KR101115718B1 - High strength steel having excellent delayed fracture resistance and elongation and method for producing the same - Google Patents

Abstract

The present invention relates to a method for producing a steel workpiece having a strength suitable for use in a high strength component and at the same time high resistance to delayed fracture and a low yield ratio, and a steel workpiece obtained through the same. In order to solve this problem, the present invention provides a manufacturing method of the present invention for achieving the above object by weight, C: 0.30 to 0.70%, Si: 2.0 to 4.0%, Mn: 0.1 to 1.0%, P: 0.01% or less , S: 0.01% or less, Se: 0.001 ~ 0.05%, Ti: 0.001 ~ 0.03%, B: 0.001 ~ 0.003%, Al: 0.002 ~ 0.01%, N: 0.004 ~ 0.008%, O: 0.005% or less Pre-rolled billet having a composition consisting of a balance Fe and other impurities, wherein Ti, N, B and Al satisfy specific relation 1, and Mn, Se and S satisfy specific relation 2. Step, winding the rolled wire rod, cooling the wound wire rod, graphitizing heat treatment of the cooled wire rod, cold forging and / or cutting the graphitized heat treatment wire rod to produce a steel workpiece a step, wherein the produced steel workpieces to _{Ac 3 - (Ac 3 -Ac 1} ) /1.3 ~ Ac 3 - (Ac 3 -Ac 1) heating to a temperature range of /5.5, wherein said heated steel workpiece Maintaining a predetermined time in the temperature range, cooling the steelwork to a temperature range of Ms + 50 ~ Ms + 110 ℃ at a cooling rate of 70 ℃ / sec or more, and isothermal heat treatment in the temperature range.

Graphitization, cold forging, machinability, graphite grain, delayed fracture resistance

Description

High-strength steel products with excellent delayed fracture resistance and elongation and manufacturing method {HIGH STRENGTH STEEL HAVING EXCELLENT DELAYED FRACTURE RESISTANCE AND ELONGATION AND METHOD FOR PRODUCING THE SAME}

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to graphite steel, a mechanical structural steel suitable for use in machine parts such as industrial machines or automobiles, and having excellent cold forging and high machinability, and at the same time having strength suitable for use in high strength parts. The present invention relates to a method for manufacturing a steel workpiece having high resistance to delayed fracture and low yield ratio, and a steel workpiece obtained through high strength parts.

Mechanical parts used in industrial machines and automobiles are usually machined into complex shapes of the parts by either cutting or cold forging. However, in the case of the cutting process, in order to process the final shape of the raw material, there is a problem that the amount of processing increases, resulting in large parts loss and excessive production cost. In addition, the cold forging process has the advantage of relatively low component loss compared to the cutting process and can be carried out through a simple process. However, if the final shape of the component is complicated, the forging operation alone cannot realize the complicated shape. . Therefore, it is most suitable in terms of material saving, production cost reduction, etc., after machining into a near net shape by cold forging and final cutting into a complicated shape. In order to perform the cutting process after such cold forging, both machinability and cold forging of the material must be combined, and such steel cannot satisfy all of these properties.

That is, in order to improve the machinability of the material, a free-cutting steel containing machinable elements such as Pb, S, Bi, etc. can be used in steel. Such free-cutting steel can be used for cutting steel such as surface roughness, chip treatment, and tool life. In terms of machinability, it is very good. However, when cold forging is performed using this, cold forging is very poor due to cracks caused by micro-deformation due to cracks caused by inclusions of machinability enhancing elements.

In addition, Pb, which has been conventionally used as a machinability granting element, emits harmful substances such as toxic fume during cutting, and thus is very harmful to the human body. Although S, Bi, Te, Sn, and the like have been proposed, Bi-added steels are easily cracked during manufacturing, which is very difficult to produce, and S, Te, Sn, etc. cause cracks during hot rolling. There is a problem.

Therefore, the free-cutting steel using the conventional machinability imparting element not only has very poor cold forging, but also has a problem of causing environmental problems and low production efficiency, making it difficult to use as a steel material having both cold forging and machinability. It was.

In addition, cold forged steel is excellent in toughness and ductility, so it is advantageous to process it by forging up to a degree similar to the final shape in the cold forging process, and to be similar to the final shape. ) There is a problem in that the machinability such as the treatment property and the surface roughness is very poor, making it difficult to use.

In order to solve the above problems, the proposed steel is graphite steel. Graphite steel is a carbon steel containing fine graphite grains inside, and has excellent impact toughness and ductility, and is excellent in cold forging, and also has good machinability as the internal fine graphite grains become a crack source during cutting. It is a river.

Japanese Unexamined Patent Publication No. 50-096416 discloses C (total): 0.45 to 0.90, Si: 1.0 to 2.5%, Mn: 0.1 to 0.7%, S: 0.015% or less, and one or two of Al and Ti 0.015 to Graphite free cutting steel containing 0.15%, containing 0.45 to 0.90% of graphite present in a distribution of 50 pieces / mm ² or more, and consisting of the balance Fe and impurities has been proposed. The invention has been proposed in terms of free machinability rather than cold forging, but can be referred to as an invention belonging to the invention of the invention for graphite-based free cutting steel. However, the composition of the present invention discloses that the silicon concentration should not exceed 2.3% and the minimum graphitization time accordingly requires about 10 hours. However, when demand is considered in the process, it is a big burden on the process to perform graphitization heat treatment for about 10 hours in order to give free machinability. It is virtually impossible to implement.

In addition, Japanese Patent Laid-Open No. 6-212351 discloses a method for producing graphite free-cutting steel that is graphitized by heat treatment in a relatively short time in order to solve the problem that the graphitization treatment, which is a problem shown in Japanese Patent Application Laid-Open No. 50-096416, takes a long time. Doing. The technique disclosed in the above publication is C: 0.3-1.0%, Si: 0.4-1.0%, Mn to shorten the graphitization treatment time; 0.3 ~ 1.0, P: 0.02% or less, S: 0.015 ~ 0.035, B: 0.001 ~ 0.004%, N: 0.002 ~ 0.008%, and hot-rolled steel bars immediately after hot rolling composed of the balance Fe and unavoidable impurities. The water cooling device installed before and after the line cools the cooling start temperature below Ac1 point, the cooling end temperature below Ms point, and the average cooling rate from 30 ° C / s to 100 ° C / s. After that, the graphitization treatment is then performed at a heating temperature of 600 to 700 ° C. This technique uses the grating distortion, a characteristic of martensitic phase, to facilitate the precipitation of carbon into graphite, which is a technical gist of at least 9-12. It was a technique that could reach 100% graphitization rate within time.

However, even with the above technique, the time required for graphitization is at least 9 hours, so that it is difficult for the demand to be applied to the process, and there is still a need for technical advancement required to shorten the graphitization time.

As a technique for further shortening the graphitization time, Japanese Patent Application Laid-Open No. 2000-063948 has a weight%, more than C: 1.00 to 1.50%, Si: 1.00 to 2.80%, Mn: 0.01 to 2.00%, and P: 0.050% or less. Cast iron or slab containing 0.10% or less, O: 0.0050% or less, N: 0.020% or less, consisting of residual iron (Fe) and unavoidable impurities, and having a graphitization index CE of 1.30 or more, 850 to 1150 ° C. After heating to a temperature in the range of, hot-rolled, and then cooled to room temperature, the hot-rolled steel thus obtained is heated to a temperature within the range of 600 to 1000 ° C. for 3 hours or less, followed by air cooling, and then averaged in the steel. the particle size or more 1.0㎛ graphite to precipitate 100 / mm ² or more, and free cutting steel to the metal structure characterized in that the tissue and the Brinell hardness consisting of or composed of more than 20% of ferrite with the balance pearlite or ferrite only to 200 bongseon There is a description of the production method is described. If the above technique is used, the time required for graphitization can be shortened to at least 30 minutes, and thus it can be referred to as a technique of drastically improving the graphitization time compared to the prior art. However, even in the case of the above technique, it is still difficult to apply to the actual processing process because the graphitization treatment time, which is at least about 30 minutes and most of the time, is required.

