Method of manufacturing high-strength steel pipes from steel compositions and parts thereof

Abstract

A method for manufacturing tubes from a well-defined steel composition, particularly for tubes used in gas storage pressurizer pressure vessels, comprising the following steps: a) producing steel tubes from the steel composition, including at least one hot rolling or hot forming process; b) subjecting the steel tubes to a cold drawing process to obtain the desired dimensions, wherein the cold drawing process includes at least two drawing operations and an intermediate austenitizing and quenching step prior to the final drawing of the cold drawing process; c) subsequently subjecting the cold-drawn steel tubes to a final recovery heat treatment in a temperature range of 200–600°C.

Description

[0001] This invention relates to a method for manufacturing high-strength steel tubes from steel compositions (such as microalloyed low-carbon steel) and tubular components thereof. The steel tubes manufactured according to this invention are particularly suitable for manufacturing components of automotive restraint systems, such as automotive airbag inflators.

[0002] The automotive industry has always sought ways to improve vehicle efficiency, with the development of engines that improve fuel efficiency and reduce weight playing a crucial role in reducing fuel consumption. Weight reduction can be achieved by decreasing the thickness of parts without compromising strength and safety requirements. Currently, advanced high-strength steels offer a high strength-to-density ratio, but they require expensive alloying processes and manufacturing cycles. Therefore, the industry has been searching for new high-strength steel products that achieve excellent final properties at a competitive cost.

[0003] The present invention relates to tubes and tubular components made of steel having improved or at least sufficient strength, ductility and toughness properties, allowing for such weight reduction, particularly for use as tubular components in airbag inflators.

[0004] EP2078764A1 (Sumitomo Metal Industries, Ltd.) discloses a seamless steel tube for use in airbag accumulators. This steel tube can be manufactured by normalizing heat treatment without quenching and tempering. The steel tube has a tensile strength of at least 850 MPa and resistance to bursting at -20°C. The composition of the steel tube, by mass%, includes: C: 0.08-0.20%, Si: 0.1-1.0%, Mn: 0.6-2.0%, P: less than 0.025%, S: less than 0.010%, Cr: 0.05-1.0%, Mo: 0.05-1.0%, Al: 0.002-0.10%, and at least one of the following: Ca: 0.0003-0.01%, Mg: 0.0003%. -0.01% and REM (rare earth metals): 0.0003-0.01%, at least one of the following: Ti: 0.002-0.1% and Nb: 0.002-0.1%, Ceq (defined according to the formula Ceq=C+Si / 24+Mn / 6+(Cr+Mo) / 5+(Ni+Cu) / 15) in the range of 0.45-0.63, the metallographic structure is a mixed structure of ferrite and bainite.

[0005] WO2005 / 035800A1 (Lopez et al.) generally discloses a low-carbon alloy steel pipe and a method for manufacturing the same, wherein the steel pipe is mainly composed of the following (in weight %): about 0.06-0.18% carbon; about 0.5-1.5% manganese; about 0.1%-0.5% silicon; up to about 0.015% sulfur; up to about 0.025% phosphorus; up to about 0.50% nickel; about 0.1-1.0% chromium; about 0.1-1.0% molybdenum; about 0.01%-0.10% vanadium; about 0.01-0.10% titanium; about 0.05-0.35% copper; about 0.010-0.050% aluminum; up to about 0.05% niobium; up to about 0.15% residual elements; and the balance being iron and incidental impurities. The manufacturing process of this steel pipe includes the following steps: steelmaking, steel casting, hot rolling, hot-rolled hollow finishing, cold drawing, heat treatment including quenching and tempering after cold drawing, and additional cold-drawn finishing. The resulting pipe has a tensile strength of 1000 MPa or higher, and therefore high burst strength.

[0006] WO2007 / 113642A2 (Lopez et al.) discloses a tube made of a similar low-carbon alloy steel composition and its improved manufacturing process, including a rapid induction austenitization / high-speed quenching step after cold drawing, preferably without tempering heat treatment.

[0007] It has now been found that tubes manufactured using Lopez’s existing technological processes either have strength at the expense of ductility, or exhibit ductility but lower strength levels, especially after tube finishing operations such as straightening and cold working.

[0008] The main objective of this invention is to provide steel pipes with improved properties, particularly with regard to the combination of strength and ductility, and more specifically, wherein the combination of strength and ductility characteristics remains unchanged or is at least less affected when the ends of the straightened and cold-formed steel pipes are subjected to finishing operations such as stretching.

[0009] Considering that manufacturing automotive parts typically involves welding steps, such as manufacturing pressure vessels for airbag inflators, another object of the present invention is to provide such steel pipes made of weldable steel components.

[0010] Another object of the present invention is to provide an alternative method for manufacturing high-strength steel tubes for use in airbag inflators.

[0011] The inventors have now discovered that a novel manufacturing process for producing steel pipes from specific steel compositions provides an advantageous combination of strength and ductility. Summary of the Invention

[0012] A method for manufacturing steel pipes from the steel composition according to the invention, particularly for use in pressure vessels for airbag inflators, as defined in claim 1.

[0013] The method includes the following steps:

[0014] a) Steel pipes are produced from the steel composition described below, including at least one hot rolling or hot forming process;

[0015] b) The steel pipe is subjected to a cold drawing process to obtain the required dimensions, wherein the cold drawing process includes at least two drawing operations and an intermediate austenitizing and quenching step prior to the final drawing operation of the cold drawing process.

[0016] c) After the final drawing in the cold drawing process, the cold-drawn steel pipe is subjected to a final recovery heat treatment at a temperature range of 200-600℃.

[0017] In step b) of the method according to the invention, the intermediate austenitizing and quenching steps, wherein the cold-drawn steel tube is heated at least once to a temperature of at least Ac3 to promote a fine-grained microstructure, typically rapidly heated such as by induction heating over a time span of a few seconds, and then quenched before final drawing, ensure that the tube is predominantly martensitic with sufficient strain hardening capacity after cold drawing, and that the subsequent cold drawing or drawing applies sufficient deformation to strain hardening, thereby obtaining excellent strength properties.

[0018] The inventors have discovered that there are significant differences in the sensitivity to strength and ductility properties among tubular products manufactured using different methods.

