KR20090078807A - Low yield ratio dual phase steel linepipe with superior strain aging resistance - Google Patents

Abstract

A steel composition and method from making a dual phase steel therefrom. In at least one embodiment, the dual phase steel comprises carbon in an amount of about 0.05% by weight to about 0.12 wt%; niobium in an amount of about 0.005 wt % to about 0.03 wt%; titanium in an amount of about 0.005 wt% to about 0.02 wt%; nitrogen in an amount of about 0.001 wt% to about 0.01 wt%; silicon in an amount of about 0.01 wt% to about 0.5 wt%; manganese in an amount of about 0.5 wt% to about 2.0 wt%; and a total of molybdenum, chromium, vanadium and copper less than about 0.15 wt%. The steel has a first phase consisting of ferrite and a second phase comprising one or more constituents selected from the group consisting of carbide, pearlite, martensite, lower bainite, granular bainite, upper bainite, and degenerate upper bainite. A solute carbon content in the first phase is about 0.01 wt% or less.

Description

LOW YIELD RATIO DUAL PHASE STEEL LINEPIPE WITH SUPERIOR STRAIN AGING RESISTANCE}

The present invention relates generally to linepipes and, more particularly, to low yield ratio, composite tissue steel linepipes with excellent strain aging resistance and methods for making the same.

This application claims the priority of US Provisional Application No. 60 / 850,216, the contents of which are hereby incorporated by reference.

Natural gas is becoming an increasingly important energy source. Usually the world's major natural gas producers are far from major markets. As such, pipelines must traverse long distances either onshore or underwater, which can cause severe deformations on the pipeline. Seismic active areas and / or arctic areas that undergo frost-heave and thawing settlement cycles can cause severe deformations in the pipeline. In addition, pipelines placed over the sea floor are subject to severe deformation due to positional fluctuations or warpage caused by water flow.

Thus, linepipes used in these environments require excellent uniform elongation in the longitudinal direction of the pipe and excellent deformation capabilities such as low yield to tensile strength ratio or yield ratio (YR) to ensure mechanical integrity. Composite steel (DP) steel has a relatively soft phase and a relatively hard phase such as a ferrite phase. Harder phases usually have more than one component. Composite tissue steels (ie, steels with composite microstructure (DP) microstructures) provide high uniform elongation and low yield ratios, thus providing excellent strain capability. For these reasons, DP river linepipes are attractive as installations in seismic active regions or in arctic regions subject to semi-permanent icing environments or in other situations requiring high strain capacity.

Typically DP steels are processed according to a series of steps. For example, steel slabs are typically reheated to an austenite temperature range of about 1000 ° C. to 1250 ° C. and co-rolled within the recrystallization temperature range to refine grain size. The roughly rolled steel is then finish rolled within the non-recrystallization temperature range, cooled to a temperature below Ar ₃ to form ferrite, and then accelerated cooled to a temperature below 400 ° C. The plate is then typically U-shaped, then O-shaped, seam welded and stretched to the desired outer diameter (known as UOE pipe making process). Arc welding, resistance welding or laser welding or the like can be used in the seam welding step of the UOE process.

Then, the outer diameter of the pipe is typically covered to provide protection against corrosion. Typically melt adhesive epoxy (FBE) coatings are used for this purpose. During the FBE coating process, the pipe is heated to elevated temperature and covered with polymer.

Due to the linepipe manufacturing and coating processes, linepipe steels, for the most part, including DP steel, are subject to strain aging. Strain aging leads to a drop in strain capacity and is a kind of behavior generally associated with yield point phenomena, which increases the flow or yield strength of metals and reduces ductility upon heating after cold working, such as during the FBE coating process. In other words, strain aging refers to the hardening of the metal and the corresponding reduction in ductility.

Strain aging can be caused by the interaction between the strain field of the solute atoms in the steel and the strain field of the dislocations. Formation of solute atmospheres around the dislocations (“Cottrell atmospheres”) increases resistance to dislocation movements upon subsequent loading. In general, the ductility of a metal is proportional to the ease with which dislocations move in that metal. As a result, higher forces or stresses are required to tear dislocations from the Cottrell atmosphere, leading to an increase in yield strength, a decrease in ductility and an increase in ductile-to-brittle transition temperature. As a result, strain aging reduces strain capacity. Therefore, steels having higher resistance to strain aging or components made from the steels substantially maintain their deformation capacity after aging following cold work.

The aging process is thought to occur in two stages. The first step is to diffuse the solute species into dislocations to form an atmosphere. In the second step, precipitates form on the solute species dislocation. Such precipitates contribute to an increase in the overall strength of the material, but lower the elongation to break. Usually, only the first stage occurs when the concentration of solute species is low.

Typically the elements that result in relatively low temperature (<300 ° C.) strain aging in steel are carbon and nitrogen, which are invasive solute elements in steel. Such elements have a low equilibrium solubility and significantly higher diffusivity compared to substituted solutes in steel, such as chromium, vanadium, molybdenum, copper and magnesium, to name just a few. However, carbide and nitride forming substituted alloy elements, such as chromium, vanadium, molybdenum, indirectly affect the strain aging sensitivity by increasing the equilibrium solubility of carbon and nitrogen. With such a position, solute carbon and nitrogen tend to migrate to the potential on the ferrite phase forming the Cottrell atmosphere. As mentioned above, this Cottrell atmosphere tends to limit the movement of dislocations and therefore impairs the deformation capacity of the steel.

Similarly, it is believed that the yield strength of composite tissue steel linepipes can be increased during post-forming treatments such as FBE coating processes. As mentioned, typical FBE coating processes require heat. The thermal exposure required by the FBE coating process provides enough energy for the solid carbon and / or nitrogen atoms in the linepipe to move to the potential on the ferrite. Such movement impairs the linepipe's deformation capacity for the reasons mentioned above.

Therefore, there is a need for composite tissue steels and linepipes made therefrom with low yield ratios, high uniform elongation and excellent work hardening properties for high deformation capacity to ensure mechanical integrity in applications in aggressive environments. . There is also a need for new steel chemistries that give excellent strain aging resistance to steel and the products produced therefrom. In addition, there is a need for a method of processing composite tissue steel that provides excellent strain aging resistance properties to linepipes and precursor steels produced for excellent strainability, especially after thermal treatment processes such as FBE coating processes.

