High-performance-uniformity subsea pipeline steel and manufacturing method therefor

Abstract

A high-performance-uniformity subsea pipeline steel and a manufacturing method therefor. The chemical components of the high-performance-uniformity subsea pipeline steel comprise, in percent by weight: C: 0.02-0.06%, Si: 0.10-0.30%, Mn: 1.20-1.65%, Cu≤0.30%, Ni≤0.30%, Cr: 0.15-0.30%, Mo≤0.15%, Nb: 0.03-0.05%, V≤0.05%, Ti: 0.008-0.020%, Ca: 0.0012-0.0030%, and Al: 0.015-0.045%, with the balance comprising Fe and inevitable impurities, where it is required to control impurity elements P≤0.01%, S≤0.001%, B≤0.0004%, N≤0.005%, O≤0.0025% and H≤0.0002%. The manufacturing method uses an austenite recrystallization zone and non-recrystallization zone two-stage controlled rolling technology, spread ratio control, and controlled cooling, such that in addition to high performance, obtained subsea pipeline steel also has good performance uniformity.

Description

High-performance uniform submarine pipeline steel and manufacturing method thereof Technical Field

[0001] The present invention belongs to the field of pipeline steel manufacturing, and in particular relates to high-performance uniform submarine pipeline steel and a manufacturing method thereof. Background Art

[0002] Submarine pipeline construction has a long history internationally. Since the first submarine pipeline was laid by Brown & Root Offshore Engineering in the Gulf of Mexico in 1954, numerous submarine pipelines of various types and diameters have been successfully laid in coastal waters around the world, including the North Sea, the Gulf of Mexico, the Mediterranean Sea, Australia, and the South China Sea. In recent years, with the gradual depletion of onshore oil and gas resources, oil and gas extraction has expanded to areas with harsh geological conditions, such as the oceans and polar regions, creating a stronger demand for high-performance submarine pipeline steel.

[0003] During installation, submarine pipelines are subjected to significant compression and tension loads, even bending deformation. During service, harsh operating conditions such as deepwater high pressure, swells, strong undercurrents, and seabed movement can cause significant plastic deformation and even fracture, potentially causing serious environmental damage. These complex operating conditions place even more stringent demands on deepwater pipelines than onshore pipelines, and as water depth increases, the requirements for raw materials continue to rise.

[0004] Submarine pipeline construction is characterized by high risk and high investment. Compared to onshore pipeline construction, it places strict demands on construction quality, environment, and technology. To mitigate the risks of difficult repairs and environmental pollution associated with failure of submarine pipelines, stricter technical requirements are often imposed to improve pipeline quality. Process management is also strengthened to enhance product reliability. Consequently, requirements are placed on the isotropy of steel plates (a more uniform strength and toughness distribution along the length and thickness), as well as excellent weldability, low-temperature fracture toughness, and strain aging performance. Furthermore, because submarine pipelines are subject to significant axial forces and displacements during installation and service, controlling longitudinal properties can ensure their load-bearing and strain capacity. Therefore, longitudinal tensile properties, particularly a low longitudinal yield ratio, are also required.

[0005] Chinese patent application CN201210123639.X discloses a "heavy-wall submarine pipeline steel plate and its production method." The steel plate is composed of the following components by weight: C: 0.05-0.07%, Si: 0.15-0.25%, Mn: 1.42-1.48%, P ≤ 0.010%, S ≤ 0.002%, Ni: 0.13-0.18%, Nb: 0.043-0.048%, Al: 0.020-0.040%, Ti: 0.014-0.024%, Mo: 0.13-0.18%, with the balance being Fe and unavoidable impurities. The plate exhibits high low-temperature impact toughness, good tear resistance, high strength, a moderate yield-to-tensile ratio, and good elongation. However, the pipeline steel in this patent can only meet the toughness requirements at -20℃, but cannot meet the toughness requirements at extremely low temperatures. It is also difficult to meet the tensile performance requirements after aging, especially the low yield ratio requirements. Most importantly, the realization of performance uniformity is not considered.

[0006] Chinese patent application CN201310457780.8 discloses a "low-cost, thick-gauge submarine pipeline steel plate and its manufacturing method." The steel plate's composition, by weight, is as follows: C: 0.04-0.06%, Si: 0.15-0.20%, Mn: 1.43-1.47%, P ≤ 0.008%, S ≤ 0.003%, Ni: 0.10-0.15%, Nb: 0.030-0.040%, Ti: 0.012-0.023%, Alt: 0.015-0.025%, N ≤ 0.006%, Nb+V+Ti ≤ 0.15%, with the remainder being iron and unavoidable impurities. The resulting steel plate exhibits a dual-phase microstructure, achieving high strength and a low yield ratio, making it suitable for submarine pipelines requiring high deformation resistance while also exhibiting good toughness. However, the pipeline steel in this patent can only meet the low-temperature toughness requirement of -30°C, cannot meet the toughness requirement at extremely low temperatures, and is difficult to meet the tensile performance requirements after aging, especially the low yield ratio requirement, and this patent does not consider the realization of performance uniformity.

