KR100951249B1 - Steel palte with high sohic resistance and low temperature toughness at the h2s containing environment and manufacturing - Google Patents

Abstract

The present invention relates to hydrogen stress cracking (SOHIC) is excellent in resistance and low-temperature toughness of steel plates and a method of manufacturing the same, and more particularly H ₂ S content of oil or gas in the high source resistance is excellent in stress cracking and hydrogen below -20 ℃ The present invention relates to a steel applied to pipes of API X80 grade or less that can be used without deteriorating toughness, and a method of manufacturing the same.

In one aspect of the present invention, the plate steel having excellent low temperature toughness and hydrogen stress cracking resistance is C: 0.02 to 0.15% by weight, Si: 0.1 to 1.0% by weight, Mn: 0.5 to 1.8% by weight, Al: 0.001 to 0.1% by weight, P: 0.012 wt% or less, S: 0.003 wt% or less, Cu: 0.05-0.3 wt%, Ni: 0.05-0.6 wt%, Ti: 0.005-0.05 wt%, Ca: 0.0005-0.005 wt%, balance Fe and others having a component system consisting of impurities inevitably included, the grain size of the ferrite phase having a 20㎛ than tissue, is characterized in that the non-metallic inclusions less than the equivalent circle diameter distribution 20㎛ to 20 or less per 1cm ^2.

Hydrogen sulfide, hydrogen stress crack, low temperature toughness, crude oil, API

Description

Steel plate with excellent hydrogen stress crack resistance and low temperature toughness and its manufacturing method {STEEL PALTE WITH HIGH SOHIC RESISTANCE AND LOW TEMPERATURE TOUGHNESS AT THE H2S CONTAINING ENVIRONMENT AND MANUFACTURING}

The present invention relates to hydrogen stress cracking (SOHIC) is excellent in resistance and low-temperature toughness of steel plates and a method of manufacturing the same, and more particularly H ₂ S content of oil or gas in the high source resistance is excellent in stress cracking and hydrogen below -20 ℃ The present invention relates to a steel applied to pipes of API X80 grade or less that can be used without deteriorating toughness, and a method of manufacturing the same.

Recently, due to the depletion of energy, energy mining sources are often faced with the need to mine and transport crude oil at very low temperatures, such as in the mountains of Siberia and the Middle East.

In this case, steels such as line pipes used for mining and transporting crude oil are also exposed to low temperatures, and thus the steels have been desperately required to have low-temperature toughness that does not cause brittle fracture even at lower temperatures. .

In general, steels have low toughness as the use temperature decreases, and thus, brittle fractures are easily generated due to cracking and propagation even in a weak impact. The brittle fracture is a phenomenon that occurs very suddenly, and thus the stability of the structure or the pipe for transportation has a great influence, and thus it is necessary to improve the low temperature toughness of the steel.

In order to improve the low temperature toughness of steel, it is necessary to control the component or microstructure, but in order to control the component or microstructure which improves low temperature toughness of steel, it is necessary to understand in detail the phenomenon that brittle fracture occurs.

In particular, the toughness of steel is expressed in terms of impact toughness, which determines whether brittle fracture occurs when a certain impact energy is applied.In the case that the brittle fracture does not occur in steel even when a large impact energy is applied, the toughness of steel This can be said to be excellent. The phenomenon that brittle fracture occurs when the impact energy is applied to the steel can be roughly divided into a phenomenon in which the crack is initiated and a phenomenon in which the disclosed crack propagates along the steel.

Therefore, in order to improve the impact toughness of steel materials, it is necessary to make the initiation of a crack or the propagation of a crack difficult.

The tendency for crack initiation to occur can be measured by the Charpy impact energy value. The larger the Charpy impact energy value that the base material can absorb, the more difficult crack initiation occurs. In general, in order to increase the Charpy impact energy, it is necessary to minimize the addition of impurities such as sulfur and phosphorus, and it is necessary to appropriately control the amount of alloying elements added such as Ni.

The propagation of cracks can be used as a measure of propagation performance by measuring the area of the part determined to be a soft wave surface through the DWTT (Drop Weight Tear Test) test. In order to make the propagation of a crack difficult, it is preferable to form many grain boundaries which become an obstacle of propagation. In other words, the finer the grains, the more difficult the propagation of the cracks, which helps to improve low-temperature toughness.

In addition, apart from the low temperature toughness, the steel is required to have high resistance to stress-oriented hydrogen induced crack (SOHIC). In other words, due to the continuous crude oil mining activities, the stock of high quality crude oil containing less harmful components such as hydrogen sulfide is gradually depleted, and thus the cases of the processing of mining and processing low grade crude oil containing a large amount of hydrogen sulfide gradually increase. Although it is a trend, there is a fear that SOHIC is generated by sour components such as hydrogen sulfide contained in such lower crude oil.

Hydrogen stress cracking (SOHIC) refers to a phenomenon in which a crack occurs in the steel by the hydrogen penetrating the steel in the environment under stress, leading to the final failure. Cracks generated by sour gases such as hydrogen sulfide can be classified into hydrogen organic crack (HIC) and hydrogen stress crack (SOHIC). Hydrogen organic crack refers to a hydrogen atom generated by corrosion in a gaseous environment including a gas such as hydrogen sulfide, invades the material from the outside, and the gaseous hydrogen atoms accumulate and gasify at a certain place, and the pressure of the gas is thresholded. If it is abnormal, it means a phenomenon in which cracks are formed, grown, and destroyed. Hydrogen stress cracking, on the other hand, refers to a phenomenon in which destruction by hydrogen occurs in an environment under stress.