FIELD OF THE INVENTION The present invention relates to composite steels capable of providing line pipes and, more particularly, low yield ratio, composite tissue steel line pipes with excellent strain aging resistance, and a method for manufacturing the same.

In one embodiment, the present invention is directed to a steel composition and a method for making composite tissue steel therefrom. In at least one embodiment, the composite tissue steel comprises about 0.05% to about 0.12% carbon; From about 0.005% to about 0.03% niobium; About 0.005% to about 0.02% titanium; About 0.001% to about 0.01% nitrogen; About 0.01% to about 0.5% silicon; About 0.5% to about 2.0% manganese; And less than 0.15% by weight of molybdenum, chromium, vanadium and copper. The steel is from a first phase consisting of ferrite and from a group consisting of carbide, pearlite, martensite, lower bainite, granular bainite, upper bainite and degenerate upper bainite. Has a second phase comprising at least one component selected. The solute carbon content in the first phase is about 0.01% by weight or less.

In one exemplary embodiment, a method for manufacturing composite tissue steel includes heating the steel slab to a reheating temperature of about 1000 ° C. to about 1250 ° C. to provide a steel slab consisting essentially of austenite phase. The steel slabs are pressed to form a plate in at least one hot rolling pass at a first temperature sufficient to recrystallize the austenite phase. The plate is pressed into one or more hot rolling passes at a second temperature at which austenite forms a rolled plate without recrystallization. The second temperature is lower than the first temperature. The rolled plate is then cooled to a first cooling temperature sufficient to induce the transformation of austenite to ferrite and then reduce the cluster forming atoms in the ferrite.

In another embodiment, a method for manufacturing composite tissue steel includes heating the steel slab from about 1000 ° C. to about 1250 ° C. to provide a steel slab consisting essentially of austenite phase. The steel slab is pressed to form a plate in one or more hot rolling passes at a temperature such that recrystallization of the austenite phase to produce a particulate austenite phase. The plate is further pressed into one or more hot rolling passes at a temperature lower than the temperature at which austenite does not recrystallize. The plate is cooled to a first cooling temperature sufficient to induce austenite to ferrite transformation and is quenched to a second temperature at a rate of at least 10 ° C. (18 ° F./sec) per second. It is then cooled at a rate that reduces the solute carbon in the ferrite.

The features and advantages of the present invention will become apparent to those skilled in the art from the following description of the preferred embodiments when taken in conjunction with the accompanying drawings. Within the spirit of the invention, numerous modifications can be made by those skilled in the art.

These drawings illustrate certain aspects of some of the embodiments of the invention and should not be used to limit or define the invention.

1-4 illustrate changes in the mechanical properties of certain illustrative steels made in accordance with one or more embodiments of the present invention.

5 shows the relationship between yield ratio (%) as a function of heat treatment temperature for steel produced according to one or more embodiments of the present invention and conventional steel.

6A is a scanning electron microscope (SEM) image of a conventional steel sheet that has been heat treated.

FIG. 6B is a transmission electron microscope (TEM) image of the heat treated conventional steel sheet shown in FIG. 6A.

7A is an SEM image of a steel sheet at a quarter thickness made in accordance with one or more embodiments of the present invention. The image mainly shows a steel having a granular bainite (GB), upper bainite (UB) or pearlite and a second phase which is some lath martensite (LM).

FIG. 7B is a TEM image at a quarter thickness of the steel sheet shown in FIG. 7A. The image shows an entangled or wave shaped dislocation that shows little or no carbon and / or nitrogen supersaturation.

The present invention generally relates to linepipes and more particularly to low yield ratios, composite tissue steel linepipes having excellent strain aging resistance and methods for making them.

Provided are high strength, composite steel (DP) steels with low yield-to-tensile ratio, high uniform elongation and high work hardening coefficients and methods for making them. Such steel can be post-treated without adversely affecting the deformation capacity. Among many other applications commonly known for steel, the steel is suitable for linepipes, offshore structures, oil and gas production facilities and pressure vessels.

In one or more embodiments, the steel includes a balance of alloying elements that reduce the supersaturation of carbon and nitrogen on the ferrites of iron and steel, thereby providing resistance to strain aging. Preferably, the content of solute carbon on the ferrite is less than 0.01% by weight, more preferably less than 0.005% by weight. In at least one embodiment, the solute carbon content is 0.005% to 0.01% by weight. In one or more embodiments, the solute carbon content is about 0.006%, about 0.007%, about 0.008%, or about 0.009% by weight. Preferably the solute nitrogen content on the ferrite is less than 0.01% by weight, more preferably less than 0.005% by weight. In at least one embodiment, the solute nitrogen content is 0.005% to 0.01% by weight. In one or more embodiments, the solute nitrogen content is about 0.006%, about 0.007%, about 0.008%, or about 0.009% by weight.

Preferably, the steel is formulated to have a tensile strength of more than 500 MPa, more preferably 520 MPa or greater, before and after heating for treatment processes such as anti-corrosion coating treatment processes. The steel is also formulated to have a minimum yield strength of about 400 MPa, more preferably a minimum yield strength of about 415 MPa. The steel is also a precursor having a yield to tensile strength (YTS) ratio or yield ratio (YR) of about 0.90 or less, preferably about 0.85 or less, even more preferably about 0.8 or less, both before and after heating for the treatment process. It is formulated to provide steel and line pipes made therefrom. In one embodiment, YR is 0.89 or 0.88 or 0.87 or 0.86 or 0.85. The steel is also formulated to have a minimum uniform elongation of greater than about 8% and preferably greater than about 10% in precursor steel and linepipes made therefrom, both before and after heating for the treatment process. In addition, the steel is greater than about 120 J at Charpy-V-Notch test at -12 ° C, preferably greater than about 200 J at Charpy-V-Notch test at -12 ° C, even more It is preferably formulated to have high toughness, such as greater than about 250 J in Charpy-V-notch testing at -12 ° C.

