JP2010065273A - High-strength steel sheet and method for manufacturing the same - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a high-strength steel sheet which is superior in workability and has a tensile strength (TS) of 980 MPa or more. <P>SOLUTION: This steel sheet has a composition comprising, by mass%, 0.17% to 0.73% C, 3.0% or less Si, 0.5% to 3.0% Mn, 0.1% or less P, 0.07% or less S, 3.0% or less Al, 0.010% or less N and the balance Fe with unavoidable impurities, while Si+Al is controlled to satisfy 0.7% or more; and has a metallurgical structure in which lower bainite and a total martensite are controlled to occupy 10% to 90% by an area rate in total with respect to the whole metallurgical structure, retained austenite is controlled to occupy 5% to 50% by an amount and bainitic ferrite in upper bainite is controlled to occupy 5% or more by an area rate with respect to the whole metallurgical structure, and in which an amount of martensite in as-quenched state out of the total amount of the lower bainite and the total martensite is controlled to satisfy 75% or less, the area rate of polygonal ferrite with respect to the whole metallurgical structure is controlled to satisfy 10% or less (including 0%) and an average C content in the retained austenite is controlled to 0.70% or more. <P>COPYRIGHT: (C)2010,JPO&INPIT

Description

  The present invention relates to a high-strength steel sheet having a tensile strength (TS) of 980 MPa or more excellent in workability, particularly ductility and stretch flangeability, used in industrial fields such as automobiles and electricity, and a method for producing the same.

  In recent years, improving the fuel efficiency of automobiles has become an important issue from the viewpoint of global environmental conservation. For this reason, efforts are being made to reduce the thickness of the vehicle body parts by increasing the strength of the vehicle body material and to reduce the weight of the vehicle body itself.

  Generally, in order to increase the strength of a steel sheet, it is necessary to increase the ratio of hard phases such as martensite and bainite to the entire structure of the steel sheet. However, since increasing the strength of the steel sheet by increasing the proportion of the hard phase causes a decrease in workability, development of a steel sheet having both high strength and excellent workability is desired. To date, various composite steel sheets such as ferrite-martensitic duplex steel (DP steel) and TRIP steel utilizing transformation-induced plasticity of retained austenite have been developed.

  When the ratio of the hard phase is increased in the composite structure steel plate, the workability of the steel plate is strongly influenced by the workability of the hard phase. This is because when the ratio of hard phase is small and soft polygonal ferrite is large, the deformability of polygonal ferrite dominates the workability of the steel sheet, and even when the hard phase has insufficient workability. Workability such as ductility was ensured, but when the ratio of the hard phase is large, the deformability of the hard phase itself directly affects the formability of the steel sheet, not the deformation of polygonal ferrite, and the hard phase itself This is because if the workability is insufficient, the workability of the steel sheet is significantly deteriorated.

  For this reason, in the case of a cold-rolled steel sheet, after performing annealing and adjusting the amount of polygonal ferrite generated in the subsequent cooling process, the steel sheet is water-quenched to generate martensite, and the steel sheet is raised again. By heating and holding at a high temperature, martensite has been tempered to generate carbides in the martensite that is a hard phase, thereby improving the workability of martensite. However, such martensite quenching and tempering requires special production equipment such as continuous annealing equipment having a water quenching function. Therefore, after water quenching the steel sheet, when using normal production equipment that cannot be heated again and kept at a high temperature, the steel sheet can be strengthened, but the workability of the martensite that is the hard phase Could not be improved.

  Further, as a steel plate having a hard phase other than martensite, there is a steel plate in which the main phase is polygonal ferrite, the hard phase is bainite or pearlite, and carbides are generated in these hard phases bainite or pearlite. This steel sheet is a steel sheet that not only improves the workability with polygonal ferrite alone, but also improves the workability of the hard phase itself by generating carbides in the hard phase, and in particular, improves the stretch flangeability. . However, as long as the main phase is polygonal ferrite, it is difficult to achieve both high strength and workability of 980 MPa or higher in tensile strength (TS). Also, even if the hard phase itself is improved in workability by generating carbides in the hard phase, the processability of polygonal ferrite is inferior, so the tensile strength (TS) is higher than 980 MPa. If the amount of polygonal ferrite is reduced in order to achieve the above, sufficient workability cannot be obtained.

Patent Document 1 proposes a high-tensile steel plate that is excellent in bending workability and impact properties by defining alloy components and making the steel structure fine and uniform bainite having retained austenite.
JP-A-4-235253

Patent Document 2 proposes a composite structure steel plate having excellent bake hardenability by defining predetermined alloy components, making the steel structure bainite having retained austenite, and defining the amount of retained austenite in bainite. ing.
JP 2004-76114 A

In Patent Document 3, a predetermined alloy component is defined, the steel structure is 90% or more in area ratio of bainite having retained austenite, the amount of retained austenite in bainite is 1% or more and 15% or less, and the hardness of bainite. By defining (HV), a composite structure steel plate excellent in impact resistance has been proposed.
JP-A-11-256273

However, the above-described steel sheet has the following problems.
In the component composition described in Patent Document 1, it is difficult to ensure a stable amount of retained austenite that exhibits the TRIP effect in a high strain region when strain is applied to the steel sheet, and bendability is obtained. However, the ductility until plastic instability occurs is low, and the stretchability is inferior.

  Although the steel sheet described in Patent Document 2 has bake hardenability, even when trying to increase the tensile strength (TS) to 980 MPa or higher, or even 1050 MPa or higher, bainite or martensite mainly composed of ferrite is suppressed as much as possible. Therefore, it is difficult to ensure workability such as ductility and stretch flangeability at the time of securing strength or increasing strength.

  The steel sheet described in Patent Document 3 is mainly intended to improve impact resistance, and has a main phase of bainite having a hardness of HV 250 or less, specifically a structure containing more than 90%. It is difficult to set the tensile strength (TS) to 980 MPa or more.