[0019] Tubular products that are cold-drawn and then quenched (i.e., without further heat treatment or cold drawing) acquire high strength but suffer a significant loss of ductility when stretched. Quenched tubular products are not used in this way; instead, they typically undergo further processing, particularly straightening and edge cold forming, to transform them into finished products, such as those ready for assembly into automotive airbag inflators. Both of these processes involve cold deformation following heat treatment, causing changes in the microstructure of the steel tubular product, most notably increasing hardness by increasing the number of dislocations, but simultaneously reducing ductility and toughness. This embrittlement is exacerbated by aging, as shown in laboratory simulations at 250°C for 1 hour (considered to represent aging at room temperature for months or more). Aging promotes the accumulation of interstitial carbon (i.e., carbon in the solid solution) at these dislocations, further impairing ductility. The more carbon in the solid solution, the higher the dislocation density, and the worse the embrittlement.

[0020] Compared to tubular products that are cold-drawn and then quenched, tubular products that are cold-drawn, quenched, and then tempered (i.e., without further cold drawing) are less sensitive to the loss of ductility after stretching (and aging), but have lower strength properties. Tempering after quenching, to some extent, reduces internal micro-strain by promoting microstructural transformations such as carbide precipitation and dislocation recovery, thereby eliminating internal stress and restoring plasticity and toughness.

[0021] The tubular products of the present invention, subjected to cold drawing, intermediate austenitization followed by quenching, cold redrawing, and recovery treatment, exhibit higher strength compared to cold-drawn, quenched, and tempered steel pipes, and show less impact on ductility compared to cold-drawn and quenched tubular products, particularly after post-drawing (straightening and cold forming, especially at the ends). The recovery treatment following the final drawing in the cold drawing process, within the range of 200-600°C, such as 300-600°C, is sufficient to ensure uniform precipitation of carbides. This serves to increase formability.

[0022] Furthermore, any post-recovery heat treatment performed at much lower temperatures has a negligible effect on the microstructure. It is also assumed that sensitivity to aging, which is related to the diffusion of free interstitial elements (primarily carbon), is suppressed in this invention.

[0023] Therefore, compared to tubular products that are cold-drawn and then quenched, the tubular products produced according to the present invention have similar (or even higher) strength and good elongation properties, but are much less sensitive to loss of ductility due to stretching. Compared to tubular products that are cold-drawn, quenched, and then tempered, the tubular products produced according to the present invention have much higher strength and similar elongation properties at the same recovery and tempering temperatures, respectively. The higher strength properties allow for the use of tubular components with smaller wall thicknesses, resulting in lighter weight components in their final applications.

[0024] In the method according to the invention, at least one cold drawing is performed after the intermediate austenitizing and quenching steps. Preferably, the total area reduction of one or more drawing operations after the intermediate austenitizing and quenching steps is at least 10%, more preferably at least 15%, and more preferably at least 20%, thereby ensuring sufficient strain hardening after the intermediate austenitizing and quenching steps. For example, a 20% total area reduction can be achieved by a penultimate drawing with a 10% area reduction and a final drawing with an 11% area reduction after the intermediate austenitizing and quenching steps.

[0025] In a preferred embodiment, the intermediate austenitization and quenching steps are performed between the penultimate and final drawing steps of cold drawing (b). Advantageously, the deformation, measured by area reduction, in the final drawing of the cold drawing process is at least 10%, preferably at least 15%, and more preferably at least 20%.

[0026] It should be noted that EP2650389A2 (Tenaris Connections BV) discloses a method for manufacturing steel pipes and rods suitable for mining applications, aiming for high wear resistance, high impact toughness, and good dimensional tolerances. The steel composition in EP2650389A2 comprises approximately 0.18-0.32 wt% carbon, approximately 0.3-1.6 wt% manganese, approximately 0.1-0.6 wt% silicon, approximately 0.005-0.08 wt% aluminum, approximately 0.2-1.5 wt% chromium, approximately 0.2-1.0 wt% molybdenum, with the balance including iron and impurities. The pipe can be cold-drawn in a first cold-drawing operation to achieve an area reduction of approximately 15%-30%, followed by heat treatment to an austenitizing temperature between approximately 50°C above AC3 and less than approximately 150°C above AC3, and then quenched to approximately room temperature at a rate of at least 20°C / s. The pipe can then be cold-drawn a second time to achieve an area reduction of approximately 6%-14%. A second heat treatment can be performed on the tube by heating it to a temperature of approximately 400-600°C for about 15-60 minutes to provide stress relief. The tube can then be cooled to approximately room temperature.

[0027] The steel composition used in the method according to the invention, excluding Fe and unavoidable impurities, comprises, by weight percent,

[0028] C: 0.04-0.15;

[0029] Mn: 0.90-1.60;

[0030] Si: 0.10-0.50;

[0031] Cr: 0.05-0.80;

[0032] Al 0.01-0.50;

[0033] N 0.0035-0.0150

[0034] And if necessary, one or more optional elements are described below.

[0035] The process steps and composition of the method according to the present invention are explained in more detail below.

[0036] process

[0037] Step a) typically includes the following sub-steps: preparing the steel composition, casting the composition into a billet, piercing the billet at an elevated temperature, and hot rolling the pierced billet in at least one hot rolling process, optionally including an intermediate reheating step between two hot rolling processes, heating to a temperature above Ac3.

[0038] For example, a starting product derived from a low-carbon steel composition according to the invention, typically a pierceable solid steel bar or billet cast in a steel mill, is formed into a hollow (seamless) tube of a certain length. The solid billet is, for example, circular, and its diameter is, for example, about 148 mm. The solid billet is then heated and pierced, for example using the Mannesmann tube rolling process, and subsequently hot-rolled in a hot rolling mill in one or more subsequent hot rolling processes, during which the outer diameter and wall thickness are significantly reduced, while the length is significantly increased.

[0039] Advantageously, the billet is heated to a temperature in the range of 1250-1300°C. The temperature difference is maintained at or below 50°C during piercing. The shrinkage during piercing is preferably 2% or greater (RR ≥ 2%), for example, the outer diameter of the pierced hollow billet is 147 mm, and the wall thickness is 13 mm. The reduction in cross-sectional area, measured as the ratio of the cross-sectional area of ​​the solid billet to the cross-sectional area of ​​the hot-rolled hollow tube, helps to achieve the desired microstructure.

[0040] The hot rolling in step a) is performed in several steps. Advantageously, the mandrel rolling temperature in the first step is at least 1150°C. It is also advantageous that the shrinkage in each step, including the final step, is 3% or more (RR ≥ 3%). Preferably, the total minimum cross-sectional area is reduced by 15% or more, more preferably by 20% or more, and most preferably by 25% or more. For example, the hot-rolled tube has an outer diameter of 42.4 mm and a wall thickness of 2.8 mm.

