US20080251168A1

US20080251168A1 - Bake-Hardenable Cold Rolled Steel Sheet With Superior Strength and Aging Resistance, Gal-Vannealed Steel Sheet Using the Cold Rolled Steel Sheet and Method For Manufacturing the Cold Rolled Steel Sheet

Info

Publication number: US20080251168A1
Application number: US12/067,617
Authority: US
Inventors: Seong-Ho Han
Original assignee: Posco Co Ltd
Current assignee: Posco Holdings Inc
Priority date: 2005-09-23
Filing date: 2006-09-22
Publication date: 2008-10-16
Also published as: JP4834733B2; JP2009509046A; EP1937853B1; KR100685037B1; EP1937853A4; CN101310031B; EP1937853A1; WO2007035059A1; CN101310031A

Abstract

A bake-hardenable cold rolled steel sheet with high strength and superior aging resistance used for outer panels of an automobile body, a galvannealed steel sheet using the cold-rolled steel sheet, and a method for manufacturing the cold-rolled steel sheet are disclosed. The steel sheet comprises, by weight %, C: 0.0016˜0.0025%, Si: 0.02% or less, Mn: 0.2˜1.2%, P: 0.05˜0.11%, S: 0.01% or less, Sol. Al: 0.08˜0.12%, N: 0.0025% or less, Ti: 0˜0.003%, Nb: 0.003˜0.011%, Mo: 0.01˜0.1%, B: 0.0005˜0.0015%, the balance of Fe and other unavoidable impurities. The steel sheet has superior bake hardenability, aging resistance at room temperature, and secondary work embrittlement resistance.

Description

TECHNICAL FIELD

The present invention relates to a cold-rolled steel sheet for outer panels and the like of an automobile body, a galvannealed steel sheet using the cold-rolled steel sheet, and a method for manufacturing the cold-rolled steel sheet. More particularly, the present invention relates to a high strength cold-rolled steel sheet with superior bake hardenability, aging resistance at room temperature and secondary work embrittlement resistance, a galvannealed steel sheet using the cold-rolled steel sheet, and a method for manufacturing the cold-rolled steel sheet.