Preferred alloying elements and preferred ranges are described in more detail below. For example, preferably the steel has a carbon content of less than 0.12% by weight, more preferably less than 0.10% by weight and most preferably less than 0.08% by weight. In at least one embodiment, the carbon content ranges from a low of about 0.05%, 0.06%, 0.07% by weight to a high point of about 0.10%, 0.11%, 0.12% by weight. Preferably, the steel has a carbon content of 0.05% to 0.12% by weight.

In one or more embodiments above or elsewhere herein, the steel may comprise silicon (Si). Silicon may be added for de-acidic purposes. Silicon is also a strong matrix strengthening agent, but has a strong adverse effect on both base steel and HAZ toughness. Therefore, an upper limit of 0.5% by weight with respect to silicon is preferred. Silicon increases the driving force for carbon migration to untranslated austenite during cooling (quenching) of the steel sheet from high temperatures, and in this sense reduces the interstitial content of ferrite and improves its flow and ductility. This beneficial effect of silicon must be counterbalanced with the original effect of silicon that degrades the toughness of the steel. Due to this balancing force, the optimum silicon addition in the alloy of the present invention is about 0.01% to 0.5% by weight.

In one or more embodiments above or elsewhere herein, the steel may comprise manganese (Mn). Manganese can be a matrix strengthening agent in steel and more importantly, can contribute to hardenability. Manganese is an inexpensive alloying additive that prevents excessive ferrite formation that can lead to a decrease in plate strength in thick cross-section plates, especially at the thickness-middle positions of such plates. Manganese, through its powerful effect of retarding austenite's ferrite, pearlite, granular bainite and upper bainite transformation products while the austenite cools, can produce fines such as las martensite, lower bainite and modified upper bainite. Provides processing flexibility to produce alternative strong second phases in the tissue. However, because too much manganese is detrimental to the toughness of the steel sheet, an upper limit of about 2.0% by weight manganese is preferred. This upper limit is also desirable to substantially minimize the center segregation that tends to occur in high manganese and continuously cast steel slabs and incidental bad microstructure and toughness properties at the center of the plates produced from the slabs. Preferably, the steel has an Mn content of about 0.5% to about 2.0% by weight.

Preferably, the rest is minimized. For example, preferably the sulfur (S) content is less than about 0.01% by weight. Preferably the phosphorus (P) content is less than about 0.015% by weight. More preferably, the P content is less than 0.01% by weight. In one or more embodiments the P content, if present, is from 0.0001% to 0.009% by weight.

In one or more embodiments above or elsewhere herein, the steel may comprise niobium (Nb). Niobium can be added to promote grain refinement during hot rolling of steel slabs into plates, which in turn improves both the strength and toughness of the steel sheet. Niobium carbide precipitation during hot rolling serves to retard recrystallization and suppress grain growth, thereby providing a means of austenite grain refinement. For these reasons, at least 0.005% by weight of niobium is preferred. Niobium is also a strong hardening enhancer and provides precipitation strengthening in HAZ through the formation of niobium carbide or carbonitelite. This effect of adding niobium to the steel is useful for minimizing HAZ softening in high strength steel weldments, especially after the melting line. For this reason, a minimum amount of niobium of 0.01% by weight is more preferred in steel sheets to be welded during manufacture of useful objects such as linepipes. However, more niobium can lead to excessive precipitation hardening and consequently lower toughness in both base steel and especially HAZ. For these reasons, an upper limit of 0.03% by weight is preferable. More preferably, the upper limit is 0.02 weight%.

In one or more embodiments above or elsewhere herein, the steel may comprise titanium (Ti). Titanium is effective in forming fine titanium nitride (TiN) precipitates that refine the grain size in both the rolled structure of the steel and the HAZ. Thus, the toughness of the steel and HAZ is improved. A minimum value of 0.005% by weight titanium is preferred for this purpose. Titanium is added to the steel in an amount such that the weight ratio of Ti / N is preferably about 3.4. Excessive titanium addition to steel tends to lower the toughness of the steel by forming coarse TiN particles or titanium carbide particles. Therefore, the upper limit of titanium is preferably 0.02% by weight.

In one or more embodiments above or elsewhere herein, the steel may comprise aluminum (Al). Aluminum may be added primarily for deoxidation of the steel. At least 0.01% by weight of aluminum is preferred for this purpose. Small amounts of aluminum in the steel also benefit HAZ properties by tying up free nitrogen resulting from the decomposition of nitride and carbon nitride particles in the coarse grain HAZ due to the intense thermal cycle of the welding process. However, aluminum is similar to silicon in reducing the deformation and toughness properties of the matrix. In addition, the addition of higher amounts of aluminum leads to excessive, coarse aluminum-oxide inclusions in steels that degrade toughness. Therefore, an upper limit of 0.1% by weight is preferred for aluminum addition to the steel.

In one or more embodiments above or elsewhere herein, the steel may comprise nitrogen (N). Nitrogen can form TiN precipitates to inhibit coarsening of the austenite grains during slab reheating and in the HAZ, thereby improving the low temperature toughness of the base metal and HAZ. For this effect, a minimum amount of 0.0015% by weight nitrogen is preferred. However, too much nitrogen addition can lead to excessive free nitrogen in HAZ and degrade HAZ toughness. Excessive free nitrogen may also increase the tendency for strain aging in line pipes. For this reason, the upper limit for nitrogen is preferably 0.01% by weight, more preferably 0.005% by weight.

In one or more embodiments above or elsewhere herein, the steel has a nitrogen content of less than 0.01%, more preferably less than 0.0075%, and most preferably less than 0.005% by weight. Preferably, the nitrogen content ranges from a low point of about 0.0025%, 0.0035%, or 0.0045% by weight to a high point of 0.0050%, 0.0075%, or 0.01% by weight. More preferably, the steel has a nitrogen content of about 0.0025% to about 0.0095% by weight.

