KR20100034983A - Bake hardened cole-rolled steel sheet, galvanized steel sheet with excellent surface properties and aging resistance and manufacturing method thereof - Google Patents

Abstract

PURPOSE: A bake hardened cold-rolled steel sheet, a galvanized steel sheet with excellent surface properties and aging resistance and a manufacturing method thereof are provided to improve the amount of bake hardening and the tensile strength by additionally performing a melting zinc coating processing step. CONSTITUTION: A bake hardened cold-rolled steel sheet comprises C 0.0025~0.0035 weight%, Si 0.02 weight%, Mn 0.2~1.2 weight%, P 0.01~0.05 weight%, S 0.01 weight%, al 0.08~0.12 weight%, N 0.0025 weight%, Ti 0.005~0.018 weight%, Mo 0.01~0.1 weight%, B 0.0005~0.0015 weight%, the rest which is Fe. The baking hardness lecture grain size is ASTM No.9 or greater. A plating layer is formed on the surface of the baked hardening cold rolled steel sheet.

Description

Bake Hardened Cole-Rolled Steel Sheet, Galvanized Steel Sheet with Excellent Surface Properties and Aging Resistance and Manufacturing Method Thereof}

The present invention relates to a hardened hardened cold rolled steel sheet having both surface properties and excellent aging resistance, a plated steel sheet using the same (hereinafter referred to collectively as a hardened hardened steel), and a method of manufacturing the same. 340 ~ 390MPa grade hardened hardened steel with excellent hardening resistance at room temperature with a hardening value of 30MPa or more and an aging index (AI, Aging Index) value of 30MPa or less, and excellent secondary work brittleness. will be. In particular, the plated steel sheet according to the present invention is characterized by strictly controlling the components and hot rolling conditions in order to suppress surface defects that may occur during manufacture.

In recent years, the use of high-strength steel sheet in the vehicle body for the purpose of improving the fuel efficiency of the automobile and reducing the weight of the vehicle body has further increased the demand for reducing the plate thickness and improving the dent resistance. In general, when the high strength characteristics are given to the steel sheet, since workability is lowered, demand for steel that can satisfy both of these characteristics is increasing.

Steels that can satisfy both of these conditions include a composite structured cold rolled steel sheet and a hardened hardened steel. Relatively easy to manufacture composite tissue steel has a tensile strength of more than 390MPa class as a material used in automobiles has the advantage of high elongation compared to high tensile strength. However, there is a problem that the average r value indicating the press formability of the automobile is low and expensive alloying elements such as Mn and Cr are often added excessively, leading to an increase in manufacturing cost.

On the contrary, because the hardened hardened steel has a yield strength close to the soft steel sheet with a tensile strength of 390 MPa or less during press molding, it has excellent ductility, and after press molding, the yield strength increases by itself during the coating hardening process. Increasing is attracting attention as a very ideal steel compared to the conventional cold rolled steel sheet is deteriorated formability.

Baking hardening is a type of strain aging that occurs when carbon or nitrogen, the solid solution dissolved in steel, adheres to the potential generated during the deformation process. As the solid solution C (carbon) and N (nitrogen) increase, the amount of hardening hardening increases. However, due to the excessive number of employment elements, the process of controlling the appropriate employment element is very important because it leads to deterioration of the moldability with room temperature aging.

In general, known as a method for producing a cold-rolled steel sheet having a hardening hardenability is that the low-carbon P-added Al-killed steel is simply wound at low temperatures, that is, by cold winding at a hot rolled winding temperature of 400-500 ℃ temperature range by the annealing method In this case, steel of about 40-50 MPa of hardening of hardening was mainly used. It is known that such a manufacturing method can improve both moldability and baking hardenability by annealing.

However, P-added Al-Killed steel, which requires continuous annealing, uses a relatively fast cooling rate, which makes it easier to secure baking hardenability compared to the prior art, whereas moldability is deteriorated by rapid heating and short annealing. There is a problem that is limited only to the automobile shell, in which workability is not required.

In addition, a number of surface-treated steel sheets are recently used in the production of automotive parts. In the case of hot-dip galvanized steel sheets coated with hardened steel, the scratches on the surface of the steel sheet after the plating process and the customer's sheet metal are insufficient if the integrity of the steel sheet surface is not sufficiently secured. Surface defects observed in the form of shiny defects after machining are very likely to occur.

These defects, in particular, during the hot rolling process of the steel sheet with excessive addition of Al, P, etc., the complex oxides of the Al and P main streams formed within several micrometers of the pole surface layer of the steel sheet form a selective oxide along the grain boundary or the amorphous grain interface. This is usually the case.

Therefore, various technologies have been developed to take advantage of the hardened hardened steel while solving the problems of the hardened hardened steel. In recent years, due to the rapid development of steelmaking technology, it is possible to control the appropriate amount of solid solution in steel, and for automotive exterior materials that require dent resistance by using Al-Killed steel sheet containing strong carbonitride-forming elements such as Ti or Nb. It has been shown a technique for producing a small hardened steel excellent in moldability can also be used as.

Japanese Laid-Open Patent Publication No. 61-026757 discloses a technique related to ultra low carbon cold rolled steel with a Ti and Ti-Nb composite having a C: 0.0005-0.015% and an S + N content ≤005%, and Japanese Laid-Open Patent No. 57-089437 No. introduces a method for producing steel having a hardening hardening amount of about 40 MPa or more using Ti-added steel of less than C: 0.010%. This method is to control the addition amount of Ti and Nb or the cooling rate at the time of annealing to appropriately adjust the amount of solid solution in the steel to prevent material deterioration and to give the hardening hardening, but in the case of Ti or Ti-Nb composite additive steel In order to secure an appropriate baking hardening amount, a strict control such as Ti, N, S, etc. is required during the steelmaking process, resulting in a cost increase. Furthermore, the Nb-added steel of the prior arts may exhibit problems of deterioration of workability due to high temperature annealing and an increase in manufacturing cost due to the addition of special elements.

