KR102437909B1 - Cold heading steel material and manufacturing method thereof - Google Patents

Abstract

The present invention relates to a steel material for cold rolling used in steel materials for automobile parts, including steel materials for steel structures and mechanical structures, and the like, and a method for manufacturing the same.
The steel for cold rolling of the present invention is nickel (Ni) 0.2-0.6 wt%, chromium (Cr) 0.2-1.1 wt%, molybdenum (Mo) 0.08-0.25 wt%, titanium (Ti) 0.01-0.04 wt%, vanadium (V) at least one selected from 0.01 to 0.04 wt%, niobium (Nb) 0.001 to 0.2 wt%, aluminum (Al) 0.001 to 0.01 wt%, carbon (C) 0.15-0.50 wt%, silicon (Si) 0.15-0.3% by weight or less, manganese (Mn) 0.5-1.50% by weight, phosphorus 0.04% by weight or less, sulfur 0.05% by weight or less, boron (B) 0.001 to 0.004% by weight, nitrogen (N) 0.004 to 0.015% by weight, oxygen It contains 0.004% by weight or less, and it is a technical idea to include the remainder iron (Fe) and other unavoidable impurities.
Therefore, the present invention improves work softening and work hardening balance elongation by excellent austenite average grain size, subcrystal grain fraction, yield ratio, etc. There is a very useful effect of having moldability and remarkably increasing the deformation range.

Description

Steel material for cold rolling and its manufacturing method

The present invention relates to a steel material for cold rolling and a method for manufacturing the same, and more particularly, to a steel material for cold rolling used in steel materials for automobile parts, including steel materials for steel structures and mechanical structures, and the like, and a method for manufacturing the same.

In order to respond to environmental regulations facing the world recently, the automobile industry is striving to reduce the weight of automobile parts, and in the automobile parts and materials industry, various process routes are available to save energy and enable eco-friendly production throughout the production process. is developing

In particular, in the field of manufacturing materials for cold forging for automobiles and industrial machines, there is an urgent need to develop eco-friendly steel materials that can shorten the production process, reduce energy, improve productivity, and remove environmental pollution.

Existing automotive parts materials for cold forging are mainly heat treated steel, which secures strength and ductility through quenching-tempering (QT) heat treatment after cold forging. The disadvantage of having a high manufacturing cost because it includes problems such as quenching cracks, dimensional non-uniformity, and torsion It is difficult to use efficiently, and there is a problem in that the installation cost and maintenance cost for the heating furnace are additionally generated.

Recently, smart factories are being built as an extension of the concept of unmanned production facilities and management automation according to factory automation, but the smart factory departs from the factory automation concept that achieves optimization for each unit process in the past, and the entire process can be organically linked with each other. By sharing data between each process through the Internet of Things (IoT), active decision-making generated by comprehensive analysis of collected data is delivered to the devices of each process in real time and implemented, production process flexibility suitable for multi-variety complex production is achieved. will secure

In the smart factory, reduction of the factory scale is important in order to facilitate data exchange and analysis between each process, and it is inevitable to reduce the number of production lines and processes in order to reduce the factory scale.

In addition, oil is applied to the surface of the steel to prevent deterioration of surface hardness and corrosion resistance as decarbonization occurs during the reheating process of steel. It became a cause of environmental pollution along with the treatment of oil. At this time, if the heat treatment process is omitted, the production line of the factory is reduced, the efficiency of space utilization of the factory is improved, and the environmental pollution that occurs during the treatment of oil vapor or waste oil generated by heating the oil used in the heat treatment can be prevented. In addition, by saving electrical or chemical energy used for heating, production costs are reduced, and the number of processes can be reduced to further improve production speed, allowing one step forward in building a smart factory.

Bijo was developed to overcome the disadvantages of the heat treatment process after cold forging, space efficiency of the production plant, and the additional cost of installation and maintenance for the heating furnace in the field of parts and materials for automotive cold rolling mentioned above. It is non-heat treated steel.

Non-tempered steel refers to steel that obtains an effect equivalent to that of tempering by effectively using processing or heat and cooling after heating during rolling or hot forging without tempering.

Steel materials that can omit the heat treatment process have been developed by virtue of the above advantages, and the development process of the heat treatment step omitted steel is largely divided into four stages.

The first generation heat treatment omitted steel is a medium carbon steel with added vanadium (V), which is hot-formed to form a ferrite-pearlite structure and then air-cooled. , applied to crankshafts or connecting rods for automobiles that do not receive excessive shock or load during operation.

The second generation of heat treatment omitted steels lower the carbon content and increase the silicon (Si) content to compensate for the low impact toughness of the first generation heat treatment omitted steels. The cooling rate was increased, and the impact toughness was improved by refining the crystal grains of pearlite by adding molybdenum (Mo) and titanium (Ti).

In order to further improve the impact toughness and strength of the 3rd generation steel material without heat treatment, niobium (Nb) and molybdenum (Mo) are added to increase the martensite end temperature to 200°C by the mass effect, and through controlled cooling immediately after hot forming The carbide was uniformly dispersed to form a complex structure of bainite and martensite.

In order to improve impact toughness, strength, and formability to a similar level to that of heat-treated steel, development of the 4th generation non-heat-treated steel is in progress in the direction of increasing the texture fraction of bainite through controlled cooling immediately after hot forming.

However, non-refined steel has a cost increase factor due to the addition of trace alloying elements, and it is difficult to secure the homogeneity of controlled rolling and controlled cooling to control the microstructure, and even if a microstructure with fine grains is implemented, there are limitations in physical properties. .

On the other hand, pre-heat treated steel is a material in which high strength and cold forming properties of steel are improved by controlling the microstructure using alloy design and high frequency induction heating heat treatment method as shown in FIG. 1 .

Since the additional quenching-tempering (QT) heat treatment, which is additionally performed to increase strength after cold forming, can be omitted, so it not only prevents dimensional change and distortion of products, saves energy and improves the working environment, but also uses a simple rolling process to achieve the same level as before. Since it is possible to manufacture mechanical parts having the above strength and toughness, quality control is easy, and process management is simplified, thereby improving productivity.

In general, the limiting deformation amount of the material during cold forging is used in the amount of deformation greater than or equal to the yield strength and less than or equal to the tensile strength. This is because the amount of deformation greater than the tensile strength adversely affects the material properties of the bar, which is the area where material necking occurs.

However, the product (n × YS) of YS (yield strength) and n value (work hardening index) as a new material parameter that affects the cold-rolling characteristics of steel wire or steel bar is 4.0 ~ 11.0Kgf/mm². Tempered steel wire with excellent cold forging characteristics that does not require separate tempering [quenching, tempering] heat treatment after cold forging by maintaining the range, or only work softening in the balance between work hardening and work softening after maximum tensile strength There are disadvantages such as instability of molding dimensions for induction reasons and many restrictions in increasing the molding amount.

2 is a graph according to the balance between work hardening and work softening among tensile properties, (a) is the case where work hardening dominates the deformation as the amount of deformation increases after the maximum tensile strength, (b) is work hardening and work softening When the balance of is maintained, (c) shows the case where work softening dominates the deformation.

When the amount of deformation increases beyond the yield stress, dislocation generation and dislocation recovery (or extinction) coexist in the material, but if dislocation generation dominates, work hardening occurs.

Therefore, work hardening refers to a phenomenon in which the resistance of plastic deformation increases as the generation of dislocation is amplified as the amount of deformation increases. As a concept, as the amount of deformation increases, the recovery of dislocation is amplified and the resistance of plastic deformation decreases.

