MXPA97003873A - Steels with secondary hardness, ultra-high deresistence, with firmness and superior solditization - Google Patents

Abstract

The present invention relates to a method for producing low alloy, high strength steel, comprising a primary martensite / bainite microstructure, comprising: a) heating a steel ingot at a temperature sufficient to dissolve substantially all of the carbonitrides of vanadium and niobium carbonitrides b) reduce the ingot to form a plate, in one or more passes, in a first temperature range in which the austenite is recrystallized, c) laminate to finish the plate in one or more passes; ) laminate to terminate the plate in one or more passes in a second temperature range below the recrystallization temperature of the austenite, and above the transformation point Ar3; d) cool the completed laminated plate with water at a rate of at least 30 ° C / second, from a temperature above Ar3 up to a temperature < = 400 ° C and e) temper the plate cooled with water at a temperature no higher than the Ac1 transformation point, for a period of time sufficient to cause precipitation of the E-copper and the carbides or vanadium, niobium and molybdenum carbonitrides

Description

STEELS WITH SECONDARY HARDENING. OF ULTRA-HIGH RESISTANCE. WITH SUPERIOR STRENGTH AND WELFARE FIELD OF THE INVENTION This invention relates to a steel plate line pipe with ultra-high strength, which has superior properties of weldability, heat affected area resistance (HAZ) and firmness at low temperatures. More particularly, this invention relates to high strength, low alloy steels, of line pipes with secondary hardening, where the resistance of the HAZ zone is substantially the same as that in the rest of the line pipe, and to a process for manufacture plates that are precursors of the line tubes.

BACKGROUND OF THE INVENTION Currently, the highest elastic resistance of pipe lines, commercially available, is 556 MPA. While higher strength steel has been experimentally produced, for example up to about 696 MPA, several problems remain to be solved before this steel can be used safely as a line pipe. One such problem is the use of boron as a component of steel. While boron can increase the strength of the material, steels containing boron are difficult to process, producing inconsistent products as well as increased susceptibility to steel corrosion cracks. Another problem relating to high strength steels, ie steels having an elastic strength greater than about 556 MPa, is the softening of the HAZ zone after welding. The HAZ zone undergoes a local phase transformation or annealing during the thermal cycles induced by the welding, which produce a significant softening, up to about 15% or more, of the HAZ area in comparison with the base metal. Accordingly, it is an object of this invention to produce a low alloy, ultra-high strength steel, for use in a line pipe, with a thickness of at least 10 mm, preferably 15 mm, more preferably 20 mm, having an elastic strength of at least about 835 MPA and a tensile strength of at least 904 MPA, while maintaining consistent product quality, substantially eliminating or at least reducing the loss of strength in the HAZ zone during the thermal cycle induced by the welding, and that it has enough firmness at room temperature or low temperatures. A further object of this invention is to provide a manufacturer-friendly steel, with a unique secondary hardening response, to accommodate a wide variety of hardening parameters, for example time and temperature.