[0007] Chinese patent application CN201210492187.2 discloses “thick-walled high DWTT performance X65-70 submarine pipeline steel and manufacturing method”, whose composition by weight percentage is: C 0.03-0.05%, Si≤0.25%, Mn 1.47-1.70%, P≤0.010%, S≤0.001%, Ti 0.006-0.010%, Cr 0.10-0.20%, Cu 0.12-0.20%, Ni 0.36-0.45%, Al 0.025-0.045%, Ca 0.0008-0.0025%, N≤0.0035%, O≤0.0025%, Nb 0.040~0.050%, the rest are Fe and inevitable impurities; and, Ceq=0.34~0.040, Pcm=0.13~0.17; among them, Ceq=C+Mn / 6+(Cr+Mo+V) / 5+(Ni+Cu) / 15; Pcm=C+Si / 30+Cu / 20+Ni / 60+Cr / 20+Mo / 15+V / 10, the rest are iron and inevitable impurities. The steel plate has thick specifications, high strength, excellent low-temperature impact toughness and weldability and good DWTT performance. However, the pipeline steel in this patent can only meet the -20℃ low-temperature toughness requirements, cannot meet the low yield ratio requirements and the tensile performance requirements after aging, and does not consider the realization of performance uniformity.

[0008] Chinese patent application CN201410515400.6 discloses "an extra-thick X70 pipeline steel and its manufacturing method." Its chemical composition, by mass percentage, is C: 0.03-0.06%, Si: 0.1-0.30%, Mn: 1.30-1.60%, P ≤ 0.010%, S ≤ 0.0050%, Nb: 0.030-0.050%, Ti: 0.008-0.020%, Ni: 0.10-0.30%, Cr: 0.10-0.30%, Cu: 0.10-0.30%, with the balance being Fe and unavoidable impurities. The resulting pipeline steel features thick thickness, high strength within a narrow strength range, excellent low-temperature toughness, and both low-temperature aging resistance and HIC resistance. However, the pipeline steel described in this patent only meets the -30°C low-temperature toughness requirement, failing to meet the toughness requirements at extremely low temperatures, and fails to address the achievement of uniform performance.

[0009] Chinese patent application CN201510383013.6 discloses a "Method for Producing Low-Yield Ratio, Acid-Resistant Submarine Pipeline Steel." The chemical composition is as follows: C: 0.03-0.05%, Si: 0.10-0.30%, Mn: 1.15-1.35%, P ≤ 0.010%, S ≤ 0.0010%, Alt: 0.01-0.06%, Nb: 0.04-0.07%, Ti: 0.008-0.020%, Ni: 0.10-0.25%, Cr: 0.10-0.30%, with the balance being Fe and unavoidable impurities. This patent utilizes advanced smelting and continuous casting technology to control the narrow fluctuation range of carbon equivalent and the number, size, and morphology of inclusions. Furthermore, a relaxed phase transformation process is used to control the microstructure of the steel plate, and a novel water-cooling control mode is employed to improve cooling uniformity. This effectively ensures the steel plate's low yield ratio and excellent acid resistance. However, the patent does not consider the realization of performance uniformity of pipeline steel.

[0010] Chinese patent application CN202011039510.1 discloses "a submarine pipeline steel coil with excellent low-temperature CTOD performance and its production process." Its chemical composition and mass percentages are as follows: C: 0.030-0.070%, Si: 0.10-0.20%, Mn: 1.10-1.60%, Nb: 0.020-0.050%, V: 0.020-0.040%, Ti: 0.010-0.030%, Mo: 0.10-0.20%, Alt ≤ 0.06%, S ≤ 0.0080%, P ≤ 0.018%, with the remainder being Fe and unavoidable impurities. This gives the submarine pipeline steel coil excellent CTOD performance, but it fails to consider the uniformity of the pipeline steel's performance.

[0011] Existing patent research shows that although existing submarine pipeline steels take strength and low-temperature toughness into consideration, not much research has been done on how to ensure that pipeline steels have good performance uniformity on the basis of high performance.

[0012] Summary of the Invention

[0013] The object of the present invention is to provide a high-performance uniform submarine pipeline steel and a manufacturing method thereof, wherein the pipeline steel has high strength and good low-temperature toughness, while ensuring the uniformity of the pipeline steel performance.

[0014] To achieve the above object, the technical solution of the present invention is:

[0015] A high-performance uniform submarine pipeline steel having the following chemical composition by weight: C: 0.02-0.06%, Si: 0.10-0.30%, Mn: 1.20-1.65%, Cu≤0.30%, Ni≤0.30%, Cr: 0.15-0.30%, Mo≤0.15%, Nb: 0.03-0.05%, V≤0.05%, Ti: 0.008-0.020%, Ca: 0.0012-0.0030%, Al: 0.015-0.045%, the balance including Fe and other unavoidable impurities, impurity elements P≤0.01%, S≤0.001%, B≤0.0004%, N≤0.005%, O≤0.0025%, H≤0.0002%; and simultaneously meeting the following requirements:

[0016] Ti / N≥2;

[0017] (Nb / 7.7C+V / 4.3C)≥0.10;

[0018] 1.5≤Ca / S≤4;

[0019] 0.30≤Ceq≤0.39, Ceq=C+Mn / 6+(Cr+V+Mo) / 5+(Cu+Ni) / 15;

[0020] Welding cold crack sensitivity index Pcm≤0.17, Pcm=C+Si / 30+(Mn+Cu+Cr) / 20+Ni / 60+Mo / 15+V / 10+5B.

[0021] Furthermore, the balance is Fe and other inevitable impurities.

[0022] The microstructure of the submarine pipeline steel of the present invention is ferrite+bainite, wherein the volume content percentage of the ferrite is 20-40%.

[0023] The grain size grade of the submarine pipeline steel of the present invention is ≥ grade 9.

