CN109154045B

CN109154045B - Plated steel sheet and method for producing same

Info

Publication number: CN109154045B
Application number: CN201780029820.5A
Authority: CN
Inventors: 高岛克利; 泽西央海; 谷口公一; 小林崇; 田川哲哉; 池田伦正
Original assignee: JFE Steel Corp
Current assignee: JFE Steel Corp
Priority date: 2016-05-25
Filing date: 2017-05-11
Publication date: 2020-10-09
Anticipated expiration: 2037-05-11
Also published as: KR102145293B1; JPWO2017203994A1; JP6424967B2; WO2017203994A1; CN109154045A; KR20180133508A; MX2018014261A

Abstract

The invention provides a plated steel sheet having an extremely high tensile strength and excellent resistance to delayed fracture and resistance to solder cracking. Provided is a plated steel sheet having the following structure: contains 35 to 70% of ferrite by volume fraction, 12% or less of retained austenite by volume fraction, 15 to 60% of martensite by volume fraction, 30% or less of bainite by volume fraction as the remaining part, and 5% or less of unrecrystallized ferrite by volume fraction, wherein the ferrite has an average crystal grain size of 5 μm or less, the retained austenite has an average crystal grain size of 2 μm or less, the martensite has an average crystal grain size of 2 μm or less, the bainite has an average crystal grain size of 3 μm or less, and the microstructure is a microstructure of 100 μm or less²Contains 30 or more Ti or Nb precipitates having an average grain diameter of 0.10 μm or less on average.

Description

Plated steel sheet and method for producing same

Technical Field

The present invention relates to a plated steel sheet and a method for producing the same, and particularly to a plated steel sheet suitable for use as a member of a structural member of an automobile or the like.

Background

In recent years, CO has become serious as environmental problems have become serious₂Emission regulations are becoming more stringent, and in the automotive field, it is a challenge to reduce the weight of a vehicle body in order to improve fuel efficiency. Therefore, the use of steel sheets having a Tensile Strength (TS) of 980MPa or more has been advanced for thinning automobile parts by applying high-strength steel sheets to automobile parts. In addition, from the viewpoint of corrosion resistance, a plated steel sheet having a hot-dip galvanized layer is used for a portion exposed to rainwater.

Patent document 1 discloses a technique for improving delayed fracture resistance (hydrogen embrittlement resistance) by controlling iron-based carbides containing Si or Si and Al in a steel sheet structure.

Patent document 2 discloses a technique for improving surface cracking during resistance welding by controlling the amounts of Si, Al, and Mn.

Documents of the prior art

Patent document

Patent document 1: japanese patent No. 4949536

Patent document 2: japanese patent No. 3758515

Disclosure of Invention

However, in the technique described in patent document 1, even if the iron-based carbide can serve as a trap site for hydrogen, there is a possibility that the improvement effect of the resistance welding cracking resistance is not obtained, and the iron-based carbide exists in the grain boundary to promote cracking during resistance welding.

In addition, in the technique described in patent document 2, it is difficult to achieve high strength of 980MPa or more and to obtain excellent delayed fracture resistance. As described above, in the case of the plated steel sheet of 980MPa or more, it is difficult to improve the resistance to both delayed fracture and resistance to solder cracking, and in fact, no steel sheet having both of these properties has been developed, including steel sheets other than the plated steel sheet.

Further, a high-strength steel sheet used for structural members or reinforcing members of automobiles may be subjected to delayed fracture (hydrogen embrittlement) by hydrogen entering the steel sheet due to the use environment. In addition, in many cases, high-strength steel sheets are produced by combining press-formed members by resistance welding (spot welding). However, in such spot welding, zinc on the surface of the steel sheet is melted and residual stress is generated in the vicinity of the welded portion, whereby liquid metal embrittlement may occur, which may cause cracking of the steel sheet. Therefore, in order to apply a high-strength hot-dip galvanized steel sheet, it is necessary to have both excellent delayed fracture resistance and excellent resistance to solder cracking.

The present invention has been made in view of the above circumstances, and an object thereof is to solve the problems of the prior art described above, and to provide a plated steel sheet excellent in delayed fracture resistance and resistance to solder cracking, and a method for producing the same.

The present inventors have conducted extensive studies to improve both the delayed fracture resistance and the solder crack resistance. As a result, they have found that excellent delayed fracture resistance and solder crack resistance can be obtained at the same time by controlling the volume fraction of the steel sheet structure of ferrite, retained austenite, martensite, bainite, and unrecrystallized ferrite at a specific ratio and by making the average crystal grain size of each steel sheet structure fine to generate fine carbides of Ti or Nb series in the steel sheet structure. The present invention is based on the above findings.

The delayed fracture is caused by hydrogen entering the steel sheet, cracking and progressing to cause fracture. For example, in a hot-dip galvanized steel sheet which can also be used as a thin steel sheet for automobiles, the coating may be scratched for some reason, and the iron-based surface may be exposed. In such a situation, if rainwater or the like adheres to the surface of the steel sheet, zinc acts as an anode, and iron acts as a cathode, thereby accelerating the generation of hydrogen on the iron surface. Therefore, it is necessary for the plated steel sheet to consider delayed fracture characteristics of the steel sheet, not the composition and composition of the plating layer.

Further, regarding cracking caused by brittleness due to liquid metal during resistance welding, internal stress is generated due to Zn melted (melted) during welding, and cracking occurs in a Heat Affected Zone (HAZ) in the vicinity of a weld spot. Conventionally, in a plated steel sheet, if welding is performed at a high current value at which spatters (dross) are generated during welding, cracking due to brittleness of liquid metal may occur on the surface on the electrode contact side. However, if the tensile strength is high in the order of 980MPa, cracking (internal cracking) may occur on the surfaces where the steel sheets overlap each other even in an appropriate current range where no spattering occurs. In particular, if welding is performed with the welding electrode at an angle to the steel sheet, internal stress increases, and cracking is likely to occur. Here, the resistance welding crack in the present invention indicates the internal crack. If such an internal crack occurs, the fatigue strength of the welded portion may be particularly reduced. Therefore, for use in automobiles and the like, it is necessary to avoid such an internal crack. When the internal crack portion is observed, it is found that cracks due to intergranular fracture occur at a position where the heat-affected zone (HAZ) becomes a martensite single phase after welding.