COMPENDIUM OF THE INVENTION According to this invention, a balance between steel chemistry and process technique is achieved, thus enabling the manufacture of a high strength steel having a specified minimum elastic strength (SMYS) of > 696 MPA, preferably > 725 MPA, more preferably 791 MPA, from which a line pipe can be prepared, and which, after welding, maintains the resistance of the HAZ zone substantially at the same level as the rest of the line pipe. In addition, this ultra-high strength, low alloy steel does not contain boron, ie, less than 5 ppm, preferably less than 1 ppm and more preferably no boron, and the product quality of the line pipe remains consistent and not too susceptible to the formation of cracks by stress corrosion. The preferred steel product has a substantially uniform microstructure, comprised primarily of martensite and fine-grained hardened bainite, which may be hardened secondarily by precipitates of the e-copper and the carbides or nitrides or vanadium, niobium and molybdenum carbonitrides. These precipitates, especially vanadium, reduce to a minimum the softening of the HAZ zone, probably preventing the elimination of dislocations in the regions heated to temperatures no greater than the transformation point Ac? or inducing the hardening of precipitation in regions heated to temperatures higher than the Ac transformation point?, or both. The steel plate of this invention is made by preparing a steel ingot in the usual manner and having the following chemical analysis, in weight percent: 0.03 to 0.12% C, preferably 0.05 to 0.09% 0.10 to 0.50% of Yes. 0.40 to 2.0% Mn 0.50 to 2.0% Cu, preferably 0.6 to 1.5% 0.50 to 2.0% Ni 0.03 to 0.12% Nb, preferably 0.04 to 0.08% 0.03 to 0.15% V, preferably 0.04 to 0.08% 0.20 a 0.80% Mo, preferably 0.3 to 0.6% 0.30 to 1.0% Cr, preferably for environments containing hydrogen 0.005 to 0.03% Ti 0.01 to 0.05% Al Pcm <; 0.35 the sum of vanadium + niobium > 0.1%. the rest being the Faith and incidental impurities.

Additionally, well known contaminants of N, P and S are minimized, although some N is convenient, as explained below, to supply titanium nitride particles that inhibit grain growth. Preferably, the concentration of N is around 0.001 to 0.01%, S not more than 0.01% and P not more than 0.01%. In this chemical analysis, the steel is substantially free of boron and the same is not added, and the concentration of this boron is < 5 ppm, preferably less than 1 ppm.

DESCRIPTION OF L08 DRAWINGS Figure 1 is a projection of the tensile strength (ksi or 6.96 MPA) of the steel plate (ordered) vs. the tempering temperature (abscissa) in itself. The figure also reveals, schematically, the additive effect of hardening / strengthening associated with the precipitation of e-copper and the carbides and carbonitrides of molybdenum, vanadium and niobium. Figure 2 is a bright field transmission electron micrograph, which reveals the granular microstructure of the plate bainite, as cooled, of the A2 Alloy. Figure 3 is an electron micrograph of bright field transmission, which reveals the martensitic lath microstructure of the plate, as cooled, of the Alloy Al. Figure 4 is a transmission electron micrograph of the bright field of Alloy A2 cooled and quenched at 600SC for 30 minutes. The dislocations, when cooled, are retained substantially after tempering, which indicates the remarkable stability of this microstructure. Figure 5 is a transmission electron micrograph of the dark field of the high amplification precipitate of the Al Al, cooled and quenched at 6002C for 30 minutes, which reveals complex mixed precipitation. The thickest globular particles are identified to be e-copper, while the finer particles are of the type (V, Nb) (C, N). The fine needles are of the type (Mo, V, Nb) (C, N) and these needles decorate and hold several of the dislocations. Figure 6 is a graph of the microhardnesses (Vickers hardness number, VHN, in the ordinate), through the heat affected zone (HAZ), of welding, for the steels in the abscissa Al (squares) and A2 (triangles) for a heat input of 3 kilo-Joules / mm. Typical microhardness data for commercial line tube steel, of lower strength, xlOO, is also projected for comparison (dotted line).

The steel ingot is processed by: heating this ingot to a temperature sufficient to dissolve substantially all, and preferably all, the vanadium carbonitrides and niobium carbonitrides, preferably in the range of 1100 to 12502C; a first hot rolling of the ingot at a rolling reduction of 30 to 70% to form a plate in one or more passes at a first temperature regime, in which the austenite recrystallizes; a second hot rolling at a reduction of 40 to 70% in one or more passes, at a second temperature regime somewhat lower than the first, and in which the austenite does not recrystallize and above the transformation point Ar3; hardening the laminated plate by cooling with water at a rate of at least 2Q2C / second, preferably at least about 3oac / second, from a temperature not less than the transformation point Ar3 at a temperature not higher than 400ac; and hardening the hardened laminated plate at a temperature no greater than the transition point Ac? for a sufficient time to precipitate at least one or more of the e-copper and the carbides or nitrides or carbonitrides of vanadium, niobium and molybdenum.

