JP7835518B2 - Ultra-high-strength steel sheet with excellent bending properties and method for manufacturing the same - Google Patents

Description

This invention relates to steel sheets suitable for use in automotive structural components and the like, and more specifically, to ultra-high-strength steel sheets with excellent bending properties and a method for manufacturing the same.

Recently, in the automotive sector, particularly in developed countries such as those in Europe, research into reducing vehicle weight is actively underway due to fuel efficiency regulations and performance improvements. In the steel sector, in particular, efforts are being made to meet these weight reduction demands from automakers by increasing strength and further reducing steel plate thickness at the same grade compared to competing materials (such as Mg, Al, and CFRP (carbon fiber reinforced plastic)). In other words, along with weight reduction, there is a growing demand for greater stability and strength in vehicle body materials due to stricter safety regulations for passengers and pedestrians.

On the other hand, to improve the stability and impact characteristics of the vehicle body, the use of high-strength steel with superior yield strength is increasing in BIW (Body-In-White) structural members. Such structural members have the characteristic that a higher yield strength relative to tensile strength, i.e., a higher yield ratio (yield strength/tensile strength, YR), is advantageous for absorbing impact energy.

Therefore, a common method for increasing the yield strength of steel is to utilize water cooling during continuous annealing. Specifically, this method involves annealing cold-rolled steel sheets in either the two-phase or single-phase region, rapidly cooling them to approximately room temperature, and then performing processes such as tempering to produce ultra-high-strength steel.

However, while the ultra-high-strength steel produced by this method has an extremely high yield ratio, it suffers from a problem where the shape quality of the coil deteriorates due to temperature variations in the width and length directions. This can lead to problems such as material defects and reduced workability depending on the part during processing by methods such as roll forming. Furthermore, since the elongation rate generally decreases as the strength of the steel increases, there is a problem of reduced formability.

To overcome this, a hot press forming (HPF) method has been developed and applied, in which the material is formed at a relatively high temperature where molding is easier, and then strength is ensured through water cooling between the die and the material (see Patent Document 1).

When applying the HPF (High-Performance Fiber) construction method, higher strength can be ensured compared to materials of the same thickness. Therefore, component development using the HPF method is being carried out, primarily in Europe.

However, the HPF method requires excessive capital investment and has led to problems such as increased process costs. Therefore, there is a real need for the development of materials for cold stamping.

In other words, there is a need to develop steel sheets that are suitable for use as cold stamping materials, possess high strength and a high yield ratio to ensure collision performance characteristics, and have excellent formability and other properties.

International Publication No. 2021/084303

One aspect of the present invention provides a steel sheet suitable for cold stamping, particularly an ultra-high-strength steel sheet with excellent bending properties, and a method for manufacturing the same, while being suitable for automotive structural components.

The problems addressed by this invention are not limited to those described above. The problems addressed by this invention can be understood from the entirety of this specification, and any person with ordinary skill in the art to which this invention pertains will have no difficulty understanding the additional problems addressed by this invention.

One aspect of the present invention provides an ultra-high-strength steel sheet with excellent bending properties, comprising, by weight percent, carbon (C): 0.1-0.3%, manganese (Mn): 1.0-2.3%, silicon (Si): 0.05-1.0%, phosphorus (P): 0.1% or less (excluding 0%), sulfur (S): 0.03% or less (excluding 0%), aluminum (Al): 0.01-0.5%, and two or more of chromium (Cr): 0.01-0.2%, molybdenum (Mo): 0.01-0.2%, and boron (B): 0.005% or less, one or more of titanium (Ti): 0.1% or less, and niobium (Nb): 0.1% or less, with the remainder being Fe and unavoidable impurities, satisfying the following relational formula 1, and containing martensite and/or tempered martensite phase in an area fraction of 99% or more as a microstructure.

[Relationship 1]

(Here, Ceq1=C+(Mn/20)+(Si/30)+(2P)+(4S), Ceq2=C+(Mn/6)+(Si/30)+(Cr+Mo+V+Nb)/5+(Cu+Ni)/15.)

Another aspect of the present invention includes the steps of: heating a steel slab satisfying the above-described alloy composition and relational formula 1 to a temperature range of 1100 to 1300°C; finishing hot-rolling the reheated steel slab at Ar3 or higher to produce a hot-rolled steel sheet; winding the hot-rolled steel sheet at a temperature of 700°C or lower; cold-rolling the wound hot-rolled steel sheet with a total reduction ratio of 30 to 80% to produce a cold-rolled steel sheet; continuously annealing the cold-rolled steel sheet at Ac3 or higher for 30 seconds or more; primary cooling after the continuous annealing to a temperature range of 550 to 750°C at an average cooling rate of 1 to 10°C/s; secondary cooling after the primary cooling to a temperature of Ms-190°C or lower at an average cooling rate of 20 to 80°C/s; and reheating and then overaging after the secondary cooling.
The above reheating and overaging steps are characterized by heating to a temperature range that satisfies the following relational equation 3, providing a method for manufacturing ultra-high-strength steel sheets with excellent bending properties.

