JP4529549B2 - Manufacturing method of high-strength cold-rolled steel sheets with excellent ductility and hole-expansion workability - Google Patents

Description

  The present invention relates to a method for producing a high-strength cold-rolled steel sheet suitable mainly for use in automobile bodies, and more particularly to improvement of ductility and hole expansion workability.

  In recent years, exhaust gas regulations for automobiles have been carried out from the viewpoint of conservation of the global environment. For this reason, weight reduction of automobile bodies has become an extremely important issue. In order to reduce the weight of automobile bodies, it is effective to reduce the thickness of parts, and high-strength steel sheets have been used as material steel sheets. However, generally, when the strength of the steel plate is increased, the workability decreases. For this reason, a high-strength steel sheet excellent in ductility is demanded for automobile bodies.

  Conventionally, as a high-strength steel sheet excellent in ductility, a composite structure steel sheet (DP steel sheet) is well known. A so-called DP steel sheet is a steel sheet having a composite structure composed of a soft ferrite phase and a hard martensite phase and having both high ductility and high strength.

  Further, as a high-strength steel plate having further excellent ductility, a steel plate using transformation induced plasticity (TRIP) of an austenite phase is known. The steel sheet using the transformation-induced plasticity of the austenite phase has a structure including a retained austenite phase in the ferrite phase or the bainite phase. This steel sheet is a steel sheet in which the cooling history from a high temperature is controlled, C is concentrated in the austenite phase by the formation of a ferrite phase or a bainite phase, and the austenite phase remains to room temperature. This steel sheet can utilize the phenomenon that the retained austenite phase is transformed by the induction of strain at the time of deformation and becomes a hard martensite phase, thereby preventing the concentration of strain and particularly improving the uniform elongation greatly.

  As a method for industrially manufacturing such a steel sheet with a low alloy composition, for example, Patent Document 1 discloses that steel containing an appropriate amount of C, Si, and Mn is wound at 650 ° C. or lower after hot rolling, and then cooled. Hot rolled, subsequently annealed at a two-phase temperature, then cooled to a temperature range of 350-500 ° C. at a cooling rate of 1-400 ° C./s, aged at that temperature range and cooled to room temperature with good ductility A method for producing a high-strength steel sheet has been proposed.

  Patent Document 2 discloses that a cold-rolled steel sheet containing a proper amount of C, Si and Mn and having Al increased to 0.25 to 1.5% is annealed in a two-phase region and then cooled to 350 to 600 ° C. And a method for producing a high-strength steel sheet with good press formability, having a structure containing 3 to 20% retained austenite, which is maintained at that temperature and then cooled to 250 ° C. or less at a cooling rate of 5 ° C./s or more. Has been proposed.

Patent Document 3 discloses that a slab adjusted to an appropriate amount of C, Si, Mn, P, S, Al, and N is finish-rolled at an Ar ₃ transformation point or higher, and then wound at 500 to 750 ° C. Cold rolled sheet with ~ 85% cold rolling is annealed at 750 ~ 900 ° C for 10s ~ 3min, then gradually cooled to 550 ~ 700 ° C at 1 ~ 10 ° C / s, 200 ~ 450 ° C Is cooled to 10 to 200 ° C / s, held at 350 to 450 ° C for 1 to 10 minutes, cooled to 200 ° C or less at a cooling rate of 5 ° C / s or more, 50% or more ferrite, 10% or more Has been proposed for producing a low-yield ratio type high-strength steel sheet having excellent ductility and having a structure composed of residual austenite, martensite of 5% or less, and the remaining bainite.

  However, although a steel sheet containing such a retained austenite phase has excellent ductility, it has a problem that it is inferior in hole expanding workability because it is a composite structure. This is because, during punching shear, microvoids due to hardness differences occur at the boundary between the ferrite phase, which is a soft phase, and the low-temperature transformation phase, such as the bainite phase, which is a hard phase, and the martensite phase. The bond is said to be connected.

  Therefore, as a method for improving the hole expansion workability of the steel sheet containing the retained austenite phase, for example, Patent Document 4 contains an appropriate amount of C, Si, Mn, Al, and Ni, and further contains Nb: 0.020 to 0.070%. , P, S, N are adjusted to an appropriate range, and the steel pieces containing Si + Al: 0.50 or more and Mn + (1/3): 1.0 or more are hot-rolled and wound at 300 to 720 ° C. After that, it is cold-rolled at a total reduction ratio of 30 to 80%, and then heated to a two-phase region, and is kept in the temperature range of 550 to 350 ° C. for 30 seconds or more during the cooling, or the temperature range is 100 ° C. / A method for producing a high-ductility, high-hole-expandability, high-tensile steel sheet that is annealed at a temperature below min has been proposed. According to this technology, NbC can be precipitated in the ferrite to increase the hardness of the ferrite ground, reduce the difference in hardness from the hard phase, suppress the generation of voids, and improve the hole expansion processability. Yes.

  Further, for example, Patent Document 5 includes appropriate amounts of C, Si, Mn, P, S, Al, and N, and further includes Ti so as to satisfy Ti / S: 5 or more, or further includes Nb, Mo, A slab containing one or more types of V is hot-rolled at a finishing delivery temperature of (900 + 50Si) ° C. or less, and cold-rolled sheets subjected to 50-80% cold rolling are treated at 700-900 ° C. After annealing for 10 s to 5 min at the phase temperature, cool to 250 to 500 ° C. with an average cooling rate between 700 to 500 ° C. being 1 to 120 ° C./s, and reheat as necessary to 250 to 600 ° C. Hold for 30 s to 10 min in the temperature range, cool down, contain martensite and residual austenite in total 6% or more, and contain Ti, Nb, Mo, V according to the volume fraction of the hard second phase. A method of producing a cold rolled steel sheet having a high yield strength and a low yield ratio has been proposed. According to this technique, carbides including Ti, Nb, Mo, and V mainly containing ferrite in the vicinity of martensite are finely precipitated, and the deformation stress difference between adjacent structures can be reduced to improve the hole expanding workability.

