CN111492083A

CN111492083A - High-strength steel material having excellent hydrogen-induced cracking resistance and low-temperature impact toughness, and method for producing same

Info

Publication number: CN111492083A
Application number: CN201880077488.4A
Authority: CN
Inventors: 严庆根; 金大优
Original assignee: Posco Co Ltd
Current assignee: Posco Holdings Inc
Priority date: 2017-12-01
Filing date: 2018-07-27
Publication date: 2020-08-04
Anticipated expiration: 2038-07-27
Also published as: EP3719162A4; KR20190065040A; KR102020434B1; JP2021504582A; EP3719162A1; CN111492083B; WO2019107700A1; JP7197582B2

Abstract

The present invention provides a high strength steel material having excellent hydrogen induced cracking resistance and low temperature impact toughness, the steel material comprising in weight%: 0.02% to 0.12% C, 0.02% to 0.60% Si, 0.6% to 1.6% Mn, 0.020% or less (excluding 0%) P, 0.003% or less (excluding 0%) S, 0.002% to 0.060% soluble Al, 0.005% to 0.6% Cu, 0.005% to 1.0% Ni, 0.005% to 0.5% Cr, 0.01% to 0.3% Mo, 0.001% to 0.10% Ti, 0.001% to 0.06% Nb, 0.001% to 0.10% V, 0.001% to 0.006% N, 0.0002% to 0.0060% Ca, and the balance Fe and unavoidable impurities, wherein (Ti + V +0.5Nb +0.5Mo) is 0.05% or more and (Ti + V + 0.5) is 0.05% or more and 0.5% Mo +0.5 + C + 0.8% ferrite or more; and at least one of pearlite, cementite, and MA (martensite-austenite composite phase), wherein the fraction of ferrite is 90 area% or more (excluding 100%), and the fraction of at least one of ferrite, cementite, and MA (martensite-austenite composite phase) is 10 area% or less, and wherein the ferrite structure contains 0.03 wt% or more of fine carbonitride precipitates having an average size of 100nm or less and the ferrite structure has an average grain size of 20 μm or less.

Description

High-strength steel material having excellent hydrogen-induced cracking resistance and low-temperature impact toughness, and method for producing same

Technical Field

The present disclosure relates to a high-strength steel material used in plant facilities for mining, processing, transporting and storing crude oil, as well as shipbuilding and offshore structures, and a method for manufacturing the same, and more particularly, to a high-strength steel material having excellent hydrogen induced cracking resistance (HIC resistance) and low-temperature impact toughness, and a method for manufacturing the same.

Background

Due to depletion of petroleum resources and high oil prices, the production of low-quality crude oil having a high content of impurities (e.g., sulfur) is gradually increasing. Therefore, steels used in all plant facilities for mining, processing, transporting and storing low-grade crude oil are also required to have the property of suppressing hydrogen-induced cracking (HIC) caused by wet hydrogen sulfide generated due to sulfur components in crude oil.

In addition, environmental pollution due to recent safety accidents has become a global problem, and astronomical digital costs are required to recover it. Therefore, the level of required characteristics of steel materials used in the energy industry is gradually increasing.

As examples of materials for plant facilities for mining, processing, transporting and storing crude oil, and shipbuilding and offshore structures, high-strength normalized steel materials have been proposed.

As a method for obtaining high strength of the normalized steel material, increasing the pearlite fraction by adding C, adding solid-solution strengthening elements such as Cu and Mo, and generating carbonitride precipitates by simply adding Nb and V have been used.

Meanwhile, as a method for improving the HIC resistance of a steel material, controlling internal defects such as inclusions, voids, etc. in the steel material, which may cause hydrogen enrichment and cracking, has been widely proposed and used; the strong rolling is applied to minimize a hardened structure such as pearlite and the like in which cracks easily occur and spread or to control the shape thereof and the like.

However, as the demand for higher strength of steel materials is gradually increased, the amount of C or added solid solution strengthening component added for increasing strength is also increased. It is contradictory that an increase in the amount of C added serves to increase the pearlite fraction and make pearlite have a band structure (bonded structure), thereby promoting the expansion of HIC and significantly reducing low-temperature impact toughness.

Further, addition of a solid solution strengthening component such as copper (Cu), nickel (Ni), molybdenum (Mo), etc. has an effect of increasing strength, but may cause surface cracking due to red brittleness or may cause a significant drop in economic efficiency due to its high cost.

Further, when an excessive amount of Nb and V is added to obtain a strengthening effect by precipitates, hardenability of the weld zone is greatly increased to a level higher than the reference level, which causes deterioration of toughness of the weld zone and HIC.

As described above, the conventional method has problems in that it is difficult to simultaneously secure high strength, excellent low-temperature impact toughness, and HIC resistance characteristics of the normalized steel material.

Therefore, there is a need to develop a normalized steel material having high strength, excellent low-temperature impact toughness, and HIC resistance characteristics.

Disclosure of Invention

Technical problem

A preferred aspect of the present disclosure is to provide a high-strength steel material having excellent hydrogen induced cracking resistance (HIC resistance) and low-temperature impact toughness.

Another preferred aspect is to provide a method for manufacturing a high-strength steel material having excellent hydrogen induced cracking resistance (HIC resistance) and low-temperature impact toughness.