[Relationship Equation 3]
CT2+30℃≦A≦270℃

(Here, CT2 represents the secondary cooling completion temperature (°C), and A represents the reheating and overaging temperature (°C).)

According to the present invention, by achieving both ultra-high strength and a high yield ratio, it is possible to provide a steel sheet with improved workability. In particular, the steel sheet of the present invention is not only suitably applicable to automotive structural components, but also has the advantage of being advantageously applicable to processing such as cold stamping.

In one embodiment of the present invention, photographs of the surface microstructure of the inventive example and the comparative example are shown. In one embodiment of the present invention, photographs of the microstructure in the 1/4t region of the inventive example and comparative example are shown.

The inventors of this invention conducted extensive research to provide a steel sheet that is suitable for automotive structural components while also being advantageous for processing such as cold stamping. As a result, they confirmed that by optimizing the alloy composition and manufacturing conditions, they could provide a steel sheet with the desired microstructure and physical properties, thus completing the present invention.

The present invention will be described in detail below.

An ultra-high-strength steel sheet according to one embodiment of the present invention may contain, by weight percent, carbon (C): 0.1-0.3%, manganese (Mn): 1.0-2.3%, silicon (Si): 0.05-1.0%, phosphorus (P): 0.1% or less (excluding 0%), sulfur (S): 0.03% or less (excluding 0%), and aluminum (Al): 0.01-0.5%.

The following section will explain in detail the reasons for limiting the alloy composition of the ultra-high-strength steel sheet provided in this invention as described above.

On the other hand, unless otherwise specified in this invention, the content of each element is based on weight, and the proportion of the structure is based on area.

Carbon (C): 0.1-0.3%
Carbon (C) is an interstitial solid solution element and is the most effective and important element for improving the strength of steel. In particular, it is an essential element that must be added to martensitic steel to ensure its strength.

To obtain a steel sheet having the target strength, yield ratio, etc., as described in this invention, it is preferable to add 0.1% or more of the above-mentioned carbon (C). However, if the content exceeds 0.3%, while the martensite strength increases, the formation of carbides becomes easier during the continuous annealing process, leading to roughening and a decrease in ductility, as well as a deterioration in bending properties. Furthermore, an excessive increase in carbon content impairs weldability.

Therefore, in the present invention, the above C can be contained in an amount of 0.1 to 0.3%, and more preferably in an amount of 0.12% or more and 0.28% or less.

Manganese (Mn): 1.0–2.3%
Manganese (Mn) is an element that readily ensures the final martensite phase in composite steels by suppressing ferrite formation and promoting austenite formation.

When the Mn content exceeds 2.3%, Mn segregates in the thickness direction of the steel, easily forming manganese bands within the slab. This increases the occurrence of defects during rolling, along with continuous casting cracks. On the other hand, if the content is less than 1.0%, it becomes impossible to achieve the target level of strength.

Therefore, in the present invention, the above-mentioned Mn can be contained in an amount of 1.0 to 2.3%, more preferably in an amount of 1.2% or more and 2.1% or less. Even more preferably, it can be contained in an amount of 1.4% or more.

Silicon (Si): 0.05–1.0%
In the process of manufacturing the steel sheet to be obtained in this invention, silicon (Si) plays a role in suppressing the formation of carbides and controlling the size of the carbides during the reheating and overaging treatment stages that are performed after continuous annealing and cooling.

To fully obtain the effects described above, it is preferable to include 0.05% or more of Si. However, if the content exceeds 1.0%, ferrite may form during cooling in a continuous annealing furnace, potentially weakening the steel's strength. Furthermore, Si-based oxides may form during reheating and overaging after cooling, potentially leading to surface oxidation problems in the steel.

Therefore, in the present invention, the above-mentioned Si can be contained in an amount of 0.05 to 1.0%, more preferably in an amount of 0.09% or more and 0.8% or less. Even more preferably, it can be contained in an amount of 0.6% or less.

Phosphorus (P): 0.1% or less (excluding 0%)
Phosphorus (P) is an impurity element contained in steel, and if its content exceeds 0.1%, the weldability of the steel deteriorates and brittleness may occur. Therefore, the amount of P is limited to 0.1% or less, and 0% can be excluded considering the level that is inevitably added during the steel manufacturing process. More advantageously, the amount of P can be 0.05% or less, and even more advantageously, 0.03% or less.