In the past, many studies have been made to prevent the hydrogen organic crack. However, the research for preventing the hydrogen stress crack (SOHIC) is not much in comparison with the research for preventing the hydrogen organic crack. Hydrogen organic cracks and hydrogen stress cracks have a common point in that they are cracks generated by hydrogen, but since they are generally steels having excellent hydrogen organic crack resistance, they are not necessarily excellent in hydrogen stress crack resistance. Research and application thereof are necessary. As such, there is no correlation between the two properties because of the different metallurgical factors that determine the resistance to hydrogen stress cracking. Therefore, in order to obtain excellent hydrogen stress cracking resistance, it is necessary to secure a component system and an optimum microstructure suitable for securing hydrogen stress cracking resistance.

As described above, pipe steels such as line pipes or steels for structural bodies have to have excellent low temperature toughness and hydrogen stress cracking resistance in an increasingly harsh environment, and these properties are often required at the same time. Therefore, not only the requirements for the properties of either low temperature toughness or hydrogen stress cracking resistance are met, but steels satisfying all of these properties need to be developed.

This demand has led to the need to derive appropriate alloy component systems, microstructures and other conditions commonly applied to them in order to simultaneously ensure low temperature toughness and resistance to hydrogen stress cracking.

As examples of the conventionally proposed invention for satisfying such a request, Japanese Unexamined Patent Publications No. 1998-068019, 1995-204881 and the like can be given. Among them, Japanese Patent No. 1998-068019 has one of the main features of finishing rolling at a temperature of Ar 3-30 ° C. or higher. However, when the rolling is carried out at a lower temperature than Ar3, it is not preferable because not only the low temperature transformation phase elongated upon cooling is generated but also the texture develops and adversely affects low temperature toughness. According to the research results of the present inventors, when rolling under the above conditions, the lowest temperature at which low-temperature toughness can be guaranteed is -20 ° C, and it is difficult to guarantee low-temperature toughness at a lower temperature.

In addition, the invention described in Japanese Unexamined Patent Publication No. 1995-204881 has one of the main features of the quenching treatment at a temperature of Ac3 + 20 to Ac3 + 200 ° C and a soaking treatment at a temperature of Ac1-40 to Ac1-200 ° C. . In the case of such temper treatment, it is possible to improve both low temperature toughness and hydrogen stress cracking resistance by securing the strength and relieving the stress around the two phases such as residual stress or pearlite in the steel. Since the heat treatment process of rolling should be carried out through a separate process after rolling, it is not preferable to not only add a load to the manufacturing process but also increase the manufacturing cost.

Accordingly, the present invention is to overcome the limitations of the prior art, according to one aspect of the present invention, while exhibiting good toughness even at a low temperature of -20 ℃ or less and sufficient hydrogen stress cracking resistance in sour gas environment such as hydrogen sulfide Indicating thick plate steel is provided.

According to another aspect of the present invention, there is provided a more preferable manufacturing method among the methods for producing the thick steel plate of the advantageous invention described above.

Thick plate steel having excellent low temperature toughness and hydrogen stress cracking resistance as one aspect of the present invention for solving the above problems is C: 0.02 ~ 0.15% by weight, Si: 0.1 ~ 1.0% by weight, Mn: 0.5 ~ 1.8% by weight, Al: 0.001 to 0.1 wt%, P: 0.012 wt% or less, S: 0.003 wt% or less, Cu: 0.05 to 0.3 wt%, Ni: 0.05 to 0.6 wt%, Ti: 0.005 to 0.05 wt%, Ca: 0.0005 to 0.005 wt% %, Remainder Fe and other components inevitably included impurities, has a structure with a grain size of 20 micrometers or less in the ferrite phase, and a non-metallic inclusion having a diameter of 20 micrometers or more in a circle is distributed to 20 or less per cm ² It is done.

Another aspect of the present invention is a method for producing the thick steel plate, C: 0.02 ~ 0.15% by weight, Si: 0.1 ~ 1.0% by weight, Mn: 0.5 ~ 1.8% by weight, Al: 0.001 ~ 0.1% by weight, P: 0.012 wt% or less, S: 0.003 wt% or less, Cu: 0.05-0.3 wt%, Ni: 0.05-0.6 wt%, Ti: 0.005-0.05 wt%, Ca: 0.0005-0.005 wt%, balance Fe and others Reheating the steel slabs having a component system composed of impurities inevitably included and having 20 or less non-metallic inclusions having a diameter equivalent to 20 µm or more per 1 cm ² at a temperature of 1050 to 1250 ° C; Rough rolling the reheated slab such that an end temperature is greater than or equal to austenite recrystallization temperature + 30 ° C .; Finish rolling the slab to start the rolled slab at an austenite uncrystallized temperature or less and finish rolling at a temperature of Ar3 or more, and manufacturing a steel sheet with a rolling reduction of 65% or more; And cooling the finish-rolled steel sheet at a cooling rate of 3 to 25 ° C./sec to terminate cooling at a temperature of 300 to 600 ° C .;

At this time, the thick steel plate or steel slab for producing the same, may further comprise one or two or more selected from 0.01: 0.5% by weight, 0.01% to 0.5% by weight, and 0.00% to 0.00% by weight of Mo.

