JP2013227635A - High strength cold rolled steel sheet, high strength galvanized steel sheet, method for manufacturing high strength cold rolled steel sheet, and method for manufacturing high strength galvanized steel sheet - Google Patents

Abstract

PROBLEM TO BE SOLVED: To manufacture a high strength cold rolled steel sheet and a high strength galvanized steel sheet which are inhibited from generating yield point elongation while having high ductility in a strength level of 440 MPa or more and less than 590 MPa without remarkable increase of manufacturing cost and alloy cost.SOLUTION: A high strength cold rolled steel sheet has a predetermined composition, a structure containing ferrite of, by a volume fraction, 85% or more, and a second phase of, by a volume fraction, 15% or less, and a tensile strength of 440 MPa or more and less than 590 MPa, wherein the second phase contains martensite of, by a volume fraction, 1% or more and less than 8%, and retained austenite of, by a volume fraction, 3% or more and less than 10%, a volume fraction of bainite being 1% or more and 5% or less, a volume fraction of pearlite being 0% or more and less than 2%, and a total volume fraction of the bainite and the pearlite being below the martensite volume fraction.

Description

  The present invention relates to a high-strength cold-rolled steel sheet, a high-strength galvanized steel sheet, a method for producing a high-strength cold-rolled steel sheet, and a method for producing a high-strength galvanized steel sheet.

In recent years, in the automobile manufacturing field, from the viewpoint of improving fuel efficiency by reducing the weight of a vehicle body, it has been strongly required to reduce the amount of steel used by reducing the thickness of a steel plate. Similarly, in the field of home appliance manufacturing, from the viewpoint of reducing CO ₂ emissions and reducing costs in the manufacturing cycle of products, it is required to reduce the amount of steel used by thinning the steel sheet. However, if the steel sheet is simply thinned in order to reduce the amount of steel used, the strength is reduced, and the steel sheet is easily deformed when subjected to an external force. Therefore, it is necessary to apply a higher-strength steel sheet. On the other hand, increasing the strength of a steel sheet has the disadvantage of reducing formability.

  From such a background, especially steel plates applied to press parts such as outer panel of automobiles and housings of home appliances are required to have very high formability, so that high-strength steel plates have not been applied. For example, for outer panels such as automobile hoods, doors, and back doors, conventionally, ferritic single phase steel having a tensile strength (TS) of 270 to 340 MPa and excellent formability is used. There are steels obtained by solid solution strengthening such ferritic single-phase steels with Mn, Si, P or the like, or steels with precipitation strengthening with Ti, Nb carbides, etc., but when this steel is strengthened to the TS: 440 MPa class, The ductility is remarkably lowered when the elongation (El) is about 30% and the uniform elongation (U.El) is less than 20%. For this reason, at the time of press molding of a part, necking may occur in the design surface and the surface appearance may be impaired, or cracking may occur and press molding itself may be difficult. In addition, a steel sheet with a pearlite phase formed in the ferrite phase to increase the strength has a high yield point and poor shape freezing property, and further has a yield point elongation, so that stretcher strain is generated and the appearance quality is significantly impaired.

  Therefore, in recent years, as an approach for improving the ductility while increasing the strength of a steel sheet, a dual-phase (DP) steel sheet in which a hard martensite phase is dispersed in a soft ferrite phase has been proposed. The DP steel sheet has a high yield strength and a low yield strength. Therefore, the DP steel sheet has excellent shape freezing properties and can suppress the occurrence of yield point elongation. Furthermore, since the DP steel sheet is excellent in work hardening characteristics as compared with the solid solution strengthened steel, it exhibits high uniform elongation (see Patent Document 1). However, since the DP steel sheet is also inferior to the elongation of the 270 to 340 MPa class steel sheet in terms of the absolute value of elongation, necking and cracking easily occur depending on the part shape.

  In addition, as one of the approaches to obtain superior ductility, transformation-induced plasticity (TRansformation-Induced Plasticity), in which retained austenite is generated in the ferrite phase and the austenite transforms into martensite during deformation, exhibits high ductility. A TRIP steel sheet using the TRIP effect has been proposed. A well-known TRIP steel sheet is a Si-added TRIP steel sheet in which a large amount of Si is added in order to leave residual austenite to room temperature and C distribution between ferrite and austenite is promoted (see Patent Document 2). ). In this TRIP steel plate, a TS × El balance higher than that of a ferritic single-phase steel or DP steel plate is obtained, and very high ductility is exhibited when compared at the same strength level.

  However, since Si is a very strong solid solution strengthening element, if a large amount of Si is added to ensure the retained austenite fraction, the tensile strength becomes 590 MPa or more, and the absolute value of total elongation and uniform elongation is low. I have to be. In fact, in the conventional 590 MPa grade Si-added TRIP steel sheet, the total elongation is about 35% and the uniform elongation is less than 24%. Furthermore, since Si easily forms an oxide film on the surface of the steel sheet during slab heating, hot rolling or annealing, it causes surface defects such as scale residue, plating unevenness, and non-plating.

  Thus, a TRIP steel sheet has been proposed in which, like Si, the concentration of C into austenite is promoted, and Al is added which has less influence on strength increase and plating deterioration than Si. For example, in Patent Document 3, in a steel sheet in which the Si amount is reduced to an Al amount of 1.5 to 2% and a certain amount of retained austenite is generated, tensile strength (TS): 440 to 490 MPa, elongation (El): 36 A method for producing a steel sheet excellent in ductility and plating adhesion having ˜39% is disclosed. Patent Document 4 also describes a method for producing a highly ductile hot-dip galvanized steel sheet that utilizes retained austenite by reducing the amount of Si and adding Al.

  Patent Document 5 proposes a method for producing a hot-dip galvanized steel sheet based on a cold-rolled steel sheet having a tensile strength (TS) of 440 to 490 MPa and excellent ductility in a high Mn and high Al component steel containing Si. Has been. In addition, Patent Document 6 discloses that a high-alloy steel has a temperature of 350 to 500 ° C. in a hot rolling process without performing a pre-heat treatment other than a manufacturing method in which a pre-heat treatment is performed on a high Al component steel and then plated with a CGL (continuous hot dip galvanizing line). A steel sheet excellent in TS × El balance in which a hot rolled structure having 10% or more of a low temperature transformation phase and having 80% or more of bainite as a low temperature transformation phase is formed by performing low temperature winding, and then plated with CGL. The manufacturing method of this is proposed. Patent Document 7 discloses a method for producing a steel sheet having a high El in high Al component steel.

Japanese Patent No. 4207738 JP-A-5-255799 JP 2001-355041 A Japanese Patent No. 4333352 JP 2000-256789 A JP 2004-256836 A Japanese Patent No. 3569307

  However, in order to produce the steel sheet described in Patent Document 3, it is indispensable to perform a two-phase region annealing at 800 ° C. for about 1 hour before final annealing and to distribute Mn in advance, and further annealing + plating Also in the process, it is necessary to cool after soaking at 800 ° C. for 60 seconds and anneal for a long time of 10 minutes or less after immersion in a plating bath at 440 ° C. For this reason, it is difficult to apply the manufacturing method described in Patent Document 3 to a normal CGL line that does not have an annealing line after plating immersion, and furthermore, since the pre-heat treatment before the final annealing is indispensable, the manufacturing cost is greatly increased. To increase. In addition, as described in the examples, all the inventive steels have elongation at yield exceeding 2%, which causes stretcher strain during pressing and deteriorates the appearance quality. Similarly, the production method described in Patent Document 4 also requires annealing at 750 ° C. or higher and tempering at 250 to 550 ° C. before final annealing by CGL.