Sulfur (S): 0.03% or less (excluding 0%)
Sulfur (S), like P mentioned above, is an impurity that is inevitably contained in steel and is an element that inhibits the ductility and weldability of steel, so it is advantageous to control its content to be as low as possible. In this invention, even if the above S is contained at a maximum of 0.03%, it is not difficult to ensure the target physical properties, and the upper limit can be limited to 0.03%, and 0% can be excluded considering the level that is inevitably added during the steel manufacturing process.

On the other hand, in order to more advantageously secure the bending properties targeted by this invention, the content of S can be limited to 0.01% or less, and more advantageously, to 0.005% or less.

Aluminum (Al): 0.01–0.5%
Aluminum (Al) can be added to remove oxygen from molten steel and, like Si mentioned above, is an element that stabilizes ferrite. Furthermore, Al increases the carbon content in austenite, improving the hardening ability of the final martensitic steel.

To fully obtain the effects described above, the Al content can be 0.01% or more. However, if the content exceeds 0.5%, ferrite may form during cooling in a continuous annealing furnace, potentially weakening the strength. Furthermore, it may combine with N, which is inevitably present as an impurity in the steel, to form AlN, potentially inducing slab cracks and hindering hot rolling.

Therefore, in the present invention, the above-mentioned Al can be present in an amount of 0.01 to 0.5%.

On the other hand, the steel sheet of the present invention may further contain elements advantageous for ensuring the physical properties of steel, in addition to the alloy composition described above. Specifically, it is preferable that the steel sheet of the present invention further contains two or more elements selected from chromium (Cr), molybdenum (Mo), and boron (B), and one or more elements from titanium (Ti) and niobium (Nb).

Chromium (Cr): 0.01–0.2%
Chromium (Cr) can be added to improve the hardening ability of steel and ensure high strength. In particular, it is useful in suppressing bainite formation during cooling in a continuous annealing furnace and in producing ultra-high-strength steel sheets composed of a pure martensite phase.

To fully obtain the effects described above, the above-mentioned Cr can be added at a concentration of 0.01% or more. However, if the content exceeds 0.2%, the cost of the ferroalloy increases, leading to economic disadvantages.

Therefore, when adding Cr as described above, it can be added at a concentration of 0.01 to 0.2%.

Molybdenum (Mo): 0.01–0.2%
Molybdenum (Mo), like Cr mentioned above, is an element that improves the hardening ability of steel.

To fully obtain the hardening effect, Mo can be added at a concentration of 0.01% or more. However, if the content exceeds 0.2%, the amount of alloy added becomes excessive, leading to a problem of increased cost for the ferroalloy.

Therefore, when adding Mo as described above, it can be added at a concentration of 0.01 to 0.2%.

Boron (B): 0.005% or less. Boron (B) is an element that suppresses the transformation of austenite to ferrite during the continuous annealing process, and even in very small amounts, it is an effective element for improving hardening ability, similar to Cr and Mo. However, if its content exceeds 0.005%, there is a risk that the Fe ₂₃ (B, C) ₆ precipitate phase will precipitate at the austenite grain boundaries, thereby promoting the formation of ferrite.

Therefore, when adding B as described above, it can be added at a concentration of 0.005% or less.

Titanium (Ti): 0.1% or less. Titanium (Ti) is an element that forms fine carbides and contributes to ensuring yield strength and tensile strength. Furthermore, since Ti is an element that causes scavenging by precipitating N, which is inevitably present in steel at impurity levels, into TiN, it can be added in a content of 48/(14 × N) or more based on the chemical equivalent standard.

If the Ti content exceeds 0.1%, coarse carbides will precipitate, leading to a decrease in strength and elongation as the carbon content in the steel decreases. Furthermore, it may cause nozzle clogging during continuous casting; therefore, the Ti should be added at a concentration of 0.1% or less.

On the other hand, to maximize the effect of adding B as described above, it is advantageous to add Ti together with it.

Niobium (Nb): 0.1% or less. Niobium (Nb) is an element that segregates at austenite grain boundaries, suppressing the coarsening of austenite grains during the continuous annealing process, and contributing to improved strength by forming fine carbides.

If the Nb content exceeds 0.1%, the precipitation of coarse carbonitrides increases, potentially reducing the carbon content in the steel and lowering its strength and elongation. Furthermore, the workability of the base material decreases, leading to increased manufacturing costs.

Therefore, when adding Nb as described above, it can be added at a concentration of 0.1% or less.

The remaining component of this invention is iron (Fe). However, in the normal manufacturing process, unintended impurities may inevitably be introduced from the raw materials or the surrounding environment, and therefore cannot be eliminated. Since these impurities are easily recognizable to any technician in a normal manufacturing process, their details are not specifically mentioned in this specification.

