Patent Application
20100218859
Embodiments of the present invention relate to a high carbon steel sheet and a method of manufacturing the high carbon steel sheet. More particularly, embodiments of the present invention relate to a high carbon steel sheet that is superior in fatigue life and a method of manufacturing the high carbon steel sheet.
Recently, a high degree of safety of an automobile has been required. Thus, fatigue life of a spring is also required to manufacture a safe automobile. The fatigue life of the spring is relatively long when the spring operates in an elastic deformation region. As a result, to increase the fatigue life of the spring, yield strength of a material included in the spring is increased to increase the elastic deformation region of the spring. Thereafter, the spring preferably operates within the elastic deformation region.
However, the operation region of the spring is sometimes included in a plastic deformation region according to the kind of the spring. In this case, a crack may be rapidly generated rather than a case where the spring operates in the elastic deformation region. Here, the fatigue life is mainly determined by growth of the crack rather than formation of the crack. Thus, in the case of a plate-type band spring having an operation region included in the plastic deformation region, the plate-type band spring preferably includes a fine structure that is capable of preventing the growth of the crack in order to increase the fatigue life of the spring.
A conventional steel sheet used to form a spring includes pearlite. However, laminar carbide is relatively large and yield strength thereof is not large. Thus, the fatigue life is not long.
A method of increasing the fatigue life by enlarging the yield strength of the spring has been suggested. In the method, the yield strength of the spring is increased by transforming a fine structure into bainite at a relatively low temperature before performing a cold rolling process to a steel sheet. However, in the case that the plastic deformation region is included in an operation region of the spring, the fatigue life decreases in spite of an increase in the yield strength due to the formed and mixed bainite.
Embodiments of the present invention provide a high carbon steel sheet that is superior in terms of fatigue life.
Embodiments of the present invention provide a method of manufacturing the high carbon steel sheet.
In accordance with embodiments of the present invention, a high carbon steel sheet includes about 0.75 wt % to about 0.95 wt % of carbon, smaller than about 1.8 wt % of silicon, about 0.1 wt % to about 1.5 wt % of manganese, about 0.1 wt %˜1.0 wt % of chromium, smaller than about 0.02 wt % of phosphorus, smaller than about 0.02 wt % of sulfur, a residual amount of iron, and inevitable impurities. A layer interval between laminar carbides included in the high carbon steel sheet is smaller than about 0.5 μm.
The high carbon steel sheet may include a fine pearlite having a lamellar structure. The fine pearlite included in the high carbon steel sheet may have a volume percentage of larger than about 90%. A ratio of length to width of the lamellar structure may be larger than about 10:1
The high carbon steel sheet may further include about 0.05 wt % to about 0.25 wt % of vanadium, niobium, molybdenum, titanium, tungsten, or copper. These may be used alone or in combination. The high carbon steel sheet may further include about 30 ppm to about 120 ppm of nitrogen.
In accordance with embodiments of the present invention, a method of manufacturing a high carbon steel sheet is provided. In the method, a steel member including about 0.75 wt % to about 0.95 wt % of carbon, smaller than about 1.8 wt % of silicon, about 0.1 wt % to about 1.5 wt % of manganese, about 0.1 wt % to about 1.0 wt % of chromium, smaller than about 0.02 wt % of phosphorus, smaller than about 0.02 wt % of sulfur, a residual amount of iron, and inevitable impurities is formed. A hot rolling process, a cold rolling process, and an annealing process are performed to allow the steel member to have a spheroidized cementite and an initial ferrite. A patenting annealing process is performed to the heated steel member after heating the steel member. The patenting annealing process is performed using a solder pot having a maintained temperature of about 500° C. to about 530° C. for over about 60 seconds.
The steel member may be heated before the patenting annealing process at a temperature of about 800° C. to about 1100° C. The steel member further may include about 0.05 wt % to about 0.25 wt % of at least one selected from the group consisting of vanadium, niobium, molybdenum, titanium, tungsten, and copper. The steel member may further include about 30 ppm to about 120 ppm of nitrogen.
To manufacture the high carbon steel sheet, a cooling process may be performed after the patenting annealing process. A cold rolling process may be the performed such that a reduction ratio of the cold rolling process is over about 85%.