Technical scheme

According to a preferred aspect of the present disclosure, there is provided a high strength steel material having excellent hydrogen-induced cracking resistance and low-temperature impact toughness, comprising in weight%: 0.02% to 0.12% carbon (C), 0.02% to 0.60% silicon (Si), 0.6% to 1.6% manganese (Mn), 0.020% or less (excluding 0%) phosphorus (P), 0.003% or less (excluding 0%) sulfur (S), 0.002% to 0.060% soluble aluminum (sol.al), 0.005% to 0.6% copper (Cu), 0.005% to 1.0% nickel (Ni), 0.005% to 0.5% chromium (Cr), 0.01% to 0.3% molybdenum (Mo), 0.001% to 0.10% titanium (Ti), 0.001% to 0.06% niobium (Nb), 0.001% to 0.10% vanadium (V), 0.001% to 0.006% nitrogen (N), 0.0002% to 0.0060% calcium (Ca), 0.001% to 0.10% iron (Ca), 0.5% ferrite +0.5 + Fe (Mo), and the balance 0.5 + 5% Fe + 5 + Fe + 5% ferrite (Ti) or more, wherein the steel comprises 0.5% Mo and 0.5% Fe + Ti + 5% Fe + 5% or more of ferrite (Ti + Fe + Ti; and at least one of pearlite, cementite, and a martensite-austenite composite phase (MA), wherein a fraction of ferrite is 90 area% or more (excluding 100%), and a fraction of at least one of pearlite, cementite, and the martensite-austenite composite phase (MA) is 10 area% or less, and wherein ferrite contains 0.03 wt% or more of fine carbonitride precipitates having an average size of 100nm or less and an average grain size of ferrite is 20 μm or less.

According to another preferred aspect, there is provided a method for manufacturing a high strength steel material having excellent hydrogen-induced cracking resistance and low-temperature impact toughness, the method comprising: reheating a slab to 1080 ℃ to 1300 ℃, the slab comprising, in weight%, 0.02% to 0.12% of carbon (C), 0.02% to 0.60% of silicon (Si), 0.6% to 1.6% of manganese (Mn), 0.020% or less (excluding 0%) of phosphorus (P), 0.003% or less (excluding 0%) of sulfur (S), 0.002% to 0.060% of soluble aluminum (Sol.Al), 0.005% to 0.6% of copper (Cu), 0.005% to 1.0% of nickel (Ni), 0.005% to 0.5% of chromium (Cr), 0.01% to 0.3% of molybdenum (Mo), 0.001% to 0.10% of titanium (Ti), 0.001% to 0.06% of niobium (Nb), 0.001% to 0.10% of vanadium (V), 0.001% to 0.006% of nitrogen (N), 0.001% to 0.006% of calcium (Ca), and the balance being unavoidable impurities of Fe (Ca), wherein (Ti + V +0.5Nb +0.5Mo) is 0.05% or more and (Ti + V +0.5Nb +0.5Mo)/C is 0.8 or more;

hot rolling the reheated slab to a final hot rolling temperature of 850 ℃ or higher to obtain a hot rolled steel sheet;

cooling the hot rolled steel plate to 550-750 ℃ with water; and

the cooled steel sheet is subjected to a normalizing heat treatment at 850 ℃ to 950 ℃ for 1.3 × t [ where t denotes a thickness (mm) of the steel sheet) ] + (5 minutes to 60 minutes). Provided is a method for producing a high-strength steel material having excellent hydrogen-induced cracking resistance (HIC resistance) and low-temperature impact toughness.

The method may further include subjecting the normalized heat-treated steel to a tempering heat treatment at 500 ℃ or more for 10 minutes or more.

Advantageous effects

According to a preferred aspect of the present disclosure, there may be provided a high-strength steel material: it can be effectively applied to plant equipment for the exploitation, processing, transportation and storage of crude oil; shipbuilding and offshore structures; a pressure vessel; and so on.

Drawings

FIG. 1 is a photograph of the microstructure of a steel material (comparative example 3) produced according to a conventional method.

FIG. 2 is a photograph of the microstructure of a steel material (inventive example 1) manufactured according to the present disclosure.

Fig. 3 is a photograph showing the distribution of precipitates in a steel material (comparative example 3) produced by a conventional method.

Fig. 4 is a photograph showing the distribution of precipitates in a steel material (invention example 1) produced according to the present disclosure.

Detailed Description

Hereinafter, preferred embodiments are described.

However, the preferred embodiments of the present disclosure are provided so that this disclosure will be thorough and complete, and will fully convey the scope of the disclosure to those skilled in the art.

Furthermore, the present disclosure may be embodied in many different forms and should not be construed as limited to the embodiments set forth herein.

As used herein, the terms "comprising" and variations such as "having" or "having," and the terms "comprising" and variations such as "comprises" and "comprising," are understood to imply the inclusion of stated elements but not the exclusion of any other elements.

One of the main concepts of the present disclosure is to provide a microstructure including a ferrite portion to the maximum extent and cementite having an unfavorable microstructure and a hard phase such as pearlite, a martensite-austenite composite phase (MA), and the like to the minimum extent, so that HIC expansion resistance and strength are simultaneously ensured. For this reason, hot rolling (controlled rolling) and water cooling (controlled cooling) are performed to generate fine precipitates in nano size in ferrite grains so that growth is significantly inhibited after austenite formation during subsequent normalizing heat treatment, whereby the size of ferrite grains generated during cooling is finely controlled. This is to improve both hydrogen induced cracking resistance (HIC resistance) and low temperature impact toughness.

