Hot-dip galvanized steel sheet and method for producing same

Abstract

The invention provides a hot-dip galvanized steel sheet and a manufacturing method thereof. The present invention relates to the production of a high-strength hot-dip galvanized steel sheet having a tensile strength of 980 MPa grade or more for an automobile structural member, and provides a hot-dip galvanized steel sheet having a hole expandability (HER) of 45% or more, a relational expression HER # imgabs0 # 100 / YS of yield strength (YS) and hole expandability (HER) of 5 or more, a relational expression HER # imgabs1 # EL of elongation (EL) and hole expandability (HER) of 700 or more, a yield ratio (YS / TS) of 0.80 or more, and a tensile strength of 980 MPa grade or more. LME cracks are not generated on the surface and inside of the welded part.

Description

TECHNICAL FIELD

[0001] The present invention relates to the manufacture of hot-dip galvanized steel sheets for automotive structural components, and more particularly, to a high-strength hot-dip galvanized steel sheet having excellent ductility, hole expansion property, and weldability, in which LME cracks do not occur on the surface and inside of a welded portion, and a method for manufacturing the same. BACKGROUND ART

[0002] In recent years, regulations for protecting the global environment have been gradually strengthened in the automotive industry. Therefore, regulations on carbon emission reduction and fuel efficiency have become increasingly strict, and in order to ensure the safety of passengers in the event of an accident such as a collision, the use of high-strength steel has been increasing. In addition, high-strength steel having excellent yield strength is used in structural components such as members, seat rails, and pillars to improve the impact resistance of the vehicle body. However, the increase in the strength of the steel sheet may cause a decrease in ductility and formability. In order to solve this problem, it is necessary to develop a material that satisfies both high strength and high formability. Generally, as the strength of the steel sheet increases, the elongation rate decreases, and there is a problem in that the workability decreases, so it is necessary to develop a material that can compensate for this problem. Generally, as a method for strengthening steel, solid solution strengthening, precipitation strengthening, strengthening by grain refinement, transformation strengthening, etc. have been studied. However, in the above methods, there is a problem in that it is very difficult to manufacture high-strength steel having a tensile strength of 490 MPa or more using solid solution strengthening and grain refinement.

[0003] Precipitation-strengthened high-strength steel is a technique for ensuring strength by adding carbon and nitride forming elements such as Nb, Ti, and V, precipitating carbon and nitrides, and suppressing grain growth caused by fine precipitates to refine grains. Compared with the low manufacturing cost, the technique has the advantage of easily ensuring high strength, but the fine precipitates cause a rapid increase in the recrystallization temperature, so there is a disadvantage in that high-temperature annealing is required to cause sufficient recrystallization to ensure ductility. In addition, there is a problem in that it is difficult to obtain high-strength steel having a strength of 600 MPa or more with precipitation-strengthened steel in which carbon and nitrides are precipitated on a ferrite matrix.

[0004] In addition, various transformation-induced high-strength steels have been developed. For example, Dual Phase (DP) steel composed of a soft ferrite matrix and a hard martensite duplex, Transformation Induced Plasticity (TRIP) steel that ensures high ductility by using the transformation-induced plasticity of retained austenite, or Complexed Phase (CP) steel composed of a composite structure of ferrite and hard bainite or martensite. In recent years, automotive steel sheets require higher-strength steel sheets to improve fuel efficiency or durability. In terms of crash safety and passenger protection, the demand for ultra-high-strength steel sheets with a tensile strength of 1180 MPa or more for use as body structures or reinforcements is increasing. Among them, DP steel has excellent ductility and is the most commonly used automotive steel sheet, but it has the disadvantages of a low yield ratio (YR) and poor formability and processability. In addition, with the trend of gradually increasing the strength of steel sheets, cracks or wrinkles are generated during the stamping process of automotive parts, making it difficult to manufacture complex parts. In the case of TRIP and XF steels among ultra-high-strength steels, compared with existing DP steels, they have higher ductility and an excellent yield ratio, so they have good processability. However, in order to ensure a high elongation rate, a large amount of Si and Al are added, resulting in poor weldability. To overcome the disadvantages of these existing ultra-high-strength steels, by reducing the Si and Al component systems, steels with a good weldability component system and a yield ratio above a certain level are manufactured, thereby expanding the application of ultra-high-strength steels to more complex parts. This can be achieved by using Quenching and Partitioning (Q&P) heat treatment, which is the latest heat treatment technology that can ensure retained austenite.

[0005] As an existing technology for simultaneously ensuring the ductility and processability of the high-tensile steel sheet, the invention disclosed in Patent Document 1 can be cited. In the invention disclosed in Patent Document 1, it is necessary to control the heat treatment temperature so that retained austenite is formed by precisely controlling the slow cooling and rapid cooling temperatures to ensure a high elongation rate.

[0006] As another existing technology, the invention disclosed in Patent Document 2 can be cited. The invention disclosed in Patent Document 2 is characterized by manufacturing a steel sheet with a high yield ratio through the Q&P process and painting treatment.

[0007] As another existing technology, the invention disclosed in Patent Document 3 can be cited. The invention disclosed in Patent Document 3 provides a high-strength cold-rolled steel sheet with a high fraction of bainite manufactured by cooling to the bainite region. Compared with the Q&P process, the carbon partitioning effect is reduced, resulting in a problem of poor elongation.

[0008] [Prior Art Documents]

[0009] [Patent Documents]

[0010] (Patent Document 1) Korean Patent Publication No. KR2018 - 0165176

[0011] (Patent Document 2) Japanese Patent Publication No. JP2010 - 090432

[0012] (Patent Document 3) Japanese Patent Publication No. JP2015 - 106696 SUMMARY OF THE INVENTION

[0013] (1) Technical Problem to be Solved

[0014] An object of the present invention is to provide a hot - dip galvanized steel sheet mainly used for automotive structural components and a manufacturing method thereof.

[0015] In addition, the technical problems to be achieved by the present invention are not limited to the above - mentioned technical problems, and those skilled in the art to which the present invention pertains can clearly understand other technical problems not mentioned from the following description.

