EP4394076A1

EP4394076A1 - Cold rolled steel sheet having excellent weldability, strength, and formability, and method for manufacturing same

Info

Publication number: EP4394076A1
Application number: EP22861736.1A
Authority: EP
Inventors: Young-Roc Im; Sang-Ho Uhm; Young-Ha Kim
Original assignee: Posco Co Ltd
Current assignee: Posco Holdings Inc
Priority date: 2021-08-26
Filing date: 2022-08-25
Publication date: 2024-07-03
Also published as: KR20230033043A; WO2023027528A1; EP4394076A4; JP2024534112A; CN117897514A; US20240352551A1

Abstract

The present invention relates to: a cold rolled steel sheet having excellent weldability, strength, and formability; and a method for manufacturing same.

Description

Technical Field

The present disclosure relates to a cold-rolled steel sheet having excellent weldability, strength, and formability, and a method of manufacturing the same.

Background Art

Recently, in order to reduce the weight of a vehicle and enhance safety thereof, manufacturing technology for a high-strength steel sheet has been promoted, and, in particular, demand for a high-strength steel material having a tensile strength of 980 MPa or higher is increasing. However, it is a common phenomenon that ductility and formability deteriorate when simply improving strength. Accordingly, in order to overcome this, a high-strength steel sheet having formability for cold forming may improve fuel efficiency through a reduction in weight, improve component manufacturing/forming productivity, and is highly useful in terms of ensuring safety of a final component.
In order to improve formability of a steel material, a method of utilizing a transformation induced plasticity (TRIP) phenomenon by introducing retained austenite as a method of increasing elongation is widely used. However, in a TRIP steel sheet, the addition of Si and Al is required to introduce retained austenite, which may cause liquid metal embrittlement (LME) to occur during spot welding of the steel sheet. Accordingly, use of a plated steel sheet and use of a cold-rolled steel sheet welded with a plating material are restricted.
(Patent Document 1) Korean Patent Publication No. 2017-7015003

Summary of Invention

Technical Problem

An aspect of the present disclosure is to provide a cold-rolled steel sheet having excellent weldability, strength, and formability, and a method of manufacturing the same.
An object of the present disclosure is not limited to those described above. Any person skilled in the art to which the present disclosure pertains will have no difficulty in understanding the additional problems of the present disclosure from the overall content of the present disclosure specification.

Solution to Problem

An aspect of the present disclosure is to provide a cold-rolled steel sheet including, by weight%, C: 0.10 to 0.16%, Si: 0.3 to 0.8%, Al: 0.01 to 0.5%, Mn: 2.0 to 3.0%, Cr: 0.001 to 0.5%, Mo: 0.001 to 0.5%, B: 0.0001 to 0.001%, Nb: 0.001 to 0.05%, Ti: 0.001 to 0.05%, P: 0.04% or less (excluding 0%), S: 0.01% or less (excluding 0%), N: 0.01% or less (excluding 0%), a remainder of Fe, and other unavoidable impurities,

as a microstructure comprises, by area%, ferrite: 10% or less (excluding 0%), retained austenite: more than 1% to 5% or less, martensite: 25% or more but less than 50%, and bainite: 35% or more but less than 70%, and
an average size of martensite-austenite (MA) in the bainite is 0.35 to 0.55 µm.

In addition, although not particularly limited, in an embodiment of the present disclosure, in order to have high local formability, alloy components of the cold-rolled steel sheet may be controlled such that a value defined by the following Relationship 1 satisfies 70 or more:

234×[C] - 29×[Si] - 128×[Al] + 29×[Mn] + 10×[Cr] - 17×[Mo] - 37×[Nb] - 49×[Ti] + 100×[B]
(In Relationship 1 above, [C], [Si], [Al], [Mn], [Cr], [Mo], [Nb], [Ti], and [B] represent weight percentages of the elements in parentheses, respectively.)
In addition, although not particularly limited, in an embodiment of the present disclosure, in order to have high local formability, alloy components of the cold-rolled steel sheet may be controlled such that a value defined by the following Relationship 2 satisfies 270 or more and 330 or less:

$270 \times [C] + 90 \times [Mn] + 70 \times [Cr] + 80 \times [Mo]$
(In Relationship 2 above, [C], [Mn], [Cr], and [Mo] represent weight percentages of the elements in parentheses, respectively.)
In addition, although not particularly limited, in an embodiment of the present disclosure, in order to have resistance to welding LME, a relationship in amounts of C, Si, and Al of the cold-rolled steel sheet may be adjusted such that the value defined by the following Relationship 3 satisfies 1.8 or less:

$5 \times [C] + [Si] + 0.5 \times [Al]$
(In Relationship 3 above, [C], [Si], and [Al] represent weight percentages of the elements in parentheses, respectively.)
Another aspect of the present disclosure is to provide a method for manufacturing a cold-rolled steel sheet
including:

heating a steel slab including, by weight%, C: 0.10 to 0.16%, Si: 0.3 to 0.8%, Al: 0.01 to 0.5%, Mn: 2.0 to 3.0%, Cr: 0.001 to 0.5%, Mo: 0.001 to 0.5%, B: 0.0001 to 0.001%, Nb: 0.001 to 0.05%, Ti: 0.001 to 0.05%, P: 0.04% or less (excluding 0%), S: 0.01% or less (excluding 0%), N: 0.01% or less (excluding 0%), a remainder of Fe, and other unavoidable impurities;
Hot-rolling the heated slab in a finish hot rolling temperature range of 830 to 980°C to obtain a hot-rolled steel sheet;
coiling the hot-rolled steel sheet in a temperature range of 450 to 700°C;
cold-rolling the coiled hot-rolled steel sheet;
continuously annealing the cold-rolled steel sheet in a temperature range of 790 to 830°C;
primary cooling the continuously annealed steel sheet at an average cooling rate of less than 10°C/s to a primary cooling end temperature of 450 to 600°C;
secondary cooling the primary cooled steel sheet at an average cooling rate of 10°C/s or more to a secondary cooling end temperature of 250 to 350°C; and
reheating the secondary cooled steel sheet to a temperature in a range of 350 to 480°C,
wherein the method for manufacturing the cold-rolled steel sheet satisfies the following Relationship 4:

$V 1 / V 2 \times t > 0.5$

(In Relationship 4 above, V1 represents an average cooling rate during the primary cooling, V2 represents an average cooling rate during the secondary cooling, and t represents a thickness of the cold-rolled steel sheet.)
Although not particularly limited, in an embodiment of the present disclosure, in order to have high local formability, alloy components of the steel slab may be controlled such that a value defined by the following Relationship 1 satisfies 70 or more: 234×[C] - 29×[Si] - 128×[Al] + 29×[Mn] + 10×[Cr] - 17×[Mo] - 37×[Nb] - 49×[Ti] + 100×[B]
(In Relationship 1 above, [C], [Si], [Al], [Mn], [Cr], [Mo], [Nb], [Ti], and [B] represent weight percentages of the elements in parentheses, respectively.)
In addition, although not particularly limited, in an embodiment of the present disclosure, in order to have high local formability, the alloy component of the steel slab may be controlled such that the value defined by the following Relationship 2 satisfies 270 or more and 330 or less:

