Steel, method for producing same, and method for processing the steel

Abstract

The invention relates to a steel, in particular a linepipe steel, which has the following composition: 0.02 - 0.1 wt.% carbon, 0.01 - 0.25 wt.% silicon, 1.60 - 1.8 wt.% manganese, ≤ 0.05 wt.% aluminum, 0.01 - 0.40 wt.% copper, 0.01 - 0.30 wt.% nickel, 0.01 - 0.30 wt.% chromium, 0.005 - 0.025 wt.% niobium, 0.005 - 0.030 wt.% titanium, ≤ 0.0050 wt.% sulfur, ≤ 0.0050 wt.% calcium, ≤ 0.010 wt.% nitrogen, ≤ 0.02 phosphorus, ≤ 0.010 vanadium, 0.001 - 0.10 wt.% molybdenum, and a remainder of: iron and production-related impurities. The invention further relates to a method for producing the steel and a method for processing the steel.

Description

[0001] Description:

[0002] Steel and a process for its manufacture as well as a process for processing the steel

[0003] The invention relates to a steel, in particular linepipe steel, which is suitable for processing by welding, and to a method for producing and a method for processing the steel.

[0004] In the case of such steels, known from previous use, the mechanical properties deteriorate due to energy input during welding in the heat-affected zone of the welding process, with a significant reduction in toughness in particular. This also applies to known linepipe steels, which were developed for the manufacture of pipelines for the transport of oil, gas, and other media and are characterized by particularly high toughness, especially at low temperatures, as well as by good weldability.

[0005] The invention is based on the objective of creating a steel of the type mentioned above which, after processing by welding, has better mechanical properties than the known steels.

[0006] According to the invention, this problem is solved by a steel of the following composition:

[0007] 0.02 - 0.1 wt.% C, 0.01 - 0.25 wt.% Si, 1.60 - 1.80 wt.% Mn,

[0008] < 0.05 wt.% AI,

[0009] 0.01 - 0.40 wt.% Cu,

[0010] 0.01 - 0.30% by weight Ni, 0.01 - 0.30% by weight Cr, 0.005 - 0.025% by weight Nb, 0.005 - 0.030% by weight Ti,

[0011] < 0.0050 wt.% S,

[0012] < 0.0050 wt.% Ca,

[0013] < 0.010 wt% N,

[0014] < 0.02 wt.% P, < 0.010 wt.% V, 0.001 - 0.10 wt.% Mo,

[0015] Remainder: Iron and manufacturing-related impurities.

[0016] It has surprisingly been shown that the combination of large austenite grains, martensite-austenite components (MAs) and / or coarse Nb / Ti primary precipitates in the heat-affected area is decisive for the comparatively poor mechanical properties, especially the low toughness, of the known steels.

[0017] The problem of low toughness in the heat-affected zone is solved according to the invention by the comparatively low carbon and silicon contents. These low contents result in low martensite / austenite content (MAs), which has an embrittlement effect in steel. The similarly low niobium and titanium contents reduce the formation of coarse primary precipitates, which preferably have a size > 500 nm, particularly > 1.5 µm, to a certain extent, and produce comparatively small niobium-titanium carbonitride precipitates that contribute to grain boundary pinning. The niobium-titanium carbonitride precipitates preferably have sizes < 500 nm, preferably < 200 nm.

[0018] In a further embodiment of the invention, silicon deoxidation is carried out at Si contents of < 0.1 wt.%, preferably < 0.05 wt.%, and particularly preferably < 0.03 wt.%, which makes it possible to provide the steel with particularly low Si and Al contents and thus achieve a further reduction in the MA content. It has surprisingly been found that the steel exhibits exceptionally high toughness in the heat-affected range of welding at such low Si contents.

[0019] The steel according to the invention exhibits comparatively high toughness, particularly impact strength, after welding. Furthermore, it has been shown that, thanks to its chemical and microstructural properties, the steel exhibits a low number of martensite / austenite components and a small austenite grain size after welding, in a heat-affected zone of welding, i.e., in a zone where welding affects the mechanical properties of the steel, possibly even with high heat input, compared to known steels.

[0020] The steel is preferably a micro-alloyed steel produced by thermomechanical rolling, in particular in the present case as steel sheet.