DETAILED DESCRIPTION OF THE INVENTION Ultra-high strength steels necessarily require a variety of properties and these properties are produced by a combination of elements and thermomechanical treatments, for example, minor changes in the chemistry of steel can lead to large changes in the characteristics of the product. The role of the various alloying elements and the preferred limits in their concentrations for the present invention, are supplied below. Carbon provides the matrix strengthening in all steels and welds, of any microstructure, and also intensifies precipitation primarily through the formation of small particles or precipitates of Nb (C, N), V (C, N) and M02C, if they are sufficiently fine and numerous. In addition, the precipitation of Nb (C, N) during hot rolling serves to retard recrystallization and to inhibit grain growth, thus providing a means of refining austenite grains and leading to an improvement in both the strength as in the firmness at low temperatures. Carbon also helps in hardening capacity, that is, the ability to form harder and stronger microstructures when cooling steel. If the carbon content is less than 0.03%, the effects of the strengthening will not be obtained. If the carbon content is greater than 0.12%, the steel will be susceptible to cracking when cooling in the field welding and decreases the firmness in the steel plate and its weld zone. Manganese is a reinforcement of the matrix in steels and welds and also contributes greatly to the hardening capacity. A minimum amount of 0.4% of Mn is necessary to achieve the high strength required. Like carbon, it is detrimental to the firmness of the plates and welds when it is in too high a quantity and also causes the formation of fissures upon cooling in the field welds, so an upper limit of 2.0% of Mn is imposed. This limit is also necessary to prevent severe segregation of the center line in steels of continuously melted line tubes, which is a factor that helps cause hydrogen-induced fissures (HIC). Silicon is always added to steel for deoxidation purposes and at least 0.1% is necessary for this. It is also a reinforcement in a strong solid ferrite solution. In larger quantities, Si has an adverse effect on the firmness of the HAZ zone, which is reduced to unacceptable levels when more than 0.5% is present. Niobium is added to promote the refinement of the grain of the laminated steel microstructure, which improves both strength and firmness. The precipitation of niobium carbonitride, during hot rolling, serves to retard recrystallization and to inhibit the growth of the grains, thus providing a means of refining the austenite grain. It gives an additional strengthening when tempering through the formation of Nb precipitates (C, N). However, too much niobium will be detrimental to the welding ability and firmness of the HAZ zone, so a maximum of 0.12% is imposed. Titanium, when added as a small amount, is effective in forming fine particles of TiN, which can contribute to the refinement of the grain size in the laminated structure and also acts as an inhibitor for grain thickening in the HAZ zone of the steel. Thus, firmness is improved. The titanium is added in such an amount that the Ti / N ratio is 3.4, so that the free nitrogen is combined with the Ti to form TiN particles. A Ti / N ratio of 3.4 also ensures that finely dispersed TiN particles are formed during continuous melting of the steel ingot. These fine particles serve to inhibit grain growth during subsequent reheating and hot rolling of austenite. The excess of titanium will deteriorate the steel firmness and welds, forming thicker Ti (C, N) particles. A titanium content below 0.005% can not supply a sufficiently fine grain size, while more than 0.03% causes a deterioration in the firmness.