$270 \times [C] + 90 \times [Mn] + 70 \times [Cr] + 80 \times [Mo]$

(In Relationship 2 above, [C], [Mn], [Cr], and [Mo] represent weight percentages of the elements in parentheses, respectively.)
In addition, although not particularly limited, in an embodiment of the present disclosure, in order to have resistance to welding LME, a relationship in amounts of C, Si, and Al of the cold-rolled steel sheet may be adjusted such that the value defined by the following Relationship 3 satisfies 1.8 or less:

$5 \times [C] + [Si] + 0.5 \times [Al]$

(In Relationship 3 above, [C], [Si], and [Al] represent weight percentages of the elements in parentheses, respectively.)
In addition, an embodiment of the present disclosure may further include, as necessary, plating the reheated steel sheet in a zinc plating bath in a temperature range of 450 to 470°C, after the reheating.
In addition, an embodiment of the present disclosure may further include, as necessary, alloying and heat-treating the plated steel sheet in a temperature range of 470 to 550°C.

Advantageous Effects of Invention

According to an aspect of the present disclosure, a cold-rolled steel sheet having excellent weldability, strength, and formability, and a method of manufacturing the same may be provided.
Various advantages and effects of the present disclosure are not limited to the above-described content, and may be more easily understood through description of specific embodiments of the present disclosure.

Brief Description of Drawings

FIG. 1 is a photograph of a cross-section of a cold-rolled steel sheet obtained in Example 1 of the present application in a thickness direction, at 5,000X magnification with a scanning electron microscope (SEM), to observe martensite-austenite (MA) present in bainite.

Best Mode for Invention

Hereinafter, preferred embodiments of the present disclosure will be described. However, embodiments of the present disclosure may be modified into various other forms, and the scope of the present disclosure is not limited to embodiments described below. Additionally, embodiments of the present disclosure may be provided to more fully illustrate the present disclosure to those with average knowledge in the relevant technical field.
Meanwhile, terms used in this specification are for describing specific embodiments, and are not intended to limit the present disclosure. For example, as used herein, singular forms include plural forms unless the relevant definition clearly indicates the contrary. Additionally, the meaning of "include" or "comprise" used in the specification may be to specify a configuration and not to exclude the presence or addition of another configuration.
In the prior art, a trip steel sheet incorporating retained austenite was developed to secure a high strength of 980 MPa or higher in tensile strength and improve formability, but there was a problem that addition of Si and Al was required, which caused liquid metal embrittlement (LME) to occur during spot welding.
Accordingly, research is being conducted to obtain elongation as high as possible while limiting amounts of C, Si, and Al added to a steel sheet, and to compensate for insufficient formability by improving local formability. In order to improve the local formability, it may be effective to reduce a hardness variation between microstructures constituting a steel material. As an industrial local formability evaluation, a widely performed test may be hole expansion ratio (HER) measurement. The hole expansion ratio (HER) may be obtained by fixing a specimen having a hole having a diameter of 10 mm, perforated with a punch, to a die, pushing and expanding the hole upwardly with a conical punch, and measuring a diameter of the expanded hole at a time at which a crack entirely penetrating a thickness of the specimen occurs, as a value such as the following Relationship A. Detailed hole expansion ratio evaluation criteria may be based on ISO 16630 regulations.

$λ (HER) = (df - do) / do$

(In the above Relationship A, do represents a diameter of an initial hole, and df represents a diameter of a hole at a time of breakage in thickness.)
Accordingly, the present inventors conducted intensive research to provide a cold-rolled steel sheet capable of suppressing LME problems while securing a high strength of 980 MPa or higher in tensile strength, and securing excellent formability and an excellent hole expansion ratio, and as a result, it is discovered that the problems could be solved by precisely controlling an alloy composition and manufacturing conditions, and came to complete the present disclosure.
Hereinafter, a cold-rolled steel sheet having excellent weldability, strength, and formability, according to an embodiment of the present disclosure, and a method of manufacturing the same will be described.
First, an alloy composition of the cold-rolled steel sheet according to an embodiment of the present disclosure will be described. An amount of the alloy composition mentioned below refers to wt%.

C: 0.10 to 0.16%

Carbon (C) may be an element securing strength of a steel material through solid solution strengthening and precipitation strengthening. When the C content is less than 0.10%, it may be difficult to secure a tensile strength (TS) of about 980 MPa. When the C content exceeds 0.16%, arc weldability and laser weldability may deteriorate, and risk of LME cracking may increase. Therefore, the C content may range from 0.10% or more and 0.16% or less. Meanwhile, it is more preferable that a lower limit of the C content is 0.137%. Additionally, it is more preferable that an upper limit of the C content is 0.151%.

Si: 0.3 to 0.8%

Silicon (Si) may be a key element in transformation induced plasticity (TRIP) steel that increases a retained austenite fraction and elongation by inhibiting precipitation of cementite. When the Si content is less than 0.3%, almost no retained austenite may remain and the elongation may become too low. When the Si content exceeds 0.8%, it is impossible to prevent deterioration of properties of a welded zone due to formation of an LME crack, and surface properties and plating properties of the steel may deteriorate. Therefore, the Si content may range 0.3 to 0.8%. Meanwhile, it is more preferable that a lower limit of the Si content is 0.49%. Additionally, it is more preferable that an upper limit of the Si content is 0.70%.

Al: 0.01 to 0.50

Aluminum (Al) may be not only an element included for deoxidation of a steel material, but also an element that is effective in stabilizing retained austenite by suppressing precipitation of cementite. When the Al content is less than 0.01%, deoxidation of the steel material may not be sufficiently achieved, and clarity of the steel material may be impaired. When the Al content exceeds 0.5%, castability of the steel material may be impaired. Therefore, the Al content may range 0.01 to 0.5%. Meanwhile, it is more preferable that a lower limit of the Al content is 0.027%. Additionally, it is more preferable that an upper limit of the Al content is 0.085%.

Mn: 2.0 to 3.0%

Manganese (Mn) may be an element added to ensure strength. When the Mn content is less than 2.0%, it may be difficult to secure strength. When the Mn content exceeds 3.0%, a bainite transformation rate may decrease, and too much fresh martensite may be formed, making it difficult to obtain a high hole expansion ratio. In addition, a band structure may be formed due to segregation of Mn, which impairs uniformity and formability of a material. Therefore, the Mn content may range 2.0 to 3.0%. A lower limit of the Mn content is more preferably 2.2%, and even more preferably 2.3%. An upper limit of the Mn content is more preferably 2.8%, and even more preferably 2.7%.

Cr: 0.001 to 0.50

Chromium (Cr) may be an element added to secure strength and hardenability. When Mn is added alone, a very large amount of Mn exceeding the Mn content range of the present disclosure should be added. This problem may be solved by adding 0.001% or more of Cr. When the Cr content exceeds 0.5%, local corrosiveness may deteriorate, and oxides may be formed on a surface thereof, impairing phosphate treatment properties. Therefore, the Cr content may range 0.001 to 0.5%. A lower limit of the Cr content is more preferably 0.002%, and an upper limit of the Cr content is more preferably 0.38%.