In addition, the thick plate steel or steel slab may further include one or both of Nb: 0.01 to 0.08% by weight and V: 0.1% by weight in addition to the above composition.

In addition, the thick steel plate provided by the present invention preferably has a grain size of 20 µm or less in the ferrite phase in the structure, and an area fraction of the second phase having a hardness of 300 Hv or more in the structure other than ferrite is 5% or less.

According to the present invention, not only exhibits excellent impact toughness even at a low temperature of -20 ° C. or less, but also steel plates and the above-described steels excellent in low temperature toughness and hydrogen stress cracking resistance that do not cause hydrogen stress cracking even when used for a long time in a stressed environment. An inexpensive and preferred method for producing the same can be provided.

The inventors of the present invention have been able to derive the steel manufacturing theory as described below as a result of a deep study to solve the problems of the present invention, and led to the present invention based on this. Hereinafter, these will be described in detail.

In order to obtain low temperature toughness and hydrogen stress cracking resistance at the same time, not only an appropriate alloying system but also an appropriate control of inclusions and microstructures in the steel must be appropriately controlled. In particular, in order to improve low temperature toughness, it is more important to properly control the microstructure of the steel as described above, and to improve the hydrogen stress cracking resistance (SOHIC) resistance in addition to the microstructure to control the inclusions in the steel It is important. This is summarized as follows.

The internal structure of the steel plays an important role in securing not only low temperature toughness but also hydrogen stress cracking resistance. In the present invention, in order to ensure good low-temperature toughness and resistance to hydrogen stress cracking, it is preferable to make the first phase (main phase) a ferrite structure. Unlike the low temperature structure such as bainite or martensite, the ferrite structure is brittle and has excellent ductility. Therefore, it is preferable to secure 90% or more of the entire phase as ferrite to secure low temperature toughness and hydrogen stress cracking resistance. In addition, in order to secure low temperature toughness and hydrogen stress cracking resistance, it is very important to finely control the grain size of the tissue, in addition to selecting the type of tissue appropriately. In other words, it has already been described that when the grain size of the tissue is fine, cracks propagate and act as resistance to make brittle fracture difficult. In addition, when the structure is fine, the cracks propagate when fracture occurs due to hydrogen stress cracking, so the fine grains perform a desirable function in the hydrogen stress cracking resistance. According to the results of the inventors of the present invention, the average grain size of the ferrite phase, the main phase present in the final product, is preferably 20 µm or less. The fine grains as described above mean grains present in the steel at room temperature, but a major factor for determining them may include austenite grain size. That is, since the crystal grains such as ferrite existing at room temperature are transformed in austenite, the smaller the size of the austenite grains, the smaller the grain size. Therefore, it is desirable to provide a means for minimizing the size of austenite grains in order to set the advantageous manufacturing conditions of the present invention.

In addition, in order to secure low-temperature toughness and hydrogen stress crack resistance, it is necessary to define not only the grain size of the microstructure, but also the conditions of the second phase existing outside the ferrite phase. In other words, the second phase present in addition to the ferrite phase is often higher in hardness than the ferrite phase, but when a hard second phase is present around the ferrite phase, it easily cracks when hydrogen is accumulated or impact is applied at low temperatures. It can act as the starting point of. In addition, when the fraction of the second phase is large, there is a high possibility that the interval between the second phase and the second phase is narrow. In this case, since the generated cracks may be connected to each other and lead to fracture of the steel, the hard second phase The fraction needs to be controlled below a certain level. Therefore, considering that the hardness of the ferrite phase is 150 to 200 Hv, the hardness of the hard phase for obtaining hydrogen stress cracking resistance is preferably not more than 300 Hv, preferably not more than 10% by fraction, When considering low temperature toughness, it is more preferable that it is 5% or less.

Inclusions are conditions that have a great impact on securing hydrogen stress cracking resistance and therefore need to be strictly regulated. Although it is necessary to control the inclusion in order to secure the hydrogen organic crack resistance, the inclusion control concept for securing the hydrogen stress crack (SOHIC) of the present invention is more stringent than the inclusion control idea for securing the hydrogen organic crack resistance. . In general, inclusion control to achieve excellent hydrogen organic crack resistance is achieved by reducing the amount of inclusions acting as the starting point of the crack and spherical inclusions by adding Ca, but hydrogen stress cracking is applied to steel materials. As the crack occurs under external stress, there is a possibility of cracking even at smaller inclusions. Thus, the restrictions on inclusion size and distribution need to be much stricter than in the hydrogenorganic crack.

According to the results of the inventors of the present invention, in order not to generate a hydrogen stress crack, it is necessary to have a certain number of nonmetallic inclusions present in a critical size or more. In order to satisfy these conditions, the number of inclusions with an equivalent diameter of 20 µm or more in the area of 1 cm ² of the final steel plate, which is the final product, should not exceed 30 pieces. Inclusions of 1 μm or more in diameter correspond to similar levels in slabs.