BACKGROUND ART

For improvement in fuel efficiency and reduction in weight of automobiles, it has been increasingly demanded to improve the dent resistance of an outer panels and to reduce the thickness thereof by use of a high strength steel sheet for an automobile body.
As used for the outer panel of the automobile body, a cold-rolled steel sheet is required to have good properties in terms of tensile strength, yield strength, press formability, spot weldability, fatigue resistance, corrosion resistance, etc.
In particular, the corrosion resistance has been recently required for extension of lifetime of components for the automobile.
Steel sheets for improvement of the corrosion resistance can be generally classified into two types, i.e. a electroplated steel sheet and a galvannealed steel sheet.
In comparison to the galvannealed steel sheet, although the electroplated steel sheet has better plating characteristics and superior corrosion resistance, it is rarely used due to its high price. Therefore, the galvannealed steel sheet is generally used in the art, and required to have improved corrosion resistance.
In recent years, most of steelmakers in the world have produced the galvannealed steel sheets as materials for the automobiles, and supplied them to automobile manufacturers. Accordingly, new techniques capable of securing superior corrosion resistance above a previous level have been continuously developed and used.
Generally, the steel sheet exhibits opposite characteristics in terms of strength and formability. As steel sheets capable of satisfying both characteristics, there are multi-phase structure cold-rolled steel sheets and bake-hardenable cold-rolled steel sheets.
The multi-phase structure cold-rolled steel can be easily manufactured, and has a tensile strength in the level of 390 MPa or more. Regardless of its high tensile strength as materials for the automobiles, the multi-phase structure cold-rolled steel has a high elongation as a factor indicating stretchability. However, it has a low average r-value as a factor indicating press formability of the automobiles, and comprises excessive amounts of expensive alloying elements such as Mn, Cr and the like, which results in high manufacturing costs.
The bake-hardenable cold-rolled steel acts like mild steel in terms of yield strength upon press forming of the steel which has a tensile strength of 390 MPa or less. Thus, the bake-hardenable cold-rolled steel has excellent ductility, and spontaneously increases in yield strength upon paint baking after press forming. This steel is considered ideal in comparison to previous steel, which is generally deteriorated in formability as the strength of the steel increases.
Bake hardening is a process which employs a kind of strain aging occurring as interstitial elements dissolved in a solid solution state in the steel, such as solute nitrogen or solute carbon, fix dislocations created during deformation. When the steel has high amounts of solute carbon and nitrogen, the bake hardenability value (BH value) advantageously increases, but the natural aging property also increases due to the high amounts of solid solution elements, deteriorating the formability. Thus, it is very important to optimize the amounts of solid solution elements in the steel.
As a method for manufacturing the bake hardenable cold-rolled steel sheet, a batch annealing method and a continuous annealing method are generally used.
Generally, the bake-hardenable cold-rolled steel sheet is produced by batch annealing a low carbon, P-added, Al-killed steel through coiling of a hot-rolled steel sheet at a lower temperature. Specifically, when manufacturing the cold-rolled steel sheet using the Al-killed steel, the hot-rolled steel sheet is coiled at the low temperature of 400˜500° C., followed by batch annealing to have BH value of about 40 to 50 MPa.
This is because the batch annealing enables both of the formability and the bake-hardenability to be obtained more easily at the same time. Meanwhile, in the case of batch annealing, since the P-added Al-killed steel is cooled at a relatively high speed, it is easy to secure the bake-hardenability, but there are problems in that the formability is deteriorated due to a high heating speed and a short annealing process. Thus, the steel sheet produced by batch annealing is restricted in use for the outer panels of the automobile body, which do not require workability.
Recently, with rapid advance in steel manufacturing techniques, it becomes possible to optimize the amount of solid solution elements in the steel and to manufacture bake-hardenable cold-rolled steel sheets with superior formability through addition of intensive carbon and nitride formation elements, such as Ti or Nb, to the Al-killed steel, thereby satisfying the increasing demand for the bake-hardenable cold-rolled steel sheets, which can be used for the outer panels of the automobiles requiring the dent resistance.
Japanese Patent Publication No. (Sho) 61-026757 discloses an ultra low carbon cold-rolled steel sheet, which comprises: 0.0005˜0.015% of C; 0.05% or less of S+N; and Ti and Nb or a compound thereof. Japanese Patent Publication No. (Sho) 57-089437 discloses a method for manufacturing a bake-hardenable cold-rolled steel sheet, which uses Ti-added steel comprising 0.010% or less of C, and has BH value of about 40 MPa or more.
The methods of the disclosures are to impart the bake hardenability to the steel sheet while preventing deterioration in other properties of the steel sheet by appropriately controlling the amount of solid solution elements in the steel through control of the added amount of Ti and Nb or the cooling rate during annealing. However, for the Ti-added steel or Ti and Nb-added steel, it is necessary to strictly control amounts of Ti, N and S during manufacturing of the steel to secure an appropriate BH value, causing an increase of manufacturing costs.
Furthermore, the Nb-added steel described above has problems in that workability is degraded due to hot annealing, and in that manufacturing costs are increased due to addition of the specific elements.
On the other hand, U.S. Pat. Nos. 5,556,485 and 5,656,102 (Bethlehem Steel) disclose methods of manufacturing a bake-hardenable cold-rolled steel sheet using a Ti—V based ultra low carbon steel, which comprises 0.0005˜0.1% of C; 0˜2.5% of Mn; 0˜0.5% of Al; 0˜0.04% of N; 0˜0.5% of Ti; and 0.005˜0.6% of V.
Generally, since V is more stable than the carbon nitride formation elements such as Ti and Nb, it can reduce an annealing temperature. Hence, carbide, such as VC and the like, created during high temperature annealing can impart the bake hardenability via re-melting even with a lower annealing temperature than that for the Nb-based steel.