In accordance with embodiments of the present invention, a method of manufacturing a high carbon steel sheet is provided. In the method, a steel member including about 0.75 wt % to about 0.95 wt % of carbon, smaller than about 1.8 wt % of silicon, about 0.1 wt % to about 1.5 wt % of manganese, about 0.1 wt % to about 1.0 wt % of chromium, smaller than about 0.02 wt % of phosphorus, smaller than about 0.02 wt % of sulfur, a residual amount of iron, and inevitable impurities is formed. A hot rolling process, a cold rolling process and an annealing process are performed to allow the steel member to have a spheroidized cementite and an initial ferrite. A patenting annealing process is performed to the heated steel member after heating the steel member, the patenting annealing process being performed using a solder pot having a maintained temperature of about 500° C. to about 530° C. for over about 20 seconds.
The steel member may be heated before the patenting annealing process at a temperature of about 800° C. to about 1100° C. The steel member may further include about 0.05 wt % to about 0.25 wt % of vanadium, niobium, molybdenum, titanium, tungsten, or copper. These may be used alone or in combination. The steel member may further include about 30 ppm to about 120 ppm of nitrogen.
To manufacture the high carbon steel sheet, a cooling process may be performed after the patenting annealing process. A cold rolling process may be performed such that a reduction ratio of the cold rolling process is over about 85%.
A high carbon steel sheet according to the embodiments of the present invention has a large amount of laminar carbide having a relatively large ratio of length to width of over about 10:1. Thus, growth of a crack may be efficiently prevented.
In addition, fatigue life of the high carbon steel sheet may be improved because the growth of the crack is prevented.
Further, in the case that a spring is formed using the high carbon steel sheet according to the embodiments of the present invention, a generated crack may not easily grow. Thus, the spring may operate for a relatively long time even though the spring operates in a plastic deformation region.
In accordance with embodiments of the present invention, a high carbon steel sheet includes about 0.75 wt % to about 0.95 wt % of carbon, smaller than about 1.8 wt % of silicon, about 0.1 wt % to about 1.5 wt % of manganese, about 0.1 wt %˜1.0 wt % of chromium, smaller than about 0.02 wt % of phosphorus, smaller than about 0.02 wt % of sulfur, a residual amount of iron, and inevitable impurities.
In addition, the high carbon steel sheet may further include about 0.05 wt % to about 0.25 wt % of vanadium, niobium, molybdenum, titanium, tungsten, or copper. These may be used alone or in combination. The high carbon steel sheet may further include about 30 ppm to about 120 ppm of nitrogen.
Here, a layer interval between laminar carbides included in the high carbon steel sheet is smaller than about 0.5 μm. The high carbon steel sheet may include a fine pearlite having a lamellar structure. The fine pearlite included in the high carbon steel sheet may have a volume percentage of larger than about 90%. A ratio of length to width of the lamellar structure may be larger than about 10:1.
In accordance with embodiments of the present invention, a method of manufacturing a high carbon steel sheet is provided. In the method: i) a steel member including about 0.75 wt % to about 0.95 wt % of carbon, smaller than about 1.8 wt % of silicon, about 0.1 wt % to about 1.5 wt % of manganese, about 0.1 wt % to about 1.0 wt % of chromium, smaller than about 0.02 wt % of phosphorus, smaller than about 0.02 wt % of sulfur, a residual amount of iron, and inevitable impurities is formed; ii) a hot rolling process, a cold rolling process and an annealing process is performed to allow the steel member to have a spheroidized cementite and an initial ferrite; and iii) a patenting annealing process is performed to the heated steel member after heating the steel member at a temperature of about 800° C. to about 1100° C. The patenting annealing process is performed using a solder pot having a maintained temperature of about 500° C. to about 530° C. for over about 60 seconds. Alternatively, the patenting annealing process is performed using a solder pot having a maintained temperature of about 530° C. to about 570° C. for about 20 seconds.
The steel member further may include about 0.05 wt % to about 0.25 wt % of vanadium, niobium, molybdenum, titanium, tungsten, or copper. These may be used alone or in combination. The steel member may further include about 30 ppm to about 120 ppm of nitrogen.
To manufacture the high carbon steel sheet, a cooling process and a cold rolling process may be performed after the patenting annealing process such that a reduction ratio of the cold rolling process is over about 85%.