[0016] (2) Technical Solution

[0017] Therefore, one aspect of the present invention relates to a hot - dip galvanized steel sheet which, by weight %, comprises: carbon (C): 0.08 - 0.16%, silicon (Si): 0.8% or less (except 0%), manganese (Mn): 2.0 - 3.0%, molybdenum (Mo): 0.4% or less (except 0%), chromium (Cr): 1.0% or less (except 0%), phosphorus (P): 0.1% or less (except 0%), sulfur (S): 0.02% or less (except 0%), aluminum (sol.Al): 0.6% or less (except 0%), titanium (Ti): 0.001 - 0.04%, niobium (Nb): 0.001 - 0.04%, nitrogen (N): 0.01% or less (except 0%), boron (B): 0.01% or less (except 0%), the balance being Fe and other inevitable impurities, wherein C, Si, Mn, Cr and Mo satisfy the following Relationship 1. By area %, the hot - dip galvanized steel sheet has a microstructure comprising a total of 70% or more of bainite and tempered martensite, 10% or less of ferrite, and the balance being fresh martensite and retained austenite. The hole expansion ratio (HER) of the hot - dip galvanized steel sheet is 45% or more, and the relationship between the yield strength (YS) and the hole expansion ratio (HER) is HER 100 / YS is 5 or more, and the yield ratio (YS / TS) is 0.80 or more.

[0018] [Relationship 1]

[0019] C + Si / 30 + Mn / 20 + (Cr + Mo) / 5 ≥ 0.3

[0020] Among them, C, Si, Mn, Cr, and Mo represent the content percentages of steel components in the matrix structure at the t / 4 position of the thickness of the base steel plate.

[0021] In addition, another aspect of the present invention relates to a method for manufacturing a hot-dip galvanized steel sheet, which includes the following processes: preparing a steel billet that satisfies the steel composition components and relational expression 1, and then reheating the steel billet; hot-rolling the reheated steel billet so that the finish rolling exit temperature is Ar3 to Ar3 + 50 °C, then coiling at 400 - 650 °C and cooling to room temperature at an average cooling rate of 0.1 °C / second or less; cold-rolling the cooled hot-rolled steel sheet at a reduction rate of 40 - 70% to manufacture a cold-rolled steel sheet; continuously annealing the cold-rolled steel sheet at a temperature of Ac3 - 20 °C to Ac3 + 20 °C; performing primary cooling on the continuously annealed cold-rolled steel sheet at an average cooling rate of 10 °C / second or less, cooling to a temperature range of 500 - 550 °C, and performing secondary cooling at an average cooling rate of 5 °C / second or more, cooling to a temperature of 300 - 340 °C; reheating the cold-rolled steel sheet after the secondary cooling to a temperature above Ms, and then holding for 60 seconds or more; and cooling the held cold-rolled steel sheet to a temperature below 150 °C at an average cooling rate of 5 °C / second or more.

[0022] (III) Beneficial effects

[0023] As described above, in the present invention, the composition and manufacturing process are optimized, so as to satisfy the high ductility as the characteristics of DP steel, and compared with the existing DP steel, it has the effect of manufacturing a high-strength hot-dip galvanized steel sheet with excellent yield ratio (YS / TS) and hole expansion property. Specifically, it is possible to provide a high-strength hot-dip galvanized steel sheet with a hole expansion property (HER) of 45% or more, a relational expression HER between the yield strength (YS) and the hole expansion property (HER) of 100 / YS of 5 or more, a relational expression HER between the elongation (EL) and the hole expansion property (HER) of EL of 700 or more, and a yield ratio (YS / TS) of 0.80 or more. Therefore, by preventing processing defects such as cracks generated during stamping forming, it can be used in various ways for automotive structural components with complex shapes that require high formability. In addition, by reducing Si and Al, it also has the advantage of good weldability compared with the existing TRIP steel. 100 / YS of 5 or more, a relational expression HER between the elongation (EL) and the hole expansion property (HER) of HER EL of 700 or more, a high-strength hot-dip galvanized steel sheet with a yield ratio (YS / TS) of 0.80 or more. Therefore, by preventing processing defects such as cracks generated during stamping forming, it can be used in various ways for automotive structural components with complex shapes that require high formability. In addition, by reducing Si and Al, it also has the advantage of good weldability compared with the existing TRIP steel. Description of the drawings

[0024] Figure 1 It shows the relational expression HER between the yield strength (YS) and the hole expansion property (HER) of the inventive steel and the comparative steel according to relational expression 1 in the embodiments of the present invention Graph of the change in 100 / YS.

[0025] Figure 2 It is a graph showing the change in the hole expansion ratio (HER) according to Relationship 1 of the inventive steel and the comparative steel in the embodiments of the present invention.

[0026] Figure 3 It is a graph showing the change in the yield ratio according to Relationship 1 of the inventive steel and the comparative steel in the embodiments of the present invention.

[0027] Figure 4 It is a graph showing the relationship HER of the elongation rate (EL) and the hole expansion ratio (HER) according to Relationship 1 of the inventive steel and the comparative steel in the embodiments of the present invention of the change in EL. Best Mode

[0028] The technical terms used herein are for illustrative purposes only for specific embodiments and are not intended to limit the present invention. Unless the contrary meaning is clearly shown herein, the singular forms used herein include the plural forms. The meaning of "comprising" or "including" used in the specification specifically describes specific features, regions, integers, steps, actions, elements, and / or components, and does not exclude the existence or addition of other specific features, regions, integers, steps, actions, elements, components, and / or groups.

[0029] Unless otherwise defined, all terms, including technical and scientific terms, used herein have the same meaning as commonly understood by those of ordinary skill in the technical field to which the present invention pertains. Further interpreted as, terms defined in commonly used dictionaries have meanings consistent with the relevant technical literature and the currently disclosed content, and should not be interpreted as ideal or overly formal meanings unless explicitly defined.

[0030] The present inventors confirmed through experiments that by optimizing the composition and manufacturing process, when introducing more than 70% tempered martensite or bainite, the balance of fresh martensite, and retained austenite in the final fine structure, compared with existing DP steels, the workability can be improved by increasing the yield ratio, and less than 5% retained austenite is formed during final cooling, thereby further improving the ductility. Based on this experimental result, the present invention was completed.

