ULTRA HIGH STRENGTH COLD ROLLED STEEL SHEET HAVING EXCELLENT DUCTILITY AND DELAYED FRACTURE RESISTANCE AND METHOD FOR MANUFACTURING THE SAME

- JFE STEEL CORPORATION

An ultra.-high-strength cold-rolled steel sheet with excellent ductility and delayed fracture resistance includes 0.15% to 0.75 C. 1.0% to 3.0% Si, 1.5% to 2.5% Mn, 0.05% or less P, 0.02% or less 5, 0.01% to 0.05% Al, and less than 0.005% N on a mass ratio, the remainder being Fe and =avoidable impurities, the ultra-high-strength cold-rolled steel sheet having a metal microstructure including 40% to 85% of a tempered martensite phase and 15% to 60% of a ferrite phase on a volume fraction basis and a tensile strength of 1320 Mtn or more.

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Description
RELATED APPLICATIONS

This is a §371 of International Application No. PCT/JP2011 /065135, with an international filing date of Jun. 24, 2011 (WO 2012/002520 Al, published Jan. 5, 2012), which is based on Japanese Patent Application No. 2010-148531, filed Jun. 30, 2010, the subject matter of which is incorporated herein by reference.

TECHNICAL FIELD

This disclosure relates to an ultra-high-strength cold-rolled steel sheet which has an excellent strength-ductility balance and excellent delayed fracture resistance and which is a material suitable for use principally in ultra-high-strength automobile structural parts such as center pillars and door impact beams for automobiles and also relates to a method for manufacturing the same.

BACKGROUND

In recent years, in Europe, regulations on CO2 emissions from automobiles, which are mobile CO2 emission sources, have been tightened because of concerns about global warming due to increasing CO2 emissions and therefore the improvement of automobile fuel efficiency has been strongly required. The lightening of car bodies is effective in improving fuel efficiency. However, since the safety of occupants needs to be ensured, the crash safety of light-weight car bodies needs to be more improved than ever. To cope with two requirements, that is, the lightening of car bodies and the ensuring of crash safety, the gauge reduction of steel sheets used is being attempted using materials with high specific strength. In recent years, high-strength steel sheets with a tensile strength of 980 to 1180 MPa have been actively used for auto-mobile structural parts such as center pillars and door impact beams. However, there are increasing demands for lightweight car bodies and therefore attempts are being made to manufacture more lightweight car bodies using steel sheets stronger than 1180 MPa class steel sheets,

Since automobile structural parts are usually manufactured by press molding, the ductility of materials significantly affects the press formability thereof in view of automobile crash safety, residual ductility after press molding is important. Since the ductility of steel sheets usually decreases with an increase in strength, press formability and residual ductility after press molding decrease with an increase in strength. In high-strength materials with a tensile strength of greater than 980 MPa, there are concerns about delayed fracture due to residual ductility after press molding and hydrogen coming from surroundings. Therefore, to use high-strength cold-rolled steel sheets for the above automobile structural parts, the high-strength cold-rolled steel sheets need to have high press formability, high ductility, and excellent delayed fracture resistance

Various proposals have been made to cope with these requirements.

For example, Japanese Unexamined Patent Application Publication No. 2005-163055 discloses an example in which a steel sheet assumed to have a tensile strength of 1350 MPa and a tempered martensite single-phase microstructure is obtained by quench and tempering although the percentages of phases are not described therein. However, the total elongation of the steel sheet is small, 7%. Therefore, it is extremely difficult to manufacture automobile safety parts from the steel sheet by pressing. The martensite single-phase microstructure is probably obtained by quenching and therefore the steel sheet probably has a seriously had shape. This case needs a step of correcting the shape thereof after annealing and therefore is not preferable in terms of manufacture.

Japanese Unexamined Patent Application Publication No. 2006-307325 discloses a TRIP (Transformation-induced Plasticity) steel sheet which has high strength and ductility and which is obtained by making use of strain-induced transformation, that is, the transformation of retained austenite into martensite by strain during deformation. That steel sheet contains 0.3% to 2% Al on a mass basis to ensure the amount of retained austenite necessary to develop a TRIP effect. A large amount of Al causes a problem that casting defects are likely to he caused. To allow retained austenite to remain in a microstructure, isothermal holding needs to be performed at a temperature not lower than the its transformation temperature in the course of cooling from the annealing temperature, which results in an increased number of manufacturing steps. Since the change in rate of cooling to the temperature of isothermal holding during operation causes a significant change in material quality, operating conditions needs to be strictly controlled to stably manufacture steel sheets with a certain level of quality, which is not preferable in terms of manufacture.

It could therefore be helpful to provide an ultra-high-strength cold-rolled steel sheet which has excellent delayed fracture resistance and a tensile strength of 1320 MPa or more and which does not excessively contain a transition metal element, such as V or Mo, causing a significant increase in alloying cost or Al, which may possibly cause casting defects, and to provide a method for manufacturing the ultra-high-strength cold-rolled steel sheet.

SUMMARY

To obtain a conventional ultra-high-strength cold-rolled steel sheet with a tensile strength of 1320 MPa or more, a microstructure needs to be transformed into a martensite single-phase microstructure by quenching. In the case where a microstructure is a martensite single-phase, sufficient ductility cannot be achieved. Even it an attempt is made to increase the ductility by tempering subsequent to quenching, the strength is reduced and ductility is apt not to be increased so much because of the recovery of a dislocation microstructure in a martensite phase and the coarsening of a carbide such as Fe3C, precipitated in the martensite phase.