[0024] Transverse tensile properties of submarine pipeline steel according to the present invention: yield strength R t0.5 : 415~555MPa, tensile strength R m : 515~670MPa, yield strength ratio R t0.5 / R m ≤0.88; the strength difference between the head and tail of the steel plate is ≤50MPa;

[0025] Longitudinal tensile strength: yield strength R t0.5 : 405~530MPa, tensile strength R m : 505~650MPa, yield strength ratio R t0.5 / Rm≤0.86; the strength difference between the head and tail of the steel plate is ≤50MPa;

[0026] After 3% pre-strain and 250℃ heat preservation for 1 hour, the strain aging test of the steel plate shows the following longitudinal tensile properties: yield strength R t0.5 : 405~530MPa, tensile strength R m : 505~650MPa, yield strength ratio R t0.5 / Rm≤0.88.

[0027] The submarine pipeline steel of the present invention has a Charpy impact energy of ≥250J at -45°C, and a difference in Charpy impact energy between the center and near-surface thickness of the steel plate of ≤50J; a DWTT test SA of ≥85% at -30°C, and a difference in SA between a single-side thinned DWTT specimen and a full-thickness DWTT specimen of ≤5%; and a hardness of ≤230HV10, wherein the difference in hardness between the surface hardness of the steel plate and the center hardness of the steel plate is ≤25HV10, and preferably, the difference in hardness between the surface hardness of the steel plate and the center hardness of the steel plate is ≤15HV10.

[0028] The characteristics of the submarine pipeline steel composition design of the present invention are:

[0029] 1. Adopt low C and low Mn alloy composition system, control 0.30≤Ceq≤0.39, Pcm≤0.17, ensure good welding performance under the premise that the strength and toughness of the material meet the requirements, and at the same time improve the center segregation and banded structure to ensure the isotropy of material properties.

[0030] 2. Impurity elements P, S, B, N, O, and H are controlled at extremely low levels to reduce the number of inclusions, lower the degree of segregation, and ensure strength and toughness, which is especially beneficial to ensure the excellent low-temperature toughness of pipeline steel.

[0031] 3. Adding alloying elements Cu, Ni, Cr, and Mo plays a role in solid solution strengthening, while affecting the phase transformation behavior of the material during cooling and improving the isotropy of the steel plate performance.

[0032] 4. Add trace amounts of V-Nb-Ti, while strictly limiting Ti / N ≥ 2, (Nb / 7.7C+V / 4.3C) ≥ 0.10, to form microalloy carbonitrides, which refine the grains, strengthen precipitation, and fix C and N atoms, thereby reducing the effect of strain aging on material properties.

[0033] Carbon (C): A fundamental strengthening element, carbon dissolves in steel to form interstitial solid solutions, contributing to solid solution strengthening. It also precipitates carbides with strong carbide-forming elements, contributing to precipitation strengthening. However, increasing C content significantly reduces the steel's ductility, toughness, and weldability. It also causes central segregation in the continuously cast ingot, hindering the uniformity of the steel's properties through the thickness. Furthermore, after strain aging, free C atoms can become entangled with dislocations, hindering their motion and increasing yield strength, hindering the maintenance of a low yield strength ratio. Therefore, the present invention limits the C content to 0.02-0.06%.

[0034] Mn: A solid solution strengthening element, it is the most important and economical strengthening element in steel to compensate for the strength loss caused by reduced carbon content. Mn also expands the γ phase region, which can improve the hardenability of steel and lower the γ→α phase transformation temperature, helping to produce fine phase transformation products and improving the toughness of steel. However, Mn is also a segregating element. Excessive Mn content can easily cause center segregation in the continuous casting billet, easily forming a hard phase structure during the post-rolling cooling process, reducing the material's low-temperature toughness and detrimentally affecting the uniformity of the steel plate's properties through the thickness. Therefore, in the present invention, the Mn content is controlled to 1.20-1.65%.

[0035] Si: A solid solution strengthening element and also a deoxidizing element in steel. However, excessive Si content can deteriorate the weldability of the steel, particularly the toughness of the weld heat-affected zone (HAZ). It also hinders the removal of oxide scale during hot rolling. Therefore, the Si content in this invention is controlled to 0.10-0.30%.

[0036] Cu and Ni: Solid solution strengthening elements that improve steel toughness. They also enhance the steel's hardenability and delay pearlite transformation during cooling. However, excessive addition can negatively impact the steel's toughness and weldability. Therefore, in this invention, the Cu and Ni contents are controlled to ≤0.30%.

[0037] Cr and Mo: Solid solution strengthening elements, important for improving the hardenability of steel, promote the formation of acicular ferrite, improve steel toughness, and help ensure uniform microstructure and properties throughout the thickness of the steel plate. However, excessive amounts can promote the formation of low-temperature phase transformation structures, which is detrimental to low-temperature toughness. Therefore, in the present invention, the Cr content is controlled to 0.15-0.30%, and the Mo content is controlled to ≤0.15%.

[0038] Nb: One of the important elements of low-carbon microalloyed steel, Nb, which is dissolved during hot rolling, strain-induced precipitation forms Nb(C, N) particles, which can delay austenite recrystallization and increase the austenite recrystallization temperature of the steel. This makes two-stage controlled rolling possible, avoids the formation of mixed crystals, and is conducive to obtaining a uniform and fine structure. While improving the strength and toughness of the material, it also ensures the isotropy of the steel plate performance. In addition, during the cooling process of the dissolved Nb, the second phase particles NbC are dispersed and precipitated in the matrix, playing a precipitation strengthening role. However, excessive Nb content cannot be completely dissolved, not only does it not play a role but also increases production costs. It also causes NbC to precipitate prematurely at high temperatures, easily forming large-sized NbC precipitation at the center segregation, which is detrimental to toughness. Therefore, in the present invention, the Nb content is controlled to 0.03-0.05%.

