Patent Application
20140150934
The present disclosure relates to a wire rod used for automobile engine bolts and the like, and more particularly, to a wire rod having improved hydrogen delayed fracture resistance, a method for manufacturing the same, a high strength bolt using the same, and a method for manufacturing the bolt.
in accordance with the recent trend for weight reductions and high functionalization in automobiles, requirements for driving parts, in particular, engine parts such as bolts and the like, to have high strength, have increased in order to reduce energy consumption. Currently used high strength bolts are manufactured to have a tensile strength of 1200 MPa through quenching and tempering processes, using alloyed steels such as SCM435, SCM440, and the like. However, in the bolts having a tensile strength of 1200 MPa or greater, since delayed fractures may be easily caused due to hydrogen, the uses of wire rods for manufacturing ultrahigh strength bolts remain inadequate.
According to a bolt manufacturing process, after performing wire drawing intended for sizing, through low temperature annealing, the drawn wire may be subjected to spheroidizing heat treatment, bolt-forming, quenching and tempering processes to finally obtain a steel having a single-phase structure composed of tempered martensite. Thus, strength of the bolt may be determined depending on composing, quenching, tempering and heat treatment processes performed thereon. However, the wire rod as a raw material needs to have as little strength as possible in order to facilitate bolt-forming.
In order to highly strengthen a steel having a single-phase structure composed of tempered martensite, the addition of alloying elements, in particular, carbon elements, has been known as the most effective method; however, the addition of carbon may rapidly increase a ductile to brittle transition temperature (DETT) of a wire rod as well as increasing strength of the wire rod, and remarkably deteriorate hydrogen delayed fracture resistance. In addition thereto, work hardening may be increased, causing disadvantages in bolt-forming and a separate softening heat treatment may be required.
Bolts manufactured as described above may generally have a tempered martensite structure in which carbide precipitates are distributed in grain boundaries or gains and the basic material thereof has precipitates distributed in lath martensite. A main factor hindering the high strengthendng of the basic material may be a degradation in delayed fracture resistance due to the introduction of hydrogen, and it has been known because the introduced hydrogen may deteriorate the strength of grain boundaries. In order to use existing tempered martensite in steel for high strength bolts, an operation for improving delayed fracture resistance may he required.
Thus, in order to achieve the high strengthening of bolts, improvements in delayed fracture resistance may be unavoidably required to increase critical delayed fracture strength, and to this end, a method of generating precipitates capable of trapping diffusible hydrogen or controlling microstructure by adding certain elements while maximally suppressing phosphorus (P) and sulfide (S) brominating austenitic grain boundaries, and the like may be present.
The related art technologies for improving hydrogen delayed fracture resistance may include 1) corrosion suppression in steel, 2) minimization of an amount of introduced hydrogen, 3) suppression of diffusible hydrogen contributing to delayed fracture, 4) the use of steel having a high concentration of limited diffusible hydrogen. contained therein, 5) minimization of tensile stress, 6) stress concentration reduction, 7) miniaturization of austenite grain boundary size, and the like. As a method of achieving improvements in hydrogen delayed fracture resistance, a method of implementing a high degree of alloying, or a surface coating method or a plating method for preventing the introduction of external hydrogen has been mainly used.
However, most inventions created domestically and internationally may have disadvantages such as high manufacturing costs and complex processes required therefor, and require excessively precise rolling and cooling conditions at the time of manufacturing steel. By way of example, in order to improve delayed fracture characteristics of a high strength wire rod having a tensile strength of 1600 MPa, technologies of adding 0.5 wt % of titanium (Ti), niobium (Nb), and vanadium (V), which are grain refinement elements, and then, adding corrosion resistance elements such as molybdenum (Mo), nickel (Ni), copper (Cu), cobalt (Co), and the like and carbide elements are present, but production costs required therefor may be significantly high, Furthermore, a method of improving hydrogen brittleness using ferrite structures extracted from grain boundaries is present, but the method does not include a chemical combination, and a product manufacturing cost may also increase due to the addition of a considerable amount of molybdenum (Mo).