[0021] In one embodiment of the invention, the steel has a yield strength Rt0.5 at room temperature according to the standard DIN EN ISO 6892-1:2019 of 360 to 600 MPa, preferably of 400 to 550 MPa. Advantageously, the steel has a tensile strength R at room temperature. m according to the standard DIN EN ISO 6892-1 :2019 from 460 to 800 MPa, preferably from 500 - 760 MPa.

[0022] Preferably the steel has a Vickers hardness (HV10) according to the standard DIN EN ISO 6507-1 :2018 of 100 to 300, preferably of 150 to 250.

[0023] In one embodiment of the invention, the shear fracture fraction of the steel in the Battelle Drop Weight Tear Test (BDWTT) according to the standard API RP 5L3:2021 at -15 °C is > 75%, preferably > 85%.

[0024] Advantageously, the steel has a) an impact energy of at least 340 J, preferably at least 350 J, at a test temperature of -30 °C, b) an impact energy of at least 340 J, preferably at least 350 J, at a test temperature of -50 °C, c) an impact energy of at least 300 J, preferably at least 320 J, at a test temperature of -70 °C, d) an impact energy of at least 280 J, preferably at least 300 J, at a test temperature of -90 °C, and / or e) an impact energy of at least 150 J, preferably at least 200 J, at a test temperature of -110 °C.

[0025] Advantageously, after processing by welding with an energy input of 0.5 to 10 kJ / mm, preferably with an energy input of 1 to 5 kJ / mm, a) at a test temperature of -30 °C, an impact energy of at least 100 J, preferably at least 150 J, particularly preferably at least 200 J, b) at a test temperature of -50 °C, an impact energy of at least 30 J, preferably at least 60 J, particularly preferably at least 200 J, c) at a test temperature of -70 °C, an impact energy of at least 10 J, preferably at least 20 J, particularly preferably at least 100 J, or / and d) at a test temperature of -90 °C, an impact energy of at least 10 J, preferably at least 20 J, particularly preferably at least 100 J.

[0026] The above-mentioned values ​​for the impact strength of steels processed by welding refer to a determination on a weld seam formed during welding by means of a Charpy impact test at -90 °C, -70 °C, -50 °C or -30 °C, in particular according to the standard DIN EN ISO 148-1:2017, wherein preferably the specimen is provided with a V-notch and the specimen position is transverse to the rolling direction. They are preferably each an average value of results from the testing of at least three Charpy impact specimens.

[0027] In one embodiment of the invention, the crack tip opening displacement (CTOD) of the steel, according to DIN EN ISO 15653:2018, after processing by welding with an energy input of 0.5 to 10 kJ / mm², preferably with an energy input of 1 to 5 kJ / mm², is: a) at a test temperature of -10 °C > 0.50 mm, preferably > 0.70 mm, particularly preferably > 1.00 mm, or / and b) at a test temperature of -30 °C > 0.15 mm, preferably > 0.35 mm, particularly preferably > 1.00 mm. In a particularly preferred embodiment of the invention, the weld, which is provided for determining the impact strengths specified above, is oriented longitudinally to the rolling direction and is carried out at a preheating temperature of 125 °C to 250 °C. This preheating temperature complies with the standard DIN EN 10225-1:2019. Advantageously, the weld is designed with one side as a steep flank. The specimen orientation is preferably transverse to the rolling direction.

[0028] Charpy impact tests are expediently carried out at a distance 'A' of the sheet thickness from the sheet surface (a "quarter of the sheet"). Preferably, the position where the notch or crack is introduced is located along the melt line and within the steep flank.

[0029] Metal active gas welding is recommended for energy inputs of 0.5 to 1 kJ / mm², and submerged arc welding for energy inputs greater than 1 kJ / mm²; however, other welding processes are also conceivable. In particular, welding is carried out in accordance with the standard DIN EN 10225-1:2019, although a different energy input may be specified in deviation from the standard.