Meanwhile, there have been prior arts showing the improvement of the properties of the hardened hardened steel by the addition of new alloy elements. In Japanese Patent Laid-Open No. 5-93502, BH is improved by adding Sn, and in Japanese Patent Laid-Open No. Hei 9-249936 By adding V in combination with Nb, the ductility improvement effect through stress concentration relaxation of grain boundary is introduced. In addition, Japanese Patent Application Laid-open No. Hei 8-49038 has an effect of improving moldability by Zr, and Japanese Patent Laid-open No. Hei 7-278654 promotes formability by adding Cr to minimize the deterioration of high strength and work hardening index (N value). Doing.

However, these patents are merely a technique for improving the quench hardening property or improving the moldability, and there are P as inevitably added elements due to the deterioration of aging resistance due to the increase of quench hardening property and the high strength of the quench hardening steel. It does not mention problems such as secondary processing brittleness caused by an increase in P content. That is, for example, when the P content added to manufacture the hardened hardened steel having a tensile strength of 340 MPa is 0.07%, the DBTT (Ductile Brittle Transition Temperature), which is a criterion for judging secondary work brittleness, is set at a drawing ratio of 1.9. If the P content is added about 0.09% to produce high strength steel of -20 ℃, 390MPa grade, DBTT can be very deteriorated as 0 ~ 10 ℃. This corresponds to the steel containing B added to about 5ppm, because the P content is excessively high, it seems that there was a problem in improving DBTT by B.

Excessive addition of B in order to improve the secondary work brittle resistance is because the material degradation caused by B has a limit in addition. That is, under these conditions, DBTT, which prevents secondary processing brittleness, should be at least -20 ℃ or lower, and DBTT should be -30 ℃ or lower to secure more stable secondary airborne brittleness. There is a need for a review of new ingredients or manufacturing conditions.

The present invention is to solve the above problems and to provide a steel material that can secure both high strength and aging resistance at the same time, in particular to suppress the surface defects caused by the selective oxidation of the surface expected during the production of hot-dip steel sheet Complementary control of the content of Al and Al, and by appropriately control the grain size after annealing by optimizing the ingredients and manufacturing conditions, it is excellent in hardening hardenability, room temperature aging resistance, secondary processing brittleness, and the aging index is maintained even at room temperature for 6 months Is less than 30MPa and the hardening hardening amount of 30MPa or more to provide a tensile strength of 340 ~ 390MPa high strength hardening hardening type cold rolled steel sheet, plated steel sheet and a manufacturing method thereof.

In the present invention, by weight%, C: 0.0025 to 0.0035%, Si: 0.02% or less, Mn: 0.2 to 1.2%, P: 0.01 to 0.05%, S: 0.01% or less, Al: 0.08 to 0.12%, N: 0.0025% or less, Ti: 0.005 to 0.018%, Mo: 0.01 to 0.1%, B: 0.0005 to 0.0015% balance Fe and other unavoidable impurities,

Ti ^* (Effective Ti) = Total Ti-(48/14) N (48/32) S ≤ 0 between Ti, N and S, DBTT = 803P-24.4Mn-58 between Mn and P A small hardened type cold rolled steel sheet and a small hardened type cold rolled steel sheet characterized by satisfying a relationship of ≤ -30 (° C) and a relationship of P ≤-0.048 * log _e (Al)-0.07 between Al and P. Provided is a plated steel plate.

Furthermore, in the present invention, by weight%, C: 0.0025 to 0.0035%, Si: 0.02% or less, Mn: 0.2 to 1.2%, P: 0.01 to 0.05%, S: 0.01% or less, Al: 0.08 to 0.12%, N : 0.0025% or less, Ti: 0.005 ~ 0.018%, Mo: 0.01 ~ 0.1%, B: 0.0005 ~ 0.0015% balance Fe and other unavoidable impurities,

Ti ^* (Effective Ti) = Total Ti-(48/14) N-(48/32) S ≤ 0 between Ti, N and S, DBTT = 803P-24.4 Mn-between Mn and P A heating step of heating at a temperature of 1200 ° C. or higher for a steel slab satisfying a relationship of 58 ≦ −30 (° C.) and satisfying a relationship of P ≦ −0.048 * log _e (Al) −0.07 between Al and P, Hot rolling step of hot finishing rolling at 900 ~ 950 ° C., winding step of winding at 600 ~ 650 ° C., after removing the scale after air cooling, cold rolling step of cold rolling at 70 ~ 80% reduction rate, 750 ~ 830 ° C. It provides a method for producing a hardened hardened cold-rolled steel sheet comprising a continuous annealing step and a temper rolling step to temper rolling at a reduction ratio of 1.2 ~ 1.5%.

Furthermore, the present invention is in weight%, C: 0.0025 to 0.0035%, Si: 0.02% or less, Mn: 0.2 to 1.2%, P: 0.01 to 0.05%, S: 0.01% or less, Al: 0.08 to 0.12%, N : 0.0025% or less, Ti: 0.005 ~ 0.018%, Mo: 0.01 ~ 0.1%, B: 0.0005 ~ 0.0015% balance Fe and other unavoidable impurities,

Ti ^* (Effective Ti) = Total Ti-(48/14) N-(48/32) S ≤ 0 between Ti, N and S, DBTT = 803P-24.4 Mn-between Mn and P For a steel slab satisfying the relationship of 58 ≤ -30 (° C), a heating step of heating to a temperature of 1200 ° C or more, a hot rolling step of hot finishing rolling at 900 to 950 ° C, a winding step of winding up to 600 ° C or less, and a natural After cooling, the scale is removed, and a cold rolling step of cold rolling at a reduction ratio of 70 to 80%, an annealing step of continuous annealing at 750 to 830 ° C., and a temper rolling step of temper rolling at a reduction ratio of 1.2 to 1.5% are included. Provided is a method for producing a hardened hardened cold rolled steel sheet.