  On the other hand, there is a technology that improves the hole expansion processability by not generating a soft ferrite phase.

For example, Non-Patent Document 1 discloses that steel containing C, Si, Mn, and Al is hot-rolled and cold-rolled, heated to a 950 ° C. austenite single phase region, and then subjected to austempering at around 400 ° C. Thus, a TRIP-type bainite steel sheet is disclosed in which lath-like retained austenite is generated in the matrix of bainitic ferrite.
JP-A-62-182225 Japanese Unexamined Patent Publication No. 06-145788 JP 2002-317249 A JP 2001-207234 A Japanese Patent Laid-Open No. 2002-69574 K. Sugimoto, etal .: ISIJ International, 40 (2000), No. 9, pp. 920-926

  However, since the techniques described in Patent Document 1 and Patent Document 2 have a composite structure containing a retained austenite phase, the hole-expansion workability is low, and further, precipitate formation that greatly affects the refinement of crystal grains Since no element is contained, there is a problem that the crystal grains are coarsened and the hole expanding workability is further deteriorated. Moreover, in the technique described in Patent Document 3, in order to increase the strength of the steel sheet, a precipitate forming element is contained, and improvement of hole expanding workability due to refinement of crystal grains can be expected. However, there is a problem that drastic improvement in hole expansion workability cannot be expected because the hard phase is formed in a band shape only by adding.

  Further, in the technique described in Patent Document 4 aimed at improving the hole expansion workability, NbC is precipitated and coarsened during winding after hot rolling or during a temperature rising process after cold rolling. A large amount of Nb is required to increase the hardness of the ferrite to a level sufficient to enhance the spreading workability. As a result, there is a problem that the ductility of the ferrite base is lowered and the ductility of the steel sheet itself is also lowered.

  Further, in the technique described in Patent Document 5, as in the technique described in Patent Document 4, coarsening of the precipitate is unavoidable, and further, the larger the volume ratio of the hard second phase, the more the precipitate is formed. There was a problem that the amount of element added increased and ductility decreased.

  Furthermore, in the technique described in Non-Patent Document 1, bainitic ferrite is used as a parent phase without requiring addition of a precipitate-forming element, and the hole expanding workability is excellent. However, it is necessary to perform annealing in the austenite single phase region at 950 ° C. after cold rolling, and in addition to the problem that it is difficult to perform such annealing treatment in a normal annealing furnace widely used in industry. Furthermore, there is a problem that the ductility of the bainitic ferrite phase itself is small and the ductility of the steel sheet is lowered.

  The object of the present invention is to solve the above-mentioned problems of the prior art and to provide a method for producing a high-strength cold-rolled steel sheet having excellent ductility and hole-expanding workability. As used herein, “high-strength cold-rolled steel sheet” refers to a cold-rolled steel sheet having a tensile strength of 780 MPa or more. “Excellent ductility” refers to the case where the product TS × El of tensile strength TS and total elongation El in the tensile test is 20000 MPa% or more, and “excellent hole expansion workability” The product TS × λ of the strength TS and the hole expansion rate λ in the hole expansion processing test is 27000 MPa% or more.

In order to achieve the above-mentioned problems, the present invention has intensively studied various factors that affect ductility and hole-expanding workability. As a result, in particular, using a material having a composition containing Nb, performing hot rolling, and further performing annealing treatment after performing cold rolling, heating in a temperature range above the Ac ₁ transformation point is slow heating, In addition, the cooling from the annealing soaking temperature is a thermal history that combines slow cooling, rapid cooling, and residence in the bainite generation temperature range, so that the hard second phase including the retained austenite phase is finely and uniformly dispersed. It has been found that it is possible to form a structure, and it is possible to achieve both the contradictory properties of ductility and hole expansion workability.

  First, the experimental results on which the present invention is based will be described.

  A cold-rolled sheet containing 0.15% C-0.20-0.40% Si-1.6% Mn-0.02% P-0.001% S-0.10-1.5% Al-0.0002% N-0.03% Nb at mass% is annealed. did.

The conditions for the annealing treatment were as follows. The temperature was raised to the Ac ₁ transformation point at a heating rate of 20 ° C./s, and further heated from the Ac ₁ transformation point to the annealing soaking temperature (850 ° C.) at a heating rate of 10 ° C./s. Thereafter, a soaking process was performed for 180 s at an annealing soaking temperature (850 ° C.). In addition, the residence time at 800 ° C. or higher including tempering, soaking and cooling was 192 s. After soaking, it was gradually cooled from the annealing soaking temperature to a slow cooling stop temperature (670 ° C.) at a cooling rate of 7 ° C./s. Subsequently, it was rapidly cooled to a quenching stop temperature (470 ° C.) at a cooling rate of 30 ° C./s. Thereafter, a residence treatment for 80 s was performed at the quenching stop temperature (470 ° C.). Subsequently, it cooled to room temperature with the cooling rate of 10 degree-C / s from the rapid cooling stop temperature (470 degreeC). The residence time in the temperature range of 350 to 500 ° C. was 93 s.

  The obtained cold-rolled annealed sheet was subjected to a tensile test and a hole expansion processing test, and a tensile strength TS, an elongation El, and a hole expansion ratio λ were obtained. The obtained results are shown in FIG. 1 in relation to TS × El, TS × λ and the amount of (Si + Al / 2). In addition, Si and Al in (Si + Al / 2) are the contents (mass%) of each element.

  FIG. 1 shows that by setting (Si + Al / 2) to 0.50 or more, TS × El can be set to 20000 MPa% or more, and TS × λ can satisfy 27000 MPa% or more. As a result of various studies based on this knowledge, the present inventors contain Nb, and after adjusting (Si + Al / 2) to 0.50 or more, annealing is gradually heated, residence in a soaking temperature range, It was found that the ductility and the hole-expanding workability can be improved by using a heat history that combines the subsequent slow cooling and rapid cooling, and the retention in the bainite generation temperature range.