[0030] The energy input is conveniently calculated using the formula (Formula l ) can be calculated, where Q = heat input [kJ / mm], k = thermal efficiency, U = applied arc voltage [V], I = welding current [A], v = welding speed [mm / s]

[0031] Advantageously, after treatment by means of physical welding simulation of the coarse grain zone with an energy input of 0.5 to 10 kJ / mm, preferably with an energy input of 1 to 5 kJ / mm, a) at a test temperature of - 30 °C, the steel exhibits an impact energy of at least 50 J, preferably at least 100 J, particularly preferably at least 200 J, or / and b) at a test temperature of - 50 °C, an impact energy of at least 10 J, preferably at least 60 J, particularly preferably at least 200 J.All the above-mentioned values ​​for the impact strength of the steels, which have been treated by means of physical welding simulation of the coarse grain zone, refer to a determination by means of a Charpy impact test at -50 °C or -30 °C, in particular according to the standard DIN EN ISO 148-1:2017, wherein preferably the welding simulation is carried out with a single cycle, wherein the specimen is particularly preferably provided with a V-notch and the specimen position is transverse to the rolling direction.

[0032] It is advantageous to use a Tmax of 1350 °C for the welding simulation. The impact strengths are preferably an average of results from the testing of at least three Charpy impact test specimens.

[0033] In one embodiment of the invention, the energy input values ​​given above refer to treatment of the coarse grain zone by means of physical welding simulation by converting the desired energy input into a t8 / 5 time. For this purpose, the following two formulas are calculated and the higher value is used as the t8 / 5 time to input the desired energy:

[0034] 3D heat conduction: t 8 / 5 = (6700 - (Formula 2)

[0035] 2D heat conduction: t 8 / 5 = (4300 - (Formula 3) where

[0036] Tv = Preheating temperature [°C],

[0037] Fs, F2 = seam factor for three-dimensional and two-dimensional heat dissipation respectively, d = sheet thickness

[0038] The above-mentioned formulas ( 1 ), (2) and 3 are mentioned in the technical rule SEW 088:2017-10 (SEW 088 Supplement 1 :2017-10; SEW 088 Supplement 2:2017-10 Weldable unalloyed and low-alloy steels - Recommendations for processing, especially for fusion welding).

[0039] The temperature profile corresponding to the respective welding process was determined using the calculation method according to Hannerz:

[0040] Hannerz's equation: where

[0041] B = 1300 - 2T0,

[0042] T max = Peak temperature (°C),

[0043] To = initial temperature (°C), e = Euler number, t8 / 5 = cooling time from 800 °C to 500 °C (s)

[0044] (Source: Hannerz, NE, “Idealized thermal cycle for weld heat affected zone simulation of steel”, Perdue Thermal Physical Property Handbook)

[0045] To avoid the formation of embrittlement-causing martensite / austenite components, the steel is provided with a comparatively low content of silicon and aluminum, preferably in a content such that Al deoxidation and / or Si deoxidation takes place.

[0046] Silicon is present in such a concentration that, despite the low aluminum content and the associated comparatively low aluminum deoxidation, sufficient oxygen is bound by the silicon (silicon deoxidation). The stated minimum silicon content of 0.01 wt% is necessary to achieve adequate silicon deoxidation.

[0047] Since aluminum and silicon exhibit high solubility in ferrite, carbon is increasingly displaced from the ferrite to the austenite during the phase transition, and the driving force for cementite precipitation is significantly reduced. Retained austenite is indirectly stabilized thanks to the increased carbon content. To avoid martensite / austenite components, the steel according to the invention contains a maximum aluminum content of 0.05 wt.%, preferably 0.03 wt.%, and particularly preferably 0.01 wt.%, and a maximum silicon content of 0.25 wt.%, preferably 0.1 wt.%. It has proven particularly advantageous for the mechanical properties of the steel, especially its toughness, in the heat-affected zone of the weld joint to contain only < 0.1 wt.% silicon, preferably < 0.05 wt.% silicon, and particularly preferably < 0.025 wt.% silicon in the steel.Furthermore, particularly good mechanical properties can be achieved in the heat-affected area if the steel composition contains aluminum with a maximum of 0.015 wt.%.

[0048] Niobium in the steel serves, in particular, to prevent recrystallization at low rolling temperatures through solute drag and / or deformation-induced precipitates. The intended minimum content of 0.005 wt.% niobium, preferably 0.010 wt.% niobium, and particularly preferably 0.015 wt.% niobium, is provided to ensure that NbC forms as a deformation-induced precipitate. To prevent coarse primary precipitates, the alloy contains a maximum of 0.025 wt.% niobium, preferably a maximum of 0.022 wt.% niobium, and particularly preferably a maximum of 0.020 wt.% niobium.