However, if the balance between generation and recovery of dislocations is maintained, physical properties without work hardening and work softening can be secured.

During cold forging or forging, the material undergoes recovery strain and shrinks. In this case, the problem of deepening dimensional deviation after cold forging (or forging) can be solved by having a constant spring back regardless of the amount of deformation.

2(a) and (c) show that the springback is irregular depending on the amount of deformation, so there is a problem of deepening dimensional deviation after forging. There are limitations in manufacturing.

In particular, in (c) of FIG. 2 , there is a limit in terms of the amount of cold forging deformation that can be applied because it enters the necking region where the fracture process proceeds immediately after the maximum tensile strength.

(0001) Republic of Korea Patent Registration No. 10-0568058 (0002) Republic of Korea Patent Registration No. 10-0568058 (0003) Republic of Korea Patent No. 10-0469671 (0004) Republic of Korea Patent No. 10-0464962 (0005) Japanese Laid-Open Patent Publication No. 59-136420 (0006) Japanese Publication No. 7-54940 (0007) U.S. Patent No. 5,554,233

The present invention has been devised to solve the conventional problems as described above, and an object of the present invention is to increase the amount of cold working during cold working and to manufacture a steel grade that can be applied to parts for cold forging having uniform machining dimensions. As a result of multi-angle research on this method, a round bar or wire rod for cold forging with a yield ratio of 0.95 or more after filigree treatment and without work hardening or work softening during cold working was manufactured, and then cold worked during cold forming or working. An object of the present invention is to provide a steel material for cold rolling having increased amount (strain amount) and dimensional homogeneity, and a method for manufacturing the same.

Another object of the present invention is to optimize the boron effect, which is an element for improving the hardenability of boron, that is, to optimize the effective boron dissolved in the base material that is not dissolved in carbonitride, so that it is manufactured for cold pressing without work hardening and work softening properties. An object of the present invention is to provide a steel material for cold rolling that can improve the hardenability of the steel and a method for manufacturing the same.

Another object of the present invention is to use a steel material having a fine austenite grain size before finish rolling, by controlling the finish rolling temperature, rolling accumulative deformation, and wire rod cooling conditions to refine the grain size of the wire rod, and then work hardening and processing manufactured after heat treatment An object of the present invention is to provide a steel material for cold rolling that can contribute to improving the fraction of fine grains and sub-crystal grains in a steel material for cold rolling without softening properties, and a method for manufacturing the same.

Another object of the present invention is to prevent the distribution of undissolved pearlite after heat treatment in a steel for cold forging without work hardening and work softening, which is produced by increasing the dislocation density by hybrid drawing, and fine austenite grains and sub-crystal grains. An object of the present invention is to provide a steel material for cold forging that can increase the fraction and a method for manufacturing the same.

Another object of the present invention is to provide a steel material for cold forging that can minimize work hardening and work softening during cold forging or forging by applying a manufacturing technology that increases the fraction of fine austenite grains and subcrystal grains, and a method for manufacturing the same there is

The steel material for cold rolling of the present invention for achieving the above object is nickel (Ni) 0.2-0.6 wt%, chromium (Cr) 0.2-1.1 wt%, molybdenum (Mo) 0.08-0.25 wt%, titanium (Ti) 0.01 ~0.04 wt%, vanadium (V) 0.01~0.04 wt%, niobium (Nb) 0.001~0.2 wt%, aluminum (Al) 0.001~0.01 wt% at least one selected from, and carbon (C) 0.15-0.50 wt% %, silicon (Si) 0.15-0.3 wt% or less, manganese (Mn) 0.5-1.50 wt%, phosphorus 0.04 wt% or less, sulfur 0.05 wt% or less, boron (B) 0.001 to 0.004 wt%, nitrogen (N) 0.004 It contains ~0.015 wt%, oxygen 0.004 wt% or less, and it is a technical idea to include the remainder iron (Fe) and other unavoidable impurities.

The steel material for cold rolling of the present invention contains a trace alloy of titanium (Ti), vanadium (V), niobium (Nb) and aluminum (Al), (1.15Ti + 0.99V + 0.58Nb + 1.99Al) / The addition constant of the microalloy expressed by the formula of 7.66N can be formed within the range of 0.86 to 1.0.

At this time, the effective amount of boron expressed by the formula of [5.25B - {7.66N - (1.15Ti + 0.99V + 0.58Nb + 1.99Al)/5.25}] x 10000 may be formed within the range of 7 to 20 ppm.

The method of manufacturing a steel for cold rolling of the present invention for achieving the above object is (a) nickel (Ni) 0.2-0.6 wt%, chromium (Cr) 0.2-1.1 wt%, molybdenum (Mo) 0.08-0.25 wt% , titanium (Ti) 0.01 to 0.04 wt%, vanadium (V) 0.01 to 0.04 wt%, niobium (Nb) 0.001 to 0.2 wt%, aluminum (Al) 0.001 to 0.01 wt% At least one selected from, and carbon ( C) 0.15-0.50 wt%, silicon (Si) 0.15-0.3 wt% or less, manganese (Mn) 0.5-1.50 wt%, phosphorus 0.04 wt% or less, sulfur 0.05 wt% or less, boron (B) 0.001-0.004 wt% , Nitrogen (N) containing 0.004 to 0.015 wt%, oxygen 0.004 wt% or less, forming a billet containing the balance iron (Fe) and other unavoidable impurities; (b) heating the billet; (c) forming a rolled steel material by hot rolling the heated billet; (d) cooling the rolled steel; to include a technical idea.

The step (c) includes a process of rough rolling the heated billet, a process of precision rolling, and a finish rolling process, wherein the average austenite grain size immediately before the finish rolling process is 15 µm or less, and the temperature range of the finish rolling process In the two-phase region of ferrite and austenite, (Ae1+Ae3)/2 temperature ±10℃, the cumulative rolling strain is in the range of 0.6 to 2.0, and the microstructure immediately after the finish rolling is 5㎛ or less of the average ferrite grain size and may have an austenite grain size of 15 μm or less.

In step (d), the rolled steel may be cooled to 450 to 550° C. at a cooling rate of 5 to 20° C./s.

The method of manufacturing a steel material for cold rolling of the present invention may further include a drawing step of (e) drawing the cooled steel material through a die.

At this time, the step (e) is a primary drawing step of drawing by passing the steel material through a first die, a common rolling step of rolling the steel material passing through the first die by passing it through one or more common rolling rolls, passing through the common rolling roll It may include a secondary drawing step of drawing one steel material through a second die.

In addition, the steel is drawn with a reduction in section in the range of 10 to 40% in the first drawing step and rolled at a reduction in section in the range of 10 to 40% in the formal rolling step, and in the range of 10 to 40% in the second drawing step By drawing at a reduction in area of , it can be finally processed so that the cumulative reduction in area of the steel material is 50% or more.

The steel material having completed the secondary drawing step may have a dislocation density of 4.0x10 ⁸ mm ^-2 or more.

The method of manufacturing a steel material for cold rolling of the present invention may further include (f) atmospheric heating or induction heating to achieve austenite heat treatment on the drawn steel material.

At this time, in the austenite heat treatment, the average size of the austenite grain size is 7 μm by atmospheric heating or induction heating in the temperature range of Ac3+ (80±50)° C., which is directly above the austenite transformation point (Ac3). or less, and may include an undissolved pearlite having a fraction of 5% or less.

And, after the austenite heat treatment, the fraction of grain boundaries having a grain boundary angle of 15° or less may be 20% or more.