The steel sheet of the present invention that satisfies the alloy composition described above preferably satisfies the relationship between the content of specific elements as shown in the following relational expression 1.

[Relationship 1]

(Here, Ceq1=C+(Mn/20)+(Si/30)+(2P)+(4S), Ceq2=C+(Mn/6)+(Si/30)+(Cr+Mo+V+Nb)/5+(Cu+Ni)/15.)

Relational equation 1 expresses the influence of the alloying element content added to the steel on the welding properties as a combined relation between Ceq1 and Ceq2. When the range satisfies 0.12 to 0.28, the basic welding properties can be satisfied while advantageously securing the physical properties targeted by this invention.

Specifically, if the value of the above relational expression 1 is less than 0.12, the strength level targeted by this invention cannot be secured. On the other hand, if the value exceeds 0.28, there is a possibility that the physical properties, particularly the welding properties, will deteriorate significantly.

The value of the above relational expression 1 can more preferably be 0.15 or more and 0.27 or less, and even more preferably 0.17 or more.

The steel sheet of the present invention, satisfying the alloy composition and relational formula 1 described above, preferably contains a martensite phase as the main phase in its microstructure.

Specifically, the above-mentioned steel sheet may contain martensite and/or tempered martensite phases in an area fraction of 99% or more. In this case, the above fraction may be 100%.

If the fraction of the above-mentioned martensite and/or tempered martensite phase is 99%, the remaining 1% may be ferrite and/or bainite phase.

The steel sheet of the present invention preferably has a surface layer in a specific region, as described later, and in the remaining region excluding the surface layer (for example, the central region), the main structure is preferably martensite and/or tempered martensite phase.

On the other hand, the steel sheet of the present invention has the characteristic that the surface layer can be defined as a region ranging from a minimum of 50 μm to a maximum of 70 μm in the thickness direction from the surface, and that this surface layer contains a soft phase.

Preferably, the surface layer contains tempered martensite phase at an area fraction of 70% or less, and the remaining structure may include one or more of ferrite and bainite, which have softer properties than the tempered martensite. By softening the surface layer of the steel sheet in this way, the bending properties can be further improved.

Furthermore, the aforementioned surface layer, which contains a certain soft phase, is characterized by containing a decarburized layer with a lower carbon content than that found in the steel sheet.

Specifically, it is preferable that the carbon content ratio of the steel sheet of the present invention within the region (A) of 1 to 3 μm in the thickness direction based on the surface is 0.6 or less. Here, the carbon content ratio in region (A) means [average carbon content in region (A) / carbon content of the steel sheet].

Furthermore, it is preferable that the carbon content ratio of the region (B) in the thickness direction of 0.2 to 30 μm relative to the carbon content of the steel sheet is 0.9 or less. Here, the carbon content ratio of region (B) means [average carbon content of region (B) / carbon content of the steel sheet].

While the decarburized layer within the surface is advantageous for improving the bending properties of the steel sheet, if the carbon (C) content ratio in specific regions (A and B) within the surface exceeds 0.6 and 0.9, respectively, the target bending properties cannot be achieved.

Here, the decarburized layer may be formed to a thickness corresponding to the surface layer, or it may be formed to a thickness less than the surface layer.

In this invention, the decarburized layer can be formed by controlling the continuous annealing process during the steel sheet manufacturing process, which will be explained in detail later.

As described above, the steel sheet of the present invention, in which the microstructure is composed of a hard phase while the surface layer contains a decarburized layer, possesses ultra-high strength with a tensile strength of 1300 MPa or more, a high yield ratio of 0.72 or more, and a bending characteristic (R/t) of 3 or less.

Furthermore, the steel sheet of the present invention can satisfy the following relational equation 2 regarding the relationship between tensile strength and bending characteristics, specifically, the relationship between tensile strength (TS), a fundamental tensile property, and the maximum bending angle after a three-point bending test according to the VDA 238-100 standard.

[Relationship Equation 2]
(Tensile strength (TS) / Maximum bending angle) ≤ 25

When the value of equation 2 above is 25 or less, excellent bending properties can be ensured in ultra-high-strength steel with a tensile strength of 1300 MPa or more. However, when the value exceeds 25, although the strength remains high, the bending properties deteriorate.

The following describes in detail a method for manufacturing an ultra-high-strength steel sheet with excellent bending properties according to another aspect of the present invention.

In short, this invention allows for the production of the target steel sheet through the following steps: [steel slab heating - hot rolling - coiling - cold rolling - continuous annealing]. Each step will be described in detail below. The continuous annealing step includes a cooling step, as well as reheating and overaging steps, meaning that these steps are performed simultaneously on a continuous annealing line.

[Heating steel slabs]
First, a steel slab satisfying the aforementioned alloy composition can be prepared and then heated.