[0049] Titanium leads to the formation of high-temperature-stable precipitates that withstand high temperatures during welding and inhibit temperature-induced growth of austenite grains in the heat-affected zone ("rinsing"). Advantageously, a maximum of 0.030 wt% titanium, preferably a maximum of 0.025 wt% titanium, and particularly preferably a maximum of 0.020 wt% titanium is provided to prevent coarse primary precipitates. The composition contains at least 0.005 wt% titanium, preferably at least 0.008 wt% titanium, and particularly preferably at least 0.010 wt% titanium to reduce austenite grain growth and ensure sufficient nitrogen binding, especially to prevent aging effects.

[0050] The alloy according to the invention contains at least 0.02 wt.% carbon and at least 1.60 wt.%, preferably at least 1.65 wt.%, particularly preferably at least 1.70 wt.% manganese, so that the steel achieves the required minimum strength. The steel has a carbon content of at most 0.1 wt.%, preferably at most 0.07 wt.%, to avoid the formation of hard phase regions.

[0051] Manganese is provided in a maximum content of 1.80 wt.%, preferably 1.75 wt.%, to prevent an austenite-stabilizing effect and to avoid the formation of martensite / austenite components.

[0052] The respective minimum contents of 0.01 wt.% copper, 0.01 wt.% nickel, 0.01 wt.% chromium and 0.001 wt.% molybdenum are intended for mixed crystal strengthening.

[0053] The maximum content of 0.4 wt.%, preferably 0.3 wt.%, copper and 0.3 wt.% nickel has an austenite-stabilizing effect and contributes to the avoidance of martensite / austenite components.

[0054] Chromium is provided in the steel with a maximum content of 0.30 wt.%, 0.20 wt.%, and particularly preferably 0.10 wt.%. Molybdenum is provided with a maximum content of 0.10 wt.% Mo to prevent the formation of brittle Cr and Mo carbides.

[0055] The maximum sulfur content of 0.0050 wt%, preferably 0.0015 wt%, is provided to ensure compliance with a required degree of purity.

[0056] Calcium is specified at a maximum content of 0.0050 wt% to maintain the required purity level. Furthermore, higher concentrations negatively affect the mechanical properties.

[0057] A maximum nitrogen content of 0.010 wt% prevents aging effects.

[0058] Phosphorus is provided at a maximum of 0.02 wt.% to avoid grain boundary fracture, especially temper embrittlement.

[0059] The steel may contain vanadium as a microalloying element. To prevent the formation of brittle carbides or nitrides, a maximum vanadium content of 0.010 wt.%, preferably 0.005 wt.%, and particularly preferably 0.003 wt.%, is specified.

[0060] In one embodiment of the invention, the steel contains no boron or only as a manufacturing-related impurity. Advantageously, the steel contains no zirconium and / or tantalum or only as a manufacturing-related impurity. Preferably, the steel contains no magnesium or only as a manufacturing-related impurity.

[0061] In one embodiment of the invention, the steel has a bainitic microstructure, in particular with a bainite content of at least 90%, preferably at least 95%, and particularly preferably 100%.

[0062] In one embodiment of the invention, the mean circle-equivalent grain diameter at a tolerance angle of 5°, preferably when considering the area-weighted distribution, is < 10 m, preferably < 8 m.

[0063] Advantageously, the steel has a proportion of high-angle grain boundaries > 63.0%, preferably > 70.0%. In one embodiment of the invention, the mean dislocation density GND (Geometrical Necessary Dislocations), preferably at a tolerance angle of 5°, and particularly at a measurement point spacing of 350 nm, is less than 100 × 10 12 / m 2 , preferably less than 70 • 10 ,2 / m 2 .

[0064] The specified grain sizes are expediently determined by SEM-EBSD (scanning electron backscatter diffraction), the value preferably referring to the mean value of the area-weighted distribution of the circle-equivalent diameter at a tolerance angle of 5°. The microstructure analyses are expediently carried out on a longitudinal section at a distance of 'A' of the sheet thickness from the sheet surface ("sheet quarter").

[0065] Advantageously, the steel is available as a cast and preferably rolled semi-finished product, preferably as a slab, as a sheet or as a tube, particularly preferably as a tube for transporting hydrogen or carbon dioxide.