  Further, the steel sheet described in Patent Document 5 has a problem that the alloying temperature is much higher than that of a conventional method, rapid heating is required, and the manufacturing cost is significantly higher than that of existing CGL equipment. Further, the technology described in Patent Document 6 employs a component system containing a certain amount of Si, and it is difficult to say that the plating property is sufficiently excellent. Moreover, in the manufacturing method described in Patent Document 7, a cooling capacity of 80 ° C./s or more is essentially essential at the secondary cooling rate. In addition, when the present inventors investigated to the low cold speed range, it became clear that some of the steel plates described in the examples were decomposed into pearlite in the second phase and the characteristics were greatly deteriorated.

  In this way, at a strength level of 440 MPa or more and less than 590 MPa, high-strength cold-rolled steel sheet and high-strength steel that can be produced without significant increase in production cost and alloy cost while suppressing the occurrence of yield point elongation while having high ductility. The galvanized steel sheet and its manufacturing method are not sufficiently provided.

  The present invention has been made in view of the above problems, and its purpose is to suppress the occurrence of elongation at the yield point while having high ductility at a tensile strength level of 440 MPa or more and less than 590 MPa. An object of the present invention is to provide a high-strength cold-rolled steel sheet, a high-strength galvanized steel sheet, a method for producing a high-strength cold-rolled steel sheet, and a method for producing a high-strength galvanized steel sheet.

  As a result of intensive studies to solve the above problems, the inventors of the present invention appropriately controlled the alloy addition amount and the manufacturing method, and appropriately formed a microstructure into a microstructure, thereby setting the tensile strength level to 440 MPa. It was discovered that the yield point elongation (YPEL) of the steel sheet can be suppressed while obtaining excellent ductility while maintaining the pressure below 590 MPa.

In order to solve the above-described problems and achieve the object, the high-strength cold-rolled steel sheet according to the present invention has a component composition of mass%, C: 0.08% or more and 0.25% or less, Si: 0.50. %: Mn: 0.7% or more and less than 2.0%, P: 0.1% or less, S: 0.01% or less, Cr: 1.0% or less (including 0%), Al: 0. 40% or more and 2.00% or less, B: 0.0050% or less (including 0%), the total content of Si and Al is 0.7 to 2%, and is shown in the following formula (1) Mn equivalent Mn _eq is 0.8 ≦ Mn _eq ≦ 2.0 and Mn _eq + 1.3 × (Si content + Al content) ≧ 2.6, with the balance being iron and inevitable impurities The structure has a ferrite with a volume fraction of 85% or more and a second phase with a volume fraction of 15% or less, and the second phase has a volume fraction of 1% or more and 8%. Containing less than 10% martensite and less than 10% residual austenite, bainite has a volume fraction of 1% to 5%, pearlite has a volume fraction of 0% to less than 2%, bainite and pearlite The total volume fraction is less than or equal to the martensite volume fraction, and the tensile strength is 440 MPa or more and less than 590 MPa.

  The high-strength cold-rolled steel sheet according to the present invention is, in the above-described invention, as a component composition, in mass%, Ti: 0.02% or less, Mo: less than 0.1%, V: 0.02% or less, Nb: 0 0.02% or less, Ni: 0.2% or less, and Cu: less than 0.1%, further containing one or more elements.

  The high-strength cold-rolled steel sheet according to the present invention is, in the above-described invention, as a component composition in mass%, Sb: 0.2% or less, Sn: 0.2% or less, Ca: 0.01% or less, REM: 0 Further, one or more elements out of 0.01% or less are further contained.

In the high-strength cold-rolled steel sheet according to the present invention, in the above-described invention, a value S / L ² obtained by dividing the volume fraction S of the second phase by the square value of the peripheral length L of the second phase is less than 0.00015. It is characterized by that.

  In order to solve the above-described problems and achieve the object, the high-strength galvanized steel sheet according to the present invention is characterized by including a zinc-based plating film on the surface of the high-strength cold-rolled steel sheet according to the present invention.

In order to solve the above-mentioned problems and achieve the object, a method for producing a high-strength cold-rolled steel sheet according to the present invention includes hot rolling and cold rolling on a steel slab having the above component composition, followed by continuous annealing. Upon application, after heating to 750 ° C. to 950 ° C. and holding for 20 seconds or more, the average of 5 ° _{C./s to} 40 ° _C./s to the cooling limit temperature T _crit (° C.) defined by the following formula (2) After cooling at a cooling rate, it is cooled to 230 ° C. or higher and lower than 410 ° C. at an average cooling rate of 20 ° C./s or higher, reheated to 410 ° C. or higher and 500 ° C. or lower, and held for 10 seconds or longer and 180 seconds or shorter.

  In order to solve the above-described problems and achieve the object, a method for producing a high-strength galvanized steel sheet according to the present invention is a method for producing a high-strength cold-rolled steel sheet according to the present invention, wherein zinc-based plating is performed in the course of continuous annealing. By applying, a zinc-based plating film is formed on the surface of the high-strength cold-rolled steel sheet.

  According to the present invention, high-strength cold-rolled steel sheet that suppresses the occurrence of yield point elongation while having high ductility at a tensile strength level of 440 MPa or more and less than 590 MPa, and can be produced without significant increase in production cost or alloy cost, A manufacturing method of a high-strength galvanized steel sheet, a high-strength cold-rolled steel sheet, and a manufacturing method of a high-strength galvanized steel sheet can be provided.

FIG. 1 is a diagram showing the results of investigating the relationship between the Mn equivalent of steel, the contents of Al and Si and materials, and the presence or absence of pearlite generation during continuous cooling. FIG. 2 is a diagram showing an example of a change in yield point elongation when the structure of steel is changed. FIG. 3 is a diagram showing the results of investigating the temperature at which pearlite formation occurs when steel plates having different Mn equivalents and different total contents of Al and Si are continuously cooled at various average cooling rates. FIG. 4 is a diagram illustrating an example of the relationship between the primary cooling stop temperature, total elongation, and uniform elongation. FIG. 5 is a diagram illustrating an example of the relationship between the secondary cooling stop temperature, S / L ² , total elongation, and uniform elongation.

  Hereinafter, the high-strength cold-rolled steel sheet and the high-strength galvanized steel sheet according to the present invention will be described in detail according to their component composition, structure, and manufacturing method.

(Component composition)
First, the component composition of the high-strength cold-rolled steel sheet and the high-strength galvanized steel sheet according to the present invention will be described. In the following, “%” of the component amount means “% by mass” unless otherwise specified.

(C content)
C (carbon) is an inexpensive and very strong austenite stabilizing element, and is an extremely important element for remaining austenite to room temperature. In the heat treatment process described later, C is discharged from ferrite and bainitic ferrite to austenite to stabilize austenite. In order to improve ductility, 3% or more of austenite in which C is sufficiently concentrated is necessary. If C is less than 0.08%, the final retained austenite amount is less than 3%, and the ductility is not sufficiently improved. As the amount of C increases, the amount of retained austenite produced and its stability increase. However, if the amount of C exceeds 0.25%, the second phase fraction increases too much and TS increases and ductility decreases. Furthermore, weldability deteriorates. Therefore, the C content is set to 0.08% or more and 0.25% or less. In order to obtain a steel sheet with higher ductility, the upper limit of the C content is preferably less than 0.20%.

(Si content)
Since Si (silicon) suppresses cementite precipitation from austenite, it is an extremely effective element for promoting C concentration of austenite. However, since Si has a very high solid solution strengthening ability, if it is contained in a large amount, it becomes difficult to maintain the tensile strength below 590 MPa. In addition, since Si has a high affinity with oxygen, it is easy to form an oxide film on the surface of the steel sheet, and even a small amount causes a scale residue during hot rolling and non-plating due to the formation of an oxide film during continuous hot dip galvanization. Accordingly, the Si content is desirably as small as possible, and is less than 0.50%. From the viewpoint of improving the plating quality and chemical conversion property, the Si content is preferably less than 0.20%, and more preferably less than 0.05%. In order to obtain particularly excellent plating quality, the Si content is preferably less than 0.03%.