In order to manufacture a high-strength normalized steel material having excellent HIC resistance and low-temperature impact toughness, the present disclosure controls the content of C to 0.12 wt% or less (lower range compared to the conventional method), and controls the microstructure to include 90% or more of ferrite. Further, the present disclosure controls cooling conditions after hot rolling such that ferrite has 0.03 wt% or more of fine precipitates having a size of 100nm or less in ferrite grains such that growth is significantly inhibited after austenite formation, such that ferrite grains having an average grain size of 20 μm or less are generated during cooling, and a fraction of at least one of pearlite, cementite, and martensite-austenite composite phase MA is 10 area% or less. Furthermore, the large number of fine precipitates not only provides increased strength but also provides sites for absorbed hydrogen. Accordingly, the present disclosure provides a high strength steel material having excellent hydrogen induced cracking resistance (HIC resistance) and low temperature impact toughness, and a method of manufacturing the same.

Hereinafter, a high-strength steel material having excellent HIC resistance and low-temperature impact toughness according to a preferred aspect is described in detail.

A high-strength steel material having excellent HIC resistance and low-temperature impact toughness according to a preferred aspect comprises, in wt%: 0.02% to 0.12% C, 0.02% to 0.60% Si, 0.6% to 1.6% Mn, 0.020% or less (excluding 0%) P, 0.003% or less (excluding 0%) S, 0.002% to 0.060% soluble Al, 0.005% to 0.6% Cu, 0.005% to 1.0% Ni, 0.005% to 0.5% Cr, 0.01% to 0.3% Mo, 0.001% to 0.10% Ti, 0.001% to 0.06% Nb, 0.001% to 0.10% V, 0.001% to 0.006% N, 0.0002% to 0.0060% Ca, the remainder being Fe and unavoidable impurities, wherein (Ti + V +0.5Nb +0.5Mo) is 0.05% or more and (Ti + V + 0.5) is 0.05% or more and 0.05% Mo +0.5 Cr) is contained in the ferrite microstructure; and at least one of pearlite, cementite, and a martensite-austenite composite phase (MA), wherein the fraction of ferrite is 90 area% or more (excluding 100%), and wherein the ferrite contains 0.03 wt% or more of fine carbonitride precipitates having an average size of 100nm or less and the average grain size of the ferrite is 20 μm or less.

Hereinafter, the components and their ranges are described.

C: 0.02 to 0.12% by weight (hereinafter, "%")

In the normalized steel, carbon (C) is generally added to secure tensile strength by forming pearlite, cementite, or MA phase, however, in the present disclosure, C is also added as an important element for generating fine precipitates having a size of 100nm or less, and is preferably added in an amount of 0.02% to 0.12%. When the content of C is less than 0.02%, the fraction of carbonitride precipitates and pearlite decreases, thereby decreasing the tensile strength of the matrix. When the content of C exceeds 0.12%, a hard phase such as pearlite is excessively generated and exists in the form of a band during a subsequent rolling process, thereby reducing not only low-temperature impact toughness but also HIC resistance. More preferably, the C content is 0.03% to 0.10%, and even more preferably, the C content is 0.05% to 0.08%.

Si: 0.02 to 0.60 percent

Silicon (Si) is an element added for solid solution strengthening in addition to the deoxidation and desulfurization effects. Si is preferably added in an amount of 0.02% or more to ensure yield strength and tensile strength.

Meanwhile, when Si is added in an amount exceeding 0.60%, weldability and low-temperature impact toughness are reduced, and the surface of the manufactured steel sheet is easily oxidized, thereby excessively forming an oxide film to deoxidize the steel. In this regard, the amount of Si is preferably 0.02% to 0.60%, more preferably 0.05% to 0.45%, most preferably 0.05% to 0.35%.

Mn: 0.6 to 1.6 percent

Manganese (Mn) is an element having a high strength increasing effect by solid solution strengthening, and is preferably added in an amount of 0.6% or more. Meanwhile, when Mn is excessively added, severe segregation occurs at the center in the thickness direction of the steel sheet and MnS (non-metallic inclusions) are caused to be formed together with the segregated S. The MnS inclusions generated at the center are elongated by the subsequent rolling and thus greatly reduce low-temperature impact toughness and HIC resistance. In this regard, the content of Mn is preferably 1.6% or less, more preferably 0.8% to 1.45%, most preferably 1.1% to 1.3%.

P: 0.020% or less (excluding 0%)

Phosphorus (P) is an element effective for increasing strength when added, but greatly deteriorates low-temperature toughness due to grain boundary segregation, particularly in heat-treated steel. Therefore, it is desirable to control the amount of P as low as possible. Meanwhile, it is too expensive to remove P excessively during steel manufacturing, and therefore, it is preferable to limit the amount of P to 0.020% or less (excluding 0%).

S: 0.003% or less (excluding 0%)

Sulfur (S) is a representative element that bonds with Mn to generate MnS inclusions in the center of the steel sheet in the thickness direction, and thus not only deteriorates low-temperature toughness but also causes HIC and expansion. Therefore, in order to ensure HIC resistance, it is necessary to eliminate S as much as possible during the steel manufacturing process, however, this may be expensive. Therefore, the content of S is preferably limited to 0.003% or less, more preferably 0.001% or less.

Soluble Al: 0.002% to 0.060%

Aluminum (Al) is added as a strong deoxidizer in addition to Si and Mn during the steel manufacturing process, and the deoxidation effect can be maximally obtained by adding aluminum in an amount of at least 0.002% when used alone or in composite deoxidation. Meanwhile, when added in an amount exceeding 0.060%, the effect is saturated and one of the oxidizing inclusions, Al, is generated due to deoxidation₂O₃Increases the fraction of (a) more than necessary, resulting in coarse size and not easily removed during refining. Therefore, the content of soluble Al is preferably 0.002% to 0.06%, more preferably 0.005% to 0.04%, most preferably 0.01% to 0.025%.