[0066] In a particularly preferred embodiment of the invention, the steel is a welded structural steel, in particular a linepipe steel.

[0067] The aforementioned process for producing the steel is characterized by a process involving Al and / or Si deoxidation. The steel is expediently cast using a continuous casting process. A semi-finished product formed in this way, in particular a slab, is preferably thermomechanically rolled.

[0068] In one embodiment of the invention, the casting process involves forming the semi-finished product, in particular the slab, with a thickness of 200–600 mm, preferably 250–400 mm. Preferably, the semi-finished product, in particular the slab, is formed with a length of 1000–5200 mm, particularly preferably 1500–4000 mm.

[0069] Advantageously, the semi-finished product, in particular the slab, is reheated in a furnace to a temperature higher than the NbC solubility temperature. Preferably, it is reheated to a temperature between 950 and 1250 °C, more preferably between 1050 and 1150 °C.

[0070] In one embodiment of the invention, the formed semi-finished product, in particular the slab, is thermomechanically rolled in at least two rolling phases. Thermomechanical rolling is advantageous for the formation of a fine-grained final structure. In a preferred embodiment of the invention, the degree of deformation per rolling phase is preferably > 0.25, particularly preferably > 0.3, where degree of deformation = ^us a to a sdicke a v Final thickness J

[0071] In a particularly preferred embodiment of the invention, in the last

[0072] During the rolling phase, a degree of deformation > 1, preferably > 1.3, can be provided, which can increase the elongation of the austenite grains in the non-recrystallizing area, thereby further improving the fine grain structure of the final microstructure.

[0073] A rolling temperature of 700 to 850 °C is recommended.

[0074] Preferably, the material is rolled to a final thickness of 14 to 45 mm, preferably 20 to 35 mm.

[0075] In one embodiment of the invention, after rolling, accelerated, continuous cooling takes place with a cooling rate of at least 8 K / s.

[0076] It is advisable to use a cooling rate to maintain a temperature between 300 and

[0077] Cooled to 600 °C, preferably between 400 and 500 °C.

[0078] Preferably, cooling takes place immediately after rolling.

[0079] It is advantageous to process the steel further after rolling without reheating and / or annealing. Accordingly, secondary cooling is also unnecessary.

[0080] In one embodiment of the invention, the formed semi-finished product, in particular a sheet rolled from the slab, is welded and preferably retains a sufficiently high toughness, in particular the aforementioned values ​​of the impact strength.

[0081] The invention is explained in more detail below with reference to exemplary embodiments and the accompanying tables.

[0082] Table 1 shows the composition of steels A and B according to the invention. Steel C has a conventional composition and serves as a reference.

[0083] Compositions are given in wt.%.

[0084] Table 2 lists the rolling parameters used to process the steels. The following tests and measurements were carried out on a sheet produced from melt A according to the invention as described in rolling table 1, on a sheet produced from melt B according to the invention as described in rolling table 2, and on a sheet produced from reference melt C as described in rolling table 3.

[0085] Results of tensile tests and hardness measurements as well as microstructure analyses are shown in Table 3.

[0086] In the tensile tests carried out according to DIN EN ISO 6892-1:201, it was shown that the reference steel C has both a higher yield strength Rt0.5 and a higher tensile strength than steels A and B. The yield strength and tensile strength of steel B are higher than those of steel A.

[0087] The elongation A5 of the reference steel C is lower than that of steels A and B.

[0088] Reference steel C has a higher Vickers hardness (HV10), determined according to DIN EN ISO 6507-1:2018, than steels A and B. The Vickers hardness of steel B is higher than that of steel A.

[0089] The proportion of high-angle grain boundaries in steel C is 62.5%. The proportion in steel A is higher at 71.2%, and that in steel B is higher at 63.5%.

[0090] The mean grain size (5° Area Fraction; EBSD) in steel C is 6.21 µm. The mean grain sizes in steel A and steel B are larger, at 7.23 µm and 7.50 µm respectively.

[0091] For all Charpy impact tests described below, a Charpy impact test according to DIN EN ISO 148-1:2017 was performed. This standard test was carried out using a pendulum impact tester to determine the impact energy. The specimen was notched with a V-notch and positioned perpendicular to the rolling direction. Three measurements were taken, the individual values ​​of which are given, and the mean value of these individual values ​​is also presented.