In addition, a technology of improving delayed fracture characteristics of a high strength wire rod having a tensile strength of 1600 MPa or greater, using complete pearlite is present. However, in such a technology, 0.2 wt % or more of chrome needs to be added in order to improve tensile strength through wire drawing and to secure drawability during wire drawing intended for sizing after the production of a wire rod, and lead patenting for isothermal transformation may necessarily be required. Thus, such a technology may have disadvantages such as high manufacturing costs and complex processes and have limitations such as the requirement for excessively precise rolling and cooling conditions at the time of manufacturing steel.
Moreover, through a technology of finally securing a tensile strength of 1200 to 1500 MPa using a ferrite-pearlite dual phase microstructure, the tensile strength may be secured without a final heat treatment, unlike in other technologies. However, since the technology basically aims at improving hydrogen delayed fracture resistance by adding a great quantity of molybdenum (Mo), it may be disadvantageous in terms of high manufacturing costs.
As described above, limitations to a decrease in hydrogen delayed fracture resistance as compared to an improvement in tensile strength in heat-treated and non heat-treated carbon steels having a tensile strength of 1200 MPa or greater have not yet been overcome, the securing of price competitiveness may not be available due to the addition of expensive alloying elements, and in particular, the stable securing of data regarding delayed fracture characteristics due to hydrogen may be defective.
An aspect of the present disclosure provides a wire rod having superior hydrogen delayed fracture resistance while securing ultrahigh strength through a heat treatment, and a method for manufacturing the same.
An aspect of the present disclosure also provides a high strength bolt having superior hydrogen delayed fracture resistance using the wire rod, and a method for manufacturing the same.
According to an aspect of the present disclosure, there is provided a wire rod having superior hydrogen delayed fracture resistance and including C: 0.3 to 0.7 wt %, Si: 0.05 to 2.0 wt %, Mn: 0.7 to 1.5 wt %, La: 30 to 70 ppm, Ni: 0.01 to 0.1%, and a remainder configured of Fe and inevitable impurities.
According to another aspect of the present disclosure, there is provided a method for manufacturing a wire rod having superior hydrogen delayed fracture resistance, the method including: heating steel including C: 0.3 to 0.7 wt %, Si: 0.05 to 2.0 wt %, Mn: 0.7 to 1.5 wt %, La: 30 to 70 ppm, Ni: 0.01 to 0.1%, and remainder configured of Fe and inevitable impurities to a temperature of Ae3+150° C. to Ae3+250° C.; cooling the heated steel at a rate of 5 to 15° C/s and rolling the steel at a temperature of Ae3+50° C. to Ae3+150° C. to manufacture a wire rod; and cooling the rolled wire rod to 600° C. or less at a rate of 0.5 to 3° C./s.
According to another aspect of the present disclosure, there is provided a bolt including C: 0.3 to 0.7 wt %, Si: 0.05 to 2.0 wt %, 0.7 to 1.5 wt %, La: 30 to 70 ppm, Ni: 0.01 to 0.1%, and a remainder configured of Fe and inevitable impurities, and having a tensile strength of 1200 MPa or greater and superior hydrogen delayed fracture resistance.
According to another aspect of the present disclosure, there is provided a method for manufacturing a bolt having superior hydrogen delayed fracture resistance, the method including: heating steel including C: 0.3 to 0.7 wt %, Si: 0.05 to 2.0 wt %, Mn: 0.7 to 1.5 wt %, La: 30 to 70 ppm, Ni: 0.01 to 0.1%, and a remainder configured of Fe and inevitable impurities to a temperature of Ae3+150° C. to Ae3+250° C.; cooling the heated steel at a rate of 5 to 15° C./s and rolling the steel at a temperature of Ae3+50° C. to Ae3+150° C. to manufacture a wire rod; cooling the rolled wire rod to 600° C. or less at a rate of 0.5 to 3° C./s; and bolt-forming using the cooled wire rod; performing a heat treatment on the formed bolt at a temperature of 850 to 950° C.; and performing quenching after the heat treatment, and then performing tempering at a temperature of 300 to 500° C.
