EP4636111A1

EP4636111A1 - High-strength seamless steel pipe for high-pressure hydrogen vessel and manufacturing method therefor

Info

Publication number: EP4636111A1
Application number: EP24766964.1A
Authority: EP
Inventors: Naho INOUE; Vanadia Irisca Yussalla; Hiroshi Okano; Kenichiro Eguchi
Original assignee: JFE Steel Corp
Current assignee: JFE Steel Corp
Priority date: 2023-03-07
Filing date: 2024-02-27
Publication date: 2025-10-22
Also published as: AU2024231754A1; JP7697601B2; CN120787267A; WO2024185593A1; JPWO2024185593A1; EP4636111A4; KR20250141195A; CL2025002634A1

Abstract

An object of the present invention is to provide a high-strength seamless steel pipe for high-pressure hydrogen container that has excellent hydrogen embrittlement resistance, and a method for manufacturing the same.

A high-strength seamless steel pipe for high-pressure hydrogen container has a specific composition; includes a microstructure containing tempered martensite with an area fraction of 95% or more; is such that 50% by mass or more of the molybdenum contained in the steel is present in precipitates and
50% by mass or more of the molybdenum present in precipitates is present in precipitates having a diameter of 50 nm or less; and
has a tensile strength TS of 850 MPa or more.

Description

Technical Field

The present invention relates to a high-strength seamless steel pipe for high-pressure hydrogen container and to a method for manufacturing the same.

Background Art

Fuel cell vehicles that use hydrogen as fuel emit no carbon dioxide (CO₂) and also excel in energy efficiency. Thus, they are expected as vehicles that can solve CO₂ emission problems and energy problems. The proliferation of fuel cell vehicles requires strong and durable containers that can safely store high-pressure hydrogen at 35 MPa or more, particularly about 70 MPa or more, in hydrogen stations where hydrogen is supplied to fuel cell vehicles or on fuel cell vehicles running on hydrogen. The development of such containers is underway.
Vehicle-mounted pressure accumulators are required to be lightweight. To address this requirement, liners made of a light material, such as aluminum or resin, are coated with a carbon fiber-reinforced plastic (CFRP). For example, Patent Literature 1 describes an Al-Mg-Si alloy liner with excellent fatigue properties.
On the other hand, lightness in weight is less required for pressure accumulators used at hydrogen stations than vehicle-mounted pressure accumulators. Thus, there have been proposed pressure accumulators entirely formed of a steel material(for example, Patent Literature 2) and pressure accumulators in which a Cr-Mo steel liner is coated with carbon fibers or glass fibers (for example, Patent Literature 3).
On the other hand, low-alloy steels, such as Cr-Mo steel, are known to be embrittled by hydrogen. Thus, the materials for 35 MPa or higher pressure hydrogen accumulators are limited to aluminum alloys and SUS316 that are less susceptible to material quality deterioration by hydrogen.
However, stainless steels, such as SUS316, are low in strength. Due to this fact, containers for containing hydrogen at, for example, 70 MPa are to be extremely large in wall thickness and therefore have an increased weight. Thus, the size of storage containers is limited and the amount of hydrogen that can be stored in the container is reduced. In addition, the material costs are so increased that the economic efficiency is deteriorated.
A lot of research has been done to construct high-pressure hydrogen storage containers with less expensive low-alloy steels in place of austenitic stainless steels. Patent Literature 4 proposes a steel for high-pressure hydrogen environment in which fine V-Mo precipitates are used as sites for trapping diffusible hydrogen in the steel, thereby suppressing embrittlement caused by diffusible hydrogen.
Patent Literature 5 proposes a low-alloy high-strength steel with excellent resistance to embrittlement in a high-pressure hydrogen environment. In this steel, the tensile strength is controlled within a very narrow range of 900 to 950 MPa by tempering a Cr-Mo steel at a relatively high temperature in refining of the steel.
As described in these conventional techniques, low-alloy steels with enhanced hydrogen embrittlement resistance are widely studied for use as infrastructure materials in high-pressure hydrogen gas environments in order to achieve cost reduction. As used herein, a high-pressure hydrogen gas environment means an environment at a total pressure of 1 MPa or more and a hydrogen partial pressure of 1 MPa or more.

Citation List

Patent Literature

PTL 1: Japanese Unexamined Patent Application Publication No. 2009-024225
PTL 2: Japanese Unexamined Patent Application Publication No. 2010-037655
PTL 3: Japanese Unexamined Patent Application Publication No. 2009-293799
PTL 4: Japanese Unexamined Patent Application Publication No. 2009-074122
PTL 5: Japanese Unexamined Patent Application Publication No. 2009-046737

Summary of Invention

Technical Problem

As known in the art, low-alloy steels used for hydrogen pressure vessels or the like undergo hydrogen embrittlement to suffer a decrease in the tensile strength of the materials or to exhibit a significantly accelerated fatigue crack propagation rate. Enhancing the material strength allows pressure vessels, such as high-pressure hydrogen accumulators, to be designed with a small wall thickness. Thus, high-strength steel materials are increasingly needed for the purposes of increasing the storage capacity and reducing the product weight. When, on the other hand, a steel material is strengthened to 850 MPa or more, the hydrogen embrittlement phenomenon becomes noticeable and the steel may be fractured below the maximum tensile strength tested in the air.
The present invention has been made in consideration of the circumstances discussed above. It is therefore an object of the present invention to provide a high-strength seamless steel pipe for high-pressure hydrogen container that has excellent hydrogen embrittlement resistance, and a method for manufacturing the same.

Solution to Problem

Based on the fact that a desired high strength and hydrogen embrittlement resistance need to be satisfied concurrently in order to achieve the above objective high-strength seamless steel pipe for high-pressure hydrogen container (hereinafter, also written simply as the high-strength seamless steel pipe), the present inventors carried out extensive studies on various factors that would affect strength and hydrogen embrittlement resistance. As a result, the present inventors have found that the hydrogen embrittlement resistance of a high-strength steel pipe with a tensile strength TS of 850 MPa or more is significantly affected by the presence or absence of intergranular fracture. The present inventors have then found that fine Mo precipitates that are dispersed in the steel microstructure serve as powerful hydrogen trapping sites to reduce hydrogen accumulation at grain boundaries and to suppress the occurrence of intergranular fracture. Furthermore, the present inventors have found that the hydrogen embrittlement resistance is enhanced with increasing amount of Mo precipitates having a size of 50 nm or less.
Based on the above findings, the present inventors have come to the realization that a high-strength seamless steel pipe will be further enhanced in hydrogen embrittlement resistance by increasing the amount of molybdenum contained in precipitates and, in particular, controlling the amount of molybdenum contained in 50 nm and smaller precipitates to or above an appropriate level.