In one or more embodiments above or elsewhere herein, the steel may comprise nickel (Ni). Nickel can improve the toughness of base steel as well as HAZ. A minimum value of 0.1% by weight nickel, and more preferably a minimum value of 0.3% by weight nickel, is preferred to create a significantly beneficial effect on HAZ and base steel toughness. Although not to the same extent as manganese and molybdenum additions, the addition of nickel to steel improves the hardenability and therefore the uniformity through thickness in the microstructure and the properties of thick cross sections (20 mm and larger). However, excessive nickel addition can compromise field weldability (cause cold cracking), promote hard microstructures, reduce HAZ toughness, and increase the cost of steel. Preferably, the steel has a nickel content of about 1% by weight or less.

In one or more embodiments above or elsewhere herein, the steel has or does not have a reduced amount of substituted alloy elements such as, for example, chromium, molybdenum, vanadium and copper. Such elements lower the activity of carbon and nitrogen on the ferrite of steel or result in excessive precipitation hardening, which increases the tendency for strain aging. The combined amount of molybdenum, chromium, vanadium and copper is about 0.20% or less, about 0.15% or less, about 0.12% or less or about 0.10% or less.

In one or more embodiments above or elsewhere herein, the steel may comprise boron (B). Boron can greatly increase the hardenability of the steel very inexpensively and can promote the formation of steel microstructures of lower bainite, ras martensite at thick cross sections (> 16 mm). Boron enables the design of steels with low alloys and Pcm (welcome hydrogen crack sensitivity parameters based on composition) as a whole, thereby improving HAZ softening resistance and weldability. Boron in excess of about 0.002% by weight may enhance the formation of embrittling particles of Fe ₂₃ (C, B) ₆ . Therefore, when boron is added, an upper limit of 0.002% by weight boron is preferable. Boron also increases the hardenability effects of molybdenum and niobium.

In one or more embodiments above or elsewhere herein, the steel may comprise chromium (Cr). Chromium may have a strong effect of increasing hardenability of the steel when directly quenched. As such, chromium is a cheaper alloying additive than molybdenum, particularly for improving hardenability in added boron-free steels. Chromium improves corrosion resistance and hydrogen induced crack resistance (HIC). Similar to molybdenum, excessive chromium tends to cause cold cracking in weldments and worsens the toughness of the steel and its HAZ. Chromium can lower the carbon activity in ferrite and thereby lead to an increase in the amount of carbon in solid solution, which can increase the steel tendency for strain aging. Thus, when chromium is added a maximum of 0.2% by weight is preferred and even a maximum of 0.1% by weight is even more preferred.

In one or more embodiments above or elsewhere herein, the steel may comprise REM (rare earth metal). Calcium and REM form sulfides to inhibit the generation of elongated MnS and to improve the properties of the steel sheet (eg lamellar cracking properties). However, addition of Ca and REM in excess of 0.01% forms CaO-CaS or REM-CaS, resulting in poor steel cleanliness and field weldability. Preferably, up to 0.02% by weight of REM is added.

In one or more embodiments above or elsewhere herein, the steel may comprise magnesium (Mg). In general, magnesium forms finely dispersed oxide particles, which can inhibit grain coarsening and / or enhance the formation of intra-granular ferrite in HAZ, thereby improving HAZ toughness. . At least about 0.0001 wt.% Mg is preferred for the addition of Mg to be effective. However, when the Mg content exceeds about 0.006% by weight, coarse oxides are formed and the toughness of the HAZ deteriorates. Therefore, the content of Mg is preferably less than 0.006% by weight.

In one or more embodiments above or elsewhere herein, the steel may comprise copper (Cu). Copper can contribute to steel reinforcement by increasing hardenability and through strong precipitation reinforcement by ε-copper precipitates. In larger amounts, copper can lead to excessive precipitation hardening and, if not properly controlled, can degrade the toughness of the base steel as well as the HAZ. In addition, more copper may cause embrittlement during slab casting and hot rolling, and it is necessary to add nickel together to alleviate this. When copper is added to the steel of the present invention, an upper limit of 0.2% by weight is preferable, and an upper limit of 0.1% by weight is even more preferable.

In one or more embodiments above or elsewhere herein, the steel may comprise vanadium (V). Vanadium is substantially similar to niobium but does not have as strong an effect as niobium. However, the addition of vanadium has a significant effect when added with niobium. The combined effect of vanadium and niobium significantly minimizes HAZ softening during welding, where high heating inputs, such as seam welding in pipeline manufacturing, are input. Like niobium, excessive vanadium can degrade the toughness of both HAZ as well as the base steel through excessive precipitation hardening. Moreover, vanadium, like chromium and molybdenum, has a strong affinity for carbon and nitrogen. In other words, vanadium can lower the carbon activity in ferrite and cause an increase in the amount of carbon and nitrogen in solid solution, which can increase the tendency of the steel to strain aging. Thus, vanadium, when added to steel, is preferably less than about 0.1 weight percent or more preferably less than about 0.05 weight percent or even more preferably less than about 0.03 weight percent.

In one or more embodiments above or elsewhere herein, the steel may comprise zirconium (Zr), hafnium (Hf) and / or tantalum (Ta). Zirconium (Zr), hafnium (Hf) and / or tantalum (Ta) are elements that are effective in forming carbides and nitrides and increasing strength, such as niobium (Nb). However, the effect cannot be realized with an addition of less than 0.0001% by weight. However, when more than 0.05% by weight, the toughness of the steel sheet is poor. Accordingly, the Ta content is preferably less than about 0.03% by weight, the Zr content is preferably less than about 0.03% by weight and the Hf content is preferably less than about 0.03% by weight.

Preferably, the steel has a Pcm less than 0.220 but greater than 0.150. Pcm refers to a method of measuring the hardenability and weldability of steel based on its chemical composition. Higher concentrations of carbon and other alloying elements (eg, Mn, Cr, Si, Mo, V, Cu, Ni) tend to increase hardness and reduce weldability of steel. Since each of these materials tends to affect steel hardness and weldability in different sizes, Pcm is a means of determining the difference in hardness / welding between alloys made from different alloying elements.

Commonly used formulas to calculate Pcm are:

Pcm = C + Si / 30 + (Mn + Cu + Cr) / 20 + Ni / 60 + Mo / 15 + V / 10 + 5B.