On the other hand, to develop high ductility, many TRIP steels have been invented by making use of file strain-induced transformation of a retained austenite. However, to develop a TRIP effect, a large amount of an alloying element needs to be used to increase the stability of austenite and isothermal holding needs to be precisely performed at temperature not lower than the Ms transformation temperature in the course of cooling from the annealing temperature, which is not preferable in terms of manufacturing stability and manufacturing costs.

In view of delayed fracture resistance, hydrogen-trapping sites, which cause delayed fracture, are preferably diminished as much as possible. Martensite phases are preferably diminished as much as possible because a large number of dislocations serving as hydrogen-trapping sites are introduced into the martensite phases during crystallographic transformation from austenitic phases. Retained austenite, which contributes to an increase in ductility, is known to serve as a hydrogen-trapping site like a dislocation and is present on a grain boundary in the form of a film. Therefore, the penetration of hydrogen into retained austenite may possibly cause grain boundary fracture to reduce delayed fracture resistance. Thus, it is not preferred that a metal microstructure contains retained austenite,

We discovered that the balance between tensile strength and ductility can be controlled such that a microstructure is converted into a microstructure containing a tempered martensite phase and a ferrite phase and the volume fraction of the tempered martensite phase is varied. We discovered a technique in which a steel sheet with ultra-high strength is obtained such that the hardness of the tempered martensite phase and that of the ferrite phase are increased by the addition of C and Si the volume fraction of an untempered martensite phase is reduced. We found that an ultra-high-strength steel sheet with high ductility can be obtained.

In addition, we discovered that the density of dislocations in a microstructure can be significantly reduced as compared with a smart-ensile single-phase microstructure by precipitating a ferrite phase containing substantially no dislocation in the microstructure and the amount of hydrogen permeating through steel can be significantly reduced by diminishing hydrogen-trapping sites. We thus found that delayed fracture resistance can be increased.

Furthermore, we found that in view of manufacturing steps, it is effective the annealing temperature and the course of cooling are appropriately controlled during annealing and cooling subsequent to cold rolling and tempering heat treatment is performed at a temperature of 100° C. to 300° C.

We thus provide:

    • (1) An ultra-high-strength cold-rolled steel sheet with excellent ductility and delayed fracture resistance contains 0.15% to 0.25% C. 1.0% to 3.0% Si, 1.5% to 2.5% Mn, 0.05% or less P. 0.02% or less S. 0.01% to 0.05% Al, and less than 0.005% N on a mass ratio, the remainder being Fe and unavoidable impurities, and has a metal microstructure containing 40% to 85% of a tempered martensite phase and 15% to 60% of a terrific phase on a volume fraction basis and a tensile strength of 1320 MPa or more.
    • (2) The ultra-high-strength cold-roiled steel sheet with excellent ducti thy and delayed fracture resistance specified in item (1) further contains one or more of 0.1% or less Nb, 0.1% or less Ti, and 5 ppm to 30 ppm B on a mass ratio,
    • (3) The ultra-high-strength cold-rolled steel sheet with excellent ductility and delayed fracture resistance specified in Item (1) or (2) has a total elongation of 12% or more.
    • (4) A method for manufacturing an ultra-high-strength cold-rolled steel sheet having excellent ductility and delayed fracture resistance includes heating a steel slab having the chemical composition specified in Item (1) or (2) to 1200° C. or higher; hot-rolling the steel slab at a finish rolling end temperature of 800° C. or higher; pickling the steel; cold-rolling the steel; continuously annealing the steel in such a manner that the steel is held at a temperature ranging from the Ac 1 transforrnation temperature to Ac3 transformation temperature thereof for 30 s to 1200 s, is cooled to a temperature of 600° C. to 800° C. at an average cooling rate of 100° C./s or less, and is then cooled to 100° C. or lower at an average cooling rate of 100° C./s to 1000° C./s; and tempering the steel in such a manner that the steel is reheated and is held at a temperature of 100° C. to 300° C. for 120 s to 1800 s.

Our cold-rolled steel sheet has extremely high tensile strength, high ductility, and therefore excellent workability. Parts fortned from the cold-rolled steel sheet have resistance to delayed fracture due to hydrogen coming from surroundings, that is excellent delayed fracture resistance. For example, a tensile strength of 1320 MPa or more, a total elongation of 12% or more, and such delayed fracture resistance that fracture does not occur for 100 hours in a 25° C. hydrochloric acid environment with a pH of 3 can be readily achieved, Furthermore, a cold-rolled steel sheet having such excellent properties as described above can be stably manufactured by our method.

The following sheet can be stably manufactured: an ultra-high-strength cold-rolled steel sheet which has a tensile strength of 1320 MPa or more and which exhibits excellent workability during forming. Parts formed from the cold-rolled steel sheet by press molding have resistance to delayed fracture due to hydrogen coming from surroundings, that is, excellent delayed fracture resistance. Ultra-high-strength parts, such as automobile safety parts including center pillars and impact beams, resistant to delayed fracture can be provided.

BRIEF DESCRIPTION OF THE DRAWING

FIG. 1 is a schematic view of a 180-degree bent specimen subjected to stress by bolting.

 REFERENCE SIGNS LIST

  • 1 specimen
  • 2 bolt

DETAILED DESCRIPTION

An ultra-high-strength cold-rolled steel sheet has a specific chemical composition and a microstructure as described below. The chemical composition of the cold-rolled steel sheet is first described.