[1] A high-strength seamless steel pipe for high-pressure hydrogen container, wherein
the high-strength seamless steel pipe has a composition including, in mass%:
- C: 0.20 to 0.50%,
- Si: 0.05 to 2.00%,
- Mn: 0.30 to 1.50%,
- P: 0.015% or less,
- S: 0.005% or less,
- Al: 0.005 to 0.150%,
- N: 0.006% or less,
- Cr: more than 0.2% and 1.7% or less,
- Mo: more than 1.0% and 3.0% or less,
- Nb: 0.001 to 0.020%,
- B: 0.0003 to 0.0030%,
- O: 0.0030% or less, and
- Ti: 0.003 to 0.025%,
- the balance being Fe and incidental impurities,
- the ratio of the Mo content to the C content, Mo/C, being in a range of more than 2.0 to 12.0;
- the high-strength seamless steel pipe includes a microstructure containing tempered martensite with an area fraction of 95% or more;
- 50% by mass or more of the molybdenum contained in the steel is present in precipitates;
- 50% by mass or more of the molybdenum present in precipitates is present in precipitates having a diameter of 50 nm or less; and
- the high-strength seamless steel pipe has a tensile strength
- TS of 850 MPa or more.
[2] The high-strength seamless steel pipe for high-pressure hydrogen container according to [1], wherein the composition further includes, in mass%, one, or two or more selected from:
- V: 0.30% or less,
- Cu: 1.00% or less,
- Ni: 2.0% or less, and
- W: 3.0% or less.
[3] The high-strength seamless steel pipe for high-pressure hydrogen container according to [1] or [2], wherein the composition further includes, in mass%:
H: 0.0010% or less.
[4] The high-strength seamless steel pipe for high-pressure hydrogen container according to any one of [1] to [3], wherein the composition further includes, in mass%:
Ca: 0.0005 to 0.005%.
[5] A method for manufacturing the high-strength seamless steel pipe for high-pressure hydrogen container described in any one of [1] to [4], the method including:
- casting a steel pipe material having the composition described into a billet, followed by heating the billet at a temperature in a range of 1050 to 1350°C;
- hot rolling the billet to form a seamless steel pipe with a predetermined shape;
- cooling the seamless steel pipe from the hot rolling, at an average cooling rate equal to or faster than air cooling to such a temperature that the surface temperature reaches 200°C or below;
- reheating the steel pipe cooled, to a temperature equal to or higher than Ac₃ transformation temperature and equal to or lower than 1000°C;
- quenching the steel pipe at least one time to such a temperature that the surface temperature reaches 200°C or below; and
- tempering the steel pipe quenched, by heating the steel pipe to a tempering temperature of 600 to 740°C in such a manner that the average heating rate until the tempering temperature is reached is 0.5°C/min or more and the steel pipe is held at the tempering temperature for a holding time of 10 minutes or more and less than 60 minutes.
[6] The method for manufacturing the high-strength seamless steel pipe for high-pressure hydrogen container according to [5], wherein the casting speed in the casting is 1.8 m/min or less.

Advantageous Effects of Invention

According to the present invention, a high-strength seamless steel pipe for high-pressure hydrogen container that has a high tensile strength TS of 850 MPa or more and excellent hydrogen embrittlement resistance can be manufactured easily and inexpensively. The present invention thus achieves significant industrial effects. Furthermore, the manufacturing method involves appropriate amounts of appropriate alloying elements and promotes the formation of Mo precipitates. In this manner, the manufacturing method can stably produce high-strength seamless steel pipes that have a high strength desired for pressure vessels and also have excellent hydrogen embrittlement resistance.

Description of Embodiments

[Chemical composition]

First, the reasons will be described as to why the composition of the high-strength seamless steel pipe for high-pressure hydrogen container (hereinafter, also written simply as the high-strength seamless steel pipe) of the present invention is limited. In the following, mass% in the composition is simply written as %.

C: 0.20 to 0.50%

Carbon contributes to increasing the strength of steel by solid solution strengthening and enhances the hardenability of steel so as to contribute to the formation of a microstructure mainly composed of a martensite phase at the time of quenching. In order to obtain these effects, the C content needs to be 0.20% or more. The C content is preferably 0.22% or more, more preferably 0.25% or more, and still more preferably 0.28% or more. If, on the other hand, the C content is more than 0.50%, cracking occurs during quenching to significantly deteriorate the manufacturability. Thus, the C content is limited to 0.50% or less. The C content is preferably 0.45% or less, and more preferably 0.40% or less. The C content is still more preferably 0.35% or less.

Si: 0.05 to 2.00%

Silicon is added for deoxidization, but the deoxidization effect is insufficient if the content is less than 0.05%. Thus, the Si content is limited to 0.05% or more. The Si content is preferably 0.10% or more, more preferably 0.20% or more, and still more preferably 0.30% or more. If, on the other hand, the Si content is more than 2.00%, the effect is saturated. Thus, the Si content is limited to 2.00% or less. The Si content is preferably 1.00% or less, and more preferably 0.80% or less. More than 0.50% silicon deteriorates toughness and weldability. Thus, the Si content is more preferably 0.50% or less.

Mn: 0.30 to 1.50%

Similarly to carbon, manganese is an element that enhances the hardenability of steel and contributes to increasing the strength of steel. In order to obtain these effects, the Mn content is limited to 0.30% or more. The Mn content is preferably 0.40% or more, more preferably 0.45% or more, and still more preferably 0.50% or more. On the other hand, manganese is an element that segregates in steel to harden the steel locally. When added in a large amount, manganese forms localized hard regions to disadvantageously lower the hydrogen embrittlement resistance. Thus, the Mn content in the present invention is limited to 1.50% or less. The Mn content is preferably 1.20% or less, more preferably 1.00% or less, and still more preferably 0.80% or less.

P: 0.015% or less

Phosphorus is an element that segregates at grain boundaries in the steel microstructure to cause grain boundary embrittlement, and also segregates to harden the steel locally. In the present invention, phosphorus is an incidental impurity. While it is preferable to remove as much phosphorus as possible, up to 0.015% phosphorus is acceptable. Thus, the P content is limited to 0.015% or less. The P content is preferably 0.008% or less. The P content is more preferably 0.005% or less, and still more preferably 0.003% or less. While a lower content is more preferable, from the point of view of refining costs, the P content is preferably 0.0001% or more, and more preferably 0.001% or more.

S: 0.005% or less

Sulfur is an incidental impurity. In steel, most of sulfur is present as sulfide inclusions that reduce ductility, toughness, and SSC resistance. While it is preferable to remove as much sulfur as possible, up to 0.005% sulfur is acceptable. Thus, the S content is limited to 0.005% or less. The S content is preferably 0.003% or less, and more preferably 0.002% or less. While a lower content is more preferable, from the point of view of refining costs, the S content is preferably 0.0002% or more. The S content is more preferably 0.001% or more.

Al: 0.005 to 0.150%

Aluminum is added as a deoxidizing agent but produces no effects when the amount is less than 0.005%. Thus, the Al content is limited to 0.005% or more. The Al content is preferably 0.010% or more, and more preferably 0.020% or more. On the other hand, more than 0.150% aluminum lowers the cleanliness of steel and deteriorates toughness. Thus, the Al content is limited to 0.150% or less. The Al content is preferably 0.130% or less, more preferably 0.100% or less, and most preferably 0.080% or less.

N: 0.006% or less

In steel, nitrogen is present as an incidental impurity. Nitrogen has an effect of reducing the size of crystal grains to enhance toughness by combining with aluminum to form AlN, and, when titanium is present, by forming TiN. Thus, the N content is preferably 0.0005% or more, and more preferably 0.001% or more. If, however, the N content is more than 0.006%, coarse nitrides are formed to cause a significant decrease in toughness. Thus, the N content is limited to 0.006% or less. The N content is preferably 0.005% or less, more preferably 0.004% or less, and still more preferably 0.003% or less.

Cr: more than 0.2% and 1.7% or less

Chromium is an element that increases the strength of steel by enhancing hardenability and also enhances corrosion resistance. Furthermore, chromium is an element that combines with carbon during tempering to form precipitates, such as M₃C, M₇C₃, and M₂₃C₆ (M is the metal element), and enhances temper softening resistance. This element is necessary especially for increasing the strength of steel pipes. In particular, M₃C precipitates are highly effective in enhancing temper softening resistance. In order to obtain these effects, the Cr content is limited to more than 0.2%. The Cr content is preferably 0.3% or more, and more preferably 0.5% or more. On the other hand, more than 1.7% chromium forms large amounts of M₇C₃ and M₂₃C₆, which serve as hydrogen trapping sites to reduce hydrogen attack resistance. Furthermore, adding much chromium results in the coarsening of Mo precipitates. Fine Mo precipitates are coarsened by being aggregated and coalesced. Consequently, the number density of fine Mo precipitates is lowered to cause a decrease in hydrogen embrittlement resistance. Thus, the Cr content is limited to 1.7% or less. The Cr content is preferably 1.5% or less, more preferably 1.0% or less, and still more preferably 0.8% or less.