However, for the steel of the present invention, unlike hydrogen organic cracking, there is no need to specifically limit the microstructure and hardness of the steel core. In other words, the propagation of hydrogen stress crack is caused by vertical cracks connected to each other by stresses, which are generated vertically in the direction of stress. When the gaps of the horizontal cracks are wide, it is difficult to connect the cracks. This is, in turn, expressed as good hydrogen stress cracking resistance. In other words, in order to improve the hydrogen stress cracking resistance, the size of the horizontal crack perpendicular to the stress direction is not important, but the interval between the horizontal cracks is more important. Microstructure non-uniformity due to the central segregation present in the center of the steel can generate a very large hydrogen organic crack, but does not cause the fracture of the steel due to the hydrogen stress crack. Therefore, there is no need to make a special restriction on the microstructure and hardness for hydrogen stress cracking resistance. In addition, in order to secure the hydrogen organic crack resistance, it is preferable to suppress the generation of the band structure of the second phase such as pearlite, but in the hydrogen stress crack, the band structure affects the generation of the horizontal crack, so the interval between the bands is wider. If the crack spacing is wide, the hydrogen stress cracking resistance is not significantly affected. Therefore, in the case where the ratio of the ferrite phase is high, as defined in the present invention, even if pearlite is generated as the second phase, the spacing between the pearlite bands is widened so as not to cause any particular problem in the hydrogen stress crack.

Hereinafter, the reason for setting the preferred component system for constituting the steel material of the present invention described above will be described.

C: 0.02-0.15 wt%

The C is effective to add 0.02% by weight or more as an element added to improve the strength. Increasing the content can improve the hardenability and improve the strength. However, if the content exceeds 0.15% by weight, the toughness of the steel may be impaired and the second phase having a high hardness may be easily formed. desirable.

Si: 0.1-1.0 wt%

The Si should be added more than 0.1% by weight because it acts as a deoxidizer, but if the content exceeds 1.0% by weight, there is a problem of inhibiting the toughness and weldability and increasing the amount of oxidation inclusions in the steel to reduce the SOHIC resistance, It is limited to 1.0% by weight or less.

Mn: 0.5-1.8 wt%

The Mn is an effective ingredient for increasing the strength without deteriorating the toughness, and as the content increases, the hardenability increases due to the increase in the hardenability. Therefore, at least 0.5 wt% or more should be added. However, Mn is an element that facilitates segregation during solidification and promotes band structure. When Mn is added in an amount of 1.8 wt% or more, the width of the band structure is narrowed, thereby inhibiting SOHIC resistance. Therefore, it is preferable to limit the content to 1.8 wt% or less.

Al: 0.001-0.1 wt%

Al is an element that is essentially added for deoxidation during steelmaking, and improves shock absorption energy but, like Si, reacts with oxygen to form oxide inclusions. If the content of Al is less than 0.001% by weight, deoxidation is not sufficiently achieved. If the content of Al exceeds 0.1% by weight, not only the impact toughness is impaired, but also a large amount of inclusions is formed to inhibit SOHIC resistance, so that the content is 0.001. It is preferable to limit it to -0.1 weight%.

P: 0.012% by weight or less

P is an element contained in steel inevitably during steelmaking, and not only inhibits weldability and toughness, but also inhibits toughness and SOHIC resistance as an element easily segregates at the center of the slab and austenite grain boundaries during solidification, so that its content is less than 0.012% by weight. It is desirable to limit.

S: 0.003 wt% or less

S is generally stretched during rolling by reacting with Mn to form MnS, which acts as a starting point for SOHIC generation and inhibits low temperature toughness. Therefore, it is desirable to reduce S as much as possible, but the range is set to 0.003% by weight or less due to process constraints for S removal.

Cu: 0.05-0.3 wt%

Cu is added to improve the strength, toughness and corrosion resistance of the steel. Cu must be added at least 0.05% by weight since it is dissolved in steel to improve strength and to form a protective film on the surface in an atmosphere containing hydrogen sulfide, which lowers the corrosion rate of the steel and reduces the amount of hydrogen diffused into the steel. . However, Cu is an element that inhibits surface quality by causing cracks on the surface during hot rolling, so it is preferable to limit the upper limit to 0.5% by weight or less.

Ni: 0.05-0.6 wt%

Ni is an element that improves the toughness of steel and is added to reduce surface cracks generated during hot rolling of Cu-added steel. In order to reduce surface cracks caused by the addition of Cu, it may be added up to 1.5 times or more of the Cu-added steel. Therefore, the lower limit of Ni is 0.05% by weight or more, similarly to the lower limit of Cu, but the upper limit thereof is 0.6% by weight. In addition, addition of Ni by 0.6% by weight or more hinders the improvement of hydrogen embrittlement characteristics by addition of Cu, and excessively increases the cost.

Ti: 0.005-0.05 wt%

Ti is an element that forms carbides or nitrides, and inhibits the grain growth of the austenite phase upon reheating, thereby improving the low temperature toughness because it finally forms a fine homogeneous ferrite. In order to suppress such austenite grain growth, the lower limit of Ti is preferably limited to 0.005%. Finely dispersed Ti (C, N) precipitates reduce the diffusion coefficient of hydrogen and increase the resistance to hydrogen organic cracks. However, if the addition amount is increased, Ti (C, N) precipitate becomes coarse, so it becomes hydrogen concentration site and rather lowers the low temperature toughness and SOHIC resistance, so it is desirable to limit the upper limit to 0.06% by weight or less.

