ULTRA-HIGH-STRENGTH STEEL WIRE HAVING EXCELLENT RESISTANCE TO DELAYED FRACTURE AND MANUFACTURING METHOD THEREOF

Abstract

An ultra-high-strength steel wire rod having excellent resistance to delayed fracture includes, by wt %, 0.7-1.2% C, 0.25-0.5% Si, 0.5-0.8% Mn, 0.02-0.1% V and a balance of Fe and inevitable impurities. The method includes the steps of heating the above steel composition to 1100° C. or lower and hot rolling at a temperature of 900-1000° C., followed by cooling to 600-650° C. at a prescribed rate, followed by cold drawing at a reduction ratio of 60-80%.

Description

TECHNICAL FIELD

The present invention relates to a steel wire rod which is used for the manufacturing of automotive engine bolts requiring ultra-high strength, and more particularly to a steel wire rod having excellent resistance to hydrogen delayed-fracture and a manufacturing method thereof.

BACKGROUND ART

In recent years, as the demand for lightweight and high-performance automobiles has increased, the high-strength requirement for engine parts such as bolts has increased, in order to reduce the consumption of energy. These days, high-strength bolts are being manufactured as bolts having a strength of about 1200 MPa from an alloy steel such as SCM435 or SCM440 through quenching and tempering. However, a steel having a tensile strength of 1300 MPa is likely to undergo delayed fracture by hydrogen, and thus has not been used for the manufacturing of super-high-strength bolts.

The biggest issue to consider in the development of high-strength bolts is delayed fracture. The term delayed fracture refers to a phenomenon in which bolts suddenly fracture when a specific tensile strength (about 1200 MPa) is applied thereto. This phenomenon occurs mainly at the notch or head portions of bolts and is known to be attributable to hydrogen embrittlement in a triaxial stress state. Thus, in the development of high-strength bolts having a strength of about 1200 MPa or higher, it is required to ensure the safety of the bolts by increasing the resistance to delayed fracture thereof.

A Japanese steel company has developed a high-strength pearlite steel based on pearlite, which has improved resistance to delayed fracture through hydrogen trap sites formed at the pearlite/cementite and maintains the characteristic strength of pearlite. This pearlite steel is being supplied to some automobile companies.

However, in the above-mentioned pearlite steel, more than 0.2 wt % Cr should be added during a drawing process for sizing after the production of the steel in order to improve the tensile strength and ensure the drawability thereof, and isothermal transformation is necessarily required. Thus, this pearlite steel has the disadvantages of a high production cost and a complex manufacturing process. Another disadvantage is that very accurate cooling conditions are required for the production of the steel.

Also, in an attempt to improve the delayed-fracture resistance of high-strength steel wire rods having a strength of 1200 MPa or higher, there is a technique in which each of grain refining elements, including Ti, Nb and V, is added in an amount of 0.5 wt % or higher, and in which corrosion-resistant elements, such as Ni, Cu, Co and the like, and carbide elements are added. However, this technique is disadvantageous in that the production cost of the steel is very high, because lead patenting is necessarily required to ensure the transformation stability of pearlite.

Meanwhile, microalloyed steel for substituting for the bolt alloy steel that has been used to date in automobile engines has been much studied in terms of reducing the production cost by omitting heat treatment. However, in recent years, complex forging designs have been used in order to impart lightweight and high-strength properties to automobiles and reduce the number of automobile parts, and such complex forging designs may cause deformation in the microalloyed steel when a conventional annealing and tempering process is applied. For this reason, it is substantially impossible to apply the microalloyed steel.

Thus, there has been research into a method of improving toughness through austenite grain refinement by reducing the content of C and adding a trace amount of Ti, and there has been research into a method of achieving high strength by forming acicular ferrite through the addition of a small amount of Mo. However, these methods have the problem of increasing production costs due to the addition of, relatively expensive alloying elements.

Also, there have been suggested methods of improving toughness by reducing the content of C to 0.1% and adding Cr and Mo and of transforming the microstructure of steel to martensite through controlled cooling. However, these methods have problems associated with a decrease in toughness, the addition of expensive elements such as Cr and Mo, and the embodiment of equipment for controlled cooling.

Meanwhile, as mentioned above, limitations in further improving the tensile strength of an alloy steel having a tensile strength of about 1200 MPa have not yet been overcome. Also, although a few technologies related to ultra-high-strength wire rods were suggested in Japan, these technologies necessarily require the addition of expensive alloying elements and lead patenting, making it impossible to ensure price competitiveness. Particularly, it is actually difficult to ensure stable data on delayed-fracture characteristics caused by hydrogen.