[0031] First, the inventor made a slab with the C+Si / 30+Mn / 20+(Cr+Mo) / 5 of the steel components in the matrix structure at the t / 4 position of the base steel plate controlled below 0.3, hot-rolled it so that the finish rolling outlet temperature was Ar3 to Ar3+50°C, coiled it at 400 - 650°C, and then cooled it to room temperature at an average cooling rate of 0.1°C or less, thereby manufacturing a hot-rolled steel plate with carbides finely dispersed during annealing, and the carbides were the nucleation sites of austenite. In addition, the hot-rolled steel plate was cold-rolled at a reduction rate of 40 - 70% to manufacture a cold-rolled steel plate, then continuously annealed at a temperature of Ac3 - 20°C to Ac3+20°C, and then the continuously annealed steel plate was cooled at an average cooling rate of 10°C or less for the first cooling to a temperature range of 500°C to 550°C, cooled at an average cooling rate of 5°C or more for the second cooling to a temperature range of 300°C to 400°C, and fresh martensite was introduced. After that, it was reheated to a temperature above Ms, and then held for 60 seconds or more to form tempered martensite and bainite, with carbon enriched in the surrounding non-transformed austenite, and then cooled at an average cooling rate of 5°C or more to a temperature below 150°C, so that fine fresh martensite was introduced into the remainder.

[0032] By controlling the above components and manufacturing process, compared with the existing DP steel, the fractions of ferrite and fresh martensite are reduced. By introducing tempered martensite, bainite, and retained austenite, compared with the existing DP steel, the yield ratio is increased, so that workability can be ensured. In addition, a large number of mobile dislocations are generated around the retained austenite during plastic deformation, which helps to improve ductility. Compared with the existing DP steel, this precisely controlled composite structure steel can maintain a high yield ratio while ensuring ductility. Thus, an ultra-high strength hot-dip galvanized steel plate with excellent ductility, hole expansion property, and weldability can be manufactured.

[0033] Hereinafter, with reference to various embodiments and drawings, the technical features according to the present invention will be described in more detail.

[0034] First, in terms of weight %, the hot-dip galvanized steel sheet with excellent ductility, hole expansion property, and weldability of the present invention contains: carbon (C): 0.08 - 0.16%, silicon (Si): less than 0.8% (except 0%), manganese (Mn): 2.0 - 3.0%, molybdenum (Mo): less than 0.4% (except 0%), chromium (Cr): less than 1.0% (except 0%), phosphorus (P): less than 0.1% (except 0%), sulfur (S): less than 0.02% (except 0%), aluminum (sol.Al): less than 0.6% (except 0%), titanium (Ti): 0.001 - 0.04%, niobium (Nb): 0.001 - 0.04%, nitrogen (N): less than 0.01% (except 0%), boron (B): less than 0.01% (except 0%), and the balance of Fe and other inevitable impurities. Hereinafter, the characteristics of each alloying element and the critical significance of the composition range will be briefly described. At this time, unless otherwise specified, the content of each component represents weight %.

[0035] C: 0.08 - 0.16%

[0036] Carbon (C) is a very important element added for the strengthening of the phase transformation structure. Carbon increases the strength by promoting the formation of hard martensite in the duplex steel. As the carbon content increases, the amount of martensite increases. However, when the amount of carbon (C) exceeds 0.16%, the strength of martensite increases, but the strength difference from ferrite with a low carbon concentration increases. Due to this strength difference, the phase interface is prone to failure during plastic deformation, so the ductility and work hardening rate decrease. In addition, due to poor weldability, welding defects occur during the processing of the components of the customer company. On the other hand, when the carbon content is as low as less than 0.08%, it is difficult to ensure the desired strength, so the content of carbon is preferably limited to 0.08 - 0.16%. More preferably, the content of the carbon can be limited to the range of 0.09 - 0.15%.

[0037] Si: less than 0.8% (except 0%)

[0038] Silicon (Si) is a ferrite stabilizing element and is an element that promotes ferrite phase transformation and enriches carbon in the untransformed austenite in the Q&P process, thus contributing to the formation of retained austenite. In addition, the silicon (Si) is a useful element that effectively reduces the hardness difference between phases by solid solution strengthening to improve the strength of ferrite, and can ensure the strength without reducing the ductility of the steel sheet. However, when the silicon (Si) exceeds 0.8%, it causes surface scale defects, has an adverse effect on the surface quality, and reduces the weldability and chemical conversion treatment property. Therefore, the addition amount of silicon is preferably limited to less than 0.8%, and more preferably, it can be limited to less than 0.6%.

[0039] Mn: 2.0 - 3.0%

[0040] Manganese (Mn) is an element that refines particles without reducing ductility, completely precipitates sulfur (S) in steel as MnS, prevents hot brittleness caused by the formation of FeS, and strengthens the steel. At the same time, in duplex steel, it plays a role in reducing the critical cooling rate for obtaining the martensite phase, making it easier to form martensite. When the content of manganese is less than 2.0%, it is difficult to ensure the strength desired in the present invention. On the other hand, when the content of manganese exceeds 3.0%, there is a high possibility of problems such as weldability and hot rolling properties, and excessive martensite is formed, making the material unstable and forming manganese bands (Mn-Bands, bands of Mn oxides) in the structure, so there is a problem of an increased risk of processing cracks and plate fracture. In addition, Mn oxides are dissolved on the surface during annealing, so there is a problem of greatly hindering the surface quality. Therefore, in the present invention, it is preferable to limit the content of Mn to the range of 2.0 - 3.0%, and more preferably, it can be limited to the range of 2.2 - 2.8%.