In other words, by using a material containing Nb and having a composition of (Si + Al / 2) over a predetermined amount, the ferrite grains of the hot-rolled sheet are refined, and the number of austenite nucleation sites increases during annealing after cold rolling. Can be made. Furthermore, by performing gradual heating in the temperature range above the Ac ₁ transformation point, the austenite generation can be uniformly dispersed, and by gradually cooling from the annealing soaking temperature, the C concentration in the austenite is increased. At the same time, miniaturization and fragmentation of the austenite phase are promoted. By performing rapid cooling following gradual cooling and holding in the bainite formation temperature range, the formation of cementite and pearlite is suppressed to promote the formation of bainite, the C concentration in the austenite is further increased, and a fine austenite phase is formed. The hard second phase to be contained can be finely and uniformly dispersed. Thereby, it can be set as the cold rolled steel plate which has high ductility and high hole expansion workability.

The present invention has been completed based on the above findings and further studies. That is, the gist of the present invention is as follows.
(1) In mass%, C: 0.05 to 0.20%, Si: 0.2 to 1.5%, Mn: 0.5 to 3.0%, P: 0.05% or less, S: 0.01% or less, Al: 0.2 to 2.0%, N: 0.01 % Or less, Nb: 0.005 to 0.20%, and Si and Al are expressed by the following formula (1)
Si + 1/2 (Al) ≧ 0.50% (However, except for Si + Al: 1.0-2.5%) ……… (1)
(Where Si, Al: content of each element (mass%))
In order to satisfy the above, a hot rolling process to hot-roll the steel material consisting of the remaining Fe and inevitable impurities into a hot-rolled sheet, and cold-rolled the hot-rolled sheet to obtain a cold-rolled sheet After sequentially performing the cold rolling process, the cold rolled sheet is subjected to an annealing treatment to form a cold rolled annealed sheet,
The hot rolling is hot rolling in which the finish rolling finish temperature is not less than the Ar ₃ transformation point and the winding temperature is 400 to 650 ° C., and the annealing treatment is average heating from the Ac ₁ transformation point to the annealing soaking temperature. speed heating to annealing soaking temperature in the temperature range of 800 to 900 ° C. in the following heating rate 20 ° C. / s, allowed to stay between 82 ～300S in a temperature range of the 800 to 900 ° C., 600 to 700 ° C. from 800 ° C. Is gradually cooled at an average cooling rate of 1 to 10 ° C./s to the slow cooling stop temperature in the temperature range of 15 to 200 ° C./s from the slow cooling stop temperature to the quenching stop temperature in the temperature range of 350 to 500 ° C. A method for producing a high-strength cold-rolled steel sheet excellent in ductility and hole-opening workability, characterized in that it is rapidly cooled at an average cooling rate and retained for 30 to 300 seconds in the temperature range of 350 to 500 ° C.
(2) In (1), in addition to the above composition, mass% is further selected from Cr: 0.1 to 1.0%, Ni: 0.1 to 1.0%, Mo: 0.1 to 1.0%, B: 0.0005 to 0.0030% A method for producing a high-strength cold-rolled steel sheet, characterized by comprising a composition containing one or more kinds.
(3) In (1) or (2), in addition to the above composition, it further contains one or two kinds selected from Ti: 0.01 to 0.20% and V: 0.01 to 0.20% in mass%. A method for producing a high-strength cold-rolled steel sheet, characterized by comprising a composition.

  According to the present invention, a hard second phase containing a retained austenite phase can be finely and uniformly dispersed, and a high-strength cold-rolled steel sheet excellent in both ductility and hole expansion workability can be manufactured at low cost and easily. It has a remarkable industrial effect.

  First, the reasons for limiting the composition of the steel material used in the present invention will be described. Hereinafter, mass% in the composition is simply expressed as%.

C: 0.05-0.20%
C is an element that stabilizes austenite, and can be concentrated in austenite so that austenite remains stably up to room temperature. This retained austenite can be strain-induced transformed into hard martensite when strain is introduced by press working or the like, and the workability can be improved by suppressing strain concentration. In order to obtain such retained austenite, it is necessary to contain C: 0.05% or more. In order to promote the generation of hard low-temperature transformation phases such as bainite and martensite and to increase the strength of the steel sheet, C is preferably contained in an amount of 0.10% or more. On the other hand, if the content exceeds 0.20%, a large amount of cementite or pearlite is generated, and both ductility and hole-expanding workability deteriorate. Further, a large amount of C content causes deterioration of weldability. For this reason, in this invention, C was limited to 0.05 to 0.20% of range. In addition, Preferably it is 0.15% or less.

Si: 0.2-1.5%
Si is an element that stabilizes ferrite, and has the effect of concentrating C in austenite by promoting the formation of ferrite. Furthermore, Si also has the effect | action which reduces the amount of C solid solution in a ferrite, and can further raise C density | concentration in austenite. Such an effect becomes remarkable when the content is 0.2% or more. On the other hand, a large amount of Si exceeding 1.5% deteriorates weldability and promotes the formation of firelite on the surface of the slab during hot rolling, thereby promoting the generation of a so-called red scale surface pattern. Moreover, the Si oxide produced | generated on the surface at the time of cold rolling annealing deteriorates chemical conversion property, and induces non-plating at the time of hot dip galvanization. For these reasons, Si was limited to 0.2 to 1.5%. In addition, Preferably it is 0.4% or more. In particular, in the case of a steel plate or hot-dip galvanized steel plate that is required to have excellent surface properties, the content is preferably 0.7% or less.

Mn: 0.5-3.0%
Mn is an element that stabilizes austenite, lowers the martensite formation start temperature Ms point, and suppresses the transformation of austenite to martensite during cooling in the annealing process. In addition, it has the effect of delaying the formation of pearlite and allowing austenite to remain stably up to room temperature. Furthermore, Mn effectively contributes to the strengthening of the steel sheet as a solid solution strengthening element. Such an effect is recognized when the content is 0.5% or more, preferably 1.0% or more. On the other hand, a large amount of Mn containing more than 3.0% deteriorates the weldability of the steel sheet, suppresses bainite transformation, and suppresses C concentration in austenite with the progress of bainite transformation during cooling in the annealing process. For this reason, the desired amount of retained austenite cannot be secured. For these reasons, Mn is limited to 0.5 to 3.0%. In addition, Preferably it is 2.0% or less.