However, although V can create the carbide such as VC, since it does not sufficiently improve the formability due to its significantly low re-melting temperature, Ti is added at an amount of about 0.02% or more in order to enhance the formability, as disclosed in the publications.
Thus, the methods of U.S. Pat. Nos. 5,556,485 and 5,656,102 suffers from deterioration in aging resistance caused by coarse crystal grains as well as increase in manufacturing costs caused by addition of large amounts of Ti.
Meanwhile, various methods of manufacturing a bake-hardenable cold-rolled steel sheet through addition of alloying elements are disclosed in Japanese Patent Laid-open Publication Nos. (Hei) 5-93502, (Hei) 9-249936, (Hei) 8-40938 and (Hei) 7-278654.
Japanese Patent Laid-open Publication No. (Hei) 5-93502 discloses a method for enhancing the bake hardenability by addition of Sn, and Japanese Patent Laid-open Publication No. (Hei) 9-249936 discloses a method for enhancing the ductility of steel by relieving stress concentration on grain boundaries through addition of V and Nb.
Japanese Patent Laid-open Publication No. (Hei) 8-49038 discloses a method for enhancing the formability through addition of Zr, and Japanese Patent Laid-open Publication No. (Hei) 7-278654 discloses a method for enhancing the formability by increasing the strength while minimizing deterioration of work hardening index (N-value) through addition of Cr.
However, these methods only give attention to improvement in the bake hardenability or the formability, and do not disclose the problem of deterioration in aging resistance resulting from improvement of the bake hardenability, and the problem of the secondary work embrittlement resulting from increase in content of P, which is necessarily added to make the bake hardenable steel high in strength.
Generally, the increase in bake hardenability causes the deterioration in aging resistance at room temperature. In particular, inventors of the present invention have found that, as the content of P added for high strength of the steel increases, the steel is degraded in the secondary work embrittlement resistance even for the bake hardenable steel which comprises solute carbon dissolved in the solid solution state in the steel, and the more the severe the degradation.
For example, when P was added at an amount of 0.07% to produce bake hardenable steel of a tensile strength in the level of 340 MPa, the ductility-brittleness transition temperature (DBTT) of the steel as a reference to judge the secondary work embrittlement was −20° C. at a draw ratio of 1.9. In addition, when P was added at an amount of about 0.09% to produce high strength steel in the level of 390 MPa, the DBTT of the steel was 0˜10° C., which can be considered to be a significantly deteriorated value.
In the methods described above, although boron (B) is added at an amount of about 5 ppm and expected to improve the secondary work embrittlement resistance, the excessive P content in the steel limits the improvement of the secondary work embrittlement resistance through addition of B.
Furthermore, if B is excessively added to the steel to improve the secondary work embrittlement resistance, the secondary work embrittlement resistance is deteriorated due to the excessive content of B. Thus, there is a certain limit in amount of B which can be added to the steel.
Since the steel must has a DBTT of at least −20° C. to prevent the secondary work embrittlement, there are needs to investigate new compositions other than B for the bake hardenable steel and new manufacturing conditions therefor.

DISCLOSURE OF INVENTION

Technical Problem

Therefore, the present invention has been made in view of the above problems, and an object of the present invention is to provide a high strength cold-rolled steel sheet with superior bake hardenability, aging resistance at room temperature and secondary work embrittlement resistance, and a galvannealed steel sheet using the cold-rolled steel sheet.
It is another object of the present invention to provide a method for manufacturing a high strength cold-rolled steel sheet with superior bake hardenability, aging resistance at room temperature and secondary work embrittlement resistance.

Technical Solution

In accordance with one aspect of the present invention, the above and other objects can be accomplished by the provision of a bake-hardenable cold-rolled steel sheet with high strength and superior aging resistance, comprising, by weight %: C: 0.0016˜0.0025%; Si: 0.02% or less; Mn: 0.2˜1.2%; P: 0.05˜0.11%; S: 0.01% or less; Sol. Al: 0.08˜0.12%; N: 0.0025% or less; Ti: 0˜0.003%; Nb: 0.003˜0.011%; Mo: 0.01˜0.1%; B: 0.0005˜0.0015%; and the balance of Fe and other unavoidable impurities, wherein the steel sheet satisfies Equation 1:
C*[amount of solute carbon in grain boundaries (GB-C)+amount of solute carbon in crystal grains (G-C)]=Total C (ppm)−C in NbC=8˜15 ppm (1)
[in Equation 1, GB-C (that is, the amount of solute carbon in the grain boundaries) is 5˜10 ppm, and G-C (that is, the amount of solute carbon in the crystal grains) is 3˜7 ppm], and wherein the steel sheet has an ASTM grain size (hereinafter referred to as “ASTM No.” of 9 or more, a bake hardening degree (BH) of 30 MPa or more, an aging index (AI) of 30 MPa or less, and a DBTT of −30° C. or less at a draw ratio of 2.0, the bake hardenability value (BH) and the aging index (AI) satisfying Equations 2 and 3:
BH=50−(885×Ti)−(1589×Nb)+(62×Al) (2)
AI=44−(423×Ti)−(2119×Nb)−(125×Mo) (3)
In accordance with another aspect of the invention, a galvannealed steel sheet using the cold-rolled steel sheet according to the present invention is provided.
In accordance with yet another aspect of the invention, a method for manufacturing a bake-hardenable cold-rolled steel sheet with high strength and superior aging resistance, comprising: performing homogenization heat treatment for an Al-killed steel slab at 1200° C. or more, the steel slab comprising, by weight %: C: 0.0016˜0.0025%, Si: 0.02% or less, Mn: 0.2˜1.2%, P: 0.05˜0.11%, S: 0.01% or less, Sol. Al: 0.08˜0.12%, N: 0.0025% or less, Ti: 0˜0.003%, Nb: 0.003˜0.011%, Mo: 0.01˜0.1%, B: 0.0005˜0.0015%, and the balance of Fe and other unavoidable impurities; hot rolling the steel slab with finish rolling at a finish rolling temperature of 900˜950° C. to form a hot-rolled steel sheet, followed by coiling the hot-rolled steel sheet at a temperature of 580˜630° C.; cold rolling the hot-rolled steel sheet at a reduction ratio of 75˜80%; continuously annealing the cold rolled steel sheet at a temperature of 770˜830° C.; and temper rolling the anneal steel sheet at a reduction ratio of 1.2˜1.5%.