(Mn content)
Mn (manganese) is an important element for suppressing the transformation of austenite to pearlite or bainite. If the Mn content is less than 0.7%, it becomes easy to decompose from austenite to pearlite or bainite during cooling after annealing, and it becomes difficult to secure stable retained austenite. On the other hand, if the Mn content is 2.0% or more, the ferrite transformation in the primary cooling described later and the bainite transformation in the secondary cooling are delayed and the martensite fraction increases, so the tensile strength becomes 590 MPa or more. Ductility decreases. Therefore, the Mn content is 0.7% or more and less than 2.0%, preferably less than 1.7%, more preferably less than 1.6%.

(P content)
P (phosphorus), like B, has the effect of suppressing pearlite transformation even when added in a small amount. However, P is a very strong solid solution strengthening element. If it is excessively contained, the strength increases more than necessary. Further, P invites generation of surface defects due to uneven plating or segregation due to a delay in alloying. Therefore, the P content is 0.1% or less, more preferably 0.05% or less, and still more preferably 0.03% or less.

(S content)
By including an appropriate amount of S (sulfur), it is possible to improve the peelability of the primary scale and improve the final plating appearance quality of the steel sheet. It is preferable to contain 0.001% or more. However, when S is present in a large amount, the hot ductility of the steel is reduced, and the surface quality is deteriorated by cracking on the surface of the steel sheet during hot rolling. Furthermore, S forms coarse MnS. As a result, the ductility of the steel sheet is reduced. For this reason, content of S shall be 0.01% or less.

(Cr content)
Cr (chromium) is an optional additive element. However, since the solid solution strengthening of Cr is small, the effect of increasing the Mn equivalent is large, and therefore Cr is preferably added in an amount of 0.1% or more in order to suppress pearlite transformation in the cooling process. On the other hand, excessive addition of Cr not only increases the cost, but also significantly delays the ferrite / bainite transformation and makes it difficult to obtain retained austenite. Therefore, the Cr content is 1.0% or less (including 0%).

(Al content)
Al (aluminum) is an essential element for increasing the C concentration of austenite in the present invention because it has an effect of suppressing precipitation of carbides in austenite and has a lower solid solution strengthening ability than Si. Moreover, since Al is a strong ferrite stabilizing element, the _Ae3 line is _shifted to the high C side, and the C concentration of austenite coexisting with ferrite can be increased, so that the stability of retained austenite is further increased. If the Al content is less than 0.40%, the effect of suppressing the formation of carbides cannot be sufficiently obtained. On the other hand, if the Al content exceeds 2.00%, an oxide layer is formed on the surface of the steel sheet, and the plating properties and chemical conversion treatment properties are significantly deteriorated. Furthermore, Al combines with N during slab casting to form AlN inclusions, which lowers the castability. In addition, the ferrite band structure is easily formed, and the ductility deteriorates because the structure becomes non-uniform. Therefore, the Al content is set to 0.40% or more and 2.00% or less. In order to exhibit the above-described effect of Al more effectively, the Al content is preferably 0.60% or more. Moreover, in order to avoid the said inconvenience more reliably, Al content is preferably 1.80% or less.

(B content)
B (boron) is an optional additive element. However, since B has an effect of suppressing pearlite transformation at the time of cooling by addition of a very small amount, it is an element that has the effect of avoiding pearlite transformation and causing bainite transformation in secondary cooling described later. However, excessive addition of B not only precipitates B carbide and lowers the hardenability, but also increases the hot deformation resistance and makes hot rolling difficult. Therefore, the B content is 0.0050% or less (including 0%).

(Total content of Al and Si)
Both Al and Si are elements that suppress carbide precipitation in austenite. As the total content of Al and Si increases, retained austenite tends to be generated. This effect is obtained when the total content of Al and Si is 0.7% or more. However, excessive addition of Al and Si promotes the ferrite band of the steel sheet and lowers the ductility. Further, since both Al and Si are easily oxidizable elements, excessive addition significantly deteriorates the plating quality and chemical conversion treatment. Therefore, the total content of Al and Si is set to 0.7% or more and 2% or less.

(Mn equivalent)
In order to obtain austenite that is stable even at room temperature, it is necessary to reduce the formation of pearlite during the cooling process. Therefore, the main alloy element in the steel of the present invention having the effect of delaying the ferrite / pearlite transformation is strictly managed as Mn equivalent (Mn _eq ). As a result of investigating the influence of various alloy elements on the formation of pearlite during cooling after annealing, Mn, Cr, P, and B have the effect of delaying the formation of pearlite, and Mn as shown in the following formula (1) It was found to be expressed as an equivalent formula. Furthermore, since the Mn equivalent also delays ferrite formation, it greatly affects the second phase fraction, and is also an important factor that almost determines the tensile strength (TS) of the steel of the present invention.

  On the other hand, in order to suppress the formation of pearlite and excessive bainite in the annealing process and to secure retained austenite, the results of investigating the influence of alloy elements on these transformation behaviors showed that in addition to the above Mn equivalent, Al and Si It was found to have a great inhibitory effect. Therefore, the inventors of the present invention investigated the influence of the Mn equivalent and Al and Si contents on the material of the steel plate after annealing. FIG. 1 shows C: 0.09 to 0.171%, Si: 0.01 to 0.80%, Mn: 0.3 to 2.35%, P: 0.01 to 0.04%, S: 0.002 to 0.009%, Al: 0.50 to 1.66%, Cr: 0.01 to 0.65%, N: 0.002%, B: 0 to 0.001% It is a figure which shows the result of having investigated the presence or absence of the pearlite generation | occurrence | production at the time of continuous cooling, and the relationship between Mn equivalent, content of Al and Si, and a material. In the figure, when TS is 440 MPa or more and less than 590 MPa, U.S. A steel sheet that achieves excellent ductility of El ≧ 25% and El ≧ 38% is ◯, and TS is 440 MPa or more and less than 590 MPa. A steel sheet that does not meet El ≧ 25% and El ≧ 38% is □, and TS ≧ 590 MPa. A steel sheet that did not reach El ≧ 25% and El ≧ 38% was indicated by Δ.

The method for producing the test piece is as follows. That is, a 27 mm-thick slab having the above components was heated to 1200 ° C., hot-rolled to a thickness of 4 mm at a finish rolling temperature of 950 ° C., immediately subjected to water spray cooling, and subjected to a winding process at 580 ° C. for 1 hour. . This hot-rolled sheet was cold-rolled to a thickness of 0.80 mm at a rolling rate of 80% to obtain a cold-rolled steel sheet. This was annealed at 800 to 850 ° C. for 100 seconds and then cooled at an average cooling rate of 15 ° C./s, Mn equivalent, CR (average cooling rate in primary cooling), Al amount, Si Cool down to an appropriate primary cooling stop temperature range T _crit (° C) determined by the amount, then cool to 350 ° C at an average cooling rate of 45 ° C / s, and immediately hold isothermal at 450 ° C for 30 seconds Then, the steel was cooled to 200 ° C. at an average cooling rate of 20 ° C./s, and then to room temperature at an average cooling rate of 10 ° C./s, and subjected to temper rolling with an elongation of 0.5%.

A JIS No. 5 tensile test piece was collected from the steel sheet thus obtained, and a tensile test was performed according to the method described in JIS Z2241 (1998). Further, the microstructure of the steel sheet was observed with a scanning electron microscope at a magnification of 3000 times to confirm the presence or absence of pearlite and bainite. 1 that Mn _eq needs to be 2.0 or less in order to promote the ferrite transformation of the steel sheet and reduce the strength to TS <590 MPa. It can be seen that the lower the Mn _{eq, the} higher the ferrite fraction and the lower the strength of the steel sheet. However, when the Mn _eq is less than 0.8, the pearlite transformation cannot be suppressed. Therefore, the range of Mn _eq is set to 0.8 ≦ Mn _eq ≦ 2.0 from the viewpoint of obtaining a delay effect of pearlite generation during cooling by Mn _eq while setting TS <590 MPa. Further, from the viewpoint of keeping the strength low, Mn _eq is preferably 1.9 or less.