Mo: 0.001 to 0.50

Molybdenum (Mo) may be an element added to secure strength and hardenability. When Mn is added alone, a very large amount of Mn exceeding the Mn content range of the present disclosure should be added. This problem may be solved by adding 0.001% or more of Mo. When the Mo content exceeds 0.5%, phase transformation may be suppressed, making it difficult to introduce a bainite structure, and as an expensive element, economic feasibility of a steel sheet may deteriorate. Therefore, the Mo content may range 0.001 to 0.5%. Meanwhile, it is more preferable that a lower limit of the Mo content is 0.07%. Additionally, an upper limit of the Mo content is more preferably 0.3%, and most preferably 0.21%.

B: 0.0001 to 0.001%

Boron (B) may be an element added to secure hardenability. When Mn is added alone, a very large amount of Mn exceeding the Mn content range of the present disclosure should be added. This problem may be solved by adding 0.0001% or more of B. However, when the B content exceeds 0.0001%, B may be excessively accumulated on a surface, impairing plating adhesion of a plating material. Therefore, the B content may range 0.0001 to 0.001%. A lower limit of the B content is more preferably 0.00010%, and an upper limit of the B content is more preferably 0.0005%.

Nb: 0.001 to 0.05%

Niobium (Nb) may be an element added to secure strength of a steel sheet and refine a structure thereof. When less than 0.001% of Nb is added, it may be difficult to achieve strength improvement and structure refinement effects, and when the Nb content exceeds 0.05%, recrystallization may be delayed due to local grain fixation, thereby impairing uniformity of the structure. Therefore, the Nb content may range 0.001 to 0.05%. Meanwhile, it is more preferable that a lower limit of the Nb content is 0.015%. In addition, it is more preferable that an upper limit of the Nb content is 0.031%.

Ti: 0.001 to 0.05%

Titanium (Ti) may be an element added to secure strength of a steel sheet and refine a structure thereof. When less than 0.001% of Ti is added, it may be difficult to achieve strength improvement and structure refinement effects. When the Ti content exceeds 0.05%, castability may be impaired due to excessive formation of TiN, and recrystallization may be delayed due to local grain fixation, thereby impairing uniformity of the structure. Therefore, the Ti content may range 0.001 to 0.05%. A lower limit of the Ti content is more preferably 0.015%, or an upper limit of the Ti content is more preferably 0.03%.

P: 0.04% or less (excluding 0%)

Phosphorus (P) may exist as an impurity in steel, and it may be advantageous to control its content as low as possible. Therefore, a lower limit of the P content may exclude 0% (i.e., exceed 0%), taking into account cases where P is inevitably included. P may be sometimes added intentionally to increase strength of the steel. When P is added excessively, toughness of the steel may deteriorate. Therefore, to prevent this, the present disclosure may restrict an upper limit thereof to 0.04%. The lower limit of the P content is more preferably 0.002%, or the upper limit of the P content is more preferably 0.0173%.

S: 0.01% or less (excluding 0%)

Sulfur (S), like P, may exist as an impurity in steel, and it may be advantageous to control its content as low as possible. Therefore, a lower limit of the S content may exclude 0% (i.e. exceed 0%), taking into account cases where S is inevitably included. Since S deteriorates ductility and impact properties of the steel, an upper limit thereof may be restricted to have 0.01%. The lower limit of the S content is more preferably 0.0009%, or the upper limit of the S content is more preferably 0.0021%.

N: 0.01% or less (excluding 0%)

In the present disclosure, nitrogen (N) may be included in a steel material as an impurity, and it may be advantageous to control its content as low as possible. Therefore, a lower limit of the N content excludes 0% (i.e. exceeds 0%), taking into account cases where N is inevitably included. An upper limit of the N content may be restricted to have 0.01%. It is more preferable that the lower limit of the N content is 0.0005%. Additionally, the upper limit of the N content is more preferably 0.007%, even more preferably 0.006%, and most preferably 0.0052%.
In addition to the steel composition described above, the remainder may include Fe and inevitable impurities. The inevitable impurities may be unintentionally mixed in a normal steel manufacturing process, and may not thus be completely excluded, and any engineer in the normal steel manufacturing field may easily understand their meaning. In addition, the present disclosure may not completely exclude addition of compositions other than the steel compositions mentioned above.
According to an embodiment of the present disclosure, although not particularly limited, the cold-rolled steel sheet may further include optionally one or more selected from the group consisting of Cu: 0.1% or less (excluding 0%) and Ni: 0.1% or less (excluding 0%).

Cu: 0.1% or less (excluding 0%), Ni: 0.1% or less (excluding 0%)

Copper (Cu) and nickel (Ni) may be elements that increase strength of a steel material. The elements may be elements that increase strength and hardenability of the steel material. When the elements are added in excessive amounts, the elements may exceed a target strength grade. Since the elements may be expensive elements, from an economic standpoint, upper limits thereof may be limited to 0.1% or less, respectively. Since Cu and Ni act as solid solution strengthening elements, when adding one or more of Cu and Ni, a solid solution strengthening effect may be minimal when less than 0.03% is added. Therefore, it is preferable to add 0.03% or more of each.
According to an embodiment of the present disclosure, although not particularly limited, the cold-rolled steel sheet may optionally further include V: 0.05% or less (excluding 0%).

V: 0.05% or less (excluding 0%)

Vanadium (V) may increase strength of a steel material even with addition of a small amount thereof, but its effect on improving elongation may not be significant. Therefore, it is desirable to control its content to 0.05% or less. The V content is more preferably 0.04% or less, and even more preferably 0.03% or less.
A microstructure of a cold-rolled steel sheet according to an embodiment of the present disclosure may include, by area%, ferrite: 10% or less (excluding 0%), retained austenite: more than 1% to 5% or less, martensite: 25% or more but less than 50%, and bainite: 35% or more but less than 70%.
Although not particularly limited, according to an embodiment of the present disclosure, the cold-rolled steel sheet may be provided to secure excellent formability even at a tensile strength (TS) of 980 MPa or higher, and in particular, to obtain high local formability, the steel sheet should reduce a difference in hardness between microstructure phases. In the present disclosure, when the composition is controlled such that a value defined by the following Relationship 1 satisfies 70 or more while satisfying the above-described alloy composition under typical annealing heating conditions, it is confirmed that an austenite single phase is obtained and a ferrite fraction may be reduced to 10 area% or less. When the ferrite fraction exceeds 10 area%, there may be risks that yield strength is lowered and a hole expansion ratio deteriorates. More preferably, in terms of securing the above-mentioned high yield strength and an excellent hole expansion ratio, a lower limit of the ferrite fraction may be 2 area%, and an upper limit of the ferrite fraction may be 7 area%. 234×[C] - 29×[Si] - 128×[Al] + 29×[Mn] + 10×[Cr] - 17×[Mo] - 37×[Nb] - 49×[Ti] + 100×[B]
(In Relationship 1 above, [C], [Si], [Al], [Mn], [Cr], [Mo], [Nb], [Ti], and [B] represent weight percentages of the elements in parentheses, respectively.)
The value defined by Relationship 1 may be satisfied to have 70 or more, to avoid a soft ferrite phase. When a bainite phase, which may be soft, next to ferrite, is not sufficiently introduced, it may be difficult to secure ductility of a steel material. In terms of further improving the above-described effect, more preferably, a lower limit of the value defined by Relationship 1 may be 75.7, or an upper limit of the value defined by Relationship 1 may be 90.
Therefore, although not particularly limited, according to an embodiment of the present disclosure, as a condition for sufficiently introducing the bainite phase even under normal continuous plating annealing conditions, a value defined by the following Relationship 2 may be controlled to satisfy 270 or more and 330 or less:

$270 \times [C] + 90 \times [Mn] + 70 \times [Cr] + 80 \times [Mo]$

(In Relationship 2 above, [C], [Mn], [Cr], and [Mo] represent weight percentages of the elements in parentheses, respectively.)
A cold-rolled steel sheet according to the present disclosure may be mainly composed of martensite and bainite, and when a difference in hardness between main phases is large, local formability may deteriorate. Since a bainite structure usually has lower strength than a martensite structure, a method of improving strength of the bainite structure is required to reduce hardness deviation.
As such, the present inventors conducted intensive researches for improving properties by reducing the hardness deviation between the martensite and the bainite, and as a result, it was discovered that a difference in hardness between the two main phases could be dramatically reduced by controlling an average size of martensite-austenite (MA) present in the bainite in an appropriate range.
Specifically, in the cold-rolled steel sheet according to the present disclosure, the average size of martensite-austenite (MA) present in the bainite may be 0.35 to 0.55 um.
When the average size of the martensite-austenite (MA) present in the bainite is less than 0.35 um, a difference in hardness with the martensite phase may increase, as the strength of the bainite structure decreases, making it difficult to secure a high hole expansion ratio. When the average size of the martensite-austenite (MA) present in the bainite exceeds 0.55 um, effects of hard martensite-austenite may increase to occur problems of causing brittleness and lowering a hole expansion ratio.
According to an embodiment of the present disclosure, more preferably, in terms of improving the above-mentioned effects, a lower limit of the average size of the martensite-austenite (MA) present in the bainite may be 0.4 um, or an upper limit of the average size of martensite-austenite (MA) present in the bainite may be 0.5 µm.
In this case, in this specification, the average size of martensite-austenite (MA) present in the above-described bainite may represent a measurement of an average size of martensite-austenite (MA) completely contained in the entire bainite, based on a cross-section of the steel sheet cut in a thickness direction. In addition, the average size of the martensite-austenite (MA) means an average value of maximum lengths penetrating an internal space of the martensite-austenite (MA).
In addition, although not particularly limited, according to an embodiment of the present disclosure, the value defined by Relationship 2 may be controlled to satisfy 270 or more and 330 or less while satisfying the above-described alloy composition. For this reason, even under normal annealing conditions, bainite, which has an MA phase (martensite-austenite aggregate) as a second phase, may be formed in 35 area% or more to less than 70 area%, making it possible to further improve hole expansion ratio. It is believed that the reason why strength of the bainite phase is secured close to that of martensite is because it contains a relatively hard second phase MA phase internally through carbon distribution.
When the value defined by Relationship 2 above exceeds 330, it may be difficult to secure a sufficient bainite fraction of 35 area% or more, resulting in excessively high strength, and poor elongation and HER values. Conversely, when the value defined by Relationship 2 is less than 270, ductility may be sufficient, but the steel sheet becomes too soft, making it difficult to obtain a tensile strength of 980 MPa or more. In this case, in terms of further improving the above-mentioned effect, more preferably, a lower limit of the value defined by Relationship 2 is 286, or an upper limit of the value defined by Relationship 2 is 311.
In the present disclosure, the retained austenite may be a structure that increases elongation of the steel material through a TRIP effect. As a fraction thereof increases, elongation thereof may increase, and a fraction of the retained austenite may exceed 1 area% to obtain the required level of elongation. To obtain austenite exceeding 5 area%, a large amount of C and Si should be added, and in this case, spot welding LME resistance may deteriorate. Therefore, in the present disclosure, the fraction of the retained austenite may be controlled to 5 area% or less. In this case, in terms of further improving the above-mentioned effect, more preferably, a lower limit of the fraction of the retained austenite is 2 area%, or an upper limit of the fraction of the retained austenite is 4 area%.
Additionally, in the present disclosure, a fraction of the martensite may be 25 area% or more and less than 50 area%. When the fraction of the martensite is less than 25 area%, a problem may arise in which overall tensile strength of a steel material may be insufficient, and when the fraction of the martensite is 50 area% or more, strength may increase excessively, causing a problem of low hole expansion ratio. In this case, more preferably in terms of further improving the above-mentioned effect, a lower limit of the fraction of the martensite may be 29 area%, or an upper limit of the fraction of the martensite may be 49 area%.
Additionally, in the present disclosure, a fraction of the bainite may be 35 area% or more and less than 70 area%. When the fraction of the bainite is less than 35 area%, the fraction of martensite or ferrite may be relatively high, which may cause a problem of low hole expansion ratio, and when the fraction of the bainite is 70 area% or more, the fraction of the martensite may be lowered, which may cause a problem of low insufficient strength. In this case, more preferably in terms of further improving the above-mentioned effect, a lower limit of the fraction of the bainite is 45 area%, or an upper limit of the fraction of the bainite is 63%.
According to an aspect of the present disclosure, the cold-rolled steel sheet may further include other phases in addition to the microstructure described above. In this case, the other phases may include martensite-austenite (MA) or the like, and for example, martensite-austenite (MA) or the like present in bainite may exist.
According to an embodiment of the present disclosure, when a large amount of an alloy element such as C, Si, Al, or the like is added, spot weldability may deteriorate, and in particular, when spot welding is performed on a galvanized steel sheet, liquid metal embrittlement (LME) may be caused. In general, spot welding of a steel material may be performed on or below a minimum current at which expulsion occurs, and the minimum current at which expulsion occurs may be referred to as a condition that provides the highest amount of heat input when performing actual spot welding. When the LME resistance is high, LME may not occur even at welding current values above this minimum current value for scattering generation, and in this case, an AE value, defined as a difference between a minimum current value for generating LME minus a minimum current value for generating scattering, may be a positive value. For example, welding may be performed on or below the minimum current value for generating scattering during actual spot welding, and at this time, when LME does not occur, it may be determined that the AE value is 0 or more. Meanwhile, the AE value may have a unit of kA.
Although not particularly limited, according to an embodiment of the present disclosure, as a result of evaluating the spot weldability of a steel sheet having various alloy components and having a 980 MPa tensile strength, an alloy component condition having excellent LME resistance, i.e., the AE value is 0 or more, may be derived, and as a result, it was recognized that a content relationship between C, Si, and Al needed to be controlled such that a value defined by Relationship 3 below satisfies 1.8 or less:

$5 \times [C] + [Si] + 0.5 \times [Al]$

(In Relationship 3 above, [C], [Si], and [Al] represent weight percentages of the elements in parentheses, respectively.)
According to an embodiment of the present disclosure, the above-described cold-rolled steel sheet may have a tensile strength (TS) of 980 MPa or more (preferably 980 to 1150 MPa, more preferably 980 to 1075 MPa), a yield strength (YS) of 740 to 950 MPa (more preferably, 790 to 920 MPa), a hole expansion ratio (HER) of 45% or more (more preferably, 50 to 65%), and an elongation (El) of 12% or more (more preferably, 12 to 20%), to secure excellent strength, ductility, and a hole expansion ratio at the same time.
The cold-rolled steel sheet of the present disclosure may have a hot-dip galvanized layer formed on at least one surface. In the present disclosure, there may be no particular limitation on the composition of the hot-dip galvanized layer, and any hot-dip galvanized layer commonly applied in the technical field may be applied to the present disclosure. Additionally, the hot-dip galvanized layer may be an alloyed hot-dip galvanized layer alloyed with some alloy components of the steel sheet.
Hereinafter, a method for manufacturing a cold-rolled steel sheet having excellent weldability, strength, and formability, according to an embodiment of the present disclosure will be described. This does not mean that the cold-rolled steel sheet of the present disclosure should be manufactured only by the following manufacturing method.
First, a slab having the above-described alloy composition may be heated. When heating the slab, a heating temperature may be 1150 to 1250°C. When the slab heating temperature is less than 1150°C, it is not possible to perform hot-rolling, which may be the next step. When the slab heating temperature exceeds 1250°C, a lot of energy may be unnecessarily consumed to increase the slab temperature. Therefore, the slab heating temperature ranges from 1150 to 1250°C. A lower limit of the slab heating temperature is more preferably 1170°C, and even more preferably 1180°C. An upper limit of the slab heating temperature is more preferably 1230°C, and even more preferably 1220°C.
Thereafter, the heated slab may be subjected to a finishing hot-rolling at 830 to 980°C to obtain a hot-rolled steel sheet. When the finishing hot-rolling temperature (hereinafter also referred to as 'FDT') is less than 830°C, rolling load may increase, shape defects may increase, and productivity may deteriorate. When the finishing hot-rolling temperature exceeds 980°C, surface quality may deteriorate due to an increase in oxides due to excessively high temperature work. Therefore, the finishing hot-rolling temperature may range from 830 to 980°C. A lower limit of the finish hot-rolling temperature is more preferably 880°C. An upper limit of the finish hot-rolling temperature is more preferably 950°C, and even more preferably 930°C.
Subsequently, the hot-rolled steel sheet may be coiled at 450 to 700°C. When the coiling temperature (hereinafter also referred to as 'CT') exceeds 700°C, there may be a disadvantage in that coarse internal oxidation of hot-rolling occurs and surface properties deteriorate. When the coiling temperature is less than 450°C, it may correspond to a transition boiling range, which has a disadvantage of worsening controllability of the coiling temperature and deteriorating a shape of the steel sheet. A lower limit of the coiling temperature is more preferably 480°C, and even more preferably 500°C. An upper limit of the coiling temperature is more preferably 670°C, and even more preferably 640°C.
After the finishing hot-rolling, cooling may be performed to the coiling temperature at an average cooling rate of 10 to 100°C/s. When the average cooling rate is less than 10°C/s, hot-rolling productivity may be low and a cooling medium with low cooling ability should be deliberately selected during actual production, and when the average cooling rate exceeds 100°C/s, there may be disadvantageous that a temperature deviation in the steel sheet is not uniform to deteriorate a shape and excessively increase strength of the steel sheet. Therefore, the average cooling rate may range 10 to 100°C/s.
Thereafter, the coiled hot-rolled steel sheet may be cold-rolled. During the cold-rolling, a cold-rolling reduction rate may be 30 to 60%. When the cold-rolling reduction rate is less than 30%, it may be difficult to secure target thickness accuracy and it may be difficult to correct a shape of the steel sheet. When the cold-rolling reduction ratio exceeds 60%, possibility of cracks occurring at an edge of the steel sheet may increase, and the cold-rolling load may excessively increase. Therefore, the cold-rolling reduction ratio may be 30 to 60%.
Subsequently, the cold-rolled steel sheet may be continuously annealed in the range of 790 to 830°C. The continuous annealing may be to heat the steel sheet to an austenite single phase region, to form austenite close to 100%, and use the same for subsequent phase transformation. When the continuous annealing temperature (hereinafter referred to as 'SS') is less than 790°C, sufficient recrystallization and austenite transformation may not occur, making it impossible to secure desired martensite and bainite fractions after annealing. When the continuous annealing temperature exceeds 830°C, productivity may decrease, coarse austenite may be formed and the material may deteriorate, and surface quality such as peeling of the plating material may deteriorate. Additionally, the continuous annealing may be performed in a continuous alloying hot-dip plating furnace.
During the continuous annealing, an atmosphere may be controlled in a continuous annealing furnace with a gas consisting of 95% or more nitrogen by volume%, and a balance of hydrogen. When a fraction of the nitrogen is less than 95% and a ratio of the hydrogen does not increase, an oxidizing atmosphere may be formed in the furnace and oxides may be formed on a surface of the steel sheet, to deteriorate a quality of the surface. When the ratio of the hydrogen is high, difficulties in process such as explosion prevention or the like may increase.
Thereafter, the continuously annealed steel sheet may be cooled to a primary cooling end temperature of 450 to 600°C (hereinafter also referred to as 'SCS') at an average cooling rate of less than 10°C/s (more preferably, between from 1°C/s or more to less than 10°C/s). The primary cooling end temperature may be defined as a time point at which secondary cooling (quick cooling) is initiated by additionally applying quenching equipment that was not applied in primary cooling. When the cooling process is divided into the primary and secondary cooling and carried out in stages, temperature distribution of the steel sheet may be made uniform in a slow cooling stage, a final temperature and a material deviation may be reduced, and a necessary phase composition may be obtained. In particular, the bainite structure of the present disclosure should be actively formed from the first cooling stage to obtain target elongation. When the primary cooling end temperature is less than 450°C, a fraction of the bainite may be excessively high, and it may be difficult to cool down to below 450°C at a cooling rate of less than 10°C/s due to an actual equipment length. When the primary cooling end temperature exceeds 600°C, an amount of cooling to the secondary cooling end temperature may increase, resulting in poor shape of the steel sheet and a lower fraction of the bainite, compared to a target level. When a primary cooling rate is less than 1°C/s, an amount of precipitation of the ferrite phase increases during cooling, making it difficult to obtain high-strength steel, and when a primary cooling rate exceeds 10°C/s, an amount of cooling in the secondary cooling may increase to increase a final temperature deviation and a material deviation. More preferably, in terms of improving the above-mentioned effect, a lower limit of the primary cooling rate may be 3°C/s, and an upper limit of the primary cooling rate may be 8°C/s.
Thereafter, the primary cooled steel sheet may be secondary cooled to a secondary cooling end temperature of 250 to 350°C (hereinafter referred to as 'RCS') at an average cooling rate of 10°C/s or more. The secondary cooling end temperature may be set to be on or below an Ms temperature of the steel sheet, such that martensite transformation occurs during cooling, and this martensite ultimately becomes a tempered martensite phase through a reheating operation, which is a post-process. Since an Ms temperature of a high-stretch steel sheet having 980 MPa may be mostly 400°C or lower, in the present disclosure, the secondary cooling end temperature was controlled to be 250 to 350°C. When the secondary cooling end temperature is less than 250°C, an initial martensite transformation amount may be too large, resulting in high yield strength and poor formability. When the secondary cooling end temperature exceeds 350°C, martensite may not be generated during cooling, making it difficult to obtain high yield strength and hole expansion ratio. When the secondary cooling rate is less than 10°C/s, even when a target secondary cooling end temperature reaches, high-temperature phase transformation may occur during cooling, making it impossible to obtain a target fraction of the martensite and high strength. More preferably, in terms of improving the above-mentioned effect, a lower limit of the secondary cooling rate is 11°C/s, and an upper limit of the secondary cooling rate is 30°C/s.
As mentioned above, the secondary cooling may additionally apply a quenching facility that was not applied in the primary cooling, and the present disclosure does not specifically limit a type of the quenching facility, but a hydrogen quenching facility may be used as a preferred example. More specifically, the hydrogen quenching facility may use a gas consisting of 5 to 80% hydrogen by volume and a balance of nitrogen. When a fraction of the hydrogen exceeds 80%, there may be a disadvantage in that it becomes difficult to manage equipment such as explosion control or the like, and when a fraction of the hydrogen is less than 5%, there may be a disadvantage in that it becomes difficult to utilize efficient heat transfer characteristics of hydrogen, which is a light element.
Thereafter, the secondary cooled steel sheet may be reheated to 350 to 480°C. Through the above process, interphase carbon distribution and additional bainite phase transformation, necessary for stabilization of retained austenite, may be obtained. In the present disclosure, an end point temperature of the heating section may be referred to as a reheating temperature (hereinafter also referred to as 'RHS') for convenience. When the reheating temperature is less than 350°C, strength may be excessively high and elongation may deteriorate. When the reheating temperature exceeds 480°C, the austenite phase may not be transformed and may remain, but may be fresh martensite during final cooling, impairing a hole expansion ratio and elongation. A so-called nose temperature at which bainite transformation is most active may be about 400 to 420°C. In consideration of this, a lower limit of the reheating temperature is more preferably 411°C, or an upper limit of the reheating temperature is more preferably 440°C.
Although not particularly limited, according to an embodiment of the present disclosure, an average temperature increase rate during reheating may be 0.5 to 2.5°C/s. When the average temperature increase rate is less than 0.5°C/s, an overall process time may be too long, which may cause excessive heat treatment, and when the average temperature increase rate exceeds 2.5°C/s, there may be a risk that it is difficult to secure desired physical properties in the present disclosure.
In addition, the present inventors conducted a research and precisely controlled conditions of the above-mentioned primary cooling and secondary cooling to satisfy the following Relationship 4. Therefore, it was discovered that the hole expansion ratio could be improved by reducing a difference in hardness between phases, such that a sufficient bainite structure could be obtained in the primary cooling and secondary cooling sections.