The wire rod according to the present disclosure may be a high strength wire rod used for the coupling of automobile components or used in such automobile components, and the method of manufacturing the wire rod may be advantageous in that a wire rod having high strength of 1200 MPa to 2000 MPa and superior hydrogen delayed fracture resistance, even in a case in which a tiny amount of lanthanum and nickel is added or even in a case in which a martensite microstructure is present after the final heat treatment, may be manufactured with low manufacturing costs.
In accordance with the development of a wire rod for bolts having superior hydrogen delayed fracture resistance and high strength, the stability of a steel structure may be increased due to a reinforcement of coupling force and a reduction of vacancies in a coupling part at the time of coupling the bolts, and an amount of steel used may be reduced due to a decrease in the number of coupled bolts. In addition, in terms of automobile components, the development of the wire rod for bolts as described above may contribute to lightening of the automobile components. Due to the lightening of automobile components, various automobile assembling device designs may be enabled and compactness of automobile assembling devices may be allowed.
Hereinafter, exemplary embodiments of the present disclosure will be described in detail.
First, a wire rod according to an exemplary embodiment of the present disclosure will be described in detail. Here, a compositional range of the wire rod according to the exemplary embodiment of the present disclosure will be described (hereinafter, referred to as weight percentage (wt %)).
Carbon (C) may be included in the wire rod in an amount of 0.3 to 0.7 wt %, When carbon (C) is included in an amount exceeding 0.7 wt %, although the wire rod may be frequently used in the form of a high carbon wire rod formed using common cold wire drawing, in a case in which the wire rod is subjected to a heat treatment suggested in the exemplary embodiment of the present disclosure, film shaped carbides may be frequently eluted in austenite grain boundaries to thereby deteriorate hydrogen delayed fracture resistance. Thus, an amount of carbon (C) exceeding 0.7 wt % may not be preferable. Meanwhile, when carbon, (C) is included in an amount less than 0.3 wt %, since tensile strength of a bolt may be insufficiently secured through quenching and tempering heat treatments, carbon (C) may be added in an amount of 0.3 wt % or greater in order to secure a sufficient degree of strength.
Silicon (Si) may be included in the wire rod in an amount of 0.05 to 2.0 wt %. When silicon (Si) is included in an amount exceeding 2.0 wt %, a work hardening phenomenon may be rapidly generated during a cold forging process for manufacturing bolts to deteriorate processability. When silicon (Si) is included in an amount less than 0.05 wt %, a sufficient degree of strength may not be secured and spheroidization of cementite may also he adversely affected.
Manganese (Mn) may be included in the wire rod in an amount of 0.7 to 1.5 wt %. Manganese (Mn), an element forming a substitutional solid solution in a base structure to perform solid solution reinforcement, may be very useful in high tension bolt characteristics. When manganese (Mn) is included in an amount exceeding 1.5 wt %, a heterogeneous structure caused by manganese segregation may have a negative influence on bolt characteristics, rather than having solid solution reinforcement effects. That is, during the coagulation of steel, macroscopic and microscopic segregation may be easily generated according to a segregating device, and manganese (Mn) may aggravate a segregation area due to the diffusion coefficient thereof relatively being lower than that of other elements and the consequent hardenability improvements may be a main factor generating a core low temperature structure (for example, core martensite). That is, an increase in local quenching properties and the formation of a segregation area caused by manganese segregation during casting may intensify dual phase properties of the structure.