Accordingly, there has been a need for a technology of manufacturing ultra-high-strength steel wire rods, which reduces the number of necessary processes by omitting basic heat treatment, as in the case of microalloyed steel, ensures price competitiveness through the use of trace amounts of alloying elements and ensures resistance to delayed fracture.

DISCLOSURE

Technical Problem

An aspect of the present invention provides a steel wire rod having both ultra-high strength and excellent resistance to delayed fracture, and a manufacturing method thereof.

Technical Solution

According to an aspect of the present invention, there is provided an ultra-high-strength steel wire rod having excellent resistance to delayed fracture, the wire rod including, by wt %, 0.7-1.2% C, 0.25-0.5% Si, 0.5-0.8% Mn, 0.02-0.1% V and a balance of Fe and inevitable impurities.

According to another aspect of the present invention, there is provided a method for manufacturing an ultra-high-strength steel wire rod having excellent resistance to delayed fracture, the method including the steps of: heating a steel, which includes, by wt %, 0.7-1.2% C, 0.25-0.5% Si, 0.5-0.8% Mn, 0.02-0.1% V and a balance of Fe and inevitable impurities, to 1000-1100° C., and hot-rolling the heated steel at a temperature of 900˜1000112; cooling the rolled steel to 600˜650° C. to a rate of 5˜10° C./s; and cold-drawing the cooled steel at a reduction ratio of 60˜80%.

Advantageous Effects

In the steel wire rod according to the present invention, the strength of pearlite can be increased due to the precipitation hardening effect according to the addition of V, and diffusible hydrogen-trap sites can be increased through the formation of V(C,N) precipitates, thereby the hydrogen delayed-fracture resistance of the wire rod. Thus, when the wire rod of the present invention is used for the manufacturing of automobile bolts and the like, it can contribute to a decrease in weight and an increase in performance of automobiles.

Also, the manufacturing method of the present invention offers excellent price competitiveness by omitting lead patenting and expensive alloying elements, and can act as the basis of novel manufacturing methods having no limitation in process conditions.

DESCRIPTION OF DRAWINGS

The above and other aspects, features and other advantages of the present invention will be more clearly understood from the following detailed description taken in conjunction with the accompanying drawings, in which:

FIGS. 1A and 1B are photographs showing the results of observing the microstructures of steel wire rods according to a conventional example and inventive example 1, respectively;

FIGS. 2A and 2B are photographs showing the results of observing the microstructures of steel wire rods according to a conventional example and inventive example 1, respectively;

FIGS. 3A through 3F are a set of photographs showing the results of observing the microstructures of steel wire rods according to a conventional example, inventive examples and comparative examples;

FIGS. 4A and 4B schematically show the microstructures of steel wire rods according to a conventional example and inventive example 1, respectively;

FIG. 5 is a graphic diagram showing the relationship between the fracture stress and diffusible hydrogen content of steel wire rods according to a conventional example and inventive example; and

FIG. 6 is a graphic diagram showing the change in tensile strength according to diameter during drawing for steel wire rods of a conventional example, inventive examples and comparative examples.

BEST MODE

Exemplary embodiments of the present invention will now be described in detail with reference to the accompanying drawings. The invention may, however, be embodied in many different forms and should not be construed as being limited to the embodiments set forth herein. Rather, these embodiments are provided so that this disclosure will be thorough and complete, and will fully convey the scope of the invention to those skilled in the art. In the drawings, the thicknesses of layers and regions are exaggerated for clarity. Like reference numerals in the drawings denote like elements, and thus their description will be omitted.

The inventors of the present invention have conducted many studies to solve the delayed-fracture problem that is the biggest issue in developing high-strength bolts for automobiles. As a result, the present inventors have found that, when vanadium carbonitride is formed in the ferrite matrix of pearlite by adding a trace amount of vanadium, it increases the strength of pearlite and acts as a diffusible hydrogen-trap site to improve hydrogen delayed-fracture resistance, thereby reaching the present invention.

Hereinafter, the steel wire rod of the present invention will be described in detail.

First, the composition of the steel wire rod of the present invention will now be described (compositional amounts are hereinafter expressed as wt %).

The most important alloying element in the wire rod of the present invention is vanadium (V). The wire rod of the present invention has a V content of 0.02˜0.1%. V forms V(C,N) precipitates in the ferrite matrix. The precipitates increase the strength of pearlite and also act as diffusible hydrogen-trap sites.