$V 1 / V 2 \times t > 0.5$

(In Relationship 4 above, V1 represents an average cooling rate during primary cooling, V2 represents an average cooling rate during secondary cooling, and t represents a thickness of the cold-rolled steel sheet.)
In addition, according to an embodiment of the present disclosure, after reheating, as necessary, the reheated steel sheet may be additionally subjected to a hot-dip galvanizing process, an alloyed hot-dip galvanizing process, and a temper-rolling process. Specifically, plating the reheated steel sheet in a zinc plating bath in a temperature range of 450 to 470°C may be further included.
In addition, according to an embodiment of the present disclosure, as necessary, alloying and heat-treating the plated steel sheet in a temperature range of 470 to 550°C may be further included. The alloying and heat-treating may be to obtain an appropriate alloying level, and a temperature thereof may be determined according to a surface condition of the steel sheet. By controlling the surface condition of the steel, an alloying heat treatment temperature should not exceed 550°C to prevent softening of the steel sheet and loss of retained austenite due to excessive tempering. To quickly proceed with alloying, the alloying heat treatment temperature may be higher than a hot-dip galvanizing temperature. Therefore, a lower limit may be controlled to 470°C. In addition, after the alloying heat treatment, to correct a shape of the steel sheet and adjust yield strength, cooling the alloyed heat-treated steel sheet to room temperature and then temper-rolling at a reduction rate of less than 1% may be further included.

Mode for Invention

Hereinafter, the present disclosure will be described in more detail through Examples. However, it should be noted that the following Examples are only for illustrating and embodying the present disclosure and may not be intended to limit the scope of the present disclosure. This may be because the right scope of the present disclosure may be determined by matters stated in the claims and matters reasonably inferred therefrom.

(Examples)

After preparing a slab having an alloy composition illustrated in Table 1 below, the slab was reheated at 1180 to 1220°C, and subjected to hot-rolling, coiling, annealing, primary cooling, secondary cooling, reheating, and hot-dip galvanizing (GI) under conditions illustrated in Table 2 below, to manufacture a cold-rolled steel sheet. In addition, some steel sheets were alloyed and heat treated under alloying heat treatment temperature (GA) conditions listed in Table 2 below. In this case, after finishing hot-rolling, a cooling rate was 30 to 50°C/s, a cold-rolling reduction rate was 33 to 55%, gas used during continuous annealing was 95% by volume N - 5% by volume H, and gas used during secondary cooling was 75% by volume N - 25% by volume H.
Evaluation results for a tensile property, a hole expansion ratio, and spot weld LME of the steel sheet manufactured in this manner were illustrated in Table 3 below. Tensile strength (TS), yield strength (YS), and elongation (EL) were measured through a tensile test in a direction, perpendicular to rolling. A specimen regulation in which a gauge length was 50 mm and a width of the tensile specimen was 25 mm was used. A hole expansion ratio was measured according to ISO 16330 standards, and a hole was sheared with a clearance of 12% using a 10 mm diameter punch.
An AE value was measured by spot welding the plated steel sheet, and results therefrom were illustrated in Table 3 below. The AE value means a value obtained by subtracting a minimum current value for scattering generation from a minimum current value for LME generation. In the spot welding test, a current value increased in 0.5kA increments from a low current value, but a short cooling time was given between each current value to prevent excessive heat input to the material. As the current value increased in this manner, a minimum current value at which a nugget of a welded portion is scattered (expulsions) was measured, at the same time, a minimum current value at which LME occurs was measured from observation of a surface and a cross-section of the welded portion, and results therefrom were listed in Table 3 below. Occurrence of LME was considered as passed when there was no crack due to LME were observed with naked eyes in observing the surface of the welded portion at 10X magnification and the cross-section at 100X magnification.
In addition, results of measuring a microstructure of the manufactured cold-rolled steel sheet and calculation results of Relationships 1 to 3 used in the present disclosure were illustrated in Table 4.