P: 0.05% or less P is an element effective for solid solution strengthening, and also has an action of promoting C concentration in austenite as a ferrite stabilizing element. It becomes remarkable by the above content. On the other hand, when P is contained in a large amount, it segregates at the grain boundary, lowers the ductility and toughness of the steel sheet, and degrades the weldability. For this reason, P is made 0.05% or less.

S: 0.01% or less S significantly lowers the hot ductility, induces hot cracking, and significantly deteriorates the surface properties. In addition, S hardly contributes to the strength, but also reduces the ductility and particularly the hole-opening workability by forming coarse MnS as an impurity element. For this reason, it is desirable to reduce S as much as possible, but it is acceptable up to 0.01%. For this reason, S was limited to 0.01% or less. In addition, from the viewpoint of improving ductility and hole expanding workability, S is preferably limited to 0.005% or less.

Al: 0.2-2.0%
Al is an element that stabilizes ferrite, and has the effect of concentrating C in austenite by promoting the formation of ferrite. Moreover, Al contributes effectively to high strength as a solid solution strengthening element. For this reason, in the present invention, Al is contained in an amount of 0.2% or more. On the other hand, a large amount of Al exceeding 2.0% raises the Ar ₃ transformation point of the steel and makes it difficult to finish hot rolling in the austenite single phase region. Therefore, Al is limited to the range of 0.2 to 2.0%. In addition, Preferably it is 0.3 to 1.0%.

N: 0.01% or less
If the N content exceeds 0.01%, slab cracking occurs during hot rolling, and the risk of generating surface defects increases. For this reason, N was limited to 0.01% or less. In addition, Preferably it is 0.004% or less.

Nb: 0.005-0.20%
Nb is one of the most important elements in the present invention, has an action of suppressing recrystallization of austenite processed during hot rolling, promotes ferrite transformation from unrecrystallized austenite, and hot-rolled sheet This has the effect of increasing the nucleation sites of recrystallized ferrite during the annealing process after cold rolling. Nb also forms fine carbonitrides and has the effect of suppressing the grain growth of recrystallized ferrite, further increasing the nucleation sites of austenite during heating during annealing, and as a result, residual austenite is reduced. Finely disperse. Moreover, Nb contributes to strength increase by refining the structure and forming carbonitride.

  Such an effect is recognized when Nb content is 0.005% or more. On the other hand, when the content is larger than 0.20%, carbonitrides cannot be completely dissolved at the time of reheating in the normal hot rolling process, coarse carbonitrides remain, and fine dispersion of residual austenite. And strength increase cannot be expected. In addition, when the slab is cooled from the continuous casting and then re-heated, and after the continuous casting, hot rolling is started as it is, even if the content exceeds 0.20%, the strength is further increased. Can not be expected, resulting in an increase in material costs. Furthermore, a large amount of Nb content not only causes an increase in load due to cumulative strain during hot rolling, but also increases the strength of the hot-rolled sheet, increases the rolling load load in cold rolling, and reduces the rolling efficiency. However, there is a problem that the manufacturing cost increases. For these reasons, Nb was limited to 0.005 to 0.20%. In addition, Preferably it is 0.10% or less.

In the present invention, the content of Si and Al is within the above range and the following formula (1)
Si + 1/2 (Al) ≥ 0.50% (except Si + Al: 1.0 to 2.5%)
Adjust to satisfy.

  The formula (1) is a limitation for generating stable retained austenite with respect to the initial processing strain to improve the ductility, and is an experimental formula obtained by various studies by the present inventors. Si and Al have the effect of concentrating C in austenite by stabilizing ferrite and promoting the formation of ferrite as described above, but further concentrating C in austenite and initial processing. In order to produce retained austenite that is stable against strain, improve workability, and obtain TS × El: 20000 MPa% or more, which is the ductility target of the present invention, Si, It is necessary to adjust the amount of Al.

  In the present invention, in addition to the basic composition described above, 1 selected from Cr: 0.1 to 1.0%, Ni: 0.1 to 1.0%, Mo: 0.1 to 1.0%, B: 0.0005 to 0.0030% as necessary. 1 type or 2 types or more and / or 1 type or 2 types chosen from Ti: 0.01-0.20% and V: 0.01-0.20% can be contained.

One or more selected from Cr: 0.1-1.0%, Ni: 0.1-1.0%, Mo: 0.1-1.0%, B: 0.0005-0.0030%
Cr, Ni, Mo, and B all have an action of increasing the strength, and can be selected and contained as necessary.

  Cr, Ni, and Mo have the effect of suppressing the bainite transformation, the effect of increasing the martensite fraction and increasing the strength of the steel sheet. Such an effect is recognized when Cr, Ni and Mo are contained in amounts of 0.1% or more. On the other hand, when Cr, Ni, and Mo each contain a large amount of more than 1.0%, the material cost increases and C concentration in austenite due to bainite transformation is suppressed, thereby inhibiting the formation of retained austenite. For this reason, Cr, Ni, and Mo are each preferably limited to a range of 0.1 to 1.0%. In addition, when using for the application which performs hot dip galvanization, since the oxide of Cr produced | generated on the surface induces non-plating, it is more preferable that Cr shall be 0.5% or less. Cr also has the effect of suppressing the formation of cementite in ferrite and tends to increase the solute C in ferrite. However, when the amount of solute C in ferrite is large, the yield strength is increased by age hardening of the steel sheet. As a result, the shape freezing property during press working tends to deteriorate. When shape freezing deterioration due to such solute C becomes a problem, the amount of Cr added to suppress the formation of cementite is decreased, and the amount of Si added to reduce the amount of solute C in ferrite is increased. Therefore, the ratio of Cr content to Si content: Cr / Si is preferably less than 0.3.