If the V content is less than 0.02%, the solid solubility with nitrogen and carbon will decrease, making it difficult to effectively form precipitates, and if the V content is more than 0.1%, the precipitation of V in the ferrite matrix will be excessive, and thus it will cause fractures in the steel during rolling and drawing and rapidly reduce the cold forgeability of the steel.

The content of carbon (C) in the steel is preferably 0.7˜1.2%. C is an essential alloying element that is generally added to ensure the strength of steel. If the content of C is less than 0.7%, it cannot ensure sufficient strength, thus making it impossible to ensure an ultra-high-strength steel. If the C content is more than 1.2%, it can cause cracks or fractures during rolling and drawing processes.

The content of manganese (Mn) in the steel is preferably 0.5˜0.8%. Mn is an alloying element that increases the strength of the steel and influences the impact properties of the steel. Also, it increases the rolling properties of the steel and reduces the embrittlement of the steel. If the content of Mn is less than 0.5%, the strength reinforcing effect will be insignificant, and if it is more than 0.8%, it will result in severe hardening. For this reason, the content of Mn is preferably limited to 0.5˜0.8 wt %.

The content of silicon (Si) in the steel is preferably 0.025˜0.5%. Si forms a solid solution in the ferrite of pearlite to increase the strength of the steel. If the content of Si is less than 0.25%, the effect of increasing the strength of the steel will be insufficient, and if it is more than 0.5%, it will increase the hardening of the steel during cold forging to reduce the toughness of the steel.

The content of phosphorus (P) in the steel is preferably 0.02% or less. Because P can be segregated in the grain boundary to reduce the toughness of the steel, the content of P is preferably as low as possible. For this reason, the upper limit of the P content is preferably limited to 0.02%.

The content of sulfur (S) in the steel is preferably 0.02% or less. S, which is a low-boiling-point element, can bond with Mn to reduce the toughness of the steel and can also adversely affect the properties of the high-strength wire rod, and for this reason, the content thereof is preferably as low as possible. Thus, the upper limit of the S content is preferably limited to 0.02% in view of inevitable problems occurring in a refining process.

In addition to the above-described elements, 60 ppm of nitrogen (N) may be added. N forms VN, but corresponds to an impurity that is incorporated into molten steel. For this reason, the content of N is preferably not more than 60 ppm.

Meanwhile, in the present invention, the precipitate-firming elements Ti and Nb, other than V, are not positively added, except for the case in which they are added as inevitable elements. This is because, if Ti is added in combination with V, nitrogen in the molten steel will first react with Ti to form a TiN precipitate, such that a V precipitate cannot be effectively formed, whereby the effect of improving the delayed-fracture resistance of the steel by the V precipitate cannot be obtained. Also, if V is added in combination with Nb, the austenite grains can be refined, but the price of the steel will be inevitably increased, and Nb will interfere with the formation of a V precipitate, because it is highly reactive with nitrogen.

In addition, the steel wire rod of the present invention comprises Fe and inevitable impurities.

Hereinafter, the microstructure of the steel wire rod of the present invention will be described in detail.

In the steel wire rod of the present invention, V(C,N) precipitates are preferably distributed in the ferrite structure of pearlite. The V(C,N) precipitates prevent the precipitation of film-like cementite and are distributed in the ferrite structure of pearlite to act as strong hydrogen-trap site, thereby improving the hydrogen delayed-fracture resistance of the steel.

The average particle size of the V(C,N) precipitates is preferably 30 nm or less, and the number of the V(C,N) precipitates is preferably 1×10⁹/mm² or more.

If the size of the V(C,N) precipitates is more than 30 nm, these precipitates will not be finely distributed in the ferrite matrix of pearlite, and thus the effect of increasing the strength of the steel through the uniform distribution of the precipitates will not be obtained. On the other hand, if the V(C,N) precipitates are coarse, they can form coarse precipitates in ferrite to cause fractures, rather than improving the tensile strength of the steel by suppressing the movement of dislocations. For these reasons, the size of the precipitates is preferably 30 nm or less.

Further, the reason for which the number of the precipitates must be 1×10⁹/mm², or more is that, if the number of the precipitates is less than 1×10⁹/mm², it will be difficult to ensure a precipitation hardening effect by the V precipitate, and thus the strength sought in the present invention cannot be achieved. If the number of the precipitates is too large, the precipitation hardening effect can be maximized to cause wire breakage during wire drawing; however, in the present invention, the number of precipitates are not specifically limited, because the content of V is limited.

Also, the steel wire rod of the present invention has a pearlite structure. As the lamellar spacing of the pearlite structure decreases, the tensile strength and ductility of the wire rod increase. The lamellar spacing of the pearlite structure of the steel wire rod according to the present invention is preferably 150-300 M.