Accordingly, the inventors have made extensive studies and found that by controlling the size and number of fine Ti or Nb-based carbides, hydrogen trap sites can be generated and delayed fracture resistance can be improved. Further, it was found that the susceptibility to brittleness of the liquid metal during welding can be improved by making the crystal grains of the steel sheet finer. Further, it was found that strength, delayed fracture resistance and resistance to solder cracking can be improved by controlling the volume fraction of the steel sheet structure. Fine Ti or Nb-based carbides not only serve as hydrogen trap sites, but also inhibit the growth of ferrite and austenite nuclei during recrystallization in annealing to promote the formation of ferrite and austenite nuclei. Therefore, fine Ti or Nb-based carbide is very effective for refining the steel sheet structure. As described above, it was found that by refining the steel sheet structure, the crystal grains are not coarsened during resistance welding, and the toughness of the steel sheet is improved, so that resistance welding cracking can be suppressed.

The present invention has been made in view of the above-described new findings, and has the following configuration.

1. A plated steel sheet having a composition of: contains, in mass%, C: 0.05-0.22%, Si: 0.05-1.80%, Mn: 1.45% -3.35%, P: 0.05% or less, S: 0.005% or less, Al: 0.01% -0.10%, N: 0.010% or less and B: 0.0002% to 0.0045%, and further contains a titanium compound selected from the group consisting of Ti: 0.005% -0.090% and Nb: more than 1 of 0.005-0.090%, the balance being Fe and inevitable impurities,

further, the tissue has the following structure: contains 35 to 70% of ferrite by volume fraction, 12% or less of retained austenite by volume fraction, 15 to 60% of martensite by volume fraction, and 30% or less of bainite by volume fraction and 5% or less of unrecrystallized ferrite by volume fraction as the remaining portion,

the ferrite has an average crystal grain diameter of 5 μm or less,

the retained austenite has an average crystal grain diameter of 2 [ mu ] m or less,

the martensite has an average crystal grain diameter of 2 [ mu ] m or less,

the average crystal grain diameter of the bainite is less than 3 mu m,

the tissue is at every 100 μm²Contains 30 or more Ti or Nb precipitates having an average grain diameter of 0.10 μm or less on average.

2. The plated steel sheet according to the above 1, wherein the composition further contains, in mass%, a component selected from the group consisting of V: 0.10% or less, Cu: 0.50% or less, Ni: 0.50% or less, Mo: 0.50% or less, Cr: 0.80% or less and Ca and/or REM: 0.0050% or less, and 1 or 2 or more.

3. A method for producing a plated steel sheet, comprising subjecting a billet having a composition of hot rolling under conditions such that the finish temperature of finish rolling is 850 to 950 ℃ to produce a hot rolled steel sheet,

cooling the hot-rolled steel sheet at a 1 st average cooling rate of 75 ℃/s or more to 680 ℃ or less and at a 2 nd average cooling rate of 5 ℃/s or more to a temperature in the range of 400 to 580 ℃, then coiling the steel sheet, cold-rolling the coiled steel sheet to obtain a cold-rolled steel sheet,

annealing the cold-rolled steel sheet by heating the steel sheet to a temperature range of 760 to 900 ℃ at an average heating rate of 3 to 30 ℃/s, soaking the steel sheet in the temperature range of 760 to 900 ℃ for 15 seconds or more, cooling the steel sheet to a temperature range of 600 ℃ or less at an average cooling rate of 3 to 30 ℃/s,

performing plating treatment on the annealed cold-rolled steel sheet;

the steel billet comprises the following components in percentage by mass: 0.05-0.22%, Si: 0.05-1.80%, Mn: 1.45% -3.35%, P: 0.05% or less, S: 0.005% or less, Al: 0.01% -0.10%, N: 0.010% or less and B: 0.0002% to 0.0045%, and further contains, in mass%, a titanium compound selected from the group consisting of Ti: 0.005% -0.090%, Nb: 0.005-0.090%, and the balance Fe and inevitable impurities.

4. The method for producing a plated steel sheet according to the above 3, wherein the composition further contains, in mass%, a component selected from the group consisting of V: 0.10% or less, Cu: 0.50% or less, Ni: 0.50% or less, Mo: 0.50% or less, Cr: 0.80% or less and Ca and/or REM: 0.0050% or less, and 1 or 2 or more.

5. The method for producing a plated steel sheet according to the above 3 or 4, wherein after the plating treatment, alloying treatment for plating is performed in a temperature range of 450 to 600 ℃.

According to the present invention, the alloy material has an extremely high tensile strength, has excellent delayed fracture resistance characteristics in which delayed fracture due to hydrogen entering from the environment does not occur even after a member is formed, and has excellent resistance to solder cracking characteristics in which cracking does not occur even during resistance welding. For example, a high-strength plated steel sheet excellent in strength, delayed fracture resistance, and resistance to solder cracking can be stably obtained, which has a tensile strength of 980MPa or more, does not crack after U-bending in a hydrochloric acid immersion environment of pH 1.5 at 20 ℃ for 100 hours, and does not crack when welded with an electrode forming an angle (about 0.5 to 10 °) with respect to the steel sheet.

Detailed Description

Hereinafter, a plated steel sheet according to an embodiment of the present invention will be described. First, the reasons for the limitations of the composition of the steel will be described. In the present specification, "%" indicating the content of each component element means "% by mass" unless otherwise specified.