In addition, even in the case of manufacturing graphite steel, it is necessary to give the strength to meet the requirements of the steel workpiece after manufacturing the steel workpiece through cold forging and cutting operations using such graphite steel. In particular, parts made of graphite steel that have undergone the cold forging and cutting operations are often used for parts requiring high strength, such as automobile parts. The strength must be increased again through heat treatment. However, when the strength is increased, there is a problem in that the delayed fracture resistance is deteriorated. Therefore, it is impossible to use more than 130kg / mm ² grade with the currently disclosed technology alone.

When the material is heat-treated by the conventional heat treatment method, most of the microstructure is a quasi single phase structure of tempered martensite, and Fe-based precipitates are distributed at grain boundaries and lath martensite. It is a general feature that the precipitates are distributed in the base metal of). However, since grain boundary precipitates act as a trap site of hydrogen and degrade the strength of grain boundaries, it is known that delayed fracture resistance is lowered, and thus, high strength of the material cannot be achieved. Therefore, it is accepted that the microstructure of tempered martensite has a limit in achieving high strength.

Therefore, in order to achieve high strength without deteriorating the delayed fracture resistance, it is most important to suppress the distribution of Fe-based precipitates distributed at the grain boundaries after the heat treatment to the maximum.

In addition, when the yield ratio of the processed steel products is lowered, it can be effectively used to reduce the noise of automobiles. Thus, when manufacturing steel products having both delayed fracture resistance, high strength and resistance to yield ratio, it is highly useful to be used for various purposes. There is no known technique for manufacturing a part wool that satisfies all of these properties among parts processed into a complex shape by performing forging and / or cutting operations.

The present invention is to solve the above problems, and has a similar high machinability through a short graphitization time without adding low melting point elements such as Pb, Bi, S, Sn added to give a high machinability At the same time, an object of the present invention is to provide a high-strength steel product excellent in delayed fracture resistance manufactured from graphite steel capable of securing excellent cold workability (or cold forging) and a method of manufacturing the same.

Production method of the present invention for achieving the above object, in weight%, C: 0.30-0.70%, Si: 2.0-4.0%, Mn: 0.1-1.0%, P: 0.01% or less, S: 0.01% or less, Se: 0.001-0.05%, Ti: 0.001-0.03%, B: 0.001-0.003%, Al: 0.002-0.01%, N: 0.004-0.008%, O: 0.005% or less (excluding 0) Wire-rolling a billet having a composition consisting of Fe and other impurities, wherein Ti, N, B and Al satisfy the following relation 1, and Mn, Se and S satisfy the following relation: Winding, cooling the wound wire, graphitizing heat treatment of the cooled wire, cold forging and / or cutting the graphitized heat treatment to produce a steel product, the manufactured steel the workpiece _{Ac 3 - (Ac 3 -Ac 1} ) /1.3 to _{Ac 3 - (Ac 3 -Ac 1} ) heating to a temperature range of /5.5, maintaining a certain amount of time for the heated steel workpiece at the temperature range Step, and a step and a step of isothermal heat treatment at this temperature range for cooling the steel work piece to 70 ℃ / sec or more cooling rate in a temperature range of Ms + 50 ~ Ms + 110 ℃.

(However, Ti, N, B, Al, Mn, Se and S in the following relations 1 and 2, respectively, means the weight percent of the element.)

[Relationship 1]

2.0≤ (Ti + 5B + Al) /N≤5.5

[Relationship 2]

1.0≤ (Mn / 5 + Se) /5S≤3.0

More preferably, in addition to the above-mentioned stress, Ni: 0.05 to 1.0%, Cu: 0.01 to 0.5%, Ca: 0.0001 to 0.05%, Zr: 0.0005 to 0.008%, REM (Rare Earth Metal, Rare Earth Metal): 0.001 to 0.05 It is preferable to include one or two or more selected from the group consisting of%.

At this time, it is preferable that the austenite grain size immediately after finish rolling in the step of wire-rolling the billet of the composition is 30 μm or less.

And it is preferable to perform the said wire rod finish rolling in the temperature range of 840-900 degreeC.                     

And, the fraction of the cornerstone ferrite in the wire inner structure after the cooling step is preferably 10% or less.

In addition, the step of winding the rolled wire is preferably made in a temperature range of 750 ~ 800 ℃.

In addition, in the step of cooling the wound wire, it is preferable to cool the wound wire to a temperature range of 650 ~ 690 ℃ at a cooling rate of 2.7 ~ 3.3 ℃ / sec.

In addition, the step of graphitizing the cooled wire rod is preferably carried out in a temperature range of Ac _1- (60 ± 30) ℃.

In addition, it is preferable that the average size of the graphite grains contained in the graphite steel manufactured from the step of graphitizing the wire rod is 15 μm or less.

In addition, the _{Ac 3 - (Ac 3 -Ac 1} ) /1.3 to _{Ac 3 - (Ac 3 -Ac 1} ) the holding time of the step of holding the steel workpieces were heated to a temperature range of /5.5 is preferably not less than 20 minutes Do.

 Here, the heat treatment holding time of the isothermal heat treatment step is preferably 20 minutes or more.

Another method of the present invention for achieving the above object, in weight%, C: 0.30-0.70%, Si: 2.0-4.0%, Mn: 0.1-1.0%, P: 0.01% or less, S: 0.01% or less , Se: 0.001-0.05%, Ti: 0.001-0.03%, B: 0.001-0.003%, Al: 0.002-0.01%, N: 0.004-0.008%, O: 0.005% or less (excluding 0), It is composed of a balance Fe and other impurities, the Ti, N, B and Al satisfy the following relation 1, Mn, Se and S has a composition satisfying the following relation 2 and has a graphite grain having an average size of 15㎛ or less inside the method comprising providing a graphite steel comprises, the step of preparing the graphite steel in cold-forged and / or machined workpieces to steel, the manufacture of steel workpieces _{Ac 3 - (Ac 3 -Ac 1} ) /1.3 to Ac ₃ - Heating to a temperature range of (Ac ₃ -Ac ₁ ) /5.5, maintaining the heated steel product in the temperature range for a predetermined time, and cooling the steel product to a cooling rate of 70 ° C./sec or more, Ms + 50 to Ms. Temperature of + 110 ℃ Step of cooling to above and to include the step of isothermal heat treatment at this temperature range.

(However, Ti, N, B, Al, Mn, Se and S in the following relations 1 and 2, respectively, means the weight percent of the element.)

[Relationship 1]

2.0≤ (Ti + 5B + Al) /N≤5.5

[Relationship 2]

1.0≤ (Mn / 5 + Se) /5S≤3.0

At this time, in addition to the above-mentioned emphasis Ni: 0.05 ~ 1.0%, Cu: 0.01 ~ 0.5%, Ca: 0.0001 ~ 0.05%, Zr: 0.0005 ~ 0.008%, REM (Rare Earth Metal, Rare Earth Metal): 0.001 ~ 0.05% It is preferable to contain 1 type (s) or 2 or more types selected from the group.

In addition, the _{Ac 3 - (Ac 3 -Ac 1} ) /1.3 to _{Ac 3 - (Ac 3 -Ac 1} ) the holding time of the step of holding the steel workpieces were heated to a temperature range of /5.5 is preferably not less than 20 minutes Do.