Ca: 0.0005 to 0.005 wt%

The Ca serves to spheroidize the MnS inclusions. MnS is drawn during rolling with inclusions with a low melting point to serve as a starting point for SOHIC cracks. The added Ca reacts with MnS and surrounds the MnS, thus preventing MnS from stretching. The MnS spheroidizing effect of Ca is closely related to the amount of S. Therefore, the appropriate amount of Ca is determined by the content of S, but in general, considering the steelmaking process, it is preferable to limit the amount to 0.0005 to 0.005% by weight.

In addition to the above elements, one or more elements selected from Cr, Mo, and B or one or both elements selected from Nb and V may be further included as necessary to further improve the properties of the steel.

In order to improve the hardenability of the steel and further improve the strength of the steel, it is preferable to further include at least one element selected from Cr, Mo, and B described below.

Cr: 0.01-0.5 wt%

Cr is an element that increases strength by increasing the hardenability of steel. It should be added in an amount of 0.01% by weight or more for the effect of strength increase. Although the strength increases with the increase in the amount added, the toughness of the steel is impaired when the amount is added in an amount of 0.5% by weight or more, so the upper limit is preferably limited to 0.5% by weight or less.

Mo: 0.01-0.5 wt%

Mo, like Cr, is an element that plays a role of increasing strength by increasing the hardenability of steel, and its effect is much higher than that of Cr. In order to obtain the effect of increasing the strength by adding Mo, 0.01% by weight or more should be added. When 0.5% by weight or more is added, the cost is not excessively high, and the second phase having a very high hardness is easily generated, which inhibits SOHIC resistance. It is preferable to limit an upper limit to 0.5 weight% or less.

B: 0.003 wt% or less

B is a component effective for improving the strength of the steel by remarkably increasing the hardenability of the steel even with a small amount of addition. When the content of B exceeds 0.003% by weight, Fe ₃ B is formed to cause redembrittlement and also inhibit weldability. Therefore, the content is preferably limited to 0.003% by weight or less.

In addition, in order to form a precipitate to refine the crystal grains or increase the hydrogen stress cracking resistance, it is preferable to further include one or both elements selected from Nb and V as described below.

Nb: 0.01% to 0.08% by weight

Nb precipitates in the form of Nb (C, N) at a temperature around 1200 ° C. to increase its strength. In addition, it suppresses the recrystallization of austenite generated during the secondary hot rolling serves to refine the ferrite particles. In order to improve the strength by adding Nb, it should be added at least 0.01% by weight. However, it is preferable to limit the upper limit to 0.08% by weight or less since secondary phases containing Nb may serve as a place for promoting crack connection during SOHIC crack propagation.

V: 0.1 wt% or less

V is usually formed in the form of VC in the ferrite region, although VN may be formed when N is sufficiently present in the steel. Lower vacancy carbon concentration when transforming to austenite-ferrite, and VC provides a nucleation site for cementite formation. Therefore, rather than continuously forming Fe ₃ C at the grain boundary has a form of a discontinuous structure to increase the resistance to SOHIC. However, when added at more than 0.10% by weight, coarse V precipitates are formed, which not only impair toughness but also become hydrogen accumulation sites in the steel, thereby lowering resistance to hydrogen organic cracks. Therefore, it is desirable to limit it to 0.10 wt% or less.

In addition to the above components, the remainder comprises Fe and other unavoidable impurities.

That is, the steel of the present invention having high low temperature toughness and hydrogen stress cracking resistance is C: 0.02 to 0.15% by weight, Si: 0.1 to 1.0% by weight, Mn: 0.5 to 1.8% by weight, Al: 0.001 to 0.1% by weight, and P: 0.012% by weight or less, S: 0.003% by weight or less, Cu: 0.05-0.3% by weight, Ni: 0.05-0.6% by weight, Ti: 0.005-0.05% by weight, Ca: 0.0005-0.005% by weight, balance Fe and other inevitably having a component system consisting of contained impurities, the average crystal grain size of the ferrite phase having a 20㎛ than tissue, is characterized in that the non-metallic inclusions less than the equivalent circle diameter distribution 20㎛ to 20 or less per 1cm ^2.

And if necessary, it may further include one or two or more selected from Cr: 0.01 to 0.5% by weight, Mo: 0.01 to 0.5% by weight and B: 0.003% by weight or less, Nb: 0.01 to 0.08% by weight and V : One or both of 0.1 wt% or less may further be included.

In addition, the advantageous steel material of the present invention more preferably has an area fraction of the second phase having a hardness of 300 Hv or more in the structure other than ferrite of 5% or less.

Advantageous steels of the present invention as described above can be manufactured by those skilled in the art to which the present invention belongs, without excessive trial and error, but the present invention can effectively produce the steel in a simpler method I would like to present a method. The following manufacturing method means a valuable and preferable method for producing the steel of the present invention, and it is technically significant as such. However, the present invention is not intended to limit the manufacturing method of the steel according to the present invention. It is worth noting that.