Moreover, high total elongation and uniform elongation can be obtained in a region where Mn _eq is 0.8 or more and 2.0 or less and Mn _eq +1.3 ([% Al] + [% Si]) ≧ 2.6. I understand. In steel with Mn _eq +1.3 ([% Al] + [% Si]) <2.6, pearlite or bainite is generated during the cooling process or intermediate holding, and the amount of retained austenite is reduced, so that the ductility is lowered. In addition, steel with Mn _eq > 2.0 has low ductility because TS is too high, and steel with Mn _eq <0.8 has low ductility because pearlite is generated during air cooling. Based on the above results, Mn _eq +1.3 ([% Al] + [% Si]) ≧ 2.6 from the viewpoint of controlling the alloy elements to secure stable retained austenite and obtaining high uniform elongation. . Furthermore, from the viewpoint of obtaining a high TS × uniform elongation balance, 2.8 or more is preferable, and 3.0 or more is more preferable.

(Content of Ti, Nb, V)
Ti (titanium), Nb (niobium), and V (vanadium) all have strong affinity with N, have an effect of fixing N in the steel as a metal nitride, and can reduce the precipitation amount of AlN. . Therefore, these elements have the effect of suppressing the grain growth of the ferrite structure accompanying the fine AlN precipitation and the reduction of hot ductility. In order to obtain the effect of N fixation by Ti, Nb, and V, it is preferable to add these elements in an amount of 0.002% or more. However, since these are expensive elements, if they are added in a large amount, the cost will increase significantly, and fine carbides are likely to precipitate during annealing, and the strength of the steel sheet is increased. For this reason, Ti, Nb, and V are each preferably 0.02% or less.

(Ni content)
Since Ni (nickel) is an austenite stabilizing element, 0.05% or more can be added as necessary to suppress pearlite transformation. However, when Ni is added in a large amount, the alloy cost is increased and the ductility of the steel is lowered. Therefore, when adding Ni, the content is made 0.2% or less.

(Cu and Mo contents)
Since Cu (copper) and Mo (molybdenum) are austenite stabilizing elements, 0.02% or more can be added as needed for the purpose of suppressing pearlite transformation. Moreover, although not as much as Si and Al, these elements are also expected to have an effect of suppressing the formation of cementite. However, since both are expensive elements, the alloy cost is remarkably increased. Furthermore, since the strength of the steel sheet is increased by solid solution strengthening or by refining the steel structure, it is not preferable to contain a large amount. Accordingly, when Cu and Mo are added, their contents are each less than 0.1%, more preferably less than 0.05%.

(Contents of Sb and Sn)
Sb (antimony) and Sn (tin) can be added in a small amount to suppress oxidation and nitridation on the surface of the steel sheet, and can be added in an amount of 0.004% or more as necessary. However, if it is contained in a large amount, the strength is increased, the toughness is deteriorated, and the cost is increased. For this reason, when adding Sb and Sn, these content shall be 0.2% or less, respectively.

(Ca and REM contents)
Ca (calcium) and REM (rare earth metal) have a strong affinity with S. Therefore, S in steel is fixed, and the S inclusions that develop in steel by hot and cold rolling and become the starting point of fracture. In order to control the form, 0.002% or more can be added as needed. However, even if they are added in excess of 0.01%, the effect is saturated. For this reason, when adding Ca and REM, these content shall be 0.01% or less, respectively.

  The balance other than the above components is Fe and inevitable impurities. The N content is preferably in the following range.

(N content)
N (nitrogen) forms Al and fine AlN in the steel and lowers the grain growth of the ferrite structure, so that the structure is strengthened. In addition, when a large amount of AlN is precipitated, the hot ductility is drastically lowered, and thus the production stability in continuous casting is significantly impaired. Accordingly, N is an element that should be kept as low as possible. From such a viewpoint, the N content is preferably less than 0.004%, more preferably less than 0.0035%.

[Organization]
The high-strength cold-rolled steel sheet and the high-strength galvanized steel sheet according to the present invention have the above component composition, and further, the second phase volume fraction is 15% or less with ferrite as a parent phase, For example, martensite with a volume fraction of 1% or more and less than 8%, retained austenite with a volume fraction of 3% or more and less than 10%, bainite with a volume fraction of 1% or more and 5% or less, and a volume fraction of 0% or more and less than 2%. It has pearlite, and the total volume fraction of bainite and pearlite is less than or equal to the martensite volume fraction. Thereby, the tensile strength of 440 MPa or more and less than 590 MPa and excellent workability are obtained. Hereinafter, the structure of the high-strength cold-rolled steel sheet and the high-strength galvanized steel sheet according to the present invention will be described.

  Although ferrite and bainitic ferrite are slightly different in structure, it is not easy to distinguish between structures using an optical microscope or scanning electron microscope, and the properties are relatively close. Treat as. The second phase is a general term for phases existing in the structure other than ferrite. Pearlite refers to a layered structure composed of ferrite and cementite. Bainite refers to a hard structure formed from austenite at a relatively low temperature (above the martensitic transformation point), in which fine carbides are dispersed in acicular or plate-like ferrite. A structure in which only hard ferrite is generated without generating carbide, generally called bainitic ferrite, is included in the category of ferrite together with polygonal ferrite unless otherwise specified. Martensite may be partially tempered during cooling to become tempered martensite, but it is also hard enough to distinguish from martensite. Martensite and retained austenite are hardly corroded and difficult to distinguish with a microscope, but the volume fraction of retained austenite can be determined by X-ray diffraction. The volume fraction of martensite is obtained by excluding the volume fractions of all other phases from the second phase volume fraction. In addition,% showing the quantity of the volume fraction of a phase means the volume% unless there is particular notice.

(Second phase volume fraction)
Even if solid solution strengthening is suppressed by the alloy component, if the volume fraction of the second phase exceeds 15%, the tensile strength increases and the ductility is greatly reduced. Therefore, the volume fraction of the second phase is 15% or less, more preferably 14% or less. That is, the steel sheet according to the present invention has a ferrite with a volume fraction of 85% or more, more preferably 86% or more.

(Martensite volume fraction)
In the steel sheet according to the present invention, in order to suppress the occurrence of yield point elongation (YPEL), it is important to disperse the hard martensite appropriately enriched with C in a small amount in the ferrite. In order to suppress the generation of YPEl, at least 1% or more of martensite is required, and it is preferable to contain 3% or more of martensite. However, since excessive martensite leads to an increase in strength of the structure, the martensite volume fraction needs to be suppressed to less than 8%.

(Volume fraction of retained austenite)
In the steel sheet according to the present invention, control of the volume fraction of retained austenite is very important in order to obtain a high uniform elongation by utilizing the TRIP effect due to retained austenite. In order to obtain high uniform elongation, it is necessary to contain at least 3% or more and less than 10% of retained austenite. When the volume fraction of retained austenite is less than 3%, the uniform elongation is less than 25%. In order to obtain further excellent uniform elongation, the volume fraction of retained austenite is preferably 4% or more, and more preferably 5% or more. On the other hand, if excess retained austenite of 10% or more is present, TS ≧ 590 MPa and ductility is lowered, so the content is made less than 10%.

(Bainite volume fraction)
In the steel sheet according to the present invention, it is very important to generate a small amount of bainite in the second phase in order to promote the formation of bainitic ferrite during intermediate holding after secondary cooling. In order to promote the formation of bainitic ferrite, at least 1% bainite is required, and it is preferable to contain 2% or more bainite. However, when bainite is present in excess, the amount of C enrichment in the second phase is lowered, the structure is hardened, and ductility is lowered. Therefore, the volume fraction of bainite needs to be suppressed to 5% or less.