[0041] The hot rolling process may include an intermediate reheating step, in which the hot-rolled intermediate product is reheated to a temperature above Ac3, such as 880°C (Ac3 of the following composition) or higher.

[0042] After hot rolling, the hot-rolled tube is cooled to ambient temperature at a suitable cooling rate, advantageously in still air. This results in a predominantly ferrite-bainite microstructure while avoiding the formation of hard microstructures. The intermediate tube product thus obtained has a nearly uniform wall thickness along its length and circumference.

[0043] In the method according to the invention, the normalizing treatment, including austenitization and slow (air) cooling, can be performed in a furnace after hot rolling, or the final hot rolling process can be performed as normalizing rolling (also known as normalizing forming). In normalizing rolling, the final rolling temperature is above Ar3, preferably between Ar3 and the grain coarsening temperature, more preferably between Ar3 and 1050°C, and most preferably in the range of 850-1000°C. If normalizing is performed in a furnace after hot rolling, the normalizing temperature is above Ac3, preferably maintained between Ac3 and 1000°C for a period of time to allow the phase transformation to complete, i.e., to allow the entire cross-section of the heat-treated tube to reach a temperature within this range.

[0044] Intermediate tube products can withstand various finishing processes, such as straightening, end trimming, cutting to the required length, and non-destructive testing.

[0045] In preparation for the subsequent cold drawing process, the surface of the tubes cut to a certain length is properly treated. Typical treatment steps include pickling, such as immersion in an acid solution, and applying one or more layers of one or more lubricants, such as a combination of zinc phosphate and sodium stearate or reactive oil.

[0046] The tube, having undergone appropriate surface treatment, then undergoes a cold drawing process comprising at least two steps, wherein the outer diameter and wall thickness of the tube are further reduced during each step. According to the invention, the cold drawing process includes an intermediate austenitizing and quenching step prior to the final step of the cold drawing process. The intermediate austenitizing and quenching step between cold drawing and drawing comprises heating the cold-drawn tube (rapidly) to above Ac3 at least once, as described above (advantageously by induction heating), and rapid cooling, advantageously by water quenching, preferably at a rate of at least 50°C / s, typically measured between 800°C and 500°C, continuing forced cooling until a temperature below the martensite initiation (Ms) temperature is reached, preferably below 100°C or lower, and more preferably below 50°C, thereby achieving a transformation to produce a hard martensite structure. As already mentioned, the total area reduction after the intermediate austenitizing and quenching step is preferably at least 10%, preferably at least 15%, more preferably at least 20%. In a preferred embodiment, the area reduction in the final drawing is at least 10% (RA ≥ 10%). Advantageously, the intermediate austenitizing and quenching steps are performed between the penultimate and final cold drawing. The final dimensions of the cold-drawn tube are, for example, an outer diameter in the range of 20-60 mm and a wall thickness in the range of 1-4 mm.

[0047] Intermediate normalizing treatment can be introduced during the cold drawing process before the austenitizing and quenching steps.

[0048] Following cold drawing, a final recovery heat treatment is performed within the range of 200-600°C (e.g., 300-600°C) to reduce internal stress and dislocation density, and to stabilize the microstructure. In this final recovery heat treatment, the steel tube is stress-relieved at temperatures within this range, where the yield strength is sufficiently lower than the ambient yield strength, and the steel is restored by promoting the precipitation of fine carbides. The latter requires a minimum temperature of at least 200°C to ensure the transformation of retained austenite. If the final recovery heat treatment temperature is above 600°C, undesirable martensite recrystallization may occur. The intermediate austenitizing and quenching steps have already produced a martensitic structure (single-phase steel) in which carbon exists in a supersaturated solid solution. During the final recovery heat treatment, carbon combines with iron and any other carbide-forming alloying elements, such as chromium and molybdenum, and precipitates as carbides. These carbides stabilize the microstructure. These carbides are also believed to minimize embrittlement caused by strain aging. Unbound by any theory, it is believed that during aging, the large amount of carbon in the solid solution, such as in untempered materials like the aforementioned cold-drawn and then quenched steel, generates very strong Cotillard atmospheres around dislocations. These atmospheres disrupt dislocation movement, leading to material embrittlement. As a result of the final recovery heat treatment according to the invention, by reducing dislocation density and promoting carbide precipitation, it is assumed that this adverse phenomenon does not occur, or at least is significantly reduced. Therefore, embrittlement caused by strain aging can also be reduced.

[0049] Following the recovery process, the tubular components manufactured according to the present invention typically undergo finishing operations, such as straightening and end forming. Therefore, in one embodiment, the method further includes a cold forming step e) of cold forming the tubular product from step c), particularly its ends, optionally preceded by a straightening step d) of straightening the recovered tubular product from step c). It has been found that after applying this stretching, the tensile strength remains at the same level or slightly increases compared to cold-drawn and then quenched tubular products, while the ductility value is less affected and remains at a high level. Cold-drawn, quenched, and then tempered steel tubes exhibit a similar increase in strength upon stretching as cold-drawn and then quenched tubular products, although the strength level is lower and the increase in strength is less pronounced.

[0050] Element

[0051] The steel composition used in the method according to the invention preferably comprises, by weight percent, excluding Fe and unavoidable impurities.

[0052] C: 0.04-0.15;

[0053] Mn: 0.90-1.60;

[0054] Si: 0.10-0.50;

[0055] Cr: 0.05-0.80;

[0056] Al 0.01-0.50;

[0057] N 0.0035-0.0150.

[0058] Preferably, the composition contains one or more elements that form carbides, nitrides, or carbonitrides, in an amount sufficient to bind N in the form of (carbon)nitrides. Examples of these elements, besides Al, include V, Ti, and Nb. Preferably, these elements satisfy the equation [%Al] / 1.9 + [%Ti] / 3.4 + [%V] / 3.6 + [%Nb] / 6.6 ≥ [%N], where % is by weight. Aging is associated with the diffusion of interstitial elements (primarily carbon), but nitrogen diffusion also plays a role in aging. The above equation ensures that residual nitrogen is bound in the form of nitrides.

[0059] Additionally, this ingredient may contain optional elements, expressed in weight percent.

[0060] Mo: 0-0.50;

[0061] Ni: 0-0.50;

[0062] Cu 0-0.25;

[0063] V 0-0.40;

[0064] Nb 0-0.20;

[0065] Ti 0-0.10;

[0066] B 0-0.005.

[0067] Ca 0-0.005.