Mo: more than 1.0% and 3.0% or less

Molybdenum is an element that forms precipitates and contributes to increasing the strength of steel by precipitation strengthening, and effectively contributes to ensuring the desired high strength while reducing the dislocation density during tempering. Furthermore, molybdenum is dissolved in steel and segregates at prior austenite grain boundaries to contribute to enhancing hydrogen embrittlement resistance. Furthermore, molybdenum has effects of densifying corrosion products and suppressing the generation and growth of pits that serve as the origin of cracks. In order to obtain these effects, the Mo content is limited to more than 1.0%. The Mo content is preferably more than 1.1%, more preferably more than 1.2%, still more preferably 1.3% or more, and most preferably 1.4% or more. On the other hand, adding more than 3.0% molybdenum promotes the formation of acicular M₂C precipitates and, in some cases, Laves phase (Fe₂Mo), resulting in low hydrogen embrittlement resistance. Thus, the Mo content is limited to 3.0% or less. The Mo content is preferably 2.8% or less, more preferably 2.5% or less, still more preferably 1.8% or less, and most preferably 1.5% or less.

Nb: 0.001 to 0.020%

Niobium contributes to increasing the strength of steel through precipitation strengthening by forming precipitates or carbonitrides, and also contributes to reducing the size of austenite grains. In order to obtain these effects, the Nb content is limited to 0.001% or more. The Nb content is preferably 0.005% or more, more preferably 0.006% or more, and still more preferably 0.007% or more. On the other hand, coarse Nb precipitates tend to serve as the origin of hydrogen-induced cracking and the presence of a large amount of Nb precipitates resulting from the addition of more than 0.020% niobium leads to a significant decrease in hydrogen embrittlement resistance of high-strength steel materials. In order to achieve the desired high strength and excellent hydrogen embrittlement resistance at the same time, the Nb content in the present invention is limited to 0.020% or less. The Nb content is preferably 0.015% or less, and more preferably less than 0.010%.

B: 0.0003 to 0.0030%

Boron segregates at austenite grain boundaries and suppresses ferrite transformation from the grain boundaries. In this manner, boron, even contained in a small amount, enhances the hardenability of steel. In order to obtain these effects, the B content is limited to 0.0003% or more. The B content is preferably 0.0007% or more, and more preferably 0.0010% or more. If, on the other hand, the B content is more than 0.0030%, boron is precipitated as, for example, carbonitrides to cause a decrease in hardenability and hence a decrease in toughness. Thus, the B content is limited to 0.0030% or less. The B content is preferably 0.0025% or less. The B content is more preferably 0.0020% or less, and still more preferably 0.0015% or less.

O (oxygen): 0.0030% or less

Oxygen (O) is an incidental impurity and is present as oxide inclusions in steel. These inclusions serve as the origin of in a hydrogen gas environment and lower hydrogen embrittlement resistance. It is therefore preferable in the present invention that oxygen (O) be removed as much as possible. However, excessive deoxidization increases the refining costs. Up to 0.0030% oxygen (O) is acceptable. Thus, the O (oxygen) content is limited to 0.0030% or less. The O content is preferably 0.0025% or less. The O content is more preferably 0.0020% or less. The O content is still more preferably 0.0015% or less. Although the lower limit is not particularly limited, the O content is preferably 0.0010% or more.

Ti: 0.003 to 0.025%

At the time of the solidification of molten steel, titanium combines with nitrogen and is precipitated as fine TiN, which produces a pinning effect to contribute to reducing the size of austenite grains. In order to obtain this effect, 0.003% or more titanium needs to be added. The effect obtained by titanium is small if the content is less than 0.003%. Thus, the Ti content is limited to 0.003% or more. The Ti content is preferably 0.005% or more, and more preferably 0.010% or more. If, on the other hand, the Ti content is more than 0.025%, TiN is coarsened and fails to produce the pinning effect described above, and further the toughness is deteriorated. Furthermore, the hydrogen embrittlement resistance is lowered due to the coarse TiN. For these reasons, the Ti content is limited to 0.025% or less. The Ti content is preferably 0.020% or less, and more preferably 0.015% or less.

Mo/C: more than 2.0 to 12.0

If the ratio of the Mo content to the C content, Mo/C, is less than 2.0, the amount of Mo precipitates that are formed is reduced due to the lack of molybdenum, and consequently Mo precipitates enough to enhance hydrogen embrittlement resistance are not formed. Thus, Mo/C is limited to more than 2.0. Mo/C is preferably 2.5 or more, more preferably 3.0 or more, and still more preferably 3.5 or more. If, on the other hand, Mo/C is greater than 12.0, Mo precipitates highly tend to become coarse to deteriorate toughness and hydrogen embrittlement resistance. Because part of the coarse Mo precipitates results from the aggregation and coalescence of fine Mo precipitates, the number density of fine Mo precipitates is lowered. Thus, Mo/C is limited to 12.0 or less. Mo/C is preferably 10.0 or less, more preferably 8.0 or less, still more preferably 6.0 or less, and most preferably 5.0 or less.
The components described above are the basic components. In addition to the basic components, the composition may optionally include one, or two or more selected from V: 0.30% or less, Cu: 1.00% or less, Ni: 2.0% or less, and W: 3.0% or less; H: 0.0010% or less; Ca: 0.0005 to 0.005%; or any combination thereof.

One, or two or more selected from V: 0.30% or less, Cu: 1.00% or less, Ni: 2.0% or less, and W: 3.0% or less

Vanadium, copper, nickel, and tungsten are elements that contribute to increasing the strength of steel. One, or two or more may be selectively added as necessary.

V: 0.30% or less

Vanadium is an element that contributes to strengthening of steel by forming precipitates and carbonitrides. The V content may be 0% or more. In order to obtain the above effect, the V content is preferably 0.02% or more, and more preferably 0.03% or more. On the other hand, the effect is saturated even when more than 0.30% vanadium is added, and the addition will not produce the corresponding effect and is economically disadvantageous. Thus, when vanadium is added, the V content is limited to 0.30% or less. The V content is preferably 0.20% or less, and more preferably 0.15% or less.

Cu: 1.00% or less

Copper is an element that is effective for improving toughness and increasing strength. However, excessive addition deteriorates weldability. Thus, when copper is added, the Cu content is limited to 1.00% or less. The Cu content is preferably 0.75% or less, more preferably 0.50% or less, and still more preferably 0.25% or less. The Cu content may be 0% or more. In order to obtain the above effects, the Cu content is preferably 0.01% or more.

Ni: 2.0% or less

Nickel is an element that contributes to increasing the strength of steel and enhances toughness and corrosion resistance. In order to obtain these effects, the Ni content is desirably 0.03% or more. The Ni content is more preferably 0.1% or more. On the other hand, the effects are saturated even when more than 2.0% nickel is added, and the addition will not produce the corresponding effect and is economically disadvantageous. Thus, when nickel is added, the Ni content is limited to 2.0% or less. The Ni content is preferably 1.5% or less, more preferably 1.0% or less, and still more preferably 0.5% or less.

W: 3.0% or less

Tungsten is an element that contributes to increasing the strength of steel through precipitation strengthening by forming precipitates. Furthermore, tungsten is dissolved and segregates at prior austenite grain boundaries to contribute to enhancing hydrogen embrittlement resistance. In order to obtain these effects, the W content is desirably 0.03% or more. The W content is preferably 0.1% or more. On the other hand, the effects are saturated even when more than 3.0% tungsten is added, and the addition will not produce the corresponding effect and is economically disadvantageous. Thus, when tungsten is added, the W content is limited to 3.0% or less. The W content is preferably 2.5% or less. The W content is more preferably 2.0% or less, still more preferably 1.5% or less, and most preferably 1.0% or less.

H: 0.0010% or less

Hydrogen can be introduced into the steel material in various manufacturing steps. Heavy introduction increases the risk of cracking after solidification and deteriorates hydrogen embrittlement resistance. It is therefore important to reduce the amount of hydrogen in the steel material. The introduction does not cause problems as long as the H content is 0.0010% or less. Thus, when hydrogen is contained, the H content is limited to 0.0010% or less. The H content is preferably 0.0008% or less, more preferably 0.0005% or less, and still more preferably 0.0001% or less. Removing hydrogen to less than 0.00001% increases costs. Thus, the H content is preferably 0.00001% or more. The H content is more preferably 0.00005% or more. The amount of hydrogen is the amount of hydrogen remaining after the formation of sheet, steel pipe, or the like.