(C: 0.15% to 0.25% by miss)

C is an element which stabilizes austenite and is necessary to ensure the strength of the steel sheet. When the content of C is less than 0.15% by mass, it is difficult for a microstructure having a tempered martensite phase and a ferrite phase to stably obtain a tensile strength of 1320 MPa or more. However, when the content of C is more than 0.25% by mass, welded portions and heat-affected zones affected by welding are significantly hardened and therefore weldability is reduced. Therefore, the content of C is preferably 0.15% to 0.25% by mass and more preferably 0.18% to 0.22% by mass.

(Si: 1.0% to 3,0% by mass)

Si is a substitutional solid solution hardening element effective in hardening the steel sheet. The content of Si needs to be 1.0% by mass or more to develop this effect. When the content of Si is more than 3.0% by mass, scales are significantly formed during hot rolling and the failure rate of final products is increased, which is not economically preferred. Therefore, the content of Si is 1.0% to 3.0% by mass.

(Mn: 1.5% to 2.5% by mass)

Mn is an element which stabilizes austertite and is effective in hardening steel. When the content of Mn is less than 1.5% by mass, it is difficult to stably manufacture the steel sheet having a target strength because the hardenability of steel is insufficient, the production of a ferrite phase during cooling from the annealing temperature and the formation of pearlite and bainite begin early, and the strength is significantly reduced. However, when the content thereof is more than 2.5% by mass, segregation is serious, workability is deteriorated in some cases, and delayed fracture resistance is reduced. Therefore, the content of Mn is preferably 1.5% to 2.5% by mass and more preferably 1.5% to 2.0% by mass.

(P: 0.05% by mass or less)

P is an element conductive to grain boundary fracture due to grain boundary segregation and Therefore is preferably low. The upper limit thereof is 0.05% by mass and is preferably 0.010% by mass. In view of an increase in weldability. the upper limit thereof is more preferably 0.008% by mass or less.

(S: 0.02% by mass or less)

S forms an inclusion such as MnS, causing a reduction in impact resistance and/or delayed fracture resistance and is preferably minimized. The upper limit thereof is 0.02% by mass and preferably 0.002% by mass.

(Al: 0.01% to 0.05% by mass)

Al is an element effective in deoxidization. The content thereof needs to be 0,01% by mass or more to achieve an effective deoxidizing effect. However, when the content thereof is excessive, more than 0.05% by mass, the steel sheet contains increased amounts of inclusions and has reduced ductility. Therefore, the content of Al is 0.01% to 0.05% by mass.

(N: less than 0.005% by mass)

When the content of N is 0.005% by mass or more, the formation of nitrides causes a reduction in ductility at high temperature and low temperature. Therefore, the content of N is less than 0.005% by mass,

The steel sheet may further contain one or more of Nb, Ti, and B as required. The effect of the addition of these three elements and the preferred content thereof are described below.

(Nb and Ti: 0.1% by mass or less)

Nb and Ti are elements which have a grain-re:Fining effect and are effective in increasing the strength of the steel sheet. Hence, the content of is preferably 0.015% by mass or more. However, when the content of each of Nb and Ti is more than 0.1% by mass, the effect thereof is saturated, which is not economically preferred. Therefore, the content of each of Nb and Ti is 0.1% by mass or less.

(B: 5 ppm to 30 ppm by mass)

B is an element effective in increasing the strength of the steel sheet. The strength-in-creasing effect of B cannot be expected when the content of B is less than 5 ppm by mass. However, when the content of B is more than 30 ppm by mass, hot workability is reduced, which is not preferable in terms of manufacture, Therefore, the content of B is 5 ppm to 30 ppm by mass.

The remainders other than the above components are Fe and unavoidable impurities.

The microstructure of the cold-rolled steel sheet is described below.

We investigated how to increase ductility affecting press moldability and obtain a steel sheet exhibiting, excellent delayed fracture resistance after press molding. We found that the appropriate control of a microstructure is important in exhibiting high ductility. In particular, it is important that the microstructure contains 40% or more of a tempered martensite phase on a volume fraction basis after continuous annealing, the remainder being a ferrite phase. The microstructure is obtained by quenching from the annealing temperature and tempering subsequent to quenching. According to this method, an ultra-high-strength cold-rolled steel sheet with high ductility can be obtained without excessively using a transition metal element such as V or Mo, causing an increase in cost or an alloying element such as Al, possibly causing casting defects.

The less the amount of hydrogen permeating through steel is, the more excellent the delayed fracture resistance is An extremely large number of dislocations are introduced into the tempered martensite phase by the crystallographic transformation from an austenite phase to a martensite phase during quenching. When the microstructure contains an appropriate amount of the ferrite phase, the number of the dislocations, which serve as hydrogen-trapping sites causing delayed fracture, can be more significantly reduced as compared with a tempered martensite single-phase microstructure and therefore the amount of hydrogen permeating through can be reduced.

The tensile strength of steel with a microstructure containing a tempered martensite phase and a ferrite phase increases with an increase in volume fraction of the tempered martensite phase. This is because the hardness of the tempered martensite phase is higher than the hardness of the ferrite phase, the tempered martensite phase, which is a hard phase, exhibits resistance to deformation during tensile deformation, and the larger the volume fraction of the tempered martensite phase is, the more the tensile strength of the steel is close to the tensile strength of the tempered martensite single-phase microstructure. In the range of each steel component specified herein a tensile strength of 1320 MPa or more is not achieved when the volume fraction of the tempered martensite phase is less than 40%. Since ductility decreases with at increase in volume fraction of the tempered martensite phase, a microstructure containing more than 85% of the tempered martensite phase on a volume fraction basis cannot ensure the volume fraction of the ferrite phase that is necessary to achieve a high ductility of 12% or more in terms of total elongation and necessary to increase the delayed fracture resistance. When the volume fraction of the ferrite phase is less than 15%, a high ductility of 12% or more in terms of total elongation is not achieved or an increase in delayed fracture resistance not sufficient. However, when the volume fraction thereof is more than 60%, the volume fraction of the tempered martensite phase that is necessary to achieve a predetermined strength cannot be ensured.