The steel material is subjected to atmospheric heating or induction heating in the temperature range of Ac3+ (80±50)°C, which is directly above the austenite transformation point (Ac3) for the cold drawn material, and then 20 to 150°C until reaching room temperature. It can be obtained by tempering by quenching at a cooling rate of /sec.

The steel has a hardness of Hv320 to 385 at tempering 550 ° C, a tensile strength of 1015 to 1200 MPa, a yield strength of 980 to 1180 MPa, a yield ratio of 0.97 to 0.99, a work hardening and work softening balance elongation of 6 to 8%, 12 It is preferable to satisfy the range of uniform elongation of ~15%%, total elongation of 19~22%, and tensile reduction of area of 75~80%.

According to the present invention implemented by the above-mentioned solution, the size of the average austenite grain size, the fraction of sub-crystal grains, the yield ratio, etc. are excellent, so that the work softening and work hardening balance elongation is improved, and the uniform elongation is high compared to the prior art. As it appears, it has excellent formability in cold forming and has a very useful effect that can significantly increase the deformation range.

1 is a block diagram showing a conventional heat treatment process.
2 is a graph showing a conventional work hardening and work softening balance.
3 is a graph showing the relationship between the effective amount of boron in the steel according to the embodiment of the present invention and the ideal martensite grain size manufactured.

In order to fully understand the present invention, preferred embodiments of the present invention will be described with reference to the accompanying drawings.

Embodiments of the present invention may be modified in various forms, and the scope of the present invention should not be construed as being limited to the embodiments described in detail below. This embodiment is provided to more completely explain the present invention to those of ordinary skill in the art.

Accordingly, the shape of the components expressed in the drawings may be exaggerated to emphasize a clearer description. It should be noted that the same members in each drawing are sometimes shown with the same reference numerals. In addition, detailed descriptions of well-known functions and configurations determined to unnecessarily obscure the gist of the present invention will be omitted.

Steel for cold rolling according to an embodiment of the present invention, nickel (Ni) 0.2-0.6 wt%, chromium (Cr) 0.2-1.1 wt%, molybdenum (Mo) 0.08-0.25 wt%, titanium (Ti) 0.01-0.04 At least one selected from weight %, vanadium (V) 0.01 to 0.04 weight %, niobium (Nb) 0.001 to 0.2 weight %, aluminum (Al) 0.001 to 0.01 weight %, and carbon (C) 0.15-0.50 weight %, Silicon (Si) 0.15-0.3% by weight or less, manganese (Mn) 0.5-1.50% by weight, phosphorus 0.04% by weight or less, sulfur 0.05% by weight or less, boron (B) 0.001 to 0.004% by weight, nitrogen (N) 0.004 to 0.015 % by weight, containing up to 0.004% by weight of oxygen, and the remainder including iron (Fe) and other unavoidable impurities.

First, the action and content of essential components excluding iron and other unavoidable impurities are as follows.

Carbon is an element beneficial to securing strength, but if its content is 0.15% or less, the strength of the wire is not sufficient, so there is a problem with the strength of the product after final Q/T heat treatment. It is not preferable because the temperature control is unstable during the finish rolling for miniaturization.

Silicon, as a substitutional element, has a great influence on securing the strength of steel. When the silicon content is 0.15% or less, it is difficult to secure the strength of steel, and when it is 0.3% or more, it promotes the generation of decarburized structures during billet and wire rolling, and additional removal costs are required, so that the work hardening rate during cold forging can be increased.

Manganese forms a substitutional solid solution in the matrix structure and lowers the A1 transformation point temperature to increase the fraction of subcrystal grains in martensite or bainite structure. Below 0.5%, the effect is insufficient and sufficient hardenability is secured. At 1.5% or more, the mechanical properties are adversely affected by the tissue heterogeneity caused by manganese segregation, which is easy to cause macrosegregation and microscopic segregation depending on the segregation mechanism during solidification of steel. This is because it promotes central segregation zone due to relatively low diffusion coefficient compared to

Boron is an element that improves hardenability by delaying the precipitation of proeutectoid ferrite during quenching after heat treatment to manufacture the final product. By suppressing the formation of proeutectoid ferrite by segregating atomic boron at the austenite grain boundary and lowering the grain boundary free energy. , promote the formation of martensitic or bainitic structures.

Depending on the material volume, if the boron content is less than 0.001 wt%, the effect of promoting martensite tissue formation becomes insignificant. It is not preferable because it is easy and the formed nitride acts as a nucleation site of ferrite and deteriorates hardenability.

Therefore, in order to prevent the formation of boron nitride, titanium, vanadium, niobium, aluminum, etc. must be added to appropriately secure the effective amount of boron existing in a single atomic state without bonding with nitrogen.

Nitrogen combines with titanium, aluminum, and vanadium contained in structural fastening steel to form carbonitrides, and the formed carbonitrides refine austenite grains to improve strength and toughness of tempered martensitic steel.

At this time, when the nitrogen content is less than 0.004 wt%, it is difficult to obtain the effect of improving the strength and toughness of the tempered martensitic steel, and when the nitrogen content exceeds 0.015 wt%, the carbonitride is coarsened and does not contribute to grain refinement can not do it.

In addition, when the content of titanium, aluminum and vanadium is insufficient compared to the amount of nitrogen contained, the hardenability effect by boron is inhibited while forming boron nitride (BN), and nitrogen is dissolved in the steel material to greatly reduce toughness.

Next, in addition to the above-described essential components, the action and content of additional components that are selectively added at least one of nickel, chromium, molybdenum, titanium, vanadium, niobium and aluminum to improve performance such as strength or toughness of steel are described below. same as

The steel according to the embodiment of the present invention may additionally contain 0.2 to 0.6% by weight of nickel, which nickel improves the hardenability of the steel to stably form a martensitic structure, and without reducing the toughness of the manufactured steel. It can increase strength and improve corrosion resistance.

At this time, when the nickel content is less than 0.2% by weight, the effect of improving the strength and corrosion resistance of the steel and securing the toughness at low temperatures can be obtained. It cannot be obtained, and the manufacturing cost may increase.

The steel according to the embodiment of the present invention may contain 0.2 to 1.1 wt% of chromium, which improves the fatigue strength and wear resistance of the steel, and forms a composite carbide with molybdenum and vanadium to increase impact resistance.

This is because when the chromium content is less than 0.2% by weight, it becomes difficult to obtain the effect according to the chromium content, and when the chromium content exceeds 1.1% by weight, the effect becomes insignificant.

The steel material according to the embodiment of the present invention may additionally contain 0.08 to 0.25 wt% of molybdenum, which molybdenum increases the hardenability of the steel to obtain a martensitic structure stably, and molybdenum carbide is manufactured by refining the grain size improve the strength and toughness of the used steel.

If the cooling rate is slow during the manufacturing process of steel, the toughness of the steel manufactured by dispersing the coarse precipitates of molybdenum may decrease. have.

At this time, when molybdenum is contained together with boron, hardenability is controlled during cooling to optimize the balance between tensile strength and toughness.

In addition, when the molybdenum content exceeds 0.25% by weight, hardenability is increased more than necessary and the effect is no longer increased to a saturated state.

In particular, since boron and molybdenum are expensive elements, it is preferable not to contain more than the necessary amount because the content of more than necessary increases the production cost of the steel greatly.

The steel according to the embodiment of the present invention may additionally contain 0.01 to 0.04 wt% of titanium, and this titanium binds to nitrogen in the steel to which boron is added and fixes nitrogen, thereby suppressing the generation of boron nitride (BN). It is possible to secure the effective amount of boron in the atomic state and to improve the hardenability effect by boron.