[0041] Mo: 0.4% or less (except 0%)

[0042] Molybdenum (Mo) is an element that delays the transformation of austenite to pearlite while improving the refinement and strength of ferrite. This Mo improves the hardenability of the steel and has the advantage of being able to finely form martensite at grain boundaries to control the yield ratio. However, molybdenum (Mo) is an expensive element, and as the content of molybdenum (Mo) increases, the manufacturing cost increases, and there is a problem of being disadvantageous in terms of cost. Therefore, it is preferable to appropriately control the content of molybdenum (Mo). To obtain the above effects, it is preferably added up to 0.4%. When the content of Mo exceeds 0.4%, there is a sharp increase in alloy cost, so the economy is reduced, and due to excessive grain refinement effect and solid solution strengthening effect, there is instead a problem of reduced ductility of the steel. Therefore, in the present invention, it is preferable to limit the content of Mo to 0.4% or less, and more preferably, it can be limited to 0.2% or less. In addition, in the present invention, considering the amount inevitably added during the manufacturing process, the content of Mo excludes 0%.

[0043] Cr: 1.0% or less (except 0%)

[0044] Chromium (Cr) is an added component to improve the hardenability of the steel and ensure high strength. In addition, chromium (Cr) is an element that plays a very important role in the formation of martensite, minimizing the reduction in elongation compared to the increase in strength, and thus is also beneficial for manufacturing duplex steel with high ductility. In particular, the chromium (Cr) forms Cr during the hot rolling process 23Cr-based carbides such as C6 are partially dissolved and partially remain undissolved during annealing, and can control the amount of dissolved C in martensite below an appropriate level after cooling, thus suppressing the occurrence of yield point elongation. Therefore, the chromium (Cr) is an element beneficial to manufacturing a complex structure steel with a low yield ratio. However, when the content of the chromium (Cr) exceeds 1.0%, not only does its effect saturate, but there is a problem of poor cold rolling property due to excessive increase in hot rolling strength. Since the fraction of Cr-based carbides increases and coarsens, the size of martensite coarsens after annealing, so there is a problem of reduced elongation. Therefore, in the present invention, it is preferable to limit the content of Cr to 1.0% or less, and more preferably, it can be limited to 0.8% or less. In addition, in the present invention, excluding 0% from the content of Cr is considered in view of the amount inevitably added during the manufacturing process.

[0045] P: 0.1% or less (excluding 0%)

[0046] Phosphorus (P) is a substitution element with the largest solution strengthening effect, and is the most favorable element for improving in-plane anisotropy and ensuring strength without significantly impairing formability. However, when adding too much phosphorus (P) above a certain level, the possibility of brittle fracture increases significantly, and there is a problem that phosphorus (P) is an element that hinders the possibility of slab fracture during hot rolling and the plating surface characteristics. In the present invention, it is limited to a maximum of 0.1%, but excluding 0% is considered in view of the inevitably added level.

[0047] S: 0.02% or less (excluding 0%)

[0048] Sulfur (S) is an impurity element inevitably added to steel and is an element that reduces ductility and weldability. Therefore, it is important to control it at a low level as much as possible. In particular, there is a problem of increased possibility of red heat brittleness, so it is preferable to control the content of sulfur (S) at 0.02% or less. However, excluding 0% is considered in view of the inevitably added level during the manufacturing process.

[0049] Acid-soluble aluminum (sol.Al): 0.6% or less (excluding 0%)

[0050] Acid-soluble aluminum is an element added for the purpose of grain refinement and deoxidation of steel. Similar to Si, it is a ferrite-stabilizing element. In addition, acid-soluble aluminum is a component that distributes carbon in ferrite into austenite to improve the hardenability of martensite and effectively form retained austenite. In addition, when maintained in the bainite region during the annealing process, it effectively suppresses the precipitation of carbides in bainite, so it is a useful element that can improve the ductility of the steel plate. However, when the content of acid-soluble aluminum exceeds 0.6%, due to the grain refinement effect, it is beneficial to improve strength, but during the continuous casting operation of steelmaking, excessive inclusions are formed, the possibility of poor surface quality of the coated steel plate increases, and there is also a problem of increased manufacturing cost. Therefore, in the present invention, it is preferably to control the content of acid-soluble aluminum below 0.6%, and more preferably, it can be limited to below 0.4%.

[0051] Ti, Nb: 0.001 - 0.04% each

[0052] Titanium (Ti) and niobium (Nb) are elements that are respectively effective in increasing the strength of the steel plate and grain refinement according to the formation of nano-precipitates. When these elements are added, they combine with carbon to form very fine nano-precipitates, and these nano-precipitates play a role in reducing the hardness difference between phases by strengthening the matrix structure. When the content of the Ti and the Nb is less than 0.001% respectively, it is difficult to ensure the above effects. When the content of the Ti and the Nb exceeds 0.04% respectively, the manufacturing cost increases, and due to excessive precipitates, the ductility may be significantly reduced. Therefore, the content of the Ti and the Nb is respectively limited to 0.001 - 0.04%.

[0053] N: 0.01% or less (except 0%)

[0054] Nitrogen (N) is a component that effectively stabilizes austenite, but when nitrogen (N) exceeds 0.01%, there is a problem that the refining cost of the steel rapidly increases, and due to the formation of AlN, etc., the risk of cracking during continuous casting significantly increases. Therefore, it is preferably to limit the upper limit of nitrogen (N) to 0.01%. However, considering the inevitably added level, 0% is excluded.

[0055] B: 0.01% or less (except 0%)

[0056] Boron (B) is a component that delays the transformation of austenite to pearlite during the annealing and cooling process, and is a hardenability element that suppresses the formation of ferrite and promotes the formation of martensite. However, when the content of boron (B) exceeds 0.01%, excessive B is enriched on the surface, which may lead to poor plating adhesion. Therefore, the content of boron (B) is controlled below 0.01%, and more preferably, 0.005% or less is added. In addition, in the present invention, considering the inevitably added level, 0% is excluded from the content of B.

[0057] Sb: 0.05% or less (except 0%)

[0058] In the present invention, antimony (Sb) may also be selectively included in a range of 0.05% or less. Antimony (Sb) is distributed at grain boundaries, delaying the diffusion of oxidizing elements such as Mn, Si, and Al through grain boundaries, thereby suppressing the surface enrichment of oxides and having a significant effect on suppressing the coarsening of surface enrichment products according to temperature rise and hot rolling process changes. However, when the content of antimony (Sb) exceeds 0.05%, not only does its effect saturate, but also the manufacturing cost and workability are poor. Therefore, it is preferable to limit the content of antimony (Sb) to 0.05% or less, and more preferably, 0.03% or less may be added. Additionally, in the present invention, considering the inevitably added level, 0% is excluded from the content of Sb.