  B suppresses the transformation from austenite to ferrite, promotes the formation of a hard low-temperature transformation phase, and contributes to an increase in the strength of the steel sheet. Such an effect is recognized when the content is 0.0005% or more. On the other hand, an excessive content exceeding 0.0030% not only saturates the effect of improving hardenability but also remarkably increases the rolling load load during hot rolling due to the effect of suppressing recrystallization, and further suppresses ferrite transformation. Thus, C concentration in the austenite is suppressed, and the production of retained austenite is inhibited. For this reason, it is preferable to limit B to 0.0005 to 0.0030% of range.

One or two selected from Ti: 0.01-0.20%, V: 0.01-0.20%
Ti and V are elements that contribute to improvement of ductility and increase in strength of the steel sheet through the formation of fine carbonitrides, and can be selected and contained as necessary to improve the strength-ductility balance. .

  Ti increases the strength of the steel sheet through the formation of fine carbonitrides, and also contributes to the refinement of hot and cold rolled sheets. Such an effect is recognized when Ti: 0.01% or more is contained. On the other hand, even if it contains a large amount of Ti exceeding 0.20%, carbonitride cannot completely dissolve at the time of reheating in the normal hot rolling process, and coarse carbonitride remains, increasing strength. And refinement of the structure cannot be expected. In addition, after continuous casting, the slab is once cooled and then reheated, and after continuous casting, when hot rolling is started as it is, even if the content exceeds 0.20%, the strength is further increased. Can not be expected, resulting in an increase in material costs. For this reason, when it contains, it is preferable to limit Ti to the range of 0.01 to 0.20%. In addition, More preferably, it is 0.03-0.10%.

V increases the strength of the steel sheet through the formation of fine carbonitrides. V is an element that stabilizes ferrite, and in the temperature raising process during annealing after cold rolling, it raises the Ar ₁ transformation point of the steel and increases the frequency of nucleation of austenite generation at the time of temperature raising. The residual austenite is finely dispersed and the ductility of the steel sheet is improved. Such an effect is recognized when the content is V: 0.01% or more. On the other hand, even if the content exceeds 0.20%, a further increase in strength cannot be expected, resulting in an increase in material cost. For this reason, it is preferable to limit V to 0.01 to 0.20%. In addition, More preferably, it is 0.03-0.10%.

The balance other than the above components is Fe and inevitable impurities .

  It is preferable that the molten steel having the above composition is melted by a normal melting method such as a converter or an electric furnace, and used as a steel material such as a slab by a normal method such as a continuous casting method.

  The steel material is cast as it is without being cooled to room temperature after casting, or after being charged in a heating furnace without cooling to room temperature, and then cooled to room temperature or cooled to room temperature, preferably at 1100 to 1300 ° C. After heating, in accordance with a conventional method, after performing a hot rolling step to hot-roll and hot-rolled sheet, and a cold-rolling step to cold-roll the hot-rolled plate to form a cold-rolled plate, The cold-rolled sheet is annealed to form a cold-rolled sheet.

The heating temperature of the steel material is not particularly limited, but if it is less than 1100 ° C, the rolling load increases during the subsequent hot rolling, resulting in a reduction in rolling efficiency. Furthermore, it is difficult to finish the finish rolling at the Ar ₃ transformation point or higher. On the other hand, heating above 1300 ° C. increases oxidation loss and leads to generation of surface defects and a decrease in yield.

In the present invention, hot rolling is hot rolling at a finish rolling end temperature: Ar ₃ transformation point or higher and a winding temperature: 400 to 650 ° C.

Finishing rolling finish temperature: Ar ₃ transformation point or more If the finishing rolling finish temperature is less than the Ar ₃ transformation point, the rolling is performed in the two-phase region of ferrite and austenite, and the strain concentrates on the soft ferrite phase and is sufficient for the austenite. Is not introduced, the ferrite grains after transformation become coarse. Further, even in the ferrite phase in which strain is concentrated, the ferrite phase has a high stacking fault energy, so that recovery proceeds and ferrite recrystallization does not occur and coarse grains are expanded in the rolling direction. Such coarsening of ferrite grains in the hot-rolled sheet stage reduces the austenite nucleation sites during annealing after cold rolling, and the second phase including the retained austenite phase after annealing becomes coarse, which is good. It becomes impossible to obtain ductility and good hole expansion workability. For this reason, the finish rolling end temperature is limited to the Ar ₃ transformation point or higher. The upper limit of the finish rolling end temperature is preferably 950 ° C. or less from the viewpoint of austenite recrystallization and grain growth.

Winding temperature: 400-650 ° C
When the coiling temperature is less than 400 ° C., a low temperature transformation phase is generated and the hot rolled sheet becomes hard. For this reason, the rolling load in a cold rolling process increases, and rolling efficiency falls. Furthermore, when the coiling temperature is less than 400 ° C., the production of Nb carbonitride is suppressed, Nb is present as solid solution Nb in the hot-rolled sheet, and recrystallization is significantly suppressed during annealing after cold rolling. The Therefore, the austenite transformation from the non-recrystallized structure is promoted, and the austenite grains are coarsened, so that the second phase formed after annealing is also coarsened and the hole expandability is lowered. On the other hand, when the coiling temperature is higher than 650 ° C., coarse pearlite is generated, so in the annealing after cold rolling, the austenite transformation proceeds from the pearlite phase having a high C content, and therefore, the retained austenite phase after annealing is increased. The contained second phase is also coarsened, and good ductility and good hole expansion workability cannot be secured. In addition, when the coiling temperature is higher than 650 ° C., Nb carbonitrides are coarsened, the effect of suppressing the growth of recrystallized grains in annealing after cold rolling is reduced, austenite transformation is promoted from the coarse structure, and austenite Since the grains are coarsened, the second phase formed after annealing is also coarsened and the hole expandability is lowered. Moreover, coarse carbonitride does not contribute to an increase in strength. For this reason, the coiling temperature was limited to a range of 400 to 650 ° C.