Cu: 0.005 to 0.6 percent

Copper (Cu) is a component added due to its effect of significantly increasing strength by solid solution and precipitation and inhibiting corrosion in a wet hydrogen sulfide atmosphere. Excessive Cu addition causes cracking on the steel sheet surface. Further, in view of economic efficiency, Cu as an expensive element is preferably added in an amount of 0.005% to 0.6%, more preferably 0.02% to 0.4%, most preferably 0.05% to 0.3%.

Ni: 0.005 to 1.0%

Nickel (Ni) has little strength-increasing effect, but is effective in improving low-temperature toughness. In particular, in the case of Cu addition, Ni has the effect of: surface cracking due to selective oxidation that may occur when reheating a slab is suppressed. Ni is added to obtain such an effect. Ni is also an expensive element and thus is preferably added in an amount of 0.005% to 1.0%, more preferably 0.02% to 0.5%, most preferably 0.05% to 0.3%, in view of economic efficiency.

Cr: 0.005 to 0.5 percent

Chromium (Cr) has no significant effect of increasing yield strength and tensile strength through solid solution, but is effective in slowing down the decomposition rate of cementite during heat treatment after welding or tempering, thereby preventing strength reduction. For this reason, 0.005% or more of Cr is preferably added. Meanwhile, when the content of Cr exceeds 0.5%, not only the manufacturing cost increases, but also weldability deteriorates. Therefore, the content of Cr is preferably limited to 0.5% or less, more preferably 0.01% to 0.4%, most preferably 0.02% to 0.2%.

Mo: 0.01 to 0.3 percent

Molybdenum (Mo) is an element effective in preventing strength from being reduced during tempering or heat treatment after welding, and has an effect of preventing toughness from being deteriorated due to grain boundary segregation of impurities such as P. Mo produces nano-sized carbonitrides in the present disclosure for strength improvement. For this, Mo is added in an amount of 0.01% or more.

Meanwhile, Mo is an expensive element and excessively increases hardenability when added in an amount exceeding 0.3%, thereby causing a hard phase, thereby deteriorating toughness and HIC resistance. In this regard, the content of Mo is preferably limited to 0.01% to 0.3%, more preferably 0.02% to 0.2%, most preferably 0.05% to 0.15%.

Ti: 0.001 to 0.10 percent

Titanium (Ti) is added in an amount of 0.001% or more due to its effect of significantly increasing nano-sized carbonitride upon addition. However, when added in an amount exceeding 0.10%, Ti may generate coarse TiN-shaped hexagonal precipitates, thereby generating HIC and unfavorable results in impact experiments at low temperatures. Therefore, the content of Ti is preferably limited to 0.10% or less, more preferably 0.002% to 0.05%, most preferably 0.005% to 0.03%.

Niobium: 0.001 to 0.06 percent

Niobium (Nb) is dissolved in austenite when the slab is reheated, thereby increasing hardenability of austenite and is precipitated as carbonitride at high temperature at the time of hot rolling, aligned with the matrix, thereby inhibiting recrystallization during rolling or cooling to allow a final microstructure to be finely formed. Further, fine precipitates having a size of 100nm or less are generated during the transformation after cooling, thereby greatly increasing the strength.

Meanwhile, when Nb is excessively added, coarse precipitates may be formed at the center in the thickness direction, and the hardenability of the weld zone increases more than necessary, thereby deteriorating the low-temperature toughness and the HIC resistance. In this regard, Nb is preferably added in an amount of 0.001% to 0.06%, more preferably 0.002% to 0.05%, most preferably 0.005% to 0.04%.

V: 0.001 to 0.10 percent

Vanadium (V) has little reinforcing effect due to precipitation or solid solution at the time of rolling because almost all slabs are dissolved again during reheating, but has an effect of improving strength by precipitating as ultrafine carbonitride during tempering or at the time of post-weld heat treatment. Due to the high price of V, it is preferred to limit the content of V to 0.001% to 0.10%, more preferably 0.001% to 0.06%, most preferably 0.005% to 0.02%.

N: 0.001 to 0.006 percent

Nitrogen (N) is an element that can form precipitates together with Nb, Ti, Mo and V added and can refine grains of the steel to improve the strength and toughness of the base metal. However, an excessive amount of N may cause residual atomic states, which may cause embrittlement at high temperatures. Further, the re-solid solution fraction of carbonitride is reduced during slab reheating, thereby reducing the formation amount of fine precipitates, which is an important means of the present disclosure. In this regard, the content of N is preferably 0.001% to 0.006%, more preferably 0.002% to 0.005%, most preferably 0.002% to 0.0045%, in view of the contents of Nb, V, Mo and Ti and the reheating temperature.

Ca: 0.0002 to 0.0060%

Calcium (Ca) is combined with S existing in the form of MnS when added after Al deoxidation to inhibit the formation of MnS while forming CaS having a spherical shape, thereby exhibiting the effect of inhibiting HIC. Therefore, in order to form the added S sufficiently into the form of CaS, Ca is preferably added in an amount of 0.0002% or more. Meanwhile, excessive Ca is added, and the residual Ca combines with oxygen (O) to produce coarse oxidative inclusions, thereby causing elongation and fracture during subsequent rolling, and thus HIC. In this respect, it is preferable to limit the upper limit of the amount of Ca to 0.0060%. The amount of Ca is more preferably 0.0005% to 0.0060%, most preferably 0.001% to 0.0025%.