From the above reasons, in the microstructure of the cold-rolled steel sheet, the volume fraction of the tempered martensite phase and that of the ferrite phase are 40% to 85% and 15% to 60%, respectively, and more preferably 60% to 85% and 15% to 40%, respectively. The microstructure of the cold-rolled steel sheet may be a two-phase microstructure containing a tempered martensite phase and ferrite phase each having a desired volume fraction and may contain a constituent phase such as a retained austenite phase, a bainite phase, or a pearlite phase, other than these two phases. However, large amounts of the bainite and Pearlite phases are present, the bainite phase and the pearlite phase cause a reduction in ductility and a reduction in strength, respectively. Therefore, it is not preferable that the microstructure contains large amounts of the bainite and pearlite phases. The retained austenite phase is principally present at a grain boundary in the form of a film, serves as a hydrogen-trapping site, and therefore may possibly act as an origin of fracture due to hydrogen embrittlement. Hence, the content thereof is preferably minimized. Therefore, the volume fraction of the constituent phase (the retained austenite phase, the bainite phase, or the pearlite phase) other than the tempered martensite phase and the ferrite phase is preferably 1% or less in total.

The tensile strength and ductility (total elongation as determined by a tensile test using a ES No. 5 tensile specimen) are 1320 MPa or more and 12% or more, respectively. The total elongation corresponds to the minimum elongation capable of pressing automobile structural parts such as impact beams. Such a strength level and elongation level can be readily achieved in our steel sheets. The delayed fracture resistance is such a performance that fracture does not occur for 100 hours in a 25° C. hydrochloric acid environment with a pH of 3. Such a performance can be readily achieved in our steel sheets.

Applications of the cold-rolled steel sheet are not particularly limited. Since the cold-rolled steel sheet has the above properties, the cold-rolled steel sheet is particularly suitable for ultra-high-strength automobile safety parts such as automobile door impact beams and center pillars. Steel sheets include steel strips. The cold-rolled steel sheet may be subjected to surface treatment such as plating (electroplating or the like) or chemical conversion to be used as a surface-treated steel sheet.

A method for manufacturing the ultra-high-strength cold-rolled steel sheet will now be described.

Steel with the above composition is produced and is then continuously cast into a cast slab (slab). After being heated to 1200° C. or higher, the slab is hot-rolled at, a finish rolling end temperature of 800° C. or higher. Reasons for limiting hot rolling are described below. (Slab-beating temperature of 1200° C. or higher)

An increase in rolling load increases the risk of causing troubles during hot rolling when the heating temperature of the slab is lower than 1200° C. Thus, the heating temperature of the slab is 1200° C. or higher. An increase in oxidation causes an increase in scale loss when the heating temperature thereof is excessively high. Thus, the heating temperature of the slab is preferably 1300° C. or lower.

(Finish Rolling End Temperature of 800° C. or Higher)

A uniform hot-rolled microstructure can be obtained when the finish rolling end temperature is 800° C. or higher. When the finish rolling end temperature is lower than 800° C., the microstructure of the steel sheet is nonuniform, the ductility thereof is, and the risk of causing various failures during molding is increased. Thus, the finish rolling end temperature is 800° C. or higher. The upper limit of the finish rolling end temperature is not particularly limited and is preferably 1000° C. or lower because rolling at excessively high temperature causes scale defects.

The hot-rolled steel sheet is coiled. The coiling temperature thereof is not particularly limited. When the coiling temperature thereof is excessively high, the microstructure of the steel sheet is nonuniform and the ductility thereof is low, due to formation of coarse grains. When the coiling temperature thereof is excessively low, a detbrmed microstructure caused by hot rolling remains to increase the rolling load in cold rolling subsequent to hot rolling. Therefore, the coiling temperature thereof is preferably 600° C. to 700° C. In particular, the coiling temperature thereof is preferably 600° C. to 650° C.

The hot-rolled steel sheet is pickled, cold-rolled, continuously annealed, and then tempered. Pickling and cold rolling conditions are not particularly limited. steel sheet is continuously annealed such that the steel sheet is held at a temperature ranging from the Ac1 transformation temperature to Ac3 transformation temperature thereof for 30 s to 1200 s, cooled to a temperature of 600° C. to 800° C. at an average cooling rate of 100° C./s or less, and then cooled to 100° C. or lower at an average cooling rate of 100° C./s to 1000° C./s. The steel sheet is subsequently tempered such that the steel sheet is reheated and held at a temperature of 100° C. to 300° C. for 120 s to 1800 s. Reasons for limiting continuous annealing and tempering conditions are described below.