A brief description of a more preferred method for manufacturing the steel of the present invention can be produced by the process of reheating the slab, and then rough rolling and finish rolling and cooling the rolled steel sheet. Hereinafter, each process will be described in detail.

Reheating Temperature: 1050 ~ 1250 ℃

Mn and P segregation on the slab is alleviated by diffusion during reheating, but if the reheating temperature is low, diffusion does not occur sufficiently, resulting in segregation of Mn and P, which impairs low temperature toughness and resistance to hydrogen stress cracking (SOHIC). . In addition, in the case of Nb-added steel, Nb added in the steel is sufficiently dissolved during reheating to increase the strength by fine precipitation during rolling. It is therefore desirable to limit the lower limit of the reheating temperature to 1050 ° C. in order to alleviate segregation in the slab and to solidify Nb. When the heating temperature is high, the segregation and relaxation of Nb are easy, but at the same time, the grain size of the austenite increases, so the low-temperature toughness deteriorates. Therefore, in order to obtain excellent low temperature toughness, it is desirable to limit the upper limit to 1250 ° C.

Rough rolling finish temperature: Austenitic uncrystallized temperature + 30 ℃ or more

The temperature range in which austenitization of steel proceeds can be divided into two categories. One of them is a section where recrystallization takes place in the austenite zone, and the other is a temperature section that is lower than the section where the recrystallization takes place, where the recrystallization does not occur (the austenite unrecrystallized section). If rolling is carried out in the austenite non-recrystallized section, fine grains can be obtained since the austenite grains refined by rolling are not recrystallized. However, in the temperature range close to the austenite non-recrystallization section as the temperature slightly lower among the austenite recrystallization section, recrystallization does not occur completely but only partially. Rough rolling is not a section for obtaining a fine recrystallization, but rolling is performed to reduce the slab to a thickness that is easy to finish rolling while reducing the deformation resistance as small as possible, so it is preferable to perform rolling in the austenite recrystallization section. By the way, when the rough rolling is carried out in the austenite recrystallization section but the temperature is rather low and only partially recrystallized, rolling is carried out in a part of the crystal grains, and some of the grains are coarsened by recrystallization. The process increases the likelihood that after rough rolling, a hybrid structure is formed in which the grain size is distributed over a wide range. This mixed form of grains results in non-uniformity of the final ferrite grain size and inhibits low temperature toughness. Therefore, in order to prevent the phenomenon of partially recrystallization, the rough rolling is preferably finished at a temperature higher than 30 ° C. higher than the unrecrystallized temperature, which is a temperature separating the recrystallized section and the unrecrystallized section.

Finish rolling start temperature: Below unrecrystallized temperature

As described above, rolling near the unrecrystallized temperature at a temperature above the unrecrystallized temperature causes partial recrystallization and thus adversely affects low temperature toughness. Therefore, after rough rolling, it is preferable to start the finishing rolling after the slab temperature falls below the unrecrystallized zone temperature.

Finish rolling reduction rate: 65% or more

The most important factors for determining low temperature toughness are grain uniformity and grain size. At unrecrystallized zone temperature, rolling does not cause recrystallization, which causes deformation of austenite, and the deformation zone generated in austenite acts as a ferrite nucleation site during phase transformation. Therefore, as the strain zone is densely formed, the final ferrite grain size decreases, so the low temperature toughness is excellent. The deformation zone in austenite becomes denser as the reduction ratio increases below the unrecrystallized temperature. Therefore, in order to obtain excellent low temperature toughness, it is preferable to limit the reduction ratio during finish rolling to 65% or more. The upper limit of the reduction ratio during the finish rolling is determined by the thickness of the final target steel sheet and the capability of the equipment, and it is not necessary to necessarily determine the value.

Finish rolling finish temperature: Above Ar3 temperature

In order to obtain good low temperature toughness and resistance to hydrogen stress cracking, the fraction of the second phase, such as pearlite, in the tissue should not be high. In order to minimize this two-phase fraction, accelerated cooling should be started after finishing rolling above or under Ar3. Therefore, the lower limit of finish rolling temperature is limited to Ar3 or more.

Cooling Speed: 3 ~ 25 ℃ / sec

The microstructure and hardness can be controlled by adjusting the cooling rate. The central microstructure can vary with alloying and cooling rates. However, if the cooling rate is slower than 3 ° C / sec, the C phase is accumulated in the center during transformation, so that a large number of low temperature transformation tissues are formed in the center, so that the second phases with hardness greater than 300 Hv can be concentrated in the center. When the temperature is 25 ° C / sec or more, the entire structure is generated as a very strong martensite structure, so low temperature toughness and hydrogen stress cracking resistance is very bad, it is preferable to limit the cooling rate to 3 ~ 25 ° C / sec.

Cooling end temperature: 300 ~ 600 ℃

In order to obtain excellent low temperature toughness and resistance to hydrogen stress cracking, the cooling end temperature must be higher than the ferrite transformation end temperature. In general, the production of a second phase, except for ferrite, can be tolerated to the extent that the vicinity of 500 ° C. does not impair the transformation end temperature or the hydrogen stress cracking resistance. Therefore, in order to obtain an appropriate low temperature transformation phase fraction, it is desirable to limit the cooling end temperature to 300 ~ 600 ℃.