C：0.05％～0.22％

C is an element effective for increasing the strength of the steel sheet, and contributes to the formation of the 2 nd phase (a structure other than ferrite as the 1 st phase) of bainite, martensite, and retained austenite in the present invention. If the content is less than 0.05%, it is difficult to secure the volume fractions of bainite, martensite, and retained austenite required, and thus it is difficult to secure the strength. Preferably 0.06% or more. More preferably 0.065% or more. On the other hand, if it is excessively added, the hardness after resistance welding becomes high, the toughness at the time of resistance welding is lowered, and the resistance to solder cracking is deteriorated, so that the content thereof is 0.22% or less. Preferably 0.20% or less, and more preferably 0.18% or less.

Si：0.05％～1.80％

Si is an element effective for strengthening ferrite by solid solution. To obtain this effect, 0.05% or more needs to be added. Preferably 0.10% or more. More preferably 0.20% or more. However, when Si is excessively added, the plating property is lowered and the plating is partially not fully plated, and therefore, the content thereof is 1.80% or less. Preferably 1.60% or less. More preferably 1.50% or less.

Mn：1.45％～3.35％

Mn is an element contributing to high strength by solid-solution strengthening and formation of the 2 nd phase. Further, the element for stabilizing austenite is an element necessary for controlling the fraction of the 2 nd phase. In order to obtain this effect, the content of the compound is required to be 1.45% or more. Preferably 1.60% or more. More preferably 1.80% or more. On the other hand, if the content is excessively large, the volume fraction of the 2 nd phase becomes excessive, and if hydrogen enters the steel sheet, the sliding restriction of grain boundaries increases, and cracks in grain boundaries easily develop, so that the delayed fracture resistance is lowered. Therefore, the content is 3.35% or less. Preferably 3.20% or less. More preferably 3.0% or less.

P: less than 0.05%

P contributes to increase in strength by solid-solution strengthening, but when added excessively, segregation into grain boundaries becomes significant and grain boundaries become brittle, and therefore, resistance to solder cracking is reduced. Therefore, the content is 0.05% or less. Preferably 0.04% or less. More preferably 0.03% or less. There is no particular lower limit, but the steel-making cost is increased by the extremely low P content, so that the content is preferably 0.0005% or more. More preferably 0.0008% or more.

S: less than 0.005%

When the content of S is large, a large amount of sulfide such as MnS is produced, and cracks are produced from MnS when hydrogen intrudes, so that delayed fracture resistance is lowered. Therefore, the upper limit of the content is made 0.005%. Preferably 0.0045% or less. More preferably 0.004% or less. Although there is no particular lower limit, the steel-making cost is increased by making the S content extremely low, as in P, and therefore, it is preferable to contain 0.0002% or more. More preferably 0.0004% or more.

Al：0.01％～0.10％

Al is an element necessary for deoxidation, and in order to obtain this effect, 0.01% or more needs to be contained. Preferably 0.015% or more. On the other hand, even if the content exceeds 0.10%, the effect is saturated, and therefore, the content is 0.10% or less. Preferably 0.06% or less. More preferably 0.05% or less.

N: 0.010% or less

Since N forms coarse nitrides to deteriorate delayed fracture resistance, the content thereof needs to be suppressed. In particular, when N exceeds 0.010%, this tendency becomes remarkable, and therefore, the content of N is made 0.010% or less. Preferably 0.008% or less. More preferably 0.006% or less.

B：0.0002％～0.0045％

B is an element that improves hardenability, contributes to high strength by forming phase 2, and does not lower the martensite transformation start point while securing hardenability. Further, since grain boundary strength is improved by segregation in grain boundaries, it is effective for delayed fracture resistance. In order to exhibit this effect, the content is 0.0002% or more. Preferably 0.0003% or more. However, excessive addition of the metal compound causes deterioration of toughness, and thus lowers the resistance to solder cracking, so that the content thereof is 0.0045% or less. Preferably 0.0035% or less. More preferably 0.0030% or less.

Selected from the group consisting of Ti: 0.005% -0.090% and Nb: more than 1 of 0.005-0.090%

Ti is an element that can contribute to an increase in strength by forming fine carbonitrides. In addition, fine carbonitride of Ti is effective for hydrogen trap sites and for grain refinement, and therefore, is also effective for resistance weld cracking suppression. In order to exert such an effect, the lower limit of the content of Ti is set to 0.005%. The preferred lower limit is 0.008%. The more preferable lower limit is 0.010%. On the other hand, if Ti is added in a large amount, ductility is significantly reduced, and therefore the content thereof is 0.090% or less. Preferably 0.080% or less. More preferably 0.070% or less.

Similarly to Ti, Nb also contributes to an increase in strength and serves as a hydrogen trap site by forming fine carbonitride, and is effective for grain refinement. In order to exert such an effect, the lower limit of the content of Nb is set to 0.005%. The preferred lower limit is 0.008%. The more preferable lower limit is 0.010%. On the other hand, if Nb is added in a large amount, not only ductility is significantly reduced, but also the recrystallization rate is significantly reduced, and therefore unrecrystallized ferrite increases. Therefore, the content thereof is 0.090% or less. Preferably 0.080% or less. More preferably 0.070% or less.

The essential components of the present invention are explained above. The balance other than the above components is Fe and inevitable impurities, and in the present invention, 1 or 2 or more of the following components may be added in addition to the above basic components.

V: less than 0.10%

V can contribute to an increase in strength by forming fine carbonitrides. In order to have such an effect, it is preferable to contain 0.01% or more of V. More preferably 0.02% or more. On the other hand, even if a large amount of V is added, the strength-increasing effect of the portion exceeding 0.10% is small, and the alloy cost is increased. Therefore, the content of V is preferably 0.10% or less. More preferably 0.08% or less.