BRIEF DESCRIPTION OF THE DRAWINGS

The above and other objects, features and other advantages of the present invention will be more clearly understood from the following detailed description taken in conjunction with the accompanying drawings, in which:

FIG. 1 is a graph describing influence of a grain size on bake hardenability and aging index;

FIG. 2 is a graph describing influence of an amount of solute carbon in steel on the bake hardenability; and

FIG. 3 is a graph describing a result of an internal friction test to Inventive steel No. 6.

BEST MODE FOR CARRYING OUT THE INVENTION

Preferred embodiments of the invention will now be described in detail.
Carbon or nitrogen in steel generally combines with precipitate formation elements such as Al, Ti, Nb, etc. in the steel during hot rolling, forming carbides and nitrides such as TiN, AlN, TiC, Ti₄C₂S₂. NbC, etc. However, when carbon or nitrogen does not combine with the precipitation formation elements in the steel, it exists as solid solutions of carbon or nitride (hereinafter, solute carbon or solute nitrogen) in the steel, and influences the bake hardenability and the aging resistance of the steel.
In particular, since nitrogen has a higher diffusion rate than that of carbon, it is very detrimental to the aging resistance in comparison to an improvement of the bake hardenability. Hence, it is general in the art to remove nitrogen from the steel as much as possible. In particular, since Al or Ti is precipitated along with nitrogen prior to carbon at high temperatures, it can be generally concluded that there is substantially no influence of nitrogen in the steel on the bake hardenability and the aging resistance.
Meanwhile, carbon is an essential element for the steel, and determines the characteristics of the steel depending on the carbon content in the steel. For the bake hardenable steel sheet according to the present invention, carbon has a very important role, and only a small amount of solute carbon is allowed to remain in the steel as an attempt to improve the bake hardenability and the aging resistance.
However, influence of the solute carbon on the bake hardenability and the aging resistance can be changed according to locations of the solute carbon in the steel, that is, whether the solute carbon resides in grain boundaries or in crystal grains.
That is, with an internal friction test to detect the solute carbon, it can be found that the solute carbon generally exists in the crystal grains and moves relatively freely. Thus, the solute carbon can combine with movable dislocations, and affect aging properties.
As a factor capable of evaluating the aging properties, an aging index (AI) is generally used in the art.
Generally, if the steel has an AI of 30 MPa or more, aging can occur within six months after maintaining the steel at room temperature, and causes severe defects upon press forming.
However, when the solute carbon resides in the grain boundaries which are a relatively stable region, it is difficult to detect such solute carbon via a vibration test such as the internal friction test.
In the grain boundaries, the solute carbon exists in a relatively stable state. Therefore, although the solute carbon of the grain boundaries rarely affects aging at low temperatures such as AI test, it is activated in a baking condition of high temperatures, and affects the bake hardenability.
Hence, it can be concluded that the solute carbon in the crystal grains can affect both of the aging properties and the bake hardenability at the same time, whereas the solute carbon the grain boundaries affects only the bake hardenability.
In this regard, reports say that, since the grain boundaries are the relatively stable region, not all solute carbon in the grain boundaries affects the bake hardenability, but only about 50% of the solute carbon in the grain boundaries affects the bake hardenability.
Accordingly, it is possible to secure both of the bake hardenability and the aging resistance via an appropriate control of the locations of the solute carbon in the steel, that is, by controlling the solute carbon to reside in the grain boundaries rather than in the crystal grains as much as possible.
For this purpose, it is important to control the grain size with an added amount of carbon in the steel. This is because it is difficult to secure both of the bake hardenability and the aging resistance at the same time even with control on the location of the solute carbon, if the added amount of the carbon is excessively high or low in the steel.
FIG. 1 shows the BH value and the aging index (AI) in relation to variation in grain size, which is a result of investigation by the inventors of the present invention.
As shown in FIG. 1, it can be appreciated that, as ASTM No. of the crystal grains increases, that is, as the grains become finer and finer, the decrease of AI becomes severe compared with that of the BH value, which causes a gradual increase in value (BH−AI) obtained by subtracting the AI from the BH value, and thus the aging resistance is excellent.
On the basis of the result shown in FIG. 1, the inventors of the invention try to decrease the grain size of an annealed sheet to a suitable level or less to make the solute carbon distributed as much as possible into the grain boundaries of the steel sheet.
According to the result of investigation, the inventors of the present invention have found that the grain size is desirably controlled to ASTM No. 9 or more to maximize the aging resistance while minimizing deterioration of the bake hardenability.
Meanwhile, even when a great amount of solute carbon is distributed into the grain boundaries, it is necessary to strictly control a total amount of carbon in the steel. This is because an excessive amount of carbon in the steel causes an amount of solute carbon in the crystal grains to increase in proportion to the total added amount of carbon irrespective of fine crystal grains, deteriorating the bake hardenability.
According to the present invention, the total amount of carbon is set to 16˜25 ppm to satisfy the above conditions.
However, when Nb is added to the steel, it combines with C to precipitate a carbide such as NbC, thereby reducing the amount of solute carbon in the steel.
Thus, for Nb-added steel, a precipitation ratio of Nb/C is determined depending on the content of Nb and C in the steel. Furthermore, while some of the solute carbon in the steel combines with Nb to precipitate NbC, remaining carbon exist in the solid solution state, and affects the bake hardenability and the aging resistance.
From the above result, it can be understood that control of the solute carbon in the steel is more important than control of the added amount of Nb or the carbon content.
Meanwhile, in order to satisfy the above requirements while satisfying the aging resistance, it is important to make the solute carbon reside in the grain boundaries rather than in the crystal grains.
With investigation about influence of the solute carbon on the steel to satisfy both of the bake hardenability and the aging resistance at the same time under the above conditions, the inventors of the present invention have obtained a result as shown in FIG. 2 when the steel has a very fine grain size of ASTM No. 9 or more.
According to the result of investigating the bake hardenability related to variation in amount of the solute carbon in Nb-added ultra low carbon steel which has fine crystal grains as shown in FIG. 2, it has been found that an amount of solute carbon in the grain boundaries is about 3˜7 ppm, which satisfies a bake hardening degree of 30˜50 MPa set in consideration of the aging resistance.
Furthermore, it has also been found that a total amount of solute carbon is about 8˜15 ppm, wherein the total amount of solute carbon is obtained by excluding the NbC precipitates in consideration of the added amount of Nb and the carbon content.
With the above results, it is possible to obtain a condition which can satisfy both the bake hardenability and the aging resistance at the same time, that is, Equation 1:
C*[amount of solute carbon in grain boundaries (GB-C)+amount of solute carbon in crystal grains (G-C)]=Total C (ppm)−C in NbC=8˜15 ppm (1)
where GB-C (that is, the amount of solute carbon in the grain boundaries) is 5˜10 ppm, and G-C (that is, the amount of solute carbon in the crystal grains) is 3˜7 ppm.
In Equation 1, the term “C in NbC” means an amount of carbon precipitated in the form of the NbC precipitate.
It is possible to secure the bake hardenability and the aging resistance at the same time by controlling the locations of the solute carbon in the steel such that the total amount of solute carbon in the steel is about 8˜15 ppm, GB-C (that is, the amount of solute carbon in the grain boundaries) is 5˜10 ppm, and G-C (that is, the amount of solute carbon in the crystal grains) is 3˜7 ppm, as shown in Equation 1.
According to the present invention, an effect of AlN precipitates through addition of Al is also considered as well as the addition of Nb in order to more stably secure the bake hardenability and the aging resistance.
In Ti-added steel, since most of nitrogen is precipitated to coarse TiN at high temperatures of 1,300° C. or more, nitrogen has an insignificant influence on a solid solution effect or a grain refining effect in the steel.
However, if Ti is added at a very small amount of 30 ppm or less to the steel, there occurs AlN precipitation by Sol. Al.
AlN precipitates have an effect of removing the solute nitrogen in the steel.
According to results of various investigations with the bake hardenable steel of the present invention, since the carbon content of the inventive steel is restricted narrowly in the range of 16˜25 ppm, the bake hardenable steel of the present invention has the bake hardenability and the aging resistance in a narrow range.
Since customers demand the bake hardenable steel to have a higher BH value and aging resistance of 6 months or more, there are needs of techniques which can improve the bake hardenability without reducing the aging resistance as much as possible.