(Perlite volume fraction)
Since pearlite is generated from the second phase in which C is concentrated and consumes C, it is preferable not to generate pearlite as much as possible in order to reduce the stability of retained austenite and reduce ductility. For this reason, it is desirable that the volume fraction of pearlite be 0% or more and less than 2%, more preferably less than 1%.

(Total volume fraction of pearlite and bainite)
When a second phase accompanied by precipitation of carbides such as pearlite and bainite is generated, the amount of martensite generated decreases, and the effect is that the martensite is introduced between the martensite and ferrite, and the martensite introduces strain into the ferrite and eliminates YPEl. Is inhibited. Then, the influence which the structure | tissue structure in a 2nd phase has on YPEl was investigated. FIG. 2 shows the component composition of C: 0.09 to 0.171%, Si: 0.01 to 0.80%, Mn: 0.3 to 2.35%, P: 0.01 to 0.04% , S: 0.002 to 0.009%, Al: 0.22 to 1.66%, Cr: 0.01 to 0.65%, N: 0.002%, B: 0 to 0.001% Furthermore, the amount of Si + Al: 0.23 to 1.7%, Mneq: 0.5 to 2.9, and Mneq + 1.3 (Si + Al): 2.2 to 4.9, when the steel structure was changed. It is a figure which shows the change of YPEl. FIG. 2 shows the effect of the difference between martensite volume fraction (Vm) and pearlite and bainite total volume fraction (Vp + b) on YPEl.

  The method for producing the test piece is as follows. That is, a 27 mm-thick slab having the above composition is heated to 1200 ° C., hot rolled to a thickness of 4 mm at a finish rolling temperature of 870 to 970 ° C., immediately subjected to water spray cooling, and 1 in a temperature range of 450 to 650 ° C. A time winding process was performed. This hot-rolled sheet was cold-rolled to a thickness of 0.80 mm at a rolling rate of 80% to obtain a cold-rolled sheet. This was annealed at 720 to 900 ° C. for 100 seconds, then stopped at the primary cooling to a temperature range of 500 to 820 ° C. at an average cooling rate of 1 to 50 ° C./s, and then averaged to 150 to 500 ° C. After cooling at a rate of 2 to 80 ° C./s, immediately isothermally holding at 450 ° C. for 30 s, then cooling to 200 ° C. at an average cooling rate of 20 ° C./s, and then cooling to room temperature at an average cooling rate of 10 ° C./s Then, temper rolling with an elongation of 0.5% was performed.

  A JIS No. 5 tensile test piece was collected from the steel sheet thus obtained, and a tensile test was performed according to the method described in JIS Z2241 (1998). Moreover, the following measurements were performed in order to obtain the volume fraction of the microstructure of the steel sheet. That is, the L cross section (vertical cross section parallel to the rolling direction) of the steel sheet was polished and then corroded with nital, and 10 structural photographs were taken with a scanning electron microscope at a magnification of 3000 times. The area ratio of the second phase was measured by counting the number of pixels in the region corresponding to the phase and calculating the ratio to the number of pixels of the entire photograph. The area ratio of the L cross section obtained in this way showed almost the same value as the area ratio obtained from the vertical cross section parallel to the direction perpendicular to the rolling direction of the steel sheet. In this case, the area ratio of the L cross section was defined as the volume ratio of the second phase.

  In the structure photograph, ferrite is a dark color contrast area. The area where carbides are observed in a lamellar shape in ferrite is pearlite, the area where carbides are observed in a row in ferrite is bainite, and the other ferrites. Regions with brighter contrast were martensite or retained austenite. The volume fractions of the areas recognized as pearlite, bainite, martensite, and retained austenite were measured, and the total volume fraction was taken as the volume fraction of the second phase.

  The volume fraction of retained austenite was determined by the following method. By grinding and chemical polishing, a ¼ part thickness of the steel sheet is exposed, and an iron ferrite phase is formed by an X-ray diffractometer (apparatus: RINT2200 manufactured by Rigaku) using Mo-Kα rays as a radiation source and an acceleration voltage of 50 keV. Measure the integrated intensity of the X-ray diffraction lines of the {200} plane, {211} plane, {220} plane, and the {200} plane, {220} plane, and {311} plane of the austenite phase. The volume fraction of retained austenite was determined using the calculation formula described in non-patent literature (Rigaku Corporation: X-ray diffraction handbook (2000), p26, 62-64). The volume fraction of martensite was obtained by subtracting the volume fraction of retained austenite measured by the X-ray diffraction method from the total volume fraction of martensite and retained austenite measured from the above microstructure.

  From FIG. 2, it was found that when the total volume fraction of pearlite and bainite is equal to or less than the martensite volume fraction, that is, when the difference is equal to or greater than 0, YPEl is suppressed to 0.2% or less. For this reason, as the second phase structure, the total volume fraction of pearlite and bainite needs to be suppressed below the martensite volume fraction. Further, from the viewpoint of excellent ductility, the total volume fraction of pearlite and bainite is preferably 3% or less, more preferably 2% or less.

(Second phase volume fraction S / (second phase circumference L) ² )
Even in steel plates with the same second phase volume fraction, the C concentration from bainitic ferrite to each retained austenite grain is promoted when the second phase is eroded by a larger number of finer bainitic ferrites. As a result, it was found that the ductility was improved (the experimental results will be described later with reference to FIG. 5 under the manufacturing conditions). That is, at the same second phase volume fraction, the retained austenite becomes more stable as the area where the second phase contacts the ferrite is larger.

As an index for evaluating such structural characteristics, a value obtained by dividing the total volume S of the second phase in the microstructure of the steel sheet cross section by the square of the total L of the circumference of the second phase (the second phase (Proportional to the circularity when all the particles are assumed to be one particle) S / L ² can be appropriately arranged. The smaller the S / L ² is, the more the second phase is eroded by the ferrite, and the ductility is improved. When S / L ² is 0.00015 or more, there is a concern that ductility is insufficient because C concentration to residual austenite is insufficient. Accordingly, S / L ² is preferably less than 0.00015. More preferably, it is 0.00012 or less.

  In the present invention, by using the above component composition and microstructure, a high strength cold rolling with a tensile strength (TS) of 440 MPa or more and less than 590 MPa that is excellent in ductility and generation of yield point elongation (YPEL) is suppressed. Steel sheets and high-strength galvanized steel sheets can be obtained. Strength and ductility (TS x El) balance and strength-uniform elongation (TS x U. El) balance are obtained by multiplying the absolute value of ductility and uniform elongation by tensile strength as indicators of steel sheets with high strength and excellent ductility. Used. From the viewpoint of ensuring high strength and excellent press formability, the TS × El balance is preferably 19000 MPa ·% or more. More preferably, it is 20000 MPa *% or more, More preferably, it is 21000 MPa *% or more. Similarly, TS × U. The El balance is preferably 12000 MPa ·% or more, more preferably 13000 MPa ·% or more.

  Furthermore, it is necessary to suppress the occurrence of stretcher strain in the steel sheet from the viewpoint of maintaining excellent appearance quality even after pressing. If the yield point elongation (YPEL) exceeds 0.2%, a clear stretcher strain may occur in the press-formed product, so that YPEl is preferably 0.2% or less. More preferably, it is 0.1% or less, More preferably, it is 0%.

〔Production method〕
In the present invention, a high-strength cold-rolled steel sheet having excellent workability can be obtained by controlling a predetermined component steel to the above-described structure. A manufacturing method for obtaining such a high-strength cold-rolled steel sheet is described below. Will be described.