And, the heat treatment holding time of the isothermal heat treatment step is preferably 20 minutes or more.

The steel workpiece of the present invention for achieving the above object is a high-strength steel workpiece excellent in delayed fracture resistance having a composite structure composed of ferrite and residual austenite by the above production method.

And, it is preferable that the fraction of retained austenite in the structure of the steelwork is 15% or more.

In addition, it is preferable that the ratio of strip like type residual austenite in the residual austenite is 70 to 90%.

Hereinafter, the present invention will be described in detail.

In the present invention, the term 'steel processed' refers to a workpiece obtained by cutting and cold working the wire rod in a predetermined form, and includes all steel products which are applied to a use capable of utilizing physical properties of the steel of the present invention. In addition, the term “graphite steel” referred to in the present invention refers to a steel in which fine graphite grains are distributed therein and the steel processed into steel by heat treatment.

The inventors of the present invention have repeatedly studied at various angles to solve the above problems and to provide a high-strength steel product and a method for manufacturing the same, which have excellent delayed fracture resistance desired in the present invention. In order to fundamentally suppress the precipitation of grain boundary precipitates (Fe-based) which are harmful to resistance, the critical delay control strength is maintained at 150kg / mm ² or more, and the microstructure control method has been studied in various angles to develop steels with excellent mechanical properties. Result The steel workpiece is heat-treated in the ideal zone to prepare an appropriate amount of the ideal zone ferrite and the rest of the austenite composite, and then quenched to the temperature range of Ms + 50 ° C to Ms + 110 ° C and isothermally heat-treated. And morphology of residual austenite have lamellar distribution similar to that of pearlite. It is possible to manufacture complex tissue forms (collectively called strip-like type bainite) and complex tissue forms (collectively referred to as blocky type bainite) in which residual austenite is dispersed in ferrite. In this case, while achieving a high strength of 150kg / mm ² or more critical delay fracture strength, it was confirmed that the excellent mechanical properties, in particular, the cross-sectional reduction rate and impact toughness can be significantly improved and completed the present invention.

Hereinafter, the conditions and manufacturing method of the steel used in the present invention from the above point of view will be described in detail.

[Lecture composition]

The reason for limiting the alloying elements in the present invention is as follows.

Carbon (C): 0.3-0.7 wt%

Carbon is an essential element for the formation of graphite, and is an important element for securing the strength of mechanical parts, but its content is less than 0.3% by weight, and its strength during heat treatment for manufacturing ferritic and residual austenite composite steel This is because it is difficult to control the ferrite structure fraction, and the effect is not only saturated at 0.7 wt% or more, but also the properties such as proper cross-sectional reduction rate, elongation rate and impact toughness are lowered after heat treatment, and segregation and surface defects are generated during wire fabrication. It is limited to 0.3 ~ 0.7% by weight because it affects deep decarburization at the time of charging furnace, permanent deformation of material after heat treatment, static fatigue characteristics and shape and size of microcomposite.

Silicon (Si): 2.0-4.0 wt%

Silicon is a necessary component as a deoxidizer in molten steel and is an element that makes carbon carbide precipitate by making iron carbide (cementite) in steel unstable. In addition, silicone is actively added because it is a component that improves strength. Less than 2.0% by weight, the effect is insufficient, and the mechanical and thermal stability of the retained austenite after ferrite transformation decreases, making it difficult to secure the cornerstone ferrite + residual austenite complex structure and to obtain the appropriate amount of retained austenite. It is difficult to secure the strength due to the insufficient effect. Even when adding more than 4.0% by weight of silicon, the effect of promoting graphitization is saturated and the temperature range where liquid phase is generated becomes narrow, so that the appropriate temperature range is narrowed during hot rolling. In order to deteriorate the microstructure, impact toughness, fatigue characteristics, etc., it was limited to the range of 2.0 to 4.0% by weight.

Manganese (Mn): 0.1-1.0 wt%

Manganese is an effective element to secure the strength of steel and is also useful as a deoxidizer in molten steel manufacturing. In addition, it combines with S to form MnS to contribute to the improvement of machinability. However, when the content is less than 0.1% by weight, the effect of improving strength is small, and when the content is more than 1.0% by weight, the toughness is deteriorated. Therefore, the content is limited to the range of 0.1 to 1.0% by weight.

Phosphorus (P): 0.01 wt% or less

Phosphorus not only inhibits graphitization, but also segregates at the austenite grain boundary during the quenching treatment, thereby lowering the grain boundary strength, so that quenching cracks are easily generated.

Sulfur (S): 0.01 wt% or less

Sulfur forms MnS to improve chip breaking during cutting to improve machinability, especially as nucleation of graphitization to promote graphitization, but if the amount is more than 0.01% by weight, the effect is not only saturated. On the other hand, the upper limit was limited to 0.01% by weight because it delayed the graphitization and drastically lowered the toughness of the steel, thus adversely affecting the cold forging.

Selenium (Se) : 0.001 ~ 0.05 wt%

Selenium improves chip breaking by forming MnSe in conjunction with manganese. At the same time, MnSe is an element that improves machinability by acting as a graphitization nucleus to promote graphitization. This effect was limited to 0.001% by weight to 0.05% by weight or less, and 0.05% by weight or more because the effect is saturated.

Titanium (Ti): 0.001-0.03 wt%

Titanium combines with nitrogen in steel to form TiN, making cementite unstable and at the same time becoming a nucleation site of graphite to promote graphitization. Moreover, since it functions effectively also as a deoxidizer, it adds actively. To attain the above useful effects, at least 0.001% by weight should be added. However, at 0.03% or more by weight, the graphitization is rather hindered, so it is limited to 0.001 to 0.03% by weight.

Boron (B): 0.001 to 0.003% by weight

Boron is an element that is actively added because it combines with N to form BN and acts as a crystal nucleus of graphite while interfering with stabilization of cementite, thereby promoting graphitization and enhancing quenchability. If it is less than 0.001% by weight, the effect of addition is insufficient, so it is necessary to add more than 0.001% by weight. On the contrary, when it is added more than 0.003% by weight, the effect can not be increased any more. The addition range was limited to 0.001 to 0.003% by weight because it lowered the hot workability.

Aluminum (Al) : 0.002 ~ 0.03% by weight

Aluminum is a powerful deoxidation element and a useful element that not only contributes to deoxidation but also promotes graphitization. In the graphitization heat treatment, the decomposition of cementite is promoted, and at the same time, it binds with nitrogen to form AlN, thereby preventing the cementite from stabilizing. In addition, aluminum oxide produced in the steel by the addition of aluminum is effective in that it also becomes a precipitation nucleus of BN and promotes crystallization of graphite. In the present invention, aluminum is actively added, but if the content is less than 0.002% by weight, it is difficult to expect the effect of addition, and at 0.03% or more by weight, the graphitization promoting action is saturated and the hot deformation is significantly lowered. Limit to the range of.

Nitrogen (N): 0.004-0.008 wt%

Nitrogen is actively added because it combines with titanium, boron and aluminum to form nitrides and use these as nuclei to promote the crystallization of graphite. On the other hand, in order to form nitrides effective for promoting graphitization, it is preferable to add in stoichiometric amounts almost equivalent to those of titanium, boron, and aluminum, but in order to uniformly finely disperse these nitrides, it is preferable to add slightly higher than chemical equivalents. . In addition, it is advantageous to add a little excessively because nitrogen improves chip treatability by dynamic strain aging. For this reason, it is necessary to add more than 0.004% by weight, but when added more than 0.008% by weight because the effect is saturated, it was limited to 0.004 ~ 0.008% by weight.