This process is performed to ensure a smooth subsequent hot rolling process and to obtain the desired physical properties of the steel sheet. This invention does not particularly limit the conditions of this heating process; ordinary conditions are acceptable. For example, the heating process can be performed in a temperature range of 1100 to 1300°C. If the heating temperature is below 1100°C, there is a problem of a rapid increase in load during the subsequent hot rolling process. On the other hand, if the temperature exceeds 1300°C, there is a problem of increased surface scale and a decrease in material yield.

[Hot rolling]
The heated steel slab can be hot-rolled to produce a hot-rolled steel sheet, and in this process, finish hot rolling can be performed in a temperature range of Ar3 or higher.

If the temperature during the finish hot rolling process is below Ar3, rolling will occur in the ferrite-austenite two-phase region or the ferrite region, resulting in the formation of a mixed-grain structure. Furthermore, fluctuations in the hot rolling load may lead to manufacturing errors.

More specifically, the above-mentioned finish hot rolling can be performed in a temperature range of 800 to 1000°C.

[Rewind]
The hot-rolled steel sheet produced by the above method can be wound into a coil shape.

The above winding process can be performed in a temperature range of 700°C or lower. If the winding temperature exceeds 700°C, excessive oxide film formation may occur on the steel sheet surface, potentially causing defects.

On the other hand, a lower winding temperature results in higher strength for the hot-rolled steel sheet, which has the disadvantage of increasing the rolling load in the subsequent cold-rolling process. Therefore, the lower limit of the winding temperature can be restricted to 100°C.

[Cold rolling]
The hot-rolled steel sheet wound up as described above can be cold-rolled to produce a cold-rolled steel sheet, and in the present invention, the cold rolling can be performed at a cold reduction ratio of 30 to 80%.

If the cold reduction ratio during cold rolling is less than 30%, it may not only fail to achieve the target thickness, but hot-rolled grains may remain, potentially affecting austenite formation and the final physical properties during the subsequent continuous annealing process. On the other hand, if the cold reduction ratio exceeds 80%, the reduction amount in the length and width directions will become uneven due to work hardening during cold rolling, potentially leading to variations in the material properties of the final steel sheet. Furthermore, achieving the target thickness may be difficult depending on the rolling load.

On the other hand, prior to the cold rolling process, a pickling step can be further performed to remove the oxide layer formed on the surface of the hot-rolled steel sheet obtained by hot rolling. The conditions for this pickling step are not particularly limited and can be carried out under commonly used conditions.

[Continuous annealing]
It is preferable to subject the cold-rolled steel sheet produced as described above to continuous annealing. This continuous annealing can be carried out, for example, in a continuous annealing furnace (CAL).

The above continuous annealing process can be carried out in a step of heat treatment at a temperature of Ac3 or higher for 30 seconds or more. This is to ensure 100% austenite fraction through annealing in the austenite single-phase region.

Here, Ac3 can be calculated using the following formula.

[Formula] Ac3=910-203√C-15.2Ni+44.7Si+104V+31.5Mo+13.1W

(In the formula, each element is represented by its weight content.)

In this invention, it is preferable to control the dew point temperature inside the annealing furnace to 0 to 20°C when performing continuous annealing under the conditions described above. By controlling the dew point temperature in this way, a decarburized layer can be formed on the steel surface during the continuous annealing process.

Normally, the dew point in a continuous annealing furnace is around -50°C. However, when humid nitrogen ( _N₂ + _H₂O ) is introduced to raise the dew point temperature above 0°C, the partial oxygen pressure increases. This causes the carbon (C) in the steel to react with the oxygen (O) in the annealing furnace, releasing CO gas, which then causes decarburization to occur in the surface layer.

If the dew point temperature inside the annealing furnace is below 0°C, a sufficient decarburized layer will not form on the steel surface. Conversely, if it exceeds 20°C, problems arise regarding equipment lifespan and productivity.

Thus, by forming a decarburized layer on the steel surface during the continuous annealing process, and softening only the surface layer, it is possible to further improve the bending properties of ultra-high-strength steel.

[Gradual cooling]
As mentioned earlier, the target microstructure can be formed by cooling the cold-rolled steel sheet that has undergone continuous annealing treatment as described above, and it is preferable to perform this cooling stepwise.

In this invention, the above-mentioned stepwise cooling can be performed in a primary-secondary cooling sequence. Specifically, after the continuous annealing described above, primary cooling can be performed to a temperature range of 550-750°C at an average cooling rate of 1-10°C/s, followed by secondary cooling to a temperature range of Ms-190°C or lower at an average cooling rate of 20-80°C/s.