Hereinafter, a chemical composition of the high carbon steel sheet is disclosed.
The content of carbon (C) is about 0.75 wt % to about 0.95 wt %. In the case that the content of carbon is lower than about 0.75 wt %, hardness may increase by a quenching process. Thus, it is difficult to obtain superior endurance. In case that the content of the carbon is over about 0.95 wt %, residing austenite may be easily formed. In addition, a crack may be generated by a stress induced transformation when a cold rolling process is performed. In addition, toughness and fatigue life of the steel sheet may be deteriorated.
The content of silicon (Si) is smaller than about 1.8 wt %. In the case that the content of silicon increases, strength and resistance with respect to plastic deformation may increase. However, in the case that the content of silicon is larger than about 1.8 wt %, the resistance with respect to the plastic deformation may decrease and decarburization may easily occur when a thermal treatment process is performed. In addition, quality of a surface may be deteriorated by an increase in scale defects.
The content of manganese (Mn) is about 0.1 wt % to about 1.5 wt %. In the case that the content of manganese (Mn) is lower than about 0.1 wt %, red brittleness may occur due to sulfur-iron (FeS) including sulfur (S) and iron (Fe) that are inevitably impurities. On the other hand, in the case that the content of manganese (Mn) is larger than about 1.5 wt %, toughness may decrease. In addition, in the case that the content of manganese (Mn) is larger than about 1.5 wt %, hardenability may exceedingly increase so that a proceeding velocity of the steel sheet is required to be reduced to obtain a fine structure. Thus, yield of the steel sheet may decrease.
The content of chromium (Cr) may be about 0.1 wt % to about 1.0 wt %. Chromium (Cr) has substantially the same effect as manganese (Mn). Particularly, chromium increases hardenability and strength. In addition, chromium may prevent decarburization and graphitization. In the case that the content of chromium (Cr) is lower than about 0.1 wt %, it is difficult to obtain sufficient hardenability. In addition, the decarburization may not be effectively prevented. On the other hand, in the case that the content of chromium (Cr) is larger than about 1.0 wt %, the hardenability may be exceedingly increased.
The content of sulfur (S) is smaller than about 0.02 wt %. In case that the content of sulfur (S) is larger than about 0.02 wt %, toughness may decrease due to grain boundary segregation.
The content of phosphorus (P) is smaller than about 0.02 wt %. In the case that the content of phosphorus is larger than about 0.02 wt %, toughness may decrease due to grain boundary segregation.
Molybdenum (Mo), niobium (Nb), titanium (Ti), vanadium (V), or tungsten (W) may be combined with carbon (C) or nitrogen (N) in the steel sheet to generate precipitation hardening. Copper (Cu) may independently generate the precipitation hardening. The content of molybdenum (Mo), niobium (Nb), titanium (Ti), vanadium (V), or tungsten (W) may be about 0.05 wt % to about 0.25 wt %. These may be used alone or in combination. The independent or interdependent precipitation hardening due to the above elements may increase strength of the steel sheet. However, in the case that the above elements are exceedingly included, a rolling property may decrease due to an exceedingly increased hardenability because an effect of the above elements tends to be saturated. Thus, it is desired to use the above elements selectively. In the case that the content of the above element is smaller than about 0.05 wt %, an effect of the precipitation hardening may decrease. On the other hand, in the case that the content of the above element is larger than about 0.25 wt %, brittleness of the steel sheet may increase when a hot rolling process is performed.
The content of nitrogen (N) is about 30 ppm to about 120 ppm. In the case that the content of nitrogen is smaller than about 30 ppm, an educed amount of carbon nitride is not sufficient. Thus, improvement of strength and resistance with respect to plastic deformation is small. On the other hand, in the case that the content of nitrogen (N) is larger than about 120 ppm, an effect of the precipitation hardening may be saturated and an induced material is exceedingly saturated in a matrix. Thus, toughness of the steel sheet may decrease.
Hereinafter, a fine structure and an effect of preventing a crack from growing in a high carbon steel sheet are described with reference to
Intensity of the carbide is larger than that of a peripheral ferrite. Thus, the crack of the steel sheet may not grow through the carbide. As a result, the crack tends to grow between the laminar carbides 101. However, the high carbon steel sheet includes a high density of the laminar carbide having the relatively large ratio of the length to the width. Thus, the crack may not grow between the carbides. As a result, the crack formed at a surface of an edge may not easily grow because the crack has to grow along a complex “A” path “A.”