Copper is added to supply the robustness of the precipitation when tempering the steel, after rolling, by the formation of fine copper particles in the steel matrix. Copper is also beneficial for corrosion resistance and HIC resistance. Too much copper will cause excessive hardening of the precipitation and poor firmness. Also, higher copper will make the steel more prone to surface cracks during hot rolling, so a maximum of 2.0% is specified. Nickel is added to counteract the detrimental effect of copper on the formation of surface cracks during hot rolling. It is also beneficial to the firmness of steel and its HAZ. Nickel is generally a beneficial element, except for its tendency to promote the formation of sulphide stress cracks, when more than 2% is added. For this reason, the maximum amount is limited to 2.0%. Aluminum is added to these steels for the purpose of deoxidation. At least 0.01% of the Al is required for this purpose. Aluminum also plays an important role in providing the firmness of the HAZ zone by eliminating the free nitrogen in the coarse-grained HAZ region, where the heat of the weld allows the TiN to partially dissolve, thereby releasing the nitrogen . If the content of the aluminum is too high, that is to say above 0.05%, there will be a tendency to form inclusions of type AI2O3, which are harmful in the steel firmness and its HAZ zone. The Vanadium is added to supply the robustness of the precipitation, they form fine particles of VC in the steel when tempering and its zone DO when cooling after the welding. When dissolved in austenite, vanadium has a strong beneficial effect on the hardening capacity. Thus, vanadium will be effective in maintaining the strength of the HAZ zone in a high strength steel. There is a maximum limit of 0.15%, since excessive vanadium will help to cause the formation of fissures when cooling in the welding of the field and will also deteriorate the firmness of the steel and its HAZ zone. Molybdenum increases the hardening capacity of a steel in direct cooling, so a strong matrix microstructure is produced and it also gives a fortification of the precipitation when tempering, by the formation of particles of M02C and NbMo carbide. Excessive molybdenum helps to cause the formation of fissures in cooling in the welding of the field and also deteriorates the strength of the steel and its HAZ zone, so a maximum of 0.8% is specified.