Ca: 0.0005 to 0.005%

Calcium is an element that combines with sulfur to form CaS and effectively acts in controlling the shape of sulfide inclusions. Through the control of the shape of sulfide inclusions, calcium contributes to enhancing toughness and hydrogen embrittlement resistance. When calcium is added in order to obtain these effects, the Ca content needs to be 0.0005% or more. The Ca content is preferably 0.001% or more. On the other hand, the effects are saturated even when more than 0.005% calcium is added, and the addition will not produce the corresponding effect and is economically disadvantageous. Thus, when calcium is added, the Ca content is limited to 0.005% or less. The Ca content is preferably 0.004% or less, more preferably 0.003% or less, and still more preferably 0.002% or less.
The balance after the components described above consists of Fe and incidental impurities. As incidental impurities, for example, Mg: 0.0008% or less and Co: 0.0008% or less are acceptable.

[Steel microstructure]

The high-strength seamless steel pipe of the present invention has the composition described hereinabove and further has a microstructure in which the main phase is tempered martensite. The microstructure contains precipitates, and the precipitates include precipitates having a diameter of 50 nm or less.

Main phase: tempered martensite phase

In the high-strength seamless steel pipe of the present invention, the microstructure is based on a martensite phase in order to ensure high strength with a tensile strength TS of 850 MPa or more. In order to maintain ductility and toughness required for structures, the main phase is a tempered martensite phase resulting from tempering of the martensite phase. As used herein, the term "main phase" indicates that the tempered martensite phase is the single phase representing 100% by area, or the microstructure includes 95% or more of the tempered martensite phase and a second phase with an area fraction of 5% or less without adversely affecting the characteristics. The tempered martensite phase preferably represents 97% or more, and more preferably represents 98% or more. As described above, the tempered martensite phase may represent 100%. The second phase in the present invention may be, for example, a bainite phase, a retained austenite phase, pearlite, or a mixture thereof.
The above-described microstructure of the high-strength seamless steel pipe for high-pressure hydrogen container of the present invention may be controlled by appropriately selecting the heating temperature during quenching and the cooling rate during cooling in accordance with the chemical composition of the steel.
If the grain size index of prior austenite grains is less than 8.5, the martensite phase that is formed contains coarse substructures and toughness is lowered. Thus, the grain size index of prior austenite grains is preferably 8.5 or more. The grain size index of prior austenite grains is preferably 9.0 or more, more preferably 9.6 or more, and still more preferably 10.0 or more. The grain size index of prior austenite grains is most preferably 12.0 or more. The upper limit is not particularly limited but is preferably small. For reasons, such as because the observation and appropriate evaluation are difficult, the grain size index is preferably 18.0 or less. The grain size index used herein is a value measured in accordance with the provisions of JIS G 0551.
In the present invention, the grain size index of prior austenite grains may be controlled by controlling the heating rate, the heating temperature, and the holding temperature during quenching, and also by changing the number of times the steel is quenched.
In the high-strength seamless steel pipe for high-pressure hydrogen container of the present invention, the concentration of Mo precipitates is controlled within an appropriate range depending on the size in order to attain enhanced hydrogen embrittlement resistance. The Mo precipitates were identified by the extraction method with filter filtration described in Patent Literature 6 and Reference 1. To measure the Mo concentration in precipitates, a 10 mm square sample taken from a cross section perpendicular to the rolling direction of the steel pipe (a cross section perpendicular to the pipe axis direction: C cross section) was electrolyzed in an electrolytic solution. Precipitates that had been attached to the steel surface were added to a dispersive liquid and ultrasonicated, thereby extracting the precipitates in the aqueous solution. The aqueous solution containing the extracted precipitates was filtered to classify the precipitates by size. The precipitates of the respective sizes were each dissolved into a solvent, and the Mo concentration was analyzed by ICP to calculate the Mo content in the precipitates of each size group. In ICP concentration analysis, a solution is introduced into plasma to obtain an element-specific spectrum, and the concentration of an element in the solution can be determined from the emission intensity of the light. The concentration (mass%) of molybdenum in the precipitates can be calculated in this manner. The content of molybdenum in all the precipitates can be calculated by the above technique. From the value obtained and the Mo content in the steel, the proportion (mass%) of molybdenum contained in the precipitates relative to the molybdenum contained in the steel can be determined. Furthermore, the electrolyzed solution was analyzed by ICP in accordance with Patent Literature 7 to determine the concentration of molybdenum dissolved in the steel. Furthermore, the content of molybdenum in the precipitates remaining on the filter is analyzed by ICP to determine the content of molybdenum in the precipitates larger than 50 nm. The content of molybdenum in the precipitates larger than 50 nm is subtracted from the content of molybdenum in all the precipitates to determine the proportion (mass%) of molybdenum contained in the precipitates having a diameter of 50 nm or less relative to the molybdenum contained in the precipitates.

[Patent Literature 6] Japanese Unexamined Patent Application Publication No. 2010-127791
[Patent Literature 7] Japanese Unexamined Patent Application Publication No. 2009-031269
[Reference 1] Ishida et al., Analysis of the Generation State of Fine Precipitates in Steel, Tetsu-to-Hagane Vol. 107, No. 08

50% by mass or more of the molybdenum contained in the steel is present in precipitates.

Molybdenum in the composition of the steel provides enhanced characteristics in a hydrogen environment when molybdenum is present as precipitates. If molybdenum is present as a solute, no such effects are expected even when the amount of molybdenum in the steel is increased. The hydrogen trapping ability was enhanced with increasing amount of Mo precipitates, and was enhanced when 50% or more of the molybdenum contained in the steel was present in precipitates. This fact shows that 50% by mass or more of the molybdenum contained in the steel needs to be present in precipitates. The molybdenum present in precipitates preferably represents 60% by mass or more of the molybdenum contained in the steel. The proportion is more preferably 65% by mass or more, and still more preferably 70% by mass or more. The upper limit is not particularly limited, but it is preferable that the molybdenum present in precipitates represent 95% by mass or less of the molybdenum contained in the steel. The proportion is more preferably 90% by mass or less.

50% by mass or more of the molybdenum present in precipitates is present in precipitates having a diameter of 50 nm or less.

Mo precipitates trap hydrogen in steel to inhibit the accumulation of hydrogen at grain boundaries, thereby enhancing grain boundary strength in a hydrogen environment. However, if their size is larger than 50 nm, the hydrogen trapping ability is lowered and is less effective in enhancing grain boundary strength. Thus, it is necessary that much molybdenum be contained in 50 nm and smaller precipitates. Here, the hydrogen trapping ability was enhanced with increasing amount of molybdenum contained in precipitates, and was enhanced when molybdenum present in fine precipitates having a diameter of 50 nm or less represented 50% by mass or more of the molybdenum present in precipitates. Thus, in the present invention, the proportion of molybdenum present in fine precipitates having a diameter of 50 nm or less is limited to 50% by mass or more of the molybdenum present in precipitates. The proportion of molybdenum present in fine precipitates having a diameter of 50 nm or less is preferably 60% by mass or more of the molybdenum present in precipitates. The proportion is more preferably 65% by mass or more, and still more preferably 70% by mass or more. While the upper limit is not particularly limited, the proportion of molybdenum present in fine precipitates having a diameter of 50 nm or less is preferably 95% by mass or less, more preferably 90% by mass or less, of the molybdenum present in precipitates. Because smaller Mo precipitates exhibit higher hydrogen trapping ability, the diameter of the precipitates is more preferably 20 nm or less. As already mentioned, coarse Mo precipitates result from the aggregation and coalescence of fine Mo precipitates. Thus, the coarsening leads to a decrease in the number of fine Mo precipitates. While the lower limit diameter of the precipitates of interest is not particularly limited, it is preferable that the precipitates of interest have a diameter of 1 nm or more.
Nitride inclusions and oxide inclusions can serve as the origin of fracture. In the present invention, reducing the amounts of these inclusions is also important in order to enhance hydrogen embrittlement resistance, in addition to causing Mo precipitates to occur. Of particular importance is management in the refining of molten steel. Pig iron is subjected to desulfurization and dephosphorization in pretreatment and is decarburized and dephosphorized in the converter. Ladle furnace (LF) stirred refining and RH vacuum degassing are performed in a ladle. A sufficient processing time is ensured for the ladle furnace (LF) stirred refining. In the RH vacuum degassing, a processing time is ensured and the RH circulation rate is controlled. When a billet (a steel pipe material) is formed by a continuous casting method, the steel being poured from the ladle into the tundish is sealed with an inert gas in order to reduce the amounts of nitride inclusions and oxide inclusions, and electromagnetic stirring is performed in the mold to cause the inclusions to float for separation. The refining process is not limited to the above. Appropriate management is important also in other refining processes.