  Next, the hot-rolled sheet is cold-rolled to obtain a cold-rolled sheet.

  The cold rolling reduction in cold rolling can be determined as appropriate according to the thickness of the cold rolled sheet. However, when the cold rolling reduction is less than 50%, the strain introduced into the cold-rolled sheet is small, the recrystallized grain size of ferrite during annealing is increased, and the hole expandability is lowered. Therefore, the cold rolling reduction is preferably 50% or more. On the other hand, even if cold rolling is performed exceeding 85%, the strain accumulation efficiency is lowered and the rolling load is increased. For this reason, the cold rolling reduction is preferably 85% or less.

In addition, before performing cold rolling, it is preferable to perform pickling at normal temperature and to remove the scale formed on the surface of the hot rolled steel sheet.
Next, the cold-rolled sheet is subjected to an annealing process to become a cold-rolled annealed sheet.

The annealing treatment is performed by heating to an annealing soaking temperature in a temperature range of 800 to 900 ° C. at a heating rate of 20 ° C./s or less from an Ac ₁ transformation point to an annealing soaking temperature, and the 800 to 900 ° C. It is allowed to stay for 82 to 300 s in the temperature range, gradually cooled at an average cooling rate of 1 to 10 ° C./s from the temperature range of 800 ° C. to 600 ° C. to 700 ° C., and 350 to 350 ° C. from the slow cooling stop temperature. A rapid cooling is performed at an average cooling rate of 15 to 200 ° C./s to a quenching stop temperature in a temperature range of 500 ° C., and a treatment is performed in the temperature range of 350 to 500 ° C. for 30 to 300 s.

Average heating rate between Ac ₁ transformation point and annealing soaking temperature: 20 ° C./s or less Heating during annealing is slow heating with an average heating rate of 20 ° C./s or less. The heating rate from the Ac ₁ transformation point to the annealing soaking temperature is one of the most important requirements of the present invention. By gradually heating in the temperature range above the Ac ₁ transformation point, fine austenite can be dispersed and transformed, and further C enrichment in the austenite can be promoted. This is because the austenite is not coarsened in the state where the austenite fraction is small, just above the Ac ₁ transformation point, and therefore the C-decreased region around the austenite also becomes small, resulting in many sites where C can be concentrated and austenite remains. This is because the number of nucleation sites increases. After that, by gradually heating to the annealing soaking temperature, the generation of coarse austenite during heating and soaking is suppressed, and the austenite fraction is increased while maintaining the dispersed state of austenite. C concentration can be promoted, a predetermined amount or more of dispersed austenite phase can be stably secured, and ductility and hole-expanding workability are improved.

Furthermore, it is desirable that the recrystallization of the ferrite is completed when the Ac ₁ transformation point is reached, but even if unrecrystallized ferrite remains, it is gradually heated between the Ac ₁ transformation point and the annealing soaking temperature. By performing, it is possible to suppress the formation of band-like austenite resulting from the austenite transformation of the unrecrystallized portion at the same time.

In order to obtain such an effect, during annealing, the heating rate between the Ac ₁ transformation point and the annealing soaking temperature is limited to an average heating rate of 20 ° C./s or less. In addition, Preferably, it is 10 degrees C / s or less. The lower limit of the heating rate is not particularly defined, but an extremely small heating rate is preferably 5 ° C./s or more because it causes a decrease in operation efficiency. Further, although the heating rate from room temperature to the Ac ₁ transformation point is not particularly defined, it is preferable to raise the temperature range at 20 ° C./s or less in order to promote recrystallization of ferrite.

Annealing temperature: 800-900 ℃
The annealing soaking temperature is set to a two-phase region of austenite + ferrite to promote an increase in the austenite fraction and C concentration in the austenite, to stably secure a retained austenite phase of a predetermined amount or more and improve ductility. , 800 ° C. or higher is preferable. On the other hand, when the annealing soaking temperature exceeds 900 ° C., the austenite fraction becomes too large, the C concentration in the austenite becomes small, and it becomes difficult to leave the austenite to room temperature in the subsequent cooling process. Austenite grain size is also increased. For this reason, annealing soaking temperature shall be the temperature of the temperature range of 800-900 degreeC. Here, the annealing soaking temperature is the highest temperature achieved during annealing.

800-900 residence time in the temperature range of ℃: 82 ~300s
If the residence time staying in the temperature range of 800 to 900 ° C during annealing (heating and soaking) and cooling (annealing soaking temperature range) is less than 82 s, C concentration in the austenite becomes insufficient, and the subsequent 800 ° C It becomes difficult to leave austenite to room temperature in the following cooling process. On the other hand, if the residence time exceeds 300 s, the austenite becomes coarse and the operation efficiency decreases. For this reason, the residence time in the temperature range of 800 to 900 ° C. is limited to the range of 82 to 300 s. As a result, a dispersed retained austenite phase of a predetermined amount or more can be stably secured, and ductility and hole expansion workability are improved. Here, the retention time may be ensured (held) at the annealing soaking temperature, or may be secured by adjusting the heating and cooling rates.

Average cooling rate from 800 ° C to the annealing stop temperature in the temperature range of 600-700 ° C: 1-10 ° C / s
The cooling at 800 ° C. or lower is a slow cooling with an average cooling rate of 1 to 10 ° C./s, and the cooling is stopped at a slow cooling stop temperature in a temperature range of 600 to 700 ° C. (slow cooling stop temperature range). By gradually cooling the temperature range from 800 ° C. to the slow cooling stop temperature, it promotes ferrite formation during cooling, promotes C concentration in the austenite, and finely divides the remaining austenite phase. This improves ductility and hole expandability. In order to obtain such an effect, the cooling rate is limited to 10 ° C./s or less. On the other hand, when the cooling rate is less than 1 ° C./s, ferrite transformed from austenite grows and coarsens, and austenite remains without being finely divided. On the other hand, when the cooling rate is less than 1 ° C./s, the operation efficiency is lowered. For this reason, the cooling from 800 ° C. to the gradual cooling stop temperature was gradual cooling at an average cooling rate of 1 to 10 ° C./s. In addition, More preferably, it is 5 degrees C / s or less. Thereby, a retained austenite phase of a predetermined amount or more can be stably secured, and both ductility and hole expansion workability are improved.