Cu: less than 0.50%

Cu is an element contributing to high strength by solid-solution strengthening and contributing to high strength by generation of the 2 nd phase, and may be added as necessary. In order to exhibit such an effect, the content is preferably 0.05% or more. More preferably 0.08% or more. On the other hand, even if the content exceeds 0.50%, the effect is saturated, and surface defects due to Cu are likely to occur, so the content is preferably 0.50% or less. More preferably 0.35% or less.

Ni: less than 0.50%

Like Cu, Ni is also an element contributing to high strength by solid solution strengthening and contributing to high strength by forming the 2 nd phase, and may be added as necessary. In order to exhibit such an effect, the content is preferably 0.05% or more. More preferably 0.08% or more. Further, if added simultaneously with Cu, it has an effect of suppressing surface defects due to Cu, and therefore, it is effective when Cu is added. On the other hand, even if the content exceeds 0.50%, the effect is saturated, and therefore, the content is preferably 0.50% or less. More preferably 0.35% or less.

Mo: less than 0.50%

Mo is an element that contributes to high strength by forming the 2 nd phase and contributes to high strength by forming a partial carbide, and may be added as needed. In order to exhibit such an effect, the content is preferably 0.05% or more. More preferably 0.08% or more. On the other hand, even if the content exceeds 0.50%, the effect is saturated, and therefore, the content is preferably 0.50% or less. More preferably 0.35% or less.

Cr: less than 0.80%

Cr is an element contributing to high strength by forming the 2 nd phase, and may be added as necessary. In order to exhibit such an effect, the content is preferably 0.10% or more. More preferably 0.13% or more. On the other hand, if the content exceeds 0.80%, the hot-dip galvanizability is lowered, and therefore, the plating is partially insufficient, and therefore, the content thereof is 0.80% or less. More preferably 0.70% or less.

Ca and/or REM: 0.0050% or less in total

Ca and REM (rare earth elements) are elements that contribute to spheroidizing the shape of the sulfide to improve the adverse effect on the delayed fracture resistance, and may be added as necessary. In order to exhibit such an effect, it is preferable to contain 0.0005% or more. More preferably 0.0008% or more. On the other hand, even if the content exceeds 0.0050%, the effect is saturated, and therefore, the content is set to 0.0050% or less. More preferably 0.0035% or less.

The balance other than the above is Fe and inevitable impurities. Examples of the inevitable impurities include Sb, Zn, Co, Sn, Zr, and the like, and the allowable ranges of the contents thereof include Sb: 0.01% or less, Zn: 0.01% or less, Co: 0.10% or less, Sn: 0.10% or less, Zr: less than 0.10%. In the present invention, the effect is not lost even if Ta or Mg is contained within the range of the usual steel composition.

Next, the microstructure of the plated steel sheet of the present invention will be described in detail. In the present invention, the composition is as follows: the steel contains 35-70% of ferrite by volume fraction, 12% or less (including 0%) of retained austenite by volume fraction, 15-60% of martensite by volume fraction, 30% or less (including 0%) of bainite by volume fraction and 5% or less (including 0%) of unrecrystallized ferrite by volume fraction as the remaining part, wherein the average crystal grain size of the ferrite is5 [ mu ] m or less, the average crystal grain size of the retained austenite is 2 [ mu ] m or less, the average crystal grain size of the martensite is 2 [ mu ] m or less, and the average crystal grain size of the bainite is 3 [ mu ] m or less. The volume fraction described here is the volume fraction of the entire steel sheet, and the same applies hereinafter.

The ferrite accounts for 35 to 70 percent in volume fraction

When the volume fraction of ferrite exceeds 70%, it is difficult to achieve tensile strength of 980MPa or more. Therefore, the volume fraction of ferrite is 70% or less. Preferably 65% or less, and more preferably 60% or less. When the volume fraction is less than 35%, the 2 nd phase transformation having a high dislocation density is large, and thus the delayed fracture resistance is deteriorated. Therefore, the volume fraction of ferrite is 35% or more. In order to increase the elongation, it is preferably 40% or more.

The ferrite has an average crystal grain diameter of 5 μm or less

If the average grain size of ferrite exceeds 5 μm, the grains are further coarsened during resistance welding, which deteriorates toughness and causes internal cracking. Therefore, the ferrite crystal grain size is5 μm or less. Preferably 4 μm or less. In order to increase the elongation, it is preferably 0.5 μm or more.

The retained austenite is 12% or less in volume fraction

The retained austenite contributes to strength by performing work to induce martensitic transformation. Further, the hydrogen trap sites are effective for the delayed fracture resistance. However, if martensite transformation is performed, the dislocation density is kept high, so that cracks are generated by hydrogen intrusion, and the delayed fracture resistance is deteriorated. Therefore, the volume fraction of retained austenite is 12% or less. Preferably, it exceeds 0% and is 10% or less. More preferably 1% or more. More preferably 7% or less. The volume fraction of the retained austenite may be 0%.

The retained austenite has an average crystal grain diameter of 2 μm or less

The average crystal grain size of the retained austenite is limited to 2 μm because martensite is easily generated during press forming due to the influence of the C distribution in the retained austenite, and the delayed fracture resistance is lowered. The lower limit is not particularly limited, and 0.3 μm or more is preferable because the contribution to the elongation is large if the lower limit is 0.3 μm or more.

Martensite accounts for 15-60 percent by volume fraction

In order to secure a desired strength, the volume fraction of martensite is 15% or more. Preferably 20% or more. More preferably 23% or more. On the other hand, if the volume fraction of martensite exceeds 60%, not only crack formation is likely to occur when hydrogen enters, but also the crack growth rate increases, so the upper limit thereof is 60%. Preferably 57% or less. More preferably 55% or less.