Chromium also increases the hardness capacity in direct cooling. Improves resistance to corrosion and HIC. In particular, it is preferred to prevent the ingress of hydrogen by the formation of a Cr2? 3 rich oxide film on the steel surface. A chromium content below 0.3% can not supply a stable film of Cr2? 3 on the surface of the steel. As for molybdenum, excess chromium helps cause the formation of fissures by cooling in field welding, and also deteriorates the strength of the steel and its HAZ zone, so a maximum of 1.0% is specified. Nitrogen can not be prevented from entering and remaining in the steel during its manufacture. In this steel, a small amount is beneficial in the formation of fine particles of TiN, which prevents the growth of the grain during the hot rolling and thus promotes the refinement of the grain in the rolled steel and its HAZ zone. At least 0.001% of N is required to supply the required volume fraction of TiN. However, too much nitrogen will deteriorate the steel's firmness and its HAZ zone, so the maximum 0.01% N amount is imposed. While high-strength steels with elastic strengths of 835 MPA or more have been produced, these steels lack the strength and weldability requirements necessary for line pipes, because such materials have a relatively high carbon equivalent, that is, greater than the Pcm of 0.35, as specified herein. The first objective of the thermomechanical treatment is to achieve a sufficiently fine microstructure of tempered martensite and bainite, which is secondarily hardened by even finer dispersed precipitates of e-Cu, Mo2C, V (C, N) and Nb (C, N) . The fine strips of tempered martensite / bainite supply the material with high strength and good firmness at low temperatures. Thus, the heated grains of austenite are first made fine in size, for example <20 microns, and second, they deform and flatten, so that the thickness dimension of the austenite grains is even smaller, for example from < 8-10 microns, and third, these flattened austenite grains are filled with high dislocation density and cutting bands. This leads to a high density of the potential nucleation sites for the formation of the transformation phases when the steel ingot is cooled after completing the hot rolling. The second objective is to retain enough Cu, Mo, V and Nb, substantially in the solid solution, after the ingot is cooled to room temperature, so that the Cu, Mo, V and Nb are available, during the tempering treatment , to be precipitated as e-Cu, M02C, Nb (C, N) and V (C, N). Thus, the temperature of reheating before the hot rolling of the ingot, has to satisfy both demands to maximize the solubility of Cu, V, Nb and Mo, while preventing the dissolution of the TiN particles formed during the continuous melting of the steel and thus preventing the thickening of the austenite grains before hot rolling. To achieve both objectives for the steel compositions of the present invention, the reheat temperature before hot rolling should not be less than 100 ° C and not more than 1250 ° C. This reheat temperature used for any steel composition in the range of the present invention is easily determined by experimentation or by calculation, using suitable models. The temperature that defines the boundary between these two intervals of the same, the recrystallization interval and the non-recrystallization interval, depends on the heating temperature before laminating, the concentration of the carbon, the concentration of the niobium and the amount of reduction given in the passes of the laminate. This temperature can be determined for each steel composition or by experimentation or model calculation. The conditions of hot rolling provide, in addition to obtaining the austenitic grains of fine size, an increase in the density of dislocation through the formation of bands of deformation in the austenitic grains, thus maximizing the density of the potential sites within of the deformed austenite for the nucleation of the transformation products during cooling, after the lamination is finished. If the reduction of the laminate in the recrystallization temperature range decreases, while the reduction of the laminate in the non-recrystallization temperature range is increased, the austenite grains will be insufficiently fine in size, resulting in austenite grains. thick, thus reducing both strength and firmness and causing the highest susceptibility to stress corrosion cracking. On the other hand, if the reduction of the laminate in the recrystallization temperature range is increased while the laminate reduction in the non-recrystallization temperature range decreases, the formation of the deformation bands and the dislocation substructures in the grains of austenite become inadequate to provide sufficient refining of the products of the transformation, when the steel is cooled, after finishing the rolling. After finishing the laminate, the steel is subjected to cooling with water from a temperature not lower than the transformation temperature Ar3 and ending at a temperature not higher than 4002C. Air cooling can not be used because it would cause the austenite to be converted to ferrite / pearlite aggregates, leading to deterioration in strength. In addition, during the cooling with air, the Cu will be precipitated and over-aged, making it virtually ineffective for the strengthening of the precipitation in the tempering., The completion of the cooling with water at a temperature higher than 400dC, causes the hardening of transformation insufficient during cooling, thus reducing the strength of the steel plate. The steel plate, hot rolled and cooled with water, is then subjected to a tempering treatment, which is carried out at a temperature no higher than the Ac? This tempering treatment is conducted in order to improve the steel firmness and allow sufficient precipitation in substantially uniform form, through the microstructure of e-Cu, M02C, V (C, N) and Nb (C, N), for increased resistance. Therefore, the secondary robustness is produced by the combined effect of e-Cu, M02C, V (C, N) and Nb (C, N), precipitates. The peak hardening due to e-Cu and M02C occurs in the temperature range of 450 to 5502C, while the hardening due to V (C, N) / Nb (C, N) occurs in the temperature range of 550 to 650SC. The use of these species of precipitates to achieve secondary hardening provides a response to hardening that is minimally affected by the variation in the composition or microstructure of the matrix, thus providing a uniform hardening through the plate. In addition, the wide temperature range of the secondary hardening response means that the stiffening of the steel is relatively insensitive to the hardening temperature. Therefore, the steel is required to be tempered for a period of at least 10 minutes, preferably at least 20 minutes, for example 30 minutes, at a temperature that is greater than about 4002C and less than about 7002C, preferably from 500 to 6502C. . A steel plate produced through the described process exhibits a high strength and high firmness with high uniformity in the thickness direction of the plate, despite the relatively low concentration of carbon. In addition, the tendency for softening of the zone affected by heat is reduced by the presence of, and additional formation of, the precipitates of V (C, N) and Nb (C, N) during welding. Also, the sensitivity of the steel to the formation of fissures by hydrogen, is markedly reduced. The HAZ zone develops during the thermal cycle induced by welding and can extend by 2-5 mm from the weld fusion line. In this zone, a temperature gradient is formed, for example from about 700 to 14002C, which encompasses an area in which the following softening phenomena occur, from a lower temperature to a higher one: softening by the reaction of high temperature tempering , and softening by austenitization and slow cooling. In the first area, vanadium and niobium and their carbides or nitrides are present to prevent or minimize softening substantially, retaining the high density of dislocation and sub-structures; in the second area, the additional precipitates of vanadium carbonitride and niobium are formed and softening is minimized. The net effect during the thermal cycle induced by the welding is that the HAZ zone retains substantially all the resistance of the remaining base steel in the line pipe. The loss of strength is less than about 10%, preferably less than 5%, and more preferably the loss of strength is less than 2% relative to the strength of the base steel. That is, the strength of the HAZ zone after welding is at least 90% of the strength of the base metal, preferably at least 95% of the strength of the base metal and more preferably at least 98% of the strength of the base metal. Maintenance of resistance in the HAZ zone is primarily due to the concentration of vanadium + niobium > 0.1% and preferably each of the vanadium and niobium are present in the steel at a concentration of > 0.4%. The line tube is formed from the plate by the well known UOE process, in which: the plate is formed in a U configuration, then formed in an O configuration, and this O configuration is expanded from 1 to 3 %. The formation and expansion with its concomitant hardening effects leads to a greater resistance of the line pipe. The following examples serve to illustrate the invention described above.