Furthermore, the present invention provides a method for producing a hardened hardened plated steel sheet further comprising a hot dip galvanizing step in the hardened hardened cold rolled steel sheet.

The hardening hardening type cold rolled steel sheet and the plated steel sheet according to the present invention have excellent aging resistance at room temperature, and at the same time, the hardening hardening amount of 30MPa or more and the tensile strength of 340MPa class are suitable for use in various automotive parts.

In the present invention, by weight% C: 0.0025-0.0035%, Si: 0.02% or less, Mn: 0.2-1.2%, P: 0.01-0.05%, S: 0.01% or less, Soluble Al: 0.08-0.12%, N: 0.0025% or less, Ti: 0.005-0.018%, Mo: 0.01-0.1%, and B: 0.0005-0.0015% remainder Fe and other unavoidable impurities, wherein Ti, N and S satisfy the following Equation 1. Homogenizing heat treatment at 1200 ℃ or higher for the controlled Al-killed steel and the slab of the composition, finishing hot rolling at a temperature range of 900 to 950 ° C, winding after winding in a temperature range of 600 to 650 ° C, followed by natural cooling. It provides a method for producing hardened steel. At this time, in order to prevent surface defects by selective oxidation of the surface of the hot-rolled steel sheet, the content of P and Al is controlled through a relational expression as in Equation 2 below.

[Equation 1]

Ti ^* (Effective Ti) = Total Ti-(48/14) N-(48/32) S ≤ 0

[Equation 2]

P ≤-0.048 * log _e (Al)-0.07

However, in another aspect of the present invention, there is provided a method of winding up to 600 ° C. or less and naturally cooling between P and Al without satisfying the condition of Equation 2 above.

The steel sheet of the composition range was wound up to 600 ° C. or less after hot rolling, or naturally cooled, or water-cooled within 30 minutes after winding up to 600 to 650 ° C. to prepare a hot rolled coil. Then, the scale was removed by hydrochloric acid solution. Cold rolling is carried out at a reduction ratio of ˜80%, and the steel sheet is produced by continuous annealing at a temperature range of 750˜830 ° C. and temper rolling at a reduction ratio of 1.2 to 1.5%.

After annealing the steel sheet produced by the above method, the grain size after annealing is finely controlled to 9 or more ASTM number (hereinafter simply referred to as "No."), so that the hardening hardening amount (BH) is 30 MPa or more, and AI (Aging Index) The value has a characteristic of 30 MPa or less. Furthermore, by controlling the content of Mn and P to satisfy the following Equation 3 in order to secure excellent DBTT characteristics, it is possible to provide a high-strength hardened hardened type cold rolled steel sheet and a hot-dip coated steel sheet with surface properties and secondary work brittleness. Can be.

[Equation 3]

DBTT = 803P-24.4Mn-58 ≤ -30 (℃)

In general, when C or N is added to steel materials, carbonitrides such as TiN, AlN, TiC, Ti ₄ C ₂ S ₂ , and NbC are formed by combining with precipitate forming elements such as Al, Ti, or Nb in the hot rolling step. However, carbon or nitrogen, which cannot be combined with these carbonitride-forming elements, is present in solid solution in the steel, which affects the hardening hardening or aging. In particular, since nitrogen has a much higher diffusion rate than carbon, deterioration of aging resistance is very fatal compared to an increase in BH. Therefore, in general, it is often desirable to remove nitrogen from the steel as much as possible. In particular, since Al or Ti is precipitated with nitrogen at a higher temperature than carbon at first, there is no big problem even if it is judged that there is little BH or aging due to nitrogen in steel.

However, C is an element that is indispensable to steel and its properties are determined by its content. In particular, the role of carbon is very important in the field of hard hardened steel, and the hard hardenability and aging resistance may change due to the presence of a small amount of solid solution C present in the steel.

However, the effects on the hardening hardening and aging may vary depending on the location of the solid solution Cs present in the steel, that is, at the grain boundaries or within the grains. For example, solid solution C, which can often be measured by internal friction tests, is mainly present in grains and is relatively free to move, affecting aging characteristics in combination with operational potential. Aging index, or AI (Aging Index), is an item that evaluates the aging characteristics. In general, if the AI value is 30 MPa or more, aging defects may occur before the six months of normal temperature, and may cause serious defects in press working.

The solid solution Cs present in the grain boundaries are difficult to detect by vibration test methods such as internal friction because they exist in grain boundaries, which are relatively stable regions. Since the solid solution Cs present in the grain boundary are in a relatively stable position, they have little effect on low-temperature aging characteristics such as AI, and are activated and influence under high temperature baking conditions such as baking hardening. Therefore, employment C in grains affects both aging and harden hardening simultaneously, but employment C in grains only affects harden hardenability.

However, because the grain boundary is a relatively stable region, all the solid solution Cs present in the grain boundary do not affect the hard hardening, and when 50% of the solid solution C in the grain boundary is activated under high temperature baking conditions, the hard hard hardness is affected. It is known. If such a state of solid solution C is properly controlled, that is, the added solid solution C is controlled to exist at grain boundaries rather than within grains, it is predicted that aging resistance and baking hardening can be secured simultaneously.

To this end, it is important to control the grain size as well as to properly manage the amount of carbon added to the steel. This is because when the amount of carbon added is very large or small, it is often difficult to secure appropriate baking hardening and aging resistance even if the presence position of the solid solution C is controlled.

