WO2023101487A1 - Cold forging wire rod and steel part having improved delayed fracture resistance, and method for manufacturing same - Google Patents

Abstract

Disclosed are: a cold forging wire rod and a steel part in which the microstructure is controlled through the alloy composition and manufacturing method, thus enabling cost reduction and enhancing delayed fracture resistance; and a method for manufacturing same. The steel part having improved delayed fracture resistance according to an embodiment of the present invention contains, in weight percent, 0.18-0.25% of C, 0.30-0.50% of Si, 0.35-0.50% of Mn, more than 0% and no more than 0.03% of P, more than 0% and no more than 0.03% of S, 0.45-0.60% of Cr, 0.015-0.03% of Ti, and 0.001-0.004% of B, with the remainder comprising Fe and inevitable impurities, and may include at least 90% by volume of auto-tempered martensite.

Description

Wire rods and steel parts for cold forging with improved resistance to delayed fracture and their manufacturing method

The present invention relates to a wire rod for cold forging with improved delayed fracture resistance, steel parts, and a manufacturing method thereof, and more particularly, to a wire rod for cold forging with improved delayed fracture resistance by controlling the microstructure, steel parts and It is about their manufacturing method.

Wire rods used as fastening bolts for automobiles and structures are required to have high strength in order to reduce the weight of automobiles and reduce the size of structures. In order to increase the strength of steel, metal strengthening mechanisms such as cold working, crystal grain refinement, martensite strengthening, and precipitation hardening are utilized. However, dislocations, grain boundaries, martensitic lath boundaries, fine precipitate boundaries, etc. used as the strengthening mechanism act as traps for hydrogen in the steel and act as a cause of inferior delayed fracture. Therefore, in high-strength bolts having a tensile strength of 1 GPa or more, there is a problem in that delayed fracture is inferior.

In order to solve this problem, in the case of conventional steel for bolts of 1 GPa or more having a tempered martensitic structure, Cr-Mo alloy steel to which Mo is added has been used. However, in order to respond to the need for cost reduction according to the development of bolt manufacturing process technology, efforts have been made to replace high-strength steel with a strength of 1 GPa or more with boron-added steel. As a result, cost reduction was realized by utilizing boron-added steel, and after confirming its safety, boron-added steel is being applied to some fastening bolts of automobiles.

However, when boron-added steel is used at 1.1 GPa or more, there is a problem in that delayed hydrogen fracture occurs. (Ref. N.Uno et al., Nippon Steel Technical Report No.97 (2008)) Therefore, for high-strength steel with a strength of 1.1 GPa or higher, Mo-added standard steel is applied, or steel companies' proprietary steel types with Mo and V added are used. was using However, for cost competitiveness, there is a need to develop high-strength steel omitting expensive carbide elements such as Mo and V.

An object of the present invention for solving the above problems is to provide a wire rod for cold forging, steel parts, and a manufacturing method thereof with improved delayed fracture resistance while cost reduction is possible by controlling the microstructure through the alloy composition and manufacturing method. is to provide

Steel parts with improved delayed fracture resistance according to an embodiment of the present invention, in weight%, C: 0.18% or more and 0.25% or less, Si: 0.30% or more and 0.50% or less, Mn: 0.35% or more and 0.50% or less, P : More than 0% and 0.03% or less, S: More than 0% and 0.03% or less, Cr: 0.45% or more and 0.60% or less, Ti: 0.015% or more and 0.03% or less, B: 0.001% or more and 0.004% or less, the remaining Fe and unavoidable impurities Including, in terms of volume fraction, autotempered martensite may be 90% or more.

In addition, in the steel component having improved delayed fracture resistance according to an embodiment of the present invention, the average thickness of carbides in prior austenite crystal grains may be 15 nm or less.

In addition, the steel component having improved delayed fracture resistance according to an embodiment of the present invention may have a tensile strength of 1200 MPa or more.

