickness (1/2 position of the plate thickness) is 30 µm or less. The top 20% grain size is preferably 25 µm or less. As an example, the top 20% grain size is 15 µm or more. The top 20% grain size is a grain size corresponding to the 20% position from the largest crystal grain size when the crystal grain sizes are arranged in descending order in the crystal grain size distribution.

The measurement range of the crystal grain size was 1 mm × 1 mm at the center position of the plate thickness. More specifically, the crystal grain size was determined as follows: The microstructure at the center of the plate thickness was analyzed by an electron backscatter diffraction (EBSD) method. A boundary having an orientation difference of 15° or more was determined to be a crystal grain boundary. The equivalent circular diameter was calculated as the crystal grain size from the area of each crystal grain. In the present invention, a frequency distribution table is prepared for all crystal grains to be measured, and the crystal grain size corresponding to 20% of the cumulative relative frequency from the largest crystal grain size calculated is referred to as the "top 20% grain size".

In a fatigue crack growth test of the high strength steel plate of the present invention in a high-pressure hydrogen gas of 21 MPa, the fatigue crack growth rate is less than 2.0 × 10^-2 (mm/cycle) when the stress intensity factor range ΔK is 45 (MPa·m^1/2). The fatigue crack growth rate is preferably 1.5 × 10^-2 (mm/cycle) or less. A lower fatigue crack growth rate is more preferred. As an example, the fatigue crack growth rate is 1.0 × 10^-2 (mm/cycle) or more.

The high strength steel plate of the present invention is intended mainly for steel plates for steel pipes having a strength of API 5L grade X65 or higher, and thus has a tensile strength of 535 MPa or higher. 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 less. The high strength steel plate of the present invention may have a tensile strength of 600 MPa or less.

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

A manufacturing method and manufacturing conditions for manufacturing the above-mentioned high strength steel plate are specifically described below.

The method for manufacturing the high strength steel plate of the present invention includes heating a steel slab (slab) having the above-mentioned chemical composition, then subjecting the steel slab to hot rolling to form a steel plate (hot rolling step), and thereafter cooling the steel plate under predetermined conditions (cooling step).

A heating temperature of the steel slab (slab) of lower than 1,000°C results in insufficient dissolution of carbides to reduce the amount of steel that is subjected to solid solution strengthening by solute C and so forth, resulting in a failure to achieve the required strength. A heating temperature of the steel slab of higher than 1,250°C results in excessively coarse crystal grains to deteriorate the fatigue crack growth resistance. Thus, the heating temperature of the steel slab is 1,000°C to 1,250°C. The heating temperature of the steel slab is preferably 1,030°C or higher. The heating temperature of the steel slab is preferably 1,200°C or lower. The steel slab (slab) is heated to the heating temperature all the way to the center.

To refine the top 20% grain size in the microstructure at the center of the plate thickness, the recrystallization of crystal grains is required to be promoted during hot rolling in the recrystallization temperature range to inhibit the formation of coarse grains. When the total rolling reduction in the recrystallization temperature range is less than 35%, recrystallization is insufficient, resulting in coarse grains remaining. Thus, the total rolling reduction in the recrystallization temperature range is 35% or more, preferably 38% or more. When the total rolling reduction in the recrystallization temperature range is more than 55%, coarsening of the crystal grains can be inhibited. However, the rolling reduction in the non-recrystallization range is insufficient, resulting in a failure to refine the crystal grains. Thus, the total rolling reduction in the recrystallization temperature range is 55% or less, preferably 52% or less. Here, the lower limit temperature Tnr (°C) of the recrystallization temperature range can be calculated, for example, from the components of the steel using the following formula. The temperature in the hot rolling is the surface temperature of the material being rolled (steel slab or steel plate), and the surface temperature can be measured with a radiation thermometer or the like.$$\mathrm{Tnr}\left({}^{\circ }\mathrm{C}\right)=174\times \log \left[\%\mathrm{Nb}\right]\left[\%\mathrm{C}+\left(12/14\right)\%\mathrm{N}\right]+1444$$ where in the above formula, each [%X] is the element X content (mass%) of the steel.