Microstructure

In one or more embodiments, the steel may comprise from 10% to 90% by volume of the softer, ferrite phase or component ("first phase") and from about 10% to about 90% by volume of the stronger phase or component ("agent" Biphasic "). The second phase may comprise one or more phases or components that are not ferrite. Exemplary non-ferrite phases or components include, but are not limited to, carbides such as martensite, lower bainite, modified upper bainite, upper bainite, granular bainite, pearlite, cementite and the like and mixtures thereof.

The Ar ₁ transformation temperature refers to the temperature at which austenite to ferrite or ferrite plus cementite is completed during cooling.

Ar ₃ transformation temperature refers to the temperature at which austenite begins to transform into ferrite during cooling.

Cooling rate refers to the cooling rate in the middle or substantially in the middle of the sheet thickness.

By complex tissue is meant at least two distinguishable phases or at least two distinguishable components.

Granular bainite (GB) is a cluster of three to five relatively equiaxed bainitic ferrite grains surrounding the small "island" located in the middle of martensite-austenite (MA). Refers to. Typically the “grain” diameter is about 1-2 μm.

Upper bainite (UB) refers to a mixture of acicular or lath phases of bainitic ferrite interspersed with film or horizontal material on carbides such as cementite.

Denatured upper bainite (DUB) is a bainite product in which each colony grows into a set (packets) of parallel laths by shear stress. During and immediately after lath growth, some carbon is discarded into interlas austenite. Due to the relatively low carbon content, the carbon enrichment of entrapped austenite is not sufficient to cause cementite plate nucleation. Such nucleation occurs in medium and high carbon steels, which results in the formation of traditional upper bainite (UB). On the other hand, low carbon concentration in interlath austenite in the DUB results in the formation of martensite or martensite-austenite (MA) mixture or remains as residual austenite (RA).

DUB can be confused with traditional upper bainite (UB). The type of UB first identified decades ago in heavy carbon steels has two main characteristics; (1) a set of parallel laths growing into a packet, and (2) a cementite film at the lath boundary. UB is similar to DUB in that both contain packets of parallel classes; However, the main difference is in the interlas material. When the carbon content is about 0.15-0.40, cementite (Fe ₃ C) may be formed between the laths. Such "film" may be relatively continuous compared to intermittent MA in the DUB. For low carbon steels, no interlas cementite is formed; Rather the residual austenite is terminated with MA, martensite or RA.

Lower bainite LB has a packet of parallel laths. LB also contains less intra lath carbide precipitates. These plate-shaped particles are consistently precipitated on a single crystallographic variant oriented about 55 ° from the main lath growth direction (long dimension of the lath).

Lath martensite (LM) looks like a packet of thin parallel laths. The lath width is typically less than about 0.5 μm. The untempered colony of martensite laths is characterized by the absence of carbides, while the auto-tempered LMs represent intra lath carbide precipitates. Intra lath carbides in the auto-tempered LM are formed on more than one crystallographic variant, for example on the '110' plane of martensite. Usually in transmission electron microscopy (TEM) of autotempered LM, cementite is not aligned along one direction, but rather precipitates on multiple planes.

Pearllite is typically a two-phase lamellae mixture composed of alternating layers of ferrite and cementite (Fe ₃ C).

Grains are individual crystals in the polycrystalline material.

Grain boundaries correspond to transitions from one crystallographic direction to another and thus refer to a narrow region in the metal that separates one grain from another.

Previous austenite grain size refers to the average austenite grain size in a hot-rolled steel plate before rolling in a temperature range where austenite does not recrystallize.

Quenching refers to accelerated cooling by any means, whereby the fluid chosen for the tendency to increase the cooling rate of the steel is used in contrast to air cooling.

Accelerated Cooling End Temperature (ACFT) is the highest or substantially the highest temperature reached at the surface of the plate after quenching, because of the heat transferred from the mid-thickness of the plate.

A slab is a piece of steel with arbitrary dimensions.

The T _m temperature is a temperature below which austenite does not recrystallize.

The transverse direction refers to the direction in the plane of rolling but perpendicular to the sheet rolling direction.

Manufacturing method

In one or more embodiments, the steel composition is processed in a manner that reduces the amount of C and / or N supersaturation on the ferrite of the composite steel obtained therefrom. Preferably, the steel is processed under conditions such that C and N diffuse and / or precipitate out of the ferrite during sheet processing. Diffusion and precipitation can be performed through a high accelerated cooling end temperature while continuing to maintain all desired microstructure characteristics of the composite microstructure design (eg, the amount of softer ferrite phase, the effective previous austenite grain size, etc.). Can be. In one or more embodiments, the volume percentage of ferrite in the steel is about 10% to about 90% by weight, more preferably 30% to 80% by weight. Preferably, the ferrite is uniformly dispersed throughout the steel.

Preferably the steel composition is processed into a composite tissue plate using a two step rolling process. In one or more embodiments, the steel billets / slabs from the described compositions are first molded in a normal manner, for example through a continuous casting process. The billet / slab may then be reheated to a temperature (“reheat temperature”) in the range of about 1000 ° C. to about 1250 ° C. Preferably, the reheating temperature is (i) substantially homogenizing the steel slab, (ii) decomposing substantially all carbides and carbonnitrides of niobium and vanadium, if present, and (iii) It is enough to establish early austenite grains. The reheated slabs are then hot rolled in one or more passes at a first reduction that provides about 30% to about 90% reduction in the first temperature range where austenite is recrystallized. The pressed billet is then hot rolled in one or more passes at a second rolling reduction where austenite does not recrystallize but provides about 40% -80% reduction in the second and somewhat lower temperature ranges above the Ar ₃ transformation point. Preferably, the cumulative rolling reduction below the Tnr temperature is at least 50%, more preferably at least about 70%, even more preferably at least 75%.