In addition, in the method of manufacturing a steel part having improved resistance to delayed fracture according to an embodiment of the present invention, in weight%, C: 0.18% or more and 0.25% or less, Si: 0.30% or more and 0.50% or less, Mn: 0.35% or more 0.50% or less, P: more than 0% and 0.03% or less, S: more than 0% and 0.03% or less, Cr: 0.45% or more and 0.60% or less, Ti: 0.015% or more and 0.03% or less, B: 0.001% or more and 0.004% or less, the rest Preparing a steel material containing Fe and unavoidable impurities; preparing a wire rod by finish rolling the steel; winding the wire rod; After freshening the wound wire, subjecting it to spheroidizing heat treatment; Forming the wire rod subjected to the spheroidization heat treatment to prepare a part; A step of quenching the part after austenizing may be included.

In addition, in the method for manufacturing a steel part with improved delayed fracture resistance according to an embodiment of the present invention, the finish rolling may be performed at 880 to 980 ° C, and the winding may be performed at 830 to 930 ° C.

In addition, in the method for manufacturing a steel part having improved delayed fracture resistance according to an embodiment of the present invention, the spheroidizing heat treatment may be performed at a maximum temperature in the range of 745 to 765 °C.

In addition, in the method for manufacturing a steel component with improved delayed fracture resistance according to an embodiment of the present invention, the austenizing may be performed at 870 to 940 ° C.

In addition, in the method for manufacturing a steel part having improved resistance to delayed fracture according to an embodiment of the present invention, the quenching may be performed with a refrigerant at 10 to 80 °C.

In addition, the wire rod for cold forging according to an embodiment of the present invention, in weight%, C: 0.18% or more and 0.25% or less, Si: 0.30% or more and 0.50% or less, Mn: 0.35% or more and 0.50% or less, P: 0 % more than 0.03%, S: more than 0% and less than 0.03%, Cr: more than 0.45% and less than 0.60%, Ti: more than 0.015% and less than 0.03%, B: more than 0.001% and less than 0.004%, the remainder including Fe and unavoidable impurities, , and may have a diameter of 5.5 to 20 mm.

According to one embodiment of the present invention, by controlling the microstructure through the alloy composition and manufacturing method, it is possible to provide a wire rod for cold forging, steel parts and their manufacturing method with improved delayed fracture resistance while cost reduction is possible. .

Steel parts with improved delayed fracture resistance according to an embodiment of the present invention, in weight%, C: 0.18% or more and 0.25% or less, Si: 0.30% or more and 0.50% or less, Mn: 0.35% or more and 0.50% or less, P : More than 0% and 0.03% or less, S: More than 0% and 0.03% or less, Cr: 0.45% or more and 0.60% or less, Ti: 0.015% or more and 0.03% or less, B: 0.001% or more and 0.004% or less, the remaining Fe and unavoidable impurities Including, in terms of volume fraction, autotempered martensite is 90% or more.

The following examples are presented to sufficiently convey the spirit of the present invention to those skilled in the art. The present invention may be embodied in other forms without being limited to only the embodiments presented herein. In the drawings, in order to clarify the present invention, illustration of parts irrelevant to the description may be omitted, and the size of components may be slightly exaggerated to aid understanding.

Throughout the specification, when a certain component is said to "include", it means that it may further include other components without excluding other components unless otherwise stated.

Expressions in the singular number include plural expressions unless the context clearly dictates otherwise.

Hereinafter, the reason for limiting the numerical value of the alloy component content in the embodiments of the present invention will be described. Hereinafter, unless otherwise specified, units are % by weight.

Steel parts with improved delayed fracture resistance according to an embodiment of the present invention, in weight%, C: 0.18% or more and 0.25% or less, Si: 0.30% or more and 0.50% or less, Mn: 0.35% or more and 0.50% or less, P : More than 0% and 0.03% or less, S: More than 0% and 0.03% or less, Cr: 0.45% or more and 0.60% or less, Ti: 0.015% or more and 0.03% or less, B: 0.001% or more and 0.004% or less, the remaining Fe and unavoidable impurities can include

The content of C (carbon) may be 0.18% or more and 0.25% or less.

C is an element added to ensure product strength. Considering this, C may be added in an amount of 0.18% or more. However, if the content of C is excessive, it may cause delayed fracture due to an increase in strength. Considering this, the upper limit of the C content may be limited to 0.25%.