The present invention advantageously solves the above-mentioned problems, and provides a high-strength steel sheet having a tensile strength (TS) of 980 MPa or more, which is excellent in workability, particularly ductility and stretch flangeability, together with its advantageous production method. With the goal.
The high-strength steel sheet of the present invention includes a steel sheet obtained by subjecting the surface of the steel sheet to hot dip galvanization or galvannealing.
In the present invention, excellent workability means that TS × T. It is assumed that the value of EL satisfies 20000 MPa ·% or more and the value of TS × λ satisfies 25000 MPa ·%.

  In order to solve the above-mentioned problems, the inventors have conducted intensive studies on the component composition and microstructure of the steel sheet. As a result, the lower bainite structure and / or martensite structure is utilized to increase the strength, and the C content in the steel sheet is increased to 0.17% or more and the C content is increased, and then the upper bainite transformation is utilized. Thus, stable retained austenite advantageous for obtaining the TRIP effect can be secured, and by making a part of the martensite tempered martensite, workability, in particular, balance between strength and ductility, and strength and stretch flange. It was found that a high-strength steel sheet having a tensile strength of 980 MPa or more that is excellent in both the balance of properties can be obtained.

The present invention is based on the above findings, and the gist of the present invention is as follows.
1. C: 0.17% to 0.73% by mass%,
Si: 3.0% or less,
Mn: 0.5% to 3.0%,
P: 0.1% or less,
S: 0.07% or less,
Al: 3.0% or less and N: 0.010% or less, and Si + Al satisfies 0.7% or more, the balance is the composition of Fe and inevitable impurities,
As the steel sheet structure, the area ratio of the total amount of lower bainite and all martensite is 10% or more and 90% or less, the amount of retained austenite is 5% or more and 50% or less, and the steel structure of bainitic ferrite in the upper bainite. The area ratio with respect to the whole is 5% or more, of the total amount of the lower bainite and all martensite, 75% or less of the as-quenched martensite, and the area ratio of the polygonal ferrite with respect to the entire steel sheet structure is 10% or less (0% A high-strength steel sheet characterized in that the average C content in the retained austenite is 0.70% or more and the tensile strength is 980 MPa or more.

2. The steel sheet is further in mass%,
Cr: 0.05% to 5.0%,
2. The element according to 1 above, which contains one or more elements selected from V: 0.005% to 1.0% and Mo: 0.005% to 0.5%. High strength steel plate.

3. The steel sheet is further in mass%,
1 or 2 above, which contains one or two elements selected from Ti: 0.01% to 0.1% and Nb: 0.01% to 0.1%. High strength steel sheet as described.

4). The steel sheet is further in mass%,
B: The high-strength steel sheet according to any one of 1 to 3 above, which contains 0.0003% or more and 0.0050% or less.

5). The steel sheet is further in mass%,
1 to 4 above, which contains one or two elements selected from Ni: 0.05% or more and 2.0% or less and Cu: 0.05% or more and 2.0% or less. The high-strength steel sheet according to any one of the items.

6). The steel sheet is further in mass%,
1 to 5 above, which contains one or two elements selected from Ca: 0.001% to 0.005% and REM: 0.001% to 0.005%. The high-strength steel sheet according to any one of the items.

7). 7. A high-strength steel sheet comprising a hot-dip galvanized layer or an alloyed hot-dip galvanized layer on the surface of the steel sheet according to any one of 1 to 6 above.

8). The steel slab having the composition described in any one of 1 to 6 above is hot-rolled to form a cold-rolled steel sheet by cold rolling, and then the cold-rolled steel sheet is 600 seconds or more in the austenite single-phase region. After annealing for a second or less, the cooling stop temperature determined in the first temperature range of 350 ° C. or more and 490 ° C. or less: when cooling to T ° C., the cooling is controlled at an average cooling rate of 5 ° C./s or more until at least 550 ° C. Thereafter, holding at 15 to 1000 seconds in the first temperature range, and then holding at 15 to 1000 seconds in the second temperature range from 200 to 350 ° C. Method.

9. 9. The method for producing a high-strength steel sheet as described in 8 above, wherein a hot dip galvanizing treatment or an alloying hot dip galvanizing treatment is performed at the cooling stop temperature: T ° C. or in the first temperature range.

  According to the present invention, it is possible to provide a high-strength steel sheet having a tensile strength (TS) of 980 MPa or more that is excellent in workability, particularly ductility and stretch flangeability, together with its advantageous manufacturing method, and industrial fields such as automobiles and electricity. The utility value is extremely large, and is particularly useful for reducing the weight of automobile bodies.

The present invention will be specifically described below.
First, the reason why the steel sheet structure is limited as described above in the present invention will be described. Hereinafter, the area ratio is the area ratio relative to the entire steel sheet structure.

Area ratio of total amount of lower bainite and all martensite: 10% or more and 90% or less Lower bainite and martensite are structures necessary for increasing the strength of a steel sheet. If the area ratio of the total amount of lower bainite and all martensite is less than 10%, the tensile strength (TS) of the steel sheet does not satisfy 980 MPa. On the other hand, if the area ratio of the total amount of the lower bainite and all martensite exceeds 90%, the upper bainite decreases, and as a result, stable retained austenite enriched in C cannot be secured, so workability such as ductility is reduced. Decreasing becomes a problem. Therefore, the area ratio of the total amount of lower bainite and all martensite is set to 10% or more and 90% or less. Preferably, it is in the range of 20% to 80%. More preferably, it is in the range of 30% to 70%.