The slab heating temperature before hot rolling is set to 1100 to 1280 ° C. If heating temperature is less than 1100 degreeC, the rolling load at the time of hot rolling will increase. Further, since the composition of the present invention has a high Al content and a low Mn content, when the _Ae3 temperature is high and the slab heating temperature is less than 1100 ° C., a locally cooled portion such as an edge during hot rolling. May become a temperature below Ae _{3 and a} large amount of ferrite may be generated. Thereby, since the structure | tissue in a hot rolled sheet becomes non-uniform | heterogenous, the dispersion | variation in a material and the deterioration of a steel plate shape are caused. On the other hand, if the heating temperature is high, the alloy components and the structure can be made uniform and the rolling load can be reduced, but if it exceeds 1280 ° C., the oxide scale generated on the slab surface increases non-uniformly and the surface quality deteriorates. Therefore, the slab heating temperature is set to 1280 ° C. or less.

  Next, the hot rolling conditions are not particularly defined and may be performed according to a conventional method, but the finish rolling finish temperature (finishing temperature) is preferably 850 to 950 ° C. When the finishing temperature is less than 850 ° C., ferrite region rolling occurs, and uneven structure formation in the steel sheet and abnormal grain growth near the surface are likely to occur, making it difficult to obtain a stable material. On the other hand, when the finishing temperature exceeds 950 ° C., the generation of secondary scale is promoted and the surface quality is deteriorated. The range of the average cooling rate from hot rolling finish rolling to winding is not particularly specified. The hot-rolled sheet needs to be wound at a coiling temperature of more than 500 ° C. after receiving arbitrary cooling. This is because the hot-rolled sheet structure is a ferrite + pearlite structure, whereby the strength of the hot-rolled sheet can be reduced, and an increase in cold-rolling load in the next process can be suppressed. Furthermore, since austenite stabilizing elements such as C, Mn, and Cr once homogenized at the time of slab heating are locally distributed from the ferrite phase to the pearlite phase, they remain after annealing, and austenite stabilization is promoted. Cheap. In addition to the above, the internal oxidation of easily oxidizable elements tends to proceed in the wound coil, so that alloy elements form oxides on the surface of the steel sheet during subsequent annealing, suppressing deterioration of plating properties and chemical conversion properties There is an effect to. On the other hand, when the coiling temperature is 500 ° C. or less, the amount of low-temperature transformation phase generated increases, so the cold rolling load increases, the distribution of alloy elements decreases, and the amount of internal oxidation during hot rolling. Is reduced, and the plating property and the chemical conversion treatment property are lowered. From the viewpoint of cold rolling load, the hot rolled sheet structure is preferably 80% by volume or more of ferrite + pearlite structure, more preferably 90% by volume or more.

  After winding the hot-rolled sheet, it is subjected to pickling treatment and then cold-rolled. The conditions for cold rolling are not particularly specified and may be carried out in accordance with a conventional method. However, in order to obtain target characteristics, the cold rolling rate is preferably 40 to 90%. The cold-rolled sheet thus obtained is heated to an annealing temperature of 750 ° C. or higher and 950 ° C. or lower and held for 20 seconds or longer. Thereby, all the carbides in the structure are dissolved to generate austenite, and distribution of austenite stabilizing elements such as C, Mn, and Cr to the austenite is promoted. If the annealing temperature is less than 750 ° C. or the holding time is less than 20 seconds, undissolved carbides may remain and ductility may decrease. From the viewpoint of sufficiently dissolving the carbide, the annealing temperature is preferably set to 770 ° C. or higher. However, the operation at a soaking temperature exceeding 950 ° C. has a large load on the annealing equipment, so the annealing temperature is set to 950 ° C. or less. Further, if the soaking time exceeds 200 seconds, the annealing equipment is lengthened or the production rate is significantly reduced.

Immediately after annealing, primary cooling is performed. The primary cooling process is a very important process for determining the ferrite fraction of the steel sheet of the present invention. When the primary cooling stop temperature is lower, the ferrite fraction is increased and the ductility is improved. However, when the temperature is too low, pearlite transformation occurs, resulting in a decrease in ductility. Therefore, an appropriate primary cooling stop temperature was examined. FIG. 3 shows C: 0.12%, Si: 0.01%, Mn: 1.0 to 1.7%, P: 0.01%, S: 0.002 to 0.009%, Al: 0 .90 to 1.64%, Cr: 0.02 to 0.24%, N: 0.002%, B: 0 to 0.001% Mn equivalent Mn _eq and the total content of Al and Si The result of investigating the temperature at which pearlite formation occurs when different steel plates are continuously cooled at various average cooling rates is shown. In the figure, the pearlite generation temperature at an average cooling rate of 5 ° C./s is indicated by ◯, the pearlite generation temperature at 15 ° C./s is indicated by □, and the pearlite generation temperature at 30 ° C./s is indicated by Δ.

The method for producing the test piece is as follows. That is, a 27 mm-thick slab having the above components was heated to 1200 ° C., hot-rolled to a thickness of 4 mm at a finish rolling temperature of 950 ° C., immediately subjected to water spray cooling, and subjected to a winding process at 580 ° C. for 1 hour. . This hot-rolled sheet was cold-rolled to a thickness of 0.80 mm at a rolling rate of 80% to obtain a cold-rolled sheet. This was subjected to two-phase region annealing at 850 ° C. for 120 seconds, then cooled at an average cooling rate of 5, 15, 30 ° C./s, and water-cooled when various temperatures of 500 to 700 ° C. were reached. The microstructure of the steel sheet thus obtained was observed with a scanning electron microscope at a magnification of 3000 times to confirm the presence or absence of pearlite. From FIG. 3, the pearlite generation start temperature has a high Mn equivalent, the total content of Al and Si is small, and the lower the cooling rate (CR (° C./s)), the lower the temperature. From this result, it became clear that the pearlite generation start temperature can be arranged as the cooling limit temperature T _{crit according} to the following formula (2). When the primary cooling stop temperature is equal to or higher than T _crit , only the ferrite transformation proceeds, and the increase of the ferrite fraction and C concentration of austenite are promoted.

Furthermore, in order to investigate the influence of the primary cooling stop temperature on the total elongation (El) and uniform elongation (U. El), C: 0.123%, Si: 0.01%, Mn: 1.10%, P : 0.01%, S: 0.004%, Al: 1.50%, Cr: 0.23%, N: 0.0028%, B: 0%, and T _crit : 615 ° C. After performing cold rolling under the same conditions as in FIG. 3, a two-phase annealing is performed at 800 to 850 ° C. for 100 seconds, followed by cooling at an average cooling rate of 15 ° C./s, and a primary cooling stop temperature of 500 The temperature was changed to ˜820 ° C., followed by secondary cooling at an average cooling rate of 45 ° C./s to a secondary cooling stop temperature of 350 ° C., immediately followed by intermediate holding at 470 ° C. for 30 seconds, and then average cooling to 200 ° C. Cooling at an average cooling rate of 10 ° C./s from a rate of 20 ° C./s to a room temperature It was subjected to temper rolling of Choritsu 0.5 percent. A JIS No. 5 tensile test piece was collected from the obtained steel sheet, and a tensile test was performed according to the method described in JISZ2241 (1998).

FIG. 4 shows the relationship between the primary cooling stop temperature and the total elongation (El) and uniform elongation (U.El). FIG. 4 shows that when the primary cooling stop temperature is equal to or higher than the cooling limit temperature T _crit , pearlite generation during primary cooling is suppressed, and high uniform elongation can be obtained. Therefore, the primary cooling stop temperature is set to the cooling limit temperature T _crit or more. On the other hand, if the primary cooling stop temperature is too high, the C concentration of austenite is insufficient, so that the austenite stability at the time of rapid cooling is insufficient and a large amount of bainite is generated, and the uniform elongation is lowered. For this reason, in order to sufficiently transform the ferrite in the primary cooling step, it is preferable that the primary cooling stop temperature is T _crit + 170 ° C. or lower.