Oxygen (O): 0.005% by weight or less

In the present invention, the role of oxygen is important. Oxygen combines with aluminum to form oxides. The production of such oxides reduces the effective concentration of aluminum. As a result, the amount of AlN produced useful for the crystallization of graphite is reduced, thus causing the effect of substantially obstructing the graphitization action. In addition, alumina oxide formed by containing a large amount of oxygen damages the cutting tool during cutting, resulting in a decrease in machinability. For this reason it is desirable to manage as low as possible. However, if the oxygen is managed too low, it causes a refining load of the steelmaking process, and the upper limit is limited to 0.005% by weight or less because the problem caused by the oxygen is not so much up to 0.005% by weight.

As can be seen from above, in order to promote graphitization, first, a carbon source necessary for graphitization is sufficiently maintained in the steel, and the cementite is destabilized so that carbon in the cementite can be easily diffused into the steel. Oxygen must be able to grow in graphite on various types of nitride or compound inclusions, which are non-homogeneous nucleation sites in the river.

However, the composition alone may not provide a sufficient nucleation site for graphitization, and the nitride- or emulsion-based inclusions generated by the composition may be effectively finely dispersed in a large amount in the steel to supply a sufficient nucleation site. Accordingly, the graphitization time can be significantly shortened and fine graphite can be produced.

In the present invention, in order to significantly shorten the graphitization and graphitization time, the ratio of the composition ratio between alloy elements is 2.0≤ (Ti + 5B + 2Al) /N≤5.5, 1.0≤ (Mn / 5 + Se) /5S≤3.0 The reason for this is as follows.

It is preferable that (Ti + 5B + Al) / N ratio is 2.0 <= (Ti + 5B + Al) / N <= 5.5 .

If the ratio of (Ti + 5B + Al) / N is less than 2.0, the number of TiN and BN precipitates that contribute to the nucleation of graphite grains due to the lack of Ti, B, Al is insufficient, and the amount of (Ti + 5B + Al) / N When the ratio exceeds 5.5, Ti, B, and Al are sufficient, but the amount of nitrogen is insufficient, so the number of precipitates of TiN, BN, and AlN necessary for graphite nucleation is not increased anymore, but due to excess nitrogen, The amount of nitrogen that is added increases and adversely affects the graphitization rate.

The value of (Mn / 5 + Se) / 5S is preferably 1.0 ≦ (Mn / 5 + Se) /5S≦3.0 .

If the value of (Mn / 5 + Se) / 5S is less than 1.0, the number of MnSe precipitates contributing to nucleation of graphite grains and the effective MnS inclusions for machinability are insufficient, and when (Mn / 5 + Se) / 5S ratio exceeds 3.0, Not only the number of MnSe precipitates and MnS precipitates required for graphite grain nucleation is saturated, but the amount of sulfur dissolved in the base material increases, causing grain boundary segregation, which adversely affects mechanical properties.

One or more of the group consisting of 0.05 to 1.0% by weight of nickel, 0.01 to 0.05% by weight of copper, 0.0001 to 0.05% by weight of calcium, 0.0005 to 0.008% by weight of zirconium, and 0.001 to 0.05% by weight of rare earth metal. Two or more kinds are selectively added, and the reason for component limitation is explained.

Nickel (Ni): 0.05-1.0 wt%

Nickel is an element that forms a nickel thickening layer on the surface during heat treatment to suppress permeation of external hydrogen to improve delayed fracture resistance. If the nickel content is less than 0.05%, it is difficult to expect the effect of improving the delayed fracture resistance due to the incomplete formation of the surface thickening layer, and the heat treatment time is increased during the graphitization treatment to improve the decarburization control, toughness and cold formability, and the nickel content When the amount exceeds 1.0%, the effect is saturated and limited to 0.05 to 1.0% because it affects the proper amount, size and shape of the amount of retained austenite.

Copper (Cu): 0.01-0.5 wt%

Since copper is effective for promoting graphitization by making cementite unstable, not only improves machinability but also has an effect of increasing strength at the time of hardening of steel by the effect of enhancing the hardening of the steel and the strengthening of precipitation. In addition, when copper is added, the corrosion resistance of the steel can be improved. If it is less than 0.01% by weight, the effect of promoting graphitization and improvement of corrosion resistance is insufficient. If it exceeds 0.5% by weight, the improvement effect is saturated and the melting point is lowered at the grain boundary segregation. It was limited to 0.01 to 0.5% by weight because of the high probability of occurrence of surface defects due to grain embrittlement during charging and the impact toughness in the final product.

Calcium (Ca) : 0.0001 ~ 0.05 wt%

Calcium improves machinability by forming Ca-Al-based oxides in the steel composition of the present invention, which acts as a nucleus of graphitization to promote graphitization. This effect is less than 0.0001% by weight, and in 0.05% by weight or more coarse oxide-based non-metallic inclusions are generated in a large amount to reduce the fatigue strength of the mechanical parts were limited to the range of 0.0001 ~ 0.05% by weight.

Zirconium (Zr) : 0.0005 ~ 0.008% by weight

Zirconium finely disperses oxides such as CaO and Ti2O3 and MnS sulfides. These oxides and sulfides serve as precipitation sites of the graphite and are effective in shortening the graphitization time to finely disperse the graphite grains. However, if the amount of zirconium added is less than 0.0005% by weight, the effect is insufficient. If the amount of zirconium is more than 0.008% by weight, coarse Zr-based sulfides and carbonitrides are formed to reduce the miniaturization effect of oxides by Zr and to deteriorate fracture toughness. It was limited to the range of 0.008% by weight.

Rare Earth Metal (REM) : 0.001 ~ 0.05% by weight

Rare earth metals are added for the purpose of improving the hot workability of steel and further promoting graphitization. This action is useful to use La, Ce, etc., but the content is less than 0.001% by weight because the effect is insufficient, and at 0.05% or more because the effect is saturated, it was limited to the range of 0.001 to 0.05% by weight. .

[Manufacture process]

In the present invention, a wire rod is manufactured and graphitized, and then a steel workpiece is made through at least one step of cold forming and cutting, and the steel is heat-treated to produce a steel workpiece having excellent delayed fracture resistance. Hereinafter, the features of the present invention for each manufacturing step will be described.

[Manufacturing method of wire rod]

In the present invention, the billet obtained by billet rolling of the bloom or obtained by casting directly to the billet is obtained by rolling and cooling the billet, then the wire is fresh to the appropriate wire diameter. Graphitization heat treatment is performed before and after the fresh working.                     

In the present invention, a wire rod having the above-described steel component system and having a grain size of spherical austenite of 30 µm or less is used. If the spherical austenite grain size of the wire rod exceeds 30㎛, the size and thickness of cementite constituting the pearlite after the transformation of the pearlite becomes too large and thick, which causes a problem of lowering the graphitization rate. It is unpreferable because it becomes coarse and adversely affects cold forging property and machinability.

On the other hand, wire rod rolling for producing a wire rod having a grain size of spherical austenite of 30 μm or less is performed at a temperature range of 870 ± 30 ° C. This is because the control rolling effect is insufficient at 900 ° C or higher, and thus it is difficult to obtain austenite grain refining effect. At 840 ° C or lower, the effect is saturated, surface flaws are easily generated, wire rod gap difference is difficult to control, and rolling load is increased. It is not preferable because workability is poor.

In addition, the present invention uses a wire rod having a granite ferrite structure fraction of 10% or less after cooling the wire rod. In the case where the cornerstone ferrite structure fraction exceeds 10% after wire rod cooling, it is not preferable because the distribution of graphite grains during graphitization becomes uneven and may cause coarsening of graphite grains and adversely affect cold forging and machinability. In the ferrite region, the carbon content is very low and the carbon necessary for the formation of graphite grains during graphitization is scarce, so that the ferrite is hardly graphitized, so it must be limited to 10% or less.