(Ti + V +0.5Nb +0.5 Mo): 0.05% or more, (Ti + V +0.5Nb +0.5 Mo)/C: 0.8 or more

The reason why the contents of Ti, Nb, V and Mo, which are elements to generate denitrogenated precipitates, are limited by the above formula is that fine precipitates of nanometer size are sufficiently generated during the normalizing heat treatment, so that the growth of austenite is inhibited and ferrite grains are refined. In addition, the strength is significantly increased due to fine precipitates. Therefore, their content needs to be at least a certain level.

Further, the ratio of the important constituent element C of the carbonitride needs to be high so that fine precipitates can be sufficiently generated.

In this regard, it is preferable to limit (Ti + V +0.5Nb +0.5Mo) and (Ti + V +0.5Nb +0.5Mo)/C to 0.05% or more and 0.8 or more, respectively.

The microstructure of the high strength steel material having excellent hydrogen induced cracking resistance and low temperature impact toughness according to one preferred aspect of the present disclosure includes ferrite; and at least one of pearlite, cementite, and a martensite-austenite composite phase (MA), wherein the fraction of ferrite is 90 area% or more (excluding 100%), the area fraction of at least one of pearlite, cementite, and the martensite-austenite composite phase (MA) is 10% or less, and wherein the ferrite contains 0.03 wt% or more of fine carbonitride precipitates having an average size of 100nm or less and the average grain size of the ferrite is 20 μm or less.

The steel material contains 90% or more (excluding 100%) of ferrite; and the remaining 10% or less of at least one of pearlite, cementite, and martensite-austenite composite phase MA.

Pearlite, cementite, and martensite-austenite composite phase MA are significantly effective in improving the strength, particularly tensile strength, of the steel sheet, however, they may be enlarged or elongated by rolling in excess. In this case, low-temperature impact toughness deteriorates, and in particular, they act as generation and propagation of HIC, thereby rapidly deteriorating HIC resistance. Therefore, it is preferable to limit the total fraction of at least one of pearlite, cementite, and MA to 10 area% or less.

Pearlite may exist in a finely dispersed form in which the degree of orientation (ω — 12) defined by ASTM E1268 is 0.1 or less, instead of having a band-shaped structure. Such finely dispersed pearlite exists between fine ferrite grains.

The steel material is produced by a steel material production method including normalizing heat treatment. The fractions of ferrite, pearlite, cementite and MA in the microstructure after the normalizing heat treatment of the steel material are respectively 10% or more increased, 10% or less decreased, 5% or less increased, 3% or less decreased.

Fine carbonitride precipitates having an average size of 100nm or less are present in the ferrite in an amount of 0.03% by weight or more.

When the fine carbonitride precipitates are included in an amount of less than 0.03 wt%, austenite growth cannot be inhibited during the normalizing heat treatment, resulting in the generation of coarse austenite and the generation of coarse pearlite during the subsequent cooling. Furthermore, there is a risk that the effect of increasing the strength and fixing hydrogen is significantly reduced by the precipitates.

The ferrite has an average grain size of 20 μm or less.

When the average grain size of ferrite exceeds 20 μm, the strength and HIC resistance generated by refining the microstructure may be reduced.

The tensile strength of the steel material may be 485MPa or more, the impact absorption energy at-40 ℃ may be 120J or more, and the crack length ratio (C L R) in a Hydrophobic Interaction Chromatography (HIC) test may be 5% or less.

The thickness of the steel material may be 6mm to 133 mm.

Hereinafter, a method for manufacturing a high-strength steel material having excellent hydrogen induced cracking resistance (HIC resistance) and low-temperature impact toughness according to another preferred aspect is described.

A method for manufacturing a high strength steel material having excellent hydrogen induced cracking resistance (HIC resistance) and low temperature impact toughness according to another preferred aspect includes reheating a slab to 1080 ℃ to 1300 ℃, the slab including, in weight%, 0.02% to 0.12% of carbon (C), 0.02% to 0.60% of silicon (Si), 0.6% to 1.6% of manganese (Mn), 0.020% or less (excluding 0%) of phosphorus (P), 0.003% or less (excluding 0%) of sulfur (S), 0.002% to 0.060% of soluble aluminum (sol.al), 0.005% to 0.6% of copper (Cu), 0.005% to 1.0% of nickel (Ni), 0.005% to 0.5% of chromium (Cr), 0.01% to 0.3% of molybdenum (Mo), 0.001% to 0.10% of titanium (Ti), 0.001% to 0.06% of Nb), 0.001% to 0.5% of niobium (Cr), 0.0002% to 0.006% of calcium (Nb), and the balance of calcium (Fe) being inevitable, wherein (Ti + V +0.5Nb +0.5Mo) is 0.05% or more, and (Ti + V +0.5Nb +0.5Mo)/C is 0.8 or more;

water-cooling the hot-rolled steel plate to 550-750 ℃;

and subjecting the cooled steel sheet to a normalizing heat treatment at 850 ℃ to 950 ℃ for 1.3 t [ wherein t denotes a thickness (mm) of the steel sheet) ] + (5 minutes to 60 minutes).

Reheating

The slab having the above composition is reheated to 1080 to 1300 ℃.

When the reheating temperature is lower than 1080 ℃, it is difficult to re-dissolve carbides and the like generated in the slab during continuous casting. When the reheating temperature exceeds 1300 ℃, coarse austenite grains are generated, thereby significantly deteriorating mechanical characteristics of the steel sheet, such as tensile strength, low-temperature toughness, and the like. In particular, reheating is required to a temperature sufficient to dissolve the added Ti, Nb, V, Mo, etc. In this regard, the reheating temperature of the slab is preferably 1080 ℃ to 1300 ℃, more preferably 1150 ℃ to 1280 ℃.