Figure 1 shows the relationship between the bake hardening amount (BH) value and the aging index (AI value) according to the grain size change. As can be seen in FIG. 1, the grain size number (No) increases, so that the grain size becomes finer, the decrease in AI value against the BH value is remarkable. As a result, the BH-AI value gradually increases, and thus the aging resistance is excellent. . Based on the results of FIG. 1, the present inventors attempted to refine the annealing grain size below an appropriate level in order to control so as to distribute the solid solution C present in the steel as much as possible within the grain boundaries. As a result, the grain size for maximizing the aging resistance while minimizing the degradation of the baking hardenability was determined by ASTM No. It was found that controlling to 9 or more is preferable.

On the other hand, even if a large amount of dissolved carbon is distributed within the grain boundary, the total amount of carbon added to the steel must be strictly controlled. This is because if the added carbon content is excessively increased, even if the grain size becomes fine, the amount of solid solution carbon present in the grain increases in proportion to the total amount of carbon added, which may increase the amount of solid solution carbon in the steel and deteriorate room temperature aging resistance. to be.

However, even if the carbon content is controlled in this way, carbides such as TiC may be formed when Ti is added in an amount of more than necessary to form precipitates such as TiN or TiS in the Ti-added ultra low carbon steel as described herein. Furthermore, under such conditions, there is also a problem that it is difficult to control the appropriate amount of solid solution carbon since the amount of solid solution carbon remaining in the steel changes with the change of Ti content.

Therefore, in order to overcome this problem, the present invention is to control the amount of Ti added to less than the amount combined with S and N, as shown in Equation 1, so that all the added carbon can remain in the steel.

In addition, in the present invention, in addition to Ti addition, in order to more stably secure the baking hardening and aging resistance, the effect of improving the aging resistance by AlN precipitates and Mo addition through Al addition was considered. In general, in the Ti-added steel with low Al content, nitrogen is mostly coarsened as TiN or AlN at a high temperature of 1300 ° C. or higher, and thus does not significantly affect the solid solution effect or grain refinement. Therefore, such AlN shows only the effect of removing the solid solution nitrogen in the steel like TiN precipitate.

However, the present inventors conducted various experiments according to the present invention, when the addition of more than 0.08% Sol.Al becomes very fine precipitates of AlN and serves as a kind of barrier to prevent the growth of crystal grains during annealing recrystallization, Sol.Al It can be seen that finer grains can be obtained than Ti-added steel not added, and thus, the hardening hardening property can be increased without changing the AI value. In addition, Mo has affinity with solid solution carbon in the steel, and if an appropriate amount of Mo is added, it is possible to suppress the diffusion of the solid solution carbon into the potential even if the steel is kept for a long time, which is advantageous in aging resistance, and Mo improves the binding strength of grain boundaries. It was found to be very effective in improving the secondary brittleness. Therefore, in the present invention, by adding a certain amount of Mo to obtain the effect of improving the aging resistance and secondary processing brittleness.

Hereinafter, the component system (wt%) constituting the steel of the present invention will be described in detail.

Carbon (C) is an element showing solid solution hardening and baking hardening. If the carbon content is less than 0.0025%, the tensile strength is insufficient due to the very low carbon content, and even if the Ti addition amount satisfies Equation 1, the absolute carbon content in the steel is too insufficient to obtain sufficient improvement in the bake hardenability. In addition, there is no site competition effect between employment C-Ps, which is very deteriorating in terms of secondary processing brittleness. On the other hand, when the amount of carbon exceeds 0.0035%, the carbon content is excessively increased to increase the grain refining effect, so that the hardening hardening property is excessively high and the secondary processing brittleness may be improved, but the room temperature resistance is deteriorated due to the remaining amount of carbon solution. In the press molding, a stretcher strain may be generated to reduce moldability and ductility. Therefore, in the present invention, the total amount of carbon added to satisfy these conditions was limited to 0.0025 to 0.0035%.

Silicon (Si) is an element that increases strength, but the strength increases as the amount added increases, but ductility deterioration is remarkable. In particular, excessive addition of Si may deteriorate the hot-plating property, so it is advantageous to add it as low as possible in the present invention. . Therefore, in order to prevent such material deterioration and plating property deterioration, the amount of Si added is limited to 0.02% or less.

Manganese (Mn) is an element that refines particles without damaging ductility, and precipitates S in steel completely with MnS to prevent hot brittleness due to the formation of FeS and to strengthen the steel. If the Mn content is less than 0.2%, it is difficult to secure proper tensile strength.However, if the Mn content exceeds 1.2%, strength increases rapidly due to solid solution strengthening and moldability deteriorates. Since a large amount of oxide is formed on the surface, there is a concern that a large amount of plating defects such as deterioration of plating adhesiveness and streaks are generated, which is not good for the quality of the final product, so the amount of addition is limited to 0.2 to 1.2%.

Phosphorus (P) is a substitution-type alloy element having a high solid solution strengthening effect, and serves to improve in-plane anisotropy and strength. In addition, P refines the hot-rolled sheet grains and promotes the development of the (111) aggregate structure, which is advantageous for the improvement of the average r value in the future annealing step. In particular, as the phosphorus content increases due to the site competition effect with carbon, the baking hardenability tends to increase. However, P has two problems as follows. First, since P promotes selective oxidation along the grain boundaries of the steel sheet at high temperatures such as hot rolling, the surface is dropped during rolling when the selective oxidation is intensified. May cause defects. In addition, the selective oxidation phenomenon may increase more rapidly in the presence of Al in the steel components may cause more danger.