On the other hand, wire-controlled cooling for producing wire rods having a saltpeter ferrite structure fraction of 10% or less is wound up after rolling, and the winding temperature at this time is preferably 750-800 ° C. The coiling temperature is secured by cooling water spraying immediately after the wire is rolled. If the coiling temperature is less than 750 ℃, the possibility of winding failure due to cold winding is likely to occur. If the coiling temperature is higher than 800 ° C, it is difficult to control the cornerstone ferrite fraction.

The wound wire is cooled down to 3.0 ± 0.3 ° C / sec to 670 ± 20 ° C. The conditions of cooling temperature and cooling rate are taken into account the difference in the degree of cooling of wire rod integrated state, ie overlapping part and non-overlapping part. That is, when the wire rod coil is cooled in the cooling zone, the wire coil may be divided into the non-overlapping portion, the middle overlapping portion and the overlapping portion in order from the center to the outside. As the overlapping degree of the wire rod increases, the cooling rate decreases due to the mutual insulation effect. In the overlapping part where the overlapping degree is large, the cooling rate is large in the central part where the cooling rate is large and the degree of overlapping is small. This difference in cooling rate becomes more prominent as the total cooling rate is increased. The difference in the cooling rate as described above is not advantageous because it appears as a variation in the material, that is, the tensile strength of each coil. Therefore, in the present invention, while enjoying a sufficient cooling effect, a range between 3.0 ± 0.3 ° C / sec is set as the cooling rate range in which the material deviation of each part does not appear.

If the cooling temperature is higher than 690 ℃, the formation rate of the cornerstone ferrite increases, and if the cooling temperature is lower than 650 ℃, the perlite transformation rate decreases. In addition, when the cooling rate is faster than 3.3 ° C / sec is because the low-temperature structure (martensite or bainite) is produced, if it is slower than 2.7 ° C / sec is not preferable because the increase in the cornerstone ferrite phase fraction.

[Graphite process]

Thereafter, in the present invention, the wire rod having the spherical austenite grain size and the cornerstone ferrite fraction as described above is subjected to the graphitization heat treatment during the cooling process or before or after drawing. The graphitization heat treatment is performed at Ac _1- (60 ± 30 ° C). This temperature range is the fastest graphitization rate and shows the fastest graphitization behavior at Ac ₁ -60 ℃. When the graphitization heat treatment temperature is less than Ac ₁ -90 ℃, there is a problem in securing productivity due to the temperature range that the graphitization rate is very slow, and at a temperature above Ac ₁ -30 ℃, the graphitization heat treatment time is long and the austenite phase is precipitated. Rather, graphite particles are likely to be re-used.

Graphite grains present in the wire rod after the wire rod rolling process and the graphitization process as described above is 15㎛ or less. It is not preferable that the graphite grain size is larger than 15 mu m because it adversely affects cold forging and machinability.

[Processed Steel Products]

The graphite steel subjected to the graphitization treatment as described above is then processed into a steel workpiece of a desired shape through one or more steps during cold forging and cutting operations. The steel workpiece processing process as described above can be performed through ordinary cold forging and cutting operations due to the excellent cold forging and machinability of the graphite steel.

[Compound structure heat treatment process]

The steel workpiece that has undergone one or more steps during the cold forging and cutting operations has a desired final shape, but its strength is weak enough to be used as a high strength component. Therefore, it is necessary to secure the desired microstructure by heat treatment of the steel workpiece.

As described above, the inventors of the present invention make the final structure of the steel processed product into a composite structure composed of ferrite and residual austenite, and appropriately controlling the composite phase fraction in the obtained composite structure while ensuring the strength and resistance ratio and delayed fracture resistance. It is a major means to do it. That is, the internal structure of the steel workpiece manufactured by the present invention is composed of ferrite + residual austenite in which carbide is hardly present.

Among the ferrites contained in the internal structure, the cornerstone ferrite discontinuously disperses the austenite grain boundary, so that precipitates such as Mo, Ti, V, and Nb other than the Fe precipitates deposited at the grain boundaries cannot be continuously distributed. It is required to exist in a proper fraction as a structure necessary for reducing segregation of impurity elements such as P and S. When the cornerstone ferrite structure fraction is 5% or less, it is required to discontinue grain boundary discontinuity of austenite to improve delayed fracture resistance. Since the amount of salt-free ferrite is insufficient, the improvement effect is insufficient, and at 25% or more, there is a problem that the tensile strength is lowered due to excessive amount of the salt-free ferrite. Also, the carbon concentration of austenite is increased during the abnormal zone and then cooled. This is because it is difficult for an organization to become an ideal composite organization that is effective for resistance. More preferred saltpeter ferrite fraction is in the range of 10-15%. This is because it is the range that can improve the delayed fracture resistance due to the high strength, high elongation and uniform ferrite distribution at the same time to implement the resistance complex ratio in the final composite structure.                     

In addition, it is preferable to achieve the object of the present invention that the finally obtained composite structure composed of ferrite and residual austenite has a ratio of 70% to 90% of the strip like type residual austenite in the total residual austenite.

If the ratio of residual austenite of the strip-like type of the retained austenite is less than 70%, the yield strength is lowered, which affects the amount and size, shape, mechanical and thermal stability of the retained austenite, which is not preferable for high strength. When the retained austenite ratio of the strip-like type exceeds 90%, elongation, cross-sectional reduction rate and impact toughness are lowered, which is not preferable.

The heat treatment method for obtaining the above-mentioned preferable structure structure is as follows.

Looking at the heat treatment conditions for obtaining the tissue in more detail, first, the work piece over the river station temperature _{Ac 3 - (Ac 3 -Ac 1} ) Ac 3 in /1.3 - (Ac ₃ -Ac ₁₎ range from /5.5 It is heating inside. Here, Ac ₃ is the austenite transformation temperature when heated, Ac ₁ is an abnormal temperature (ferrite + austenite) transformation temperature when heating, Ac ₃ , Ac ₁ transformation temperature is different depending on the alloy component system.

If the heating temperature is lower than Ac _3- (Ac ₃ -Ac ₁ ) /1.3, the amount of ferrite produced during the abnormal reverse heat treatment exceeds 25%, which leads to a decrease in yield strength, elongation and impact toughness, as mentioned above. If the heating temperature exceeds Ac _3- (Ac ₃ -Ac ₁ ) /5.5, the amount of ferrite required for discontinuity of grain boundaries becomes less than 5%, and the effect is insufficient.

A more preferable heat treatment condition is over station _{Ac 3 - (Ac 3 -Ac 1} ) Ac 3 /1.7 from - inde (Ac ₃ -Ac1) /2.5 temperature range, which spherical austenite crystal grain boundary of the discontinuity, a discontinuity, the heating of the grain boundary precipitates It is a range that takes into account decarburization control and heat treatment time. In order to obtain a high-temperature composite structure of the target ferrite + austenite, heating must be performed for a sufficient time. If the heating is performed for about 20 minutes or more, the desired transformation is completed to obtain the desired tissue.

After the heating step as described above, it is necessary to quench to a temperature range of Ms + 50 ℃ ~ Ms + 110 ℃ at a cooling rate of 70 ℃ / sec or more (the Ms means the temperature at which martensite transformation starts). It is necessary to perform isothermal heat treatment for at least 20 minutes in the temperature range so that constant temperature transformation occurs after the quenching.

At this time, when the cooling rate is less than the 70 ℃ / sec, an abnormal phase ferrite (basestone ferrite) may be additionally precipitated during cooling, it is difficult to secure the strength of the steel product.