[0092] Table 4 shows results of Charpy impact tests on the (unwelded) steel sheets made of steels A, B and C at test temperatures of

[0093] -30 °C, -50 °C and -70 °C. Table 5 shows further Charpy impact tests performed at

[0094] Tests were carried out at -90 °C and -110 °C. It is shown that the steels A and B according to the invention exhibit greater impact strength in all measurements.

[0095] These results are confirmed by the Battelle Drop Weight Tear Tests (BDWTT), which were carried out in accordance with the standard API RP 5L3:2021 and are also shown in Table 5.

[0096] To test the properties of the steels after welding, welding operations were carried out on steel sheets with a thickness of 28 mm, as explained below.

[0097] The welds with an energy input of 1 kJ / mm² were performed as metal active gas (MAG) welding, and the welds with an energy input of 5 kJ / mm² as submerged arc welding (SAW). For MAG welding, an ESAB gas metal arc welding machine with a LAW51 OW power source, a MEK4 wire feeder, a Fluxofil 41 electrode, and Sanarc Gl 8 shielding gas were used.

[0098] A submerged arc welding machine (“SAW”), with which the actual welding was carried out in the form of multi-layer welding, is composed of several components, namely a UniWeld device carrier with Subarc-5 control as an up double-head welding system with the power sources lx OERLIKON TRE1004 AG and lx SAF Starmatic 1000DC.

[0099] The welding was carried out using the OK OUTROD 13.40 electrode and the OP 121 TT powder.

[0100] The energy input was determined using the formula mentioned above. For example, in submerged arc welding, the thermal efficiency k = 1, while in MAG welding it is k = 0.8.

[0101] For the sheets from melt A, produced according to rolling table 1, and for those from melt B, produced according to rolling table 2, multi-layer welding was carried out under the various energy inputs specified in the tables. Three measurements were taken for each test procedure, the individual values ​​of which are given, and the mean value of the individual values ​​is given.

[0102] Table 6 shows results of Charpy impact tests on sheets with real multi-layer welds, with an energy input of 1 kJ / mm along a melt line for the multi-layer welds, performed at -30 °C, -50 °C, -70 °C, and -90 °C. The Charpy impact energies for the sheets produced from the steels according to the invention are greater at all test temperatures than those of the sheet made from the reference steel. Furthermore, it is evident that the sheet from melt A exhibits significantly higher Charpy impact energies than the sheet from melt B in the measurements at temperatures of -30 °C, -50 °C, and -90 °C. This is due to the comparatively lower silicon and aluminum content of the steel resulting from the silicon deoxidation that occurs.

[0103] Furthermore, the determination of the crack tip opening displacements (CTOD) of the steels was carried out in accordance with the standard DIN EN ISO 15653:2018. The results of the Charpy impact tests are confirmed by the CTOD measurements at -10 °C and -30 °C, which are shown in Table 7.

[0104] Table 8 shows results of Charpy impact tests on sheets with real multi-layer welds with an energy input of 5 kJ / mm along a weld seam formed during the multi-layer welds in sheet quarters at -30 °C, -50 °C, -70 °C, and -90 °C. The Charpy impact tests were carried out with the same sheets as the Charpy impact tests with an energy input of 1 kJ / mm. Even at an energy input of 5 kJ / mm, it is evident that the Charpy impact energies for the sheets made from the steels A and B according to the invention are significantly higher at all test temperatures than those of the sheet made from the reference steel C.

[0105] Furthermore, it is shown that the sheet from melt A exhibits significantly higher impact energies than the sheet from melt B, even with an energy input of 5 kJ / mm at different temperatures. This is due to the comparatively lower silicon and aluminum content of the steel resulting from the silicon deoxidation process.

[0106] Here too, the results of the Charpy impact tests are confirmed by the measurement results of the CTOD measurements at -10 °C and at -30 °C, which are shown in Table 9.

[0107] Further impact energies were determined using a physical welding simulation of the coarse grain zone, which is explained below.