Of the total amount of lower bainite and all martensite, ratio of as-quenched martensite: 75% or less Of martensite, the ratio of as-quenched martensite is the sum of lower bainite and all martensite present in the steel sheet. If it exceeds 75% of the amount, the tensile strength becomes 980 MPa or more, but the stretch flangeability is inferior. The as-quenched martensite is extremely hard and the deformability of the as-quenched martensite itself is extremely low, so that the workability of the steel sheet, particularly the stretch flangeability, is significantly deteriorated. In addition, since the hardness difference between the as-quenched martensite and the upper bainite is remarkably large, when the amount of the as-quenched martensite is large, the interface between the as-quenched martensite and the upper bainite increases, and during punching, Microvoids are generated at the interface between the as-quenched martensite and the upper bainite, and the stretched flangeability deteriorates because the voids are easily connected and cracks are easily developed during the stretched flange forming after the punching process. Accordingly, the ratio of martensite as quenched in martensite is 75% or less with respect to the total amount of lower bainite and all martensite present in the steel sheet. Preferably it is 50% or less. The as-quenched martensite is a structure in which carbides are not recognized in the martensite and can be observed by SEM.

Residual austenite amount: 5% or more and 50% or less Residual austenite undergoes martensitic transformation by the TRIP effect during processing, and improves ductility by increasing strain dispersibility.
In the steel sheet of the present invention, utilizing the upper bainite transformation, in particular, retained austenite with an increased amount of C concentration is formed in the upper bainite. As a result, retained austenite that can exhibit the TRIP effect even in a high strain region during processing can be obtained. By utilizing such retained austenite and martensite in combination, good workability can be obtained even in a high strength region where the tensile strength (TS) is 980 MPa or more, specifically, TS × T. The value of El can be set to 20000 MPa or more, and a steel sheet having an excellent balance between strength and ductility can be obtained.
Here, the retained austenite in the upper bainite is formed between the laths of bainitic ferrite in the upper bainite and is finely distributed. Therefore, in order to obtain the amount (area ratio) by microstructure observation, a large amount of measurement is performed at a high magnification. Is necessary and accurate quantification is difficult. However, the amount of retained austenite formed between the laths of the bainitic ferrite is a certain amount commensurate with the amount of bainitic ferrite formed. Therefore, as a result of investigations by the inventors, the area ratio of bainitic ferrite in the upper bainite is 5% or more, and X-ray diffraction (XRD) which is a conventional method for measuring the amount of retained austenite is performed. If the amount of retained austenite obtained from the strength measurement, specifically the X-ray diffraction intensity ratio of ferrite and austenite is 5% or more, a sufficient TRIP effect can be obtained, and the tensile strength (TS) is 980 MPa or more. TS × T. It was found that El can achieve 20000 MPa ·% or more. It has been confirmed that the amount of retained austenite obtained by a conventional method for measuring the amount of retained austenite is equivalent to the area ratio of retained austenite to the entire steel sheet structure.
When the amount of retained austenite is less than 5%, a sufficient TRIP effect cannot be obtained. On the other hand, if it exceeds 50%, hard martensite generated after the TRIP effect appears becomes excessive, which causes problems such as deterioration of toughness. Therefore, the amount of retained austenite is in the range of 5% to 50%. Preferably, it is in the range of more than 5%, more preferably 10% or more and 45% or less. More preferably, it is the range of 15% or more and 40% or less.

Average C content in retained austenite: 0.70% or more In order to obtain excellent workability by utilizing the TRIP effect, in a high strength steel sheet having a tensile strength (TS) of 980 MPa to 2.5 GPa, The amount of C in the austenite is important. In the steel sheet of the present invention, C is concentrated in the retained austenite formed between the laths of bainitic ferrite in the upper bainite. Although it is difficult to accurately evaluate the amount of C concentrated in the retained austenite between the laths, as a result of the study by the inventors, in the steel sheet of the present invention, the conventional austenite in the retained austenite If the average C content in the retained austenite obtained from the shift amount of the diffraction peak in X-ray diffraction (XRD), which is a method for measuring the average C content (average of the C content in the retained austenite) is 0.70% or more It was found that excellent processability can be obtained.
When the average C content in the retained austenite is less than 0.70%, martensitic transformation occurs in the low strain region during processing, and the TRIP effect in the high strain region that improves workability cannot be obtained. Therefore, the average amount of C in the retained austenite is 0.70% or more. Preferably it is 0.90% or more. On the other hand, if the average C content in the retained austenite exceeds 2.00%, the retained austenite becomes excessively stable, the martensitic transformation does not occur during processing, and the TRIP effect does not appear, thereby reducing ductility. Accordingly, the average C content in the retained austenite is preferably 2.00% or less. More preferably, it is 1.50% or less.

The area ratio of bainitic ferrite in the upper bainite: 5% or more The formation of bainitic ferrite by the upper bainite transformation concentrates C in the untransformed austenite and exhibits the TRIP effect in the high strain region during processing. It is necessary to obtain retained austenite that enhances strain resolution. The transformation from austenite to bainite occurs over a wide temperature range of approximately 150 to 550 ° C., and various types of bainite are produced within this temperature range. In the prior art, such various bainite was often simply defined as bainite, but in order to obtain the target workability in the present invention, it is necessary to clearly define the bainite structure. The bainite and lower bainite are defined as follows.
The upper bainite is composed of lath-like bainitic ferrite and residual austenite and / or carbide existing between bainitic ferrite, and there is no fine carbide regularly arranged in lath-like bainitic ferrite. It is a feature. On the other hand, the lower bainite is composed of the lath-shaped bainitic ferrite and the residual austenite and / or carbide existing between the bainitic ferrites in common with the upper bainite. It is characterized by the presence of fine carbides regularly arranged in the bainitic ferrite.
That is, the upper bainite and the lower bainite are distinguished by the presence or absence of regularly arranged fine carbides in bainitic ferrite. Such a difference in the state of carbide formation in bainitic ferrite has a great influence on the concentration of C in the retained austenite. That is, when the area ratio of the bainitic ferrite of the upper bainite is less than 5%, even when the bainite transformation is advanced, the amount of C generated as carbides in the bainitic ferrite increases, resulting in a The amount of C enriched in the residual austenite present in the steel decreases, and the amount of residual austenite that exhibits the TRIP effect in the high strain region during processing decreases. Therefore, the area ratio of bainitic ferrite in the upper bainite needs to be 5% or more in terms of the area ratio with respect to the entire steel sheet structure. On the other hand, if the area ratio of the bainitic ferrite of the upper bainite to the entire steel sheet structure exceeds 85%, it may be difficult to ensure the strength.