Obtaining a hot rolled steel product

The slab reheated as above is subjected to controlled rolling at a final hot rolling temperature of 850 ℃ or more to obtain a hot rolled steel.

If the reheated slab is subjected to general rolling, hot rolling is completed at an extremely high temperature, resulting in insignificant grain refining effect. Further, if controlled rolling is performed at an extremely low temperature, re-dissolved Nb or the like precipitates as carbonitrides. The precipitates produced at higher temperatures have large sizes and therefore significantly reduce the inhibition of austenite grain growth during the subsequent normalizing process. Further, when the strength of the steel sheet increases as the rolling temperature decreases, coarse composite inclusions generated during the refinement process must undergo deformation due to rolling, whereby the inclusions are segmented into small-sized inclusions, or amorphous inclusions are elongated. Elongated or segmented inclusions are factors directly related to the development and propagation of HIC. In view of the above, the final hot rolling temperature is preferably limited to 850 ℃ or more, more preferably 850 ℃ to 1050 ℃. The final hot rolling temperature may be Ar₃+70 ℃ to Ar₃+270℃。

The thickness of the hot rolled steel may be 6mm to 133 mm.

Controlled cooling

The hot rolled steel is water-cooled to 550 ℃ to 750 ℃ and then air-cooled.

The steel subjected to controlled rolling as above is then subjected to controlled cooling involving water cooling and air cooling. Performing water cooling after the controlled rolling can prevent the generation of ferrite at a high temperature during the water cooling in most cases and serves to generate ferrite at a relatively low temperature compared to air cooling, thereby producing a large refining effect. In particular, when cooling is performed in the temperature range of 550 ℃ to 750 ℃, fine carbonitrides are generated as well as ferrite. The precipitates produced at this time were fine and 100nm or less. When the cooling temperature exceeds 750 ℃, ferrite is generated at a higher temperature and becomes coarse, and precipitates are generated to be coarse, thereby remarkably reducing the effect. Meanwhile, when the cooling temperature is less than 550 ℃, fine precipitates are not generated, and sufficient strength cannot be secured. In the case of a high cooling rate, ferrite generation is suppressed, and a hard phase such as bainite is generated, thereby increasing strength but reducing toughness at a lower temperature. Therefore, the cooling temperature of the controlled cooling is preferably limited to 550 ℃ to 750 ℃, more preferably 580 ℃ to 670 ℃, in view of the composition, rolling conditions, and the like.

The steel after controlled cooling may contain 90% or more (excluding 100%) ferrite; 10% or less of at least one of pearlite, cementite and MA in total fraction, while containing 0.03% by weight or more of fine precipitates having an average size of 100nm or less in ferrite.

Normalizing heat treatment

The steel cooled as above is kept at 850 ℃ to 950 ℃ for 1.3 × t [ where t denotes the thickness (mm) of the steel plate) ] + (5 minutes to 60 minutes) and subjected to a normalizing heat treatment involving air cooling.

When the normalizing heat treatment temperature is less than 850 ℃, it is difficult to re-dissolve solute atoms for solid solution, and thus it is difficult to secure the strength of the steel sheet. Meanwhile, when the normalizing heat treatment is performed at a temperature exceeding 950 ℃, low-temperature toughness is deteriorated due to grain growth.

The reason why the holding time is limited during the normalizing heat treatment is that if the holding time is shorter than the reference time, homogenization of the texture becomes difficult, and if it is held for a longer time, the yield is impaired.

Therefore, it is preferable to maintain the normalizing heat treatment at 1.3 × t [ where t denotes the thickness (mm) of the steel plate) ] + (5 to 60 minutes).

The microstructure of the steel material microstructure after the normalizing heat treatment may include ferrite; and at least one of pearlite, cementite, and a martensite-austenite composite phase (MA), wherein the fraction of ferrite is 90 area% or more (excluding 100%), the area fraction of at least one of pearlite, cementite, and the martensite-austenite composite phase (MA) is 10% or less, and wherein the ferrite contains 0.03 wt% or more of fine carbonitride precipitates having an average size of 100nm or less and the average grain size of the ferrite is 20 μm or less.

Regarding the fractions of ferrite, pearlite, cementite, and MA in the steel microstructure after the normalizing heat treatment, the fraction of ferrite may be increased by 10% or less, the fraction of pearlite may be decreased by 10% or less, while the fraction of cementite may be increased by 5% or less, and the fraction of MA may be decreased by 3% or less, as compared to those before the normalizing heat treatment.

As described previously, by optimally controlling the steel composition, controlled rolling, controlled cooling, and normalizing heat treatment according to the present disclosure, it is possible to efficiently manufacture a high-strength steel material having excellent low-temperature toughness and HIC resistance by refining the structure and minimizing the hardened structure of pearlite, cementite, MA, and the like, which are disadvantageous in toughness and HIC resistance, as compared to conventional methods.

Tempering heat treatment

The normalized heat-treated steel material may be subjected to a tempering heat treatment at a temperature of 500 ℃ or more for 10 minutes or more.

A tempering heat treatment may be performed to improve toughness and strength at lower temperatures. The tempering temperature is preferably 500 ℃ or higher. When the tempering temperature is less than 500 ℃, the residual stress of the steel cannot be smoothly removed or desired fine precipitates cannot be remarkably generated. When the tempering heat treatment is performed for less than 10 minutes, the generation of precipitates is reduced, and thus it is preferably performed for 10 minutes or more.