[0108] To perform the Charpy impact test according to DIN EN ISO 148-1 on a material treated by physical welding simulation, a test blank with a length of at least 55 mm in the transverse direction of the sheet and a cross-sectional area of ​​10 mm x 10 mm (rolling direction x normal direction) was first taken from the layer ^-sheet thickness (a "quarter of the sheet") of the test material. The welding simulations were performed using a Gleeble 3800 hot forming simulator and the QuickSim 2 software (version 2.5.801 1.33152). As preparation, a thermocouple pair was spot-welded to one side of these impact blanks (standard specimen without a notch) at a distance of 27.5 mm from the end face. The prepared specimens were then inserted into the hot forming simulator and clamped with minimal tension between the designated copper dies, centering them at the position of the thermocouple.During the experiment, the current flow necessary for resistance heating was supplied via these stamps, and the control was achieved via the welded-on thermocouples.

[0109] The temperature control corresponding to the respective welding process was determined using the above-mentioned integrated calculation method according to Hannerz (Formula 4).

[0110] The input parameters used were Tmax = 1350 °C (1 cycle), To = 100 °C and t 8 / 5 A time of 60 s was selected. Cooling was controlled down to a minimum temperature of 350 °C. After the welding simulation, the thermocouples were removed, a 2 mm deep V-notch was made at their location, and the length of the samples was shortened to 55 mm if necessary.

[0111] Finally, the standard test of the weld-simulated samples was carried out using a pendulum impact tester to determine the impact energy.

[0112] The energy input values ​​given below refer to a treatment using physical welding simulation of the coarse grain zone by converting the desired energy input into a t8 / 5 time using formulas 2 and 3 shown above.

[0113] A preheating temperature Tv of 200 °C is appropriately assumed, along with a seam factor of 0.9 (Fi and F2) for simulating a multi-pass weld. A sheet thickness d of 28 mm was specified.

[0114] Table 10 shows the results of Charpy impact tests on various sheets produced from melts A, B, and C according to rolled sheets 1, 2, and 3, and treated for welding simulation as described above. Table 6 shows the Charpy impact test results with a t8 / 5 time of 60 s, corresponding to an energy input of 4.3 kJ / mm², for one cycle at 1350 °C at both -50 °C and -30 °C.

[0115] For the sheets produced from the steels according to the invention, significantly higher impact energies are also obtained after welding simulation during measurements at the various temperatures. Furthermore, it is confirmed that the sheet from melt A exhibits considerably higher impact energies than the sheet from melt B.

Claims

Patent claims:

1. Steel, in particular linepipe steel, having the following composition: 0.02 - 0.1 wt.% carbon, 0.01 - 0.25 wt.% silicon, 1.60 - 1.80 wt.% manganese, < 0.05 wt.% aluminum, 0.01 - 0.40 wt.% copper, 0.01 - 0.30 wt.% nickel, 0.01 - 0.30 wt.% chromium, 0.005 - 0.025 wt.% niobium, 0.005 - 0.030 wt.% titanium, < 0.0050 wt.% sulfur, < 0.0050 wt% calcium, < 0.010 wt% nitrogen, < 0.02 wt% phosphorus, < 0.010 wt% vanadium, 0.001 - 0.10 wt% molybdenum, Remainder: Iron and manufacturing-related impurities.

2. Steel according to claim 1, characterized by at least one of the following contents: < 0.07 wt.% carbon, < 0.1 wt.% silicon, preferably < 0.05 wt.% silicon, 1.65 - 1.75 wt.% manganese, < 0.03 wt.% aluminium, preferably < 0.01 wt.% aluminium, < 0.3 wt.% copper, < 0.2 wt.% chromium, preferably < 0.1 wt.% chromium; 0.010 - 0.022 wt.% niobium, 0.008 - 0.025 wt% titanium, < 0.0015 wt% sulfur, < 0.015 wt.% phosphorus.

3. Steel according to claim 1 or 2, characterized by a bainitic microstructure.

4. Steel according to any one of claims 1 to 3, characterized by a mean grain size of < 10 µm, preferably < 8 µm.

5. Steel according to any one of claims 1 to 4, characterized by a proportion of high-angle grain boundaries > 63%, preferably > 70%.

6. Steel according to one of claims 1 to 5, characterized in that the steel is in the form of a sheet, the thickness of which is preferably 14 to 45 mm, preferably 20 to 35 mm.