Meanwhile, when manganese (Mn) is included in an amount less than 0.7 wt %, the segregation area may be barely affected by the manganese segregation, but tensile strength of a final product may not be secured through solid solution reinforcement. That is, when manganese (Mn) is included in an amount less than 0.7 wt %, improvements in quenching and permanent deformation resistance may he insufficient due to insufficient solid solution reinforcement,
Nickel (Ni) may be included in the wire rod in an amount of 0.01 to 0.1 wt %. Nickel (Ni) may be a very important element forming a compound within a grain boundary, together with lanthanum (La). Thus, when nickel (Ni) is included in an amount less than 0.01 wt %, an effective compound, in particular, precipitates, may not be completely generated, thereby leading to an inability to improve hydrogen delayed fracture resistance, When nickel (Ni) is included in an amount exceeding 0.1 wt. %, the amount of the remaining austenite may be increased to degrade impact toughness and manufacturing costs may be increased due to an excessive amount of nickel.
Lanthanum (La) may be included in the wire rod in an amount of 0.003 to 0.007 wt % (30-70 ppm), Lanthanum (La) may be a very important element forming a compound within a grain boundary, together with Nickel (Ni) and decreasing phosphorous and sulfur segregated in the grain boundary. Thus, when lanthanum (La) is included in an amount less than 30 ppm, the compound may not he effectively formed and the removal of phosphorus and sulfur segregated in the grain boundary may not be facilitated. Thus, the securing of tensile strength may be enabled but superior hydrogen delayed fracture resistance may not be expected. On the other hand, when lanthanum (La) is included in an amount exceeding 70 ppm, since manufacturing costs may be increased and the hydrogen delayed fracture resistance may not be improved due to an excessive amount of lanthanum, the upper limit of the amount of added lanthanum may be 70 ppm,
The remainder may include iron (Fe) and inevitable impurities. In addition to the composition described above, the addition of effective elements may, not be excluded.
The wire rod according to the exemplary embodiment of the present invention may include a lanthanum (La)-based, a nickel (Ni)-based, or a LaNi-based precipitate. Types of the precipitate are not particularly limited, but examples thereof may include LaNi5, LaPO4 La2O2S and the like. The precipitate may he formed in a grain or a grain boundary of a microstructure and trap hydrogen introduced into the grain or the grain boundary to prevent the introduced hydrogen from deteriorating strength of the grain boundary, thereby improving hydrogen delayed fracture resistance.
Meanwhile, hydrogen trapping effects due to the precipitates according to the exemplary embodiment of the present disclosure may be significantly superior, as compared to the related art technology intended to improve hydrogen delayed fracture resistance through molybdenum (Mc).
An aspect ratio of the precipitate may be 1.2 to 2.0. When the aspect ratio of the precipitate is less than 1.2, the securing of the compound may rarely be allowed due to the crystal structure. When the aspect ratio of the precipitate exceeds 2.0, the precipitate may be easily broken. In a case in which the precipitate is broken in a material, continuity thereof with a base may be deficient and micro-voids may be generated, thereby causing defects, Thus, breakage of the wire rod may be caused and expected hydrogen delayed fracture resistance may not be secured.
Meanwhile, in terms of a size of the precipitate, a circular-equivalent diameter of the precipitate may be 100 to 400 nm. When the diameter is less than 100 nm, the size of the precipitate may be excessively small, an amount of hydrogen trapped in the precipitate may be reduced, whereby effective hydrogen trapping effects may not be secured, When the diameter exceeds 400 nm, and is significantly large, since the number of precipitates distributed per unit area may be reduced, a decrease in a surface area of the precipitates in the overall steel may result, thereby reducing hydrogen trapping effects, the upper limit of the diameter of the precipitate may be 400 nm.
Hereinafter, a method for manufacturing a wire rod according to an exemplary embodiment of the present disclosure will be described in detail.
In order to manufacture the wire rod according to the exemplary embodiment of the present disclosure, steel satisfying the composition described above may he heated to a temperature of Ae3+150° C. to Ae3+250° C. The heating to the temperature may be intended to maintain an austenite single phase, and in a range of the temperature, austenite grain coarsening may not be generated and the remaining segregation, carbides and inclusions may be effectively dissolved. When the temperature exceeds Ae3+250° C., an austenite grain may be significantly coarse, such that a final microstructure formed after cooling may be highly coarse, resulting in an inability to secure a high strength wire rod having a high degree of toughness. Meanwhile, when the heating temperature is less than Ae3+150° C., heating effects may not be obtained and consequently, the heating temperature may be Ae3+150° C. to Ae3+250° C.