The content of Si (silicon) may be 0.30% or more and 0.50% or less.

Si can be added for deoxidation of steel. In addition, Si is an element effective in securing strength through solid solution strengthening. Considering this, Si may be added in an amount of 0.30% or more. However, when the content of Si is excessive, impact properties and formability may be inferior. Considering this, the upper limit of the Si content may be limited to 0.50%. Preferably, the content of Si may be 0.31% or more and 0.48% or less.

The content of Mn (manganese) may be 0.35% or more and 0.50% or less.

Mn is a very effective element for improving hardenability and forming a substitutional solid solution in the matrix structure to produce a solid solution strengthening effect. In addition, when the content of Mn is low, it is not sufficiently combined with S (sulfur) introduced as impurities in the steel, which may cause casting cracks and the like. Considering this, Mn may be added in an amount of 0.35% or more. However, when the Mn content is excessive, delayed fracture resistance may be deteriorated by forming coarse MnS. Considering this, the upper limit of the Mn content may be limited to 0.50%. Preferably, the content of Mn may be 0.36% or more and 0.49% or less.

The content of P (phosphorus) may be greater than 0% and 0.03% or less.

P is segregated at the grain boundary and acts as a cause of lowering toughness and reducing delayed fracture resistance. Therefore, in the present invention, impurities can be managed. In consideration of this, the upper limit of the P content may be limited to 0.03%, and the closer to 0% is preferable.

The content of S (sulfur) may be greater than 0% and 0.03% or less.

S, like P, is segregated at grain boundaries to reduce toughness and acts as a cause of inhibiting hot rolling by forming low melting point emulsifiers. Therefore, in the present invention, impurities can be managed. Considering this, the upper limit of the S content may be limited to 0.03%, and the closer to 0% is preferable.

The content of Cr (chromium) may be 0.45% or more and 0.60% or less.

Cr is a very effective element for enhancing hardenability and forming a substitutional solid solution in the base structure to produce a solid solution strengthening effect. Considering this, Cr may be added in an amount of 0.45% or more. However, when the Cr content is excessive, the c/a ratio of the corrosion pit increases due to the formation of a chromium oxide layer on the surface, so that delayed fracture resistance may be inferior due to the notch effect. Considering this, the upper limit of the Cr content may be limited to 0.60%. Preferably, the Cr content may be 0.46% or more and 0.59% or less.

The content of Ti (titanium) may be 0.015% or more and 0.03% or less.

Ti is an element effective in preventing boron (B) from combining with N by forming titanium nitride by combining with N (nitrogen) introduced into steel. Considering this, Ti may be added in an amount of 0.015% or more. However, if the content of Ti is excessive, coarse carbonitrides may be formed and delayed fracture resistance may be inferior. Considering this, the upper limit of the Ti content may be limited to 0.03%. Preferably, the content of Ti may be 0.023% or more and 0.026% or less.

The content of B (boron) may be 0.001% or more and 0.004% or less.

B is an element effective in improving hardenability. Considering this, B may be added in an amount of 0.001% or more. However, when the content of B is excessive, by forming Fe ₂₃ (CB) ₆ carbides at grain boundaries, brittleness of austenite grain boundaries may be caused, and delayed fracture resistance may be inferior. Considering this, the upper limit of the B content may be limited to 0.004%. Preferably, the content of B may be 0.0018% or more and 0.0023% or less.

The remaining component of the present invention is iron (Fe). However, since unintended impurities from raw materials or the surrounding environment may inevitably be mixed in a normal manufacturing process, this cannot be excluded. Since these impurities are known to anyone skilled in the ordinary manufacturing process, not all of them are specifically mentioned in this specification.

In the steel part having improved resistance to delayed fracture according to an embodiment of the present invention, the volume fraction of auto-tempered martensite may be 90% or more.

When the auto-tempered martensite is less than 90%, it is difficult to secure sufficient toughness and delayed fracture resistance may be inferior. Therefore, in the present invention, it is possible to control the auto-tempered martensite to be 90% or more through the alloy components and the manufacturing process. In particular, the auto-tempered martensitic structure is characterized in that it can be automatically tempered during quenching without an additional tempering heat treatment process.