The microstructure was measured using a point counting method from scanning electron microscope (SEM) photographs, and a fraction of retained austenite was measured using XRD.

[Table 1]

	Alloy Composition(wt%)
Steel	C	Si	Al	Mn	Cr	Mo	B	Nb	Ti	P	S	N
A	0.137	0.49	0.085	2.55	0.04	0.21	0.0002	0.031	0.019	0.0067	0.0015	0.0034
B	0.141	0.55	0.038	2.57	0.002	0.207	0.0001	0.019	0.018	0.0173	0.0015	0.0052
C	0.151	0.6	0.034	2.61	0.38	0.11	0.0003	0.019	0.021	0.0072	0.0021	0.0042
D	0.149	0.57	0.027	2.54	0.22	0.09	0.0002	0.021	0.022	0.0084	0.0009	0.0048
E	0.105	0.75	0.09	2.45	0.32	0.15	0.0002	0.045	0.021	0.0095	0.0014	0.0042
F	0.12	0.42	0.315	2.42	0.402	0.195	0.0002	0.03	0.021	0.012	0.0012	0.0045
G	0.102	0.15	0.1	2.44	0.2	0.05	0.0001	0.001	0.002	0.0142	0.0019	0.0044
H	0.133	0.53	0.035	2.56	0.42	0.15	0.0003	0.012	0.015	0.0093	0.0032	0.0067
I	0.166	0.43	0.025	2.71	0.14	0.05	0.0001	0.022	0.025	0.0088	0.0007	0.0072
J	0.142	0.67	0.055	2.92	0.45	0.33	0.0002	0.005	0.005	0.0065	0.0023	0.0051
K	0.125	1.23	0.15	2.57	0.33	0.21	0.0001	0.022	0.019	0.0077	0.0018	0.0038
L	0.175	1.35	0.09	2.75	0.21	0.05	0.0001	0.032	0.018	0.0076	0.0018	0.0044
M	0.192	0.77	0.45	2.62	0.44	0.12	0.0002	0.002	0.0005	0.0124	0.0022	0.0057

[Table 2]

	Steel	Hot Roll Thicknes s [mm]	Cold Roll Thickness (t) [mm]	Reductio n Rate[%]	FDT [°C ]	CT [°C ]	SS [°C ]	SCS [°C ]	V1*	RCS [°C ]	V2*	Vh*	RHS [°C ]	GI Pot [°C]	GA [°C ]
Inventive Example 1	A	2.3	1.2	48	905	605	830	549	5.0	323	11.8	1.1	425	458	517
Inventive Example 2	B	2.2	1.1	50	911	582	823	571	7.2	312	14.7	1.3	440	461	-
Inventive Example 3	C	2.1	1	52	924	616	819	490	6.2	299	11. 9	1.2	411	463	521
Inventive Example 4	D	2.6	1.4	46	895	591	823	523	4 . 7	315	12.8	1.3	432	460	519
Comparativ e Example 1	E	2.4	1.2	50	899	621	810	561	6.2	335	14.3	1.1	425	456	520
Comparativ e Example 2	F	2.1	1.2	43	933	572	825	532	5.9	304	13.4	2.1	454	455	518
Comparativ e Example 3	G	2.6	1.3	50	875	602	833	632	4.3	305	19.2	2.3	457	462	-
Comparativ e Example 4	H	2.1	1	52	872	661	841	552	5.2	392	10.1	0.9	453	466	505
Comparativ e Example 5	I	2.4	1.2	50	887	535	825	425	7.7	335	4 . 7	0.9	423	453	-
Comparativ e Example 6	J	2.1	1.3	38	933	552	823	532	5.1	325	12.6	1.1	435	462	521
Comparativ e Example 7	K	2.6	1.4	46	912	656	833	552	5.7	312	15.3	1.4	433	463	-
Comparativ e Example 8	L	2.1	1.2	43	905	618	835	565	6.8	304	14.2	1.0	392	461	532
Comparativ e Example 9	M	2.6	1.3	50	930	605	845	593	5.2	322	16.5	1.3	425	444	512

V1*: Average cooling rate during primary cooling [°C/s]
V2*: Average cooling rate during secondary cooling [°C/s]
Vh*: Average temperature increase rate during reheating [°C/s]

[Table 3]

	Steel	Microstructure Fraction [area %]				L_MA*	Relationship
	Steel	F	B	M	Retained γ	[µm]	[1]	[2]	[3]
Inventive Example 1	A	5	63	29	3	0.44	75.7	286	1.22
Inventive Example 2	B	7	45	44	4	0.40	81.6	286	1.27
Inventive Example 3	C	2	46	49	3	0.45	89.5	311	1.37
Inventive Example 4	D	3	52	43	2	0.50	87.4	291	1.33
Comparative Example 1	E	18	36	43	3	0.58	60.3	283	1.32
Comparative Example 2	F	28	40	30	2	0.57	44.3	294	1.18
Comparative Example 3	G	37	28	32	3	0.63	78.5	265	0.71
Comparative Example 4	H	3	40	53	4	0.72	86.0	308	1.21
Comparative Example 5	I	3	74	20	3	0.46	100.3	303	1.27
Comparative Example 6	J	4	15	77	4	0.33	89.9	359	1.41
Comparative Example 7	K	25	25	45	5	0.62	46.9	305	1.93
Comparative Example 8	L	5	22	65	8	0.29	69.2	313	2.27
Comparative Example 9	M	35	13	45	7	0.56	43.3	328	1.96
		F: Ferrite,
		B: Bainite,
		M: Martensite,
		γ: Austenite

L_MA*: Average size of martensite-austenite (MA) present in bainite [pm]

[Table 4]

	Steel	Mechanical Property				LME Property
	Steel	YS [MPa]	TS [Mpa]	EL [%]	HER [%]	AE	Expulsion Generated Current [kA]	LME Generated Current [kA]
Inventive Example 1	A	798	1042	15	51	1.5	10.5	12
Inventive Example 2	B	895	1071	13	65	1.0	11	12
Inventive Example 3	C	912	1060	13	63	1.0	10.5	11.5
Inventive Example 4	D	848	1035	14	59	1.0	10.5	11.5
Comparative Example 1	E	723	1012	15	38	1.0	10	11
Comparative Example 2	F	648	990	16	35	1.5	10	11.5
Comparative Example 3	G	668	848	17	27	1.0	10	11
Comparative Example 4	H	705	1108	13	37	0.5	10.5	11
Comparative Example 5	I	673	967	15	47	1.0	10.5	11.5
Comparative Example 6	J	1053	1172	11	46	0.5	11	11.5
Comparative Example 7	K	797	1045	13	44	-0.5	11	10.5
Comparative Example 8	L	989	1112	14	42	-1.0	11.5	10.5
Comparative Example 9	M	852	1251	13	39	-0.5	10.5	10