Polygonal ferrite area ratio: 10% or less (including 0%)
When the area ratio of polygonal ferrite exceeds 10%, it becomes difficult to satisfy the tensile strength (TS) of 980 MPa or more, and at the same time, strain concentrates on the soft polygonal ferrite mixed in the hard structure during processing. As a result, cracks are easily generated during processing, and as a result, desired workability cannot be obtained. Here, if the area ratio of polygonal ferrite is 10% or less, even if polygonal ferrite is present, a small amount of polygonal ferrite is isolated and dispersed in the hard phase, and strain concentration can be suppressed. Degradation of workability can be avoided. Therefore, the area ratio of polygonal ferrite is 10% or less. Preferably it is 5% or less, More preferably, it is 3% or less, and 0% may be sufficient.

  In the case of the steel sheet of the present invention, the hardness of the hardest structure in the steel sheet structure is HV ≦ 800. That is, in the steel sheet of the present invention, when there is no as-quenched martensite, either the tempered martensite or the lower bainite or the upper bainite is the hardest phase, but these structures all have HV ≦ 800. It is a phase. In addition, when there is as-quenched martensite, the as-quenched martensite becomes the hardest structure, in the steel sheet of the present invention, even if it is as-quenched martensite, the hardness is HV ≦ 800, There is no extremely hard martensite such that HV> 800, and good stretch flangeability can be secured.

  The steel sheet of the present invention may contain pearlite, Widmanstatten ferrite, or lower bainite as the remaining structure. In that case, the allowable content of the remaining tissue is preferably 20% or less in terms of area ratio. More preferably, it is 10% or less.

  The above is the basic structure of the steel sheet structure in the high-strength steel sheet of the present invention, but the following structure may be added as necessary.

Next, the reason why the component composition of the steel sheet is limited as described above in the present invention will be described. In addition,% showing the following component compositions shall mean the mass%.
C: 0.17% or more and 0.73% or less C is an element indispensable for increasing the strength of a steel sheet and securing a stable retained austenite amount, and for securing a martensite amount and retaining austenite at room temperature. It is a necessary element. If the C content is less than 0.17%, it is difficult to ensure the strength and workability of the steel sheet. On the other hand, if the amount of C exceeds 0.73%, the welded part and the heat-affected zone are hardened and the weldability deteriorates. Therefore, the C content is in the range of 0.17% or more and 0.73% or less. Preferably, it is 0.20% or more and 0.48% or less of range, More preferably, it is 0.25% or more.

Si: 3.0% or less (including 0%)
Si is a useful element that contributes to improving the strength of steel by solid solution strengthening. However, if the amount of Si exceeds 3.0%, the workability and toughness deteriorate due to the increase in the amount of solid solution in polygonal ferrite and bainitic ferrite, and the surface properties due to the occurrence of red scale, etc. In the case of deterioration or hot dipping, it causes deterioration of plating adhesion and adhesion. Therefore, the Si content is 3.0% or less. Preferably it is 2.6% or less. More preferably, it is 2.2% or less.
Si is an element useful for suppressing the formation of carbides and promoting the formation of retained austenite. Therefore, the Si content is preferably 0.5% or more, but the formation of carbides is only Al. In the case of suppressing by Si, Si does not need to be added, and the Si amount may be 0%.

Mn: 0.5% to 3.0% Mn is an element effective for strengthening steel. If the amount of Mn is less than 0.5%, carbide precipitates in a temperature range higher than the temperature at which bainite and martensite are generated during cooling after annealing, so ensure the amount of hard phase that contributes to strengthening of steel. I can't. On the other hand, if the amount of Mn exceeds 3.0%, castability is deteriorated. Therefore, the amount of Mn is set in the range of 0.5% to 3.0%. Preferably it is 1.5 to 2.5% of range.

P: 0.1% or less
P is an element useful for strengthening steel, but if the P content exceeds 0.1%, the impact resistance deteriorates due to embrittlement due to grain boundary segregation, and the steel sheet is subjected to alloying hot dip galvanizing. In this case, the alloying speed is greatly delayed. Therefore, the P content is 0.1% or less. Preferably it is 0.05% or less. The amount of P is preferably reduced, but if it is less than 0.005%, it causes a significant increase in cost, so the lower limit is preferably about 0.005%.

S: 0.07% or less
Since S generates MnS to become inclusions, causing deterioration of impact resistance and cracking along the metal flow of the weld, it is preferable to reduce the amount of S as much as possible. However, excessively reducing the amount of S causes an increase in manufacturing cost, so the amount of S is set to 0.07% or less. Preferably it is 0.05% or less, More preferably, it is 0.01% or less. In addition, since it is accompanied by a big increase in manufacturing cost to make S less than 0.0005%, the lower limit is about 0.0005% from the point of manufacturing cost.

Al: 3.0% or less Al is an element useful for strengthening steel and a useful element added as a deoxidizer in the steel making process. When the amount of Al exceeds 3.0%, the inclusions in the steel sheet increase and the ductility deteriorates. Therefore, the Al content is 3.0% or less. Preferably, it is 2.0% or less.
Al is an element useful for suppressing the formation of carbides and promoting the formation of retained austenite. In order to obtain a deoxidizing effect, the Al content should be 0.001% or more. Preferably, it is more preferably 0.005% or more. The amount of Al in the present invention is the amount of Al contained in the steel sheet after deoxidation.