If the end temperature during the primary cooling process is below 550°C, phases such as ferrite and bainite may form, potentially reducing the material's strength. Conversely, if the temperature exceeds 750°C, not only will the annealing furnace's lifespan be shortened, but excessive cooling will be required during the subsequent secondary cooling process. This could lead to problems in the actual production line, such as defects in the sheet metal's shape and difficulties in controlling its meandering.

Furthermore, if the average cooling rate during the primary cooling is less than 1°C/s, a ferrite phase will form during cooling, making it impossible to achieve the target level of strength. On the other hand, if it exceeds 10°C/s, the average cooling rate during the subsequent secondary cooling will decrease, increasing the proportion of other low-temperature transformation phases besides martensite, ultimately making it impossible to achieve the target level of strength.

As mentioned above, after the primary cooling is complete, rapid cooling (secondary cooling) can be performed at an average cooling rate above a certain level.

In particular, in this invention, in order to secure the martensite and/or tempered martensite phase as the main structure, it is advantageous to rapidly cool the material to a temperature below Mf (martensitic transformation completion temperature) during secondary cooling.

Specifically, by cooling to a temperature of Ms-190°C or lower, a sufficiently hard martensite structure can be formed, and an increase in yield strength due to carbide precipitation can be obtained during the subsequent reheating (tempering) process. If the temperature at which the above cooling ends exceeds Ms-190°C, it becomes difficult to secure the desired level of strength in this invention, and the subsequent reheating temperature may become excessively high, which may lead to a deterioration in the bendability of the steel. Furthermore, the intended fraction of the structure (martensite and/or tempered martensite) may not be sufficiently secured.

Therefore, in this invention, by limiting the termination temperature during secondary cooling, the tempering effect can be sufficiently induced without excessively increasing the subsequent reheating temperature, thereby ensuring the bending characteristics.

While there are no specific limitations on the lower limit of the completion temperature for the secondary cooling process described above, it can be restricted to approximately 50°C, taking into consideration the characteristics of the equipment.

Here, Ms (the transformation initiation temperature of martensite) can be calculated using the following formula.

[Formula] Ms=539-423C-30.4Mn-7.5Si+30Al-17.7Ni-12.1Cr-7.5Mo

(In the formula, each element is represented by its weight content.)

If the average cooling rate during the secondary cooling process is less than 20°C/s, there is a risk of partial formation of bainite structure during the secondary cooling process. On the other hand, if it exceeds 80°C/s, the rapid transformation rate of martensite at the secondary cooling stage can degrade the surface shape of the steel sheet, leading to material variation in the width direction.

[Reheating and over-aging]
In this invention, the toughness of steel can be improved by transforming the martensite phase, which has a high potential density and is hard, formed during secondary cooling into tempered martensite through reheating and overaging treatment.

Specifically, the above reheating and overaging treatment preferably involves heating the cold-rolled steel sheet, which has been cooled in stages as described above, to a temperature range that satisfies the following relational equation 3, and then maintaining that temperature for 1 to 20 minutes.

[Relationship Equation 3]
CT2+30℃≦A≦270℃

(Here, CT2 represents the secondary cooling completion temperature (°C), and A represents the reheating and overaging temperature (°C).)

In other words, to ensure a sufficient tempering effect, the lower limit of the reheating temperature is restricted to a temperature of 30°C or higher relative to the secondary cooling completion temperature (CT2). While the yield strength of the steel increases due to the fine carbides formed during the reheating process of this invention, if the temperature at this time is less than CT2 + 30°C, the tempering effect becomes insufficient. On the other hand, if the temperature exceeds 270°C, the carbides become coarser, leading to a deterioration of bending properties.

Furthermore, if the holding time during the overaging treatment is less than one minute after reheating to the aforementioned temperature range, the martensite will not be sufficiently transformed into tempered martensite, making it difficult to obtain the desired tempering effect. On the other hand, if the holding time exceeds 20 minutes, the carbides formed by overaging will become coarse, reducing the bending properties and potentially adversely affecting the material.

The steel sheet manufactured according to the above method possesses ultra-high strength of 1300 MPa or more due to its microstructure composed of martensite and/or tempered martensite. Furthermore, by controlling the temperature during the continuous annealing process, the cooling process, and the reheating process, an excellent yield ratio can be ensured. Moreover, the formation of a decarburized layer on the surface during the continuous annealing process results in excellent bending properties.

The present invention will be described in more detail below with reference to examples. However, these examples are for illustrative purposes only and do not limit the present invention. The scope of the present invention is determined by the matters described in the claims and matters that can be reasonably inferred therefrom.

(Examples)
Hot-rolled steel sheets were produced by heating steel slabs having the alloy composition shown in Table 1 below to 1100-1300°C, followed by finish hot rolling at 850-950°C, a temperature of Ar3 or higher. Subsequently, each hot-rolled steel sheet was wound at 300-700°C, and then cold-rolled at a cold reduction ratio of 45-65% to produce cold-rolled steel sheets.