The second rolling reduction is completed at the "end rolling temperature". In one or more embodiments, the finish rolling temperature is higher than 700 ° C, preferably higher than 720 ° C, more preferably higher than 770 ° C. In one or more embodiments, the finish rolling temperature is about 700 to 800 ° C. Thereafter, the hot rolled plate is cooled (eg, cooled in air) to a controlled cooling start temperature ("ACST") or to a first cooling temperature that is sufficient to induce austenite to ferrite transformation, and then Control cooling at a rate of at least 10 ° C. per second is followed by a second cooling temperature or controlled cooling end temperature (“ACFT”). After ACFT, the steel sheet can be cooled to room temperature (ie, ambient temperature) in ambient air. Preferably, the steel sheet can be allowed to cool itself to room temperature.

In one or more embodiments, the ACST is at least about 600 ° C, at least about 650 ° C, at least about 700 ° C, or at least about 730 ° C. In one or more embodiments, the ACST may be about 600 ° C to about 800 ° C. In one or more embodiments, the ACST may be between about 650 ° C and about 750 ° C. Preferably, the ACST is a low point of about 650 ° C, 660 ° C or 690 ° C to a high point of about 700 ° C, 730 ° C or 750 ° C. In one or more embodiments, the ACST is about 650 ° C, about 660 ° C, about 670 ° C, about 680 ° C, about 690 ° C, about 700 ° C, about 710 ° C, about 720 ° C, about 730 ° C, about 740 ° C or about 750 ° C. May be ° C.

In one or more embodiments, the ACFT may be about 400 ° C to about 700 ° C. In one or more embodiments, the ACFT may be about 450 ° C to about 650 ° C. Preferably, the ACFT is from a low point of about 400 ° C, 450 ° C, or 500 ° C to a high point of about 550 ° C, 600 ° C or 650 ° C. For example, the ACFT is about 505 ° C, about 510 ° C, about 515 ° C, about 520 ° C, about 525 ° C, about 530 ° C, about 535 ° C, about 540 ° C, about 545 ° C, about 550 ° C, or about 575 ° C. Can be. In one or more embodiments, the ACFT may be between about 540 ° C. and about 560 ° C.

Without wishing to be bound by theory, it is believed that a high accelerated cooling termination rate ("ACFT") causes at least some of the carbon and nitrogen atoms to diffuse from the ferrite phase of the steel composition into the second phase. In addition, high accelerated cooling end rates (“ACFTs”) allow at least some of the carbon and nitrogen atoms to precipitate as carbides, carbon nitrides and / or nitrides from the ferrite phase during subsequent cooling from the ACFT to ambient temperature. I'm thinking. In that position, the amount of free C and N in the interstice on the ferrite is reduced and the amount of C and N available to move to the potential in the ferrite is reduced. Therefore, if not removed, the tendency of the steel to strain aging is reduced.

Following the rolling and cooling steps, the plates can be formed into pipes (eg linepipes). Any method for forming the pipe can be used. Preferably, the precursor steel sheet can be made into linepipe by conventional UOE processes well known in the art.

FBE process

After forming the pipe, the pipe can then be subjected to a coating / painting step to prevent corrosion and / or mechanical damage. The coating process may include one or more polymer coatings applied to at least the outer diameter or surface of the pipe. The coating can also be applied to both the inner and outer surfaces of the pipe. Exemplary coatings include, but are not limited to, melt adhesive epoxy (FBE), polypropylene, polyethylene, and polyurethane. Preferably a melt adhesive epoxy (FBE) is added. FBE is a thermosetting polymer that can be sprayed and heat cured onto a pipe using known techniques. Preferably, at least one layer of FBE is applied or sprayed onto the pipe. In one or more embodiments, each layer of the coating has a thickness of about 2 microns to 75 mm. In one embodiment, the pipe can be heated and rotated while applying the spray powder. In another embodiment, the pipe can be heated and immersed in a fluidized bed containing the polymer. Preferably the pipe is heated to a temperature of about 180 ° C to about 300 ° C. One or more other coatings may then be applied to at least a portion of the pipe over the FBE layer.

As mentioned above, post-treatment steps, such as the FBE application process, facilitate the diffusion of supersaturated carbon and nitrogen atoms, leading to the formation of a solute atmosphere around the dislocations in the steel. The formation of such a solute atmosphere (“Cottrell atmosphere”) increases the strength of the steel, but reduces the ductility, because more deformation or force is needed to escape the atmosphere from the dislocation. As a result, the steel may be unsuitable for use in areas where ductility is smaller and high strain capacity is required.

End use

Steel sheets made in accordance with the described embodiments retain all of the desired microstructure features of the composite microstructure design, while minimizing carbon supersaturation on the ferrite. Such DP steel can be easily implemented in applications where both high strength and high deformation capacity are required. For example, steel is particularly useful as a precursor for making line pipes or pressure vessels. Steel can also be used for offshore structures, including risers, oil and gas production facilities, shipbuilding, automobile manufacturing, aircraft manufacturing and power generation.

In order to facilitate a better understanding of the present invention, the following examples of certain aspects of some embodiments are given. The following examples should never be read as limiting or limiting the scope of the invention.

Yes:

The previous discussion may be discussed further with reference to the following non-limiting examples.

Four steel precursors (steels A, B, C, D and E) were prepared from heat treatments having the chemical compositions shown in Table 1. Each precursor is produced either by vacuum induction, which is melted into a 150 kg heat treatment and slab, or by continuous casting into a steel slab using a 250 ton industrial basic oxygen furnace. Steel sheets (Examples 1-8) were made from these steel precursors (steels A, B, C, D and E) according to the process conditions summarized in Table 2. Examples 1-7 show inventive steels, while Example 8 shows comparative or conventional DP steels.

After the steel is processed into precursor plates, the steel sheet is formed into line pipes. Heat-treatment was performed for 5-8 minutes at 200 ° C. to 250 ° C. for the molded pipe to simulate the FBE coating processes listed in Table 3. The term "As-UOE" as used in Table 3 refers to a linepipe at room temperature, ie without heat treatment. Mechanical properties of the pipe relative to the longitudinal direction are measured and reported in Table 3.

1-4 show changes in the mechanical properties listed in Table 3 as a function of heat treatment temperature. In particular, FIGS. 1 and 2 show that the inventive steel (Examples 3-7) showed much improved strain aging resistance than the comparative steels, Example 8, i.e. lower YR values (Figure 1) and higher uniform elongation (Figure 2). do. At all times, the inventive steels (Examples 3-7) exhibit good, consistent yield strength (FIG. 3) and tensile strength (FIG. 4). In that sense, DP steels produced according to the described examples do not undergo significant strain aging in contrast to comparative DP steels (Example 8).