N: 0.010% or less
N is an element that most deteriorates the aging resistance of steel, and is preferably reduced as much as possible. When the N content exceeds 0.010%, deterioration of aging resistance becomes remarkable, so the N content is set to 0.010% or less. Note that, if N is less than 0.001%, a large increase in manufacturing cost is caused, so that the lower limit is about 0.001% from the viewpoint of manufacturing cost.

The basic component has been described above. However, in the present invention, it is not sufficient to satisfy the above component range, and it is necessary to satisfy the following equation.
Si + Al ≧ 0.7%
As described above, both Si and Al are useful elements for suppressing the formation of carbides and promoting the formation of retained austenite. Although suppression of the formation of carbides is effective even if Si or Al is contained alone, it is necessary to satisfy 0.7% or more in total of the Si amount and the Al amount. The amount of Al in the above formula is the amount of Al contained in the steel sheet after deoxidation.

Moreover, in this invention, the component described below other than the above-mentioned basic component can be contained appropriately.
Cr: 0.05% or more and 5.0% or less, V: 0.005% or more and 1.0% or less, and Mo: 0.005% or more and 0.5% or less , V and Mo are elements having an action of suppressing the formation of pearlite during cooling from the annealing temperature. The effect is obtained when Cr: 0.05% or more, V: 0.005% or more, and Mo: 0.005% or more. On the other hand, if it exceeds Cr: 5.0%, V: 1.0%, and Mo: 0.5%, the amount of hard martensite becomes excessive and the strength becomes higher than necessary. Therefore, when Cr, V and Mo are contained, Cr: 0.05% to 5.0%, V: 0.005% to 1.0% and Mo: 0.005% to 0.5% % Or less.

One or two selected from Ti: 0.01% or more and 0.1% or less, Nb: 0.01% or more and 0.1% or less Ti and Nb are useful for the precipitation strengthening of steel. , Each content is obtained at 0.01% or more. On the other hand, when each content exceeds 0.1%, the workability and the shape freezing property are lowered. Therefore, when Ti and Nb are contained, the range is Ti: 0.01% to 0.1% and Nb: 0.01% to 0.1%.

B: 0.0003% or more and 0.0050% or less B is an element useful for suppressing the formation and growth of ferrite from the austenite grain boundary. The effect is obtained when the content is 0.0003% or more. On the other hand, if the content exceeds 0.0050%, the workability decreases. Therefore, when it contains B, it is set as B: 0.0003% or more and 0.0050% or less of range.

One or two selected from Ni: 0.05% or more and 2.0% or less and Cu: 0.05% or more and 2.0% or less. Ni and Cu are effective elements for strengthening steel. Moreover, when performing hot dip galvanization or alloying hot dip galvanization to a steel plate, the internal oxidation of a steel plate surface layer part is accelerated | stimulated and plating adhesiveness is improved. These effects are obtained when the respective contents are 0.05% or more. On the other hand, when each content exceeds 2.0%, the workability of the steel sheet is lowered. Therefore, when Ni and Cu are contained, the range is Ni: 0.05% to 2.0% and Cu: 0.05% to 2.0%.

One or two types selected from Ca: 0.001% or more and 0.005% or less and REM: 0.001% or more and 0.005% or less Ca and REM spheroidize the shape of the sulfide, and stretch flange Useful to improve the negative effects of sulfides on sex. The effect is obtained when each content is 0.001% or more. On the other hand, if the respective contents exceed 0.005%, inclusions and the like increase, causing surface defects and internal defects. Therefore, when Ca and REM are contained, the range is Ca: 0.001% to 0.005% and REM: 0.001% to 0.005%.

  In the steel plate of the present invention, components other than those described above are Fe and inevitable impurities. However, as long as the effects of the present invention are not impaired, the inclusion of components other than those described above is not rejected.

Next, the manufacturing method of the high strength steel plate of this invention is demonstrated.
A steel slab adjusted to the above preferred component composition is manufactured, then hot-rolled, and then cold-rolled to obtain a cold-rolled steel sheet. In the present invention, these treatments are not particularly limited, and may be performed according to ordinary methods.
The preferred production conditions are as follows. After heating the steel slab to a temperature range of 1000 ° C. or higher and 1300 ° C. or lower, hot rolling is finished in a temperature range of 870 ° C. or higher and 950 ° C. or lower, and the obtained hot rolled steel sheet is heated to a temperature of 350 ° C. or higher and 720 ° C. or lower. Take up in the area. Next, the hot-rolled steel sheet is pickled and then cold-rolled at a rolling reduction in the range of 40% to 90% to obtain a cold-rolled steel sheet.
In the present invention, it is assumed that the steel sheet is manufactured through normal steelmaking, casting, hot rolling, pickling and cold rolling processes. However, for example, the steel plate is heated by thin slab casting or strip casting. You may manufacture by omitting a part or all of a hot rolling process.

The obtained cold-rolled steel sheet is subjected to the heat treatment shown in FIG. Hereinafter, a description will be given with reference to FIG.
Annealing is performed for 15 seconds to 600 seconds in an austenite single phase region. The steel sheet of the present invention is mainly composed of upper bainite, lower bainite and martensite which are transformed from untransformed austenite in a relatively low temperature range of 350 ° C. or more and 490 ° C. or less. Preferably, annealing in the austenite single phase region is necessary. The annealing temperature is not particularly limited as long as it is in the austenite single phase region, but if the annealing temperature exceeds 1000 ° C., the growth of austenite grains is remarkable, which causes coarsening of the constituent phases caused by subsequent cooling, and deteriorates toughness and the like. Let On the other hand, the annealing temperature is of less than A ₃ point (austenitic transformation point) is already generated by the polygonal ferrite in the annealing step, in order to suppress the growth of polygonal ferrite during cooling than 500 ° C. It becomes necessary to cool the temperature range very rapidly. Therefore, the annealing temperature is required to be A ₃ point (austenitic transformation point) ° C. or higher, preferably set to 1000 ° C. or less.
Moreover, when annealing time is less than 15 second, the reverse transformation to austenite may not fully advance, or the carbide | carbonized_material in a steel plate may not fully melt | dissolve. On the other hand, if the annealing time exceeds 600 seconds, the cost increases due to the great energy consumption. Accordingly, the annealing time is in the range of 15 seconds to 600 seconds. Preferably, it is the range of 60 seconds or more and 500 seconds or less. Here, A ₃ points are
A ₃ points (° C.) = 910−203 × [C%] ^1/2 + 44.7 × [Si%] − 30 × [Mn%]
+ 700 × [P%] + 130 × [Al%] − 15.2 × [Ni%]
−11 × [Cr%] − 20 × [Cu%] + 31.5 × [Mo%]
+ 104 × [V%] + 400 × [Ti%]
Can be calculated. Note that [X%] is the mass% of the component element X of the steel sheet.