In this point of view, Al is very effective.
Specifically, when Sol. Al is added at a typical amount of 0.02˜0.06% to the steel, it serves simply to fix the solute nitrogen. However, when Sol. Al is added at an amount of 0.08% or more, the AlN precipitates become very fine, and act as a kind of barrier which obstructs growth of grains upon recrystallization annealing, so that the grains of the steel become finer than the Nb-added steel to which Sol. Al is not added, thereby providing an effect of improving the bake hardenability without changing the Al. The following Equation 2 shows influence of Sol. Al added in the range of the present invention with respect to the improvement of the bake hardenability in a statistical manner, in which Sol. Al is in the range of 0.08˜0.12% to provide the effect by Al.
BH=50−(885×Ti)−(1589×Nb)+(62×Al) (2)
However, even when the carbon content and the added amounts of Sol. Al and Nb are controlled, a hot coiling temperature is very important for the Nb-added ultra low carbon steel.
Specifically, when attempting to improve the bake hardenability and the aging resistance through the grain refining effect by addition of Nb, if a coiling temperature is excessively high, the grains become coarse during hot rolling. As a result, the steel has a grain size of ASTM No. 9 or less upon the recrystallization annealing, so that the Al exceeds 30 MPa, which is the upper limit of the present invention.
Meanwhile, in terms of secondary work embrittlement, it can be considered that components of automobiles are generally formed to desired shapes through several repetitions of press forming by automobile manufacturers. In this regard, the secondary work embrittlement means that cracks occur during a process performed after primary press working.
When P resides in the grain boundaries of the steel, it weakens the bonding force between the grains so that the cracks propagate along the grains boundaries, causing fracture of the steel.
Basically, it is desirable that P is not added to the steel in order to prevent the secondary work embrittlement. However, P has merits in that it serves to increase the strength while suppressing reduction of elongation, and in that it is very low in price.
Accordingly, although it is considered that P is basically added for the strength of the steel, there are investigations to increase the strength of the steel through addition of other elements instead of P in order to prevent the secondary work embrittlement notwithstanding a slight increase in manufacturing costs.
From the results of the investigation, however, it is expected that P will be used as a strengthening element of the steel for the time being.
As a method of improving the secondary work embrittlement resistance in such P-added steel, there is an attempt to promote a site competition effect between boron and phosphorus or increase the bonding force between the grains boundaries by allowing solid solution elements to remain in the steel or by adding boron and the like. It has also been attempted to minimize boundary diffusion of P by lowering the coiling temperature to a predetermined temperature or less in the hot rolling process. However, these methods do not completely solve the problem of the secondary work embrittlement.
In this regard, the present invention suggests addition of Mo to improve the secondary work embrittlement resistance more stably. From the results of the investigation, since Mo improves the bonding force between the grain boundaries, it is very advantageous to improve the secondary work embrittlement resistance.
In addition, since Mo has an affinity to the solute carbon in the steel, it suppresses diffusion of the solute carbon into the dislocations when being maintained for a long period of time at room temperature, thereby providing an advantageous effect in terms of aging resistance.
The following Equation 3 shows an effect of improving the aging resistance by Mo in a statistical manner.
AI=44−(423×Ti)−(2119×Nb)−(125×Mo) (3)
The inventors of the present invention deduce an optimal composition for the steel by suitably using the characteristics of Mo without deterioration in properties of the steel possibly caused by excessive addition of Mo.
Furthermore, the inventors try to maximize the effect of improving the secondary work embrittlement resistance through addition of a suitable amount of B and selection of a suitable coiling temperature at the same time among various methods conventionally used to improve the secondary work embrittlement resistance.
The bake-hardenable steel according to the present invention will be described in detail hereinafter.
Carbon (C) is an element used for solid solution strengthening and bake hardening.
If the carbon content is less than 0.0016%, the tensile strength of the steel is lowered due to such a low content of carbon, and a sufficient bake hardenability cannot be obtained due to a low absolute amount of carbon in the steel even with an attempt to achieve the grain refining effect through addition of Nb.
Furthermore, the secondary work embrittlement resistance is significantly deteriorated since the site competition effect between solute carbon and P is removed.
If the carbon content exceeds 0.0025%, the aging resistance at room temperature cannot be obtained irrespective of the BH value due to excessive quantities of solute carbon above 3˜7 ppm of the present invention in the crystal grains of the steel, and thus stretcher strain occurs during press working, causing deterioration in formability and ductility of the steel.
Silicon (Si) is an element for increasing the strength of the steel. However, as the silicon content is increased, the ductility is significantly deteriorated. In addition, since silicon deteriorates galvannealing capability, it is advantageous to minimize the added amount of silicon.
According to the invention, in order to prevent deterioration of the properties including coating properties of the steel due to Si, the added amount of silicon is preferably 0.02% or less.
Manganese (Mn) is an element used for preventing hot embrittlement caused by formation of FeS by completely precipitating sulfur in the steel into MnS while refining the crystal grains without deteriorating the ductility, and for strengthening the steel. According to the present invention, if the Mn content is less than 0.2%, a suitable tensile strength cannot be obtained, whereas if the Mn content exceeds 1.2%, the formability is deteriorated with a rapid increase in strength due to solid solution strengthening. In particular, when manufacturing a galvannealing steel sheet using such steel, a great amount of oxides, such as MnO, and a number of coating defects, such as a stripe pattern, are formed on the surface of the steel sheet during annealing, thereby deteriorating coating adherence and other properties of the steel. Accordingly, the Mn content is preferably in the range of 0.2˜1.2%.
Phosphorus (P) is a substitutional alloying element which provides the highest solid solution strengthening effect, and serves to enhance the anisotropy while increasing the strength of the steel.
From the results of the investigation, P causes the crystal grains of a hot-rolled steel sheet to become finer, thereby promoting development of the (111) texture, which is advantageous to improve an average r-value, during subsequent annealing. In particular, it has been found that, due to the site competition effect between P and carbon in terms of influence on the bake hardenability, the greater the amount of P the higher the bake hardenability. However, increase of P content causes a problem of deteriorating the secondary work embrittlement resistance by weakening the bonding force between the grain boundaries.
However, if the P content is less than 0.05%, the secondary work embrittlement resistance can be improved due to such a low content of P in the grain boundaries, but it is difficult to sufficiently obtain the effect of improving the other properties of the steel through the grain refining effect of P. On the other hand, if the P content is above 0.11%, it occurs a rapid increase of the strength compared with an improved degree of the formability. In addition, with such a high content of P, there is increased possibility of the secondary work embrittlement caused by segregation of P in the grain boundaries. Thus, the P content is preferably in the range of 0.05˜0.11%.
Sulfur (S) is an element, which precipitates into the sulfide such as MnS and prevents the hot embrittlement caused by FeS. However, if the S content is excessive, some of S remaining after precipitation of MnS makes the grain boundaries brittle, possibly causing the hot embrittlement.
In addition, even when S is added at an amount to allow complete precipitation of MnS, if the added amount of S is excessive, it occurs deterioration in properties of the steel due to the excessive precipitates. Thus, S is preferably in the range of 0.01% or less.
Aluminum (Al) is an element which is generally added for deoxidization of the steel. However, according to the present invention, aluminum is used for an effect of improving the grain refining effect and the bake hardenability via precipitation of AlN.
In other words, according to the present invention, although the grain refining effect is generally obtained using NbC precipitates by addition of Nb, aluminum serves to further improve the grain refining effect by the AlN precipitates, thereby improving the bake hardenability without deteriorating the aging resistance.
According to Equation 2, increase of the Al content is advantageous in view of the bake hardenability.
Considering the other properties of the steel, however, it is necessary to have a suitable content of Al.
Al must be added at an amount of at least 0.08% or more in order to achieve advantageous effects of the present invention.
When the Al content is above 0.12%, oxide inclusions are increased during steel making process and causes degradation of surface quality along with deterioration of the formability. Furthermore, the excessive content of Al results in high manufacturing costs. Thus, the Al content is preferably in the range of 0.08˜0.12%.