The ductility and strength of pearlite depend on the lamellar spacing of the pearlite. Particularly, the yield strength of pearlite depends on the lamellar spacing thereof, and this can be expressed by the Hall-Petch relationship. Thus, the lamellar spacing needs to be maintained at a suitable level, because a decrease in the lamellar spacing leads to an increase in strength and ductility.

If the lamellar spacing is less than 150 nm, the strain hardening rate of the wire rod will be excessively increased to cause wire breakage during wire drawing. On the other hand, if the lamellar spacing is more than 300 nm, shear failures, such as cleavage fractures, will be highly likely to occur, and it will be difficult to ensure the strength described below.

Also, the content of diffusible hydrogen in the wire rod of the present invention is preferably limited to 0.6˜0.9 ppm. The term content of diffusible hydrogen refers to the highest concentration at which steel can contain hydrogen. The content of diffusible hydrogen varies depending on a matrix structure. If the content of diffusible hydrogen in the steel of the present invention is less than 0.6 ppm, the effect of improving resistance to delayed fracture by hydrogen trapping cannot be obtained. The reason for which the content of diffusible hydrogen in the steel wire rod of the present invention is limited to 0.9 ppm is that it is not easy to ensure a diffusible hydrogen content of more than 9 ppm in pearlite steels containing V precipitates, as in the case of the present invention.

Hereinafter, the method for manufacturing the steel wire rod of the present invention will be described in detail.

First, a steel satisfying the above-described composition is heated before rolling. Herein, the heating temperature is 1100° C. or lower, and preferably 1000˜1100° C.

The heated steel is subjected to hot rolling. Herein, a process ranging from rough rolling to finish rolling is carried out at a temperature of 1050˜800° C.

The rolled steel is cooled to 650˜600° C. at a rate of 5˜10° C./s. If the cooling rate is less than 5° C., proeutectoid cementite will be precipitated to cause anisotropy, and if the cooling rate is more than 10° C./s, martensite, a low-temperature structure, will be formed. The steel cooled after hot rolling has a tensile strength of 1100˜1300 MPa. After the cooling process, the steel is subjected to cold drawing. The cold drawing is preferably carried out at a reduction ratio of 60˜80%. In order to ensure the tensile strength of the steel by work hardening through the cold drawing process, the steel is cold drawn at a reduction ratio of 60% or higher. If the reduction ratio is higher than 80%, the cold forgeability of the steel will be deteriorated. For this reason, the upper limit of the reduction ratio is preferably 80%. The cold-drawn wire rod has a tensile strength of 1550˜1650 MPa.

Hereinafter, the present invention will be described in detail with reference to examples, but the scope of the present invention is not limited to these examples.

EXAMPLES

Each of steels satisfying the compositions shown in Table 1 below was heated at 1100° C., after which a strain of 0.6 was applied thereto at a strain rate of 10/s at 950° C. Then, the steels were cooled at a rate of 7.5° C./s and drawn to 10˜90%, thereby manufacturing wire rods. In Table 1 below, the inventive examples are steels to which V has been added within the content range specified in the present invention, the conventional example is a steel to which Cr had been added. Meanwhile, comparative examples 1 and 2 are steels that are out of the V content of the present invention, and comparative examples 3 and 4 are steels to which Al had been added in place of V.

The wire rods manufactured as described above were measured for tensile strength, elongation and surface roughness, and the results of the measurement are shown in Table 2 below. In addition, the manufactured wire rods were measured for microstructure, fracture stress according to diffusible hydrogen content, and the change in tensile strength according to the amount of drawing, and the results of the measurement are shown in FIGS. 1 to 6.

	TABLE 1
	C	Si	Mn	Cr	Al	V
Conventional example	0.82	0.25	0.8	0.2	—	—
Inventive example 1	0.82	0.25	0.8	—	—	0.05
Inventive example 2	0.82	0.25	0.8	—	—	0.1
Comparative example 1	0.82	0.25	0.8	—	—	0.15
Comparative example 2	0.82	0.25	0.8	—	—	0.2
Comparative example 3	0.82	0.25	0.8	—	0.04	—
Comparative example 4	0.82	0.25	0.8	—	0.08	—

	TABLE 2
			Surface
	Tensile strength	Elongation	roughness
	(TS, MPa)	(El, %)	(RA, %)
Conventional example	1051	7	—
Inventive example 1	1062.8	7.5	16.7
Inventive example 2	1106.8	6.7	12.2
Comparative example 1	1070.3	7.3	15.2
Comparative example 2	1146.5	5.1	14
Comparative example 3	990.9	6.8	14.5
Comparative example 4	1027.8	7.4	15.1