[0068] If present, the amount of unavoidable impurities is

[0069] As 0-0.05;

[0070] Sb 0-0.05;

[0071] Sn 0-0.05;

[0072] Pb 0-0.05;

[0073] Bi 0-0.005;

[0074] S 0-0.015;

[0075] P 0-0.025.

[0076] The remainder of the composition is iron (Fe).

[0077] Advantageous

[0078] [%Sn]+[%Sb]+[%Pb]+[%As]+[%Bi]≤0.10%;

[0079] and / or

[0080] 0.3≤Ceq≤0.7, where

[0081] Ceq = [%C] + [%Mn] / 6 + ([%Cr] + [%Mo] + [%V]) / 5 + ([%Ni] + [%Cu]) / 15, and / or

[0082] [%AI] / 1.9+[%Ti] / 3.4+[%V] / 3.6+[%Nb] / 6.6≥[%N], where [%] is by weight, and the preferred steel composition satisfies all three equations.

[0083] The steel composition (preferably low-carbon steel composition considering weldability), and preferably the (microalloyed) steel composition, contains one or more elements that form carbides, nitrides or carbonitrides, as described above, to ensure that N is combined in the form of (carbon)nitrides to take advantage of the effect of (carbon)nitrides on grain refinement.

[0084] This composition is very low in alloying elements, especially since it does not require a minimum amount of molybdenum and / or vanadium. This composition ensures a minimum N content associated with nitride-forming elements such as Al, Nb, Ti, and V, so that sufficient (carbon)nitrides can be present during austenitization to improve grain size control.

[0085] The following explanations are given regarding the various elements in the low-carbon microalloy composition. The ranges in parentheses are preferred ranges and represent a balance between cost and beneficial effects on structure, process, and / or performance.

[0086] Carbon (C): 0.04-0.15 (0.06-0.12)

[0087] Carbon is needed to strengthen steel by precipitating very fine carbides in the final stage of the transformation; however, excessive carbon will cause a significant increase in internal stress during quenching, making welding impractical or completely impossible. Therefore, the carbon content is 0.04-0.15%, preferably 0.06-0.12%.

[0088] Manganese (Mn): 0.90-1.60 (1.00-1.40)

[0089] Mn is an important alloying element with various functions. During austenite cooling, it lowers the transformation temperature from austenite to ferrite; therefore, during normalizing, it increases the nucleation-to-growth ratio and ultimately leads to refined grain size.

[0090] Conversely, during quenching, Mn increases the hardenability of the material, ensuring a fully martensitic structure is obtained over a larger cross-section. However, excessive Mn can lead to an undesirable large amount of retained austenite after quenching. Furthermore, Mn is known to reduce intergranular fracture strength, and therefore, excessive Mn can affect impact toughness. Therefore, the Mn content is 0.90-1.60, preferably 1.00-1.40.

[0091] Silicon (Si): 0.10-0.50 (0.20-0.35)

[0092] The presence of silicon (Si) is for deoxidizing the steel. However, excessive amounts negatively impact toughness. Furthermore, Si increases susceptibility to temper embrittlement by enhancing P segregation at grain boundaries. Therefore, the Si content is 0.10-0.50%, preferably 0.20-0.35%.

[0093] Chromium (Cr): 0.05-0.80 (0.30-0.60)

[0094] Cr effectively improves the hardenability of steel and, as a carbide former, allows bainite to form during continuous cooling. Very high Cr content reduces hardening effectiveness and unnecessarily increases steelmaking costs. Therefore, the Cr content is 0.05-0.80%, preferably 0.30-0.60%.

[0095] Aluminum (Al): 0.01-0.50 (0.015-0.030)

[0096] Al is a deoxidizing element and a nitride formation product. A minimum amount is required to ensure sufficient deoxidation and allow for the binding of residual nitrogen. Excessive amounts may result in a large number of nonmetallic inclusions. Therefore, the Al content is 0.01-0.50, preferably 0.015-0.030.

[0097] Nitrogen (N): 0.0035-0.0150 (0.006-0.010)

[0098] On the one hand, nitrogen (N) is an unavoidable residual element in steelmaking. However, a small amount is actually preferable because N can be used to control grain size by promoting the precipitation of nitride and (carbon)nitride forming elements (e.g., Al, Ti, Nb, or V). Therefore, a minimum content is required for grain size control. On the other hand, free N (in interstitial solid solutions) needs to be avoided because it increases the aging effect and promotes the formation of Lüders bands, ultimately reducing the cold formability of the product. Therefore, the N content is 0.0035-0.0150, preferably 0.006-0.010. According to the stoichiometric formula [%Al] / 1.9+[%Ti] / 3.4+[%V] / 3.6+[%Nb] / 6.6≥[%N], the preferred formula is [%Al] / 1.9+[%Ti] / 3.4+[%V] / 3.6+[%Nb] / 6.6≥1.1[%N], where [%] is by weight, and the available combined amounts of Al, Ti, Nb and V need to be sufficient to bind any residual N.

[0099] Molybdenum (Mo): 0-0.50 (0.10-0.20)

[0100] Mo is highly effective in improving the hardenability of steel, and as a strong carbide former, it allows bainite to form during continuous cooling. Furthermore, Mo improves tempering resistance, thus allowing for the maintenance of desirable strength levels while increasing toughness and reducing internal stress. However, excessive Mo content is undesirable for cost reasons, and also because it lowers the martensitic transformation temperature and may lead to a greater amount of retained austenite during quenching. Therefore, the Mo content is 0-0.50%, preferably 0.10-0.20%.

[0101] Nickel (Ni): 0-0.50 (0-0.20)

[0102] Ni is an austenite stabilizer that allows for refinement of ferrite grain size by lowering the transformation temperature in a manner similar to Mn. Ni also improves toughness.

[0103] However, Ni increases the amount of retained austenite during quenching and therefore needs to be limited. Furthermore, Ni is generally expensive, and similar effects can be achieved through other means. Therefore, the Ni content is 0-0.50%, preferably 0-0.20%.

[0104] Copper (Cu): 0-0.25 (0-0.20)

[0105] Cu slightly improves hardenability and is inevitably present in scrap steel. However, a large amount of Cu can cause hot brittleness; this reduces the surface quality of hot-finished products (increases roughness) and can also lead to serious and irreparable defects. Therefore, the Cu content is limited to 0-0.25%, preferably 0-0.20%.