TiN, the most stable among the grain refining element compounds, has low solid solubility at high temperature and slow grain growth rate, which can contribute to grain refinement.

When the titanium content is less than 0.01% by weight, the above effect becomes insignificant, and when the titanium content exceeds 0.04% by weight, the improvement effect reaches a saturated state, so that the improved effect is not exhibited, and the toughness of the manufactured steel is reduced.

The steel material according to the embodiment of the present invention may additionally contain 0.01 to 0.04 wt % of vanadium, which improves the strength of the steel through fine carbide formed by bonding with carbon in the steel, and nitrogen at a temperature of 900 ° C. or higher. It binds with and forms vanadium carbonitride (VC, VCN) to suppress the formation of boron nitride (BN) to control the appropriate effective amount of boron, and to prevent grain growth of the austenite phase to secure an appropriate martensite structure. It plays an important role in entry control.

Vanadium improves the strength of the manufactured steel through precipitation strengthening and dispersion strengthening while forming fine precipitates of vanadium carbonitride inside the ferrite structure while the steel is cooled.

At this time, when the content of vanadium is lower than 0.01 wt %, the effect caused by the vanadium content is insignificant, and when the content of vanadium exceeds 0.04 wt %, the improvement effect does not increase any more, and the saturation state is reached.

In particular, when coarse carbonitride is formed according to an excessive amount of vanadium, the toughness of the manufactured steel is reduced and the steel is brittle, so it is preferable to contain an appropriate amount so as not to deviate from the appropriate content range.

The steel material according to the embodiment of the present invention may additionally contain 0.01 to 0.04 wt% of niobium, and this niobium has a fixing effect while niobium carbonitride (NbC, NbN) is precipitated at the grain boundary during a hot forming process such as rolling or forging of the steel material. appears to refine crystal grains and improve the strength and toughness of the martensitic structure.

In addition, since niobium suppresses the formation of boron nitride (BN) while bonding with nitrogen, the effective amount of boron is controlled to improve hardenability of steel.

At this time, when the content of niobium is contained in less than 0.01 wt%, it is difficult to compensate for the fixing effect due to niobium carbonitride and the effect of improving hardenability accompanied by lowering the carbon content, martensite transformation is not easy, and the content of niobium When it exceeds 0.04 wt%, it is preferable to contain an appropriate amount of niobium since coarse niobium carbonitride is formed and the toughness of the martensite structure may be reduced.

The steel material according to an embodiment of the present invention may additionally contain 0.001 to 0.01% by weight of aluminum. This aluminum acts as a strong deoxidizer to remove oxygen contained in the steel while forming aluminum oxide, thereby refining martensite grains. do

In addition, aluminum suppresses the formation of boron nitride (BN) while bonding with nitrogen, thereby improving the hardenability of the steel by controlling an appropriate effective amount of boron.

At this time, when the content of aluminum is less than 0.001 wt%, the effect of deoxidation or martensite grain refining action is reduced, which is undesirable, and when the content of aluminum exceeds 0.04 wt%, the amount of aluminum oxide non-metallic inclusions increases. It may cause a decrease in the toughness of the manufactured steel or clogging of the nozzle during casting.

In the steel material according to the embodiment of the present invention, the addition constant of the trace alloy containing titanium (Ti), vanadium (V), niobium (Nb) and aluminum (Al) is preferably in the range of 0.8 to 1.0, these trace amounts The addition constant of the alloy is expressed by the formula below.

Trace alloy addition constant = (1.15Ti + 0.99V + 0.58Nb + 1.99Al) / 7.66N

When the microalloy addition constant is less than 0.8, the effect of boron addition decreases, making it difficult to secure a martensitic structure. When the microalloy addition constant exceeds 1.0, the effect of boron addition reaches a saturated state, As this is formed, the toughness of the manufactured steel may be reduced.

In the case of including titanium (Ti), vanadium (V), niobium (Nb) and aluminum (Al) in the steel material according to an embodiment of the present invention, forming an effective boron amount satisfying the range of 7 to 20 ppm Preferably, such an effective boron amount is derived from the following formula.

Effective boron amount (ppm) = [5.25B - {7.66N - (1.15Ti + 0.99V + 0.58Nb + 1.99Al)/5.25}] x 10000

The relationship between the effective amount of boron in steel and the ideal manufactured martensite grain size is shown as in the graph of Fig. 3, which can induce stable martensite transformation by most effectively delaying ferrite transformation at the effective boron content within the box range indicated by the dashed line in this graph. If the effective boron amount according to the effective boron content formed in the broken-line box of FIG. 3 is obtained through the above formula, a value in the range of 7 to 20 ppm can be obtained.

At this time, when the effective amount of boron is less than 7 ppm, the effect of adding boron becomes insignificant, and when the effective amount of boron exceeds 20 ppm, hardenability of the spring steel decreases and it becomes difficult to secure a martensitic structure.

The method for manufacturing a steel material for cold forging configured as described above includes (a) forming a billet including the above-described essential components and additional components, (b) heating the formed billet to be maintained in an austenite single-phase region; (c) hot-rolling the heated billet to increase the sub-crystal grain fraction in ferrite and forming it into a rolled steel, (d) cooling the formed rolled steel, (e) reducing the section of the cooled steel by 10 to 40 The step of drawing to maintain %, (f) consists of a step of atmospheric heating or induction heating the drawn steel material.

The step (b) of forming a steel material including the above-described essential and additional components into a billet and heating the billet again is preferably maintained within 90 minutes at that temperature after heating to a temperature range of 900 to 1050 °C. That is, to maintain the billet in the austenite single phase region, the temperature range is a range in which the austenite grains are not coarsened, and is effective in removing the remaining coarse cementite.

This is because if the heating temperature in step (b) exceeds 1050 ° C, the austenite grains are coarsened and it is difficult to secure ultra-fine ferrite after subsequent wire rolling. This is because it is inferior and there is a problem in that the rolling load increases.

And the heating holding time of step (b) is preferably in the range of 90 to 120 minutes. If the heating temperature holding time is less than 90 minutes, the target temperature may not be reached due to insufficient heating to the center of the steel billet. Deteriorated tissue may increase in the

Step (c) forms a rolled steel material including a round bar or wire rod by hot rolling the billet heated in step (b).

In step (c), it is preferable to perform hot rolling of the billet in a two-phase region of ferrite and austenite in order to increase the subgrain fraction in ferrite of the microstructure.

The hot rolling of step (c) may include a process of rough rolling the heated billet, a precision rolling process, and a finish rolling process.

Here, in the finish rolling process, it is preferable to proceed in a temperature range of (Ae1 + Ae3)/2, which is a two-phase region of ferrite and austenite, and to carry out the cumulative deformation by rolling to generate 0.6 or more.

In this finishing rolling process, when the inlet temperature of the rolling mill is less than (Ae1 + Ae3)/2, supercooling or rapid cooling occurs during rolling, and decarburization structures such as ferrite and surface defects may rapidly increase on the surface of the manufactured round bar. When the temperature of (Ae1 + Ae3)/2 exceeds (Ae1 + Ae3)/2, the fraction of sub-crystal grains (modified ferrite) may decrease, so that the target grain size may not be secured.

In addition, when the cumulative deformation amount of the finish rolling is less than 0.6, the driving force required for ferrite nucleation for grain refinement is lowered. Meaning, the sub-crystal grains act as nucleation sites for austenite grains for reverse transformation during the final heat treatment.