(Annealing Temperature: Holding at Temperature Ranging from Ac1 Transformation Temperature to Ac3 Transformation Temperature for 30 s to 1200 s)

When the annealing temperature is lower than the Aci transformation temperature, an austenite phase (transformed into a martensite phase after quenching) necessary to ensure a predetermined strength is not produced during annealing and therefore such a predetermined strength cannot be achieved even if quenching is performed subsequently to annealing. Even if the annealing temperature is higher than the Ac3 transformation temperature, 40% or more of the martensite phase can be obtained on a volume fraction basis by controlling a ferrite phase precipitated during cooling from the annealing temperature. In the ease of performing annealing at a temperature higher than the Ac3 transformation temperature, a desired microstructure is unlikely to be obtained. Therefore, the annealing temperature ranges from the Ac1 transformation temperature to the Ac3 transformation temperature. In view of stably ensuring the equilibrium volume fraction of the austenite phase to be 40% or more within this temperature range, 760° C. or higher is preferred and 780° C. or higher is more preferred. When the holding time (annealing time) at the annealing temperature is excessively short, a microstructure is not sufficiently annealed, a nonuniform microstructure in which a deformed microstructure caused by hot rolling is present is caused, and the ductility is reduced. However, when the holding time is excessively long, an increase in manufacturing time is caused, which is not preferable in terms of manufacturing costs. Therefore, the holding time is 30 seconds to 1200 seconds. In particular, the holding time is preferably 250 seconds to 600 seconds.

(Cooling (Annealing) to Temperature of 600° C. to 800° C. at Average Cooling Rate of 100° C./s or Less)

The steel sheet is cooled (the term “cool” is hereinafter referred to as “anneal” in some cases) to a temperature (annealing end temperature) of 600° C. to 800° C. from the annealing temperature at an average cooling rate of 100° C./s or less. The ferrite phase is precipitated during annealing from the annealing temperature and the strength-ductility balance ran be thereby controlled. When the annealing end temperature is lower than 600° C., a large amount of pearlite is formed in the microstructure to cause a significant reduction in strength and therefore a tensile strength of 1320 MPa cannot be achieved. A sufficient amount of the ferrite phase cannot be precipitated during annealing from the annealing temperature and therefore sufficient ductility cannot be achieved when the annealing and temperature is higher than 800° C. Therefore, the annealing end temperature is 600° C. to 800° C. The annealing end temperature is preferably 700° C. to 750° C. to suppress a change in material quality due to an operational change in annealing end temperature.

When the average annealing rate during annealing is more than 100° C./s a sufficient amount of the ferrite phase is not precipitated and therefore predetermined ductility cannot be achieved. The ductility of the microstructure, which contains the tempered martensite phase and the ferrite phase results from high work hardenability developed by the coexistence of the tempered martensite phase, which is hard, and the ferrite phase, which is soft. The concentration of carbon in the austenite phase during annealing is insufficient and therefore a hard martensite phase cannot be obtained during quenching when the average annealing rate is more than 100° C./s. As a result, the work hardenability of a final microstructure is reduced and therefore sufficient ductility is not achieved. Therefore, the average annealing rate during annealing is 100° C./s or less. The average annealing rate is preferably 5° C./s or less to sufficiently concentrate carbon in the austenitic phase.

(Cooling (Quenching) to 100° C. or Lower at Average Cooling Rate of 100° C./s to 1000° C./s)

Subsequently to annealing, the steel sheet is cooled (the term “cool” is hereinafter referred to as “quench” in some cases) to a temperature (cooling end temperature) of 100° C. or lower at an average cooling rate of 100° C./s to 1000° C./s. Quenching subsequent to annealing is performed for the purpose of transforming the austenite phase into the martensite phase. When the average cooling rate is less than 100° C./s, the austenite phase is transformed into the ferrite phase, a bainite phase, or a pearlite phase during cooling and therefore a predetermined strength cannot be achieved. However, when the average cooling rate is more than 1000° C./s, shrinkage cracks may possibly be induced in the steel sheet by cooling. Therefore, the average cooling rate during quenching is 100° C./s to 1000° C./s. The steel sheet is preferably cooled by water quenching.

The cooling end temperature is preferably 100° C. or lower, When the cooling end temperature is higher than 100° C., the volume fraction of the martensite phase is reduced because of the insufficient transformation of austenite phase into martensite phase during quenching and a reduction in material strength is caused by the self-tempering of the martensite phase produced by quenching, which is not preferable in terms of manufacture.

(Tempering: Holding at Temperature of 100° C. to 300° C. for 120 Seconds to 1800 Seconds)

Subsequent to quenching, the steel sheet is tempered for the purpose of tempering the martensite phase such that the steel sheet is reheated and then held at a temperature of 100° C. to 100° C. for 120 seconds to 1800 seconds. The tempering thereof softens the martensite phase to increase the workability. The softening of martensite is insufficient and therefore the effect of increasing the workability cannot he expected when performing tempering at lower than 100° C. Performing tempering at higher than 300° C. increases manufacturing costs for reheating causes a significant reduction in strength, and is incapable of achieving a useful effect.

When the holding time is less than 120 s, martensite phase is not sufficiently softened at a holding temperature and therefore the effect of increasing the workability cannot be expected. When the holding, time is more than 1800 s, the strength is significantly reduced because of the excessive softening of martensite phase and manufacturing costs are increased because of an increase in reheating time, which is not preferable.

The ultra-high-strength cold-rolled steel sheet can he manufactured through the above manufacturing steps. Since the ultra-high-strength cold-rolled steel sheet has excellent shapeahility (flatness) after annealing, a step of correcting the shape of the steel sheet by roiling, leveling, or the like is not necessarily needed. In view of adjusting the quality and/or surface roughness thereof, the annealed steel sheet may he rolled with an elongation of several percent.