[Manufacturing method]

Next, a method for manufacturing a high-strength seamless steel pipe for high-pressure hydrogen container according to the present invention will be described.
In the present invention, a steel pipe material having the composition described above is heated and hot-rolled to form a seamless steel pipe with a predetermined shape. The seamless steel pipe for high-pressure hydrogen container is preferably applied to a hydrogen container with a hydrogen pressure of 1 MPa or more, more preferably a hydrogen pressure of 20 MPa or more. The upper limit of the hydrogen pressure is not particularly limited, but the hydrogen pressure of interest is preferably 120 MPa or less.
The steel pipe material (hereinafter, also written simply as the steel material) used in the present invention is preferably a billet (a round billet) produced by smelting a molten steel having the above-described composition with a common smelting technique, such as a converter, and casting the steel by a common casting method, such as a continuous casting method. The billet may be further hot-rolled into a round bar having a predetermined shape, or a round bar may be produced by ingot making-blooming.
The following describes the manufacturing method assuming that the steel pipe is a seamless steel pipe as an example. It is needless to mention that an electric resistance welded pipe or a UOE steel pipe can be manufactured by processing the steel so that the steel will have a similar thermal history. For example, an electric resistance welded pipe having the similar properties may be produced by hot rolling a steel sheet at a temperature equal to or higher than the Ac₃ transformation temperature and equal to or lower than 1000°C, quenching the steel sheet at least one time to such a temperature that the surface temperature reaches 200°C or below, tempering the quenched steel sheet by heating the steel sheet to a temperature in the range of 600 to 740°C in such a manner that the average heating rate until the tempering temperature is reached is 0.5°C/min or more and the steel sheet is held at the tempering temperature for a holding time of 10 minutes or more and less than 60 minutes, and welding the steel sheet.
The steel pipe of the present invention may be manufactured by sequentially performing the following steps (1) to (3).

(1) A step in which the composition of the steel pipe material is adjusted and the steel pipe material is cast.
(2) A rolling step in which the billet (the cast material) is heated and rolled to give a steel pipe.
(3) A step in which the steel pipe obtained in the rolling step is cooled and tempered.

The steps will be individually described below. Unless otherwise specified, the temperature in the following description indicates the temperature on the surface of the billet or the steel pipe.

[Casting step]

Casting speed: 1.8 m/min or less

Excessively high casting speed increases the amount of inclusions and deteriorates hydrogen embrittlement resistance. Thus, the casting speed is preferably 1.8 m/min or less. The lower the casting speed, the greater the reduction in hydrogen concentration and in the amount of inclusions in the steel. This effect is more marked at the casting speed 1.0 m/min or less. Thus, the casting speed is more preferably 1.0 m/min or less, still more preferably 0.5 m/min or less, and most preferably 0.1 m/min or less. The lower limit is not particularly limited, but the casting speed is preferably 0.01 m/min or more to avoid difficult device control.

[Heating step]

The billet having the above-described chemical composition is heated for hot rolling. The billet is not particularly limited and may be, for example, a billet obtained by a usual continuous casting method.

Heating temperature: 1050 to 1350°C

If the heating temperature is below 1050°C, precipitates in the steel pipe material are not dissolved sufficiently. Thus, the heating temperature is limited to 1050°C or above. The heating temperature is preferably 1100°C or above, and more preferably 1150°C or above. On the other hand, heating at above 1350°C increases the grain size and also results in coarsening of precipitates, such as TiN, that occur during solidification, as well as coarsening of cementite, thereby lowering the toughness of the steel pipe. Furthermore, heating to a temperature above 1350°C forms a thick scale layer on the surface of the steel pipe material to cause problems, such as surface defects at the time of rolling, and also increases the energy loss and is disadvantageous from the point of view of energy saving. For these reasons, the heating temperature is limited to 1350°C or below. The heating temperature is preferably 1300°C or below. The heating temperature is more preferably 1250°C or below.

[Hot rolling step]

Next, the billet heated in the heating step is rolled to give a steel pipe with a predetermined shape. The rolling may be hot rolling including piercing by a usual Mannesmann-plug mill process or Mannesmann-mandrel mill process. The predetermined shape may be, for example, a hollow cylindrical steel pipe shape, with examples including tapered steel pipes with a smaller end diameter than the center diameter, and gas cylinder shapes represented by pressure vessels. For example, a hollow cylindrical steel pipe is preferably 200 to 600 mm in outer diameter and 500 to 12000 mm in steel pipe length in the pipe axis direction. A tapered steel pipe with a smaller end diameter than the center diameter is preferably 200 to 600 mm in outer diameter of the central portion, 50 to 550 mm in end diameter, and 500 to 12000 mm in steel pipe length in the pipe axis direction. A gas cylinder-shaped steel pipe is preferably 200 to 600 mm in outer diameter and 500 to 12000 mm in cylinder length in the pipe axis direction. The hot rolling also includes a process that performs a step of rolling a billet into a steel pipe shape (a hot working step) and an expansion step simultaneously. Where necessary, a sizing step for adjusting the wall thickness may be carried out after the reheating step described later.
After the completion of hot rolling, the seamless steel pipe obtained is subjected to a cooling treatment in which the seamless steel pipe is cooled at a cooling rate equal to or faster than air cooling until the surface temperature reaches 200°C or below.

Cooling treatment after the completion of hot rolling: average cooling rate: equal to or faster than air cooling, cooling stop temperature: 200°C or below

The chemical composition of the present invention can give a microstructure containing a martensite phase as the main phase when the steel pipe from the hot rolling is cooled at an average cooling rate equal to or faster than air cooling. If the air cooling (the cooling) is stopped when the surface temperature is still above 200°C, the transformation may not be fully completed. Thus, the cooling treatment after the hot rolling is to be performed at an average cooling rate equal to or faster than air cooling until the surface temperature reaches 200°C or below. In the present invention, the "cooling rate equal to or faster than air cooling" indicates 0.1°C/s or more. If the average cooling rate is less than 0.1°C/s, the metallic microstructure after the cooling becomes nonuniform, and the subsequent heat treatment gives a nonuniform metallic microstructure. The average cooling rate is preferably 1.0°C/s or more, and more preferably 10.0°C/s or more. Although the upper limit is not particularly limited, the average cooling rate is preferably 1000.0°C/s or less. The average cooling rate is the average of the cooling rates from Ac₃ transformation temperature to 200°C.

[Heat treatment step]

Reheating temperature for quenching: equal to or higher than Ac₃ transformation temperature and equal to or lower than 1000°C