Rolling in the partial recrystallization range needs to be started in a uniform grain state free of coarse grains by ensuring a sufficient rolling reduction in the final rolling pass in the recrystallization temperature range to sufficiently promote recrystallization, in addition to the total rolling reduction in the recrystallization temperature range of 35% or more and 55% or less. When the rolling reduction in the final rolling pass in the recrystallization temperature range is less than 10%, recrystallization is insufficient, and the grains grow into coarse grains during the holding time between rough rolling and the start of finish rolling. Thus, the rolling reduction of the final rolling pass in the recrystallization temperature range is 10% or more, preferably 11% or more. The upper limit of the rolling reduction in the final rolling pass in the recrystallization temperature range is not limited, and a higher rolling reduction is more preferred. As an example, the rolling reduction in the final rolling pass in the recrystallization temperature range is 20% or less.

Even after the rolling in the recrystallization temperature range is completed, recrystallization occurs partially. Thus, recrystallization can be promoted by further increasing the rolling reduction in the temperature range of (the lower limit temperature in the recrystallization temperature range - 80°C) or higher and lower than the lower limit temperature in the recrystallization temperature range. This makes it possible to effectively refine the top 20% grain size in the microstructure at the center of the plate thickness. Thus, the rolling reduction in the final rolling pass in the temperature range of (the lower limit temperature in the recrystallization temperature range - 80°C) or higher and lower than the lower limit temperature in the recrystallization temperature range is 15% or more, preferably 16% or more. The upper limit of the rolling reduction in the final rolling pass in the above temperature range is not particularly limited, and a higher rolling reduction is more preferred. As an example, the rolling reduction in the final rolling pass in the above temperature range is 25% or less.

Rolling at a temperature lower than (the lower limit temperature in the recrystallization temperature range - 80°C) is effective in refining grains because more strain is introduced when the rolling is performed at a lower temperature. For this reason, rolling is preferably performed at a low temperature lower than (the lower limit temperature of the recrystallization temperature range - 80°C) within a range in which the cooling start temperature can be observed.

In the hot rolling step, a lower finish rolling temperature is more preferred in order to refine the crystal grains. From the viewpoint of ensuring the HISC resistance in a high-pressure hydrogen environment, the cooling start temperature in the cooling step after the hot rolling step needs to be higher than or equal to the Ar₃ transformation point in terms of a surface temperature of the steel plate. In consideration of this, the finish rolling temperature needs to be set. The term "Ar₃ transformation point" refers to the ferrite transformation start temperature during cooling, and can be calculated, for example, from the components of the steel using the following formula. The surface temperature of the steel plate can be measured using a radiation thermometer or the like. Ar₃ transformation point (°C) = 910 - 310[%C] - 80[%Mn] - 20[%Cu] - 15[%Cr] - 55[%Ni] - 80[%Mo] where in the above formula, each [%X] is the element X content (mass%) of the steel, and an element that is not contained is defined as 0.

After the hot rolling step, the steel plate is cooled (controlled cooling). When the surface temperature of the steel plate at the start of cooling is lower than the Ar₃ transformation point (°C), ferrite forms before cooling, resulting in a significant decrease in strength. For this reason, the surface temperature of the steel plate at the start of cooling is higher than or equal to the Ar₃ transformation point (°C). The surface temperature of the steel plate at the start of cooling is the temperature of the surface region of the steel plate where the cooling start temperature is the lowest. Specifically, the surface temperature of the steel plate at the start of cooling is the surface temperature of the steel plate at its trailing end portion when, for example, the steel plate is cooled while traveling in one direction with respect to a cooling device. For example, in the case where cooling is performed individually for each specific region of the entire steel plate and where different cooling start times are set for different regions, the surface temperature of the steel plate is the surface temperature of the region that is cooled last. As an example, the upper limit of the surface temperature of the steel plate at the start of cooling is the above-mentioned finish rolling temperature.