The heating may be undertaken for 30 minutes to one and a half hours. When the heating is performed for less than 30 minutes, the entire temperature may not be uniform. When the heating is performed for more than one and a half hours, possibility that the austenite grain may be coarse may be higher and productivity may be significantly reduced.
The heated steel may be cooled and be subjected to hot rolling. The cooling may be performed at a cooling rate of 5 to 15° C./s and the rolling may be performed at a temperature of Ae3+50° C. to Ae3+150° C., to thereby manufacture a wire rod.
The cooling may be intended to perform controlling aiming at minimizing the transformation of the microstructure. When the cooling rate is less than 5° C./s before the rolling, productivity may be decreased, an additional device may be required in order to maintain a slow cooling rate, and further, strength and toughness of the wire rod may be deteriorated after the hot rolling, similarly to the case in which the heating is maintained for long hours, On the other hand, when the cooling rate exceeds 15° C./s, since driving force of the transformation in steel before the rolling may be increased, the possibility that a new microstructure may emerge during the rolling may be increased, such that a lower rolling temperature may need to be reset.
In addition, the rolling temperature may be a temperature at which the emergence of a microstructure caused by the transformation during the rolling may be inhibited, recrystallization may not be generated, and only sizing rolling may be enabled. When the rolling temperature is less than Ae3+50° C., it may be close to the dynamic recrystallization temperature, such that the securing of the microstructure may be unavailable and general soft ferrite may be highly secured. Meanwhile, when the rolling temperature is greater than Ae3+150° C., since reheating may be required after the cooling, the upper limit of the rolling temperature may be set as described above.
The wire rod manufactured through the rolling as described above may be cooled to 600° C. or less at a cooling rate of 0.5 to 3° C./s. The cooling rate may refer to a cooling rate at which the diffusion of carbon may be suppressed by the addition of manganese, and the wire rod may be effectively generated while pearlite is incompletely generated and a sufficient area fraction is secured. When the cooling rate is less than 0.5° C./s, the cooling rate may be extremely low, thereby degrading productivity to a degree to which actual work becomes infeasible. When the cooling rate exceeds 3° C./s, hardenability may be improved due to overlapping effects of the added elements, such that ferrite-pearlite transformation may be delayed and a low temperature structure such as martensite or bainite may be generated,
Hereinafter, a bolt according to an exemplary embodiment of the present disclosure and a method for manufacturing the same will be described in detail.
The bolt manufactured using the wire rod according to the embodiment of the present disclosure may have ultrahigh strength and at the same time, may have superior hydrogen delayed fracture resistance due to the precipitate. The bolt according to the exemplary embodiment of the present disclosure may have ultrahigh strength of 1200 MPa or greater and at the same time, may have superior hydrogen delayed fracture resistance.
The manufacturing method of a wire rod having tensile strength of 1200 MPa or greater may be performed according to the following operations. First, bolt-forming may be performed using the wire rod according to the embodiment of the present disclosure, and a heat treatment may be performed on the formed bolt at a temperature of 850 to 950° C. The heat treatment may be intended to achieve homogenization of the structure through austenizing. When the temperature is less than 850° C., a sufficient amount of homogenization may not be performed, while when the temperature is greater than 950° C., no further effects derived from an increase in temperature may be secured and ductility may be deteriorated due to the coarsening of grains. Thus, the upper limit of the temperature may be 950° C.
After tine heat treatment, quenching may be performed and tempering may be undertaken at a temperature of 300 to 500° C. The structure homogenized through rapid cooling may form a low temperature transformation structure such as a martensite structure to thereby improve strength of the bolt.