As can be seen from Tables 1 to 4 above, Inventive Examples 1 to 4 manufactured using steel A to D satisfied the alloy composition, Relationship 1, and the manufacturing conditions proposed by the present disclosure. Therefore, it can be confirmed that the present disclosure acquired the desired microstructure to secure a target tensile strength of 980 to 1150 MPa, a yield strength of 740 to 950 MPa, a hole expansion ratio (HER) of 45% or more, and an elongation of 12% or more, and at the same time, the LME properties were also excellent.
In particular, to observe martensite-austenite (MA) present in bainite of a cold-rolled steel sheet obtained in Inventive Example 1 of the present application, a photograph of a cross-section in a thickness direction was observed at 5,000X magnification with a scanning electron microscope (SEM) is illustrated in FIG. 1. In FIG. 1, the martensite-austenite (MA) present in the bainite is indicated by an arrow.
In Comparative Examples 1 and 2, it can be confirmed that a fraction of a ferrite phase exceeded 10%, a tensile strength, a hole expansion ratio, and a hardness deviation were inferior.
Additionally, in Comparative Examples 3 and 6, a fraction of ferrite or martensite phase was outside the defined range of the present disclosure, so the required material was not obtained.
In addition, in Comparative Examples 4 and 5, required materials could not be obtained as they were outside the defined process range of the present disclosure.
In addition, in Comparative Examples 7 and 8, a large amount of Si was contained beyond the defined component range of the present disclosure, and in Comparative Example 9, an amount of C or the like was outside the defined range. In LME evaluation, all of these steels show that a minimum current at which LME occurs may be lower than a minimum current at which expulsion occurs, making spot weld LME vulnerable, respectively.

Claims

A cold-rolled steel sheet comprising:
by weight%, C: 0.10 to 0.16%, Si: 0.3 to 0.8%, Al: 0.01 to 0.5%, Mn: 2.0 to 3.0%, Cr: 0.001 to 0.5%, Mo: 0.001 to 0.5%, B: 0.0001 to 0.001%, Nb: 0.001 to 0.05%, Ti: 0.001 to 0.05%, P: 0.04% or less (excluding 0%), S: 0.01% or less (excluding 0%), N: 0.01% or less (excluding 0%), a remainder of Fe, and other unavoidable impurities,

as a microstructure comprises, by area%, ferrite: 10% or less (excluding 0%), retained austenite: more than 1% to 5% or less, martensite: 25% or more but less than 50%, and bainite: 35% or more but less than 70%, and

an average size of martensite-austenite (MA) in the bainite is 0.35 to 0.55 um.
The cold-rolled steel sheet of claim 1, wherein a value defined by the following Relationship 1 satisfies 70 or more: 234×[C] - 29×[Si] - 128×[Al] + 29×[Mn] + 10×[Cr] - 17×[Mo] - 37×[Nb] - 49×[Ti] + 100×[B] where [C], [Si], [Al], [Mn], [Cr], [Mo], [Nb], [Ti], and [B] represent weight percentages of the elements in parentheses, respectively.
The cold-rolled steel sheet of claim 1, wherein a value defined by the following Relationship 2 satisfies 270 or more and 330 or less: $270 \times [C] + 90 \times [Mn] + 70 \times [Cr] + 80 \times [Mo]$
where [C], [Mn], [Cr], and [Mo] represent weight percentages of the elements in parentheses, respectively.
The cold-rolled steel sheet of claim 1, wherein a value defined by the following Relationship 3 satisfies 1.8 or less: $5 \times [C] + [Si] + 0.5 \times [Al]$
where [C], [Si], and [Al] represent weight percentages of the elements in parentheses, respectively.
The cold-rolled steel sheet of claim 1, wherein the microstructure comprises, by area%, 2 to 7% of the ferrite.
The cold-rolled steel sheet of claim 1, wherein the microstructure comprises, by area%, 2 to 4% of the retained austenite.
The cold-rolled steel sheet of claim 1, wherein the microstructure comprises, by area%, 45 to 63% of the bainite.
The cold-rolled steel sheet of claim 1, wherein the microstructure comprises, by area%, 29 to 49% of the martensite.
The cold-rolled steel sheet of claim 1, wherein the cold-rolled steel sheet has a tensile strength of 980 to 1150 MPa and a yield strength of 740 to 950 MPa.
The cold-rolled steel sheet of claim 1, wherein the cold-rolled steel sheet has a hole expansion ratio (HER) of 45% or more.
A method for manufacturing a cold-rolled steel sheet, comprising:
heating a steel slab including, by weight%, C: 0.10 to 0.16%, Si: 0.3 to 0.8%, Al: 0.01 to 0.5%, Mn: 2.0 to 3.0%, Cr: 0.001 to 0.5%, Mo: 0.001 to 0.5%, B: 0.0001 to 0.001%, Nb: 0.001 to 0.05%, Ti: 0.001 to 0.05%, P: 0.04% or less (excluding 0%), S: 0.01% or less (excluding 0%), N: 0.01% or less (excluding 0%), a remainder of Fe, and other unavoidable impurities;

hot-rolling the heated slab in a finish hot rolling temperature range of 830 to 980°C to obtain a hot-rolled steel sheet;

coiling the hot-rolled steel sheet in a temperature range of 450 to 700°C;

cold-rolling the coiled hot-rolled steel sheet;

continuously annealing the cold-rolled steel sheet in a temperature range of 790 to 830°C;

primary cooling the continuously annealed steel sheet at an average cooling rate of less than 10°C/s to a primary cooling end temperature of 450 to 600°C;

secondary cooling the primary cooled steel sheet at an average cooling rate of 10°C/s or more to a secondary cooling end temperature of 250 to 350°C; and

reheating the secondary cooled steel sheet to a temperature in a range of 350 to 480°C,

wherein the method for manufacturing the cold-rolled steel sheet satisfies the following Relationship 4: $V 1 / V 2 \times t > 0.5$
where V1 represents an average cooling rate during the primary cooling, V2 represents an average cooling rate during the secondary cooling, and t represents a thickness of the cold-rolled steel sheet.
The method of claim 11, wherein, during the cold-rolling, a cold-rolling reduction rate is 30 to 60%.
The method of claim 11, further comprising plating the reheated steel sheet in a zinc plating bath in a temperature range of 450 to 470°C.
The method of claim 13, further comprising alloying and heat-treating the plated steel sheet in a temperature range of 470 to 550°C.
The method of claim 14, further comprising cooling the alloy heat-treated steel sheet to room temperature and then temper-rolling the cooled alloy heat-treated steel sheet at a reduction ratio of less than 1%.
The method of claim 11, wherein an average temperature increase rate during the reheating is 0.5 to 2.5°C/s.