The martensite has an average crystal grain diameter of 2 μm or less

If the average grain size of martensite exceeds 2 μm, the grains are further coarsened during resistance welding, which deteriorates the toughness and causes internal cracking. Therefore, the average crystal grain size of martensite is 2 μm or less. Preferably 1.8 μm or less. The martensite here refers to martensite generated after annealing, and includes self-tempered (auto-temper) martensite in which martensite is transformed during cooling of annealing, tempered martensite in which tempering is performed after martensite transformation, and fresh martensite (fresh martentite) in which martensite is transformed from austenite without tempering.

The remainder being bainite in a volume fraction of 30% or less

Bainite contributes to an increase in strength, but when the volume fraction exceeds 30%, delayed fracture resistance is deteriorated because it contains a high dislocation density. Therefore, the upper limit is 30%. Preferably, it exceeds 0% and is 25% or less. More preferably 5% or more. More preferably 20% or less. The volume fraction of bainite may be 0%.

The average crystal grain diameter of bainite is less than 3 μm

When the average grain size of bainite exceeds 3 μm, the crystal grains are further coarsened during resistance welding, which causes deterioration in toughness and internal cracking, and therefore the average crystal grain size of bainite is 3 μm or less. Preferably 2.5 μm or less.

The remaining portion of unrecrystallized ferrite is 5% or less in volume fraction

Further, the unrecrystallized ferrite contributes to the increase in strength, but since it includes a high dislocation density as in the case of bainite, the upper limit thereof is 5%. Preferably more than 0% and not more than 3%. More preferably 1% or less. The volume fraction of unrecrystallized ferrite may be 0%.

In the present invention, pearlite may be formed in addition to ferrite, bainite, martensite, retained austenite and unrecrystallized ferrite, but the effects of the present invention can be achieved if the volume fraction of ferrite, bainite, martensite, retained austenite and unrecrystallized ferrite, the average crystal grain size of ferrite, bainite, martensite and retained austenite, and the distribution state of Ti or Nb-based precipitates (carbides) described above satisfy the ranges specified in the present invention. Among them, the volume fraction of pearlite is preferably 5% or less, and more preferably 3% or less.

Ti or Nb precipitates having an average grain size of 0.10 μm or less per 100 μm²The average number of the particles is more than 30

In the present invention, it is required to be in the range of 100 μm per unit²Contains 30 or more Ti or Nb precipitates having an average grain diameter of 0.10 μm or less on average. This is because Ti or Nb precipitates serve as trap sites for hydrogen, thereby improving the delayed fracture resistance, and further, are effective for grain refinement, thereby improving the resistance to solder cracking. If the grain size exceeds 0.10. mu.m, or if the above precipitates are per 100. mu.m²If the average number of the particles is less than 30, the delayed fracture resistance and the solder resist cracking resistance are lowered. Preferably at every 100 μm²More than 50 in the total number. Further preferably in an amount of 100 μm²More than 60 in the total number. Specific examples of the Ti or Nb-based precipitates include carbides.

Next, the production conditions of the plated steel sheet of the present invention will be described.

A billet having the above-mentioned composition (chemical composition) is hot-rolled under the condition that the finish temperature of finish rolling is 850 to 950 ℃, is cooled 1 times to 680 ℃ at a 1 st average cooling rate of 75 ℃/s or more, is then cooled 2 times to 400 to 580 ℃ at a 2 nd average cooling rate of 5 ℃/s or more, is then coiled, is subjected to pickling and cold rolling, is then, in an annealing step, a cold-rolled steel sheet is heated to a temperature region of 760 to 900 ℃ at an average heating rate of 3 to 30 ℃/s as a 1 st soaking temperature, is held for 15 seconds or more in a temperature region of 760 to 900 ℃ and is then cooled to a temperature region of 600 ℃ or less at an average cooling rate of 3 to 30 ℃/s to be annealed, is then subjected to hot dip galvanizing treatment, and cooling to room temperature.

In the hot rolling step, the slab is preferably hot-rolled at 1150 to 1300 ℃ without reheating after casting, or hot-rolled after reheating to 1150 to 1300 ℃. The billet used is preferably produced by a continuous casting method in order to prevent macro-segregation of the components, but may be produced by an ingot casting method or a thin slab casting method. In the present invention, in addition to the conventional method of cooling down to room temperature once after billet production and then reheating, an energy saving process such as direct rolling or direct rolling in which a billet is directly charged into a heating furnace in a state of a warm sheet without cooling, or rolling is performed immediately after holding heat, or rolling is performed directly after casting can be applied without any problem.

[ Hot Rolling Process ]

Finish rolling finish temperature: 850-950 DEG C

In order to improve the delayed fracture resistance and the resistance to solder cracking after annealing by uniformizing the structure in the steel sheet and reducing the anisotropy of the material, it is necessary to finish hot rolling in an austenite single-phase region. Therefore, the finish rolling finishing temperature is 850 ℃ or higher. On the other hand, if the finish rolling temperature exceeds 950 ℃, the hot rolled structure becomes coarse, and the crystal grains after annealing also become coarse. Therefore, the finish rolling temperature is 850 to 950 ℃.

Cooling conditions after finish Rolling

Cooling to 680 deg.C or below at 1 st average cooling rate of 75 deg.C/s or above as 1 cooling, and cooling to 400-580 deg.C at 2 nd average cooling rate of 5 deg.C/s or above as 2 cooling

In the present invention, the microstructure of the steel sheet after annealing is controlled by controlling the precipitation form of precipitates of Ti or Nb during hot rolling, and therefore, cooling after finish rolling is an important step. After the hot rolling is completed, austenite undergoes ferrite transformation during cooling, but the ferrite coarsens at a high temperature. Therefore, by rapidly cooling the steel sheet after the hot rolling is completed, the formation of precipitates is controlled while the structure is homogenized as much as possible. Therefore, as 1 cooling, the cooling is performed at the 1 st average cooling rate of 75 ℃/s or more to 680 ℃ or less.