[0106] Vanadium (V): 0-0.40 (0-0.10)

[0107] V is a strong carbide and nitride formant, and its presence is intended to improve hardenability, achieve precipitation hardening, and refine austenite grain size. Its effectiveness as a refining element is limited by its solubility in austenite at higher temperatures. Therefore, the V content is 0-0.20, preferably 0-0.10.

[0108] Niobium (Nb): 0–0.20 (0–0.05) and titanium (Ti): 0–0.10 (0–0.05) are both strong carbide and nitride formants. They function similarly to niobium (V) in controlling austenite grain size, and are more effective than niobium (V) due to their lower solubility in austenite. Titanium is more effective than Nb at higher temperatures (above approximately 1100 °C), while Nb typically results in finer precipitate dispersions, thus allowing for the finest prior austenite grain size.

[0109] Tin (Sn): 0-0.05 (0-0.03), Antimony (Sb): 0-0.05 (0-0.01), Arsenic (As): 0-0.05 (0-0.03), Lead (Pb): 0-0.05 (0-0.01) and Bismuth (Bi): 0-0.005.

[0110] These unavoidable impurities negatively affect the toughness of steel. Therefore, their content is limited. Advantageously, [%Sn]+[%Sb]+[%Pb]+[%As]+[%Bi]≤0.10%, where [%] is by weight%.

[0111] Phosphorus (P): 0-0.025, preferably 0-0.02; Sulfur (S): 0-0.015, preferably 0-0.005. P and S are also unavoidable elements, and their contents are limited as follows.

[0112] Calcium (Ca) 0-0.005; REM: 0-0.005.

[0113] Ca and rare earth metals (REMs) can be used for inclusion control. Ca and REMs form complex oxides with Al and Mg. These complex oxides have lower melting points. They promote flotation, leading to a reduction in inclusion content. Furthermore, the shape of residual nonmetallic inclusions becomes spherical, reducing their embrittlement effect. Although most of the Ca and Mg remain in the slag thus formed, a residual amount of Ca is unavoidable in the treated steel.

[0114] Boron (B) 0-0.005 (0-0.0005)

[0115] Boron (B) increases hardenability to approximately 0.0020% (depending on the actual carbon content). Boron can also negatively impact toughness by promoting the formation of boron nitride, the precipitation of which can only be suppressed by the action of Ti exceeding approximately 3.4 times the amount of N. Intentional addition of B is not strictly required to achieve the desired hardenability; furthermore, especially without the addition of Ti, the B content should be limited to ensure optimal toughness.

[0116] Furthermore, it is advantageous to impose a constraint on hardenability measured according to carbon equivalent (IIW formula): 0.3 ≤ Ceq ≤ 0.7, where

[0117] Ceq = [%C] + [%Mn] / 6 + ([%Cr] + [%Mo] + [%V]) / 5 + ([%Ni] + [%Cu]) / 15, where [%] is by weight.

[0118] Steelmaking process and inclusion content

[0119] Steelmaking processes are typically carried out using clean operating conditions to achieve very low sulfur and phosphorus contents. Low S and P contents are crucial for achieving mechanical properties, particularly ductility and toughness.

[0120] The steel is produced according to clean practices, ensuring very low levels of non-metallic inclusions. Therefore, it is advantageous to apply inclusion levels according to ASTM E45 standard – Worst-Case Field Method (Method A):

[0121] Inclusion type Thin Heavy A 0.5 1 B 1.5 1 C 0 0 D 1.5 0.5

[0122] Furthermore, cleanroom practices allow for the attainment of extremely large inclusions with a size of 30 ppm or smaller. Therefore, the total oxygen content is limited to 20 ppm.

[0123] As an example of extremely clean practices in secondary metallurgy, the use of bubbling inert gas in a ladle furnace is introduced. The bubbling gas forces non-metallic inclusions and impurities to float on the molten steel. Producing a fluid slag capable of absorbing these inclusions and impurities, and adding silicon and calcium to the molten steel to alter the size and shape of the inclusions, facilitates the preparation of microalloyed low-carbon steel with the desired low inclusion content.

[0124] microstructure

[0125] The hollow core after optional normalizing treatment as described above preferably has a fine-grained microstructure, which consists of ferrite (polygonal, acicular and / or Wiedmanstattern), bainite, preferably >20% (area) of bainite and preferably <5% of pearlite.

[0126] The microstructure is homogeneous to reduce the segregation of residual elements that is unavoidable during the casting process. The hollow structure has good tensile hardening ability to ensure the quality of the cold-drawn tube, especially its mechanical properties.

[0127] The intermediate austenitizing and quenching steps, which are part of the method according to the invention (which are carried out before the final cold drawing in a multi-stage cold drawing process), transform the microstructure of the hot-rolled tube undergoing cold drawing to a predominantly martensitic structure, which consists of martensite and small amounts of bainite and ferrite, with bainite preferably equal to or less than 20% and ferrite preferably equal to or less than 5%.

[0128] The final microstructure obtained by the final recovery heat treatment after cold drawing comprises 80% or more tensile hardened and recovered martensite and lower bainite, and a small amount of coarse bainite and ferrite, preferably as low as possible. Preferably, the microstructure comprises 90% martensite and lower bainite (determined by hardness (HRC) > 27 + 58 x [%C] measured after quenching and before further cold drawing), more preferably 95% or more martensite and lower bainite (determined by hardness (HRC) > 29 + 59 x [%C] measured after quenching and before further cold drawing).

[0129] Advantageously, the final microstructure has a grain size number (ASTM E112) of 9 or higher, preferably 10 or higher. The higher the grain size number, the finer the microstructure.

[0130] characteristic

[0131] The method according to the invention allows for the manufacture of tubular products having one or more of the following mechanical properties:

[0132] Yield strength (YS): ≥896MPa (130ksi);

[0133] Tensile strength (TS): ≥1103MPa (160ksi);

[0134] Total elongation (A5D): ≥9%;

[0135] DBTT: ≤-60℃;

[0136] Bursting: Exhibits significant ductility (>50%) at -60°C.

[0137] Yield strength, tensile strength, and elongation are determined according to ASTM E8.

[0138] The burst test is performed through the sealed end of the pipe, for example by welding a flat steel plate or flange to the end of the pipe. Internal pressure is then applied to the pipe using a suitable fluid until it fails. The test can be conducted in a temperature-controlled chamber at the desired temperature, or by adjusting the fluid temperature.

[0139] Advantageously, the resulting product has a combination of at least two of the above-mentioned properties, and more preferably a combination of all of the above-mentioned properties.