Nitrogen (N) exists in the solid solution state before or after annealing, and deteriorates the formability of the steel. Furthermore, since nitrogen has a higher aging ability than other interstitial solid solution elements, it is necessary to fix nitrogen by use of Ti or Al.
As in the invention wherein a suitable amount of Nb is added to the steel along with a small amount of Ti, excessive addition of nitrogen causes generation of solute nitrogen in the steel.
Since nitrogen has a higher diffusion rate than carbon, when nitrogen exists as solute nitrogen in the steel, the aging resistance at room temperature is deteriorated significantly more than the case by solute carbon.
In addition, since the yield strength and the r-value of the steel are lowered due to the solute nitrogen, it is preferable to have a nitrogen content of 0.0025% or less.
Titanium (Ti) is added to the steel as a carbide and nitride forming element, and forms nitrides such as TiN, sulfides such as TiS or Ti₄C₂S₂, and carbides such as TiC, in the steel.
According to the invention, Ti is added at an amount of 0.003% or less so as to fix a small amount of nitrogen.
The reason of adding such a small amount of Ti is that, generally, when manufacturing the steel in practice, an ultra low amount of Ti is contained with other components to the steel for the purpose of satisfying the properties of the steel, and that, when it is simultaneously added to steel slabs several times for the purpose of continuous casting of the steel, Ti of a previously added steel slab can be transferred into a subsequently added steel slab of the present invention.
However, as in the present invention, when Nb is added as a major element for improving the aging resistance, Ti is not necessarily added to the steel, and, if any, Ti is added at an ultra low amount of 0.003% or less under consideration of practical manufacturing conditions irrespective of reduction in bake hardenability thereby.
Niobium (Nb) is a very important element together with Al and Mo in the present invention.
Generally, Nb is an intensive carbide and nitride former, and serves to control an amount of solute carbon in the steel by pinning carbon of the steel into NbC precipitates. In particular, since the NbC precipitates are very fine compared with other precipitates, it acts as an intensive barrier to impede the grain growth during recrystallization annealing.
That is, in the present invention, the grain refining effect of Nb is obtained by use of such effect of the NbC precipitates. However, the present invention is suggested in an attempt to improve the bake hardenability using the solute carbon, which is allowed to remain in the steel.
For this purpose, the Nb content is preferably in the range of 0.003˜0.011% under consideration of the carbon content of 16˜25 ppm according to the present invention to obtain both of the bake hardenability and the aging resistance by allowing an amount of solute carbon of about 3˜7 ppm to remain in the crystal grains of the steel while providing the grain refining effect by the NbC precipitates.
Molybdenum (Mo) is another very important element of the present invention.
Mo exists in the solid solution state in the steel, and serves to enhance the strength of the steel or to form an Mo-based carbide.
In particular, Mo serves to increase the bonding force of the grain boundaries while existing as a solid solution element in the steel, so that fracture of the grain boundaries due to phosphorus is prevented, that is, the secondary work embrittlement resistance is improved. In addition, since Mo has an affinity to carbon, it serves to suppress the diffusion of carbon, improving the aging resistance. Equation 3 represents the effect of improving the aging resistance by Mo in a quantitative manner. For this purpose, it is necessary to add a suitable amount of Mo.
If Mo is added at an amount of 0.01% or less, the above effects cannot be obtained.
Thus, considering the manufacturing costs and the effect obtainable through addition of Mo, the Mo content is preferably in the range of 0.01˜0.1%.
Boron (B) is an interstitial element residing in the steel. B is dissolved as a solid solution element in the grain boundaries, or combines with nitrogen to form the nitride such as BN.
Since B has a highly significant influence on the properties of the steel compared with an added amount, it is necessary to precisely control the amount added. That is, when even a small amount of B is added, B is segregated in the grain boundaries and improves the secondary work embrittlement resistance.
However, when B is added at a predetermined amount or more, the steel is significantly deteriorated in ductility along with increase of the strength, and thus it is necessary to add an appropriate amount of B.
According to the invention, considering these characteristics of B and capability of manufacturing the steel through addition of B, the content of B is preferably in the range of 0.0005˜0.0015%.
A method for manufacturing steel of the invention will now be described.
After manufacturing a steel slab having the composition as described above, the steel slab is reheated at a temperature of 1,200° C. or more, where austenite structure before hot rolling can be sufficiently homogenized. The reheated steel slab is then subjected to hot-rolling with finish rolling at a finish rolling temperature of 900˜950° C., which is just above the Ar temperature, thereby providing a hot rolled steel sheet.
If the steel slab is reheated at a temperature less than 1,200° C., the structure of the steel is likely to become multi-phase structure, and cannot have homogeneous austenite crystal grains, causing deterioration in properties of the steel.
If the finish hot rolling temperature is less than 900° C., a top portion, a tail portion, and edges of a hot-rolled coil become single-phase regions, thereby increasing the anisotropy while deteriorating the formability of the steel. If the finish hot rolling temperature is above 950° C., crystal grains of the steel become remarkably coarsened, causing defects such as orange peel to be formed on the surface of the steel sheet after machining.
According to the present invention, it is important to suitably control a coiling temperature.
If the coiling temperature is less than 580° C., the steel sheet has refined crystal grains, which improve the aging resistance and the secondary work embrittlement resistance, but it suffers from an excessive increase in yield strength along with deterioration of the formability due to the excessively refined grain size.
Meanwhile, with an excessively high coiling temperature, the total amount of the solute carbon in the steel, the amount of the solute carbon in the crystal grains, and the amount of the solute carbon in the grain boundaries do not satisfy the above Equation 1. Thus, the coiling temperature is preferably controlled to 630° C. or less.
In this manner, according to the present invention, the coiling temperature is controlled to be in the range of 580˜630? as an exemplary means to make the amount of the solute carbon in the grains, and the amount of the solute carbon in the grain boundaries satisfy Equation 1.
After the hot rolled steel sheet is subjected to acid pickling in a typical manner, cold rolling of the hot rolled steel sheet is performed at a reduction rate of 75˜80%.
Such a high reduction rate of 75% or more is set for the purpose of enhancing the formability of the steel sheet, in particular, the r-value, together with the aging resistance through the grain refining effect.
If the reduction rate is above 80%, the steel sheet has a high grain refining effect. However, such an excessive reduction rate results in gradual decrease of the r-value, and excessive reduction in grain size, which makes the steel sheet hard.
After cold rolling, the steel sheet is subjected to continuous annealing by a typical method at a temperature of 770˜830° C.
Since the Nb-added steel has a higher recrystallization temperature than that of the Ti-added steel, annealing of the steel sheet is preferably performed at a temperature of 770° C. or more. That is, when the annealing is performed at a temperature less than 770° C., non-recrystallized crystal grains exist in the steel sheet, causing an increase of the yield strength while reducing the elongation and the r-value.
On the other hand, when annealing is performed at a temperature above 830° C., the formability can be enhanced. In this case, however, since the steel sheet has a grain size less than ASTM No. 9, which is the desired ASTM grain size in the present invention, the steel sheet has the AI of 30 MPa or less, and thus is deteriorated in aging resistance.
Then, for the purpose of improving the aging resistance at room temperature along with the suitable bake hardenability in the bake hardenable cold-rolled steel sheet produced by the above method, the cold-rolled steel sheet is subjected to temper rolling at a reduction ratio of 1.2˜1.5%, which is more or less higher than a typical temper rolling reduction ratio.
The reason of such a higher reduction ratio of 1.2 or more for the temper rolling is to prevent the aging resistance from being deteriorated due to the solute carbon in the steel.
However, if the reduction ratio of the temper rolling is set to an excessively high value above 1.5%, work hardening occurs and deteriorates the properties of the steel sheet irrespective of improved aging resistance. In particular, when manufacturing a galvannealed steel sheet using the bake hardenable cold-rolled steel sheet of the present invention, excessive temper rolling results in deterioration of the coating adherence, thereby causing separation of a coated layer. Thus, the temper rolling is preferably performed at the reduction ratio of 1.2˜1.5%, which is a suitable condition to solve the above problems.
The invention will be described in detail with reference to examples.