  In the steel sheet according to the present invention, when the average primary cooling rate is less than 5 ° C./s, the pearlite transformation is likely to occur at a high temperature, so that the uniform elongation tends to decrease. Accordingly, the average primary cooling rate is set to 5 ° C./s or more. In order to sufficiently avoid the pearlite transformation, the average primary cooling rate is more preferably more than 10 ° C./s. On the other hand, if it exceeds 40 ° C./s, the cooling rate is too fast and the ferrite transformation does not proceed sufficiently. Therefore, the primary cooling average speed is set to 40 ° C./s or less. Preferably it is 20 degrees C / s or less. Subsequently, secondary cooling is performed at an average cooling rate of 20 ° C./s or higher from the primary cooling stop temperature to a temperature range of 230 ° C. or higher and lower than 410 ° C. This was determined from the following results of investigating the influence of the secondary cooling stop temperature on the total elongation (El) and uniform elongation (U. El). That is, after carrying out the same component steel as in the case of the example shown in FIG. 4 until cold rolling, performing a two-phase region annealing at 850 ° C. for 100 seconds, and thereafter cooling at an average cooling rate of 15 ° C./s, The primary cooling stop temperature was set to 650 ° C., then the secondary cooling was performed at an average cooling rate of 25 ° C./s to the secondary cooling stop temperature of 200 to 500 ° C., and immediately after intermediate holding at 450 ° C. for 30 seconds, The steel was cooled to an average cooling rate of 20 ° C./s up to 0 ° C. and then cooled to an average cooling rate of 10 ° C./s from there to room temperature, and subjected to temper rolling with an elongation of 0.5%.

  A JIS No. 5 tensile test piece was collected from the obtained steel sheet, and a tensile test was performed according to the method described in JISZ2241 (1998). In addition, the microstructure is measured by measuring the volume fraction S of the second phase in the same manner as described with reference to FIG. 2, and the circumference L of the second phase is the same structure photograph as that measured by the volume fraction. Using the function of the image analysis software (Image Pro Plus ver.4.0), the number of pixels at the position corresponding to the outermost circumference of the second phase area is measured, and at the same time, the actual length per pixel is calculated from the photo data. By measuring, the number of pixels at the outermost periphery of the second phase region was converted to the actual length, and the peripheral length L was obtained.

The relationship between the secondary cooling stop temperature and S / L ² , total elongation (El), and uniform elongation (U.El) is shown in FIG. From FIG. 5, the effect of improving the ductility is recognized at the secondary cooling stop temperature lower than the intermediate holding temperature range described later. This is thought to be due to the effect of promoting the formation of fine bainitic ferrite by bainite since S / L ² is lowered. However, when the secondary cooling stop temperature is too low, excessive bainite is generated, resulting in a hardened structure and a decrease in retained austenite. FIG. 5 shows that the preferred range of the secondary cooling stop temperature is 230 ° C. or higher and lower than 410 ° C. Further, if the average cooling rate of the secondary cooling is less than 20 ° C./s, part of the austenite undergoes pearlite transformation, so that the ductility is lowered. Therefore, it is set to 20 ° C./s or more from the viewpoint of securing excellent ductility.

  Subsequent to secondary cooling, the steel sheet is heated to an intermediate holding temperature range of 410 ° C. or higher and 500 ° C. or lower and then held for 10 seconds or longer to generate bainitic ferrite, further increasing C concentration to austenite. Facilitate. If the holding time is less than 10 seconds, the formation of bainitic ferrite does not proceed sufficiently, the retained austenite fraction and its stability are insufficient, and the ductility decreases. On the other hand, even if holding is performed in this temperature range for more than 180 seconds, the austenite is decomposed into pearlite or bainite, so that the ductility is significantly reduced. For this reason, the holding time in the intermediate holding temperature region is set within 180 seconds. From the viewpoint of obtaining high ductility, the holding time is preferably within 140 seconds.

  When intermediate holding is started in a temperature range exceeding 500 ° C., a large amount of pearlite is generated in a short time, the volume fraction of retained austenite is reduced, the ductility is remarkably lowered, and yield point elongation (YPEL) is also generated. Moreover, if it is less than 410 degreeC, a bainite will produce | generate abundantly and a ductility will also fall remarkably and YPEl will generate | occur | produce. Accordingly, the intermediate holding temperature is set to 410 ° C. or more and 500 ° C. or less. More preferably, it is 420 degreeC or more and 490 degrees C or less.

  In this way, the intermediately held steel sheet is cooled at an average cooling rate of 10 ° C./s or more. Alternatively, after the intermediately held steel plate is immersed in a hot dip galvanizing bath to form galvanizing, it may be cooled at an average cooling rate of 10 ° C./s or more. Further, if necessary, the alloying treatment can be performed by raising the temperature to 490 to 600 ° C. and holding it for 3 to 100 seconds. In the alloying treatment, alloying does not proceed sufficiently if the alloying temperature is less than 490 ° C. When the alloying temperature exceeds 600 ° C., alloying is remarkably promoted, the galvanizing is hardened and easily peeled off, and austenite is transformed into pearlite, resulting in reduced ductility and generation of YPEl.

  This galvanized steel sheet may be subjected to temper rolling for the purpose of adjusting the surface roughness and flattening the shape of the steel sheet. However, since excessive pressure regulation reduces uniform elongation, it is preferable that the elongation ratio in pressure regulation be 0.1% or more and 0.7% or less. By using the manufacturing method as described above, a high-strength cold-rolled steel sheet excellent in ductility and a high-strength hot-dip galvanized steel sheet excellent in ductility and plating properties can be manufactured without using a complicated process. .

〔Example〕
Examples of the present invention will be described below.

  Table 1 shows the chemical composition of the test steel (remainder: iron and inevitable impurities, N corresponds to inevitable impurities), and Table 2 shows the production conditions. Steels having the chemical composition shown in Table 1 were melted in a vacuum melting furnace and cast into slabs. This was reheated to a temperature range of 1200 to 1250 ° C. and then subjected to rough rolling to obtain a rough bar having a thickness of about 27 mm. Subsequently, hot finish rolling was performed to a thickness of 4 mm at a finishing temperature of 900 ° C., and then held in a heating furnace at 450 to 650 ° C. for 1 hour, followed by furnace cooling to perform a winding equivalent process to obtain a hot rolled sheet. After removing the scale of the hot-rolled sheet by pickling, it was cold-rolled at a rolling rate of 80% to obtain a cold-rolled sheet having a thickness of 0.8 mm. The cold-rolled sheet thus obtained was subjected to annealing heating, primary cooling, secondary cooling, and intermediate holding in accordance with the temperature conditions shown in Table 2, and then an average cooling rate of 20 ° C./s to 200 ° C. To room temperature at an average cooling rate of 10 ° C./s to obtain a cold-rolled steel sheet. Some steel sheets are subjected to intermediate holding, then immersed in a hot-dip galvanizing bath at 460 ° C., cooled at an average cooling rate of 10 ° C./s to obtain galvanized steel sheets, and some steel sheets at 510 ° C. for 20 seconds. The alloying treatment was performed to obtain an alloyed hot-dip galvanized steel sheet. These steel sheets were subjected to temper rolling with an elongation of 0.5%.

  A JIS No. 5 tensile test piece was collected from the steel plate thus obtained so that the longitudinal direction was perpendicular to the rolling direction, and mechanical properties (tensile strength TS, yield point) were determined by a tensile test in accordance with JIS Z2241 (1998). Elongation YPEl, uniform elongation U.El, total elongation El) were evaluated. In addition, the volume fraction of the microstructure was measured. The volume fraction of the microstructure was measured as follows. That is, the L cross section (vertical cross section parallel to the rolling direction) of the steel sheet was polished and then corroded with nital, and 10 structural photographs were taken with a scanning electron microscope at a magnification of 3000 times. The area ratio was measured by counting the number of pixels in the region corresponding to the phase and calculating the ratio to the total number of pixels in the photo. The area ratio of the L cross section obtained in this way showed almost the same value as the area ratio obtained from the vertical cross section parallel to the direction perpendicular to the rolling direction of the steel sheet. Here, the area ratio of the L cross section was defined as the volume fraction. In addition, the circumference L of the second phase is the same as that measured for the volume fraction, and the outer circumference of the second phase region is set using the function of image analysis software (Image Pro Plus ver.4.0). By measuring the number of pixels at the corresponding position and simultaneously measuring the actual length per pixel from the photo data, the number of pixels on the outermost periphery of the second phase region is converted to the actual length, and the perimeter L is Asked.