[0059] [Relationship 1]

[0060] C + Si / 30 + Mn / 20 + (Cr + Mo) / 5 ≥ 0.3

[0061] Among them, C, Si, Mn, Cr, and Mo represent the content percentages of steel components in the matrix structure at the t / 4 position of the thickness of the base steel plate.

[0062] In the present invention, it is important to control the contents of C, Si, Al, Cr, and Mo in the alloy components in the matrix structure at the t / 4 position of the thickness of the base steel plate to satisfy the said Relationship 1.

[0063] Si and Al are ferrite stabilizing elements and are elements that promote ferrite phase transformation, promote C enrichment in untransformed austenite, and thus contribute to the formation of retained austenite and martensite. C is also an element that contributes to the formation of martensite and the adjustment of the martensite fraction by promoting C enrichment in untransformed austenite. However, C, Si, and Al have an adverse effect on weldability, thus causing cracks on the surface and inside of the welded part.

[0064] On the other hand, Mn, Cr, Mo, and B are elements that contribute to improving hardenability, but compared with C, Si, and Al, their effect of enriching C in austenite is relatively low. Therefore, it is very important to well adjust the ratio of C, Si, Al and other hardenability elements Mn, Cr, Mo, and B.

[0065] When the value defined by the said Relationship 1 is 0.3 or more, the relationship between the yield strength (YS) and the hole expansion ratio (HER) of HER The 100 / YS can be ensured to be above 5. In addition, while ensuring that the fractions of bainite and tempered martensite are above 70%, the hardness difference between phases is reduced, so that the hole expansion value can be ensured to be above 45%. In addition, the yield strength ratio (YS / TS) can also be ensured to be above 0.8. However, when the value defined by the above relationship 1 is less than 0.3, the above effects may not be expected.

[0066] In the present invention, in addition to the above components, it may be composed of the balance of Fe and other inevitable impurities.

[0067] In addition, the hot-dip galvanized steel sheet of the present invention is a composite structure steel sheet. Compared with the existing DP steel, by increasing the yield strength ratio, workability can be improved while ensuring a certain degree of ductility. For this reason, in addition to the above alloy composition, the following fine microstructure and phase fraction control conditions must also be satisfied.

[0068] Specifically, in terms of area%, the hot-dip galvanized steel sheet of the present invention has a fine microstructure including a total of 70% or more of bainite and tempered martensite, 10% or less of ferrite, and the balance of fresh martensite and retained austenite.

[0069] In order to satisfy the lean composition system and high ductility as the characteristics of DP steel while increasing the yield strength ratio (YS / TS) compared with the existing DP steel, it is necessary to control the structure and composition and perform careful heat treatment. First of all, it is important to introduce a small amount of retained austenite. Retained austenite contributes to improving the ductility of the steel sheet by inducing transformation-induced plasticity. In order to introduce this retained austenite, during rapid cooling, it is cooled to a temperature below Ms to form fresh martensite, and then reheated to a temperature above Ms and subjected to the Quenching&Partitioning (Q&P) process. At this time, a large amount of tempered martensite and bainite are formed, C is stably partitioned, and it helps to form a small amount of retained austenite in the final structure. In order to ensure high hole expansion, it is important to control so that other structures except tempered martensite and bainite are not formed as much as possible. By single-phase region annealing, the formation of ferrite during annealing is prevented, and by adding hardenability elements, the formation of ferrite in the slow cooling range is prevented. In addition, partial bainite is introduced in the slow cooling range to improve ductility. In addition, by the Q&P process of rapid cooling to below the Ms temperature and reheating to above the Ms temperature, the formation of fresh martensite is suppressed as much as possible. In addition, by precipitating fine nano-precipitates in ferrite, the hardness difference between phases is further reduced, so that workability can be improved. Finally, a small amount of fresh martensite is introduced during final cooling, so that the desired strength can be ensured.

[0070] Therefore, in the microstructure of the hot-dip galvanized steel sheet constituting the present invention, the sum of bainite and tempered martensite is controlled to be 70% or more by area. When the sum of bainite and tempered martensite is less than 70% by area, there are problems in that the desired yield ratio and hole expansion property cannot be ensured. More preferably, the sum of bainite and tempered martensite can be limited to 75% or more by area.

[0071] In addition, in the present invention, the ferrite fraction in the microstructure of the hot-dip galvanized steel sheet is preferably limited to 10% or less. When the fraction of the ferrite exceeds 10%, there are problems in that the desired yield ratio and hole expansion property cannot be ensured. More preferably, the fraction of the ferrite structure can be limited to 8% or less by area.

[0072] In addition, the hot-dip galvanized steel sheet of the present invention is composed of the balance of fresh martensite and retained austenite.

[0073] At this time, in the present invention, the area fraction of the retained austenite is preferably limited to less than 5% (excluding 0%). In order to make the fraction of the retained austenite 5% or more, the content of Si must be increased, and at this time, LME problems may occur. More preferably, the fraction of the retained austenite can be limited to the range of 1-4% by area.

[0074] Due to the effect of tissue homogenization, the hot-dip galvanized steel sheet of the present invention having the above-described microstructure has an excellent yield ratio compared with the existing DP steel, and thus can improve workability and formability. Specifically, thereby, a high-strength hot-dip galvanized steel sheet having excellent ductility, hole expansion property, and weldability and a tensile strength of 980 MPa or more can be provided. The hole expansion property (HER) of the hot-dip galvanized steel sheet is 45% or more, and the relationship between the yield strength (YS) and the hole expansion property (HER) is HER 100 / YS is 5 or more, and the relationship between the elongation (EL) and the hole expansion property (HER) is HER EL is 700 or more, the yield ratio (YS / TS) is 0.80 or more, and no LME cracks are generated on the surface and inside of the welded portion.

[0075] Next, the manufacturing method of the hot-dip galvanized steel sheet of the present invention will be described in detail.