EXAMPLES

After hot rolling steel slabs having compositions as shown in Table 1 at a finish rolling delivery-side temperature of 900˜910° C. to form hot rolled steel sheets, the hot rolled steel sheets were coiled at a coiling temperature of 610˜630° C. Then, the hot rolled steel sheets were subjected to cold rolling at a reduction rate of 75˜78%, followed by continuous annealing at an annealing temperature of 800˜820° C. Then, the annealed cold-rolled steel sheets were subjected to hot-dipping at a temperature of 450˜470° C. and galvannealing at a temperature of 500˜530° C., followed by temper rolling at a temper rolling reduction ratio of about 1.5%. Next, BH value, aging index (AI), grain size, and ductility-brittleness transition temperature (DBTT) at a draw ratio of 2.0 for evaluation of secondary work embrittlement were measured with respect to final steel sheets. Results thereof are shown in Table 2. In addition, for Inventive Steel No. 6 of Table 1, an amount of solute carbon in crystal grains was measured, of which result is shown in FIG. 3.
In FIG. 3, the amount of solute carbon was measured using an internal friction tester (Horizontal type, 10 KHz).

	TABLE 1

	Composition (wt %)

C

Mn

P

S

Sol. Al

Ti

Nb

N

Mo

B

Remark

1	0.0023	0.30	0.060	0.0075	0.087	0	0.007	0.0022	0.034	0.0005	IS
2	0.0022	0.43	0.068	0.0062	0.108	0	0.009	0.0024	0.048	0.0005	IS
3	0.0023	0.55	0.062	0.0058	0.098	0.0015	0.0082	0.0021	0.051	0.0006	IS
4	0.0024	0.71	0.071	0.0068	0.118	0.0025	0.0073	0.0015	0.059	0.0005	IS
5	0.0021	1.01	0.10	0.0063	0.118	0	0.006	0.0021	0.052	0.0008	IS
6	0.0023	1.08	0.091	0.0057	0.104	0.001	0.011	0.0013	0.088	0.0009	IS
7	0.0054	0.42	0.059	0.0071	0.082	0	0.0105	0.0017	0.041	0.0007	CS
8	0.0011	0.75	0.072	0.0072	0.095	0.001	0.009	0.0021	0.069	0.0006	CS
9	0.0021	1.03	0.086	0.0089	0.023	0.001	0.035	0.0017	0.021	0.0006	CS
10	0.0022	0.79	0.062	0.0066	0.091	0.001	0.008	0.0022	0	0.0007	CS
11	0.0023	1.08	0.071	0.0078	0.098	0.035	0.022	0.0023	0.031	0	CS
12	0.0019	0.72	0.085	0.0069	0.110	0.002	0.011	0.010	0.052	0.0006	CS

IS: Inventive steel
CS: Comparative steel

TABLE 2

	Coiling	Annealing			Solute
	Temper-	Temper-			carbon in	Grain size
Steel	ature	ature	BH	AI	grains	(ASTM	DBTT
No.	(° C.)	(° C.)	(MPa)	(MPa)	(ppm)	No.)	(° C.)	Remark

1	620	805	44.3	24.9	5.7	10.1	−60	IS
2	620	810	42.4	18.9	3.9	9.7	−40	IS
3	620	815	41.7	19.6	3.5	9.9	−50	IS
4	610	800	43.5	20.1	5.2	10.4	−50	IS
5	620	820	47.8	24.8	6.6	10.8	−50	IS
6	620	800	38.1	9.3	3.1	9.9	−40	IS
7	620	810	68.0	61.5	10.5	11.7	−60	CS
8	620	810	0	0	0.1	8.0	20	CS
9	630	800	0	0	0	9.1	10	CS
10	720	810	42.1	26.6	4.1	10.9	20	CS
11	630	820	0	0	0	9.8	20	CS
12	620	810	83.0	81.1	11.3 (Solute	11.4	−30	CS
					carbon)

IS: Inventive steel
CS: Comparative steel

As shown in Table 2, it can be appreciated that Inventive steel Nos. 1 to 6 have an amount of solute carbon of 3.1˜6.6 ppm in crystal grains, and thus satisfy the condition for the amount of solute carbon in the grains according to the present invention, which is in the range of 3˜7 ppm.
As shown in FIG. 3, it can be appreciated that the inventive steels do not have solute nitrogen in the steel and has an amount of solute carbon of 3.1 ppm or more in the grains.
This is because nitrogen combines with Al added at a high content to the steel to form AlN precipitates, which contributes to refinement of the grains, and because remaining solute carbon, which is not used to form NbC precipitates, exists in the grains so as to be detected as above.
It is considered that the solute carbon in the crystal grains influences the bake hardenability.
As shown in Table 2, it can be appreciated that the Inventive steels Nos. 1 to 6 have a grain size of ASTM No. 9.8˜11.5 (average grain size of 6.7˜12.0 μm), and thus satisfy the condition for the grain size of the present invention, which is ASTM No. 9 or more.
As shown in Table 2, the Inventive steels Nos. 1 to 6 have fine crystal grains. That is, since the Inventive steel Nos. 1 to 6 have higher Al contents than a typical Al content, fine AlN precipitates are actively formed in the steel and obstruct the grain growth along with the NbC precipitates upon recrystallization annealing, which results in such fine crystal grains.
Thus, due to such a grain refining effect and an appropriate control of the solute carbon in the steel, the inventive steels have a bake hardening degree of 38.1˜47.9 MPa, and an AI of 9.3˜28.3 MPa, which is used for indicating the aging resistance at room temperature. Accordingly, it can be appreciated that balance between the bake hardenability and the aging resistance is excellent.
Meanwhile, the Inventive steels Nos. 1 to 6 have a relatively low AI in comparison to a relatively high bake hardening degree. It is considered that this phenomenon is based on a retarding effect of the solute carbon in the steel through addition of Mo in addition to the grain refining effect by the AlN precipitates.
Additionally, for the secondary work embrittlement, it can be seen that the DBTT at the draw ratio of 2.0 is in the range of −40˜−60° C.
Comparative steel No. 7 has a carbon content of 0.0054%, which is higher than the carbon content of the present invention in the range of 0.0016˜0.0025%, but satisfies the conditions for the high coiling temperature and the annealing temperature of the present invention. In addition, the Comparative steel No. 7 has a very fine grain size of ASTM No. 11.7 that satisfies the condition for the grain size of the present invention, and is excellent in view of DBTT and BH value. However, due to a high content of solute carbon in the steel, the Comparative steel No. 7 has an AI of 30 MPa or more, which indicates a significantly low aging resistance.
Unlike the Comparative steel No. 7, Comparative steel No. 8 has a significantly low carbon content of 0.0011%, which forms the NbC precipitates without remaining in a solid solution state in the steel. Thus, it can be found that the Comparative steel No. 8 does not exhibit the bake hardenability at all, and has coarsened grains and a low DBTT due to the low carbon content.
Comparative steel No. 9 has a soluble Al content of 0.023%, which is lower than the Al content of the present invention in the range of 0.08˜0.12%, and an Nb content of 0.035%, which is higher than the Nb content of the invention.
Thus, for the Comparative steel No. 9, the grain refining effect and the BH value increasing effect by the AlN precipitates were not achieved. Furthermore, since the high Nb content causes all carbon in the steel to be precipitated into NbC, the Comparative steel No. 9 hardly exhibits the bake hardenability and has a low DBTT due to a low site competition effect between C and P caused by reduction in amount of solute carbon in the steel.
Comparative steel No. 10 does not comprise Mo, and thus cannot be expected to have improved secondary work embrittlement resistance by Mo.
Furthermore, in view of manufacturing conditions, the Comparative steel No. 10 was coiled at a temperature of 720° C., which is higher than the coiling temperature of the present invention, causing a high possibility of activating movement of P.
Thus, due to non-addition of Mo and the high coiling temperature, the Comparative steel No. 10 is significantly deteriorated in DBTT irrespective of excellent bake hardenability and high AI value.
Comparative steel No. 11 has a very high Ti content of 0.035%, which causes no solute carbon to remain in the steel. Thus, the BH value and the AI of Comparative steel No. 11 are zero.
In addition, since the Comparative steel No. 11 does not comprise B at all, the secondary work embrittlement by P added at an amount of 0.071% cannot be prevented.
In this regard, it is considered that, even though it satisfies the condition of the present invention in terms of the Mo content, the Comparative steel No. 11 does not comprises the solute carbon, which is advantageous to improve the secondary work embrittlement resistance, and cannot be expected to have an increased bonding force of the grain boundaries by addition of B, so that the DBTT of the Comparative steel No. 11 is deteriorated.
Comparative steel No. 12 sufficiently satisfies the requirement for the composition suggested by the present invention except for a high content of nitrogen.
Nitrogen is the element which causes a detrimental problem in terms of bake hardenability and aging resistance.
For the Comparative steel No. 12, the excessively high content of nitrogen results in a high amount of solute nitrogen of 11.3 ppm in the grains and high AI value, which is detrimental to the bake hardenability and the aging resistance.