Each cold-rolled steel sheet produced as described above was subjected to continuous annealing at a temperature range of 800-900°C for 100-400 seconds, followed by stepwise cooling under the conditions shown in Table 2 below. Subsequently, reheating and overaging treatment were performed under the conditions shown in Table 2 below to produce the final steel sheet. The dew point temperature inside the annealing furnace during the continuous annealing process is also shown in Table 2 below.

Subsequently, the carbon content in the surface layer of the manufactured steel sheets was measured using GDS, and the physical properties were determined through material evaluation. At this time, yield strength, tensile strength, yield ratio, total elongation, and uniform elongation were measured by processing each steel sheet to JIS standards (gauge length width × length: 25 × 50 mm, total specimen length: 200-260 mm) and then performing a tensile test at a test speed of 28 mm/min.

Furthermore, the bending properties (R/t) were determined by processing the same steel plate into a 100 mm wide x 30 mm long specimen and performing a 90° bending test at a test speed of 100 mm/min. Afterward, cracks in the bent area were checked using a microscope, and the R/t value was calculated by dividing the minimum bending radius (R value of the die) without cracking by the thickness of the specimen (t, mm). For the three-point bending test, the maximum bending angle was measured by processing the same steel plate into a 60 mm wide x 30 mm long specimen and performing the test at a test speed of 20 mm/min and a punching radius of 0.4R, as per the VDA238-100 standard, at the maximum load at which cracking occurred.

The microstructure of each steel sheet was then observed using a scanning electron microscope (SEM), and each fraction was measured.

As shown in Tables 1-4 above, Invention Examples 1-3, which satisfy all of the alloy compositions and manufacturing conditions proposed in this invention, exhibited excellent bending properties due to the sufficient formation of a decarburized layer within the surface layer. Furthermore, it was confirmed that the steel sheet possessed ultra-high strength because its main structure was formed from martensite/tempered martensite.

On the other hand, Comparative Examples 1 to 11, which satisfied the alloy composition of the present invention but whose manufacturing conditions, particularly the annealing or reheating conditions, did not meet the requirements of the present invention, exhibited inferior bending properties because a decarburized layer was not formed within the surface layer.

Comparative Example 12 showed low yield strength and tensile strength because, after continuous annealing, the secondary cooling completion temperature was high, and the temperature during reheating did not rise sufficiently. Although a decarburized layer was formed, the tempering effect was insufficient.

Comparative Examples 13 and 14, despite both satisfying the alloy composition requirements of the present invention, exhibited a decarburized layer in the surface layer due to annealing conditions (dew point temperature conditions) that did not meet the criteria of the present invention. This resulted in poor bending properties (approximately 4) and insufficient maximum bending angles during the bending test, thus confirming a deviation from relational equation 2.

Comparative Example 15 showed that, after continuous annealing, the secondary cooling completion temperature was high, and the temperature during reheating was not sufficiently raised. As a result, although a decarburized layer was formed, the tempering effect was insufficient, leading to low yield strength and a poor yield ratio.

Comparative Example 16 did not satisfy the alloy composition system of the present invention, and therefore exhibited inferior yield strength and yield ratio.

Comparative Example 17 also failed to satisfy the alloy composition system of the present invention. Due to insufficient martensite (+ tempered martensite) phase in the steel sheet microstructure, it exhibited inferior yield strength and tensile strength.

Comparative Example 18 is an example that deviates from relational formula 1 of the present invention. Despite the application of the annealing conditions of the present invention, the strength was extremely poor because almost no martensite (+ tempered martensite) phase was formed not only in the surface layer but also in the core.

Figure 1 shows SEM-measured photographs of the microstructure of the surface layer cross-section (up to approximately 80 μm in the thickness direction) of Invention Example 1 and Comparative Example 1.

As shown in Figure 1, in the case of Invention Example 1, it can be confirmed that a decarburized layer containing a soft phase is formed in the surface layer, whereas in Comparative Example 1, it can be seen that a dense hard phase is formed.

Figure 2 shows SEM-measured photographs of the cross-sectional microstructure of the 1/4t region (t: steel plate thickness (mm), based on 1.4 mm) of Invention Example 1 and Comparative Example 1.

As shown in Figure 2, it can be confirmed that both Invention Example 1 and Comparative Example 1 have a martensite (or tempered martensite) phase as the main structure.