When the difference in cooling start time for the entirety of the steel plate is more than 50 seconds, the difference in the temperature of the steel plate is large, resulting in a large variation in the temperature of the steel plate when cooling is stopped. This increases the variation in Vickers hardness at 0.25 mm below the surface of the steel plate and deteriorates the HISC resistance. Therefore, the difference in cooling start time for the entirety of the steel plate is within 50 seconds, preferably within 45 seconds. Specifically, for example, when the steel plate is cooled while traveling in one direction with respect to the cooling device, the difference in cooling start time between the leading end of the steel plate and the trailing end of the steel plate is within 50 seconds. For example, in the case where cooling is performed individually for each specific region of the entire steel plate and where different cooling start times are set for different regions, the difference in cooling start time between the first region and the last region is within 50 seconds. When the entire steel plate can be cooled at once, the difference in cooling start time for the entire steel plate may be 0 seconds.

To achieve high strength while obtaining excellent HISC resistance, it is necessary to control the cooling rates at 0.25 mm below the surface of the steel plate and at the center of the plate thickness.

It is important to minimize the average cooling rate from 750°C to 550°C in terms of a steel plate temperature at 0.25 mm below the surface of the steel plate to thereby form granular bainite. The temperature range of 750°C to 550°C is an important temperature range for bainite transformation. It is thus important to control the cooling rate in this temperature range. When the average cooling rate in the temperature range is more than 50 °C/s, the hardness may vary, and the HISC resistance after pipe production deteriorates. Therefore, the average cooling rate is 50 °C/s or less, preferably 45 °C/s or less. When the cooling rate is too low, ferrite and pearlite are formed, resulting in insufficient strength. From the viewpoint of inhibiting this, the average cooling rate in the above-mentioned temperature range is 15 °C/s or more, preferably 17 °C/s or more. With regard to cooling in a temperature range of 550°C or lower in terms of the steel plate temperature at 0.25 mm below the surface of the steel plate, when the cooling rate is low, cooling in a stable nucleate boiling state may fail to be achieved, resulting in a variation in hardness in the extreme surface layer portion of the steel plate. Therefore, the average cooling rate in the temperature range of 550°C to the cooling stop temperature in terms of the steel plate temperature at 0.25 mm below the surface of the steel plate is preferably 150 °C/s or more. The average cooling rate is preferably 250 °C/s or less from the viewpoint of more easily inhibiting the variation in hardness.

When the average cooling rate from 750°C to 550°C at the center of the plate thickness is less than 15 °C/s, a granular bainite structure is not formed, resulting in a reduction in strength. Therefore, the average cooling rate from 750°C to 550°C at the center of the plate thickness is 15 °C/s or more. The average cooling rate is preferably 17 °C/s or more from the viewpoint of inhibiting the variation of the microstructure. The average cooling rate is 50 °C/s or less, preferably 45 °C/s or less, in order to inhibit a variation in grain size. The cooling in the temperature range of 550°C or lower in terms of the steel plate temperature at the center of the plate thickness is not particularly limited, but the average cooling rate in the above-mentioned temperature range is preferably 15 °C/s or more from the viewpoint of inhibiting the variations of the microstructure and grain size. From the above viewpoint, the average cooling rate in the above temperature range is preferably 50 °C/s or less.

The steel plate temperature at 0.25 mm below the surface of the steel plate and at the center of the plate thickness cannot be measured directly and physically. However, based on the surface temperature at the start of cooling measured by a radiation thermometer and a target surface temperature at the stop of cooling, the temperature distribution in the thickness section is calculated by difference calculation using, for example, a process computer, and the temperature distribution can be obtained in real time from the results. The temperature at 0.25 mm below the surface of the steel plate in the temperature distribution is defined as the "temperature of the steel plate at 0.25 mm below the surface of the steel plate" in the present specification. The temperature at the center of the plate thickness in the temperature distribution is defined as the "steel plate temperature at the center of the plate thickness" in the present specification.

When the cooling stop temperature is higher than 550°C in terms of the steel plate temperature at 0.25 mm below the surface of the steel plate and at the center of the plate thickness, bainite transformation is incomplete, resulting in a failure to obtain sufficient strength. Thus, the cooling stop temperature is 550°C or lower, preferably 500°C or lower. When the cooling stop temperature is lower than 250°C, the hardness increases to thereby deteriorate the HISC resistance. Thus, the cooling stop temperature is 250°C or higher, preferably 300°C or higher.