When the 1 st average cooling rate is less than 75 ℃/s, ferrite coarsens, and therefore the steel sheet structure of the hot-rolled steel sheet becomes inhomogeneous, and the resistance to solder cracking is lowered. In addition, a desired average crystal grain size cannot be obtained for ferrite, martensite, and bainite. Preferably 85 ℃/s or more.

If the temperature of cooling in 1-time cooling exceeds 680 ℃, pearlite excessively forms in the steel sheet structure of the hot-rolled steel sheet, and the steel sheet structure becomes inhomogeneous, and therefore, the resistance to solder cracking is reduced. In addition, in the case of martensite, a desired average crystal grain size cannot be obtained. Preferably 650 ℃ or lower. Further, since martensite excessively increases in the steel sheet structure of the hot-rolled steel sheet, it is preferably 400 ℃ or higher.

The subsequent 2 times of cooling is carried out at a 2 nd average cooling rate of 5 ℃/s or more to a temperature in the range of 400 to 580 ℃. When the temperature is cooled to less than 5 ℃/s or more than 580 ℃, ferrite or pearlite is excessively generated in the steel sheet structure of the hot-rolled steel sheet, and the resistance to solder cracking after annealing is lowered. Further, martensite cannot have a desired average crystal grain size, nor cannot be obtained at 100 μm per unit²Thereby obtaining Ti or Nb precipitates having an average crystal grain size of not more than 30 on average and not more than 0.10 μm. Preferably 10 ℃/s or more. Further, Ti and Nb are dissolved in an excessive amount, and therefore, it is preferably 65 ℃/s or less.

Coiling temperature: 400-580 deg.C

When the coiling temperature exceeds 580 ℃, ferrite and pearlite in the steel sheet structure of the hot-rolled steel sheet are excessively generated. In addition, the residual austenite cannot have a desired average crystal grain size, nor can it be 100 μm²Thereby obtaining Ti or Nb precipitates having an average crystal grain size of not more than 30 on average and not more than 0.10 μm. Therefore, the upper limit of the winding temperature is preferably 580 ℃. More preferably 550 ℃ or lower. Further, when the coiling temperature is less than 400 ℃, Ti and Nb precipitates cannot be sufficiently precipitated and become a solid solution state, and therefore, the effect of refining after annealing cannot be expected. In addition, it cannot be set at every 100 μm²Thereby obtaining Ti or Nb precipitates having an average crystal grain size of not more than 30 on average and not more than 0.10 μm. Therefore, the coiling temperature is preferably 400 ℃ or higher. More preferably 420 ℃ or higher.

[ Pickling step ]

After the hot rolling step, it is preferable to perform an acid pickling step to remove scale on the surface layer of the hot-rolled sheet. The acid washing step is not particularly limited as long as it is carried out according to a conventional method.

[ Cold Rolling Process ]

A cold rolling step of rolling the steel sheet into a cold rolled sheet having a predetermined thickness is performed. The cold rolling step is not particularly limited as long as it is carried out by a conventional method. The preferable range of the reduction ratio in the cold rolling is 30% to 95%.

[ annealing step ]

In the annealing step, recrystallization is performed and fine bainite, retained austenite, and martensite are formed in the steel sheet structure for higher strength. Therefore, in the annealing step, the steel sheet is heated to a temperature range of 760 ℃ to 900 ℃ at an average heating rate of 3 to 30 ℃/s, is held in the temperature range of 760 ℃ to 900 ℃ as a soaking temperature for 15 seconds or more, and is then cooled to a temperature range of 600 ℃ or less at an average cooling rate of 3 to 30 ℃/s.

Note that temper rolling may be performed after annealing. The preferable range of the elongation is 0.05% to 2.0%.

Average heating rate: 3-30 ℃/s

By setting the average heating rate to 3 to 30 ℃/s, the crystal grains after annealing can be refined. If the heating is rapidly performed, recrystallization is difficult to proceed. In addition, the bainite does not have a desired average crystal grain size, and a desired volume fraction and a desired amount of unrecrystallized ferrite per 100 μm are not obtained²In which 30 or more Ti or Nb precipitates having an average crystal grain size of 0.10 μm or less are contained in the alloy. Therefore, the upper limit of the average heating rate is 30 ℃/s. Since the amount of unrecrystallized ferrite increases, it is preferably 25 ℃/s or less.

If the heating rate is too low, ferrite and martensite grains become coarse and a predetermined average grain size cannot be obtained. In addition, it cannot be set at every 100 μm²Thereby obtaining Ti or Nb precipitates having an average crystal grain size of not more than 30 on average and not more than 0.10 μm. Therefore, an average heating rate of 3 ℃/s or more is required. Preferably 5 deg.CMore than s.

Soaking temperature (holding temperature): 760 ℃ to 900 DEG C

The soaking temperature is a temperature region belonging to a 2-phase region of ferrite and austenite or a single-phase region of austenite. When the temperature is less than 760 ℃, the ferrite fraction increases, and thus it is difficult to secure strength. In addition, a desired average crystal grain size cannot be obtained for ferrite and martensite. Therefore, the lower limit of the soaking temperature is 760 ℃. Preferably above 780 ℃. If the soaking temperature is too high, the grain growth of ferrite, martensite and austenite becomes remarkable, and the resistance to solder cracking is lowered due to coarsening of the grains. In addition, it cannot be set at every 100 μm²Thereby obtaining Ti or Nb precipitates having an average crystal grain size of not more than 30 on average and not more than 0.10 μm. Therefore, the upper limit of the soaking temperature is 900 ℃. Preferably below 880 ℃.

Soaking time: for 15 seconds or more

At the soaking temperature, the soaking time needs to be maintained for 15 seconds or more for recrystallization to proceed and for some or all of the austenite to be transformed. The volume fraction of unrecrystallized ferrite is preferably increased to 20 seconds or more. The upper limit is not particularly limited, but is preferably within 600 seconds.