[0140] Example

[0141] The microalloy steel compositions listed in Table 1 were prepared under clean practices and cast into round billets with a diameter of approximately 148 mm. The billets underwent a process including the following steps: induction heating to a temperature of 870°C (i.e., above Ac3), piercing, hot rolling using a floating mandrel technique, intermediate reheating and final stretching and reduction rolling, cooling, and furnace normalizing.

[0142] Table 1 Chemical Composition

[0143]

[0144]

[0145] Example 1 (Comparison)

[0146] Hot-rolled hollow tubes, thus obtained from composition A, with an outer diameter (OD) of 42.4 mm and a wall thickness (WT) of 2.9 mm, were cold-drawn to a size of 30 x 1.85 mm (OD x WT) in two drawing operations. They were then heat-treated in the range of 900-1030°C and quenched using a 10-water spray. The resulting tubular product was then subjected to a stretching to 25 mm OD simulated by cold forming (mandrelless cold drawing) to mimic the effect of a finishing forming operation. No recovery treatment was applied.

[0147] Example 2 (Comparison)

[0148] In another embodiment, in addition to quenching and tempering heat treatment at 400°C before simulated stretching (mandrelless cold drawing), the same component A is also used to manufacture tubes under the same conditions according to a similar process.

[0149] Table 2 below lists the properties measured using the corresponding standards ASTM E8 and ASTM E10 for the products obtained before simulation (“as is”) and for the products after cold working (“stretched”) during simulation straightening and stretching in these examples.

[0150] Table 2

[0151]

[0152]

[0153] As can be seen from the comparison of these examples, Example 1 (stretching-quenching-re-stretching) is superior to Example 2 (stretching-quenching and tempering-re-stretching) in almost every respect, except for the reduction in elongation (A 5D).

[0154] Example 3 (Invention).

[0155] Following the process outlined in Example 1, tubular products are made from steel composition B, but intermediate austenitization and quenching are introduced prior to final cold drawing, followed by a final recovery heat treatment at 430°C after final cold drawing. Austenitization is performed by induction heating to 950°C and a soaking time of 5 seconds, followed by quenching to room temperature using external water spray (cooling rate exceeding 50°C / s). The hollow dimensions after hot rolling are 48.3 x 3.4 mm (OD x WT). The final dimensions of the cold-drawn product are 35 x 2 mm.

[0156] The obtained product has the following metallurgical and mechanical properties:

[0157] UTS: 1248MPa (182ksi);

[0158] YS: 1228MPa (178ksi);

[0159] Total elongation: 10%;

[0160] Grain size class number (ASTM E112): 13;

[0161] Hardness HV 10 394;

[0162] Bursting at ambient temperature: 1731-1738 bar (25.1-25.2 ksi);

[0163] -69℃ burst fracture appearance: >50% shear area.

[0164] Example 4 (Invention)

[0165] Following the process outlined in Example 1, tubular products are made from steel composition C, but intermediate austenitization and quenching are introduced again before the final cold drawing, followed by a final recovery heat treatment at 400°C after the final cold drawing. Austenitization is performed by induction heating to 900-1030°C, followed by quenching to room temperature using external water spray (cooling rate exceeding 50°C / s). The hollow dimensions after hot rolling are 38.0 x 2.9 mm. A 29% reduction is achieved in the first cold drawing, resulting in a hollow dimension of 34.5 x 2.25 mm. A 26% reduction is achieved after the second cold drawing, resulting in a final cold-drawn product dimension of 30 x 1.92 mm.

[0166] The resulting product has the following metallurgical and mechanical properties:

[0167] UTS: 1262MPa (183ksi);

[0168] YS: 1172MPa (170ksi);

[0169] Total elongation: 16.8%;

[0170] Grain size class number (ASTM E112): 11-12;

[0171] Hardness HV 10 :428;

[0172] Bursting at ambient temperature: average 1972 bar (28.6 ksi);

[0173] -60℃ burst fracture appearance: >50% of shear area.

[0174] Example 5 (Comparison)

[0175] Using steel composition D, Example 1 was repeated, except that cold drawing included a single drawing step, followed by a quenching step. The hollow dimensions after hot rolling were 38.1 x 2.7 mm. After a single cold drawing step, the hollow dimensions were reduced by 32%, to 33.2 x 2.08 mm.

[0176] This product has the following metallurgical and mechanical properties:

[0177] UTS: 1277MPa (183ksi);

[0178] YS: 992MPa (170ksi);

[0179] Total elongation: 15%;

[0180] Grain size class number (ASTM E112): 11-12;

[0181] Hardness HV 10 :413;

[0182] Example 6 (Comparison)

[0183] Example 2 was repeated using steel composition E, except that cold drawing included a single drawing, followed by quenching and tempering at 380°C. The hollow dimensions after hot rolling were 38.1 x 2.7 mm. After a single cold drawing step, the hollow dimensions were reduced by 33%, to 32 x 2.15 mm.

[0184] This product has the following metallurgical and mechanical properties:

[0185] UTS: 1084MPa (183ksi);

[0186] YS: 911MPa (170ksi);

[0187] Total elongation: 13%;

[0188] Grain size class number (ASTM E112): 11-12;

[0189] Hardness HV 10 :not applicable

[0190] The tubular products from Examples 4-6 underwent stretching simulated by cold forming (mandrel-less cold drawing), resulting in a 17% reduction in area. Table 3 below summarizes the results, where "as is" refers to the tubular products manufactured according to these examples, and "stretched" refers to the tubular products after simulated stretching.

[0191] Table 3 Experimental Data Examples 4-6

[0192]

[0193] As can be seen from the table, the tensile strength of Example 4 according to the invention is higher than that of Example 6 when stretched. This also applies to elongation. Although the strength of Example 5 is higher than that of Example 4, the elongation value of Example 4 according to the invention is higher for both the original tubular product and the stretched product. Therefore, the advantageous combination of strength and ductility properties of the product manufactured according to the invention remains unchanged during cold working, allowing for proper product finishing.

[0194] Furthermore, it has been found that the dislocation density in Example 4 according to the present invention is significantly lower than that in Example 5, such as... Figure 1 As shown, Figure 1 The average microstrain ε is shown (note that the dislocation density ρ is related to ε). 2 Proportional (ρ=A*(ε) 2 (where A is a material constant). It can be seen that the dislocation density in this invention is much lower than that in the embodiment of Example 5. Furthermore, during cold working (i.e., stretching), the dislocation density in this invention remains almost the same, while the material of Example 5 exhibits microstrain and therefore a significant increase in dislocation density. The increase in dislocation density increases hardness and strength, but reduces ductility and toughness properties. It can be assumed that stretching has a lesser effect on the strength and elongation of the steel pipe according to the invention, and therefore on its formability, than it does on the material of Example 5.