In addition, the steel parts with improved resistance to delayed fracture according to an embodiment of the present invention can be automatically tempered during quenching without an additional tempering heat treatment process, so that the average thickness of carbides in old austenite grains is 15 nm or less You can control it.

The surface direction of the carbide precipitated during auto-tempering in a plate-type is not effective as a hydrogen trap due to its high coherency, and the side part of the plate-type has low coherency, making it a non-diffusible hydrogen trap. By acting, it is known to improve hydrogen retardation resistance.

Therefore, when the average thickness of the carbide in the prior austenite crystal grain is small, the coherency of the carbide interface is increased, and it may be difficult to improve the hydrogen delayed fracture resistance. On the other hand, when the average thickness of carbides in old austenite crystal grains is large, the number of carbides decreases, making it difficult to improve delayed fracture resistance. Therefore, in the present invention, the delayed fracture resistance is improved by controlling the average thickness of carbides in the prior austenite crystal grains to be 15 nm or less.

In addition, the steel component having improved delayed fracture resistance according to an embodiment of the present invention may have a tensile strength of 1200 MPa or more by controlling the alloy composition and manufacturing method.

Next, a method for manufacturing a steel part having improved resistance to delayed fracture according to another aspect of the present invention will be described.

In the method for manufacturing a steel part with improved delayed fracture resistance according to an embodiment of the present invention, in weight%, C: 0.18% or more and 0.25% or less, Si: 0.30% or more and 0.50% or less, Mn: 0.35% or more 0.50% Below, P: more than 0% and 0.03% or less, S: more than 0% and 0.03% or less, Cr: 0.45% or more and 0.60% or less, Ti: 0.005% or more and 0.03% or less, B: 0.001% or more and 0.004% or less, the remainder Fe and Preparing a steel material containing unavoidable impurities; preparing a wire rod by finish rolling the steel; winding the wire rod; After freshening the wound wire, subjecting it to spheroidizing heat treatment; Forming the wire rod subjected to the spheroidization heat treatment to prepare a part; A step of quenching the part after austenizing may be included.

The reasons for limiting the range of components of each alloy composition are as described above, and each manufacturing step will be described in detail below.

First, after preparing a steel material satisfying the alloy composition, a series of finish rolling, winding, spheroidization heat treatment, forming, austenizing and quenching may be performed.

First, the steel material may be finish-rolled at 880 to 980 ° C to prepare a wire rod, and the wire rod may be wound at 830 to 930 ° C.

When the finish rolling temperature or coiling temperature is low, since the surface layer is a quasi-two-phase inverse, a surface ferrite decarburized layer may be formed by phase transformation. Therefore, when the finish rolling temperature or coiling temperature is low, a ferrite decarburized layer is formed on the surface even during heat treatment of the steel part, and delayed fracture resistance may be inferior. Considering this, the finish rolling temperature may be 880°C or higher, or the coiling temperature may be 830°C or higher.

On the other hand, when the finish rolling temperature or coiling temperature is high, decarburization is accelerated by diffusion and a ferrite decarburized layer is formed on the surface, so that delayed fracture resistance may be inferior. Considering this, the finish rolling temperature may be 980°C or less, or the coiling temperature may be 930°C or less.

Next, after the wound wire is fresh for the purpose, spheroidizing heat treatment may be performed at a maximum temperature in the range of 745 to 765 ° C.

When the maximum temperature of the spheroidization heat treatment is too low or too high, the hardness of the spheroidization heat treatment material increases due to the low spheroidization rate, thereby causing cracks due to poor formability during processing of steel parts. In consideration of this, the maximum temperature of the spheroidization heat treatment may be performed at 745 to 765 ° C.

The wire rod subjected to the spheroidizing heat treatment may be molded to suit the purpose and prepared as a steel part, and then austenizing may be performed at 870 to 940 ° C.

When the austenizing temperature is low, since reverse austenite transformation does not sufficiently occur, the martensitic structure after quenching may be non-uniform, resulting in inferior toughness. Considering this, the austenizing temperature may be 870°C or higher. However, when the austenizing temperature is high, the grain size of austenite becomes coarse, and delayed fracture resistance may deteriorate. Considering this, the upper limit of the austenizing temperature may be limited to 940°C.