[0076] The manufacturing method of the hot-dip galvanized steel sheet of the present invention includes the following processes: preparing a steel billet that meets the above steel composition and relational expression 1, and then reheating the steel billet; performing hot rolling on the reheated steel billet so that the finishing mill exit temperature is Ar3 to Ar3 + 50 °C, then coiling at 400 - 650 °C and cooling to room temperature at an average cooling rate of 0.1 °C / second or less; performing cold rolling on the cooled hot-rolled steel sheet at a reduction rate of 40 - 70% to manufacture a cold-rolled steel sheet; continuously annealing the cold-rolled steel sheet at a temperature of Ac3 - 20 °C to Ac3 + 20 °C; performing primary cooling on the continuously annealed cold-rolled steel sheet at an average cooling rate of 10 °C / second or less, cooling to a temperature range of 500 - 550 °C, and performing secondary cooling at an average cooling rate of 5 °C / second or more, cooling to a temperature of 300 - 340 °C; reheating the cold-rolled steel sheet after the secondary cooling to a temperature above Ms, and then holding for 60 seconds or more; and cooling the held cold-rolled steel sheet to a temperature below 150 °C at an average cooling rate of 5 °C or more.

[0077] First, in the present invention, the steel billet configured as described above is reheated under normal conditions. The slab reheating process is a process of heating the steel billet in order to smoothly perform the subsequent rolling process and fully obtain the desired physical properties of the steel sheet. In the present invention, such reheating conditions are not particularly limited as long as they are normal reheating conditions. As an example, reheating is performed in the temperature range of 1100 - 1300 °C.

[0078] Next, in the present invention, the reheated steel billet is hot rolled so that the finishing mill exit temperature is Ar3 to Ar3 + 50 °C, and then, after coiling at 400 - 650 °C, it is cooled to room temperature at an average cooling rate of 0.1 °C / second or less.

[0079] Under normal conditions, the reheated steel billet is hot finish rolled above the Ar3 transformation point. The present invention does not limit the specific hot rolling conditions at this time, and normal hot rolling temperatures can be used. As an example, hot finish rolling can be performed in the temperature range of 800 - 1000 °C.

[0080] After that, in the present invention, by coiling the hot finish rolled steel sheet in the temperature range of 400 - 650 °C and then cooling to room temperature at an average cooling rate of 0.1 °C / second or less, a hot-rolled steel sheet in which carbides serving as austenite nucleation sites are finely dispersed is manufactured. By uniformly dispersing fine carbides during such hot rolling, carbides are dissolved and austenite is finely dispersed and formed during the subsequent annealing process, and as a result, fine martensite can be uniformly dispersed after annealing.

[0081] In addition, in the present invention, the cooled hot-rolled steel sheet is cold-rolled at a reduction ratio of 40-70% to produce a cold-rolled steel sheet.

[0082] When the cold-rolling reduction ratio is less than 40%, it is difficult to ensure the desired thickness and it is also difficult to correct the shape of the steel sheet. However, when the cold-rolling reduction ratio exceeds 70%, there is a high possibility of cracks occurring in the edge portion of the steel sheet, and there is a problem of bringing about a cold-rolling load. Therefore, in the present invention, it is preferable to limit the cold-rolling reduction ratio to 40-70%.

[0083] Next, in the present invention, the cold-rolled steel sheet is continuously annealed at a temperature of Ac3 - 20°C to Ac3 + 20°C.

[0084] In the present invention, the cold-rolled steel sheet is continuously annealed within a temperature range of Ac3 - 20°C to Ac3 + 20°C. The purpose of the continuous annealing step is to form a single-phase austenite and achieve tissue homogenization. When the continuous annealing temperature is lower than Ac3 - 20°C, it is difficult to ensure a sufficient austenite fraction, and the ferrite fraction increases after annealing, making it impossible to ensure a homogenized fine structure. On the other hand, when the continuous annealing temperature exceeds Ac3 + 20°C, productivity decreases, excessive austenite is formed, and since the grain size becomes too large, it is difficult to ensure the desired strength. In addition, elements such as Si, Mn, and B that deteriorate the surface quality cause increased surface enrichment, and the surface quality may decrease. Considering this, in the present invention, it is preferable to limit the continuous annealing temperature to Ac3 - 20°C to Ac3 + 20°C. More preferably, continuous annealing can also be performed within a temperature range of 810 - 850°C.

[0085] Furthermore, in the present invention, the cold-rolled steel sheet after the continuous annealing is first cooled at an average cooling rate of 10°C / second or less to a temperature range of 500 - 550°C, and then second-cooled at an average cooling rate of 5°C / second or more to a temperature of 300 - 340°C, thereby introducing fresh martensite. At this time, in the present invention, the steel sheet can also be cooled with H2 gas during the second cooling.

[0086] In the present invention, it is very important to control the slow cooling temperature during the first cooling to 500 - 550°C. When the slow cooling temperature exceeds 550°C, bainite will not be introduced in the slow cooling interval, so the desired ductility cannot be ensured. In addition, when the slow cooling temperature is controlled to be lower than 500°C, the fraction of bainite increases excessively, thus making it impossible to ensure the desired strength.

[0087] In addition, it is very important that the rapid cooling temperature during the secondary cooling is controlled at 300 - 340°C. 300 - 340°C is below Ms, which is the martensite formation temperature. When the rapid cooling temperature exceeds 340°C, the fraction of martensite formed initially is very small or it is difficult to form martensite, and it is difficult to form the desired fraction of tempered martensite and bainite during the final cooling, so the desired hole expansion property cannot be obtained. In addition, when the rapid cooling temperature is reduced to below 300°C, the fractions of tempered martensite and bainite are too high, and it is difficult to form the desired fraction of fresh martensite during the final cooling, so the desired strength cannot be obtained.

[0088] Next, in the present invention, the cold-rolled steel sheet after the secondary cooling is reheated to a temperature above Ms and held for 60 seconds or more.

[0089] In the present invention, it is important to form the desired fine microstructure by controlling the slow cooling temperature, rapid cooling temperature, and reheating temperature in the Q&P process of reheating to above Ms after the above-mentioned primary slow cooling process and secondary rapid cooling process of cooling to below Ms. By holding the reheated cold-rolled steel sheet for 60 seconds or more, tempered martensite and bainite are formed, and carbon is enriched in the surrounding untransformed austenite.