[0195] airbag inflator pressure vessel

[0196] Seamless tubes manufactured according to the present invention are cut to a certain length and then cold-formed using known techniques, such as flanging, forging, etc., to the desired shape. Alternatively, welded tubes processed according to the present invention can be used. End caps and diffusers are welded to each end of the cold-formed tube using known techniques, such as friction welding, arc welding, and laser welding, thereby producing airbag inflator pressure vessels.

[0197] This invention is also embodied in the following provisions:

[0198] 1. A method for manufacturing pipes from steel components, particularly for pipes used in pressure vessels of gas storage pressurizers, comprising the following steps:

[0199] a) The production of steel pipes from steel components includes at least one hot rolling or hot forming process;

[0200] b) The steel pipe is subjected to a cold drawing process to obtain the required dimensions, wherein the cold drawing process includes at least two drawing operations and an intermediate austenitizing and quenching step prior to the final drawing operation of the cold drawing process.

[0201] c) After the final drawing in the cold drawing process, the cold-drawn steel pipe undergoes a final recovery heat treatment at a temperature range of 200-600℃. The steel composition, by weight %, comprises...

[0202] C: 0.04-0.15;

[0203] Mn: 0.90-1.60;

[0204] Si: 0.10-0.50;

[0205] Cr: 0.05-0.80;

[0206] Al 0.01-0.50;

[0207] N 0.0035-0.0150;

[0208] The remainder consists of iron and unavoidable impurities.

[0209] 2. The method according to Clause 1, wherein, after the intermediate austenitizing and quenching steps, the total area of ​​one or more drawing operations is reduced by at least 10%.

[0210] 3. The method according to Clause 2, wherein, after the intermediate austenitizing and quenching steps, the total area of ​​one or more drawing operations is reduced by at least 15%.

[0211] 4. The method according to Clause 2, wherein, after the intermediate austenitizing and quenching steps, the total area of ​​one or more drawing operations is reduced by at least 20%.

[0212] 5. The method according to Clause 1, wherein the intermediate austenitizing and quenching steps are performed between the penultimate and final drawing steps of the cold drawing process.

[0213] 6. The method according to Clause 1, wherein the intermediate austenitizing and quenching steps include quenching at a quenching rate of at least 50°C / s.

[0214] 7. The method according to Clause 1, wherein step a) of producing the steel pipe comprises the following sub-steps: preparing a steel composition, casting the composition into a steel billet, piercing the steel billet at an elevated temperature, and hot rolling the pierced steel billet in at least one hot rolling process, optionally including an intermediate reheating step between two hot rolling processes to a temperature above Ac3.

[0215] 8. The method according to Clause 1, wherein the shrinkage in each hot rolling process is at least 3%.

[0216] 9. The method according to Clause 1, wherein, in step b), the intermediate austenitizing and quenching steps include heating to a temperature above Ac3.

[0217] 10. The method according to Clause 9, wherein, in step b), the intermediate austenitizing and quenching steps include heating in the range of 880-1050°C.

[0218] 11. The method according to Clause 1, wherein the method further comprises normalizing heat treatment, which includes heat treating the hot-rolled tube at a temperature above Ac3 after hot rolling or normalizing it in a final hot rolling process at a temperature above Ar3.

[0219] 12. The method according to Clause 11, wherein the normalizing heat treatment comprises heat treating the hot-rolled tube at a temperature between Ac3 and 1000°C after hot rolling.

[0220] 13. The method according to Clause 11, wherein the normalizing heat treatment comprises normalizing rolling in a final hot rolling process at a temperature between Ar3 and the grain coarsening temperature.

[0221] 14. The method according to Clause 13, wherein the normalizing heat treatment comprises normalizing rolling in a final hot rolling process at a temperature between Ar3 and 1050°C.

[0222] 15. The method according to Clause 13, wherein the normalizing heat treatment comprises normalizing rolling in the final hot rolling process at a temperature in the range of 850-1000°C.

[0223] 16. The method according to Clause 1 further includes a cold forming step e) of cold forming the tubular product from step c), particularly cold forming its ends, optionally preceded by a straightening step d) of straightening the recovered tubular product from step c).

[0224] 17. The method according to Clause 1 further includes one or more elements that form carbides, nitrides or carbonitrides, in an amount sufficient to bind N in the form of (carbon)nitrides.

[0225] 18. The method according to Clause 1, wherein the steel composition further comprises one or more optional elements Mo: 0-0.50;

[0226] Ni: 0-0.50;

[0227] Cu: 0-0.25;

[0228] V 0-0.40;

[0229] Nb 0-0.20;

[0230] Ti 0-0.10;

[0231] B 0-0.005;

[0232] Ca 0-0.005.

[0233] 19. The method according to Clause 1, wherein unavoidable impurities include

[0234] As 0-0.05;

[0235] Sb 0-0.05;

[0236] Sn 0-0.05;

[0237] Pb 0-0.05.

[0238] Bi 0-0.005;

[0239] S 0-0.015;

[0240] P 0-0.025.

[0241] 20. The method according to Clause 19, wherein [%Sn]+[%Sb]+[%Pb]+[%As]+[%Bi]≤0.10%, where [%] is by weight%.

[0242] 21. The method described in accordance with Clause 18, wherein

[0243] 0.3≤Ceq≤0.7, where

[0244] Ceq=[%C]+[%Mn] / 6+([%Cr]+[%Mo]+[%V]) / 5+([%Ni]+[%Cu]) / 15.

[0245] 22. The method described in accordance with Clause 18, wherein

[0246] [%AI] / 1.9 + [%Ti / 3.4] + [%V] / 3.6 + [%Nb] / 6.6 ≥ [%N], where [%] is by weight.

[0247] 23. The method according to Clause 1, wherein the steel component comprises, in weight percent,

[0248] C: 0.06-0.12;

[0249] Mn: 1.00-1.40;

[0250] Si: 0.20-0.35;

[0251] Cr: 0.30-0.60;

[0252] Al 0.015-0.030;

[0253] N 0.006-0.010.

[0254] 24. The method according to Clause 22, wherein [%AI] / 1.9+[%Ti] / 3.4+[%V] / 3.6+[%Nb] / 6.6≥1.1[%N], where [%] is by weight%.