As described above, the present invention can provide a thick plate steel having excellent low temperature toughness and hydrogen stress crack resistance at the same time by optimizing the steel components and limiting the manufacturing process.

Hereinafter, the present invention will be described in more detail with reference to Examples. However, it should be noted that the following examples are only intended to illustrate the present invention and are not intended to limit the scope of the present invention. This is because the scope of the present invention is determined by the matters described in the claims and the matters reasonably inferred therefrom.

(Example)

The steel slab having a composition as shown in Table 1 was rolled under the conditions described in Table 2 to prepare a thick steel plate. Table 1 also shows the size of inclusions with a diameter of 20 µm or more included in the steel slab.

Low temperature toughness through -40 ° C DWTT test and hydrogen stress cracking resistance through stress corrosion test of the manufactured steel sheet were shown in Table 2. If the SOHIC break time is 0 in Table 2, it means that no break has occurred even after 720 hours have elapsed. The average grain size of each of the steels and the fraction of the second phase showing a hardness of 300Hv or more are shown in Table 3 below. A microhardness and an optical microscope were used to measure the hardness and total fraction (second phase area / observation area) of the second phase present in the microstructure of each steel sheet.

Emphasis Heating temperature Unrecrystallized station temperature Rough rolling end temperature Finish rolling start temperature Finish Rolling Rate Ar3 Finish Rolling End Temperature Cooling end temperature Cooling rate DWTT Compound Wavelength (-40 ℃) SOHIC Break Time Inventive Example A 1096 926 980 900 75 752 800 485 9 99 O B 1154 815 1050 810 65 758 760 505 7 95 O C 1120 872 1030 870 70 754 790 430 12 99 O D 1198 865 1090 860 70 775 780 380 15 90 O E 1150 838 1060 830 65 770 770 550 7 90 O F 1139 973 1020 950 75 778 860 500 9 99 O G 1201 949 1100 830 65 783 780 470 15 85 O H 1120 880 990 880 70 791 800 350 22 95 O I 1085 917 960 910 75 799 810 580 6 99 0 J 1134 870 1040 870 70 773 775 470 11 95 O K 1220 1016 1080 970 80 788 880 430 14 89 O Comparative example L 1150 913 1020 950 75 789 820 500 12 70 580 M 1145 868 1000 960 80 738 830 480 13 75 648 N 1150 995 1000 950 70 804 830 510 12 70 500 O 1208 916 1020 940 60 789 780 650 2 65 360 P 1198 830 1000 940 60 752 760 510 4 65 450 Q 1100 971 980 930 70 763 790 480 9 65 535 R 1095 866 960 930 75 732 780 475 12 60 610 S 1126 915 990 950 60 799 740 520 6 70 285 T 1210 1062 1050 950 60 690 750 350 17 65 340 A 1270 926 1130 1050 80 752 900 580 14 60 170 B 1215 815 1070 970 65 758 820 480 11 60 247 D 1220 865 1090 1000 80 775 890 400 17 55 75 E 1060 838 950 880 60 770 820 510 9 70 115 I 1150 917 1030 900 75 799 810 650 5 75 276 K 1030 1016 920 870 80 788 760 310 17 70 95

Ferrite Hardness (Hv)  Average grain size (mm) Second phase area fraction (%) with hardness above 300 (Hv) Invention A 158 12.4 0.8 B 162 11.9 1.0 C 161 14.2 1.1 D 168 17.5 0.6 E 165 13.4 1.8 F 171 10.2 1.2 G 159 9.8 2.1 H 184 16.4 2.5 I 175 14.2 1.9 J 168 11.7 2.4 K 166 18.5 2.7 Comparative example L 194 19.4 7.9 M 189 21.4 6.8 N 196 15.7 8.1 O 191 16.9 10.3 P 179 25.4 7.1 Q 187 24.6 6.5 R 193 17.4 5.4 S 195 16.3 6.4 T 172 21.4 7.9 A 165 23.7 4.9 B 159 18.9 8.6 D 171 24.6 5.1 E 175 22.1 6.8 I 179 17.5 9.4 K 162 20.3 8.7

Low-temperature toughness evaluation of the inventive steel and the comparative steel was evaluated by measuring the ductile fracture rate of the wavefront after the DWTT test at -40 ℃. If the ductility is 80% or more at -40 ℃, it can be judged that low temperature toughness is very good. As can be seen in Table 2, in the case of the invention steel manufactured according to the conditions of the present invention, it was found that the ductile fracture rate was good at low temperature toughness of 80% or more. However, in the case of the comparative steel, it was hardly observed that the ductile fracture rate exceeding 80% exceeded 80%, and thus, the low-temperature toughness was poorer than that produced in accordance with the conditions of the present invention.