If the reheating temperature for performing quenching is below Ac₃ transformation temperature, the steel is not heated to the austenite single phase region and consequently the microstructure that is obtained will not contain a martensite phase as the main phase. Thus, the reheating temperature is limited to equal to or higher than Ac₃ transformation temperature. The reheating temperature is preferably Ac₃ temperature + 30°C or above, and more preferably Ac₃ temperature + 50°C or above. However, the reheating temperature is not set to Ac₃ temperature + 30°C or above or to Ac₃ temperature + 50°C or above when "Ac₃ temperature + 30°C" or "Ac₃ temperature + 50°C" for the system exceeds 1000°C. On the other hand, reheating at above 1000°C is disadvantageous in that, for example, grains are coarsened to cause a decrease in toughness, and oxide scales on the surface are increased in thickness and easily come off to cause defects on the steel sheet surface. Furthermore, such reheating applies an excessively high load to the heat treatment furnace and is problematic from the point of view of energy saving. For these reasons and also from the point of view of energy saving, the reheating temperature for quenching is limited to 1000°C or below, and is preferably 980°C or below, and more preferably 950°C or below.
The reheated steel is quenched. The cooling in the quenching treatment is effected by rapidly cooling the steel until the surface temperature is lowered to 200°C or below. The term quenching or rapid cooling means that the average cooling rate from Ac₃ transformation temperature to 200°C is 2.0°C/s or more. The average cooling rate is preferably 5.0°C/s or more, and more preferably 10.0°C/s or more. The upper limit is not particularly limited, but the average cooling rate is preferably 1000.0°C/s or less. In addition to satisfying the above, water cooling preferably lowers the temperature at the center of the wall thickness from Ac₃ transformation temperature to 400°C or below at an average cooling rate of 2.0°C/s or more. The upper limit is not particularly limited, but this average cooling rate is preferably 1000.0°C/s or less. The cooling is preferably effected until the surface temperature is lowered to 100°C or below. The surface temperature after the cooling is preferably low and is preferably room temperature. The quenching treatment may be repeated two or more times.
The Ac₃ transformation temperature used here is a value calculated from the following expression. Ac₃ transformation temperature (°C) = 937 - 476.5C + 56Si - 19.7Mn - 16.3Cu - 4.9Cr - 26.6Ni + 38.1Mo + 124.8V + 136.3Ti + 198Al + 3315B (Here, C, Si, Mn, Cu, Cr, Ni, Mo, V, Ti, Al, and B indicate the contents (mass%) of the respective elements.)
When any of the elements described in the above expression is not contained in the steel, the content of that element is taken as 0% in the calculation of Ac₃ transformation temperature.
After cooled at a cooling rate equal to or faster than air cooling, the steel is tempered. The tempering treatment heats the steel to a temperature in the range of 600 to 740°C.

Tempering temperature: 600 to 740°C

Tempering is performed for the purposes of reducing the dislocation density and precipitating Mo precipitates, thereby enhancing toughness and hydrogen embrittlement resistance. The reduction in dislocation and the precipitation of Mo precipitates are insufficient at a tempering temperature of below 600°C, and the treatment fails to ensure excellent hydrogen embrittlement resistance. Thus, the tempering temperature is limited to 600°C or above. The tempering temperature is preferably 620°C or above, more preferably 640°C or above, and still more preferably 660°C or above. On the other hand, tempering at a temperature above 740°C significantly softens the microstructure and the desired high strength cannot be ensured. Thus, the tempering temperature is limited to 740°C or below. The tempering temperature is preferably 710°C or below. The tempering temperature is more preferably 700°C or below, and still more preferably 680°C or below.

Average heating rate until the tempering temperature is reached: 0.5°C/min or more

Mo precipitates are precipitated and grow during the heating process at the time of tempering. If the heating rate is slow until the predetermined temperature is reached in the tempering treatment, the size of precipitates is excessively increased and the desired hydrogen embrittlement resistance cannot be obtained. Thus, the average heating rate until the tempering temperature is reached is limited to 0.5°C/min or more, and is preferably 1.0°C/min or more, more preferably 2.0°C/min or more, and most preferably 5.0°C/min or more. There is no particular upper limit, but excessively rapid heating produces nonuniform temperature distribution and results in inhomogeneous material microstructure. Thus, the average heating rate is preferably 50.0°C/min or less.

Holding time at the tempering temperature: 10 minutes or more and less than 60 minutes

Mo precipitates occur most during temper holding. If this time is short, their precipitation is insufficient and the desired hydrogen embrittlement resistance cannot be obtained. The holding time at the tempering temperature is limited to 10 minutes or more. The holding time at the tempering temperature is preferably 15 minutes or more, and more preferably 20 minutes or more. If the holding time at the tempering temperature is too long, the size of precipitates is excessively increased. Thus, the holding time is limited to less than 60 minutes. Because holding time is a cost increasing factor in terms of energy, the tempering time is preferably less than 50 minutes, more preferably less than 40 minutes, and still more preferably less than 30 minutes.
In order to stably ensure the desired properties, the steel after being hot-rolled is subjected to the cooling treatment in which the steel is cooled at a cooling rate equal to or faster than air cooling, and is reheated and quenched at least one time by water cooling or the like, and is tempered in the manner described above. There is no particular upper limit to the number of quenching treatments, but it is preferable that the quenching be performed five times or less.
After the quenching and tempering treatments, a warm or cold straightening treatment may be performed as required to correct shape defects of the steel pipe.

EXAMPLES

Hereinbelow, Examples demonstrating the advantageous effects of the present invention will be described. The following description is illustrative of some preferred examples of the present invention, and the present invention is not limited to those Examples. In the following Examples, manufacturing methods and properties of seamless steel pipes for actual steel structures were examined. Table 1 describes the chemical compositions of steels Nos. 1 to 24. Table 2 describes the tempering conditions, the tempered martensite area fraction, the prior austenite grain size index, the proportion by mass% of molybdenum contained in precipitates relative to the molybdenum contained in the steel, the proportion of molybdenum contained in precipitates having a diameter of 50 nm or less relative to the molybdenum contained in precipitates, TS, and the relative reduction of area (RRA) for each of Nos. 1 to 24.

[Table 1]

Billets No.	Chemical composition (mass%)														Mo/C	Ac₃ transformation temperature (°C)
Billets No.	C	Si	Mn	P	S	Al	N	Cr	Mo	Nb	B	O	Ti	Others	Mo/C	Ac₃ transformation temperature (°C)
1	0.20	0.20	1.05	0.010	0.003	0.008	0.003	1.10	1.1	0.010	0.0010	0.0010	0.010	-	5.5	875.0
2	0.20	0.22	0.70	0.010	0.003	0.100	0.003	0.80	2.4	0.006	0.0003	0.0010	0.010	V 0.10, Cu 0.10, Ni 1.0, W 0.1	12.0	934.2
3	0.22	0.20	0.77	0.010	0.003	0.005	0.003	1.50	2.9	0.006	0.0010	0.0010	0.020	-	13.2	938.4
4	0.38	0.20	0.55	0.008	0.004	0.008	0.003	1.30	1.1	0.003	0.0020	0.0017	0.010	-	2.9	801.4
5	0.35	0.28	0.51	0.010	0.003	0.009	0.003	0.50	1.5	0.006	0.0010	0.0010	0.010	-	4.3	837.0
6	0.34	0.20	0.62	0.010	0.003	0.005	0.004	0.44	2.2	0.006	0.0012	0.0020	0.010	-	6.5	862.0
7	0.38	1.25	0.78	0.010	0.002	0.008	0.003	1.10	2.9	0.006	0.0010	0.0021	0.020	-	7.6	923.3
8	0.32	0.22	0.66	0.010	0.003	0.008	0.003	1.20	1.8	0.006	0.0010	0.0010	0.010	Ni 0.2, H 0.00100, Ca 0.0010	5.6	847.5
9	0.49	0.19	0.70	0.012	0.003	0.008	0.003	0.90	0.9	0.008	0.0010	0.0010	0.010	-	1.8	736.5
10	0.45	0.30	1.30	0.010	0.004	0.090	0.003	1.80	1.3	0.006	0.0015	0.0010	0.010	-	2.9	778.6
11	0.46	0.22	0.30	0.010	0.003	0.010	0.003	1.20	1.1	0.006	0.0010	0.0010	0.025	-	2.4	769.0
12	0.49	0.26	0.88	0.015	0.003	0.008	0.003	0.30	1.8	0.015	0.0010	0.0010	0.010	-	3.7	774.1
13	0.44	0.21	0.79	0.010	0.003	0.007	0.003	0.28	2.9	0.006	0.0005	0.0003	0.025	Ni 0.8	6.6	817.8
14	0.49	0.20	0.71	0.010	0.005	0.008	0.003	0.70	3.1	0.006	0.0010	0.0010	0.010	-	6.3	821.7
15	0.55	0.10	0.30	0.009	0.003	0.009	0.003	0.55	1.7	0.020	0.0022	0.0008	0.003	-	3.1	746.2
16	0.36	2.10	1.00	0.006	0.004	0.008	0.003	1.55	3.0	0.018	0.0010	0.0015	0.020	-	8.3	977.7
17	0.25	0.50	1.60	0.010	0.003	0.008	0.003	1.29	1.3	0.008	0.0005	0.0013	0.005	V 0.20, Ni 0.1, H 0.00080	5.2	883.8
18	0.29	0.85	0.40	0.010	0.003	0.008	0.003	0.99	1.9	0.030	0.0008	0.0022	0.015	-	6.6	912.4
19	0.34	1.05	1.40	0.010	0.002	0.120	0.003	1.37	2.3	0.011	0.0012	0.0005	0.030	Cu 0.05, W 1.0, Ca 0.0020	6.8	918.1
20	0.28	0.95	0.51	0.010	0.003	0.009	0.003	0.65	1.5	0.009	0.0010	0.0010	0.015	-	5.4	907.8
21	0.23	1.08	0.62	0.010	0.003	0.005	0.003	0.80	2.2	0.005	0.0012	0.0020	0.010	V 0.10, Cu 0.10, Ni 0.1	9.6	970.1
22	0.29	0.55	0.55	0.013	0.005	0.023	0.005	0.99	1.1	0.015	0.0005	0.0011	0.012	Ni 0.5	3.8	850.4
23	0.35	0.20	1.05	0.010	0.003	0.008	0.003	1.10	1.8	0.010	0.0010	0.0010	0.010		5.1	830.2
24	0.26	0.20	0.62	0.010	0.003	0.005	0.004	0.44	1.7	0.006	0.0012	0.0020	0.010		6.5	881.0