[0090] After that, in the present invention, fine fresh martensite is introduced by cooling the held cold-rolled steel sheet to a temperature below 150°C at an average cooling rate of 5°C / second or more.

[0091] In addition, in the present invention, if necessary, the cooled steel sheet can be subjected to hot-dip galvanizing treatment to manufacture a hot-dip galvanized steel sheet. Specifically, the cooled cold-rolled steel sheet is immersed in a zinc pot at 440 - 480°C, and then heated and cooled in a GA furnace at 500 - 540°C to manufacture an alloyed hot-dip galvanized steel sheet, or it can be cooled without alloying treatment after being immersed in the zinc pot, thereby manufacturing a hot-dip galvanized steel sheet.

[0092] In addition, in the present invention, skin pass rolling of less than 1% can also be performed on the cold-rolled steel sheet or the hot-dip galvanized steel sheet. Detailed Description of the Invention

[0093] Hereinafter, the present invention will be described in detail by preferred embodiments. However, it should be noted that the following embodiments are only used to illustrate the present invention for more detailed description and are not used to limit the scope of the rights of the present invention.

[0094] (Example)

[0095] Prepare a billet with the composition shown in Table 1 below, then heat it to a temperature range of 1050 - 1250 °C, and perform hot finish rolling in the temperature range of Ar3 + 50 °C above the Ar3 transformation point temperature. Coil the hot-rolled steel sheet as described above at 400 - 650 °C, and then cool it at a cooling rate of 0.1 °C per second or less to produce a hot-rolled steel sheet. After pickling the hot-rolled steel sheet, perform cold rolling with a reduction rate of 40 - 70%, then perform continuous annealing under the same conditions as Table 2 below, and perform Q&P heat treatment. Next, cool it to a temperature below 150 °C at an average cooling rate of 5 °C or more, then immerse it in a molten zinc pot at 460 °C for 5 seconds, and subsequently heat it in a GA furnace at 520 °C for about 2 seconds and then cool it to produce an alloyed hot-dip galvanized steel sheet. In addition, perform skin pass rolling of less than 1% on this alloyed hot-dip galvanized steel sheet. In addition, when performing Q&P heat treatment in this experiment, the cooling rate during the first cooling is uniformly applied at 6 °C / second, the cooling rate during the second cooling is uniformly applied at 15 °C / second, and the holding time after reheating is uniformly applied at 200 seconds.

[0096] For each steel sheet manufactured as described above, evaluate the mechanical properties and microstructure properties, and show the results in Table 3 below. At this time, for each test piece, perform a tensile test in the C direction using JIS standards to evaluate the tensile physical properties. For the fraction of the microstructure, analyze the matrix structure at the t / 4 position of the plate thickness of the annealed steel sheet and use the results. Specifically, after etching with a nitric acid alcohol solution (Nital), measure the fractions of ferrite, bainite, martensite, and austenite using FE-SEM and an image analyzer. In addition, measure the hole expansion property using a hole expansion tester. In addition, evaluate whether spot welding LME cracks occur.

[0097] [Table 1]

[0098]

[0099] The balance components in Table 1 are Fe and unavoidable impurities.

[0100] Relationship 1 is C + Si / 30 + Mn / 20 + (Cr + Mo) / 5.

[0101] [Table 2]

[0102]

[0103] [Table 3]

[0104]

[0105] In Table 3, F refers to ferrite, B+TM refers to bainite + tempered martensite, FM refers to fresh martensite, and RA refers to retained austenite.

[0106] As shown in Tables 1 to 3 above, in the case of Invention Examples 1 to 4 where the steel composition range, manufacturing conditions, and fine microstructure of the steel satisfy the scope of the present invention, the hole expansion ratio (HER) is 45% or more, and the relationship between the yield strength (YS) and the hole expansion ratio (HER) is HER 100 / YS is 5 or more, and the relationship between the elongation (EL) and the hole expansion ratio (HER) is HER EL is 700 or more. In addition, no LME cracks are generated on the surface and inside of the welded part, and the yield ratio (YS / TS) is 0.80 or more, indicating that the material properties, hole expansion ratio, and weldability of the steel plate desired by the present invention can be ensured.

[0107] In contrast, in the case of Comparative Examples 1 to 10 where the steel composition (Relationship 1) and / or the steel manufacturing process deviate from the scope of the present invention or the fine microstructure fraction of the steel deviates from the scope of the present invention, it can be seen that the relationship between the yield strength (YS) and the hole expansion ratio (HER) is HER 100 / YS is less than 5, or the relationship between the elongation (EL) and the hole expansion ratio (HER) is HER EL is less than 700, the yield ratio (YS / TS) is less than 0.80, and the hole expansion ratio is less than 45%. Therefore, it can be seen that the strength, ductility, hole expansion ratio, and weldability of the steel plate desired by the present invention cannot be ensured simultaneously.

[0108] Specifically, in Comparative Example 1, the annealing temperature is too high, so the elongation is low and the hole expansion ratio cannot be ensured. In Comparative Example 2, the annealing temperature is too low, forming too much two-phase region ferrite, so the target strength cannot be obtained.

[0109] In Comparative Example 3, since the primary cooling termination temperature is too high, bainite cannot be ensured, so YS is too high and the elongation is not excellent.

[0110] In Comparative Example 4, since the primary cooling termination temperature is too low, the target strength cannot be ensured. In Comparative Example 5, since the secondary cooling termination temperature is too high, tempered martensite cannot be sufficiently ensured, so the hole expansion ratio cannot be ensured.

[0111] In addition, in Comparative Example 6, since the secondary cooling termination temperature is too low, too much tempered martensite and bainite are formed, so the strength deteriorates.

[0112] In addition, Figure 1 shows the relationship between the yield strength (YS) and the hole expansion ratio (HER) of the inventive steel and the comparative steel in the examples of the present invention according to Relationship 1, HER Graph of the change in 100 / YS Figure 2 It is a graph showing the change in the hole expansion ratio (HER) according to Relationship 1 of the inventive steel and the comparative steel in the embodiments of the present invention. Figure 3 It is a graph showing the change in the yield ratio according to Relationship 1 of the inventive steel and the comparative steel in the embodiments of the present invention. Figure 4 It is a graph showing the relationship HER between the elongation (EL) and the hole expansion ratio (HER). EL. Additionally, in the Figures 1 to 4 the inventive steel refers to the steel corresponding to Invention Examples 1 to 4, and the comparative steel refers to the steel corresponding to Comparative Examples 7 to 10.