Hot rolling is performed so that the fraction of sub-crystal grains in the ferrite phase of the round bar prepared in this way is 20-50 area%. If it is less than that, there may be a problem in that the crystal grain refining effect is not secured during the heat treatment process.

In addition, the ferrite average grain size of the manufactured round bar is preferably formed to be 8 μm or less, so that it is possible to secure austenite fine grains of 15 μm or less, and the fine grain boundaries affect delayed fracture. The amount of hydrogen and impurities are relatively reduced to maintain the grain boundary strength.

This is because, when the average ferrite grain size exceeds 8 μm, it is difficult to secure a target sub-crystal grain fraction.

Meanwhile, the austenite grain size immediately before the finish hot rolling is preferably 20 μm or less because it affects the austenite grain size after the wire rod is cooled.

Since the austenite grain size just before finish hot rolling is similar to the austenite grain size after the wire rod is cooled through finish hot rolling, the finer the austenite grain size, the more the finish rolling temperature and cumulative deformation are optimized to induce a lot of deformation during rolling. This is because the ferrite nucleation site can be maximized, and thus the grains are refined to obtain ferrite with an average grain size of 8 μm or less.

The steel manufacturing method for cold rolling of the present invention may include the step (d) of cooling the round bar or wire rod formed by the finish hot rolling under the above-described conditions. In this step (d), ferrite having an average grain size of 8 μm or less is formed. After this, it is preferable to perform cooling at a cooling rate of 5 to 20 °C/s.

When the cooling rate of the hot-rolled round bar is less than 5°C/sec, grain coarsening occurs due to deformation and transformation heat, and pearlite fragments cannot be sufficiently induced. It is preferable that the cooling of the hot-rolled round bar is made within a rate range of ~ 20 °C/sec.

At this time, it is preferable to end the transformation in the cooling zone, and to limit the temperature at which cooling is terminated so as to secure the required microstructure within the range of 200 ~ 600 ℃.

The present invention finely forms the average grain size of austenite before finishing hot rolling to secure fine grain FGS of 9 μm or less, and reduces the ferrite nucleation site by maximizing the ferrite nucleation site during finish hot rolling. can

In the subsequent step (e), drawing is performed to maintain the reduction rate of the section of the cooled steel material at 10 to 40%. In this step (e), the steel material including the cooled round bar or wire rod is passed through the first die to draw the first die. It includes a drawing step, a ball rolling step in which the steel that has passed through the first die is passed through at least one ball rolling roll to be rolled, and a secondary drawing step in which the steel material passed through the ball rolling roll is passed through a second die to be drawn.

In order to maintain the balance between work hardening and work softening, the fraction of fine austenite grains and subcrystal grains must be properly combined. Heat treatment is required at AC3+ (30~80℃), and it is to generate a significant amount of dislocation density in order to secure an austenite single phase without undissolved pearlite at a low heat treatment temperature, and promote atom diffusion during heat treatment become a means to

Here, the reason for limiting the cross-sectional reduction rate of the so-called hybrid drawing process including primary drawing, ball rolling, and secondary drawing is as follows.

It is preferable that the reduction ratio of the primary drawing area is 10 to 40%. If the area reduction rate is less than 10%, it is difficult to secure the proper dislocation density required for crystal grain refinement. If it exceeds 40%, chevron cracks occur due to the increase in central triaxial stress and microbands are generated due to mechanical processing instability. Because it adversely affects the physical properties.

The purpose of giving the ball rolling process to the primary drawn material is to provide a strain path that is almost perpendicular to the deformation direction by drawing. In the drawing process and ball rolling, the direction of the strain path is almost vertical to a large extent. Drawing processing has a large amount of processing in the linear direction (X-axis), and ball rolling has a large amount of strain in the Y-axis and Z-axis directions.

Ball rolling having a deformation direction different from the drawing deformation direction promotes recovery and disappearance of a significant amount of dislocations generated by the drawing process.

Due to this, since the triaxial stress and work hardening rate applied to the center of the material during the primary drawing process are significantly lowered, and even if the secondary drawing process is given, chevron cracks and micro shear bands in the center are not generated. it is possible

The reduction in area of ball rolling refers to the time when the wire rod passes through an elliptical ball rolling roll with an elliptical cross section preferentially, and then finally passes through a round ball rolling roll with a round cross section.

It is preferable that the area reduction ratio of ball rolling be 10 to 40%. Cold rolling is performed while passing through a pair of horizontal and vertical rolling rolls (oval and round type), and when the reduction in area is less than 10%, it is difficult for the material to penetrate the round type rolling roll. As the strain path changes, the effect of recovering and accelerating the disappearance of dislocations generated by the drawing process is insufficient. This is because the reduction effect is insufficient.

This is because, when the cross-sectional reduction rate of ball rolling exceeds 40%, the effect of removing the triaxial stress applied to the center during the primary drawing process and the effect of reducing the work hardening rate are saturated.

It is preferable that the reduction ratio of the secondary drawing area be 10 to 40%. If the area reduction rate is less than 10%, it is difficult to secure the proper dislocation density required for crystal grain refinement. If it exceeds 40%, chevron cracks occur due to the increase in central triaxial stress and microbands are generated due to mechanical processing instability. Because it adversely affects the physical properties.

This is because it is difficult to secure the proper dislocation density required for grain refinement below the lower limit of the reduction in area, and above the upper limit, chevron cracks occur due to an increase in central triaxial stress and micro-shear bands are generated due to mechanical processing instability, resulting in mechanical properties because it has a detrimental effect on

It is preferable that the dislocation density distributed in the material according to the hybrid drawing process be 4x10 ⁸ mm ^-2 or more.

When the dislocation density is less than about 4x10 ⁸ mm ^-2 , in the range of Ac3~Ac3+80℃, which is directly above the austenite transformation point (Ac3), pearlite is partially composed of the initial microstructure of the drawn workpiece during austenite heat treatment in normal induction and atmospheric heating methods. This is because (ferrite + pearlite) is not completely re-dissolved into austenite during heat treatment, and some remains, adversely affecting the toughness of mechanical properties.

In addition, if the dislocation density is less than 4x10 ⁸ mm ^-2 , even if grain refinement is achieved, it is difficult to secure an appropriate fraction of sub-crystal grain boundaries with an effective grain boundary angle of 15° or less for work hardening and work softening balance.

The last step (f) proceeds is a step of atmospheric heating or induction heating of the drawn steel material.

This step (f) is a step in which austenite heat treatment of the steel is performed after step (e), and the austenite heat treatment is preferably performed in the range of Ac3 ~ Ac3+80° C., which is directly above the austenite transformation point (Ac3). Heat treatment below the austenite transformation point (Ac3) is because the cold working microstructure is not completely austenitized, and at Ac3+80℃ or higher, the grains grow rapidly and it is impossible to secure the austenite grains below 7㎛. Because.

The reason for limiting the undissolved pearlite fraction during austenite heat treatment is as follows.

During austenite heat treatment of the wire rod, it is preferable to limit the undissolved pearlite to have a fraction of 5% or less. This is because, when the undissolved pearlite fraction exceeds 5%, coarse cementite, a constituent factor of undissolved pearlite distributed in the heat-treated material, has a negative effect on toughness.

The reasons for limiting the austenite grain boundary properties after heat treatment are as follows.

After the austenite heat treatment of the wire rod, it is preferable that the fraction of grain boundaries having small-angle grain boundaries of 15° or less is 20% or more. This is because when the fraction of grain boundaries having small-diameter grain boundaries of less than 15° is less than 20%, the effect of controlling the work hardening and work softening balance is insufficient.