On the other hand, the inventors have found that such surface defects are closely related to the hot rolling temperature. As a result of the study of the present inventors, it was found that when wound and cooled at a high temperature of 750 ° C., a large amount of P or Al-based fine oxides existed directly under the cross section, and these oxides acted as a cause of linear defects during hot dip plating. Therefore, when winding a steel material added a large amount of P and Al at the same time, it is necessary to limit the content of P and Al. Furthermore, in the case of performing water cooling rather than slow cooling in the hot rolled winding process or when the coiling temperature is within 600 ° C., since the growth time of grain boundary oxide is limited, the development degree of grain boundary oxide is weak and surface embrittlement phenomenon is alleviated. In this case, it is possible to suppress surface scratches due to surface layer dropping and dropping material during cold rolling.

Therefore, in the case of not cooling the water after the hot rolled up in the manufacturing process proposed in the present invention, the content of P is limited by the following equation.

[Equation 2]

P ≤-0.048 * log _e (Al)-0.07

In addition, P has a problem in that the secondary work brittleness deteriorates due to the weakening of the binding force of the grain boundary when the amount of addition increases more than a certain level. FIG. 2 shows observations of surface defects that may occur when the P and Al contents are simultaneously manufactured by hot-dip steel sheets using a scanning electron microscope (SEM). As shown in FIG. 2, the defects existed in the form of a surface drop off in a point or linear form. Observing the surface of these defects precisely, it can be seen that the grain boundary is relatively weak. As a result of the present inventors, it was found that such surface defects are closely related to the hot rolling temperature. That is, as can be seen in Figure 3, when the slow cooling by winding at a high temperature of 750 ℃ a large amount of P or Al-based fine oxides are generated directly under the cross-section, these oxides act as a cause of linear defects during the hot dip plating. Therefore, it is necessary to control the content of P and Al when winding the steel at the same time a large amount of P and Al is added, it was possible to derive the relationship between the P and Al content as shown in Equation 3. Apart from this, in the case of performing water cooling rather than slow cooling in the hot rolled winding process, even though steel having a high P and Al content, as shown in FIG. 4 (a), almost no fine oxides were found directly under the surface, and thus the plating surface. No defects were seen either. FIG. 5 shows X when the embrittlement of the surface of the hot rolled steel sheet was observed at a coiling temperature of 620 ° C. with respect to the specimens in which P and Al components were variously changed. In order to be safe from surface embrittlement, it was found that the composition of Al and P should be appropriately controlled. By such experiment, in the present invention, the composition of Al-P is regulated as in Equation 2 above.

In general, the molding of the parts performed in the automobile company goes through several times of repeated press processing, and the second cracking occurs in the processing cracks after the first press processing. . These cracks occur due to the presence of phosphorus (P) in the grain as a grain boundary, which weakens the binding force of the grains. Thus, cracking occurs around grain boundaries. It is preferable to add. However, compared to the increase in strength, the least decrease in elongation is the employment element P, and above all, the cost is low. Therefore, in order to improve the strength of the steel to some extent, but in order to solve the secondary work brittleness problem caused by P in the present invention, the P content was limited to 0.01 ~ 0.05% level, at this time to reduce the strength due to P decrease The content of Mn was considered together for the purpose of supplementation. Figure 5 shows the relationship between the DBTT characteristics by the addition of Mn and P, in order to ensure the DBTT -30 ℃ or less to be secured in the present invention, the P content is expressed in relation to the content of Mn and P as shown in Equation 2 It can be seen that it must be satisfied.

[Equation 3]

DBTT = 803P-24.4Mn-58 ≤ -30 (℃)

Sulfur (S) is an element that should be precipitated as sulfide of MnS at high temperature to prevent hot brittleness by FeS. However, if the S content is excessive, the precipitated MnS and the remaining S may embrittle the grain boundary, causing hot brittleness. In addition, even if the amount of S added is enough to completely precipitate the MnS precipitates, if the amount is too large, material degradation may occur due to excessive precipitates, so the amount is limited to 0.01% or less.

Aluminum (Al) is usually added for deoxidation of steel, but in the present invention steel, the grain refining effect and Al hardenability by AlN precipitation are improved. In general, in the Ti-added steel, nitrogen precipitates mostly coarse with TiN at a high temperature of 1300 ° C. or higher. However, AlN precipitation by Sol. Al occurs in steel in which Ti is added in a small amount of 30 ppm or less. As a result of various experiments using the inventive steel, when Sol.Al is present in the range of 0.02% to 0.06%, which is a normal level, it simply serves to fix solid nitrogen, but when added to 0.08% or more, As the precipitate becomes very fine and acts as a kind of barrier to prevent grain growth during annealing recrystallization, grains become finer than Nb-added steel without Sol.Al. You can get the effect. Therefore, in order to obtain such an effect, Al content is added at least 0.08% or more. However, when Al is excessively added in excess of 0.12%, the surface quality may be degraded due to the increase of oxidation inclusions during steelmaking, and the production cost may be increased due to excessive addition of Al. Therefore, the addition amount is limited to 0.08 to 0.12%. do.

Nitrogen (N) is present in solid solution before or after annealing, thereby degrading the formability of the steel, and it is necessary to fix with Ti or Al because the aging deterioration is much larger than other invasive elements. When Nb is properly added together with a small amount of Ti as in the present invention steel, excessive addition of nitrogen causes generation of solid solution nitrogen in the steel. In general, nitrogen has a very fast diffusion rate compared to carbon, so when present as solid nitrogen, the deterioration of room temperature aging resistance is very serious compared to solid solution C. In addition, since the yield strength increases and elongation and r value deteriorate due to the remaining of solid solution nitrogen, the content thereof needs to be limited to 0.0025% or less as in the present invention.