Hydrogen stress cracking resistance (SOHIC) resistance was evaluated as the time at which fracture occurred after applying a constant stress to the steel in a 5% NaCl solution saturated with hydrogen sulfide. In other words, in a solution saturated with H ₂ S, a load (yield stress × 0.9 × tensile specimen cross-sectional area) corresponding to 90% of the yield stress was applied to each steel material using a lever and a weight. SOHIC resistance was evaluated by observing the fracture of steel over time. In general, steel that does not break for 720 hours is evaluated to have excellent SOHIC resistance. Failure does not occur is a steel that does not occur even after 720 hours elapsed hydrogen is very excellent in stress cracking resistance. As can be seen in Table 2, in the case of the invention steel manufactured under the conditions of the present invention, even after 720 hours of warning, the fracture by the hydrogen stress crack did not occur, indicating that the hydrogen stress cracking resistance was excellent. However, in the case of the comparative steel, it was found that the hydrogen stress cracking resistance was poor compared to the inventive steel according to the present invention because all fractures were reached before 720 hours, which is a measure of excellent hydrogen stress cracking resistance.

As described in Table 1, the difference between the inventive steel and the comparative steel is that the number of inclusions having a circular equivalent diameter of 20 µm or more of the comparative steel does not satisfy the conditions of 30 or less specified in the present invention, or as described in Table 3. It is believed that the grain size or hard phase fraction is due to the result of not meeting the conditions defined in the present invention.

Therefore, the advantageous effect of the characteristic on the structure of this invention was confirmed.

Claims (10)

C: 0.02 to 0.15% by weight, Si: 0.1 to 1.0% by weight, Mn: 0.5 to 1.8% by weight, Al: 0.001 to 0.1% by weight, P: 0.012% by weight or less, S: 0.003% by weight or less, Cu: 0.05 to 0.05 0.3% by weight, Ni: 0.05-0.6% by weight, Ti: 0.005-0.05% by weight, Ca: 0.0005-0.005% by weight, balance of Fe and other unavoidable impurities, and have an average grain size of 20 It has a ㎛ than tissue, the non-metallic inclusions less than the equivalent circle diameter distribution 20㎛ and below 20 per 1cm ^2, when the hydrogen sulfide inflicts hang a load corresponding to 90% of the yield stress of the steel material in the saturated solution in the steel Steel plate with excellent hydrogen stress crack resistance and low temperature toughness, characterized in that fracture does not occur for 720 hours. The hydrogen stress crack resistance and low temperature according to claim 1, further comprising at least one selected from Cr: 0.01 to 0.5% by weight, Mo: 0.01 to 0.5% by weight, and B: 0.003% by weight or less. Tough plate steel with excellent toughness. The thick steel sheet having excellent hydrogen stress crack resistance and low temperature toughness according to claim 1 or 2, further comprising one or both of Nb: 0.01 to 0.08% by weight and V: 0.1% by weight or less. . The thick sheet steel according to claim 1 or 2, wherein an area fraction of the second phase having a hardness other than ferrite in the structure of 300 Hv or more is 5% or less. 4. The thick steel plate having excellent hydrogen stress crack resistance and low temperature toughness according to claim 3, wherein an area fraction of the second phase having a hardness of 300 Hv or more in the structure is 5% or less. C: 0.02 to 0.15% by weight, Si: 0.1 to 1.0% by weight, Mn: 0.5 to 1.8% by weight, Al: 0.001 to 0.1% by weight, P: 0.012% by weight or less, S: 0.003% by weight or less, Cu: 0.05 to 0.05 0.3% by weight, Ni: 0.05-0.6% by weight, Ti: 0.005-0.05% by weight, Ca: 0.0005-0.005% by weight, balance of Fe and other unavoidable impurities, and has an equivalent diameter of 20 µm or more Reheating the steel slabs having up to 20 non-metallic inclusions per cm ² to a temperature of 1050-1250 ° C .; Rough rolling the reheated slab such that an end temperature is greater than or equal to austenite recrystallization temperature + 30 ° C .; Finish rolling the slab to start the rolled slab at an austenite uncrystallized temperature or less and finish rolling at a temperature of Ar3 or more, and manufacturing a steel sheet with a rolling reduction of 65% or more; And Cooling the finished rolled steel sheet at a cooling rate of 3 to 25 ° C./sec to finish cooling at a temperature of 300 to 600 ° C .; The steel plate produced by the manufacturing method is hydrogen, characterized in that the fracture does not occur for 720 hours when a load corresponding to 90% of the yield stress of the steel in the hydrogen sulfide saturated solution to the steel Manufacturing method of thick steel plate with excellent stress cracking resistance and low temperature toughness. 7. The hydrogen stress crack resistance and low temperature according to claim 6, further comprising at least one selected from Cr: 0.01 to 0.5% by weight, Mo: 0.01 to 0.5% by weight, and B: 0.003% by weight or less. Method for producing thick steel plate with excellent toughness. The thick plate steel according to claim 6 or 7, further comprising one or both of Nb: 0.01 to 0.08% by weight and V: 0.1% by weight or less. Manufacturing method. According to claim 6 or 7, characterized in that the manufactured steel plate has a grain size of 20 μm or less in the ferrite phase in the structure, and an area fraction of the second phase having a hardness of 300 Hv or more in the structure other than ferrite is 5% or less. Manufacturing method of thick steel plate with excellent hydrogen stress crack resistance and low temperature toughness. [9] The hydrogen stress cracking resistance of claim 8, wherein the manufactured steel plate has a grain size of 20 µm or less in the ferrite phase of the structure, and an area fraction of the second phase having a hardness of 300 Hv or more in the structure of the ferrite phase is 5% or less. Method for producing thick steel plate with excellent low temperature toughness.