In general, the process of forming grain boundaries through recrystallization of the material during heat treatment after plastic deformation is as follows.

Dislocations are generated due to plastic deformation and they become tangled with each other, and the dislocations entangled during heat treatment undergo a recovery process in which the dislocations are organized, forming a dislocation cell, and dislocations distributed within the cell Crystal grains are formed through annihilation process. When the dislocation density is sufficient, a grain having a high angle grain boundary with a large grain boundary angle is formed, and when the dislocation density is insufficient, a low angle grain boundary is formed. is formed

When these grain boundaries are plastically deformed, high-hardness grain boundaries increase work hardenability as dislocations are accumulated (file-up). It has the effect of increasing the work softening property.

In the present invention, as a means for controlling the balance between work hardening and work softening, work softening was imparted by distributing a fraction of small-angle grains by 20% or more while appropriately controlling work hardenability by refining the size of the high-angle grain boundary.

The reason for limiting the austenite grain size after heat treatment is as follows.

It is preferable that the average grain size after the austenite heat treatment be 7 μm or less. In order to control the work hardening and work softening balance, it is required to set the fraction of grain boundaries having a grain boundary angle of 15° or less to 20% or more. This is because it is difficult to obtain a fraction of

Hereinafter, the configuration and effect of the present invention will be described in more detail through Examples, Comparative Examples and Experimental Examples.

<Example>

A billet having the ingredients shown in Table 1 below was prepared.

Component content (wt%) of billet by example C Si Mn Cr B Ni Mo Ti V Nb Al N Example 1 0.15 0.20 1.00 - 0.0018 - - 0.025 0.025 0 0 0.007 Example 2 0.30 0.20 0.85 - 0.0025 - - 0.02 0.01 0.01 0.005 0.007 Example 3 0.40 0.30 1.10 0.0038 0.031 0.03 0.029 0.009 0.010 Example 4 0.25 0.25 1.30 0.002 0.034 0.014 0.01 0.001 0.008 Example 5 0.48 0.20 1.40 0.0015 0.015 0.01 0.01 0.01 0.007 Example 6 0.40 0.20 0.90 0.50 0.0013 0.03 0.024 0.02 0.01 0.008 Example 7 0.20 0.25 0.91 0.80 0.0015 0.01 0.02 0 0.01 0.007 Example 8 0.45 0.20 0.90 0.45 0.004 0.3 0.1 0.03 0.02 0.02 0.01 0.013 Example 9 0.16 0.20 0.90 0.40 0.002 0.45 0.1 0.029 0.02 0 0 0.007 Example 10 0.30 0.20 0.90 0.40 0.0035 0.45 0.1 0.03 0.03 0.03 0.01 0.0151

<Comparative example>

A billet having the ingredients shown in Table 2 below was prepared.

Component content of billet by comparative example (wt%) C Si Mn Cr B Ni Mo Ti V Nb Al N Comparative Example 1 0.16 0.23 1.05 - 0.0018 - - 0.015 0.025 0 0 0.007 Comparative Example 2 0.31 0.21 0.84 - 0.026 - - 0.02 0.01 0.01 0.001 0.007 Comparative Example 3 0.41 0.32 1.11 0.0036 0.03 0.03 0.002 0.01 0.010 Comparative Example 4 0.25 0.25 1.30 0.002 0.01 0.001 0 0.005 0.009 Comparative Example 5 0.48 0.20 1.40 0.0015 0.01 0.01 0.01 0.01 0.007 Comparative Example 6 0.40 0.20 0.90 0.50 0.0013 0.01 0.001 0.01 0.01 0.009 Comparative Example 7 0.20 0.25 0.91 0.80 0.0017 0.01 0.02 0 0.007 0.007 Comparative Example 8 0.45 0.20 0.90 0.45 0.0042 0.3 0.1 0.05 0.02 0 0.01 0.012 Comparative Example 9 0.16 0.20 0.90 0.40 0.0022 0.45 0.1 0.018 0.02 0 0 0.007 Comparative Example 10 0.30 0.20 0.90 0.40 0.0035 0.45 0.1 0.03 0.02 0.03 0.01 0.014

The billets of Examples and Comparative Examples were subjected to ultra-fine rolling and cooling steps shown in Tables 3 and 4 below to prepare 18mmΦ wire rods, respectively.

At this time, in manufacturing each wire rod, the austenite average grain size (AGS) immediately before hot rolling, the ferrite average grain size (FGS) and the austenite average grain size (AGS) after the wire rod cooling were calculated by 1/2D ( Here, D means the diameter of the wire rod), and the measurement method was measured according to the KS standard (KS D 0205).

For the hybrid drawing process, 18 mmΦ wire rod was used for each condition, and primary drawing with a 13-40% reduction in area was applied using a drawing die under normal conditions, and the reduction in area was in the range of 0 to 38% under normal cold rolling conditions. Afterwards, the secondary drawing was carried out in the range of 14 to 33% of the reduction in area under normal conditions, and the total cumulative reduction in area was carried out at 20 to 72%.

The method and evaluation of processing dislocation density after hybrid drawing and normal drawing processing were evaluated by measuring the number of dislocations crossing per unit area using a transmission electron microscope. After austenite heat treatment, the fraction of undissolved pearlite tissue was evaluated using an image analyzer.

The sub-crystal grain fraction was evaluated using EBSD (Electron Backscatter Diffraction), which can measure the azimuth difference angle between crystals. At this time, since the grain boundary angle is 15° or less, the grain boundary angle change due to grain refinement was evaluated based on 15°.

In order to investigate the tensile properties and the reduction in area of the materials prepared as described above, a KS standard (KS B0801) tensile test specimen was used, and the tensile test was performed at a cross head speed of 5 mm/min.

The finish hot rolling of the examples was (Ae1 + Ae3)/2 ± 10 ℃ temperature range, so that the cumulative hot rolling deformation amount was included in the range of 0.6 ~ 1.9, the comparative examples, the finish hot rolling temperature is Ae3 + (160 ~ 187) ℃ range The cumulative deformation amount of hot rolling was included in the range of 0.2~0.4. Here, Ae1 and Ae3 mean A1 and A3, which are the phase transformation temperatures of the steel during cooling, respectively.

Tables 3 and 4 below are results of measuring ferrite grains and austenite grain sizes according to effective boron and AGS before finish rolling, finish rolling temperature, finish rolling cumulative deformation amount, and wire rod cooling conditions for Examples and Comparative Examples.

Example: Wasterite grain size and austenite grain size after rolling boron addition
optimization rolling condition wire rod
Cooling
Condition
(℃/s) After cooling the wire rod
grain size a very small amount
alloy
adding
a constant available
boron
(ppm) Wrap-up
before rolling
AGS
(μm) Ae1 Ae3 (Ae1+Ae3)
/2 Wrap-up
rolling temperature
range Wrap-up
rolled stacked
deformation amount
(ε) FGS
(μm) AGS
(μm) Example 1 1.00 18 18 718 825 772 772±10 1.9 14±5 3±2 7±2 Example 2 0.91 16 16 720 795 757 757±10 14 12±5 3±2 7±2 Example 3 0.86 18 18 720 777 748 748±10 0.9 11±5 4±2 9±2 Example 4 0.99 19 20 716 798 757 757±10 0.8 10±5 5±2 11±2 Example 5 0.99 14 18 714 749 732 732±10 0.6 13±5 7±2 13±2 Example 6 0.72 8 19 728 768 748 748±10 0.8 11±5 5±2 12±2 Example 7 0.95 10 15 734 806 770 770±10 1.3 9±5 3±2 7±2 Example 8 0.86 14 17 722 769 746 746±10 0.9 12±5 4±2 10±2 Example 9 0.99 19 15 718 818 768 768±10 1.6 13±5 3±2 7±2 Example 10 0.88 9 18 718 794 756 756±10 0.8 14±5 5±2 13±2