Ti is a carbonitride forming element and forms nitrides such as TiN, sulfides such as TiS or Ti ₄ C ₂ S ₂ and carbides such as TiC in the steel. In the present invention, the Ti content is controlled as in Equation 1 within 0.005% to 0.018% to maintain solid solution carbon in the steel. If the Ti content is less than 0.005%, Equation 1 is easily satisfied, but the Ti content is so small that the grain size increases, so that it is difficult to obtain the effect of improving the aging resistance due to the grain refinement effect pursued by the present invention, and the solid carbon in steel. As a result, moldability conditions such as elongation and r value deteriorate. On the other hand, when the Ti content exceeds 0.018%, the conditions of Equation 1 are not satisfied, so that the solid solution carbon in the steel is reduced, which leads to a decrease in the hard hardenability.

Mo is one of the very important elements in the present invention. Mo is dissolved in steel to improve strength or to form Mo-based carbides. In particular, Mo in the present invention increases the binding force of the grain boundary when present in the solid solution state serves to improve the grain breakage problem caused by phosphorus (P), that is, secondary processing brittleness. In addition, the diffusion of carbon is suppressed by the affinity with solid solution C, and the aging resistance is also improved. Therefore, Mo is added 0.01% or more in the present invention. However, when the Mo content is excessively excessively exceeding 0.1%, the improvement effect of the secondary work brittleness and the aging resistance is saturated and the economical efficiency is low. Therefore, Mo content is limited to the range of 0.01 to 0.1%.

B exists as an invasive element in steel and is dissolved in grain boundaries or combines with nitrogen to form nitrides of BN. B is an element having a great influence of the material compared to the added amount, and it is necessary to strictly limit the added amount. In other words, even if a small amount of B is added to the steel, segregation at grain boundaries may improve secondary work brittleness, but if it is added in a predetermined amount or more, material degradation occurs due to an increase in strength and a significant decrease in ductility. Inventive textbooks are added with B limited to 0.0005 to 0.0015%.

The slab having the above composition is heated to a temperature of 1200 ° C. or higher at which austenite structure can be sufficiently homogenized before hot rolling, and the hot rolling is finished at 900 to 950 ° C. in the range immediately above Ar ₃ .

If the slab temperature is less than 1200 ℃, the structure of the steel does not become uniform austenite grains, and the mixing occurs, which may cause deterioration of the material.If the hot rolling finish temperature is less than 900 ℃, the top of the hot rolled coil may be The tail and edges become single-phase regions, which may increase the in-plane anisotropy and deteriorate the moldability. When it exceeds 950 ° C, remarkable coarse grains occur, such as orange peel or the like on the surface after processing. Defects can occur.

The first aspect of the present invention is to control the relationship between Al-P and 600 to 600 deg. A winding step of 650 ° C. is performed. In this case, when the coiling temperature exceeds 650 ° C., the grain size increases after annealing, and sufficient grain refinement effect cannot be obtained even if other conditions are satisfied, and grain boundary segregation of P is increased to deteriorate secondary work brittleness. do. On the other hand, when the coiling temperature is less than 600 ° C., the selective oxidation of the hot rolled pole surface layer by Al and P becomes small, but the hot rolled rolling load becomes high. In the case of cooling by air cooling after winding, it is important to satisfy the relationship of Expression 1, which is the Al-P relationship.

Furthermore, according to the 2nd viewpoint of this invention, a process can be performed by setting a winding temperature to 600 degrees C or less. The second aspect does not have to satisfy the relational expression of Al-P, and the surface embrittlement can be prevented by adjusting only the winding temperature. The second aspect may be advantageous because the coiling temperature is lower than the first aspect and is not limited to the Al-P relationship, but a low coiling temperature is not always preferable for the process, and thus the type or property of subsequent processes Accordingly, the bake hardened steel of the present invention can be produced under suitable winding conditions.

After hot rolling, the steel is pickled in the usual manner and then cold rolled at a high cold rolling rate of 70 to 80%. The reason for the high cold rolling rate of 70% or more in the present invention is to improve moldability, in particular, r value as well as improving aging resistance by grain refining effect. On the other hand, if the cold rolling rate exceeds 80%, the grain refining effect is large, but the grain size becomes too large due to excessive rolling rate, which causes hardening of the material, and the r value may gradually decrease due to excessive cold rolling rate increase. have.

Cold-rolled steel is subjected to a continuous annealing operation by a common method in the temperature range of 750 ~ 830 ℃. If the annealing temperature is less than 750 ℃ due to the presence of unrecrystallized grains yield strength can be increased and elongation and r value may be deteriorated. However, if the annealing temperature exceeds 830 ℃ formability can be improved, but the grain size of the ASTM No. It becomes smaller than 9, AI value is 30 MPa or less, and aging resistance deteriorates.

The temper rolling is carried out at a reduction ratio of 1.2 to 1.5%, which is somewhat higher than the usual temper rolling ratio, for the purpose of securing appropriate bake hardening resistance and room temperature aging resistance using the calcined hardened steel manufactured by the above-described manufacturing method. The reason why the temper rolling ratio is set higher than 1.2% is to prevent the deterioration of room temperature aging due to solid solution C in steel. However, if the temper rolling rate exceeds 1.5%, even though the room temperature aging resistance can be improved, the temper rolling rate is high, resulting in work hardening, resulting in deterioration of the material. Since the peeling of the plating layer may occur, it is limited to the temper rolling ratio of 1.2 ~ 1.5%.

The calcined hardened steel thus manufactured may be manufactured as a plated steel sheet by performing a hot dip galvanizing treatment on the surface layer. In this case, the hot dip galvanizing treatment is not particularly limited and may be processed by a conventional method.

The present invention will be described in detail through the following examples.

(Example)

Table 1 below shows the chemical composition of the invention steels and comparative steels that strictly control the amount of C, P, Ti, Sol.Al and Mo to satisfy the surface properties and material properties at the same time, steel 1-6 is the invention steel Steel 9-15 is comparative steel.