[0113] As described above, the preferred embodiments of the present invention have been described in the detailed description of the present invention, but those skilled in the art can make various modifications without departing from the scope of the present invention. Therefore, the scope of the rights of the present invention is not limited to the described embodiments and should be determined by the claims and their equivalents.

Claims

1. A hot-dip galvanized steel sheet, by weight%, the hot-dip galvanized steel sheet contains: carbon (C): 0.08 - 0.16%, silicon (Si): less than 0.8% and except 0%, manganese (Mn): 2.0 - 3.0%, molybdenum (Mo): less than 0.4% and except 0%, chromium (Cr): less than 1.0% and except 0%, phosphorus (P): less than 0.1% and except 0%, sulfur (S): less than 0.02% and except 0%, aluminum (sol.Al): less than 0.6% and except 0%, titanium (Ti): 0.001 - 0.04%, niobium (Nb): 0.001 - 0.04%, nitrogen (N): less than 0.01% and except 0%, boron (B): less than 0.01% and except 0%, the balance of Fe and other inevitable impurities, and the C, Si, Mn, Cr and Mo satisfy the following relational expression 1, By area%, the hot-dip galvanized steel sheet has a microstructure comprising a total of 70% or more of bainite and tempered martensite, 10% or less of ferrite, and the balance of fresh martensite and retained austenite, The hole expansion rate (HER) of the hot-dip galvanized steel sheet is more than 45%, and the relationship between the yield strength (YS) and the hole expansion rate (HER) is HER 100 / YS is 5 or more, and the relationship between the elongation (EL) and the hole expansion rate (HER) is HER EL is 700 or more, and the yield ratio (YS / TS) is 0.80 or more. [Relational expression 1] C + Si / 30 + Mn / 20 + (Cr + Mo) / 5 ≥ 0.3 Among them, C, Si, Mn, Cr and Mo represent the content percentages of the steel components in the matrix structure at the t / 4 position of the thickness of the base steel sheet.

2. The hot-dip galvanized steel sheet according to claim 1, wherein, By area%, the fraction of the retained austenite is less than 5% and except 0%.

3. The hot-dip galvanized steel sheet according to claim 2, wherein, By area%, the fraction of the retained austenite satisfies 1 - 4%.

4. The hot-dip galvanized steel sheet according to claim 1, wherein, The hot-dip galvanized steel sheet further contains antimony (Sb) of 0.05% or less and except 0%.

5. A method for manufacturing a hot-dip galvanized steel sheet, which includes the following processes: Prepare a steel billet, and then reheat the steel billet. By weight%, the steel billet contains: carbon (C): 0.08 - 0.16%, silicon (Si): less than 0.8% and except 0%, manganese (Mn): 2.0 - 3.0%, molybdenum (Mo): less than 0.4% and except 0%, chromium (Cr): less than 1.0% and except 0%, phosphorus (P): less than 0.1% and except 0%, sulfur (S): less than 0.02% and except 0%, aluminum (sol.Al): less than 0.6% and except 0%, titanium (Ti): 0.001 - 0.04%, niobium (Nb): 0.001 - 0.04%, nitrogen (N): less than 0.01% and except 0%, boron (B): less than 0.01% and except 0%, the balance of Fe and other inevitable impurities, and the C, Si, Mn, Cr and Mo satisfy the following relational expression 1; Hot-roll the reheated steel billet so that the finish rolling exit temperature is Ar3 to Ar3 + 50 °C, then coil at 400 - 650 °C and cool to room temperature at an average cooling rate of 0.1 °C / second or less; Cold-roll the cooled hot-rolled steel sheet at a reduction rate of 40 - 70% to manufacture a cold-rolled steel sheet; Continuously anneal the cold-rolled steel sheet at a temperature of Ac3 - 20 °C to Ac3 + 20 °C; The cold-rolled steel sheet after the continuous annealing is subjected to primary cooling at an average cooling rate of 10°C / second or less, cooled to a temperature range of 500 - 550°C, and then subjected to secondary cooling at an average cooling rate of 5°C / second or more, cooled to a temperature of 300 - 340°C; The cold-rolled steel sheet after the secondary cooling is reheated to a temperature above Ms and then held for 60 seconds or more; and The cold-rolled steel sheet after the holding is cooled to a temperature below 150°C at an average cooling rate of 5°C / second or more, [Relationship 1] C + Si / 30 + Mn / 20 + (Cr + Mo) / 5 ≥ 0.3 Among them, C, Si, Mn, Cr, and Mo represent the content percentages of the steel components in the matrix structure at the t / 4 position of the thickness of the base steel sheet.

6. The manufacturing method of the hot-dip galvanized steel sheet according to claim 5, wherein, Continuous annealing is carried out in the temperature range of 810 - 850°C.

7. The manufacturing method of the hot-dip galvanized steel sheet according to claim 5, wherein, The steel sheet cooled to a temperature below 150°C at an average cooling rate of 5°C / second or more is further subjected to skin pass rolling of less than 1%.

8. The manufacturing method of the hot-dip galvanized steel sheet according to claim 5, wherein, Further includes antimony (Sb): 0.05% or less and excluding 0%.

9. The manufacturing method of the hot-dip galvanized steel sheet according to claim 5, wherein, The manufacturing method further includes a process of immersing the steel sheet cooled to a temperature below 150°C at an average cooling rate of 5°C / second or more in a zinc pot at 440 - 480°C and then cooling.

10. The manufacturing method of the hot-dip galvanized steel sheet according to claim 5, wherein, The manufacturing method further includes a process of immersing the steel sheet cooled to a temperature below 150°C at an average cooling rate of 5°C / second or more in a zinc pot at 440 - 480°C and then heating in a GA furnace at 500 - 540°C and then cooling.

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