Comparative Example: Size of wastelite grains and austenite grains after rolling Optimization of boron addition rolling condition wire rod
Cooling
Condition
(℃/s) After cooling the wire rod
grain size a very small amount
alloy
adding
a constant available
boron
(ppm) Wrap-up
before rolling
AGS
(μm) Ae3 Wrap-up
rolling temperature
range Wrap-up
rolling accumulation
deformation amount
(ε) FGS
(μm) AGS
(μm) Example 1 0.78 5 20 825 927±10 0.3 14±5 11±2 21±2 Example 2 0.76 One 20 795 957±10 0.2 12±5 11±2 21±2 Example 3 1.11 36 19 777 948±10 0.4 11±5 10±2 20±2 Example 4 0.33 0 30 798 957±10 0.2 10±2 16±2 31±2 Example 5 0.88 3 28 749 932±10 0.4 13±5 14±2 29±2 Example 6 0,55 0 29 768 948±10 0.3 11±5 15±2 29±2 Example 7 0.84 One 35 806 970±10 0.4 9±5 15±2 31±2 Example 8 1.18 42 28 769 946±10 0.3 12±5 15±2 29±2 Example 9 0.75 3 28 818 968±10 0.2 13±5 16±2 25±2 Example 10 0.85 5 30 794 956±10 0.3 14±5 15±2 30±2

From the results of Tables 3 and 4, effective boron, which is very important for ensuring hardenability, was found to be superior to that of Comparative Examples in the range of 7 to 20 ppm in the Examples.

The austenite grains before finish rolling are fine in the range of 28 to 38 μm in Comparative Examples, but in the range of 8 to 15 μm in Examples, so that a lot of proeutectoid ferrite can be generated at the grain boundary by the amount of deformation during finish rolling. The grain size of the final wire rod can be made even finer by the generation and growth of ferrite.

It can be seen that the ferrite grains and austenite grain sizes of the Examples after the wire rolling and cooling were 1 to 7 μm and 5 to 15 μm, whereas the comparative examples were 13 to 18 μm and 27 to 33 μm, showing a significant difference. can

For the examples in which the processing was performed under the hybrid drawing conditions of the present invention and the comparative examples under the normal drawing conditions, heat treatment was performed at a heating temperature of Ac3+30 to Ac3+125℃ to determine the undissolved pearlite tissue fraction, Austenite average grain size and sub-crystal grain fraction were measured, and the results are shown in Tables 5 and 6 below.

Austenite grain size and sub-crystal grain fraction after wire drawing and induction heating by Example 1st drawing
*RA
(%) ball rolling
*RA
(%) 2nd drawing
*RA
(%) total accumulation
*RA
(%) dislocation density
(Number/㎟) austenite heat treatment temperature
**Ac ₃ +Temperature (℃) Unemployed pearlite fraction (%) austenite
Average grain size (μm) Subcrystal grain fraction (%) Example 1-1 20 - - 20 4.0×10 ⁸ 125 0 7 25 Example 1-2 40 - - 40 4.6×10 ⁸ 75 0 6 30 Examples 1-3 37 10 14 51 6.4×10 ⁸ 30 0 5 40 Example 2 13 38 30 62 7.3×10 ⁸ 70 0 4 35 Example 3 37 34 33 72 8.8×10 ⁸ 70 0 3 35 Example 4 13 20 30 51 6.2×10 ⁸ 70 0 5 35 Example 5 27 20 24 56 6.4×10 ⁸ 70 0 5 35 Example 6 27 28 27 62 7.6×10 ⁸ 70 0 4 35 Example 7-1 20 - - - 4.0×10 ⁸ 125 0 7 20 Example 7-2 40 - - - 4.7×10 ⁸ 75 0 6 25 Example 7-3 27 20 24 56 6.8×10 ⁸ 30 0 5 40 Example 8 27 13 22 51 6.3×10 ⁸ 70 0 5 35 Example 9 27 20 24 56 6.8×10 ⁸ 70 0 5 35 Example 10 13 20 30 51 6.5×10 ⁸ 70 0 5 35

(Note 1) * RA (Reduction of Area): Area reduction rate

(Note 2) ** A _C3 + Temperature: A _C3 is the transformation point at which it becomes austenite single phase upon heating, and there is a difference depending on the carbon content and the amount of alloying elements added _. set and heat-treated.

(Note 3) A _C3 = 910-203(C) ^1/2 -15Ni+44.5Si+104V+31.5Mo+13.1W

Austenite grain size and sub-crystal grain fraction after wire drawing and induction heating for each comparative example Area reduction ratio
*RA
(%) dislocation density
(Number/㎟) austenite heat treatment temperature
**Ac ₃ +Temperature (℃) Unemployed pearlite fraction (%) austenite
Average grain size (μm) Subcrystal grain fraction (%) Comparative Example 1-1






20±3
2.0×10 ⁸ 125 5 18 13 Comparative Example 1-2 75 10 17 15 Comparative Example 1-3 30 25 13 18 Comparative Example 2 2.3×10 ⁸ 125 5 18 7 Comparative Example 3 2.4×10 ⁸ 125 7 17 7 Comparative Example 4 2.1×10 ⁸ 125 6 28 7 Comparative Example 5 2.3×10 ⁸ 125 8 25 7 Comparative Example 6 2.2×10 ⁸ 125 8 25 7 Comparative Example 7-1
2.1×10 ⁸ 125 6 28 7 Comparative Example 7-2 75 10 24 10 Comparative Example 7-3 30 25 19 15 Comparative Example 8 2.4×10 ⁸ 125 5 25 7 Comparative Example 9 2.0×10 ⁸ 125 6 21 7 Comparative Example 10 2.3×10 ⁸ 125 5 27 7

From the results of Tables 5 and 6, the dislocation density for obtaining fine grains according to the lowering of the heating temperature without remaining undissolved pearlite is 4.0 x 10 ⁸ to 7.6 x 10 ⁸ in the examples, but compared In the case of the examples, it can be seen that the dislocation density of the examples is remarkably excellent in the range of 2.0 x 10 ⁸ to 2.4 x 10 ⁸ .

And from the results of Examples and Comparative Examples, heat treatment at Ac3+30 ~ Ac3+125℃, which is a temperature condition for grain refinement, in the Examples, no undissolved pearlite remained, but in the comparative examples, 5 to 25% range remained as

In addition, in the austenite grain size after wire drawing and induction heating, the Examples showed a range of 3 to 7 μm, but Comparative Examples appeared to have a range of 13 to 17 μm, and the austenite grain boundary angle was 15° or less. In the sub-crystal grain boundary fraction having

Compared with Examples and Comparative Examples, it can be seen that a large number of fine austenite grain sizes as well as sub-crystal grains are generated. will affect

Tables 7 and 8 below show the results of evaluating the tensile properties at a tempering temperature of 550°C after austenite quenching.