DESCRIPTION AND EXAMPLES OF MODALITIES A load of 227 kilograms of each alloy, representing the following chemical analyzes, was melted by vacuum induction, molded into ingots and forged into sheets of 1000 mm thickness and hot rolled, as is described below, for the characterization of the properties. Table 1 shows the chemical composition (% by weight) of the alloys Al and A2.

TABLE 1 Alloy Al A2 c 0.089 0.056 Mn 1.91 1.26 P 0.006 0.006 S 0.004 0.004 Yes 0.13 0.11 Mo 0.42 0.40 Cr 0.31 0.29 Cu 0.83 0.63 Ni 1.05 1.04 Nb 0.068 0.064 V 0.062 0.061 Ti 0.024 0.020 Al 0.018 0.019 (ppm) 34 34 pcm 0.30 0.22 The ingots, as they were cast, must undergo proper reheating before rolling, to induce the desired effects in the microstructure. The reheating serves in order to substantially dissolve in the austenite the carbides and carbonitrides of Mo, Nb and V, so that the elements can be reprecipitated later in the steel process in the most desired form, that is, fine precipitation in the austenite before cooling, as in the tempering and welding of the austenite transformation products. In the present invention, reheating is carried out at temperatures in the range of 1100 to 12502C and more specifically of 12400C for Alloy 1 and 11602C for Alloy 2, each for 2 hours. The design of the alloy and the thermomechanical process have been coupled to produce the following equilibrium with respect to the first strong carbonitrides, specifically niobium and vanadium: • approximately one third of these elements precipitate in austenite before rapid cooling; • approximately one third of these elements precipitate in the products of transformation of austenite, in tempering after cooling; • approximately one third of these elements are retained in solid solution, which will be available for precipitation in the HAZ zone, to improve the normal softening observed in steels that have an elastic strength greater than 557 MPA. The thermomechanical lamination program involving the initial square 100 mm plate is shown in Table 2 below for Al Alloy. The rolling program for Alloy A2 was similar, but the reheat temperature was 1160ac. TABLE 2 Starting Thickness; 100 mm Reheat Temperature: 12 QSC Pass Thickness (mm) After Temperature (SQ Pass 0 100 1240 1 85 1104 2 70 1082 3 57 1060 Delay (turning the piece over the edge (l) 4 47 899 5 38 877 6 32 852 7 25 827 8 20 799 Cooling with water at the Ambient Temperature (1) allows cooling on all sides, due to the small sample.