Cooling conditions at the time of annealing: cooling to a temperature region of 600 ℃ or lower at an average cooling rate of 3 to 30 ℃/s

After the soaking, the steel sheet needs to be cooled from the soaking temperature to a temperature range of 600 ℃ or less (cooling stop temperature) at an average cooling rate of 3 to 30 ℃/s. When the average cooling rate is less than 3 ℃/s, ferrite transformation proceeds during cooling, and the volume fraction of the 2 nd phase decreases, so that it is difficult to secure strength. In addition, the desired average crystal grain size cannot be obtained for martensite, retained austenite and bainite. On the other hand, if the average cooling rate exceeds 30 ℃/s, not only martensite is excessively generated, but also it is difficult to realize it in equipment. When the cooling stop temperature exceeds 600 ℃, pearlite is excessively produced, and therefore a predetermined volume fraction in the microstructure of the steel sheet cannot be obtained, and it is difficult to ensure strength. In addition, it cannot be set at every 100 μm²Thereby obtaining Ti or Nb precipitates having an average crystal grain diameter of 0.10 μm or less in an average of 30 or more, and reducing delayed fracture resistance and solder crack resistance. The average cooling rate is an average value of the cooling rates in the range from 600 ℃ or lower to the immersion in the plating bath, and the average cooling rate may be maintained at 3 to 30 ℃/s in the temperature range.

[ plating treatment ]

After the annealing, the plating treatment is performed, and the temperature is cooled to room temperature. The plating treatment includes hot dip galvanizing treatment, electrogalvanizing treatment, and the like. For example, in the case of hot-dip galvanizing treatment, the temperature of the steel sheet immersed in the plating bath is preferably from (hot-dip galvanizing bath temperature-40) ° c to (hot-dip galvanizing bath temperature +50) ° c. If the temperature of the steel sheet immersed in the plating bath is lower than (the temperature of the hot dip galvanizing bath-40) ° c, part of the molten zinc may solidify when the steel sheet is immersed in the plating bath, deteriorating the plating appearance. Therefore, the lower limit is set to (hot-dip galvanizing bath temperature-40) DEG C. Further, if the temperature of the steel sheet immersed in the plating bath exceeds (hot dip galvanizing bath temperature +50) ° c, the temperature of the plating bath rises, which causes a problem in mass productivity. As the other conditions, the conditions performed in the usual plating treatment may be used.

[ alloying treatment ]

After the plating, the plating may be alloyed in a temperature range of 450 to 600 ℃. By alloying in a temperature range of 450 to 600 ℃, the Fe concentration in the plating is 7 to 15 mass%, and the adhesion of the plating and the corrosion resistance after coating are improved. When the temperature is less than 450 ℃, alloying does not sufficiently proceed, and a decrease in sacrificial corrosion resistance and a decrease in sliding property are caused, and when the temperature is higher than 600 ℃, alloying proceeds excessively, and powdering resistance is lowered.

The conditions of the other production methods are not particularly limited, and when the hot dip galvanizing treatment is performed, a series of treatments such as the above-described annealing, plating treatment, alloying treatment of plating, and the like are preferably performed in a continuous hot dip galvanizing line (CGL) from the viewpoint of productivity. In addition, the hot dip galvanizing preferably uses a zinc plating bath containing 0.10 to 0.20 mass% of Al. After plating, wiping may be performed in order to adjust the weight per unit area of plating.

(examples)

Hereinafter, examples of the present invention will be described.

However, the present invention is not limited to the following examples, and can be carried out by appropriately changing the scope of the present invention, and these are included in the technical scope of the present invention.

Steels having the composition shown in table 1 were melted and cast to produce slabs, and hot-rolled at a hot-rolling heating temperature of 1250 ℃ and a finish rolling temperature (FDT) shown in table 2 to produce thicknesses: after 3.2mm hot-rolled steel sheet, it was cooled to the 1 st cooling temperature at the 1 st average cooling rate (cooling rate 1) shown in Table 2, and then cooled to the 2 nd cooling temperature at the 2 nd average cooling rate (cooling rate 2), and then coiled at the Coiling Temperature (CT). Subsequently, the obtained hot-rolled sheet was subjected to acid washing and cold rolling to produce a cold-rolled sheet (sheet thickness: 1.4 mm).

The cold-rolled steel sheets thus obtained were subjected to annealing treatment under the production conditions shown in table 2 in a continuous hot-dip galvanizing line, hot-dip galvanized treatment, and then further subjected to alloying treatment at the temperatures shown in table 2, thereby obtaining alloyed hot-dip galvanized steel sheets (GA). Here, the plating treatment is performed at a zinc plating bath temperature: al concentration of zinc plating bath at 460 ℃: 0.14 mass% (when alloying treatment was performed), 0.18 mass% (when alloying treatment was not performed), and a plating adhesion amount per one surface of 45g/m²(double-sided plating). Some of the steel sheets were not alloyed with zinc plating, and a non-alloyed hot dip galvanized steel sheet (GI) was produced.

Tensile test pieces of JIS5 were taken from the produced steel sheets so that the rolling direction at right angles became the longitudinal direction (tensile direction), and the Tensile Strength (TS) was measured by a tensile test (JIS Z2241 (1998)).

For the delayed fracture test, a test piece obtained by cutting the obtained cold-rolled steel sheet into 30mm × 100mm long sides in the rolling direction and grinding the end face was used, and the test piece was bent at 180 ° with a curvature radius of 10mm at the tip of the punch. The test piece subjected to the bending was fastened with bolts so that the inner space became 20mm, and after the test piece was subjected to a stress, the test piece was immersed in hydrochloric acid having a pH of 1.5 at 20 ℃. The test piece was good (o) when no crack occurred in 100 hours, and poor (x) when a crack occurred in the test piece.