5 shows the relationship between yield ratio (%) as a function of heat treatment temperature for steel produced according to the described embodiments (eg, Example 1-7) and comparative steel (eg, Example 8). do. Curve 510 represents Example 8 and curve 520 is Example 6. As shown in FIG. 5, the inventive steel 520 has much improved strain aging resistance compared to conventional DP steel 510, ie, a temperature range typical for FBE coating processes (eg, about 200 ° C. to about 250 ° C.). Lower yield ratio at.

6A is an SEM image of the steel produced, for example, in Example 8. FIG. 6B is a TEM image of Example 8 at a quarter thickness. In both FIGS. 6A and 6B, the steel was heat treated for FBE coating simulation according to the conditions listed in Table 3. The steel has a first phase 600 of ferrite and a second phase of mainly granular bainite (GB) 605 and modified upper bainite (DUB) 610. Referring to FIG. 6B, the dislocations 650 in the ferrite appear mainly straight, slightly twisted, indicating that these dislocations 650 have less mobility during deformation. In such a position, higher energy or greater force was needed to move or tear dislocations 650. Therefore, such additional forces increased the strength of the steel, but reduced the ductility as shown in Table 3.

FIG. 7A shows an SEM image of inventive steel, Example 5 (Steel D according to 566 ° C. cooling end temperature), at quarter thickness. 7b shows a TEM image of the same steel. In addition, both SEM and TEM are images after heating the steel and simulating the FBE coating process according to the conditions listed in Table 3. FIG. 7A shows the second phase of the steel primarily granular bainite (GB) 705, upper bainite or pearlite 710 and some las martensite (LM) 720. The TEM image (not shown) of the steel shown in FIG. 7A actually reveals the component indicated at 710, perhaps as perlite. FIG. 7B shows that the dislocations 850 are in the form of entangled, curved and / or wave, which exhibits high mobility of these dislocations during deformation. In other words, less force was needed to move C and / or N atoms from dislocations 850. Therefore, the ductility of the steel was increased and the tensile strength was not affected as shown in Table 3.

As shown in FIGS. 1-7 and Table 3, steels B, C, D and E machined in accordance with the embodiments described herein, respectively, compared to steel A machined according to Example 8, to maintain tensile strength But contains increased carbon and manganese content, but is much less affected by strain aging. Increased carbon content in steels B, C, D and E would have been expected to negatively affect strain aging. Surprisingly, I found out that the truth is the opposite. In addition, the combination of increased carbon and manganese content of the steels B, D and E would have been expected to negatively affect the strain aging even more than the increased carbon content alone. Surprisingly, I found out that the truth is the opposite. Therefore, it is believed that the absence of carbon-cluster forming alloys and / or cold end temperatures higher than 528 ° C. surprisingly provide composite steel pipes with good tensile and yield strength in addition to high resistance to strain aging.

Particular embodiments and features have been described using a set of upper and lower numerical limits. Unless otherwise indicated, it should be recognized that a range of any lower limit to any upper limit is contemplated. Certain lower limits, upper limits and ranges are set forth in one or more claims below. All numerical values are those indicated as "about" or "approximately" and take into account experimental errors and changes expected by one of ordinary skill in the art.

Various terms have been defined above. If the terms used in the claims are not defined above, such terms should be given the broadest definition given to those skilled in the art as reflected in at least one published publication or published patent. Moreover, all patents, test procedures and other documents cited in this application are fully cited to the extent such disclosure is inconsistent with this application and to all jurisdictions in which citation is permitted.

While the foregoing is directed to embodiments of the invention, other and further embodiments of the invention may be devised without departing from the basic scope thereof, and the scope thereof is determined by the claims that follow.

Therefore, the present invention is well-modified to carry out the objects and achieve the results and advantages mentioned as well as the original ones herein. Although the present invention has been depicted and described with reference to exemplary embodiments of the invention, such references do not include limitations to the invention and such limitations should not be inferred. The invention is susceptible to significant modifications, changes, and equivalents in form and function, which would be contemplated by one of ordinary skill in the art and having the benefit of this disclosure. The described and described embodiments of the invention are exemplary only and are not exhaustive of the scope of the invention. As a result, the invention is intended to be limited only by the spirit and scope of the appended claims, which fully recognize the equivalents in all respects. The terms in the claims have the plain, ordinary meaning unless otherwise specified and clearly defined by the patentee.

Claims (55)