Tensile properties at tempering temperature (550°C) after austenite quenching according to Examples Hardness
(Hv) The tensile strength
(MPa) yield strength
(MPa) yield ratio
(YS/TS) work hardening and
work softening
balance elongation Uniform elongation (%) total elongation
(%) Seal
section reduction rate
(%) Processing
hardening index Example 1-1 323 1016 986 0.97 6 12 20 77 0.03 Example 1-2 327 1026 1008 0.98 7 13 19 75 0.02 Examples 1-3 337 1050 1025 0.98 8 14 21 80 0.01 Example 2 354 1110 1088 0.98 8 15 21 80 0.01 Example 3 366 1150 1127 0.98 8 14 20 75 0.01 Example 4 347 1090 1068 0.98 8 15 22 80 0.01 Example 5 380 1190 1166 0.98 7 13 21 75 0.01 Example 6 368 1160 1148 0.99 7 14 21 80 0.01 Example 7-1 333 1042 1011 0.97 6 13 20 73 0.03 Example 7-2 338 1053 1021 0.97 7 14 21 75 0.02 Example 7-3 341 1070 1049 0.98 8 15 22 80 0.01 Example 8 381 1200 1176 0.98 7 13 19 75 0.01 Example 9 325 1020 1000 0.98 8 15 22 80 0.01 Example 10 345 1080 1058 0.98 8 15 21 80 0.01

Tensile properties at tempering temperature (550℃) after austenite quenching according to comparative examples Austenite heating temperature
(℃) Hardness
(Hv) The tensile strength
(MPa) yield strength
(MPa) yield ratio
(YS/TS) work hardening
and
work softening
balance elongation Uniform elongation (%) total elongation
(%) Seal
section reduction rate
(%) Processing
hardening index Comparative Example 1 950 274 866 771 0.89 0 5 17 70 0.32 Comparative Example 2 850 305 960 864 0.90 0 6 18 70 0.25 Comparative Example 3 850 318 1000 920 0.92 0 7 18 70 0.15 Comparative Example 4 850 297 940 846 0.90 0 7 19 70 0.25 Comparative Example 5 800 333 1040 946 0.91 0 8 18 70 0.25 Comparative Example 6 820 309 1010 919 0.91 0 8 18 70 0.25 Comparative Example 7 950 292 920 828 0.90 0 5 17 70 0.25 Comparative Example 8 820 337 1050 966 0.92 0 5 15 70 0.15 Comparative Example 9 870 276 870 792 0.89 0 7 19 70 0.33 Comparative Example 10 850 295 930 837 0.90 0 7 18 70 0.25

From the results of Tables 7 and 8, the yield ratio of the Examples was 0.97 to 0.98, which was significantly higher than that of the comparative examples, and the work hardening and work softening balance elongation-work hardening and work hardening were similar, so that the strength increased as the strain increased. It can be seen that the invention materials are significantly superior to the comparative examples at 6-8% in the section that appears constantly.

In general, as the amount of deformation increases, the necking phenomenon occurs in which the strength gradually decreases after the maximum tensile strength. The elongation before necking is called the uniform elongation, and the total elongation including the elongation from necking to fracture is called total elongation.

Considering that the uniform elongation of the Examples is at the level of 12 to 15%, but the Comparative Examples are at the level of 5 to 8%, the Examples of the present invention can significantly increase the deformation range for molding while having excellent formability in cold forming. It can be seen that

The embodiments of the present invention described above are merely exemplary, and those of ordinary skill in the art to which the present invention pertains will appreciate that various modifications and equivalent other embodiments are possible therefrom.

Therefore, it will be well understood that the present invention is not limited to the form mentioned in the above detailed description.

Accordingly, the true technical protection scope of the present invention should be defined by the technical spirit of the appended claims. Moreover, it is to be understood that the present invention covers all modifications, equivalents and substitutions falling within the spirit and scope of the invention as defined by the appended claims.

Claims (11)

delete delete delete In the manufacturing method of the steel material for cold rolling,
(a) Nickel (Ni) 0.2-0.6 wt%, Chromium (Cr) 0.2-1.1 wt%, Molybdenum (Mo) 0.08-0.25 wt%, Titanium (Ti) 0.01-0.04 wt%, Vanadium (V) 0.01-0.04 Weight %, niobium (Nb) 0.001 to 0.2 weight %, aluminum (Al) 0.001 to 0.01 weight %, and at least one selected from carbon (C) 0.15-0.50 weight %, silicon (Si) 0.15-0.3 weight % or less , Manganese (Mn) 0.5-1.50% by weight, phosphorus 0.04% by weight or less, sulfur 0.05% by weight or less, boron (B) 0.001 to 0.004% by weight, nitrogen (N) 0.004 to 0.015% by weight, oxygen 0.004% by weight or less and forming a billet containing the remainder iron (Fe) and other unavoidable impurities;
(b) heating the billet;
(c) forming a rolled steel material by hot rolling the heated billet;
(d) cooling the rolled steel;
including,
The step (c) includes a process of rough-rolling the heated billet, a process of precision rolling, and a finish-rolling process,
The austenite average grain size immediately before the finish rolling process is 15㎛ or less, and the temperature range of the finish rolling process is the two-phase region of ferrite and austenite (Ae1+Ae3)/2 The cumulative rolling deformation amount in the range of ±10℃ It ranges from 0.6 to 2.0,
The microstructure immediately after the finish rolling includes a ferrite average grain size of 5 µm or less and an austenite grain size of 15 µm or less. delete 5. The method according to claim 4,
The step (d) is a method of manufacturing a steel for cold rolling, characterized in that for cooling the rolled steel to 450 ~ 550 ℃ at a cooling rate of 5 ~ 20 ℃ / s. 7. The method of claim 6,
(e) further comprising a drawing step of drawing by passing the cooled steel material through a die,
The step (e) is a primary drawing step in which the steel is drawn by passing it through a first die, a common rolling step in which the steel that has passed through the first die is passed through one or more common rolling rolls and rolled, and the steel passed through the common rolling rolls. A second drawing step of drawing by passing through a second die,
The steel material is drawn at a reduction in section in the range of 10 to 40% in the first drawing step, rolled at a reduction in section in the range of 10 to 40% in the formal rolling step, and in the range of 10 to 40% in the second drawing step A method of manufacturing a steel for cold rolling, characterized in that the drawing is performed at a reduced rate, and finally processing is performed such that the cumulative cross-sectional reduction rate of the steel material is 50% or more. 8. The method of claim 7,
The steel material having completed the secondary drawing step has a dislocation density of 4.0x10 ⁸ mm ^-2 or more. 8. The method of claim 7,
(f) further comprising the step of atmospheric heating or induction heating to achieve austenite heat treatment on the drawn steel,
In the austenite heat treatment, the average size of the austenite grain size is 7 μm or less by atmospheric heating or induction heating in the temperature range of Ac3+ (80±50)° C., which is directly above the austenite transformation point (Ac3). to be made, including those having an undissolved pearlite fraction of 5% or less,
After the austenite heat treatment, the method of manufacturing a steel for cold forging, characterized in that the fraction of grain boundaries having a grain boundary angle of 15° or less is 20% or more. 8. The method of claim 7,
The steel material is quenched at a cooling rate of 20~150℃/sec until it reaches room temperature after atmospheric heating or induction heating in the temperature range of Ac3+(80±50)℃, which is directly above the austenite transformation point (Ac3) for the cold drawn material. A method for producing a steel material for cold rolling, characterized in that obtained by tempering. 11. The method of claim 10,
The steel has a hardness of Hv320 to 385 at tempering 550 ° C, a tensile strength of 1015 to 1200 MPa, a yield strength of 980 to 1180 MPa, a yield ratio of 0.97 to 0.99, a work hardening and work softening balance elongation of 6 to 8%, 12 A method of manufacturing a steel for cold forging, characterized in that it satisfies the ranges of uniform elongation of ~15%%, total elongation of 19~22%, and tensile area reduction rate of 75~80%.

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