[0039] V: A key microalloying element, it easily combines with carbon in steel to form VC, offering strong precipitation strengthening but weak grain-refining properties. However, excessive V content can significantly coarsen the precipitated particles, impairing the steel's low-temperature toughness. In the present invention, V primarily compensates for the strength reduction caused by other alloying deficiencies by exerting precipitation strengthening. Therefore, the V content is controlled to ≤ 0.05%.

[0040] Furthermore, to fully utilize the carbon fixation effect of Nb and V and reduce the effect of free carbon on strain aging properties, it is necessary to control (Nb / 7.7C+V / 4.3C) to be ≥ 0.10. In some embodiments, (Nb / 7.7C+V / 4.3C) is controlled within the range of 0.10 to 0.50, such as within the range of 0.10 to 0.47.

[0041] Ti: It is a strong carbonitride-forming element that can play a role in nitrogen fixation. It can form fine, high-temperature stable TiN precipitates during slab continuous casting, thereby improving the toughness of the material. In addition, the undissolved carbonitrides of Ti can prevent the growth of austenite grains when the steel is heated, and the TiN and TiC precipitated during rough rolling in the high-temperature austenite zone can effectively inhibit the growth of austenite grains. In addition, during the welding process, the TiN and TiC particles in the steel can significantly prevent the growth of grains in the heat-affected zone, thereby improving the welding performance of the steel plate. Therefore, in the present invention, the Ti content is controlled at 0.008-0.020%; at the same time, in order to give full play to the nitrogen fixation effect of Ti, it is necessary to control Ti / N≥2. In some embodiments, Ti / N is controlled within the range of 2.0-5.0.

[0042] Ca: It is primarily used to modify MnS inclusions, spheroidizing them and ensuring their uniform distribution, which improves low-temperature toughness. However, excessive amounts can easily form CaO inclusions, which negatively impact toughness. Therefore, in this invention, the Ca content is controlled between 0.0012% and 0.0030%. Furthermore, to fully leverage Ca's sulfide-modifying properties, the ratio of Ca / S is controlled to 1.5 ≤ Ca / S ≤ 4, transforming elongated MnS inclusions into dispersed, spherical CaS inclusions, thereby improving the material's low-temperature toughness.

[0043] Al: Mainly used for deoxidation of steel, adding an appropriate amount of Al helps refine grains and improve the steel's strength and toughness. However, increasing Al content can lead to the formation of Al oxides in the steel, reducing the toughness of the base metal and the weld heat-affected zone. Therefore, in the present invention, the Al content is controlled to 0.015-0.045%.

[0044] S, P, B, N, O, and H are inevitable impurity elements in steel. P is a readily segregating element in steel; S easily forms sulfides in steel, which severely reduce the steel's low-temperature toughness; B easily precipitates at grain boundaries, reducing the material's plasticity and toughness; excessive N and H content can easily cause cracks in continuous castings; and excessive O content forms various oxides, reducing hot workability, corrosion resistance, and toughness. Therefore, the present invention controls P ≤ 0.01%, S ≤ 0.001%, B ≤ 0.0005%, N ≤ 0.005%, O ≤ 0.0025%, and H ≤ 0.0002%.

[0045] Pcm: An indicator of weld cold crack sensitivity. When Pcm exceeds a certain value, the steel's hardenability becomes excessive, significantly reducing the toughness of the weld zone. The present invention controls Pcm to ≤ 0.17. In some embodiments, Pcm is within the range of 0.12 to 0.17.

[0046] The present invention utilizes a low-C, low-Mn alloy composition system to reduce central segregation while meeting the required strength of the steel plate. Combined with the effects of Cr, Mo, Cu, and Ni, which improve hardenability and inhibit ferrite transformation, the ferrite content is controlled between 20% and 40%, further reducing central segregation and achieving good uniformity across the thickness of the steel plate. The difference in SA% between single-side thinned and full-thickness DWTT samples is ≤5%, and the difference between the surface hardness and the center hardness of the steel plate is ≤25 HV10. Furthermore, through the combination of C, Mn, Cu, Ni, Cr, Mo, and V, the Ceq ratio is controlled to be 0.30 ≤ ≤ 0.39, ensuring the steel plate's strength performance.

[0047] Microalloying elements Nb and V are usually added to pipeline steel to play the role of grain refinement and precipitation strengthening. However, in the present invention, the fixing effect of Nb and V on carbon atoms is utilized, and (Nb / 7.7C+V / 4.3C)≥0.10 is designed to maximize the fixing effect of Nb and V on carbon atoms, reduce the influence of free carbon on strain aging performance, and meet the unconventional performance requirements of submarine pipeline steel for tensile performance after strain aging. After the strain aging test of the submarine pipeline steel of the present invention after 3% pre-strain and 250℃ insulation for 1 hour, the longitudinal tensile properties of the steel plate are: yield strength R t0.5 : 405~530MPa, tensile strength R m : 505~650MPa, yield strength ratio R t0.5 / Rm≤0.88.

[0048] The present invention controls the ferrite content to 20-40%. Obtaining an appropriate proportion of ferrite is key to meeting the performance requirements of submarine pipeline steel. Too low a ferrite content cannot meet the low yield ratio requirement, while too high a ferrite content results in insufficient high strength. Therefore, in terms of composition design, the addition of Cu, Ni, Cr, and Mo alloying elements achieves a certain austenite stability and can control the ferrite transformation amount within the desired range.