A steel pipe for transporting hydrogen (UOE steel pipe, electric resistance welded steel pipe, spiral steel pile, or the like) suitable for transporting high-pressure hydrogen gas can be manufactured by forming the high strength steel plate of the present invention into a tubular shape using press bend forming, roll forming, UOE forming, or the like, and then welding the butting portions. In addition, by producing a steel pipe using the high strength steel plate of the present invention, it is possible to produce a steel pipe having excellent HISC resistance even if a high hardness region is present in the weld. In the present invention, the high-pressure hydrogen refers to, for example, a hydrogen gas environment with a pressure of 15 MPa or more.

For example, a UOE steel pipe is manufactured by groove cutting the ends of a steel plate, forming the steel plate into a steel pipe shape by C press, U press, and O press, then seam welding the butting portions by inner surface welding and outer surface welding, and optionally subjecting it to an expansion process. Any welding method may be used as long as sufficient joint strength and joint toughness can be obtained, but it is preferable to use submerged arc welding from the viewpoint of excellent welding quality and production efficiency. A steel pipe produced by forming a steel plate into a pipe shape using press bend forming and then seam welding the butting portions may also be subjected to pipe expansion.

Steels having the chemical compositions given in Table 1 (steel types A to W) were formed into steel slabs (slabs) by a continuous casting method. The steel slabs were heated to the heating temperature given in Table 2 and then subjected to hot rolling and cooling under the conditions given in Table 2 to form steel plates having the final thickness given in Table 2. In the cooling step, each steel plate was subjected to controlled cooling using a water cooling-type controlled cooling device while the steel plate was traveling in one direction. Thereafter, the edges of the steel plate were subjected to groove cutting, and the steel plate was formed into a steel pipe shape by C press, U press, and O press. Subsequently, the butting portions of the inner surface and the outer surface were seam welded by submerged arc welding, and a pipe expansion process was performed to obtain a steel pipe. The Ar₃ transformation point and the lower limit temperature Tnr of the recrystallization temperature range in Table 1 were determined from the above formulae.

With respect to a section perpendicular to the rolling direction, the Vickers hardness (HV 0.5) was measured at 100 equally spaced points in the plate width direction from each of the leading end and the trailing end in the rolling direction of the steel plate in accordance with JIS Z 2244 (2009), the 100 equally spaced points being located at 0.25 mm below the surface of the steel plate. The average value and the standard deviation σ of the Vickers hardness values (HV 0.5) at 200 points in total were determined. The leading end of the steel plate in the rolling direction is a position 1 m downstream of the leading edge of the steel plate in the rolling direction. The trailing end of the steel plate in the rolling direction is a position 1 m upstream of the trailing edge of the steel plate in the rolling direction. The measurement was performed in a region excluding non-steady portions near the ends in the plate width direction. Table 3 presents the average value of the Vickers hardness at 0.25 mm below the surface of the steel plate + 3σ.

A sample for observing the metallic microstructure was taken from the central portion of the width of the steel plate obtained as described above. A section of this sample perpendicular to the plate width direction was mirror-polished and then etched with colloidal silica. Thereafter, crystal data was collected by an electron backscatter diffraction (EBSD) method in a field of view of 1 mm × 1 mm at the center of the plate thickness (measurement step: 0.8 µm). After the data collection, a boundary having an orientation difference of 15° or more was determined as a crystal grain boundary using OIM-Analysis (OIM Analysis software, manufactured by EDAX). The equivalent circular diameter was calculated as the crystal grain size from the area of each crystal grain. A frequency distribution table was prepared for all crystal grains measured, and the crystal grain size corresponding to 20% of the cumulative relative frequency from the largest crystal grain size calculated was defined as the "top 20% grain size". The measurement results are presented in Table 3.

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

A full-thickness test piece in the direction perpendicular to the rolling direction was used as a test piece for a tensile test, and the tensile test was performed in accordance with the provisions of JIS Z2241 (2011) to measure tensile strength and yield strength. Table 3 presents the results.