The steel was cooled from the finished rolling temperature to room temperature, at a cooling rate of 302C / second. This cooling regime produces the desired microstructure, such as cooled, which consists predominantly of bainite and / or martensite or preferably 100% of lath martensite. In general, when aging, the steel softens and loses hardness and strength such as after cooling, the degree of this loss of strength is a function of the specific chemistry of the steel. In the steels of the present invention, this natural loss in strength / hardness is substantially eliminated or significantly improved by a combination of the final precipitation of e-copper, VC, NbC and Mo2C. The tempering is carried out at various temperatures in the range of 400 to 700 ° C, for 30 minutes, followed by cooling with water or air, preferably with water, at room temperature. The design of the multiple secondary hardening, which results from the precipitates, as reflected in the strength of the steel, is illustrated schematically in Figure 1 for Al Alloy. This steel has a high hardness and strength, on cooling, but could soften in the absence of secondary hardening precipitates, easily in the range of the aging temperature of 400 to 700 ° C, as schematically shown by the dotted line, which declines continuously. The solid line represents the real measured properties of the steel. The tensile strength of steel is remarkably insensitive to aging in the wide temperature range of 400 to 6502C. The robustness results from the precipitation of e-Cu, M02C, VC, NbC, which occurs and is maximized at several temperature regimes in this wide range of aging and provides a cumulative resistance to compensate for the loss of resistance normally observed with aging. Simple low-alloy and martensitic carbon steels, without strong carbide formers. In Alloy A2, which has lower values of carbon and Pcm, the secondary hardening processes showed a similar behavior as Al Al, but the resistance level was lower than that in Al Al for all process conditions. An example of the microstructure, upon cooling, is presented in Figures 2 and 3, which show a bainitic and Tuesdaynsitic microstructure, predominantly granular, respectively, of these alloys. The greater hardness capacity that results from the greater Al alloy alloy, resulted in the martensitic lath structure, while Alloy 2 was characterized by predominantly granular bainite. Notably, even after annealing at 6002C, both alloys showed excellent microstructural stability, Figure 4, with negligible recovery in the dislocation substructure and little cell / ribbon / grain growth. When tempering in the range of 500 to 650 ° C, the precipitation of secondary hardening was observed in the form of e-copper precipitates. globular and needle type precipitates of M02C and (Nb, V) C. The particle size for the precipitates varied from 10 to 150 Á. The very high magnification transmission electron micrograph taken selectively to highlight the precipitates is shown in the dark field image of the precipitate, Figure 5. The data of the tensile strength at room temperature are summarized in Table 3, together with firmness at ambient and low temperatures. It is clear that the Alloy Al exceeds the minimum desired resistance to the tension of the invention, while the Alloy 2 meets this criterion. The firmness to the impact of Slot V Charpy at ambient temperature and at -402C, was performed on longitudinal and transverse samples, according to the specification of ASTM E23. For all tempering conditions, Alloy A2 had a greater impact strength, which exceeds 200 Joules at -40ac. Alloy 1 also demonstrated excellent impact strength in light of its ultra-high strength, which exceeds 100 Joules at -402C, preferably the steel strength of >; 120 Joules at -402C.

The microhardness data obtained from a simple laboratory strip in a plate welding test, are projected in Figure 6 for the steels of the present invention, together with comparable data for a commercial steel line resistance pipe. minor, the XlOO. The laboratory welding was performed at a heat input of 3kJ / mm and the hardness profiles through the welding zone HAZ are shown. The steels produced according to the present invention exhibited a remarkable resistance to softening of the HAZ zone, less than 2%, compared to the hardness of the base metal. In contrast, XlOO commercial steel, which has a base metal strength and hardness much lower than Al steel, showed a significant softening of 15% in the HAZ zone. This is even more remarkable, since it is well known that maintaining the strength of the base metal in the HAZ zone becomes even more difficult as the strength of the base metal increases. The high strength of the HAZ zone of this invention is obtained when the welding heat input varies from about 1-5 kilo Joules / mm.