[0049] The method for manufacturing high-performance uniform submarine pipeline steel of the present invention comprises the following steps:

[0050] 1) Smelting and continuous casting

[0051] The above composition is smelted and continuously cast into billets; during the continuous casting process, the tundish superheat is controlled at 15-40°C, and the continuous casting billet drawing speed is controlled at 0.40-1.10 m / min to ensure that the solidification end of the continuous casting billet is in the horizontal section of the continuous casting machine; the continuous casting machine crystallizer adopts an electromagnetic stirring process;

[0052] 2) Controlled rolling

[0053] The heating temperature is 1120-1190°C, the rough rolling start temperature is 1000-1100°C, the rough rolling finish temperature is 930-980°C, the rough rolling reduction ratio is ≥50%, and the width ratio is 1.20≤≤1.80; the finishing rolling start temperature is 790-880°C, the finishing rolling finish temperature is 780-850°C, and the finishing rolling reduction ratio is ≥70%;

[0054] 3) Controlled cooling

[0055] Relaxation time ≥ 30s, cooling start temperature 730-780℃, cooling rate 10-40℃ / s, cooling stop temperature 350-550℃, after cooling stop, the temperature difference between the upper and lower surfaces of the steel plate ≤ 30℃;

[0056] 4) After the steel plate stops cooling, air cool or pile cool it to room temperature.

[0057] Preferably, in step 4), when the cooling stop temperature is ≤ 450°C, the steel plate is tempered at 450-650°C for 10-30 minutes.

[0058] Preferably, in step 1), the fluctuation value of the superheat degree of adjacent slabs is ≤10°C and / or the fluctuation value of the drawing speed of adjacent slabs is ≤10%.

[0059] Preferably, in step 1), a dynamic soft reduction process is adopted at the end of solidification of the continuous casting billet, with a reduction amount of 2 to 3 mm.

[0060] Preferably, in step 3), the temperature difference between the head and tail of the steel plate entering water is ≤30°C.

[0061] Preferably, in step 4), after the steel plate is cooled to room temperature, the head and tail portions of the steel plate are cut off, and the head and tail cuts are each ≥ 250 mm, such as 250-350 mm. In some embodiments, the head cut is 250-350 mm, and the tail cut is 250-300 mm.

[0062] In the manufacturing method of the present invention:

[0063] In order to reduce the center segregation of the continuous casting billet and improve the low-temperature toughness and isotropy of the steel plate, the continuous casting process needs to meet the following requirements:

[0064] ① The superheat of the tundish is controlled at 15-40℃, and the superheat fluctuation of adjacent slabs is ≤10℃;

[0065] ② The continuous casting billet pulling speed is controlled at 0.40~1.10m / min, with constant pulling speed control, and the fluctuation of the pulling speed of adjacent slabs is ≤10%, ensuring that the solidification end of the continuous casting billet is in the horizontal section of the continuous casting machine;

[0066] ③ Using electromagnetic stirring technology in the continuous casting mold section;

[0067] ④ Dynamic light reduction process is adopted at the end of solidification of continuous casting billet.

[0068] This invention ensures the stability of the continuous casting process by controlling the tundish superheat and the casting speed, and strictly limiting their fluctuation range. This ensures that the final solidification of the continuous cast strand is in the horizontal section of the continuous caster, maximizing the dynamic soft reduction effect and thus effectively controlling the centerline segregation of the slab. Furthermore, the continuous casting process utilizes electromagnetic stirring, which effectively homogenizes the element distribution, thereby reducing segregation tendencies and ensuring uniform properties of pipeline steel.

[0069] The slab heating temperature is 1120-1190℃. If the heating temperature is too low, the alloy carbides will not be fully dissolved, and it will be difficult to fully exert the effects of solid solution strengthening and precipitation strengthening, and it will be unfavorable for the control of central segregation. If the heating temperature is too high, the grains will coarsen, which will be detrimental to the low-temperature toughness of the steel plate.

[0070] The starting temperature for rough rolling is 1000-1100°C, and the finishing temperature for rough rolling is 930-980°C. The rough rolling reduction is ≥50%, and the width ratio is 1.20≤≤1.80. Rough rolling is performed in the fully recrystallized austenite region to avoid the formation of mixed crystals. A sufficient rough rolling reduction can fully refine the original austenite grains, which is beneficial to the strength and toughness of the steel plate. Precise control of the width ratio also facilitates shape control and reduces differences in the transverse and longitudinal properties of the steel plate. In some embodiments, the rough rolling reduction is 50-80%, such as 50-74%.

[0071] Finish rolling begins at a temperature of 790-880°C and ends at a temperature of 780-850°C, with a finishing reduction of ≥70%. Finish rolling is performed in the non-recrystallized austenite region. A sufficient finishing reduction can fully flatten the austenite grains, providing more nucleation sites for subsequent phase transformation, thereby refining the final microstructure and improving the strength and toughness of the steel plate. In some embodiments, the roughing reduction is 70-85%.