In the resistance welding cracking test, resistance welding (spot welding) was performed using a test piece cut into 50 × 150mm long pieces in a direction perpendicular to the rolling direction of the obtained cold-rolled steel sheet. For welding, resistance spot welding was performed on a plate group on which 2 steel plates were stacked, with the plate group being tilted by 4 ° by a servomotor-pressurized single-phase direct current (50Hz) resistance welding machine attached to a welding gun. The welding conditions were set to a pressing force of 3.5kN and a holding time of 0.36 seconds. The welding current and the welding time were adjusted so that the diameter of the spot weld became 5.9 mm. After welding, the test piece was cut in half, and the cross section was observed with an optical microscope, and the resistance weld cracking resistance was good when no crack of 0.2mm or more was observed (o), and the resistance weld cracking resistance was poor when no crack of 0.2mm or more was observed (x).

The volume fractions of ferrite, martensite, and unrecrystallized ferrite in the steel sheet were determined by grinding a sheet thickness cross section parallel to the rolling direction of the steel sheet, etching the cross section with 3 vol% nitroethanol, observing the cross section with a magnification of 2000 times or 5000 times using an SEM (scanning electron microscope), measuring the area fraction by a point counting method (according to ASTM E562-83 (1988)), and determining the area fraction as the volume fraction. The average grain sizes of ferrite and martensite can be obtained by calculating the area of each phase and the equivalent circle diameter thereof by previously recognizing the picture of each ferrite and martensite crystal grain from the picture of the steel sheet microstructure using the Image-Pro input of Media Cybernetics, and averaging these values.

The volume fraction of retained austenite was determined by grinding the steel sheet to 1/4 planes in the sheet thickness direction and determining the intensity of diffracted X-rays at 1/4 planes. The integrated intensities of X-ray diffraction lines on the {200} plane, {211} plane, {220} plane of ferrite, and the {200} plane, {220} plane and {311} plane of austenite, of iron were measured by an X-ray diffraction method (apparatus: RINT2200, manufactured by Rigaku corporation) at an acceleration voltage of 50keV using Mo Kalpha rays as a radiation source, and the volume fractions of retained austenite were determined from the calculation formulas described in "X-ray diffraction Manual" (2000), physical and electrical machines corporation, p.26, and 62-64, using the measured values.

The average crystal grain size of the retained austenite was determined by observing the grain size at 5000-fold magnification using EBSD (electron beam back scattering diffraction), calculating the equivalent circle diameter using Image-Pro, and averaging the values.

Further, the steel sheet structure was observed by SEM, TEM (transmission electron microscope), and FE-SEM (field emission scanning electron microscope), and bainite was observed, and the volume fraction was determined in the same manner as described above. The average crystal grain size of bainite was also determined by calculating the equivalent circle diameter from the microstructure photograph of the steel sheet using the above Image-Pro and averaging these values.

The particle size of Ti or Nb precipitates was determined by observing the particle size at 5000, 10000, and 20000 times using SEM and TEM, and calculating the equivalent circle diameter using Image-Pro. The number of Ti or Nb precipitates was observed at 5000 times, 10000 times, and 20000 times using SEM and TEM, and the average number of 10 sites was determined.

The measurement results of the tensile properties, delayed fracture resistance, solder crack resistance, and steel sheet structure are shown in table 3.

According to the results shown in Table 3, the examples of the present invention each had a composite structure as follows: the tensile strength of 980MPa or more was confirmed to be obtained, and as a result, the steel sheet was not broken at 100 hours in the delayed fracture characteristic evaluation test, had excellent delayed fracture resistance, and was not cracked inside after resistance welding, and excellent resistance to solder resist fracture characteristics were obtained.

Claims

the ferrite has an average crystal grain diameter of 5 μm or less,

the martensite has an average crystal grain diameter of 2 [ mu ] m or less,

the average crystal grain diameter of the bainite is less than 3 mu m,

2. The plated steel sheet according to claim 1, further comprising, in mass%, a component selected from the group consisting of V: 0.10% or less, Cu: 0.50% or less, Ni: 0.50% or less, Mo: 0.50% or less, Cr: 0.80% or less and Ca and/or REM: 0.0050% or less, and 1 or 2 or more.

performing plating treatment on the annealed cold-rolled steel sheet;

the steel billet comprises the following components in percentage by mass: 0.05-0.22%, Si: 0.05-1.80%, Mn: 1.45% -3.35%, P: 0.05% or less, S: 0.005% or less, Al: 0.01% -0.10%, N: 0.010% or less and B: 0.0002% -0.0045%, and further contains, in mass%, a Ti: 0.005% -0.090%, Nb: more than 1 of 0.005-0.090%, the balance being Fe and inevitable impurities,

the plated steel sheet has the following structure: contains 35 to 70% of ferrite by volume fraction, 12% or less of retained austenite by volume fraction, 15 to 60% of martensite by volume fraction, and 30% or less of bainite by volume fraction and 5% or less of unrecrystallized ferrite by volume fraction as the remaining portion,

the ferrite has an average crystal grain diameter of 5 μm or less,

the martensite has an average crystal grain diameter of 2 [ mu ] m or less,

the average crystal grain diameter of the bainite is less than 3 mu m,

4. The method for producing a plated steel sheet according to claim 3, wherein the composition further contains, in mass%, a component selected from the group consisting of V: 0.10% or less, Cu: 0.50% or less, Ni: 0.50% or less, Mo: 0.50% or less, Cr: 0.80% or less and Ca and/or REM: 0.0050% or less, and 1 or 2 or more.

5. The method for producing a plated steel sheet according to claim 3 or 4, wherein the alloying treatment for plating is performed in a temperature range of 450 to 600 ℃ after the plating treatment is performed.