With regard to the HISC resistance, as illustrated in Fig. 1, a test piece (coupon) cut out from the obtained steel pipe was flattened, and then a test specimen of 3 mm × 10 mm × 50 mm was taken from the inner surface of the steel pipe. At this time, a test specimen containing both the weld and the base material was collected in addition to a test specimen containing only the base material without the weld. The inner surface to be tested was left intact without removing the scale in order to leave the state of the outermost layer. That is, the position 0.25 mm below the surface of the steel plate is included in the test specimen. A stress of 90% of the actual yield strength (0.5% YS) of each steel pipe was applied to the test specimen thus collected, and a four-point bending test was performed in a high-pressure hydrogen gas of 21 MPa. When no cracking was observed in both the test specimen of the base material alone without the weld and the test specimen including both the weld and the base material after the exposure for 720 hours, the HISC resistance was determined to be excellent (good) and marked with ∘. In addition, when a crack occurred in at least one of the test specimens, the test specimen was determined to be defective, and was marked with ×. Table 3 presents the results.

The target ranges of the present invention are as described below. As a high strength steel plate for a hydrogen transport steel pipe, the average value of the Vickers hardness at 0.25 mm below the surface of the plate + 3σ is 225 HV or less. The top 20% grain size in the microstructure at the center of the plate thickness is 30 µm or less. The fatigue crack growth rate is less than 2.0 × 10^-2 (mm/cycle) when the stress intensity factor range ΔK is 45 (MPa·m^1/2). The tensile strength is 535 MPa or more. In addition, no cracking is observed in the evaluation of the HISC resistance (four-point bending test). No. Steel type Final thickness (mm) Steel slab heating temperature (°C) Hot rolling step Cooling step Class Rolling in recrystallization temperature range Rolling reduction in final rolling pass in temperature range of (lower limit temperature in recrystallization temperature range - 80°C) or higher^*1 (%) Finish rolling temperature (°C) Surface temperature of steel plate at start of cooling^*2 (°C) Surface temperature of steel plate at start of cooling - Ar₃ transformation point (°C) Difference in cooling start time for entirety of steel plate (seconds) Average cooling rate from 750°C to 550°C Average cooling rate from 550°C to cooling stop temperature Cooling stop temperature Total rolling reduction (%) Rolling reduction in final rolling pass (%) 0.25 mm below surface of steel plate (°C/s) Center of plate thickness (°C/s) 0.25 mm below surface of steel plate (°C/s) Center of plate thickness (°C/s) 0.25 mm below surface of steel plate (°C) Center of plate thickness (°C) 1 A 35 1030 37 12 17 805 780 12 41 18 16 170 16 450 450 Example 2 B 30 1060 39 11 16 817 788 13 35 21 18 190 18 410 410 Example 3 C 18 1160 49 11 17 855 785 7 19 35 31 230 33 500 500 Example 4 D 27 1090 40 10 15 813 775 10 32 24 21 200 22 430 430 Example 5 E 39 1000 35 12 17 804 781 19 50 17 15 150 15 400 400 Example 6 F 21 1130 46 13 20 842 794 10 20 30 26 220 27 250 250 Example 7 G 12 1250 55 13 20 890 790 2 10 50 44 250 47 550 550 Example 8 H 24 1100 43 12 16 816 775 15 25 26 23 210 24 300 300 Example 9 I 15 1200 52 10 18 883 802 1 15 41 36 240 38 520 520 Example 10 J 27 1080 39 13 19 833 798 7 29 25 22 210 23 470 470 Comparative Example 11 K 18 1150 48 10 17 851 783 8 18 34 30 230 32 270 510 Comparative Example 12 L 30 1050 38 13 20 850 820 1 34 20 17 180 17 470 470 Comparative Example 13 M 15 1190 51 12 18 847 765 14 14 39 35 230 37 320 530 Comparative Example 14 N 12 1240 54 12 20 884 786 3 11 49 44 250 47 350 540 Comparative Example 15 O 21 1120 45 13 20 845 795 8 21 31 27 220 28 480 480 