[0255] 25. The method according to Clause 1, wherein the resulting tube has one or more of the following characteristics:

[0256] Yield strength (YS): ≥896MPa (130ksi);

[0257] Tensile strength (TS): ≥1103MPa (160ksi);

[0258] Total elongation (A5D): ≥9%;

[0259] DBTT: ≤-60℃;

[0260] Bursting: Exhibits >50% ductility at -60℃;

[0261] 26. The method according to clause 25, wherein the resulting tube has the following characteristics:

[0262] Yield strength (YS): ≥896MPa (130ksi);

[0263] Tensile strength (TS): ≥1103MPa (160ksi);

[0264] Total elongation (A5D): ≥9%;

[0265] DBTT: ≤-60℃;

[0266] Bursting: Exhibits >50% ductility at -60℃;

[0267] 27. The method according to Clause 1, wherein the resulting tube has a predominantly martensitic microstructure comprising 80% or more martensite and lower bainite, with the remainder being coarse bainite and ferrite.

[0268] 28. The method according to Clause 27, wherein the resulting tube has a predominantly martensitic microstructure comprising equal to or greater than 90% martensite and lower bainite.

[0269] 29. The method according to Clause 27, wherein the resulting tube has a predominantly martensitic microstructure comprising 95% or more martensite and lower bainite.

[0270] 30. The method according to Clause 27, wherein the resulting tube has a predominantly martensitic microstructure containing less than 5% ferrite.

[0271] 31. The method according to Clause 1, wherein the number of grain size levels in the resulting tube is 9 or higher, preferably 10 or higher.

[0272] 32. Automotive parts, particularly airbag inflator pressure vessels, including tubes of a certain length manufactured in accordance with Clause 1.

Claims

1. A method for manufacturing pipes from steel components, comprising the following steps: a) The production of steel pipes from steel components includes at least one hot rolling or hot forming process; b) The steel pipe is subjected to a cold drawing process to obtain the required dimensions, wherein the cold drawing process includes at least two drawing operations and an intermediate austenitizing and quenching step prior to the final drawing operation of the cold drawing process. c) After the final drawing in the cold drawing process, the cold-drawn steel pipe undergoes a final recovery heat treatment at a temperature range of 200-600°C, wherein the steel composition comprises, in weight %... C：0.04-0.15； Mn: 0.90-1.60; Si: 0.10-0.50; Cr：0.05-0.80； Al 0.01-0.50； N 0.0035-0.0150； Mo: 0-0.50; Ni: 0-0.50; Cu: 0-0.25; V 0-0.40； Nb 0-0.20; Ti 0-0.10; B 0-0.005； Ca 0-0.005; As 0-0.05; Sb 0-0.05; Sn 0-0.05; Pb 0-0.05; Bi 0-0.005; S 0-0.015； P 0-0.025； The remainder consists of iron and unavoidable impurities.

2. The method according to claim 1, characterized in that, The total area reduced by at least 10% after one or more drawing steps following intermediate austenitization and quenching.

3. The method according to claim 1 or claim 2, characterized in that, The intermediate austenitization and quenching steps are performed between the penultimate and final drawing steps of the cold drawing process.

4. The method according to claim 1 or claim 2, characterized in that, The intermediate austenitizing and quenching steps include quenching at a quenching rate of at least 50 °C / s.

5. The method according to claim 1 or claim 2, characterized in that, The steps for producing steel pipes a) include the following sub-steps: preparing steel components, casting the components into steel billets, piercing the steel billets at elevated temperatures, and hot-rolling the pierced steel billets in at least one hot-rolling process.

6. The method according to claim 1 or claim 2, characterized in that, The shrinkage in each hot rolling process is at least 3%.

7. The method according to claim 1 or claim 2, characterized in that, In step b), the intermediate austenitizing and quenching steps include heating to a temperature above Ac3.

8. The method according to claim 1 or claim 2, characterized in that, The method also includes normalizing heat treatment, which includes heat treating the hot-rolled tube at a temperature above Ac3 after hot rolling or normalizing it in the final hot rolling process at a temperature above Ac3.

9. The method according to claim 8, characterized in that, Normalizing heat treatment involves heat-treating the hot-rolled tube at a temperature between Ac3 and 1000°C after hot rolling.

10. The method according to claim 8, characterized in that, Normalizing heat treatment involves normalizing rolling in the final hot rolling process at a temperature between Ar3 and grain coarsening temperature.

11. The method according to claim 1 or claim 2, further comprising a cold forming step e) of cold forming the tubular product from step c).

12. The method according to claim 1 or claim 2, characterized in that, [%Sn] + [%Sb] + [%Pb] + [%As] + [%Bi] ≤ 0.10%, where [%] is weight.

13. The method according to claim 1 or claim 2, characterized in that... 0.3 ≤ Ceq ≤ 0.7, where Ceq = [%C] + [%Mn] / 6 + ([%Cr]+[%Mo]+[%V]) / 5+([%Ni]+[%Cu]) / 15, and / or [%AI] / 1.9 + [%Ti / 3.4] + [%V] / 3.6 + [%Nb] / 6.6 ≥ [%N], where [%] is weight.

14. The method according to claim 1 or claim 2, characterized in that, In the steel composition, expressed as a percentage by weight, C：0.06-0.12； Mn: 1.00-1.40; Si: 0.20-0.35; Cr：0.30-0.60； Al 0.015-0.030； N 0.006-0.010。 15. The method according to claim 1 or claim 2, characterized in that... [%AI] / 1.9 + [%Ti] / 3.4 + [%V] / 3.6 + [%Nb] / 6.6 ≥ 1.1 [%N], where [%] is by weight.

16. The method according to claim 1 or claim 2, characterized in that, The resulting tube has one or more of the following characteristics: Yield strength (YS): ≥ 896MPa (130ksi); Tensile strength (TS): ≥ 1103MPa (160ksi); Total elongation (A5D): ≥ 9%; Among them, YS, TS, and A 5D are determined according to ASTM E8. DBTT: ≤ -60℃; Bursting: Exhibits >50% ductility at -60℃.

17. The method according to claim 1 or claim 2, characterized in that, The resulting tube has a microstructure that is mainly martensite, containing 80% or more martensite and lower bainite, with the remainder being coarse bainite and ferrite.

18. The method according to claim 1 or claim 2, characterized in that, The resulting tube has a grain size grade number (ASTM E112) of 9 or higher.

19. An automotive component, comprising a tube of a certain length manufactured according to any one of the preceding claims.

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