[0072] The relaxation time is ≥30s, the start cooling temperature is 730-780℃, the cooling rate is 10-40℃ / s, and the stop cooling temperature is 350-550℃. The appropriate cooling rate and stop cooling temperature can ensure the formation of a dual-phase structure of ferrite + bainite, ensuring the strength and toughness of the steel plate while obtaining the characteristics of a low yield ratio; a certain relaxation time and start cooling temperature can ensure the generation of a certain amount of proeutectoid ferrite in the steel plate structure, thereby reducing the yield ratio while ensuring strength. The temperature difference between the head and tail of the steel plate entering the water is ≤30℃. Reducing the temperature difference between the head and tail entering the water is conducive to ensuring the isotropy of the steel plate performance. After cooling, the temperature difference between the upper and lower surfaces of the steel plate is ≤30℃. Reducing the temperature difference between the upper and lower surfaces after cooling is conducive to improving the isotropy of the performance in the thickness direction. In some embodiments, the relaxation time is 30-60s.

[0073] Higher final rolling temperature, lower final cooling temperature and faster cooling rate are beneficial to inhibiting the enrichment of alloy elements to the center of the plate thickness during phase transformation, thereby further reducing the impact of central segregation.

[0074] After the steel plate stops cooling, it is cooled to room temperature. Preferably, the steel plate can be cooled to room temperature by stack cooling.

[0075] Preferably, for steel plates with a cooling stop temperature of ≤450°C, tempering can be performed at 450-650°C for 10-30 minutes.

[0076] After the steel plate stops cooling, it can usually be cooled to room temperature on a cooling bed; or the rolled steel plates can be stacked together and cooled to room temperature.

[0077] The slower cooling rate of the stack cooling method allows the steel plate to remain at a high temperature for a longer period of time, which is conducive to the precipitation of alloy carbides and fully exerts the effects of precipitation strengthening and carbon fixation. For steel plates with lower cooling temperatures, the precipitation time of alloy carbides is insufficient. Using a short high-temperature tempering method can fully exert the effects of alloy precipitation strengthening and carbon fixation, while also facilitating the decomposition of the hard phase MA component and improving low-temperature strength and toughness.

[0078] The head and tail cutoff of the finished steel plate after rolling is ≥250mm. Due to the large differences in the deformation conditions, cooling conditions and the steel plate body at the head and tail of the steel plate, the organizational properties are quite different from those of the steel plate body. Ensuring a certain head and tail cutoff eliminates the influence of uncontrolled deformation and cooling positions, which is conducive to achieving isotropy of steel plate performance.

[0079] Compared with the prior art, the present invention has the following advantages:

[0080] In terms of composition design, the present invention adds relatively low contents of C and Mn elements, controls the impurity elements P, S, B, N, O, and H at extremely low levels, combines the addition of alloying elements Cr, Mo, Cu, and Ni, and controls 0.30≤Ceq≤0.39 and Pcm≤0.17, thereby improving central segregation, affecting the phase transformation behavior during the cooling process of the material, increasing the stability of austenite, and delaying the ferrite phase transformation process. On the basis of ensuring the mechanical properties of the steel plate, the steel plate performance has good uniformity.

[0081] The present invention utilizes the effect of adding trace amounts of V-Nb-Ti, while strictly limiting Ti / N≥2 and (Nb / 7.7C+V / 4.3C)≥0.10, to form microalloy carbonitrides, which refine the grains, strengthen precipitation, and fix C and N atoms, thereby reducing the effect of strain aging on material properties. The material still maintains high performance after strain aging. After the steel plate is subjected to a 3% pre-strain and a 1-hour heat preservation at 250°C, the longitudinal tensile properties of the steel plate are: yield strength R t0.5 : 405~530MPa, tensile strength R m : 505~650MPa, yield strength ratio R t0.5 / Rm≤0.88.

[0082] Based on the composition design, the continuous casting process strictly controls the tundish superheat, casting speed, and its fluctuation range, and utilizes dynamic soft reduction to ensure that the final solidification of the continuous cast slab is in the horizontal section of the continuous caster. This maximizes the dynamic soft reduction effect, thereby reducing segregation and ensuring uniform performance of pipeline steel. Conventional continuous casting processes, however, have no specific regulations or low requirements for tundish superheat, casting speed, and their fluctuation range. Parameters are set primarily to match the production rhythm of the preceding and subsequent steps. As a result, the final solidification of the continuous cast slab often occurs in the curved section of the continuous caster. The dynamic soft reduction applied in the horizontal section cannot fully enhance the internal quality of the slab and reduce center segregation, thus failing to meet the requirements for high performance and uniformity.

[0083] In addition, the two-stage controlled rolling technology and width ratio control in the austenite recrystallization zone and non-recrystallization zone are used to ensure a sufficient compression ratio in the two-stage rolling process, fully refine the grains, avoid mixed crystal structure, and control the rolling temperature, relaxation time, and cooling temperature from both thermodynamic and kinetic aspects to control the ferrite phase transformation process, form a two-phase structure of ferrite + bainite, and accurately control the ferrite ratio at 20-40%, ensuring that the steel plate has excellent mechanical properties such as strength and toughness and isotropy, solving the problem of mismatch between the strength and performance uniformity of submarine pipeline steel. BRIEF DESCRIPTION OF THE DRAWINGS

[0084] FIG1 is a microstructure photograph of the steel of Example 1 of the present invention. DETAILED DESCRIPTION

[0085] The present invention will be further described below with reference to the accompanying drawings.

[0086] The chemical composition of the steel according to the embodiment of the present invention is shown in Tables 1 and 2, with the remainder including Fe and unavoidable impurities; the specific process parameters are shown in Tables 3 and 4.

[0087] Tensile property test: At room temperature, according to ASTM A370-2021, the tensile property test of the steel plate samples obtained in each embodiment of the present invention was carried out to measure the yield strength, tensile strength and yield strength ratio of the steel plate of each embodiment. The longitudinal tensile properties of the steel plate after 3% prestrain and 1 hour of strain aging at 250°C were measured. The specific results are shown in Table 5.