  In the structure photograph, ferrite is a dark color contrast area. The area where carbides are observed in a lamellar shape in ferrite is pearlite, the area where carbides are observed in a row in ferrite is bainite, and the other ferrites. Regions with a brighter contrast were defined as martensite or retained austenite. The volume fractions of the regions recognized as pearlite, bainite, martensite, and retained austenite were measured, and the total volume fraction was taken as the second phase volume fraction. The volume fraction of ferrite is 100-second phase volume fraction.

  The volume fraction of retained austenite was determined by the following method. By grinding and chemical polishing, a ¼ part thickness of the steel sheet is exposed, and an iron ferrite phase is formed by an X-ray diffractometer (apparatus: RINT2200 manufactured by Rigaku) using Mo-Kα rays as a radiation source and an acceleration voltage of 50 keV. The integrated intensity of the X-ray diffraction lines of the {200} plane, {211} plane, {220} plane, and the {200} plane, {220} plane, {311} plane of the austenite phase are measured, and these measured values Was used to calculate the volume fraction of retained austenite using the calculation formula described in non-patent literature (Rigaku Corporation: X-ray diffraction handbook (2000), p26, 62-64). The volume fraction of martensite was determined by subtracting the volume fraction of residual austenite (residual γ) measured by X-ray diffraction from the total volume fraction of martensite and residual austenite measured from the above microstructure. These results are shown in Table 3.

  As can be seen from Table 3, the steel sheet of the present invention example satisfying the component composition and production conditions within the scope of the present invention has a TS of 440 MPa or more and less than 590 MPa. El is 25% or more, and El is 38% or more, and both show values satisfying the scope of the present invention. Furthermore, TS × U. El has an excellent workability with respect to the strength level of 12000 MPa ·% or more and TS × El of 19000 MPa ·% or more. Moreover, in the steel plate which plated, all the plating external appearances are favorable and non-plating is not recognized.

On the other hand, the steel plates with process numbers 1 and 37 manufactured using steels A and L with Mn _eq +1.3 ([% Al] + [% Si]) of less than 2.6 are pearlite and bainite volume fractions. Is high and the retained austenite is low. El was less than 25% and YPEL was also generated. In addition, steel plates with process numbers 33, 38, and 39 manufactured using steels H, M, and N with Mn _eq exceeding 2.0 and process numbers 35 manufactured with steel J having an excess of C Since TS has a high strength of 590 MPa or more due to an increase in volume fraction, El and U.I. The absolute value of El is very low. The steel plate with the processing number 34 manufactured using the steel I whose C is lower than the appropriate range has a small amount of retained austenite. El was less than 25%.

Even in the case of a test steel whose components are within the scope of the present invention, a steel sheet that deviates from the production conditions has any of the following inferior characteristics. In the treatment No. 31, the coiling temperature was 500 ° C. or less, the austenite stability was reduced, the amount of bainite and pearlite was increased, YPEl was generated, and the ductility was lowered. In the treatment number 24 having a low annealing temperature and the treatment number 23 having a short annealing time, the carbide of the cold-rolled base metal remained, the residual austenite fraction was less than 3%, and the ductility was lowered. Process No. 15 with a slow primary cooling rate, Process Nos. 19 and 20 with a primary cooling stop temperature less than the cooling limit temperature T _crit , and Process No. 25 with a slow secondary cooling rate are out of the appropriate manufacturing conditions. As the pearlite fraction in the second phase increased, the ductility decreased and YPEl was generated. Further, as described above with reference to FIG. 5, the ductility of the treatment numbers 7, 8, and 14 where the secondary cooling stop temperature deviated from 230 ° C. or more and less than 410 ° C. decreased. In the treatment number 26 and the treatment number 29 where the intermediate holding temperature is higher than the specified value, the desired microstructure was not obtained and the ductility was lowered, and YPEl was generated.

  In the treatment number 22 subjected to the alloying treatment, it was observed by observation with an electron microscope that fine pearlite was generated, and this decreased the retained austenite fraction and the ductility, but the prescribed components and The desired mechanical properties were obtained in the steel sheet that satisfied the production conditions.

  Although the embodiment to which the invention made by the present inventor is applied has been described above, the present invention is not limited by the description and the drawings that form a part of the disclosure of the present invention according to this embodiment. That is, other embodiments, examples, operational techniques, and the like made by those skilled in the art based on the present embodiment are all included in the scope of the present invention.

Claims (7)

As component composition, in mass%, C: 0.08% or more and 0.25% or less, Si: less than 0.50%, Mn: 0.7% or more and less than 2.0%, P: 0.1% or less, S: 0.01% or less, Cr: 1.0% or less (including 0%), Al: 0.40% or more and 2.00% or less, B: 0.0050% or less (including 0%) In addition, the total content of Si and Al is 0.7 to 2%, the Mn equivalent Mn _eq shown in the following formula (1) is 0.8 ≦ Mn _eq ≦ 2.0, and Mn _eq +1.3 × (Si content + Al content) ≧ 2.6, the balance is iron and inevitable impurities, and the structure is a ferrite with a volume fraction of 85% or more and a second phase with a volume fraction of 15% or less. The second phase is composed of martensite with a volume fraction of 1% or more and less than 8%, and residual austenite with a volume fraction of 3% or more and less than 10%. Bainite has a volume fraction of 1% to 5%, pearlite has a volume fraction of 0% to less than 2%, the total volume fraction of bainite and pearlite is less than martensite volume fraction, and tensile strength Is a high-strength cold-rolled steel sheet characterized by having a 440 MPa or more and less than 590 MPa.

  As a component composition, Ti: 0.02% or less, Mo: less than 0.1%, V: 0.02% or less, Nb: 0.02% or less, Ni: 0.2% or less, Cu: The high-strength cold-rolled steel sheet according to claim 1, further comprising one or more elements of less than 0.1%.   As a component composition, one or more elements of mass%, Sb: 0.2% or less, Sn: 0.2% or less, Ca: 0.01% or less, REM: 0.01% or less The high-strength cold-rolled steel sheet according to claim 1 or 2, further comprising: Any of claims 1 to 3 wherein the second phase value S / L ² where the volume fraction S divided by square of the length L of the outer periphery of the second phase is equal to or less than 0.00015 The high-strength cold-rolled steel sheet according to item 1.   The high-strength cold-rolled steel sheet according to any one of claims 1 to 4, comprising a zinc-based plating film on a surface thereof. The steel slab having the composition according to any one of claims 1 to 3 is subjected to hot rolling and cold rolling, and then subjected to continuous annealing, and is heated to 750 ° C to 950 ° C for 20 seconds. After being held above, after cooling at an average cooling rate of 5 ° _C./s or more and 40 ° _C./s or less to the cooling limit temperature T _crit (° C.) defined by the following formula (2), it is averaged from 230 ° C. to less than 410 ° C. A method for producing a high-strength cold-rolled steel sheet, which is cooled at a cooling rate of 20 ° C / s or higher, reheated to 410 ° C or higher and 500 ° C or lower and held for 10 seconds or longer and 180 seconds or shorter.

  The method for producing a high-strength cold-rolled steel sheet according to claim 6, wherein a zinc-based plating film is formed on the surface of the high-strength cold-rolled steel sheet by applying zinc-based plating in the course of continuous annealing. Manufacturing method of galvanized steel sheet.

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