* The balance is Fe and incidental impurities.
* Underlines indicate being outside the range of the present invention.

[Table 2]

Steel pipes No.	Billets No.	Tempering step			Tempered martensite area fraction	Prior austenite grain size index	Proportion of molybdenum contained in precipitates relative to the molybdenum contained in the steel	Proportion of molybdenum contained in precipitates having a diameter of 50 nm or less relative to the molybdenum contained in precipitates	TS	Relative reduction of area (RRA)	Remarks
		Heating rate	Holding temp.	Holding time	Tempered martensite area fraction	Prior austenite grain size index			TS	Relative reduction of area (RRA)
		°C/min	°C	min	%		%	%	MPa	%
1	1	0.5	660	30	95	12.0	70	66	914	69	Inv. Ex.
2	2	0.6	700	40	98	13.0	80	63	897	69	Inv. Ex.
3	3	0.5	680	30	95	9.0	38	75	931	41	Comp. Ex.
4	4	0.4	630	15	99	10.0	68	22	869	32	Comp. Ex.
5	5	0.5	650	55	97	8.5	72	51	893	75	Inv. Ex.
6	6	0.5	720	30	99	10.5	66	70	941	64	Inv. Ex.
7	7	0.5	680	65	95	9.0	75	34	938	32	Comp. Ex.
8	8	0.7	620	55	97	11.0	88	75	927	73	Inv. Ex.
9	9	0.6	670	40	99	9.0	54	66	897	55	Comp. Ex.
10	10	0.5	610	20	97	8.5	71	41	869	52	Comp. Ex.
11	11	0.5	600	5	98	9.5	28	88	916	30	Comp. Ex.
12	12	0.5	650	10	98	10.5	89	66	872	71	Inv. Ex.
13	13	0.7	740	40	96	8.5	74	62	892	67	Inv. Ex.
14	14	0.5	620	45	99	10.0	79	39	925	48	Comp. Ex.
15	15	0.6	650	30	99	9.5	64	66	892	54	Comp. Ex.
16	16	0.5	630	35	97	9.0	68	54	925	49	Comp. Ex.
17	17	0.5	600	15	98	8.5	66	63	930	52	Comp. Ex.
18	18	0.7	640	30	98	11.5	59	68	905	57	Comp. Ex.
19	19	0.5	620	45	96	10.0	62	71	928	48	Comp. Ex.
20	20	0.6	590	50	98	9.5	44	59	962	40	Comp. Ex.
21	21	0.5	750	40	95	9.0	69	52	839	58	Comp. Ex.
22	22	0.6	700	20	94	12.5	52	81	898	62	Inv. Ex.
23	23	0.5	730	30	95	12.0	80	72	918	78	Inv. Ex.
24	24	0.5	720	30	99	10.5	66	69	863	75	Inv. Ex.

* Underlines indicate being outside the range of the present invention or that the target was not reached.

Billets having the chemical compositions described in Nos. 1 to 24 in Table 1 were produced at a casting speed of 0.6 m/min. The billets were heated to 1250°C, hot-worked, and expanded to give seamless steel pipes. The conditions in the production of the seamless steel pipes were such that the expansion was completed at 820°C or above, and the hot-worked pipes were cooled at a cooling rate equal to or faster than air cooling to such a temperature that the surface temperature was 200°C or below. Those steel pipes having an Ac₃ transformation temperature of 950°C or below were heated and held at 950°C, and those steel pipes having an Ac₃ transformation temperature of above 950°C were heated and held at 1000°C. The steel pipes were then water-cooled at 5.0°C/s to 200°C or below, and were subsequently tempered. The tempering step was performed at the heating rate, the holding temperature, and the holding time described in Table 2. The tempering temperature was controlled so that the tensile strength would be in the range of 850 to 950 MPa. The steel pipes obtained were analyzed to evaluate the metallic microstructure and mechanical properties.
Furthermore, steel pipes Nos. 25 to 39 in Table 3 were manufactured as follows. Any of the billets Nos. 5, 8, and 12 having the chemical compositions described in Table 1 were produced at various casting speeds. The billets were heated to 1250°C and expanded to give seamless steel pipes. The conditions in the steel pipe production were such that the expansion was completed at 820°C or above, and the hot-worked pipes were cooled at a cooling rate equal to or faster than air cooling to such a temperature that the surface temperature was 200°C or below. Those steel pipes having an Ac₃ transformation temperature of 950°C or below were heated and held at 950°C, and those steel pipes having an Ac₃ transformation temperature of above 950°C were heated and held at 1000°C. The steel pipes were then water-cooled at 5.0°C/s to 200°C or below, and were subsequently tempered under conditions described in Table 3. The steel pipes obtained were analyzed to evaluate the metallic microstructure and mechanical properties.
The evaluation methods are as follows.

Measurement of the martensite area fraction

The metallic microstructure at 1/4 wall thickness on the inner side of the steel pipe was evaluated as follows. Samples were obtained from a cross section parallel to the longitudinal direction and the wall thickness direction of the steel pipe, so that positions at 1/4 wall thickness on the inner side and at the center of the wall thickness would be the observation faces. The cross sections of the samples were etched with a 3 vol% Nital solution and were photographed with a scanning electron microscope at an appropriate magnification between 1000 and 5000 times. Tempered martensite, ferrite, bainite, and pearlite were observed. Tempered martensite was identified visually by comparing the image with the microstructure image shown in Reference 2. To determine the microstructure fraction, the SEM image was divided into regions based on the above identification, and the resultant image was binarized by image analysis into martensite and other regions. The fraction of tempered martensite was determined as the area fraction of tempered martensite.
[Reference 2] The Japan Society for Heat Treatment, Nyuumon · Kinzoku Zairyou no Soshiki to Seishitsu - Zairyou wo Ikasu Netsushori to Soshiki Seigyo (Introduction to the Microstructures and Properties of Metal Materials - Heat Treatment and Microstructure Control to Make the Most of Materials), 2004

Evaluation of prior austenite grain size

To study prior austenite (γ), a cross section of a test specimen for microstructure observation that was perpendicular to the longitudinal direction of the pipe (a C cross section) was polished and etched (Picral (picric acid-ethanol mixed solution)) to expose prior γ grain boundaries. The grain boundaries were observed with an optical microscope (magnification: 1000 times) and at least three fields of view were photographed. The microstructure images obtained were analyzed by a linear intercept method in accordance with the provisions of JIS G 0551 to determine the grain size index of prior γ grains. The average of the results was taken as the grain size index of prior γ grains in each steel pipe.