7. Steel according to one of claims 1 to 6, characterized in that the steel is in the form of a tube, in particular a welded tube, which preferably has a tube wall thickness of 14 to 45 mm, preferably of 20 to 35 mm.

8. Steel according to one of claims 1 to 7, characterized by a yield strength Rt0.5 of 360 to 600 MPa, preferably of 400 - 550 MPa.

9. Steel according to one of claims 1 to 8, characterized in that the steel a) exhibits an impact energy of at least 340 J, preferably at least 350 J, at a test temperature of -30 °C, and / or b) exhibits an impact energy of at least 340 J, preferably at least 350 J, at a test temperature of -50 °C.

10. Steel according to one of claims 1 to 9, characterized in that the steel, after processing by welding with an energy input of 0.5 to 10 kJ / mm, preferably with an energy input of 1 to 5 kJ / mm, a) at a test temperature of -30 °C, an impact energy of at least 100 J, preferably at least 150 J, particularly preferably at least 200 J, and / or b) at a test temperature of -50 °C, an impact energy of at least 30 J, preferably at least 60 J, particularly preferably at least 200 J. 1 1. Steel according to one of claims 1 to 10, characterized in that the crack tip opening displacement (CTOD) after processing by welding with an energy input of 0.5 to 10 kJ / mm, preferably with an energy input of 1 to 5 kJ / mm, a) at a test temperature of - 10 °C > 0.50 mm, preferably > 0.70 mm, particularly preferably > 1.00 mm, or / and b) at a test temperature of - 30 °C > 0.15 mm, preferably > 0.35 mm, particularly preferably > 1.00 mm.

12. Steel according to one of claims 1 to 1 1 , characterized in that the steel, after treatment by means of physical welding simulation of the coarse grain zone with an energy input of 0.5 to 10 kJ / mm, preferably with an energy input of 1 to 5 kJ / mm, a) at a test temperature of - 30 °C, exhibits an impact energy of at least 50 J, preferably at least 100 J, particularly preferably at least 200 J, or / and b) at a test temperature of - 50 °C, an impact energy of at least 10 J, preferably at least 60 J, particularly preferably at least 200 J.

13. Method for producing steel, in particular linepipe steel, wherein the steel is formed with the following composition: 0.02 - 0.1 wt% carbon, 0.01 - 0.25 wt.% silicon, 1.60 - 1.80 wt% manganese, < 0.05 wt.% aluminum, 0.01 - 0.40 wt% copper, 0.01 - 0.30 wt% nickel, 0.01 - 0.30 wt% chromium, 0.005 - 0.025 wt% niobium, 0.005 - 0.030 wt% titanium, < 0.0050 wt.% sulfur, < 0.0050 wt% calcium, < 0.010 wt% nitrogen, < 0.02 wt% phosphorus, < 0.010 wt.% vanadium, 0.001 - 0.10 wt% molybdenum Remainder: Iron and manufacturing-related impurities.

14. Method according to claim 13, characterized in that casting is carried out in the continuous casting process to form a semi-finished product, in particular a slab.

15. Method according to claim 13 or 14, characterized in that the semi-finished product formed, in particular the slab, is heated to a temperature between 950 and 1250 °C, preferably between 1050 and 1150 °C.

16. Method according to one of claims 13 to 15, characterized in that the semi-finished product formed, in particular the slab, is thermomechanically rolled in at least two rolling phases, wherein the degree of deformation per rolling phase is preferably > 0.25, particularly preferably > 0.

3.

17. Method according to one of claims 13 to 16, characterized in that rolling is carried out to a final thickness of 14 to 45 mm, preferably 20 to 35 mm.

18. Method according to one of claims 13 to 17, characterized in that, that after rolling, an accelerated, continuous cooling takes place with a cooling rate of at least 8 K / s.

19. Method according to claim 18, characterized in that the cooling rate is adjusted to a temperature between 300 and 600 °C, preferably between 400 and 500 °C.

20. Method according to one of claims 13 to 1 , characterized in that cooling takes place immediately after rolling.

21. Method according to one of claims 13 to 20, characterized in that the semi-finished product formed, in particular a sheet rolled from the slab or a tube optionally formed from the sheet, is welded, in particular with an energy input of 0.5 kJ / mm to 10 kJ / mm, preferably 1 kJ / mm to 5 kJ / mm.

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