[0088] Low-Temperature Toughness Test: Impact tests were conducted on the test steel plates at -45°C using V-notch specimens in accordance with ASTM E23-2018. The -45°C Charpy impact energy of the steel plates of each example was measured, and the specific results are shown in Table 6. Drop weight tear tests were conducted on the test steel plates at -30°C using full-thickness V-notch specimens and single-side thinned V-notch specimens in accordance with API RP 5L3-2014 to measure the SA value of the steel plates of each example at -30°C under the DWTT test. The specific results are shown in Table 7.

[0089] Hardness test: According to ASTM E92, under 10 kgf, the hardness values ​​of the steel plate near the upper surface, the center of the thickness, and the near lower surface were measured. The specific results are shown in Table 8.

[0090] Grain size test method is in accordance with ASTM E112.

[0091] As can be seen from FIG1 , the microstructure of the steel plate obtained by using the composition and process of the present invention is a ferrite+bainite structure, wherein the ferrite content is 20-40%.

[0092] As can be seen from Table 5, the transverse tensile properties of the steel of the present invention: yield strength R t0.5 : 415~555MPa, tensile strength Rm: 515~670MPa, yield strength ratio R t0.5 / Rm0.88, and the strength difference between the head and tail of the steel plate is ≤50MPa.

[0093] Longitudinal tensile properties: yield strength R t0.5 :405~530MPa, tensile strength Rm:505~650MPa, yield strength ratio R t0.5 / Rm0.86, and the strength difference between the head and tail of the steel plate is ≤50MPa.

[0094] After 3% pre-strain and 250℃ heat preservation for 1 hour, the longitudinal tensile properties of the steel plate are: yield strength R t0.5 :405~530MPa, tensile strength Rm:505~650MPa, yield strength ratio R t0.5 / Rm0.88.

[0095] It can be seen from Table 6 that the -45°C Charpy impact energy of the steel plate obtained by the present invention is ≥250J, and the difference in Charpy impact energy between the center of the thickness of the steel plate and the near surface is ≤50J.

[0096] As can be seen from Table 7, the SA of the steel plate obtained in the present invention at -30°C DWTT test is ≥85%. In particular, the SA difference between the single-side thinned DWTT specimen (tested at a correspondingly lower temperature according to API RP 5L3) and the full-thickness DWTT specimen is ≤5%.

[0097] As can be seen from Table 8, the present invention obtains a steel plate hardness of ≤230HV10, and in particular, the difference between the hardness near the surface of the steel plate and the hardness at the center of the steel plate is ≤25HV10.

[0098] In Comparative Example 1, Ti / N=1.5, Ca / S=1.12, neither of which meets the requirements of Ti / N≥2, 1.5≤Ca / S≤4 of the present invention, and the P, S, and N contents are relatively high. The pulling wave velocity in the process control is not within the control range of the present invention, and cannot meet the requirement that the solidification end of the continuous casting billet is in the horizontal section of the continuous casting machine, and the superheat fluctuation value, the temperature difference between the head and tail of the steel plate entering the water, and the temperature difference between the upper and lower surfaces of the steel plate after cooling are all not within the control range of the present invention. Although the strength of the steel plate finally obtained can meet the requirements, the strength difference between the head and tail is too large, exceeding 50MPa, and the yield strength is relatively high. The difference in Charpy impact energy between the center and near surface of the steel plate thickness is 66, the SA difference between the single-sided thinning DWTT sample and the full-thickness DWTT sample is 14, and the difference between the surface hardness of the steel plate and the center hardness of the steel plate is 35. It is obvious that the performance uniformity of the steel plate obtained in Comparative Document 1 is not good.

[0099] In Comparative Example 2, Nb / 7.7C+V / 4.3C=0.09, which does not meet the requirement of (Nb / 7.7C+V / 4.3C)≥0.10 of the present invention. Furthermore, the B and O contents are relatively high. In the process control, the heating temperature, the rough rolling and finishing temperature, and the relaxation time are all outside the ranges specified in the present invention. The resulting -45°C Charpy impact energy is less than 250 J, and the DWTT test SA at -30°C is less than 85%, indicating poor low-temperature toughness of the steel plate.

[0100] In Comparative Example 3, Ceq=0.40, Pcm=0.18, and Nb / 7.7C+V / 4.3C=0.07, all of which do not meet the requirements of the present invention of 0.30≤Ceq≤0.39, Pcm≤0.17, and (Nb / 7.7C+V / 4.3C)≥0.10. The SA difference between the single-sided thinned DWTT specimen and the full-thickness DWTT specimen of the steel plate finally obtained is 13, and the Charpy impact energy difference between the thickness center and the near-surface is 115J. The performance uniformity of the steel plate obtained in Comparative Example 3 is poor.

[0101] Table 1 (Unit: weight percentage)

[0102] Table 2

[0103] Note: Ceq = C + Mn / 6 + (Cr + V + Mo) / 5 + (Cu + Ni) / 15; Pcm = C + Si / 30 + (Mn + Cu + Cr) / 20 + Ni / 60 + Mo / 15 + V / 10 + 5B.