  On the other hand, when the slow cooling stop temperature is higher than 700 ° C., the ferrite fraction is small, and C concentration in the austenite becomes insufficient. On the other hand, when the slow cooling stop temperature is less than 600 ° C., cementite and pearlite are generated, and the concentration in austenite becomes insufficient. For this reason, in cooling after soaking at the annealing soaking temperature, the slow cooling stop temperature is limited to a temperature in the temperature range of 600 to 700 ° C.

Average cooling rate from the slow cooling stop temperature to the rapid cooling stop temperature in the temperature range of 350 to 500 ° C: 15 to 200 ° C / s
Then, the cooling is rapidly performed from the slow cooling stop temperature to the rapid cooling stop temperature in the temperature range of 350 to 500 ° C (rapid cooling stop temperature range) at an average cooling rate of 15 to 200 ° C / s. By performing rapid cooling to a temperature range of 350 to 500 ° C., which is a bainite generation temperature range, generation of cementite and pearlite from austenite during cooling can be suppressed, and driving force for bainite transformation can be increased. As a result, C enrichment in the austenite phase is promoted, and austenite can remain stably up to room temperature. Thereby, the retained austenite phase more than predetermined amount can be secured stably, and ductility improves. In order to obtain the effect as described above, it is necessary to cool to a quenching stop temperature in a temperature range of 350 to 500 ° C. at a cooling rate of 15 ° C./s or more. On the other hand, even if the average cooling rate is made higher than 200 ° C./s, not only a further effect is observed, but also in the operation, the variation in the cooling stop temperature becomes large, so that the variation in the material also becomes large. For this reason, in the cooling from the slow cooling stop temperature to the rapid cooling stop temperature, the cooling rate is limited to an average cooling rate of 15 to 200 ° C./s.

  On the other hand, when the quenching stop temperature exceeds 500 ° C., the bainite transformation does not proceed and cementite and pearlite are generated, so that ductility and hole expandability are deteriorated. On the other hand, if the quenching stop temperature is less than 350 ° C., the martensitic transformation proceeds and it becomes difficult to concentrate C in the austenite. For this reason, the quenching stop temperature was limited to a temperature in the temperature range of 350 to 500 ° C.

Residence time in the temperature range of 350 to 500 ° C: 30 to 300 s
By retaining in a temperature range of 350 to 500 ° C. (quenching stop temperature range), bainite transformation can be advanced and the C concentration in austenite can be increased. Thereby, the retained austenite phase more than predetermined amount can be secured stably, and ductility improves. In order to obtain such an effect, it is preferable to retain for 30 seconds or more in this temperature range. On the other hand, if it is retained for longer than 300 s, ferrite and carbide precipitate from austenite, the austenite fraction decreases, and the C concentration in the austenite also decreases. For this reason, the residence time in the 350-500 degreeC temperature range was limited to the range of 30-300 s.

  After the above-mentioned residence time in the temperature range of 350 to 500 ° C. (quenching stop temperature range), cooling is performed. Cooling after residence in the rapid cooling stop temperature range may be either natural cooling or rapid cooling. In addition, when performing hot dip galvanization after an annealing process and also performing an alloying process, you may reheat at 500 degreeC or more for an alloying process.

  The cold-rolled steel sheet obtained by the manufacturing method described above has the above-described composition, and further, in area ratio, 50 to 90% ferrite phase, 5 to 30% bainite phase, and 5% or more residual austenite phase. In addition, it generally has a composite structure consisting of martensite phase of 30% or less, tensile strength: high strength of 780MPa or more, high ductility of 20000MPa or more in TS × El, and 27000MPa% in TS × λ A high-strength cold-rolled steel sheet having the above-described high-hole expanding workability and excellent in both ductility and hole-expanding workability is obtained.

  Hereinafter, based on an Example, this invention is demonstrated further. Needless to say, the present invention is not limited to these examples.

  Steel having the composition shown in Table 1 was melted in a vacuum melting furnace to form a small steel ingot (100 kg).

Subsequently, these steel ingots were heated to 1250 ° C. (holding: 1 h) and hot-rolled to obtain hot rolled sheets having a thickness of 3.5 mm. The finish temperature for hot finish rolling was 870 ° C., which is not lower than the Ar ₃ transformation point. Immediately after the end of hot rolling, cooling was started at a cooling rate of 20 ° C./s, and holding was performed for 1 hour at 550 ° C., which is a winding equivalent temperature. After that, it was cooled to room temperature by slow cooling at 50 ° C./h to simulate winding with an actual machine (equivalent to a winding temperature of 550 ° C.). The hot-rolled sheet cooled to room temperature was scaled with hydrochloric acid (pickling), and then cold-rolled to a sheet thickness of 1.0 mm at a cold reduction rate of 71% to obtain a cold-rolled sheet.

  These cold-rolled sheets were subjected to annealing treatment under the conditions shown in Table 2 to obtain cold-rolled annealed sheets. The annealing treatment was based on the following conditions.

(1) The temperature was raised to the Ac ₁ transformation point at a heating rate of 20 ° C./s.

(2) Further, heating was performed from the Ac ₁ transformation point to an annealing soaking temperature (T1 = 850 ° C.) at a heating rate (HR) of 10 ° C./s. Thereafter, a soaking treatment was performed at an annealing soaking temperature (850 ° C.) for 180 seconds (soaking time t0).

  (3) After soaking, the steel was gradually cooled from the annealing soaking temperature to the slow cooling stop temperature (T2 = 670 ° C.) at a cooling rate (CR1) of 7 ° C./s. In addition, the residence time (t1) at 800 ° C. or higher including tempering, soaking and cooling was 192 s. The cooling rate from 800 ° C. to T2 was 7 ° C./s.