The tempering may be intended to control strength and improve brittleness by removing residual stress generated due to the rapid cooling. When the temperature is less than 300° C., sufficient removal of residual stress may be difficult and rather, brittleness may be generated as a temper brittleness phenomenon. Thus, the temperature may be 300° C. or greater. When the temperature exceeds 500° C., the strength may be reduced due to an excessive heat treatment, thereby leading to an inability to secure a required level of strength. Thus, the tempering may be undertaken at a temperature of 300 to 500° C.
The method for manufacturing the bolt may be intended to secure a required level of strength by applying a common heat treatment thereto. The common heat treatment may be applied by controlling time and temperature in order to secure strength required by a person having ordinary skill in the art and the present disclosure is not particularly limited thereto.
Hereinafter, examples of the present disclosure will be described in detail. The following examples are merely provided for understanding of the present disclosure, and the present disclosure is not limited thereto.
Steels having compositions of Table 1 and Ae3 temperature were manufactured and then, wire rods were manufactured using the steels under conditions of Table 2. Bolts were manufactured using the wire rods manufactured as above. In this case, heat treatment conditions in a manufacturing process of the bolts were described in Table 2.
Tensile strength and hydrogen delayed fracture resistance of the respective bolts manufactured above were measured and the results thereof are shown in Table 3. The hydrogen delayed fracture resistance of the respective bolts were measured in such a manner that tensile strengths corresponding to about 0.9 times those of tensile strength of the respective bolts, measured in a state in which the respective bolts were immersed in a test solution having an acidity of about 2 and configured of H2O: 2000 cc, CH3COOH: 80 ml, and NaCl: 100 g were imparted to the bolts and then, hours after which the respective specimens were broken, were measured. Through the test, in a case in which the specimen was maintained unbroken for 100 hours or more, it was estimated that resistance to hydrogen delayed fracture was excellent.
In a case in which conditions of the present disclosure were satisfied, at the time of manufacturing the bolt, it could be confirmed that the bolt had high strength of 1200 MPa or greater, while having superior hydrogen delayed fracture resistance, However, comparative examples 9 and 10 were classified as comparative examples because they had sufficient strength and hydrogen delayed fracture resistance, but were not preferable in terms of economical feasibility due to the addition of an excessive amount of La and Ni.
Meanwhile, in a case in which carbon (C) was included in an excessively low amount as in comparative example 1, it could be confirmed that a sufficient amount of strength was not secured, and in a case in which carbon (C) was included in an excessively high amount as in comparative example 2, it could he confirmed that hydrogen delayed fracture resistance was significantly low. In the cases of comparative examples 3 to 5 having molybdenum (Mo) added therein, it could be confirmed that breakage occurred before 100 hours, the securing of sufficient hydrogen delayed fracture resistance was difficult. In the case in which only one of La and Ni was added, as in comparative examples 6 and 7, sufficient hydrogen delayed fracture resistance was not secured.
in a case in which the addition of La or Ni does not reach a set range of the present disclosure, as in examples 8 and 9, it could be confirmed that sufficient hydrogen delayed fracture resistance was not secured.
In order to determine hydrogen delayed fracture resistance depending on a size and an aspect ratio of a lanthanum (La)-based, a nickel (Ni)-based, or a LaNi-based precipitate, the size and the aspect ratio of the precipitate were varied through a heat treatment in the cases of inventive examples 1 to 3.
After the size and the aspect ratio of the precipitate were varied as described above, hydrogen delayed fracture resistance was measured in the same manner as that of the foregoing example 1 and the results thereof were shown in Table 4.
As can be seen in Table 4, it could he confirmed that when the aspect ratio of the precipitate was outside the range of the present disclosure, hydrogen delayed fracture resistance was low.
| Number | Date | Country | Kind |
|---|---|---|---|
| 10-2011-0070206 | Jul 2011 | KR | national |
| Filing Document | Filing Date | Country | Kind | 371c Date |
|---|---|---|---|---|
| PCT/KR2012/003757 | 5/14/2012 | WO | 00 | 1/14/2014 |