[0104] Table 3

[0105] Table 4

[0106] Table 5

[0107] Table 6

[0108] Table 7

[0109] Table 8

Claims

1. A high-performance uniform submarine pipeline steel, the chemical composition weight percentage of which is: C: 0.02-0.06%, Si: 0.10-0.30%, Mn: 1.20-1.65%, Cu≤0.30%, Ni≤0.30%, Cr: 0.15-0.30%, Mo≤0.15%, Nb: 0.03-0.05%, V≤0.05%, Ti: 0.008-0.020%, Ca: 0.0012-0.0030%, Al: 0.015-0.045%, the balance includes Fe and unavoidable impurities, wherein, Impurity elements need to be controlled at P≤0.01%, S≤0.001%, B≤0.0004%, N≤0.005%, O≤0.0025%, H≤0.0002%; and at the same time meet the following requirements: Ti / N≥2; (Nb / 7.7C+V / 4.3C)≥0.10; 1.5≤Ca / S≤4; 0.30≤Ceq≤0.39, Ceq=C+Mn / 6+(Cr+V+Mo) / 5+(Cu+Ni) / 15; Pcm≤0.17, Pcm=C+Si / 30+(Mn+Cu+Cr) / 20+Ni / 60+Mo / 15+V / 10 +5B.

2. The high-performance uniform submarine pipeline steel according to claim 1, characterized in that: The balance is Fe and unavoidable impurities.

3. The high performance uniform submarine pipeline steel according to claim 1 or 2, characterized in that: The microstructure of the submarine pipeline steel is ferrite+bainite, wherein the volume content percentage of the ferrite is 20-40%.

4. The high performance uniform submarine pipeline steel as claimed in claim 1, 2 or 3, characterized in that: The grain size grade of the submarine pipeline steel is ≥9.

5. The high performance uniform submarine pipeline steel according to claim 1, 2, 3 or 4, characterized in that: The transverse tensile properties of the submarine pipeline steel are: yield strength Rt 0.5 : 415~555MPa, tensile strength R m : 515~670MPa, yield strength ratio Rt 0.5 / R m ≤0.88; the difference in yield strength and tensile strength between the head and tail of the steel plate is ≤50MPa; Longitudinal tensile strength: yield strength R t0.5 : 405~530MPa, tensile strength R m : 505~650MPa, yield strength ratio R t0.5 / R m ≤0.86; The difference in yield strength and tensile strength between the head and tail of the steel plate is ≤50MPa; After 3% pre-strain and 250℃ heat preservation for 1 hour, the longitudinal tensile properties of the steel plate are: yield strength R t0.5 : 405~530MPa, tensile strength R m : 505~650MPa, yield strength ratio R t0.5 / R m ≤0.

88.

6. The high-performance uniform submarine pipeline steel according to claim 1, 2, 3, 4 or 5, characterized in that: The submarine pipeline steel has a Charpy impact energy of ≥250J at -45°C, and a Charpy impact energy difference between the center and near-surface of the steel plate thickness of ≤50J; a DWTT test SA of ≥85% at -30°C, and a SA difference between a single-sided thinned DWTT specimen and a full-thickness DWTT specimen of ≤5%; and a hardness of ≤230HV10, wherein the difference between the surface hardness of the steel plate and the hardness at the center of the steel plate is ≤25HV10.

7. The method for manufacturing high-performance uniform submarine pipeline steel according to any one of claims 1 to 6, characterized in that: The steps include: 1) Smelting and continuous casting The composition according to claim 1 or 2 is smelted and continuously cast into billets; during the continuous casting process, the superheat of the tundish is controlled at 15-40°C, and the pulling speed of the continuous casting billet is controlled at 0.40-1.10 m / min to ensure that the solidification end of the continuous casting billet is in the horizontal section of the continuous casting machine, and the continuous casting machine crystallizer adopts an electromagnetic stirring process; 2) Controlled rolling The heating temperature is 1120-1190°C, the rough rolling start temperature is 1000-1100°C, the rough rolling final rolling temperature is 930-980°C, the rough rolling reduction rate is ≥50%, and the width ratio is 1.20≤≤1.80; the finishing rolling start temperature is 790-880°C, the finishing rolling final rolling temperature is 780-850°C, and the finishing rolling reduction rate is ≥70%; 3) Controlled cooling Relaxation time ≥30s, cooling start temperature 730~780℃, cooling rate 10~40℃ / s, cooling stop temperature 350~550℃, temperature difference between upper and lower surfaces of steel plate after cooling stop ≤30℃; 4) After the steel plate stops cooling, air cool or pile cool it to room temperature.

8. The manufacturing method according to claim 7, characterized in that: In step 1), the fluctuation value of superheat degree of adjacent slabs is ≤10°C and / or the fluctuation value of drawing speed of adjacent slabs is ≤10%.

9. The manufacturing method according to claim 7 or 8, characterized in that: In step 1), a dynamic light reduction process is adopted at the end of solidification of the continuous casting billet, with a reduction amount of 2 to 3 mm.

10. The manufacturing method according to claim 7 or 8, characterized in that: In step 1), the rough rolling reduction ratio is 50-80%, and the finish rolling reduction ratio is 70-85%.

11. The manufacturing method according to claim 7, characterized in that: In step 3), the temperature difference between the head and tail of the steel plate entering water is ≤30°C.

12. The manufacturing method according to claim 7, characterized in that: In step 3), the relaxation time is 30 to 60 s.

13. The manufacturing method according to claim 7, characterized in that: In step 4), when the cooling stop temperature is ≤450°C, the steel plate is tempered at 450-650°C for 10-30 min.

14. The manufacturing method according to claim 7 or 13, characterized in that: In step 4), after the steel plate is cooled to room temperature, the head and tail parts of the steel plate are cut off, and the head and tail cutting amounts are both ≥250mm.

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