If the constant temperature of transformation is less than Ms + 50 ℃, the ratio of the strip-like residual austenite within the total residual austenite fraction when the carbide-free composite tissue is manufactured may be more than 90%, may cause problems, Ms + 110 ℃ It is not preferable to exceed the ratio of the Brachy type residual austenite in the total residual austenite exceeding 30% (ie, the strip like type is less than 70%). In addition, when the constant temperature transformation time is 20 minutes or less, it is because the time at which phase transformation occurs to the target microstructure in the present invention is insufficient. However, this process is a process that induces a constant temperature transformation, so the microstructure is no longer changed even if the heat treatment time is longer, so the treatment time upper limit is not separately defined.

The steel workpiece formed through the above process has a microstructure consisting of ferrite and residual austenite, and the proportion of retained austenite in strip-like type of austenite is 70 to 90%, and the remainder has a brachy type, a complex structure, and the like. The mechanical properties of can be significantly improved.

Hereinafter, the present invention will be described in more detail with reference to Examples.

EXAMPLE

Steels having the composition as shown in Table 1 were cast into billet-shaped ingots weighing 50 kg, and then homogenized and heat-treated at 1250 ° C. for 48 hours to hot roll to 13 mm in thickness to prepare wires. At this time, the finishing temperature was 840-900 ℃ and hot-rolled after hot rolling. The rolling ratio was 80% or more.

Steel grade C,
wt% Si,
wt% Mn,
wt% S,
wt% P,
wt% Ti,
wt% B,
ppm Al,
wt% Se,
wt% N,
ppm O,
ppm Ni,
wt% Cu,
wt% Ca,
wt% Zr,
ppm REM,
ppm foot

persons

River One 0.51 3.1 0.51 0.008 0.011 0.007 20 0.005 0.004 60 20 0.1 2 0.52 3.1 0.40 0.009 0.015 0.004 10 0.004 0.003 55 21 0.1 0.05 3 0.50 3.0 0.48 0.007 0.009 0.010 21 0.005 0.002 56 19 0.01 4 0.49 2.9 0.55 0.009 0.008 0.010 20 0.003 0.010 60 24 10 5 0.53 3.2 0.53 0.008 0.009 0.009 30 0.005 0.004 65 22 10 6 0.7 3.1 0.48 0.007 0.008 0.004 13 0.003 0.002 75 17 0.09 0.04 7 0.31 3.2 0.51 0.008 0.011 0.010 26 0.004 0.002 66 29 0.16 0.06 8 0.53 4.0 0.55 0.009 0.009 0.003 17 0.004 0.005 60 23 0.20 0.09 9 0.49 2.8 1.00 0.008 0.015 0.011 29 0.015 0.005 55 30 0.25 0.11 10 0.66 2.05 0.44 0.009 0.009 0.003 20 0.010 0.006 80 40 0.54 0.21 ratio
School
River One 0.51 3.0 0.51 0.008 0.011 0.001 10 0.003 0.004 65 19 0.13 2 0.52 3.1 0.48 0.009 0.008 0.010 20 0.005 0.002 40 23 0.12 0.04 3 0.50 3.0 0.22 0.007 0.130 0.004 10 0.004 0.001 55 17 0.02 4 0.49 2.9 0.71 0.009 0.009 0.003 17 0.004 0.021 60 21 20 5 0.53 3.2 0.19 0.008 0.129 0.001 10 0.003 0.001 65 20 10 6 0.8 3.1 0.70 0.007 0.009 0.010 20 0.005 0.019 40 28 0.21 0.10 7 0.31 3.2 0.69 0.008 0.009 0.001 10 0.003 0.020 65 15 0.51 0.23 8 0.53 4.0 0.20 0.009 0.131 0.010 20 0.005 0.001 40 34 0.24 9 0.49 1.5 0.52 0.008 0.015 0.011 15 0.015 30 25 10 0.66 2.5 0.44 0.009 0.009 0.044 21 0.03 100 32

The graphitization heat treatment of the inventive and comparative steels was air cooled after holding at 750 ° C. until graphitization was complete. In order to examine the effect of each wire structure on the size of graphite grain, an analysis of the cornerstone ferrite area fraction was carried out for the area of 300mm ² per specimen using an image analyzer for the pre-graphitized wire structure. Table 2 summarizes the results of the analysis of the cornerstone ferrite area fraction and the relational expressions 1 and 2. In addition, the austenite grain size of the inventive steel and the test steel was measured by the KS standard (KS D 0205), and the results are also shown in Table 2.

Steel grade Relationship 1 Relation 2 Austenite
Average grain size (㎛) Cornerstone Ferrite Fraction in Microstructure after Hot Rolling (%) (Ti + 5B + Al) / N (Mn / 5 + Se) / 5S Inventive Steel 1 4.5 1.9 25 5 Inventive Steel 2 3.1 1.1 15 4 Inventive Steel 3 5.5 2.2 20 5 Inventive Steel 4 4.3 3.0 30 4 Inventive Steel 5 5.4 2.4 25 4 Inventive Steel 6 2.1 2.5 15 0 Inventive Steel 7 4.1 1.9 30 10 Inventive Steel 8 5.2 2.6 25 9 Inventive Steel 9 3.3 2.7 15 3 Inventive Steel 10 4.1 2.1 20 0 Comparative Steel 1 1.4 1.9 40 35 Comparative Steel 2 6.5 2.5 55 25 Comparative Steel 3 3.1 0.6 45 32 Comparative Steel 4 5.2 3.6 40 21 Comparative Steel 5 1.4 0.6 50 26 Comparative Steel 6 6.5 3.6 45 23 Comparative Steel 7 1.4 3.6 50 30 Comparative Steel 8 6.5 0.6 40 24 Comparative Steel 9 11.2 - 45 20 Comparative Steel 10 8.4 - 40 22

In order to confirm the performance of the graphite steel manufactured through the graphitization heat treatment from the viewpoint of machinability and cold forging, the following experiment was performed.

Cold forging was performed by using a test piece manufactured with a diameter of 19 mm x 25 mm in height, and the critical volume change rate until cracking occurred at room temperature compression was used as an evaluation criterion for the critical cold forging rate. At this time, the critical cold forging rate was performed 10 times, and the remaining value except the maximum minimum value was evaluated as the average value.

Automatic machinability was used for the machinability test, and the machinability of the graphite steel was determined by the chip treatment and the tool life. Chip chipping was used as the chip processing method. Chips were divided into 2 or less volumes, and chips were classified into 3 or 6 volumes. In the test of tool life, the tool life was measured by measuring the time from the tip of the die tip to the point where it became impossible to cut when cutting in the environment using cutting oil at cutting speed 150m / min and 0.20mm / rev.

The graphitization time, graphite steel internal structure, cold forging property and free machinability test results are shown in Table 3. The internal structure of the graphite steel was analyzed for an area of 300 mm ² per specimen using an image analyzer in the same manner as the cornerstone ferrite fraction analysis.

division Graphitization properties Cold Forging Machinability Graphitization Time
(min) Graphite fraction
(%) Graphite Grain Average Size
(Μm) Critical cold forging rate
(%) Chip Processing Tool life
(min) Inventive Steel 1 5 2.1 7 130 Great 120 Inventive Steel 2 6 2.2 8 140 Great 110 Inventive Steel 3 5 2.0 9 130 Great 130 Inventive Steel 4 7 1.9 7 135 Great 115 Inventive Steel 5 7 2.1 6 140 Great 140 Inventive Steel 6 6 2.5 8 130 Great 120 Inventive Steel 7 7 1.2 7 140 Great 130 Inventive Steel 8 5 2.1 7 135 Great 125 Inventive Steel 9 6 2.0 6 135 Great 140 Inventive Steel 10 5 2.7 9 140 Great 135 Comparative Steel 1 350 2.0 25 70 usually 40 Comparative Steel 2 400 2.0 30 75 usually 35 Comparative Steel 3 10 2.1 20 80 usually 50 Comparative Steel 4 9 2.2 17 90 usually 60 Comparative Steel 5 600 2.1 21 85 usually 40 Comparative Steel 6 440 2.9 33 70 usually 30 Comparative Steel 7 530 1.3 20 80 usually 15 Comparative Steel 8 380 2.0 19 80 usually 45 Comparative Steel 9 3600 2.2 24 75 usually 20 Comparative Steel 10 3900 2.6 35 65 usually 25

As shown in Table 3, when the emphasis of the present invention is used, the graphitization time is within 7 minutes, and 100% graphitization treatment is possible within the graphitization time range targeted by the present invention, while the comparative examples are performed at a minimum of 9 minutes. The deviation was up to 3900 minutes. Thus, it was possible to confirm the superiority of the graphitization performance of the steel according to the present invention.