TABLE 3: TYPICAL MECHANICAL PROPERTIES Steel Condition Properties Tensiies ü) Ownership YS MPA UTS. MPA E (%) vE20. Joul Al As cooled 904 1205 13 136 Tempering at 5502C for 30 minutes 1058 1090 15 123 Tempering at 650se for 30 minutes 1030 1038 17 157 A2 As cooled 904 1205 13 136 Tempering at 5502C for 30 minutes 1058 1090 15 123 Tempering at 6502C for 30 minutes 1030 1038 17 157 (1) Cross direction, round samples (ASTM, E8); YS - 0.2% elastic displacement resistance; UTS - Final tensile strength; EL - Elongation in a caliber length of 25.4 mm. (2) Transverse sample: vE20 - V-groove energy in test at 20 ° C; vE40 - V-groove energy in test at -40 ° C.

Claims (18)

1. CLAIMS 1. A method to produce high strength, low alloy steel, with an elastic strength of at least 835 MPA, this method comprises: (a) heating a steel ingot at a temperature sufficient to dissolve substantially all vanadium carbonitrides and niobium carbonitrides; (b) reducing the ingot to form a plate, by one or more passes, to a first temperature range in which the austenite recrystallizes; (c) further reducing the plate in one or more passes to a second temperature range, below the recrystallization temperature of the austenite, and above the transformation point Ar3; (d) cooling the reduced plate with water from a temperature above the transformation point Ar3 at a temperature < 4002C; in which the steel contains niobium and vanadium in a total concentration of > 0.1% by weight. The method of claim 1, wherein the temperature, in step (a), is approximately 1100 to 12502C. The method of claim 1, wherein the reduction, in step (b), is from about 30 to 70%, and the reduction in step (c) is about 40 to 70%. 4. The method of claim 1, wherein the plate cooled with water is tempered at a temperature not higher than the transformation point AC] _, for a sufficient period of time, to cause precipitation of the e-copper and the carbides or vanadium carbonitrides, niobium and molybdenum. The method of claim 4, wherein the tempering step is carried out at a temperature range of 400 to 7002c. 6. The method of claim 1, wherein the step of cooling with water is carried out at a rate of at least about 202c / second. The method of claim 1, wherein the plate is formed in a line tube and expanded approximately 1 to 3%. The method of claim 1, wherein the chemical analysis of the steel, in percent by weight is: 0.03 to 0.12% C 0.01 to 0.50% Si 0.40 to 2.0% Mn 0.50 to 2.0% Cu 0.50 a 2.0% Ni 0.03 to 0.12% Nb 0.03 to 0.15% V 0. 20 to 0.80% of Mo 0.005 to 0.03 of Ti 0.01 to 0.05 of Al Pcm < 0.35 and the remainder being Fe. 9. The method of claim 8, wherein the steel contains from 0.3 to 1.0% Cr. 10. The method of claim 8, wherein the concentrations of each of the vanadium and the niobium are from > 0.04% 11. A high-strength, low-alloy steel with at least an elastic strength of 835 MPA, comprising primarily a martensite / bainite phase, containing precipitates of e-copper and the vanadium carbides, nitrides or carbonitrides. , niobium and molybdenum, and where the vanadium + niobium concentrations are > 0.1% by weight. 12. The steel of claim 1, in the form of a plate with a thickness of at least about 10 mm. The steel of claim 11, wherein additional amounts of vanadium and niobium are in solution. The steel of claim 13, wherein the concentrations of each of the vanadium and the niobium are > 0.4% by weight. 15. The steel of claim 11, wherein the chemical analysis thereof, in percent by weight is: 0.03 to 0.12% C 0.01 to 0.50% Si 0.40 to 2.0% Mn 0.50 to 2.0% Cu 0.50 to 2.0 % of Ni 0.03 to 0.12% of Nb 0.03 to 0.15% of V 0.20 to 0.80% of Mo 0.005 to 0.03 of Ti 0.01 to 0.05 of Al Pcm 0-35 and the rest being Fe. 16. The steel of claim 15 , which contains from 0.3 to 1.0% Cr. 17. The steel of claim 14, in which the resistance of the area affected by heat, after welding, is at least 95% of the strength of the base metal . The steel of claim 14, wherein the strength of the area affected by heat after welding is at least 98% of the strength of the base metal.