Next, the austenized steel parts may be quenched with a refrigerant at 10 to 80°C.

When the temperature of the quenching refrigerant is low, fine quenching cracks may occur due to thermal deformation of steel parts, causing delayed fracture. Considering this, the temperature of the quenching refrigerant may be 10 °C or higher. However, when the temperature of the quenching refrigerant is high, the auto-tempering effect increases, making it difficult to achieve the desired strength. Considering this, the upper limit of the temperature of the quenching refrigerant may be limited to 80°C.

From the above-described process, the final microstructure of the steel part can realize an auto-tempered martensite structure of 90% or more without a tempering process, and a structure in which carbides with an average thickness of 15 nm or less in old austenite crystal grains are precipitated can be implemented. . Accordingly, delayed fracture resistance may be improved by controlling the microstructure.

Next, a wire rod for cold forging according to another aspect of the present invention will be described.

In the wire rod for cold forging according to an embodiment of the present invention, in weight%, C: 0.18% or more and 0.25% or less, Si: 0.30% or more and 0.50% or less, Mn: 0.35% or more and 0.50% or less, P: more than 0% 0.03% or less, S: more than 0% and 0.03% or less, Cr: 0.45% or more and 0.60% or less, Ti: 0.005% or more and 0.03% or less, B: 0.001% or more and 0.004% or less, including the remainder Fe and unavoidable impurities, diameter It may be 5.5 to 20 mm.

The reasons for limiting the range of components of each alloy composition are as described above, and the wire rod for cold forging according to an example of the present invention may be manufactured with a diameter of 5.5 to 20 mm. However, it is not limited thereto, and may be manufactured in various diameters depending on the purpose.

Hereinafter, the present invention will be described in more detail through examples. However, the description of these examples is only for exemplifying the practice of the present invention, and the present invention is not limited by the description of these examples. This is because the scope of the present invention is determined by the matters described in the claims and the matters reasonably inferred therefrom.

{Example}

For the various alloy composition ranges shown in Table 1 below, steel materials were prepared, finished rolled at 910 ° C to prepare a wire rod having a diameter of 15 mm, and then wound into a coil at 880 ° C. The wound wire rod was subjected to spheroidization heat treatment at a maximum temperature of 755 ° C, molded into a screw M12 standard bolt, austenized at 890 ° C, and quenched in a refrigerant at 60 ° C. On the other hand, the spheroidization heat treatment temperature means the highest heating temperature.

division	alloy component
	C	Si	Mn	P	S	Cr	Ti	B
Example 1	0.18	0.40	0.46	0.011	0.005	0.53	0.024	0.0018
Example 2	0.25	0.41	0.45	0.010	0.005	0.54	0.024	0.0021
Example 3	0.20	0.40	0.36	0.009	0.005	0.52	0.026	0.0023
Example 4	0.21	0.44	0.49	0.011	0.005	0.52	0.023	0.0021
Example 5	0.19	0.41	0.44	0.011	0.005	0.46	0.024	0.0020
Example 6	0.22	0.41	0.43	0.010	0.005	0.59	0.025	0.0021
Example 7	0.23	0.31	0.42	0.009	0.005	0.52	0.024	0.0022
Example 8	0.21	0.48	0.44	0.010	0.005	0.53	0.024	0.0018
Comparative Example 1	0.17	0.41	0.46	0.011	0.005	0.52	0.024	0.0021
Comparative Example 2	0.26	0.42	0.47	0.010	0.005	0.52	0.023	0.0023
Comparative Example 3	0.18	0.25	0.45	0.011	0.005	0.50	0.024	0.0022
Comparative Example 4	0.24	0.40	0.53	0.011	0.005	0.51	0.024	0.0020
Comparative Example 5	0.23	0.37	0.49	0.009	0.005	0.42	0.023	0.0019
Comparative Example 6	0.22	0.42	0.47	0.009	0.005	0.64	0.022	0.0021

Table 2 below shows the presence or absence of cracks according to the tensile strength, carbide thickness, and delayed fracture performance evaluation of the manufactured bolts. The tensile strength was measured using a Zwick/Roell ZWICK Z250 tensile tester. The tensile strength measurement test was performed using a tensile specimen having a diameter of 10 mm and a gauge part diameter of 6.25 mm.