Comparative Example 16 P 18 1170 48 11 18 864 790 4 17 36 32 220 34 290 510 Comparative Example 17 Q 39 1010 36 15 20 780 760 3 45 17 15 160 15 510 510 Comparative Example 18 R 35 1040 37 14 20 811 787 13 40 19 17 170 17 530 530 Comparative Example 19 S 24 1110 42 12 18 842 799 4 24 25 22 200 23 490 490 Comparative Example 20 T 15 1210 52 12 19 861 775 5 16 40 35 240 37 300 510 Comparative Example 21 C 35 990 38 10 15 806 780 2 39 20 17 160 17 480 480 Comparative Example 22 F 27 1260 41 12 20 829 796 12 27 23 20 200 21 430 430 Comparative Example 23 E 39 1020 33 11 16 800 777 15 47 18 16 150 16 480 480 Comparative Example 24 A 35 1020 57 11 17 808 781 13 42 17 15 160 15 540 540 Comparative Example 25 A 30 1070 37 9 17 814 782 14 36 22 19 190 19 490 490 Comparative Example 26 B 35 1030 36 10 11 819 794 19 41 19 17 180 17 520 520 Comparative Example 27 I 30 1050 38 11 15 820 791 -10 35 20 17 180 17 480 480 Comparative Example 28 D 24 1090 44 11 15 817 778 13 53 27 24 210 25 280 280 Comparative Example 29 G 39 1010 35 12 18 821 799 11 34 11 12 160 12 490 490 Comparative Example 30 H 15 1200 53 11 18 848 767 7 15 57 51 240 53 330 520 Comparative Example 31 E 18 1160 50 11 19 845 770 8 19 34 30 230 32 240 240 Comparative Example 32 C 35 1020 38 10 16 820 793 15 40 20 17 170 17 560 560 Comparative Example 33 U 39 1020 36 10 16 767 745 2 48 17 15 150 15 330 330 Example 34 V 35 1030 35 10 15 765 742 3 42 18 16 170 16 340 340 Example 35 W 20 1100 44 12 18 840 790 13 21 32 28 230 32 450 450 Example Note 1: Underlined shows outside the scope of the present invention.
*1 Rolling reduction in the final rolling pass in the temperature range of (the lower limit temperature in the recrystallization temperature range - 80°C) or higher and lower than the lower limit temperature in the recrystallization temperature range.
*2 Temperature at the trailing end of the steel plate. No. Average value of Vickers hardness +3σ^*1 (HV) Top 20% grain size^*2 (µm) Fatigue crack growth rate (mm/cycle) Yield strength (MPa) Tensile strength (MPa) HISC resistance Class 1 210 26 1.6 × 10^-2 472 561 ○ Example 2 211 23 1.4 × 10^-2 485 578 ○ Example 3 220 19 1.3 × 10^-2 477 572 ○ Example 4 215 21 1.3 × 10^-2 482 585 ○ Example 5 208 30 1.9 × 10^-2 465 550 ○ Example 6 218 19 1.1 × 10^-2 487 582 ○ Example 7 225 17 1.0 × 10^-2 486 589 ○ Example 8 217 20 1.1 × 10^-2 470 569 ○ Example 9 222 18 1.3 × 10^-2 473 565 ○ Example 10 213 22 1.3 × 10^-2 422 508 ○ Comparative Example 11 234 19 1.4 × 10-2 501 583 × Comparative Example 12 212 24 1.4 × 10^-2 420 510 ○ Comparative Example 13 234 17 1.1 × 10^-2 511 595 × Comparative Example 14 233 18 1.4 × 10-2 484 591 × Comparative Example 15 219 19 1.1 × 10^-2 432 518 ○ Comparative Example 16 232 18 1.3 × 10^-2 503 590 × Comparative Example 17 206 32 2.0 × 10^-2 471 558 ○ Comparative Example 18 210 34 2.2 × 10^-2 479 575 ○ Comparative Example 19 217 21 1.1 × 10^-2 435 513 ○ Comparative Example 20 231 18 1.3 × 10^-2 508 593 × Comparative Example 21 207 23 1.1 × 10^-2 442 511 ○ Comparative Example 22 214 35 2.2 × 10^-2 475 569 ○ Comparative Example 23 210 32 2.1 × 10^-2 488 579 ○ Comparative Example 24 209 31 2.0 × 10^-2 482 580 ○ Comparative Example 25 212 32 2.1 × 10^-2 478 571 ○ Comparative Example 26 208 31 2.0 × 10^-2 473 555 ○ Comparative Example 27 204 23 1.4 × 10^-2 443 524 ○ Comparative Example 28 239 22 1.6 × 10^-2 497 575 × Comparative Example 29 211 30 1.9 × 10^-2 439 522 ○ Comparative Example 30 235 18 1.4 × 10^-2 503 590 × Comparative Example 31 237 19 1.4 × 10^-2 505 588 × Comparative Example 32 213 28 1.8 × 10^-2 445 523 ○ Comparative Example 33 224 19 1.7 × 10^-2 516 609 ○ Example 34 225 18 1.8 × 10^-2 523 613 ○ Example 35 215 25 1.5 × 10^-2 477 570 ○ Example Note 1: Underlined shows outside the scope of the present invention.
*1 Position 0.25 mm below the surface of the steel plate.
*2 Microstructure at the center of the plate thickness.