EXAMPLES

Test Steels A to M with compositions shown in Table 1 were produced in a \laMUM and were then formed into slabs, which were hot-rolled under conditions shown in Table 2, whereby hot-roiled steel sheets with a thickness of 3.4 mm were prepared. The hot-rolled steel sheets were surface-desealed by pickling and were then cold-rolled to a thickness of 1.4 mm. The cold-rolled steel sheets were continuously annealed and tempered under conditions shown in Table 2. The Ac1 transfbrmation temperature and Ac3 transformation temperature of each steels is determined from relational equations (the following two equations) described in The Japan Institute of Metals, Tekkou Zairyou, Maruzen, 1985, p. 43 and Kinzoku Netsushori Gijutsu Binran Henshuu Enkal, Kinzoku Netsushori Gijutsu 3rd Edition, The Nikkan Kogyo Shimbun, Ltd., 1966, p. 137, the equations being involved in the dependence of transformation temperature on alloying components:


Ac1(° C.)=723−10.7×(% by mass Mn)+29.1×(% by mass Si)   (1)


Ac3(° C.)=910−203×(% by mass C)1/2+29.1×(% by mass Si)−30×(% by mass Mn)+700×(% by mass P)+400×(% by mass Al)+400×(% by mass Ti)   (2).

TABLE 1 AC1 AC3 transformation transformation Steel Composition (% by mass) temperature temperature symbol C Si Mn P S Al N Ti Nb B (° C.) (° C.) Remarks A 0.15 1.48 1.8 0.007 0.0011 0.028 0.0031 — — — 747 837 Inventive steel B 0.18 1.48 1.8 0.007 0.0008 0.031 0.0036 — — — 747 830 Inventive steel C 0.25 1.49 1.8 0.010 0.0014 0.027 0.0024 — — — 747 816 Inventive steel D 0.20 1.03 1.8 0.011 0.0008 0.027 0.0027 — — — 734 814 Inventive steel E 0.18 2.97 1.8 0.010 0.0009 0.025 0.0028 — — — 790 873 Inventive steel F 0.20 1.52 1.5 0.011 0.0007 0.033 0.0028 — — — 751 839 Inventive steel G 0.19 1.54 2.4 0.009 0.0018 0.024 0.0033 — — — 742 810 Inventive steel H 0.18 1.51 1.8 0.010 0.0009 0.026 0.0036 0.04 — — 748 847 Inventive steel I 0.18 1.50 1.8 0.010 0.0010 0.038 0.0035 — 0.04 — 747 836 Inventive steel J 0.19 1.49 1.8 0.009 0.0010 0.033 0.0029 — — 0.002 747 830 Inventive steel K 0.19 1.48 1.8 0.007 0.0012 0.035 0.0037 0.04 0.04 0.002 747 845 Inventive steel L 0.12 1.46 2.0 0.011 0.0011 0.029 0.0041 — — — 744 841 Comparative steel M 0.15 0.44 1.6 0.009 0.0010 0.022 0.0038 — — — 719 811 Comparative steel

TABLE 2 Hot rolling step Annealing step Finish Annealing Annealing Slab-heating rolling Coiling Annealing Holding average end Steel temperature temperature temperature temperature time cooling rate temperature No. symbol (° C.) (° C.) (° C.) (° C.) (s) (° C./s) (° C.) 1 A 1250 900 650 830 600 5 750 2 B 1250 900 650 800 600 14 690 3 B 1250 900 650 800 600 14 710 4 B 1250 900 650 800 600 5 750 5 B 1250 900 650 830 600 19 700 6 C 1250 900 650 800 600 22 650 7 C 1250 900 650 800 600 4 680 8 D 1250 900 650 800 600 5 700 9 E 1250 900 650 800 600 15 750 10 E 1250 900 650 800 600 13 750 11 F 1250 900 650 800 600 5 750 12 G 1250 900 650 780 600 4 620 13 H 1250 900 650 800 600 5 700 14 H 1250 900 650 800 600 12 730 15 I 1250 900 650 800 600 4 680 16 I 1250 900 650 800 600 5 710 17 J 1250 900 650 800 600 3 720 18 K 1250 900 650 800 600 5 720 19 K 1250 900 650 800 600 19 740 20 B 1250 900 650 800  30 12 700 21 B 1250 900 650 800 1200  5 700 22 F 1250 900 650 800 600 4 750 23 F 1250 900 650 800 600 4 750 24 E 1250 900 650 800  10 14 700 25 K 1250 900 650 900 600 24 800 26 L 1250 900 650 830 600 19 650 27 M 1250 900 650 800 600 11 700 28 M 1250 900 650 780 600 10 700 29 M 1250 900 650 780 600 12 750 30 A 1250 900 650 830 600 16 500 31 A 1250 900 650 780 600 20 750 32 B 1250 900 650 830 600 19 700 Annealing step Quenching Tempering step average Cooling end Tempering Holding cooling rate temperature temperature time No. (° C./s) (° C.) (° C.) (s) Remarks  1 904 25 150 1200 Example  2 751 24 150 1200 Example  3 814 22 200 1200 Example  4 973 31 300 1200 Example  5 883 28 300 1200 Example  6 885 25 200 1200 Example  7 833 20 200 1200 Example  8 910 22 200 1200 Example  9 837 21 150 1200 Example 10 767 21 200 1200 Example 11 681 22 150 1200 Example 12 753 24 200 1200 Example 13 625 19 150 1200 Example 14 869 19 200 1200 Example 15 867 23 150 1200 Example 16 855 22 200 1200 Example 17 774 22 200 1200 Example 18 864 23 150 1200 Example 19 995 23 200 1200 Example 20 887 21 200 1200 Example 21 646 20 200 1200 Example 22 984 25 150 150 Example 23 738 24 150 1800 Example 24 846 23 300 1200 Comparative Example 25 967 22 300 1200 Comparative Example 26 941 19 300 1200 Comparative Example 27 910 25 200 1200 Comparative Example 28 809 25 200 1200 Comparative Example 29 811 26 200 1200 Comparative Example 30 786 20 200 1200 Comparative Example 31  20 19 200 1200 Comparative Example 32 889 20 400 120 Comparative Example