Steel grade Chemical composition (wt%) C Mn P S Sol.Al Ti Nb N Mo B Inventive Steel 1 0.0031 0.59 0.031 0.0082 0.097 0.009 - 0.0017 0.034 0.0005 Inventive Steel 2 0.0032 0.73 0.015 0.0081 0.098 0.014 - 0.0024 0.048 0.0005 Inventive Steel 3 0.0033 0.75 0.022 0.0058 0.105 0.015 - 0.0019 0.061 0.0007 Inventive Steel 4 0.0028 0.61 0.031 0.0083 0.118 0.013 - 0.0015 0.059 0.0005 Inventive Steel 5 0.0027 0.98 0.036 0.0070 0.105 0.010 - 0.0017 0.061 0.0008 Inventive Steel 6 0.0029 1.01 0.04 0.0063 0.089 0.011 - 0.0020 0.052 0.0009 Comparative Steel 1 0.0063 0.64 0.034 0.0071 0.093 0.014 - 0.0017 0.021 0.0007 Comparative Steel 2 0.0032 0.61 0.046 0.0085 0.041 0.050 - 0.0015 0.019 0.0005 Comparative Steel 3 0.0010 0.65 0.042 0.0072 0.075 0.011 - 0.0019 0.059 0.0008 Comparative Steel 4 0.0031 0.93 0.036 0.0089 0.074 0.017 0.042 0.0017 0.021 0.0006 Comparative Steel 5 0.0032 0.049 0.072 0.0066 0.081 0.018 - 0.0022 0 0.0007 Comparative Steel 6 0.0029 0.99 0.069 0.0078 0.041 0.011 - 0.0021 0.025 0 Comparative Steel 7 0.0033 0.98 0.14 0.0078 0.098 0.012 - 0.0023 0.031 0

Hot rolling at hot rolled odor temperature of 610 ~ 640 ℃ using the steel of Table 1, rolling at 70-78% cold rolling rate and continuous annealing at annealing temperature of 780-820 ℃, plating and After alloying at about 530 ° C, rough rolling was carried out at a roughness reduction rate of about 1.5% to evaluate the plating defects, as well as to evaluate the BH value, AI value, grain size, and secondary work brittleness. One result is shown.

Steel grade CT Annealing Temperature TS BH (MPa) AI (MPa) Grain size No. DBTT (℃) Plating defect Inventive Steel 1 620 ℃ 800 ℃ 352.8 44.7 25.8 10.5 -40 ⊙ Inventive Steel 2 620 ℃ 810 ℃ 355.3 43.2 26.8 9.7 -45 ⊙ Inventive Steel 3 620 ℃ 790 ℃ 365.1 44.3 27.9 9.9 -50 ⊙ Inventive Steel 4 610 ℃ 800 ℃ 361.2 46.4 27.6 10.5 -40 ⊙ Inventive Steel 5 640 ℃ 790 ℃ 357.9 45.2 26.9 10.3 -45 ⊙ Inventive Steel 6 620 ℃ 820 ℃ 359.7 47.6 22.3 11.2 -40 ⊙ Comparative Steel 1 620 ℃ 810 ℃ 380.1 72.1 59.3 11.2 -50 ⊙ Comparative Steel 2 640 ℃ 800 ℃ 349.1 19.2 10.3 8.1 -40 ⊙ Comparative Steel 3 620 ℃ 810 ℃ 351.2 0.0 0.0 8.2 -50 ⊙ Comparative Steel 4 630 ℃ 800 ℃ 391.1 5.1 0.0 9.4 -50 ⊙ Comparative Steel 5 620 ℃ 810 ℃ 351.4 38.3 29.4 10.9 0 × Comparative Steel 6 630 ℃ 790 ℃ 367.4 39.2 27.1 9.2 -15 △ Comparative Steel 7 640 ℃ 810 ℃ 401.2 40.2 20.4 10.2 15 ×

(However, plating defects: ⊙ (within 10 defects per km), △ (10-100 defects per km), × (100 defects per km))

The grain size of the inventive steel exhibited through the above-mentioned conditions was ASTM No. 9.7 to 11.5 (average grain size 6.7-12.0 μm) as shown in ASTM No. It turns out that all the conditions which are nine or more are satisfied. In addition, the invention steels satisfying the conditions of the present invention have a BH value of 40.4 to 47.6 MPa and an AI value of 10.2 to 27.9 MPa. In 1.9 it can be seen that the DBTT conditions presented in the present invention steel satisfies less than -30 ℃. Furthermore, through the optimization of the P content, it was found that the plating defects of the hot-dip plated material had a surface defect per 10 km per coil within 10 km.

On the other hand, in the case of Comparative Steel 1, the C content was 0.0063%, which satisfies the process conditions such as hot rolling temperature and annealing temperature. The value is very high and the AI value is 59.3 MPa, which is beyond the proper range suggested by the present invention steel.

And Comparative steel 2, Sol.Al and Ti content is out of the conditions of the present invention steel. The content of Ti is such that the effective Ti content is very high, the remaining Ti reacts with N and S, and the remaining Ti reacts with C to precipitate TiC, so that almost no solid solution carbon exists in the steel and hardly harden the hard part. In addition, the material is softened by a large amount of Ti, the size of the crystal grains also appears.

Comparative steel 3 had a low C content, coarse grains, and no BH or AI properties. In addition, Comparative Steel 4 was a steel with Nb added, and the grain refinement effect and the BH value could not be improved, and the solid solution C in the steel was formed as NbC precipitates, and the BH and AI properties were not obtained. .