  The cold-rolled steel sheet after annealing is cooled to a cooling stop temperature: T ° C. determined in a first temperature range of 350 ° C. or more and 490 ° C. or less, but the average cooling rate is controlled to 5 ° C./s or more until at least 550 ° C. And cooled. When the average cooling rate is less than 5 ° C./s, excessive formation and growth of polygonal ferrite and precipitation of pearlite occur, and a desired steel sheet structure cannot be obtained. Therefore, the average cooling rate from the annealing temperature to the first temperature range is set to 5 ° C./s or more. Preferably, it is 10 ° C./s or more. The upper limit of the average cooling rate is not particularly limited as long as the cooling stop temperature does not vary. In general equipment, when the average cooling rate exceeds 100 ° C./s, the structure in the longitudinal direction and the sheet width direction of the steel sheet. The dispersion is significantly increased, so that it is preferably 100 ° C./s or less.

  The steel sheet cooled to 550 ° C. is continuously cooled to a cooling stop temperature: T ° C. The speed at which the steel sheet is cooled in the temperature range of T ° C. or more and 550 ° C. or less is not particularly limited except that the holding time in the first holding temperature range is 15 seconds or more and 1000 seconds or less, but the steel sheet is excessively slow. In the case of cooling with, a carbide is generated from untransformed austenite, and there is a high possibility that a desired structure cannot be obtained. Therefore, in the temperature range of T ° C. or higher and 550 ° C. or lower, the steel plate is preferably cooled at an average rate of 1 ° C./s or higher.

Cooling stop temperature: The steel sheet cooled to T ° C is held for a period of 15 seconds to 1000 seconds in a first temperature range of 350 ° C to 490 ° C. When the upper limit of the first temperature range exceeds 490 ° C., carbide precipitates from untransformed austenite and a desired structure cannot be obtained. On the other hand, when the lower limit of the first temperature range is less than 350 ° C., lower bainite is generated instead of upper bainite, and there is a problem that the amount of C concentration in austenite decreases. Therefore, the range of the first temperature range is 350 ° C. or more and 490 ° C. or less. Preferably, it is the range of 370 degreeC or more and 460 degreeC or less.
Further, when the holding time in the first temperature range is less than 15 seconds, the amount of upper bainite transformation is reduced, and the amount of C enrichment in untransformed austenite is problematic. On the other hand, when the retention time in the first temperature range exceeds 1000 seconds, stable retained austenite in which C is concentrated by precipitation of carbides from untransformed austenite which becomes retained austenite as the final structure of the steel sheet cannot be obtained. The desired processability cannot be obtained. Therefore, the holding time is 15 seconds or more and 1000 seconds or less. Preferably, it is the range of 30 seconds or more and 600 seconds or less.

The steel sheet that has been held in the first temperature range is cooled to a second temperature range of 200 ° C. or higher and 350 ° C. or lower at an arbitrary rate, and held in the second temperature range for 15 seconds to 1000 seconds. When the upper limit of the second temperature range exceeds 350 ° C., the lower bainite transformation does not proceed, resulting in a problem that the amount of martensite as quenched is increased. On the other hand, when the lower limit of the second temperature range is less than 200 ° C., similarly, the lower bainite transformation does not proceed and the amount of martensite as quenched is increased. Therefore, the range of the second temperature range is 200 ° C. or more and 350 ° C. or less. Preferably, it is the range of 250 degreeC or more and 340 degrees C or less.
Moreover, when holding time is less than 15 second, sufficient quantity of lower bainite cannot be obtained and desired workability cannot be obtained. On the other hand, if the holding time exceeds 1000 seconds, carbide precipitates from stable retained austenite in the upper bainite generated in the first temperature range, and as a result, desired workability cannot be obtained. Accordingly, the holding time is in the range of 15 seconds to 1000 seconds. Preferably, it is the range of 30 seconds or more and 600 seconds or less.

  In the series of heat treatments according to the present invention, the holding temperature does not have to be constant as long as it is within the predetermined temperature range described above, and the gist of the present invention is not impaired even if it fluctuates within the predetermined temperature range. The same applies to the cooling rate. Further, as long as the thermal history is satisfied, the steel sheet may be heat-treated with any equipment. Further, it is also included in the scope of the present invention that after the heat treatment, the surface of the steel sheet is subjected to temper rolling or surface treatment such as electroplating for shape correction.

  The method for producing a high-strength steel sheet according to the present invention can further include hot dip galvanizing treatment or galvannealing treatment in which alloying treatment is further added to hot dip galvanizing treatment. The hot dip galvanizing process or the alloying hot dip galvanizing process may be performed during the cooling to the first temperature range or in the first temperature range. In this case, the holding time in the first temperature range is 15 seconds or more and 1000 seconds or less including the processing time in the first temperature range of the hot dip galvanizing process or the alloying galvanizing process. The hot dip galvanizing treatment or alloying hot dip galvanizing treatment is preferably performed in a continuous hot dip galvanizing line.