The tempering heat treatment may be used to increase the fraction of ferrite in the microstructure of the steel material by 3 area% or less and to reduce the total fraction of pearlite, cementite and MA by 3 area%.

The amount of carbonitride precipitates in the steel material can be increased by 0.001 to 0.005% by weight due to the tempering heat treatment.

EMBODIMENTS FOR CARRYING OUT THE INVENTION

Hereinafter, the present disclosure is described in more detail with reference to the following exemplary embodiments.

A slab having the steel composition (chemical composition) of table 1 below was reheated under the manufacturing conditions shown in table 2 below, and then the reheated slab was subjected to controlled rolling to obtain a hot-rolled steel product having the thickness of table 2, and then subjected to controlled cooling and normalizing heat treatment, thereby manufacturing a steel product.

The normalizing heat treatment was performed for 1.3 × t [ where t denotes the thickness (mm) of the steel plate) ] +20 minutes.

Inventive materials (a to c) in table 1 below are consistent with the steel compositions of the present disclosure, and comparative materials (d to h) are outside the steel compositions of the present disclosure.

The steels manufactured as described above were measured and studied with respect to ferrite fraction (area%), average grain size (μm), amount of fine precipitates (weight%), average size of fine precipitates (nm), pearlite fraction (area%), and yield strength (MPa), tensile strength (MPa), impact absorption energy (J, -45 ℃) and HICC L R (%) after normalizing heat treatment, and the results thereof are shown in tables 3 and 4 below.

In table 3 below, the structure other than ferrite and pearlite in the microstructure is at least one of cementite and MA.

Meanwhile, in the invention example 1 and the comparative example 3, the microstructure and the precipitate distribution were observed, and the results thereof are shown in fig. 1 to 4. Fig. 1 is a photograph of a microstructure of comparative example 3, fig. 2 is a photograph of a microstructure of inventive example 1, fig. 3 is a photograph of a precipitate distribution of comparative example 3, and fig. 4 is a photograph of a precipitate distribution of inventive example 1.

[ Table 1]

[ Table 2]

[ Table 3]

[ Table 4]

As shown in tables 3 and 4, the steel materials (inventive examples 1 to 3) manufactured according to the steel composition and manufacturing conditions of the present disclosure had a tensile strength of 485MPa or more, a shock absorption energy of 120J or more at-40 ℃, and C L R of 5% or less in the HIC test.

Meanwhile, the composition of comparative example 1 falls within the range of the present disclosure, however, the slab reheating temperature is lower than the temperature range of the present disclosure, thus making it difficult to re-dissolve coarse carbonitride precipitates generated during slab manufacturing. Therefore, the fraction of fine precipitates generated during rolling and cooling is significantly reduced, resulting in almost no increase in the strength-improving effect after normalizing.

Comparative example 2 is a case where the composition falls within the range of the present disclosure but the cooling temperature after rolling is lower than the temperature range of the present disclosure, thereby causing a decrease in the rate of generation of precipitates and insufficient generation of carbonitrides. In this case, the strength after the normalization cannot be sufficiently ensured.

Comparative example 3 is the case: compositions fall within the scope of the present disclosure, however, controlled cooling is not applied after rolling and involves air cooling as in conventional manufacturing methods. Due to the significantly low cooling rate, ferrite is transformed at high temperature, thereby generating coarse ferrite. In addition, fine precipitates are not sufficiently generated. In this case, as shown in fig. 1, the final microstructure has not only coarse ferrite (white in the photograph) but also pearlite (black in the photograph) containing some M-a structure in the form of bands, resulting in insufficient strength and impact toughness and HIC characteristics.

Comparative example 4 is a case where the added elements generating carbonitrides are insufficient. Even when all the production conditions fall within the range of the present disclosure, the amount of fine precipitates generated is insufficient, thereby causing insufficient securing of strength.

In comparative example 5, the amount of C added exceeds the range of the present disclosure. The more C is contained, the more carbonitride precipitates, however, when the ratio of C is significantly higher than that of the alloying elements, coarse carbonitride is generated. This may result in a decrease in the suppression of austenite growth during subsequent normalizing, and a significant decrease in the strength-improving effect, thereby deteriorating both strength and low-temperature toughness.

In comparative example 6, the Mo addition amount exceeds the range of the present disclosure. Mo is an important element for generating fine precipitates while greatly improving hardenability of the steel material so that residual austenite is formed into an MA component during cooling after normalizing. As a result, the yield strength of the steel is reduced, the tensile strength is greatly improved, and the toughness at low temperatures is greatly reduced.

In comparative example 7, the amount of Ti added exceeded the range of the present disclosure. Ti is the most important element for generating fine precipitates, however, when it is added in excess, coarse precipitates are generated. This causes damage, eventually leading to significant deterioration in low-temperature toughness and HIC resistance.

In comparative example 8, the N addition amount exceeds the range of the present disclosure. When an excessive amount of N is added, the amount of re-solid solution of coarse carbonitrides generated during slab production is significantly reduced even when the reheating temperature is increased. Therefore, even when the subsequent manufacturing conditions fall within the range of the present disclosure, the fraction of the fine carbonitride produced is greatly reduced, so that a sufficient strengthening effect cannot be obtained.

Meanwhile, the steel material (comparative example 3) manufactured according to the conventional method had pearlite having a banded structure as shown in fig. 1, whereas the steel material (inventive example 1) manufactured according to the present disclosure had finely dispersed pearlite as shown in fig. 2, not pearlite having a banded structure. Further, the steel material (inventive example 1) (fig. 4) manufactured according to the present disclosure contains more dispersed fine precipitates than the steel material (comparative example 3) (fig. 3) manufactured according to the conventional method.