Comparative steel 5 had a higher P content than the level suggested by the present invention steel, and B was not added at all so that the secondary workability improvement effect by B did not appear. In addition, the P content was high as 0.072%, and the interaction between P and Al occurred in the inventive steel, which is a high Al addition steel, thereby increasing the surface oxide from the hot rolling step. Therefore, due to the increase in oxide, a large amount of surface defects such as linear defects and the like occurred in the production of a hot-dip plating material such as 100 or more per km.

Comparative steel 6 had a P content of 0.069%, deviating from the conditions of the present invention steel, a low addition of Sol.Al, and no addition of B. Therefore, as shown in Table 2, the BH and AI properties are satisfactory, but the DBTT characteristics deteriorate due to the high P content and no addition of B. The surface defects also exceed within 10 km / km of the coil, which is the level of the present invention.

Comparative steel 8 is 0.14% of P content, far exceeding 0.01 to 0.05% of the component range suggested by the present invention steel, and B is not added. Although small amount of DBTT property is improved by Mo, the addition amount of P is very high, and the improvement effect is limited. In particular, the addition of B has lost the effect of improving DBTT property. Due to this effect, the DBTT was very high as 15 ° C., and the surface defects of the molten plating material were greatly increased due to the excessive addition of P.

1 is a graph showing the effect of grain size on BH and AI values.

2 is a micrograph analyzing the microstructure of the linear binding.

3 is a fine oxide formed on the metal surface grain boundary surface of the 750 ℃ winding material and EDS analysis results thereof.

Figure 4 is a microphotograph showing the fine oxide distribution (a is quenching, b is 600 ℃ slow cooling, c is 750 ℃ slow cooling) of the metal surface layer according to the winding temperature and the cooling method.

5 is a graph showing the defect generation and non-occurrence zone according to the content of P and Al.

6 is a graph showing the change in secondary work brittleness characteristics according to the content of P and Mn.

7 is a graph showing the effect of drawing ratio on the secondary work brittle resistance (DBTT) for each material.

Claims (7)

In weight%, C: 0.0025 to 0.0035%, Si: 0.02% or less, Mn: 0.2 to 1.2%, P: 0.01 to 0.05%, S: 0.01% or less, Al: 0.08 to 0.12%, N: 0.0025% or less, Ti: 0.005 ~ 0.018%, Mo: 0.01 ~ 0.1%, B: 0.0005 ~ 0.0015% balance Fe and other unavoidable impurities, Between the Ti, N and S Ti ^* (Effective Ti) = Total Ti-(48/14) N-(48/32) S ≤ 0 Between Mn and P DBTT = 803P-24.4Mn-58 ≤ -30 (℃) And between Al and P P ≤-0.048 * log _e (Al)-0.07 The hardened hardened cold rolled steel sheet, characterized in that to satisfy the relationship. The grain hardening type cold rolled steel sheet according to claim 1, wherein the grain size of the hardened hardened steel is ASTM No. 9 or more. The quench hardening plated steel sheet further comprising a plating layer on the quench hardening type cold rolled steel sheet surface layer of claim 1. In weight%, C: 0.0025 to 0.0035%, Si: 0.02% or less, Mn: 0.2 to 1.2%, P: 0.01 to 0.05%, S: 0.01% or less, Al: 0.08 to 0.12%, N: 0.0025% or less, Ti: 0.005 ~ 0.018%, Mo: 0.01 ~ 0.1%, B: 0.0005 ~ 0.0015% balance Fe and other unavoidable impurities, Between the Ti, N and S Ti ^* (Effective Ti) = Total Ti-(48/14) N-(48/32) S ≤ 0 Between Mn and P DBTT = 803P-24.4Mn-58 ≤ -30 (℃) And between Al and P P ≤-0.048 * log _e (Al)-0.07 For steel slabs that satisfy the relationship of A heating step of heating to a temperature of at least 1200 ° C .; Hot rolling step of hot finish rolling at 900 ~ 950 ℃; Winding step of winding at 600 ~ 650 ° C; After removing the scale after air cooling, cold rolling step of cold rolling at a reduction ratio of 70 ~ 80%; Annealing step of continuous annealing at 750 ~ 830 ° C; And A crude rolling step of rough rolling at a reduction ratio of 1.2 to 1.5%; Method for producing a hardened hardened cold rolled steel sheet comprising a. In weight%, C: 0.0025 to 0.0035%, Si: 0.02% or less, Mn: 0.2 to 1.2%, P: 0.01 to 0.05%, S: 0.01% or less, Al: 0.08 to 0.12%, N: 0.0025% or less, Ti: 0.005 ~ 0.018%, Mo: 0.01 ~ 0.1%, B: 0.0005 ~ 0.0015% balance Fe and other unavoidable impurities, Between the Ti, N and S Ti ^* (Effective Ti) = Total Ti-(48/14) N-(48/32) S ≤ 0 Between Mn and P DBTT = 803P-24.4Mn-58 ≤ -30 (℃) For steel slabs that satisfy the relationship of A heating step of heating to a temperature of at least 1200 ° C .; Hot rolling step of hot finish rolling at 900 ~ 950 ℃; Winding step of winding up to 600 ℃; Removing the scale after natural cooling and cold rolling at a rolling reduction of 70 to 80%; Annealing step of continuous annealing at 750 ~ 830 ° C; And A crude rolling step of rough rolling at a reduction ratio of 1.2 to 1.5%; Method for producing a hardened hardened cold rolled steel sheet comprising a. The method for manufacturing a quench hardened cold rolled steel sheet according to any one of claims 4 to 5, wherein the grain size of the quenched hardened steel is ASTM No. 9 or higher. A method of manufacturing a sinter hardening plated steel sheet, further comprising a hot dip galvanizing step in the sinter hardening cold rolled steel sheet of claim 4.

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