Measurement of Mo precipitates in steel materials

The steel material taken from the steel pipe was analyzed as follows to measure Mo precipitates. Mo precipitates were identified by an extraction method in which the steel material was electrolyzed and the resultant precipitates were filtered. A 10 mm square sample taken from a cross section perpendicular to the rolling direction of the steel pipe (a cross section perpendicular to the pipe axis direction: C cross section) was electrolyzed at a constant current in a 10% AA electrolytic solution to dissolve the steel. The residue was added to a 0.05 wt% aqueous sodium hexametaphosphate solution and was ultrasonicated. Precipitates were thus isolated. The dispersion was filtered through a 50 nm mesh filter, and thereby 50 nm and smaller precipitates were obtained. The 50 nm and smaller precipitates that had passed through the filter, and more than 50 nm precipitates remaining on the filter were each thermally treated with white fuming sulfuric acid, perchloric acid, and nitric acid, and were then dissolved into hydrochloric acid. The precipitate solutions and the electrolytic solution containing the dissolved portion were each analyzed by ICP to determine the Mo concentration (mass%) in the precipitates of each size group and the concentration (mass%) of dissolved molybdenum. The amounts of molybdenum contained in all the precipitates, and the amount of dissolved molybdenum obtained as described above were combined to give the total amount of molybdenum contained in the steel. The ratios "amount of molybdenum contained in all the precipitates/total amount of molybdenum" and "amount of molybdenum contained in 50 nm and smaller precipitates/amount of molybdenum contained in all the precipitates" were determined.

Evaluation of mechanical properties

A test piece for tensile test was taken from a cross section perpendicular to the steel pipe axis (C direction) in such a manner that a position at 1/4 wall thickness on the inner side of the steel pipe would be the center of the test piece and the longitudinal direction of the test piece would be the C direction. The test piece used here was a bar-shaped test piece specified in JIS Z 2201 "Tensile test pieces for metallic materials". The test was performed using the method specified in JIS Z2241, and the maximum load was taken as TS of the steel pipe.
While it is preferable that the sampling is centered at 1/4 wall thickness, the center of sampling may be other than at 1/4 wall thickness when the steel pipe has a small wall thickness (for example, a wall thickness of 45 mm or less).
Hydrogen embrittlement resistance was evaluated based on the relative reduction of area (RRA) of a test specimen after a slow strain rate tensile test in hydrogen gas in accordance with ASTM G 142. In air, a steel material undergoes plastic deformation and the fracture surface has a small area, and thus the reduction of area φ_air is large. In hydrogen, on the other hand, a steel material exhibits less elongation and is fractured before the area is reduced, and consequently the fracture surface has a large area. Thus, the reduction of area φ_H of the fracture surface after the test in hydrogen is small, unlike the testing in air. Hydrogen embrittlement resistance was evaluated based on how much the reduction of area had decreased. The relative reduction of area (RRA) is obtained from: $Relative reduction of area (RRA) = φ_{H} / φ_{air} \times 100$ The relative reduction of area obtained by a slow strain rate tensile test (stress rate: 0.002 mm/s) at room temperature under 105 MPa hydrogen gas is described in Table 2. The larger the RRA, the higher the hydrogen embrittlement resistance. In this evaluation, 60% or higher RRA was accepted. Incidentally, φ_air is the ratio "sectional area of the test specimen after testing in air/sectional area before testing", and φ_H is the ratio "sectional area of the test specimen after testing in hydrogen/sectional area before testing".

All of Inventive Examples satisfied 850 MPa or higher TS in the tensile test in air, and 60% or higher RRA in the slow strain rate tensile test in hydrogen gas.

[Table 3]

Steel pipes No.	Billets No.	Casting speed	Tempering step			Tempered martensite area fraction	Prior austenite grain size index	Proportion of molybdenum contained in precipitates relative to the molybdenum contained in the steel	Proportion of molybdenum contained in precipitates having a diameter of 50 nm or less relative to the molybdenum contained in precipitates	TS	Relative reduction of area (RRA)	Remarks
		Casting speed	Heating rate	Holding temp.	Holding time	Tempered martensite area fraction	Prior austenite grain size index			TS	Relative reduction of area (RRA)
		(m/min)	°C /min	°C	min	%		%	%	MPa	%
25	5	0.80	0.5	650	55	97	8.5	72	59	892	75	Inv. Ex.
26	5	1.20	0.5	650	55	97	8.5	73	60	893	74	Inv. Ex.
27	5	1.50	0.5	650	55	97	8.5	72	58	893	72	Inv. Ex.
28	5	1.80	0.5	650	55	97	8.5	71	58	892	72	Inv. Ex.
29	5	2.00	0.5	650	55	97	8.5	73	59	894	68	Inv. Ex.
30	8	0.80	0.7	620	55	97	11.0	88	75	926	72	Inv. Ex.
31	8	1.20	0.7	620	55	97	11.0	88	74	926	71	Inv. Ex.
32	8	1.50	0.7	620	55	97	11.0	86	74	928	71	Inv. Ex.
33	8	1.80	0.7	620	55	97	11.0	87	75	927	70	Inv. Ex.
34	8	2.00	0.7	620	55	97	11.0	86	75	926	67	Inv. Ex.
35	12	0.80	0.5	650	10	98	10.5	88	66	870	70	Inv. Ex.
36	12	1.20	0.5	650	10	98	10.5	89	66	873	70	Inv. Ex.
37	12	1.50	0.5	650	10	98	10.5	90	67	872	69	Inv. Ex.
38	12	1.80	0.5	650	10	98	10.5	88	65	871	68	Inv. Ex.
39	12	2.00	0.5	650	10	98	10.5	89	65	873	66	Inv. Ex.

Claims

A high-strength seamless steel pipe for high-pressure hydrogen container, wherein
the high-strength seamless steel pipe has a composition comprising, in mass%:
C: 0.20 to 0.50%,

Si: 0.05 to 2.00%,

Mn: 0.30 to 1.50%,

P: 0.015% or less,

S: 0.005% or less,

Al: 0.005 to 0.150%,

N: 0.006% or less,

Cr: more than 0.2% and 1.7% or less,

Mo: more than 1.0% and 3.0% or less,

Nb: 0.001 to 0.020%,

B: 0.0003 to 0.0030%,

O: 0.0030% or less, and

Ti: 0.003 to 0.025%,

the balance being Fe and incidental impurities,

the ratio of the Mo content to the C content, Mo/C, being in a range of more than 2.0 to 12.0;

the high-strength seamless steel pipe comprises a microstructure containing tempered martensite with an area fraction of 95% or more;

50% by mass or more of the molybdenum contained in the steel is present in precipitates;

50% by mass or more of the molybdenum present in precipitates is present in precipitates having a diameter of 50 nm or less; and

the high-strength seamless steel pipe has a tensile strength TS of 850 MPa or more.
The high-strength seamless steel pipe for high-pressure hydrogen container according to claim 1, wherein the composition further comprises, in mass%, one, or two or more selected from:
V: 0.30% or less,

Cu: 1.00% or less,

Ni: 2.0% or less, and

W: 3.0% or less.
The high-strength seamless steel pipe for high-pressure hydrogen container according to claim 1 or 2, wherein the composition further comprises, in mass%:
H: 0.0010% or less.
The high-strength seamless steel pipe for high-pressure hydrogen container according to any one of claims 1 to 3, wherein the composition further comprises, in mass%:
Ca: 0.0005 to 0.005%.
A method for manufacturing the high-strength seamless steel pipe for high-pressure hydrogen container described in any one of claims 1 to 4, the method comprising:
casting a steel pipe material having the composition described into a billet, followed by heating the billet at a temperature in a range of 1050 to 1350°C;

hot rolling the billet to form a seamless steel pipe with a predetermined shape;

cooling the seamless steel pipe from the hot rolling, at an average cooling rate equal to or faster than air cooling to such a temperature that the surface temperature reaches 200°C or below;

reheating the steel pipe cooled, to a temperature equal to or higher than Ac₃ transformation temperature and equal to or lower than 1000°C;

quenching the steel pipe at least one time to such a temperature that the surface temperature reaches 200°C or below; and

tempering the steel pipe quenched, by heating the steel pipe to a tempering temperature of 600 to 740°C in such a manner that:
the average heating rate until the tempering temperature is reached is 0.5°C/min or more and the steel pipe is held at the tempering temperature for a holding time of 10 minutes or more and less than 60 minutes.
The method for manufacturing the high-strength seamless steel pipe for high-pressure hydrogen container according to claim 5, wherein the casting speed in the casting is 1.8 m/min or less.