In addition, it can be seen that the size of the graphite grains of the graphite steel graphitized by using the steel of the present invention has a graphite finer than that of the comparative example of 17 to 35 μm at about 7 to 9 μm. While the cold critical forging is shown in the range 65 ~ 90% it can be seen that the examples of the present invention is quite excellent in the range 130 ~ 140%. In addition, as a result of determining the machinability, in the case of graphite steel according to the present invention, the chip treatment property is all divided into two or less books, whereas in the comparative example, graphite chips manufactured from the steel for graphite steel according to the present invention are divided into three or more books. It was confirmed that this has a much better chip treatment, the tool life was also found that the tool life of the graphite steel made of the steel for graphite steel according to the present invention improved at least 55 minutes over the comparative example.

As an invention example of the present invention, heat treatment was performed on the inventive steel of the present invention under the conditions shown in Table 4. X in Table 4 means a value satisfying the formula 'Ac _3- (Ac ₃ -Ac ₁ ) / X = heating temperature' described in the table.

division Steel grade Ac _3- (Ac ₃ -Ac ₁ ) / X
(℃) heating
time
(min) Isothermal heating temperature
Ms + Y
(℃) Isothermal
Heat treatment time
(min) Ideal-phase Ferrite Percentage
(%) Total retained austenite percentage
(%) Strip-like type austenite ratio (%) Old Austenitic Average Crystal Grain Size (㎛) Inventive Example 1 Inventive Steel 1 X = 5.7 50 Y = 50 30 5 22 85 12 Inventive Example 2 Inventive Steel 2 X = 2.5 50 Y = 70 50 12 23 81 10 Inventory 3 Invention Steel 3 X = 2.5 50 Y = 70 50 13 26 74 14 Honorable 4 Inventive Steel 4 X = 1.7 50 Y = 70 70 20 33 77 15 Inventory 5 Inventive Steel 5 X = 2.5 40 Y = 100 20 18 15 86 13 Inventory 6 Inventive Steel 6 X = 2.5 70 Y = 50 20 10 30 83 13 Honorable 7 Inventive Steel 7 X = 2.5 80 Y = 50 20 10 19 83 12 Inventive Example 8 Inventive Steel 8 X = 2.5 40 Y = 50 20 13 20 81 9 Proposition 9 Inventive Steel 9 X = 2.5 50 Y = 50 20 13 19 77 11 Inventory 10 Inventive Steel 10 X = 2.5 120 Y = 50 20 10 33 83 13

The processing conditions of the comparative example for comparison with the inventive example of Table 4 are shown in Table 5 below. However, the processing conditions shown in Table 5 are carried out under ordinary Q / T (quenching & tempering) conditions. The cooling rate after heating of all the invention examples and comparative examples of Table 4 and Table 5 was unified at 70 ℃ / sec.

division Steel grade Heating temperature
(℃) Heating time
(min) Tempering temperature
(℃) Tempering time
(min) Old Austenitic Average Grain Size
(Μm) Comparative Example 1 Inventive Steel 1 1000 40 500 60 15 Comparative Example 2 Inventive Steel 2 1000 40 450 60 16 Comparative Example 3 Invention Steel 3 1000 40 450 60 11 Comparative Example 4 Inventive Steel 4 1000 40 500 60 19 Comparative Example 5 Inventive Steel 5 1000 40 500 60 14 Comparative Example 6 Inventive Steel 6 1000 40 450 60 16 Comparative Example 7 Inventive Steel 7 1000 40 450 60 17 Comparative Example 8 Inventive Steel 8 1000 40 500 60 13 Comparative Example 9 Inventive Steel 9 1000 40 450 60 15 Comparative Example 10 Inventive Steel 10 1000 40 450 60 16

The following tests were performed to evaluate the strength and impact toughness by using the steel products heat-treated according to the conditions of Tables 4 and 5.

The tensile test piece was tested at a cross head speed of 5 mm / min using the KS standard (KS B 0801) No. 4 test piece. The impact test piece was manufactured according to KS standard (KS B 0809) No. 3 test piece, and the notch direction was processed from the side of the rolling direction (L-T direction).

After heat treatment, the total residual austenite phase fraction in the microcomposite was measured by X-ray (Cu radiation). Respective phase fractions for the amount of retained austenite of the strip like type and the brachy type within the total residual austenite fraction were measured using a point counting method. In addition, the austenite grain size was measured by the KS standard (KS D 0205).

In order to show the effect of the present invention, the delayed fracture resistance evaluation was applied to a generally used constant load method. This evaluation method is a general method for evaluating the delayed fracture resistance by the time of additional stress or the time required to break under a specific stress. In the delayed fracture test, the test stress was determined based on the notched tensile strength. The delayed fracture tester used a constant loading type delayed fracture testing machine. The delayed fracture test specimen was prepared with a specimen diameter of 6 mmφ, a notch diameter of 4 mmφ, and a notch root radius of 0.1 mm. The test piece atmosphere solution was performed at room temperature (25 ± 5C) at a pH of 2 ± 0.5 with a wolpole buffer solution (HCl + CH3COONa). The critical delay fracture strength is the tensile strength at which the required time from failure to failure is equal to or greater than 150 hours at the same stress ratio (load stress / notch tensile strength ratio, 0.5). ÷ notch area). The number of specimens for setting the critical delay fracture strength was based on the case where more than 13 undecided specimens were used.

Table 6 shows the results of each test. As shown in Table 6, the invention examples (1 to 10) are tensile strength 160 ~ 173kg / mm ² , elongation 41 ~ 45%, cross section reduction 60 ~ 66%, impact toughness 100 ~ 124 J / cm ² and critical delay fracture strength On the contrary, in the comparative example (1-10), the tensile strength is 185-196kg / mm ² , which is much higher than necessary, but the elongation is 10-18% and the cross-sectional reduction rate is 150 ~ 155kg / mm ^2. 30 to 37%, impact toughness in the range of 20 to 40 J / cm ² is not suitable for the purpose of the present invention, the critical delay fracture strength is low 125 ~ 130kg / mm ² it turned out that the delay failure resistance is significantly lower than the invention example It became. Therefore, it can be seen that the invention examples have excellent delayed breaking strength while having a resistance ratio compared to the comparative examples.

division The tensile strength
(kg / mm ² ) Yield strength
(kg / mm ² ) Elongation
(%) Section reduction rate Impact toughness
(J / cm ² ) Critical Delay Fracture Strength
(kg / mm ² ) Inventive Example 1 164 125 44 60 124 155 Inventive Example 2 160 118 45 62 122 155 Inventory 3 161 119 41 60 100 150 Honorable 4 164 117 43 66 103 155 Inventory 5 161 124 41 63 101 150 Inventory 6 173 128 41 61 103 155 Honorable 7 170 124 41 61 113 155 Inventive Example 8 162 127 42 62 113 155 Proposition 9 161 126 43 60 111 155 Inventory 10 170 125 42 60 109 155 Comparative Example 1 190 174 13 33 35 130 Comparative Example 2 185 177 18 37 40 130 Comparative Example 3 186 176 10 34 30 125 Comparative Example 4 190 173 14 31 25 130 Comparative Example 5 187 180 10 30 20 125 Comparative Example 6 195 182 10 32 22 130 Comparative Example 7 196 184 13 35 37 130 Comparative Example 8 188 181 11 33 33 130 Comparative Example 9 189 180 12 31 30 130 Comparative Example 10 196 182 14 32 27 130

As described above, according to the present invention, not only can the cold forging property and the free machinability be greatly improved at the same time by miniaturizing graphite grains and heat treatment in a short time, but also the critical delay while achieving the high strength resistance ratio through the control of the composite structure. It is possible to provide a steel product with a significantly improved breaking strength.