As described previously, in the case of manufacturing a steel material according to the steel composition and manufacturing conditions of the present disclosure, the effects of high strength, excellent low-temperature toughness and HIC resistance can be obtained by structure refinement while minimizing the hardened structure of pearlite and the like, which are disadvantageous in toughness and HIC resistance, as compared to manufacturing a steel material according to a conventional method.

Claims

1. A high-strength steel material having excellent hydrogen-induced cracking resistance and low-temperature impact toughness, comprising, in weight%:

0.02% to 0.12% of carbon (C), 0.02% to 0.60% of silicon (Si), 0.6% to 1.6% of manganese (Mn), 0.020% or less (excluding 0%) of phosphorus (P), 0.003% or less (excluding 0%) of sulfur (S), 0.002% to 0.060% of soluble aluminum (Sol. Al), 0.005% to 0.6% of copper (Cu), 0.005% to 1.0% of nickel (Ni), 0.005% to 0.5% of chromium (Cr), 0.01% to 0.3% of molybdenum (Mo), 0.001% to 0.10% of titanium (Ti), 0.001% to 0.06% of niobium (Nb), 0.001% to 0.10% of vanadium (V), 0.001% to 0.006% of nitrogen (N), 0.0002% to 0.0060.0060.5% of calcium (Ca) and the balance of 0.5 + 5% of Fe + 5% of Fe (Mo)/5% of unavoidable impurities (Ti),

the microstructure of the steel comprises ferrite; and at least one of pearlite, cementite, and a martensite-austenite composite phase (MA), wherein a fraction of the ferrite is 90 area% or more (excluding 100%), and a fraction of at least one of pearlite, cementite, and the martensite-austenite composite phase (MA) is 10 area% or less, and wherein the ferrite contains 0.03 wt% or more of fine carbonitride precipitates having an average size of 100nm or less and an average grain size of the ferrite is 20 μm or less.

2. A high-strength steel material excellent in hydrogen-induced cracking resistance and low-temperature impact toughness according to claim 1, wherein said pearlite is in a fine distribution form in which the degree of orientation (ω — 12) defined in ASTM E1268 is 0.1 or less.

3. A high strength steel material having excellent hydrogen induced cracking resistance and low temperature impact toughness according to claim 1, wherein said steel material has a tensile strength of 485MPa or more, an impact absorption energy at-40 ℃ of 120J or more, and a crack length ratio in a Hydrogen Induced Cracking (HIC) test (C L R) of 5% or less.

4. The high strength steel material having excellent hydrogen-induced cracking resistance and low temperature impact toughness according to claim 1, wherein the thickness of the steel material is 6mm to 133 mm.

5. A method for manufacturing a high-strength steel material having excellent hydrogen-induced cracking resistance and low-temperature impact toughness, comprising:

reheating a slab to 1080 ℃ to 1300 ℃, the slab comprising in weight%: 0.02% to 0.12% carbon (C), 0.02% to 0.6% silicon (Si), 0.6% to 1.6% manganese (Mn), 0.002% to 0.060% soluble aluminum (sol.al), 0.001% to 0.06% niobium (Nb), 0.001% to 0.10% vanadium (V), 0.001% to 0.10% titanium (Ti), 0.005% to 0.6% copper (Cu), 0.005% to 1.0% nickel (Ni), 0.005% to 0.5% chromium (Cr), 0.01% to 0.3% molybdenum (Mo), 0.0002% to 0.0060% calcium (Ca), 0.001% to 0.006% nitrogen (N), 0.020% or less (excluding 0%), 0.003% or less (excluding 0%) phosphorus (P), 0.003% or less (S) sulfur (S), and the balance 0.5 + 0.5% Fe +0.5 + 5 + Mo (V), wherein the balance is 0.5 + 0.5% Mo +0.5 + Fe + 5% Mo;

cooling the hot rolled steel sheet to 550 ℃ to 750 ℃; and

the cooled steel sheet is subjected to a normalizing heat treatment at 850 ℃ to 950 ℃ for 1.3 × t [ where t denotes a thickness (mm) of the steel sheet) ] + (5 minutes to 60 minutes).

6. The method for manufacturing a high-strength steel material excellent in hydrogen-induced cracking resistance and low-temperature impact toughness according to claim 5, wherein in the step of reheating the slab, the reheating temperature is 1150 ℃ to 1300 ℃.

7. The method for manufacturing a high-strength steel material excellent in hydrogen-induced cracking resistance and low-temperature impact toughness according to claim 5, wherein the final hot rolling temperature is 850 ℃ to 1050 ℃.

8. The method for producing a high-strength steel material excellent in hydrogen-induced cracking resistance and low-temperature impact toughness according to claim 5, wherein the final hot rolling temperature is Ar₃+70 ℃ to Ar₃+270℃。

9. The method for manufacturing a high-strength steel material excellent in hydrogen-induced cracking resistance and low-temperature impact toughness according to claim 5, wherein the thickness of the hot-rolled steel sheet is 6mm to 133 mm.

10. The method for manufacturing a high-strength steel material having excellent hydrogen-induced cracking resistance and low-temperature impact toughness according to claim 5, further comprising subjecting the normalized heat-treated steel sheet to a tempering heat treatment of 500 ℃ or more for 10 minutes or more.

11. The method for manufacturing a high-strength steel material excellent in hydrogen-induced cracking resistance and low-temperature impact toughness according to claim 10, wherein the content of carbonitride precipitates in the steel sheet is increased by 0.001 to 0.005% by weight through the tempering heat treatment.