Specimens were taken from the steel sheets obtained through the above manufacturing steps, were observed (measured) for microstructure, and subjected to a tensile test. Furthermore, some of the steels were subjected to a delayed fracture test. The results are shown in Table 3.

The observation (measurement) of microstructure and performance tests were conducted as described below.

(1) Observation of Microstructure

Specimens were taken from the obtained cold-rolled steel sheets. A surface of each specimen that was parallel to the rolling direction was mirror-polished and was etched with nitai. The microstructure thereof was observed and photographed with an optical microscope or a scanning electron microscope, whereby the type of a constituent phase such as a tempered martensite phase or a ferrite phase was identify. A photograph of the microstructure was binarized, whereby the volume fraction of each of the tempered martensite phase the ferrite phase was determined. Since there was a possibility that a retained austenite phase was present in the obtained cold-rolled steel sheets, attempts were made to measure our examples for retained austenite phase by X-ray (Mo-Ka) determination. However, the amount of the retained austenite phase present therein was substantially zero and therefore was not included in the remainder shown in Table 3.

(2) Tensile test

JIS No. 5 tensile specimens were taken from the obtained cold-rolled steel sheets in a direction perpendicular to the rolling direction and were subjected to a tensile test according to SIS Z 2241, whereby the specimens were determined for tensile property (0.2% proof stress (YS)), tensile strength (TS), and total elongation (EL).

(3) Delayed Fracture Characterization Test

A specimen with a size of 30 mm×100 mm was cut out of each of the obtained cold-roiled steel sheets such that the longitudinal direction of the specimen corresponded to the rolling direction of the cold-rolled steel sheets. An end surface of the specimen was ground. The specimen was bent to 180 degrees using a punch having a tip with a radius of curvature of 10 mm. As shown in FIG. 1, the springback caused in the bent specimen was retained with a bolt 2 such that the distance between inner portions of the specimen I was 20 mm. After the specimen i was stressed, the specimen 1 was immersed in hydrochloric acid with a pH of 3 at 25° C. and was measured for up to 100 hours until the specimen I was broken. A specimen that was not broken within 100 hours was judged to be acceptable.

TABLE 3 Volume fraction of Volume Results of tempered fraction Other delayed fracture Steel martensite of ferrite constituent YS TS EL characterization No. symbol (%) (%) phase (MPa) (MPa) (%) test Remarks 1 A 85 15 — 972 1349 14 Acceptable Example 2 B 64 36 — 878 1338 14 Acceptable Example 3 B 74 26 — 1024 1363 15 Acceptable Example 4 B 81 19 — 1135 1326 14 Acceptable Example 5 B 85 15 — 1166 1359 12 Acceptable Example 6 C 60 40 — 819 1378 16 Acceptable Example 7 C 63 37 — 798 1336 17 Acceptable Example 8 D 81 19 — 1037 1347 13 Acceptable Example 9 E 45 55 — 789 1361 19 Acceptable Example 10 E 45 55 — 863 1332 19 Acceptable Example 11 F 68 32 — 911 1341 14 Acceptable Example 12 G 60 40 — 953 1401 13 Acceptable Example 13 H 61 39 — 819 1327 15 Acceptable Example 14 H 72 28 — 1092 1394 14 Acceptable Example 15 I 65 35 — 789 1325 13 Acceptable Example 16 I 74 26 — 1005 1368 15 Acceptable Example 17 J 80 20 — 1080 1384 13 Acceptable Example 18 K 60 40 — 796 1323 16 Acceptable Example 19 K 69 31 — 1008 1328 15 Acceptable Example 20 B 71 29 — 997 1349 14 Acceptable Example 21 B 75 25 — 976 1354 14 Acceptable Example 22 F 69 31 — 941 1364 14 Acceptable Example 23 F 70 30 — 1012 1322 13 Acceptable Example 24 E 32 58 Pearlite 951 1068 9 — Comparative Example 25 K 100 0 — 1348 1498 7 Unacceptable Comparative Example 26 L 52 48 — 889 1075 19 — Comparative Example 27 M 42 58 — 632 993 13 — Comparative Example 28 M 48 52 — 668 991 14 — Comparative Example 29 M 100 0 — 1136 1352 6 Unacceptable Comparative Example 30 A 24 66 Pearlite 437 658 30 — Comparative Example 31 A 0 72 Pearlite 462 578 32 — Comparative Example 32 B 72 28 — 983 1166 14 — Comparative Example

Tables 1 to 3 confirm that our Examples meet requirements specified herein and have a tensile strength of 1320 MPa or more, a total elongation of 12% or more, a high strength-ductility balance, and excellent delayed fracture resistance because the Examples were not broken for 100 hours in the delayed fracture characterization test