As can be seen in Table 2 above, inventive examples 1 and 2 could achieve excellent strength, elongation and surface roughness (RA), which were comparable to the conventional example and comparative examples 3 and 4, even when a trace amount of V was added thereto. Also, as can be seen in comparative examples 1 and 2, even when V was added in an amount higher than the upper limit of the content range specified in the present invention, strength and elongation were not improved. For this reason, the content of V was limited to 0.05˜0.1% in view of diffusible hydrogen content and fracture stress.

Meanwhile, FIGS. 1A, 1B, 2A and 2B are photographs showing the results of observing the microstructures of the conventional example and inventive example 1, respectively. As shown in FIGS. 1 and 2, the lamellar spacing of pearlite in inventive example 1 (about 184.3 nm) was about 50% smaller than the conventional example (about 368.75 nm).

FIGS. 3A through 3F show the results of observing the microstructures of the conventional example, inventive example 1, inventive example 2, comparative example 1, comparative example 2 and comparative example 4, respectively. As shown in FIG. 3. inventive examples 1 and 2 had lamellar spacing gaps of 184.3 nm and 213 nm, respectively, which were smaller than the conventional example and the comparative examples.

FIGS. 4A and 4B schematically show the microstructures of the conventional example and inventive example 1, respectively. As can be seen therein, in the conventional example, film-like cementite was precipitated, but in inventive example 1, spherical V(C,N) precipitates were distributed.

FIG. 5 shows the results of comparing hydrogen delayed-fracture resistance between inventive example 1 and the conventional example (82BC steel). As can be seen therein, at a fracture stress of about 1500 MPa, inventive example 1 had a diffusible hydrogen content of 0.87 ppm which was about 1.5 times higher than the conventional example (diffusible hydrogen content: 0.52 ppm). This is because the spherical V(C,N) precipitates shown in FIG. 4 were segregated in the ferrite-cementite of prior pearlite, whereby diffusible hydrogen was trapped by the spherical V(C,N) precipitates, thus improving the hydrogen delayed-fracture resistance of the wire rod.

FIG. 6 is a graphic diagram showing a change in tensile strength according to a decrease in diameter in various drawing amounts. As can be seen in Table 2 above and FIG. 6, the case in which V was added showed excellent tensile strength. If V is added in an amount of more than 0.1%, the increase in tensile strength caused by the formation of V precipitates can be expected, but in comparative examples 1 and 2 in which V was added in large amounts, the wire rods could be cracked or fractured due to the formation of excessive V precipitates.

In comparative examples 3 and 4 in which Al was added, the colony size of the wire rods was advantageously decreased due to grain refinement, Al precipitates were not easily formed due to high cooling rate, and thus it was not easy to correct the strength of the wire rods.

While the present invention has been shown and described in connection with the exemplary embodiments, it will be apparent to those skilled in the art that modifications and variations can be made without departing from the spirit and scope of the invention as defined by the appended claims.

Claims

1. An ultra-high-strength steel wire rod having excellent resistance to delayed fracture, the wire rod comprising, by wt %, 0.7-1.2% C, 0.25-0.5% Si, 0.5-0.8% Mn, 0.02-0.1% V and a balance of Fe and inevitable impurities.

2. The ultra-high-strength steel wire rod of claim 1, wherein the wire rod has a pearlite structure having a lamellar spacing of 150˜300 nm, in which V(C,N) precipitates are distributed in the ferrite structure of the pearlite.

3. The ultra-high-strength steel wire rod of claim 2, wherein the average particle size of the V(C,N) precipitates is 30 nm or less, and the number of the V(C,N) precipitates is 1×109/mm2 or more.

4. The ultra-high-strength steel wire rod of claim 1, wherein the wire rod has a diffusible hydrogen content 0.6˜0.9 ppm.

5. A method for manufacturing an ultra-high-strength steel wire rod having excellent resistance to delayed fracture, the method comprising:

heating a steel, which includes, by wt %, 0.7-1.2% C, 0.25-0.5% Si, 0.5-0.8% Mn, 0.02-0.1% V and a balance of Fe and inevitable impurities, to 1100° C. or lower, and hot-rolling the heated steel at a temperature of 900˜1000° C.;

cooling the rolled steel to 600˜650° C. to a rate of 5˜10° C./s; and

cold-drawing the cooled steel at a reduction ratio of 60˜80%.