Claims (18)

By weight%, C: 0.30 to 0.70%, Si: 2.0 to 4.0%, Mn: 0.1 to 1.0%, P: 0.01% (or less), S: 0.01% or less, Se: 0.001 to 0.05%, Ti: 0.001 to 0.03%, B: 0.001% to 0.003%, Al: 0.002% to 0.01%, N: 0.004% to 0.008%, O: 0.005% or less (excluding 0), and the balance is made of Fe and other impurities, and the Ti, Pre-rolling a billet having a composition in which N, B, and Al satisfy the following relation 1, and Mn, Se, and S satisfy the following relation 2; Winding the rolled wire; Cooling the wound wire rod; Graphitizing heat treatment of the cooled wire; Manufacturing the graphitized heat treated wire rod into a steel workpiece through one or more steps of cold forging and cutting; Heating to (Ac ₃ -Ac ₁₎ /5.5 temperature of the above prepared steel workpieces _{Ac 3 - - (Ac 3 -Ac} 1) /1.3 to Ac _3; Maintaining the heated steel product in the temperature range for a predetermined time; Cooling the steel workpiece to a temperature range of Ms + 50 to Ms + 110 ° C at a cooling rate of 70 ° C / sec or more; And Isothermal heat treatment in the temperature range; a method of manufacturing a high-strength steel product excellent in delayed fracture resistance comprising a. (However, Ti, N, B, Al, Mn, Se and S in the following relations 1 and 2, respectively, means the weight percent of the element.) [Relationship 1] 2.0≤ (Ti + 5B + Al) /N≤5.5 [Relationship 2] 1.0≤ (Mn / 5 + Se) /5S≤3.0 The method according to claim 1, wherein in addition to the stress, Ni: 0.05 to 1.0%, Cu: 0.01 to 0.5%, Ca: 0.0001 to 0.05%, Zr: 0.0005 to 0.008%, REM (Rare Earth Metal): 0.001 to A method for producing a high strength steel product excellent in delayed fracture resistance, characterized in that it comprises one or two or more selected from the group consisting of 0.05%. The method for producing a high strength steel product having excellent delayed fracture resistance according to claim 1 or 2, wherein the austenite grain size immediately after the finish rolling in the step of wire-rolling the billet of the composition is 30 m or less. The method of manufacturing a high strength steel product having excellent delay fracture resistance according to claim 3, wherein the wire rod finish rolling is performed at a temperature range of 840 to 900 ° C. The method according to claim 1 or 2, wherein the fraction of the cornerstone ferrite in the wire inner structure after the cooling step is 10% or less. The method of claim 5, wherein the winding of the rolled wire is performed at a temperature range of 750 to 800 ° C. The method of claim 6, wherein the step of cooling the wound wire, high strength excellent in delayed fracture resistance, characterized in that for cooling the wound wire to a temperature range of 650 ~ 690 ℃ at a cooling rate of 2.7 ~ 3.3 ℃ / sec Method of manufacturing steel products. The method of claim 1 or claim 2, wherein the step of graphitizing the cooled wire is performed in a temperature range of Ac _1- (60 ± 30) ℃ to produce a high strength steel product excellent in delayed fracture resistance Way. The method of claim 1 or claim 2, wherein the average size of the graphite grains contained in the graphite steel produced from the step of graphitizing the wire is 15㎛ or less, the production of high strength steel products excellent in delayed fracture resistance Way. The method of claim 1 or claim 2, wherein said Ac ₃ - in the step of maintaining (Ac ₃ -Ac ₁₎ After heating a steel work piece at a temperature range of _{/5.5 - (Ac 3 -Ac 1)} /1.3 to Ac ₃ The holding time is 20 minutes or more, the production method of high strength steel products excellent in delayed fracture resistance.  The method of claim 1 or 2, wherein the heat treatment holding time of the isothermal heat treatment is 20 minutes or more. By weight%, C: 0.30 ~ 0.70%, Si: 2.0 ~ 4.0%, Mn: 0.1 ~ 1.0%, P: 0.01% or less, S: 0.01% or less, Se: 0.001 ~ 0.05%, Ti: 0.001 ~ 0.03% , B: 0.001% to 0.003%, Al: 0.002% to 0.01%, N: 0.004% to 0.008%, O: 0.005% or less (excluding 0), and the balance is made of Fe and other impurities, and the Ti, N, Providing graphite steel including graphite grains having an average size of 15 μm or less therein, wherein B and Al satisfy the following relational formula 1, and Mn, Se, and S satisfy the following relational formula 2; Manufacturing the graphite steel into a steel product through one or more steps of cold forging and cutting; Heating to (Ac ₃ -Ac ₁₎ /5.5 temperature of the above prepared steel workpieces _{Ac 3 - - (Ac 3 -Ac} 1) /1.3 to Ac _3; Maintaining the heated steel product in the temperature range for a predetermined time; Cooling the steel workpiece to a temperature range of Ms + 50 to Ms + 110 ° C at a cooling rate of 70 ° C / sec or more; And Isothermal heat treatment in the temperature range; a method of manufacturing a high-strength steel product excellent in delayed fracture resistance comprising a. (However, Ti, N, B, Al, Mn, Se and S in the following relations 1 and 2, respectively, means the weight percent of the element.) [Relationship 1] 2.0≤ (Ti + 5B + Al) /N≤5.5 [Relationship 2] 1.0≤ (Mn / 5 + Se) /5S≤3.0 The method according to claim 12, wherein Ni: 0.05 to 1.0%, Cu: 0.01 to 0.5%, Ca: 0.0001 to 0.05%, Zr: 0.0005 to 0.008%, REM (Rare Earth Metal): 0.001 to A method for producing a high strength steel product excellent in delayed fracture resistance, characterized in that it comprises one or two or more selected from the group consisting of 0.05%. The method of claim 12 or claim 13, wherein said Ac ₃ - in the step of maintaining (Ac ₃ -Ac ₁₎ After heating a steel work piece at a temperature range of _{/5.5 - (Ac 3 -Ac 1)} /1.3 to Ac ₃ The holding time is 20 minutes or more, the production method of high strength steel products excellent in delayed fracture resistance. 15. The method of claim 12 or 13, wherein the heat treatment holding time of the isothermal heat treatment is 20 minutes or more. A high-strength steel product manufactured by any one of claims 1, 2, 12, or 13, and having excellent delayed fracture resistance, having a composite structure of ferrite and residual austenite. The high strength steel product having excellent delayed fracture resistance according to claim 16, wherein the fraction of retained austenite in the structure of the steel product is 15% or more. 18. The workpiece as set forth in claim 17, wherein the proportion of strip like type residual austenite in the residual austenite is 70 to 90%.