No. 24, of which the annealing time is 10 seconds and therefore is outside our range, has no predetermined strength or ductility because a Pearlite phase produced after hot rolling remains after annealing and the influence of strain due to cold rolling is not sufficiently removed. Nos. 25 and 29, each of which the annealing temperature is not lower than the Ac3 temperature, cannot precipitate any ferrite phase during annealing, have a martensite single-phase microstructure, and exhibit predetermined strength and no predetermined ductility, Nos. 26 and 27, of which steel components are outside our range, have no predetermined strength although continuous annealing and tempering were performed as specified herein. No. 30, of which the annealing end temperature is 500° C., contains a large amount of a ferrite phase precipitated therein and a pearlite phase and therefore has no predetermined strength. No. 31. of which the average cooling rate in a quenching step is 20° C./s and therefore is outside our range, cannot obtain a predetermined amount of a martensite phase and has no predetermined strength. No. 32, of which the tempering temperature is 400° C., has no predetermined strength because a martensite phase was excessively softened by tempering.

Example Nos. 1 to 23, which meet the requirements specified herein, were not broken for 100 hours in the delayed fracture characterization test. This confirms that our cold-rolled steel sheets have sufficient delayed fracture resistance, However, Comparative Example Nos. 25 and 29, each of which the microstructure is a tempered martensite single-phase which is outside our range, were broken within 100 hours and therefore failed in the delayed fracture characterization test.

INDUSTRIAL APPLICABILITY

We provide a thin steel sheet for quenching or tempering, the thin steel sheet being suitable for use principally in ultra-high-strength automobile structural parts such as door impact beams and center pillars for automobiles. In advance of manufacturing automobile parts from the steel sheet, the composition, rolling conditions, and annealing conditions are appropriately controlled. This allows the stud sheet to have a microstructure containing 40% to 85% of a tempered martensite phase and 15% to 60% of a ferrite phase on a volume fraction basis, a tensile strength of 1320 NIPa or more, a total elongation of 12% or more, an excellent strength ductility balance, and excellent delayed fracture resistance. The use of an ultra-high-strength cold-rolled steel sheet enables the pressing of automobile safety parts such as impact beams, The automobile safety parts exhibit excellent delayed fracture resistance,

Claims

1. An ultra-high-strength cold-rolled steel sheet with excellent ductility and delayed fracture resistance, comprising 0.15% to 0.25% C, 1.0% to 3.0% Si, 1.5% to 2.5% Mn, 0.05% or less P, 0.02% or less S, 0.01% to 0.05% Al, and less than 0.005% N on a mass ratio, the remainder being Fe and unavoidable impurities, the ultra-high-strength cold-rolled steel sheet having a microstructure comprising 40% to 85% of a tempered martensite phase and 15% to 60% of a ferrite phase on a volume fraction basis and a tensile strength of 1320 MPa or more.

2. The cold-rolled steel sheet according to claim 1, further comprising one or more of 0.1% or less Nb, 0.1% or less Ti, and 5 ppm to 30 ppm B on a mass ratio.

3. The cold-rolled steel sheet according to claim 1, having a total elongation of 12% or more.

4. A method for manufacturing an ulta-high-strengh cold-rolled steel sheet having excellent ductility and delayed fracture resistance, comprising;

heating a steel slab having the chemical composition specified in claim 1 to 1200° C. or higher;
hot-rolling the steel slab at a finish rolling end temperature of 800° C. or higher;
pickling the steel;
cold-rolling the steel;
continuously annealing the steel such that the steel is held at a temperature ranging from the Ac1 transformation temperature to Ac3 transformation temperature thereof for 30 a to 1200 s, cooled to a temperature of 600° C. to 800° C. at an average cooling rate of 100° C./s or less, and then cooled to 100° C. or lower at an average cooling rate of 100° C./s to 1000° C./s; and
tempering the steel such that the steel is reheated and held at a temperature of 100° C. to 300° C. for 120 s to 1800 s.

5. The cold-rolled steel sheet according to claim 2, having a total elongation of 12% or more.

6. A method for manufacturing an ultra-high-strength cold-rolled steel sheet having excellent ductility and delayed fracture resistance, comprising;

heating a steel slab having the chemical composition specified in claims 2 to 1200° C. or higher;
hot-rolling the steel slab at a finish roiling end temperature of800° C. or higher;
pickling the steel;
cold-rolling the steel;
continuously annealing the steel such that the steel is held at a temperature ranging from the Ac1 transformation temperature to Ac3 transformation temperature thereof for 30 s to 1200 s, cooled to a temperature of 600° C. to 800′C at an average cooling rate of 100° Cis or less, and then cooled to 100° C. or lower at an average cooling rate of 100° C/.s to 1000° C./s; and
tempering the steel such that the steel is reheated and held at a temperature of 100° C. to 300° C. for 120 s to 1800 a.
Patent History
Publication number: 20130087257
Type: Application
Filed: Jun 24, 2011
Publication Date: Apr 11, 2013
Applicant: JFE STEEL CORPORATION (Tokyo)
Inventors: Masataka Yoshino (Tokyo), Kohei Hasegawa (Tokyo)
Application Number: 13/805,144
Classifications
Current U.S. Class: With Working (148/648); Ferrous (i.e., Iron Base) (148/320); Beryllium Or Boron Containing (148/330)
International Classification: C21D 8/00 (20060101); C22C 38/12 (20060101); C22C 38/00 (20060101); C22C 38/04 (20060101); C22C 38/02 (20060101); C22C 38/14 (20060101); C22C 38/06 (20060101);