Moreover, in the manufacturing method of the high-strength steel sheet of the present invention, after manufacturing the high-strength steel sheet completed up to the heat treatment according to the above-described manufacturing method of the present invention, the hot-dip galvanizing treatment or the alloying hot-dip galvanizing treatment is performed again. You can add that.
Further, according to the production method of the present invention, the hot dip galvanizing treatment or the alloying hot dip galvanizing treatment can be performed after the holding in the second temperature range.

The method of performing hot dip galvanizing treatment or alloying hot dip galvanizing treatment on a steel sheet is as follows.
The steel sheet is infiltrated into the plating bath and the amount of adhesion is adjusted by gas wiping. The amount of dissolved Al in the plating bath ranges from 0.12% to 0.22% in the case of hot dip galvanizing, and ranges from 0.08% to 0.18% in the case of galvannealed alloying. It is preferable that
In the case of hot dip galvanizing, the temperature of the plating bath may be in the range of 450 ° C. or higher and 500 ° C. or lower. When further alloying is performed, the temperature during alloying is 550 ° C. or lower. It is preferable. When the alloying temperature exceeds 550 ° C., carbide precipitates from untransformed austenite or pearlite is generated in some cases, so that strength and workability or both cannot be obtained, and the powdering property of the plating layer is also low. to degrade. On the other hand, if the temperature during alloying is less than 450 ° C., alloying may not proceed.
Coating weight is preferably in a per side 20 g / ^{m 2} or more 150 g / ^{m 2} or less. If the plating adhesion amount is less than 20 g / m ² , the corrosion resistance is insufficient. On the other hand, if it exceeds 150 g / m ² , the corrosion resistance effect is saturated and only the cost is increased.
The alloying degree (Fe mass% (Fe content)) of the plating layer is preferably in the range of 7 mass% to 15 mass%. If the degree of alloying of the plating layer is less than 7% by mass, unevenness in alloying occurs and the appearance quality deteriorates, or the so-called ζ phase is generated in the plating layer and the slidability of the steel sheet deteriorates. On the other hand, when the degree of alloying of the plating layer exceeds 15% by mass, a large amount of hard and brittle Γ phase is formed, and the plating adhesion deteriorates.

  EXAMPLES Hereinafter, although an Example demonstrates this invention further in detail, the following Example does not limit this invention. In addition, changing the configuration within the scope of the gist configuration of the present invention is included in the scope of the present invention.

The slab obtained by melting the steel having the composition shown in Table 1 is heated to 1200 ° C, the hot-rolled steel sheet finished by hot rolling at 870 ° C is wound up at 650 ° C, and then the hot-rolled steel sheet is pickled. Thereafter, it was cold-rolled at a rolling rate of 65% to obtain a cold-rolled steel sheet having a thickness of 1.2 mm. The obtained cold-rolled steel sheet was heat-treated under the conditions shown in Table 2. In addition, the cooling stop temperature: T in Table 2 is a temperature at which the cooling of the steel sheet is stopped when the steel sheet is cooled from the annealing temperature.
Further, some cold-rolled steel sheets were subjected to hot dip galvanizing treatment or alloying hot dip galvanizing treatment. Here, in the hot dip galvanizing treatment, double-sided plating was performed so that the plating bath temperature was 463 ° C. and the basis weight (per one side) was 50 g / m ² . In addition, the alloying hot dip galvanizing treatment is performed by adjusting the alloying conditions so that the weight per unit area (per one side): 50 g / m ² and the degree of alloying (Fe mass% (Fe content)) is 9 mass%. Double-sided plating was applied. In addition, the hot dip galvanizing treatment and the alloying hot galvannealing hot dip galvanizing treatment were performed after cooling to T ° C shown in Table 2 once.

  When the obtained steel sheet is not subjected to plating treatment, after heat treatment, when subjected to hot dip galvanizing treatment or alloyed hot dip galvanizing treatment, after these treatments, rolling ratio (elongation rate): 0.3% Temper rolling was applied.

Various properties of the steel sheet thus obtained were evaluated by the following methods.
Samples were cut from each steel plate and polished, and the surface parallel to the rolling direction was observed with a scanning electron microscope (SEM) at 10Ox field of view at 10 000 times to measure the area ratio of each phase. The phase structure was identified.

  The amount of retained austenite was determined by measuring the X-ray diffraction intensity after grinding and polishing the steel plate to ¼ of the plate thickness in the plate thickness direction. For incident X-rays, Co—Kα is used, and from the intensity ratio of each surface of austenite (200), (220), (311) to the diffraction intensity of each surface of ferrite (200), (211), (220). The average value was calculated for the amount of retained austenite.

The average amount of C in the retained austenite is obtained by calculating the lattice constant from the intensity peaks of the (200), (220) and (311) surfaces of austenite in the X-ray diffraction intensity measurement. C amount (mass%) was calculated | required.
a ₀ = 0.3580 + 0.0033 × [C%] + 0.00095 × [Mn%]
+ 0.0056 × [Al%] + 0.022 × [N%]
However, _{a 0:} the lattice constant (nm), [X%] : % by weight of the element X. In addition, mass% of elements other than C was mass% with respect to the whole steel plate.

  The tensile test was performed according to JIS Z2241, using a JIS No. 5 test piece taken from a direction perpendicular to the rolling direction of the steel sheet. TS (tensile strength) and T.El (total elongation) were measured, and the product of strength and total elongation (TS × T.El) was calculated to evaluate the balance between strength and workability (ductility). In the present invention, the case of TS × T.E1 ≧ 20000 MPa ·% was considered good.