Claims (9)

In weight percent, it consists of carbon (C): 0.1-0.3%, manganese (Mn): 1.0-2.3%, silicon (Si): 0.05-1.0%, phosphorus (P): 0.1% or less (excluding 0%), sulfur (S): 0.03% or less (excluding 0%), aluminum (Al): 0.01-0.5%, and two or more of the following: chromium (Cr): 0.01-0.2%, molybdenum (Mo): 0.01-0.2%, and boron (B): 0.005% or less, and one or more of the following: titanium (Ti): 0.1% or less, and niobium (Nb): 0.1% or less, with the remainder being Fe and unavoidable impurities, and satisfies the following relational formula 1.
The microstructure contains martensite and/or tempered martensite phases in an area fraction of 99% or more.
An ultra-high-strength steel sheet with excellent bending properties, having a bending characteristic (R/t) of 3 or less .
[Relationship 1]
(Here, Ceq1=C+(Mn/20)+(Si/30)+(2P)+(4S), Ceq2=C+(Mn/6)+(Si/30)+(Cr+Mo+V+Nb)/5+(Cu+Ni)/15.) The steel sheet is an ultra-high-strength steel sheet with excellent bending properties, as described in claim 1, wherein the ratio of carbon content within a region (A) of 1 to 3 μm in the thickness direction based on the surface (average carbon content in region (A) / carbon content of the steel sheet) is 0.6 or less relative to the carbon content. The steel sheet is an ultra-high-strength steel sheet with excellent bending properties, as described in claim 1, wherein the ratio of carbon content within a region (B) from 0.2 to 30 μm in the thickness direction based on the surface (average carbon content in region (B) / carbon content of the steel sheet) is 0.9 or less relative to the carbon content. The ultra-high-strength steel sheet with excellent bending properties according to claim 1, wherein the content of ferrite and/or bainite phase in the steel sheet is 1% or less (including 0%) by area fraction . The steel sheet, as described in claim 1, is an ultra-high-strength steel sheet with excellent bending properties, wherein the microstructure of the surface layer, which is a region from the surface in the thickness direction from a minimum of 50 μm to a maximum of 70 μm, is composed of tempered martensite with an area fraction of 70% or less (excluding 0%) and the remainder being one or more of ferrite and bainite. The steel plate is an ultra-high-strength steel plate with excellent bending properties as described in claim 1, having a tensile strength of 1300 MPa or more and a yield ratio of 0.72 or more . The steel plate is an ultra-high-strength steel plate with excellent bending properties as described in claim 1, satisfying the following relational equation 2.
[Relationship Equation 2]
(Tensile strength (TS) / Maximum bending angle) ≤ 25 A steel slab containing, by weight percent, carbon (C): 0.1-0.3%, manganese (Mn): 1.0-2.3%, silicon (Si): 0.05-1.0%, phosphorus (P): 0.1% or less (excluding 0%), sulfur (S): 0.03% or less (excluding 0%), aluminum (Al): 0.01-0.5%, two or more of chromium (Cr): 0.01-0.2%, molybdenum (Mo): 0.01-0.2%, and boron (B): 0.005% or less, one or more of titanium (Ti): 0.1% or less, and niobium (Nb): 0.1% or less, with the remainder being Fe and unavoidable impurities, and satisfying the following relational formula 1, is heated in a temperature range of 1100-1300°C;
The step of manufacturing a hot-rolled steel sheet by finishing hot-rolling the heated steel slab with Ar3 or higher;
The step of winding the hot-rolled steel sheet at a temperature of 700°C or lower;
The step of cold-rolling the wound hot-rolled steel sheet at a total reduction ratio of 30-80% to produce cold-rolled steel sheet;
The step of continuously annealing the cold-rolled steel sheet at Ac3 or higher and a dew point temperature of 0 to 20°C for 30 seconds or more;
Following the continuous annealing, a primary cooling step is performed to a temperature range of 550 to 750°C at an average cooling rate of 1 to 10°C/s;
The process includes the steps of: a first cooling followed by a second cooling to a temperature of Ms-190°C or lower at an average cooling rate of 20 to 80°C/s; and a second cooling followed by reheating and then overaging treatment.
The method for manufacturing an ultra-high-strength steel sheet with excellent bending properties according to claim 1 , characterized in that the reheating and overaging steps are performed by heating to a temperature range that satisfies the following relational expression 3.
[Relationship 1]
(Here, Ceq1=C+(Mn/20)+(Si/30)+(2P)+(4S), Ceq2=C+(Mn/6)+(Si/30)+(Cr+Mo+V+Nb)/5+(Cu+Ni)/15.)
[Relationship Equation 3]
CT2+30℃≦A≦270℃
(Here, CT2 represents the secondary cooling completion temperature (°C), and A represents the reheating and overaging temperature (°C).) The method for manufacturing an ultra-high-strength steel sheet with excellent bending properties, as described in claim 8, wherein the over-aging treatment is performed for 1 to 20 minutes.