About 0.05% to about 0.12% carbon; From about 0.005% to about 0.03% niobium; About 0.005% to about 0.02% titanium; About 0.001% to about 0.01% nitrogen; About 0.5% to about 2.0% manganese; About 0.01% to about 0.5% silicon; Less than about 0.15% by weight of molybdenum, chromium, vanadium and copper; A first phase consisting of ferrite; A second phase comprising one or more components selected from the group consisting of carbide, pearlite, martensite, lower bainite, granular bainite, upper bainite and modified upper bainite; And And about 0.01% by weight or less of solute carbon content in the first phase. The composite tissue steel according to claim 1, wherein said second phase comprises pearlite. The composite tissue steel of claim 1, wherein nickel is present in an amount of less than about 1.0 weight percent. 2. The composite tissue steel of claim 1, wherein boron is present in an amount of less than about 0.02% by weight. The composite tissue steel of claim 1, wherein the steel has a Pcm of less than about 0.22. The composite tissue steel according to claim 1, wherein the steel has a tensile strength of at least 500 MPa. The composite tissue steel of claim 1, wherein the steel has a tensile strength of at least 520 MPa. The steel of claim 1, wherein the steel has a minimum uniform elongation of at least 8%. The steel of claim 1, wherein the steel has a minimum uniform elongation of at least 10%. The composite tissue steel according to claim 1, wherein the steel has a yield ratio of less than 0.90. The composite tissue steel according to claim 1, wherein the steel has a yield ratio of less than 0.85. Heating the steel slab to a reheating temperature of about 1000 ° C. to about 1250 ° C .; Pressing the steel slab to form a plate in at least one hot rolling pass at a first temperature; Pressing the plate in at least one hot rolling pass at a second temperature; Cooling the plate to a first cooling temperature sufficient to transform austenite into ferrite; And Reducing cluster forming atoms in the ferrite, The heating of the steel slab at the reheat temperature provides a steel slab consisting essentially of austenite phase; The first temperature is such as to recrystallize the austenite phase; The austenite phase does not recrystallize at the second temperature; And And said second temperature is lower than said first temperature. 13. The method of claim 12, wherein the cluster forming atoms comprise carbon. 13. The method of claim 12, wherein the cluster forming atoms comprise nitrogen. 13. The method of claim 12, wherein the cluster forming atoms comprise carbon and nitrogen. The method of claim 12, wherein reducing the cluster forming atoms in the ferrite comprises quenching the cooled plate to a second cooling temperature at a rate of at least 10 ° C. per second. The method of claim 12, wherein the first cooling temperature is about 650 ° C. to about 750 ° C. 13. The method of claim 12, wherein the first cooling temperature is about 660 ° C. to about 750 ° C. 13. The method of claim 12, wherein the first cooling temperature is about 670 ° C. to about 740 ° C. 13. The method of claim 12, wherein the first cooling temperature is about 730 ° C. 13. The method of claim 16, wherein the second cooling temperature is about 400 ° C. to about 700 ° C. 18. The method of claim 16, wherein the second cooling temperature is about 450 ° C. to about 650 ° C. 18. The method of claim 16, wherein the second cooling temperature is about 500 ° C. to about 600 ° C. 18. The method of claim 16, wherein the second cooling temperature is about 560 ° C. 18. The method of claim 12, wherein the rolled plate comprises from about 10% to about 90% by volume of ferrite. The method of claim 12, wherein the rolled plate comprises from about 10% to about 90% by volume of the second phase. 27. The method of claim 26, wherein the second phase comprises one or more components selected from the group consisting of carbide, pearlite, martensite, lower bainite, granular bainite, upper bainite and modified upper bainite. Method for manufacturing composite tissue steel. 17. The method of claim 16, further comprising cooling the steel sheet to ambient temperature after quenching to the second cooling temperature. 13. The method of claim 12, further comprising forming the cooled plate into linepipes using UOE techniques. 30. The method of claim 29, further comprising applying a coating for corrosion resistance to at least a portion of said linepipe. 31. The method of claim 30, wherein the coating comprises at least one melt adhesive epoxy compound. Heating the steel slab to about 1000 ° C. to about 1250 ° C. to provide a steel slab consisting essentially of the austenite phase;  Pressing the steel slab to form a plate with at least one hot rolling pass at a temperature sufficient to recrystallize the austenite phase to produce an austenite phase of microparticles; Pressing the plate with at least one hot rolling pass at a temperature lower than the temperature at which austenite does not recrystallize; Cooling the plate to a first temperature sufficient to transform austenite into ferrite; Quenching the plate to a second temperature at a rate of at least 10 ° C. per second (18 ° F./sec); And Cooling the plate at a rate such that the solute carbon in the ferrite is reduced. 33. The method of claim 32, wherein the second temperature is such as to diffuse the carbon from the ferrite into the second phase. 33. The method of claim 32, wherein the second temperature is such that the carbon in the ferrite is precipitated with one or more carbides. 34. The method of claim 33, wherein the second phase comprises one or more components selected from the group consisting of carbide, pearlite, martensite, lower bainite, granular bainite, upper bainite and modified upper bainite. Method for manufacturing composite tissue steel. 33. The method of claim 32, further comprising forming the cooled plate into linepipes using UOE techniques. 33. The method of claim 32, wherein the first temperature is about 650 ° C to about 750 ° C. 33. The method of claim 32, wherein the first temperature is about 670 ° C to about 740 ° C. 33. The method of claim 32, wherein the second temperature is about 400 ° C to about 700 ° C. 33. The method of claim 32, wherein the second temperature is about 450 ° C to about 650 ° C. 33. The method of claim 32, wherein the second temperature is about 500 ° C to about 600 ° C. 33. The method of claim 32, wherein the second temperature is about 560 ° C. 33. The method of claim 32, Forming a line pipe from the plate; Heating the line pipe to a temperature of about 180 ° C. to 300 ° C .; And And applying at least one coating to at least a portion of said linepipe. 44. The method of claim 43, wherein said at least one coating comprises one or more melt adhesive epoxy compounds. About 0.05% to about 0.12% carbon; From about 0.005% to about 0.03% niobium; About 0.005% to about 0.02% titanium; About 0.001% to about 0.01% nitrogen; About 0.5% to about 2.0% manganese; About 0.01% to about 0.5% silicon; Less than about 0.10 weight percent of molybdenum, chromium, vanadium and copper; A first phase consisting of ferrite; A second phase comprising one or more components selected from the group consisting of carbide, pearlite, martensite, lower bainite, granular bainite, upper bainite and modified upper bainite; And And about 0.01% by weight or less of solute carbon content in the first phase. 46. The composite tissue steel according to claim 45, wherein nickel is present in an amount of less than about 1.0 weight percent. 46. The composite tissue steel according to claim 45, wherein boron is present in an amount of less than about 0.02% by weight. 46. The composite tissue steel of claim 45, wherein the steel has a Pcm of less than about 0.22. 46. The composite tissue steel of claim 45, wherein the steel has a tensile strength of at least 500 MPa. 46. The composite tissue steel of claim 45, wherein the steel has a tensile strength of at least 520 MPa. 46. The composite tissue steel of claim 45, wherein the steel has a minimum uniform elongation of at least 8%. 46. The composite tissue steel of claim 45, wherein the steel has a minimum uniform elongation of at least 10%. 46. The composite tissue steel of claim 45, wherein the steel has a yield ratio of less than 0.90. 46. The composite tissue steel of claim 45, wherein the steel has a yield ratio of less than 0.85. 46. The composite tissue steel according to claim 45, wherein said second phase comprises pearlite.

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