The stretch flangeability was evaluated in accordance with Japan Iron and Steel Federation standard JFST1001. Each steel plate obtained was cut to 100 mm × 100 mm, a hole with a clearance of 12% of the plate thickness and a diameter of 10 mm was punched out, and then pressed with a wrinkle holding force of 88.2 kN using a die with an inner diameter of 75 mm. In this state, a 60 ° conical punch was pushed into the hole, the hole diameter at the crack initiation limit was measured, and the critical hole expansion ratio λ (%) was obtained from the equation (1).
Limit hole expansion rate λ (%) = {(D _f −D ₀ ) / D ₀ } × 100 (1)
However, _{D f} is the hole diameter at crack initiation (mm), _{D 0} is the initial hole diameter (mm).
The product of strength and limit hole expansion rate (TS × λ) was calculated using λ measured in this manner, and the balance between strength and stretch flangeability was evaluated.
In the present invention, when TS × λ ≧ 25000 MPa ·%, the stretch flangeability is good.

  Further, the hardness of the hardest structure in the steel sheet structure was determined by the following method. That is, when martensite is observed as-quenched as a result of structure observation, these martensite as-quenched is measured at 10 points at a load of 0.02N with ultra micro Vickers, and the average value thereof is measured in the steel sheet structure. The hardness of the hardest tissue. In addition, when martensite is not recognized as quenched, any of the structures of tempered martensite, upper bainite or lower bainite is the hardest phase in the steel sheet of the present invention. These hardest phases were HV ≦ 800 in the case of the steel sheet of the present invention.

  The above evaluation results are shown in Table 3.

  As is clear from the table, all the steel sheets of the present invention satisfy the tensile strength of 980 MPa or more, and the value of TS × T.El is 20000 MPa ·% or more and TS × λ ≧ 25000 MPa ·%. It was confirmed that it had high strength and excellent workability, especially excellent stretch flangeability.

In contrast, sample no. No. 1 has an average cooling rate of up to 550 ° C. that is outside the proper range, so that a desired steel sheet structure cannot be obtained and TS × λ ≧ 25000 MPa ·% is satisfied, but tensile strength (TS) ≧ 980 MPa and TS × T. EL ≧ 20000 MPa ·% was not satisfied. Sample No. No. 2 is a sample No. 2 because the holding time in the first temperature range is outside the proper range. 5, since the annealing temperature is less than ₃ points _A, Sample No. 6 is the cooling stop temperature: T is outside the first temperature range. Since the holding temperature in the second temperature range is outside the proper range, No. 11 has a holding time in the second temperature range that is outside the appropriate range, so that a desired steel sheet structure cannot be obtained, and although tensile strength (TS) ≧ 980 MPa is satisfied, TS × T. Either EL ≧ 20000 MPa ·% or TS × λ ≧ 25000 MPa ·% was not satisfied. Sample No. 30 to 34, since the component composition is outside the proper range, the desired steel sheet structure cannot be obtained, and the tensile strength (TS) ≧ 980 MPa, TS × T. Any one or more of EL ≧ 20000 MPa ·% and TS × λ ≧ 25000 MPa ·% were not satisfied.

It is the figure which showed the temperature pattern of the heat processing in the manufacturing method according to this invention.

Claims (9)

C: 0.17% to 0.73% by mass%,
Si: 3.0% or less,
Mn: 0.5% to 3.0%,
P: 0.1% or less,
S: 0.07% or less,
Al: 3.0% or less and N: 0.010% or less, and Si + Al satisfies 0.7% or more, the balance is the composition of Fe and inevitable impurities,
As the steel sheet structure, the area ratio of the total amount of lower bainite and all martensite is 10% or more and 90% or less, the amount of retained austenite is 5% or more and 50% or less, and the steel structure of bainitic ferrite in the upper bainite. The area ratio with respect to the whole is 5% or more, of the total amount of the lower bainite and all martensite, 75% or less of the as-quenched martensite, and the area ratio of the polygonal ferrite with respect to the entire steel sheet structure is 10% or less (0% A high-strength steel sheet characterized in that the average C content in the retained austenite is 0.70% or more and the tensile strength is 980 MPa or more. The steel sheet is further in mass%,
Cr: 0.05% to 5.0%,
2. One or more elements selected from V: 0.005% or more and 1.0% or less and Mo: 0.005% or more and 0.5% or less are contained. High strength steel sheet as described. The steel sheet is further in mass%,
3. One or two elements selected from Ti: 0.01% or more and 0.1% or less and Nb: 0.01% or more and 0.1% or less are contained. The high-strength steel sheet described in 1. The steel sheet is further in mass%,
The high-strength steel plate according to any one of claims 1 to 3, wherein B: 0.0003% or more and 0.0050% or less is contained. The steel sheet is further in mass%,
5. One or two elements selected from Ni: 0.05% or more and 2.0% or less and Cu: 0.05% or more and 2.0% or less are contained. The high-strength steel sheet according to any one of the above. The steel sheet is further in mass%,
6. One or two elements selected from Ca: 0.001% or more and 0.005% or less and REM: 0.001% or more and 0.005% or less are contained. The high-strength steel sheet according to any one of the above.   A high-strength steel plate comprising a hot-dip galvanized layer or an alloyed hot-dip galvanized layer on the surface of the steel plate according to any one of claims 1 to 6.   The steel slab having the composition according to any one of claims 1 to 6 is hot-rolled to form a cold-rolled steel sheet by cold rolling, and then the cold-rolled steel sheet is at least 15 seconds in the austenite single-phase region. After annealing for 600 seconds or less, cooling stop temperature determined in the first temperature range of 350 ° C. or more and 490 ° C. or less: When cooling to T ° C., cooling is performed by controlling the average cooling rate to 5 ° C./s or more until at least 550 ° C. And holding for 15 seconds to 1000 seconds in the first temperature range, and then holding for 15 seconds to 1000 seconds in the second temperature range of 200 ° C. to 350 ° C. Production method.   The method for producing a high-strength steel sheet according to claim 8, wherein hot-dip galvanizing treatment or alloying hot-dip galvanizing treatment is performed at the time of cooling to T ° C or in the first temperature range.

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