As presented in Table 2, No. 1 to No. 9 and No. 33 to No. 35 are Examples in which the chemical composition and production conditions satisfy the appropriate ranges of the present invention. As presented in Table 3, in each of No. 1 to No. 9 and No. 33 to No. 35, the average value of the Vickers hardness + 3σ was 225 HV or less at 0.25 mm below the surface of the steel plate as a high strength steel plate. The top 20% grain size in the microstructure at the center of the plate thickness was less than 30 µm. The fatigue crack growth rate was less than 2.0 × 10^-2 (mm/cycle) when the stress intensity factor range ΔK was 45 (MPa·m^1/2). The tensile strength was 535 MPa or more. Furthermore, the HISC resistance was also good.

In contrast, the chemical composition of the steel plate of each of No. 10 to No. 20 is outside the scope of the present invention. In each of No. 10, No. 12, No. 15, and No. 19, the solid solution strengthening was insufficient, resulting in insufficient strength. In each of No. 11, No. 13, No. 14, No. 16, and No. 20, the Vickers hardness at 0.25 mm below the surface of the steel plate was increased, resulting in poor HISC resistance. In each of No. 17 and No. 18, the grain growth was not sufficiently inhibited by precipitates, resulting in poor fatigue crack growth resistance.

No. 21 to No. 32 are Comparative Examples in which the chemical compositions are within the scope of the present invention, but the production conditions are outside the scope of the present invention. In No. 21, since the heating temperature of the steel slab (slab) was low, carbides were not sufficiently dissolved, resulting in low strength. In No. 22, since the heating temperature of the steel slab was high, the crystal grains were coarsened, resulting in a deterioration in fatigue crack growth resistance. In No. 23, since the total rolling reduction in the recrystallization temperature range was insufficient, coarse grains remained, resulting in a deterioration in fatigue crack growth resistance. In No. 24, the top 20% grain size in the microstructure at the center of the plate thickness was large due to excessive total rolling reduction in the recrystallization temperature range, resulting in a deterioration in the fatigue crack growth resistance. In No. 25, since the rolling reduction in the final rolling pass in the recrystallization temperature range was insufficient, coarse grains remained, resulting in a deterioration in fatigue crack growth resistance. In No. 26, since the rolling reduction in the final rolling pass in the temperature range of (the lower limit temperature in the recrystallization temperature range - 80°C) or higher and lower than the lower limit temperature in the recrystallization temperature range was insufficient, the top 20% grain size in the microstructure at the center of the plate thickness was large, resulting in a deterioration in fatigue crack growth resistance. In No. 27, since the cooling start temperature was low, ferrite was partially formed, resulting in low strength. In No. 28, since the difference in cooling start time was large for the entirety of the steel plate, the variation in Vickers hardness at 0.25 mm below the surface of the steel plate was large, resulting in a deterioration in HISC resistance. In No. 29, since the average cooling rate from 750°C to 550°C was low, ferrite is partially formed, resulting in low strength. In No. 30, since the average cooling rate from 750°C to 550°C was high, the variation in Vickers hardness at 0.25 mm below the surface of the steel plate was large, resulting in a deterioration in HISC resistance. In No. 31, since the cooling stop temperature was low, the variation in Vickers hardness at 0.25 mm below the surface of the steel plate was large, resulting in a deterioration in HISC resistance. In No. 32, since the cooling stop temperature was high, ferrite was partially formed, resulting in low strength.

According to the present invention, the high strength steel plate for a hydrogen transport steel pipe can be provided, the high strength steel plate having excellent HISC resistance and fatigue crack growth resistance in a high-pressure hydrogen environment.