INDUSTRIAL APPLICABILITY

As apparent from the above description, according to the present invention, the bake-hardenable cold-rolled steel sheet and the galvannealed steel sheet using the same have excellent bake hardenability, aging resistance, and secondary work embrittlement resistance, as well as high strength.
It should be understood that the embodiments and the accompanying drawings as described above have been described for illustrative purposes and the present invention is limited by the following claims. Further, those skilled in the art will appreciate that various modifications, additions and substitutions are allowed without departing from the scope and spirit of the invention as set forth in the accompanying claims.

Claims

1. A bake-hardenable cold-rolled steel sheet with high strength and superior aging resistance, comprising, by weight %: C: 0.0016˜0.0025%; Si: 0.02% or less; Mn: 0.2˜1.2%; P: 0.05˜0.11%; S: 0.01% or less; Sol. Al: 0.08˜0.12%; N: 0.0025% or less; Ti: 0˜0.003%; Nb: 0.003˜0.011%; Mo: 0.01˜0.1%; B: 0.0005˜0.0015%; and the balance of Fe and other unavoidable impurities, wherein the steel sheet satisfies Equation 1:

C*[amount of solute carbon in grain boundaries (GB-C)+amount of solute carbon in crystal grains (G-C)]=Total C (ppm)−C in NbC=8˜15 ppm (1)

[in Equation 1, GB-C (that is, the amount of solute carbon in the grain boundaries) is 5˜10 ppm, and G-C (that is, the amount of solute carbon in the crystal grains) is 3˜7 ppm], and

wherein the steel sheet has a grain size of ASTM No. 9 or more, the bake hardenability value (BH) of 30 MPa or more, aging index (AI) of 30 MPa or less, and a DBTT of −30° C. or less at a draw ratio of 2.0, the bake hardenability value BH) and the aging index (AI) satisfying Equations 2 and 3, respectively:

BH=50−(885×Ti)−(1589×Nb)+(62×Al) (2)

AI=44−(423×Ti)−(2119×Nb)−(125×Mo) (3).

2. A galvannealed bake hardenable steel sheet with high strength and superior aging resistance, comprising, by weight %: C: 0.0016˜0.0025%; Si: 0.02% or less; Mn: 0.2˜1.2%; P: 0.05˜0.11%; S: 0.01% or less; Sol. Al: 0.08˜0.12%; N: 0.0025% or less; Ti: 0˜0.003%; Nb: 0.003˜0.011%; Mo: 0.01˜0.1%; B: 0.0005˜0.0015%; and the balance of Fe and other unavoidable impurities, wherein the steel sheet satisfies Equation 1:

wherein the steel sheet has a grain size of ASTM No. 9 or more, the bake hardenability value (BH) of 30 MPa or more, an aging index (AI) of 30 MPa or less, and a DBTT of −30° C. or less at a draw ratio of 2.0, the bake hardenability value (BH) and the aging index (AI) satisfying Equations 2 and 3, respectively:

BH=50−(885×Ti)−(1589×Nb)+(62×Al) (2)

AI=44−(423×Ti)−(2119×Nb)−(125×Mo) (3)

3. A method for manufacturing a bake-hardenable cold-rolled steel sheet with high strength and superior aging resistance, comprising: performing homogenization heat treatment for an Al-killed steel slab at 1200° C. or more, the steel slab comprising, by weight %: C: 0.0016˜0.0025%, Si: 0.02% or less, Mn: 0.2˜1.2%, P: 0.05˜0.11%, S: 0.01% or less, Sol. Al: 0.08-0.12%, N: 0.0025% or less, Ti: 0˜0.003%, Nb: 0.003˜0.011%, Mo: 0.01˜0.1%, B: 0.0005˜0.0015%, and the balance of Fe and other unavoidable impurities;

hot rolling the steel slab with finish rolling at a finish rolling temperature of 900˜950° C. to form a hot-rolled steel sheet, followed by coiling the hot-rolled steel sheet at a temperature of 580˜630° C.;

cold rolling the hot-rolled steel sheet at a reduction ratio of 75˜80%;

continuously annealing the cold-rolled steel sheet at a temperature of 770˜830° C.; and

temper rolling the anneal steel sheet at a reduction ratio of 1.2˜1.5%.