Carbide thickness measurements were made with a FEI Tecnai OSIRIS transmission electron microscope (TEM). At this time, the carbide thickness was measured by randomly measuring 5 places on the replica specimen and expressed as an average thickness, and the short axis of the carbide formed in a plate-type was defined as the thickness and measured.

The test method for evaluating the delayed failure performance is to fasten the bolt to the structure with the fastening force of the yield strength, and then observe the presence or absence of cracks in the screw thread, which is the stress concentration part, before / after immersing the bolt in 5% hydrochloric acid + 95% distilled water solution for 10 minutes. The delayed failure simulation method was used.

As a result of the delayed fracture performance evaluation, when cracks occurred, 'O' was marked, and when cracks did not occur, 'X' was marked.

division	Tensile strength (MPa)	Carbide thickness (nm)	presence or absence of cracks
Example 1	1213	12	X
Example 2	1680	13	X
Example 3	1365	11	X
Example 4	1391	11	X
Example 5	1265	12	X
Example 6	1354	13	X
Example 7	1566	14	X
Example 8	1421	10	X
Comparative Example 1	1185	13	X
Comparative Example 2	1693	16	O
Comparative Example 3	1193	16	O
Comparative Example 4	1655	14	O
Comparative Example 5	1586	13	O
Comparative Example 6	1543	15	O

Referring to Table 2, Examples 1 to 8 satisfied the alloy composition, component range and manufacturing process presented in the present invention. Therefore, Examples 1 to 8 satisfied tensile strength of 1200 MPa or more, carbide thickness of 15 nm or less, and no cracks occurred as a result of delayed fracture performance evaluation. However, Comparative Example 1 had a low C content, so the tensile strength 1200MPa or more was not satisfied.

In Comparative Example 2, the thickness of the carbide exceeded 15 nm due to the high C content, and cracks occurred as a result of the delayed fracture performance evaluation.

In addition, Comparative Example 3 did not satisfy the tensile strength of 1200 MPa or more due to the low Si content.

In Comparative Example 4, coarse MnS was formed due to the high Mn content, resulting in cracks as a result of the delayed fracture performance evaluation.

In addition, in Comparative Example 5, a bainite-mixed structure was formed in the microstructure because the Cr content was low, and cracks occurred as a result of the delayed fracture performance evaluation.

In addition, Comparative Example 6 had a high Cr content, and sharp corrosion pits were formed when corroded by hydrochloric acid, resulting in cracks as a result of delayed fracture performance evaluation.

Next, in Table 3 below, with respect to the alloy components of Example 5 of Table 1, steel materials were prepared, and bolts were prepared at the finish rolling temperature, coiling temperature, maximum spheroidizing heat treatment temperature, and austenizing temperature shown in Table 3 below. After manufacturing, as a result of the delayed fracture performance evaluation, the presence or absence of cracks was shown.

division	Temperature (℃)				presence or absence of cracks
	finish rolling	winding	Spheroidization heat treatment	austenizing
Example 5	910	880	755	890	X
Comparative Example 7	985	935	755	890	O
Comparative Example 8	865	825	755	890	O
Comparative Example 9	910	880	755	950	O
Comparative Example 10	910	880	755	860	O
Comparative Example 11	910	880	740	890	O
Comparative Example 12	910	880	770	890	O

In Comparative Example 7, the finish rolling temperature and coiling temperature were high, and the grain size of prior austenite increased, resulting in cracks as a result of the delayed fracture performance evaluation. Due to this formation, cracks occurred as a result of delayed fracture performance evaluation.

In Comparative Example 9, the austenizing temperature was high, and the prior austenite crystal grain size increased, so cracks occurred as a result of the delayed fracture performance evaluation.

In Comparative Example 10, the austenizing temperature was low and entered the quasi-2 phase region, and a ferrite decarburized layer was formed during heating, resulting in cracks as a result of the delayed fracture performance evaluation.