  (4) Subsequently, rapid cooling was performed to the quenching stop temperature (T3 = 470 ° C.) at a cooling rate (CR2) of 30 ° C./s. Thereafter, a residence treatment for 80 s was performed at the quenching stop temperature (470 ° C.).

  (5) Then, it was cooled from the rapid cooling stop temperature (470 ° C.) to room temperature at a cooling rate of 10 ° C./s. The residence time (t2) in the temperature range of 350 to 500 ° C. (quenching stop temperature range) was 93 s.

The Ar ₃ transformation point and Ac ₁ transformation point were determined from the thermal expansion curves during cooling and temperature elevation.

Based on these conditions, the heating rate (HR) from the Ac ₁ transformation point to the annealing soaking temperature, the annealing soaking temperature (T1), the soaking time at the annealing soaking temperature (t0), and the temperature of 800 to 900 ° C. Residence time (t1) in region, cooling rate of slow cooling from annealing soaking temperature (CR1), slow cooling stop temperature (T2), cooling rate of rapid cooling (CR2), quenching stop temperature (T3), 350-500 The residence time (t2) in the temperature range of ° C. was variously changed.

  The resulting cold-rolled annealed plate was examined for structure observation, tensile properties, and hole-expansion workability. The test method was as follows.

(A) Microstructure investigation From the obtained cold-rolled annealed plate, a test piece is collected, and a cross section parallel to the rolling direction (cross section in the L direction) is corroded with a nital liquid and the structure is imaged using a scanning electron microscope. Using an image analysis apparatus, the tissue fraction of each phase was determined. In addition, about each cold-rolled annealing board, three visual fields were observed by 1000 time magnification, the fraction of each phase in each visual field was calculated | required, and the average value of each visual field was made into the structure | tissue fraction in each cold-rolled annealing board. In addition, the martensite phase and the retained austenite phase were further tempered at 200 ° C. for 2 hours, and the structure was observed (magnification: 5000 times) for distinction. The phase in which the carbide precipitated after the tempering heat treatment was martensite, and the phase in which no carbide precipitation was observed was the retained austenite phase.

(B) Tensile properties After subjecting the obtained cold-rolled annealed sheet to temper rolling with an elongation of 0.5%, the direction perpendicular to the rolling direction (C direction) is the tensile direction. A JIS No. 5 tensile test piece was cut out in accordance with the regulations, and a tensile test up to fracture was performed in accordance with the regulations of JIS Z 2241 to determine the yield stress YS, tensile strength TS, and total elongation El.

(C) Hole expansion workability test A test piece (size: 100 × 100 mm) was taken from the obtained cold-rolled annealed sheet, and an initial hole with a diameter of d ₀ (= 10 mmφ) was punched with a punch in the center. this initial hole, inserting a conical punch of apex angle 60 °, to expand, seeking hole diameter d _f when cracks passing through the thickness occurs, the following formula
λ (%) = {(d _f −d ₀ ) / d ₀ } × 100
Was used to determine the hole expansion ratio λ.

  The obtained results are shown in Tables 4 and 5.

  Each of the inventive examples has a high tensile strength of TS: 780 MPa or more, an excellent ductility of TS × El: 20000 MPa% or more, and an excellent hole expanding workability of TS × λ: 27000 MPa% or more. It is a high-strength cold-rolled steel sheet with excellent ductility and hole-expansion workability. On the other hand, the comparative example out of the scope of the present invention is that the strength is low, the ductility is lowered, the hole expansion rate is decreased, or all of these are decreased, and the ductility and the hole expansion processing are performed. Both properties are not excellent.

It is a graph which shows the relationship between elongation El, hole expansion ratio (lambda), and the quantity of (Si + Al / 2).

Claims (3)

mass%
C: 0.05 to 0.20%, Si: 0.2 to 1.5%,
Mn: 0.5 to 3.0%, P: 0.05% or less,
S: 0.01% or less, Al: 0.2-2.0%,
N: 0.01% or less, Nb: 0.005-0.20%
And a steel material containing Si and Al so as to satisfy the following formula (1), and a hot-rolling step of subjecting the steel material consisting of the remaining Fe and inevitable impurities to hot rolling to form a hot-rolled sheet, After performing the cold rolling step to cold-rolled sheet by subjecting the hot-rolled sheet to cold rolling, the cold-rolled sheet is subjected to annealing treatment to make a cold-rolled annealed sheet,
The hot rolling is hot rolling in which the finish rolling finish temperature is Ar ₃ transformation point or higher and the winding temperature is 400 to 650 ° C.,
The annealing treatment is heated to an annealing soaking temperature in a temperature range of 800 to 900 ° C. at a heating rate of 20 ° C./s or less at an average heating rate from Ac ₁ transformation point to annealing soaking temperature, and the 800 to 900 ° C. In the temperature range of 82 to 300 s, gradually cooled at an average cooling rate of 1 to 10 ° C./s from the 800 ° C. to the 600 ° C. to 700 ° C. slow cooling stop temperature, and then from the slow cooling stop temperature to 350 ° C. Ductility characterized by rapid cooling at an average cooling rate of 15 to 200 ° C./s to a quenching stop temperature in a temperature range of ˜500 ° C. and retaining for 30 to 300 s in the temperature range of 350 to 500 ° C. And manufacturing method of high-strength cold-rolled steel sheet with excellent hole-opening workability.
Record
Si + 1/2 (Al) ≧ 0.50% (However, except for Si + Al: 1.0-2.5%) ……… (1)
Here, Si, Al: Content of each element (mass%)   In addition to the above composition, in mass%, one or more selected from Cr: 0.1-1.0%, Ni: 0.1-1.0%, Mo: 0.1-1.0%, B: 0.0005-0.0030% The method for producing a high-strength cold-rolled steel sheet according to claim 1, wherein the composition contains 2. The composition according to claim 1, wherein the composition further comprises one or two kinds selected from Ti: 0.01 to 0.20% and V: 0.01 to 0.20% in mass% in addition to the composition. The manufacturing method of the high intensity | strength cold-rolled steel plate of 2.

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