In Comparative Examples 11 and 12, since the maximum temperature of the spheroidization heat treatment was low or high, respectively, the spheroidization heat treatment was not sufficiently performed, resulting in poor formability. Therefore, in Comparative Examples 11 and 12, cracks were formed during thread forming, and as a result of the delayed fracture performance evaluation, cracks occurred.

	Austenizing temperature (℃)	Quenching oil temperature (℃)	Otomartensite fraction (%)	carbide thickness (nm)	delayed destruction presence or absence of cracks
Example 5	910	60	95%	12	X
Comparative Example 13	910	85	89%	16	O
Comparative Example 14	910	95	86%	18	O
Comparative Example 15	910	105	85%	21	O
Comparative Example 16	910	115	85%	22	O

Comparative Examples 13 to 16 used steel materials having the same alloy components and composition ranges as Example 5, and compared to Example 5, the auto martensite fraction and carbide thickness did not satisfy the ranges of the present application, resulting in delayed fracture performance evaluation. It can be seen that cracks occur.

According to the present invention, by controlling the microstructure through the alloy composition and manufacturing method, it is possible to provide a wire rod and steel parts for cold forging with improved resistance to delayed fracture while reducing costs, and thus industrial applicability is recognized. .

Claims (9)

1. In weight percent, C: 0.18% or more and 0.25% or less, Si: 0.30% or more and 0.50% or less, Mn: 0.35% or more and 0.50% or less, P: 0% or more and 0.03% or less, S: 0% or more and 0.03% or less, Cr : 0.45% or more and 0.60% or less, Ti: 0.015% or more and 0.03% or less, B: 0.001% or more and 0.004% or less, including the remainder Fe and unavoidable impurities,

   Steel parts with improved resistance to delayed fracture, with an autotempered martensite of 90% or more by volume fraction.
2. The method of claim 1,

   Steel parts with improved resistance to delayed fracture, with an average thickness of carbides in the old austenite crystal grains of 15 nm or less.
3. The method of claim 1,

   Steel parts with improved resistance to delayed fracture with a tensile strength of 1200 MPa or more.
4. In weight percent, C: 0.18% or more and 0.25% or less, Si: 0.30% or more and 0.50% or less, Mn: 0.35% or more and 0.50% or less, P: 0% or more and 0.03% or less, S: 0% or more and 0.03% or less, Cr : 0.45% or more and 0.60% or less, Ti: 0.015% or more and 0.03% or less, B: 0.001% or more and 0.004% or less, preparing a steel material containing the remaining Fe and unavoidable impurities;

   preparing a wire rod by finish rolling the steel;

   winding the wire rod;

   After freshening the wound wire rod, spheroidization heat treatment step;

   Forming the wire rod subjected to the spheroidization heat treatment to prepare a part;

   A method of manufacturing a steel part having improved delayed fracture resistance, comprising the step of quenching the part after austenizing.
5. The method of claim 4,

   The finish rolling is performed at 880 to 980 ° C,

   The winding is performed at 830 to 930 ° C., a method of manufacturing a steel part with improved delayed fracture resistance.
6. The method of claim 4,

   The spheroidizing heat treatment is a method of manufacturing a steel part having improved delayed fracture resistance, which is performed in a range of a maximum temperature of 745 to 765 ° C.
7. The method of claim 4,

   The austenizing is performed at 870 to 940 ° C., a method of manufacturing a steel part with improved delayed fracture resistance.
8. The method of claim 4,

   The quenching is performed with a refrigerant of 10 to 80 ° C., a method of manufacturing a steel part with improved delayed fracture resistance.
9. In weight percent, C: 0.18% or more and 0.25% or less, Si: 0.30% or more and 0.50% or less, Mn: 0.35% or more and 0.50% or less, P: 0% or more and 0.03% or less, S: 0% or more and 0.03% or less, Cr : 0.45% or more and 0.60% or less, Ti: 0.015% or more and 0.03% or less, B: 0.001% or more and 0.004% or less, including the remainder Fe and unavoidable impurities,

   A wire rod for cold forging with a diameter of 5.5 to 20 mm.