That is, the laminar carbide 101 in
Hereinafter, a method of manufacturing a high carbon steel sheet is described.
A steel member is formed. The steel member includes about 0.75 wt % to about 0.95 wt % of carbon (C), less than about 1.8 wt % of silicon (Si), about 0.1 wt % to about 1.5 wt % of manganese (Mn), about 0.1 wt % to about 1.0 wt % of chromium (Cr), less than about 0.02 wt % of phosphorus (P), less than about 0.02 wt % of sulfur (S), a residual amount of iron (Fe), and inevitable impurities. The steel member may further include a predetermined element. The predetermined element may be about 0.05 wt % to about 0.25 wt % of vanadium (V), about 0.05 wt % to about 0.25 wt % of niobium (Nb), about 0.05 wt % to about 0.25 wt % of molybdenum (Mo), about 0.05 wt % to about 0.25 wt % of titanium (Ti), about 0.05 wt % to about 0.25 wt % of tungsten (W), about 0.05 wt % to about 0.25 wt % of copper (Cu), or about 30 ppm to about 120 ppm of nitrogen (N). These may be used alone or in combination. A chemical composition of the steel sheet is previously described. Thus, any further explanation will be omitted.
A hot rolling process, a cold rolling process, and an annealing process are performed on the steel member so that a steel sheet having spheroidized cementite and ferrite may be formed. The steel sheet is then heated at a temperature of about 800° C. to about 1100° C. In the case that the steel sheet is heated at a temperature of lower than about 800° C., the spheroidized cementite may not be fully dissolved in a quenching process. Thus, intensity of a product formed using the steel sheet may decrease after thermal treatment. On the other hand, when the steel sheet is heated at a temperature of over about 1100° C., surface decarburization of a spring steel may occur. In addition, a grain size of an austenite phase may increase so that hardening may be exceedingly required. As a result, it is difficult to obtain a fine structure.
Thereafter, a patenting annealing process is performed on the steel sheet at a temperature, i.e., a temperature of a solder pot, maintained between about 500° C. to about 530° C. for about 60 seconds. Alternatively, the patenting annealing process is performed on the steel sheet at a temperature maintained between about 530° C. to about 570° C. for about 20 seconds. In the case that the temperature and the time required for the patenting annealing process are not adequate, the austenite formed in the quenching process may not be transformed into fine pearlite in the solder pot. In this case, the austenite may be transformed into martensite. Alternatively, the austenite may reside. The residing austenite may be changed into a stress-induced martensite and generate a crack, thereby reducing fatigue life in the cold rolling process. In addition, a rolling property may be deteriorated because the stress-induced martensite generates the crack in the cold rolling process. On the other hand, in the case that an annealing process is performed on the steel sheet at a temperature maintained over about 570° C., an interval between layers of the fine pearlite increases. Thus, it is difficult to form a fine structure that is capable of preventing growth of a fatigue crack by using a subsequent cold rolling process. In addition, intensity may not be easily improved by work hardening.
As described above, the high carbon steel sheet including a fine structure that is superior in fatigue life may be formed by adjusting conditions of the patenting annealing process.
Hereinafter, embodiments of the present invention may be more fully described with reference to examples. It is to be understood that the foregoing examples are illustrative of the present invention and are not to be construed as limited to the specific embodiments disclosed.
Experiment
A carbon steel sheet used for forming a high intensity spring and having the above-described composition was prepared. A rolling process was then performed on the high carbon steel sheet in which a spheroidization annealing process was performed so that a plate-shaped coil having a thickness of about 1.3 mm to about 1.6 mm was formed. Thereafter, the plate-shaped coil was heated at a quench temperature of about 750° C. to about 1200° C. for about 2 minutes. A patenting annealing process, i.e., an austempering process, was then performed using a solder pot having a temperature of about 300° C. to about 650° C. Thereafter, a plate-shaped coil having a thickness of about 0.23 mm was formed by a rolling process. As a result, a material having a uniform thickness in spite of a different reducing ratio in a cold rolling process was formed. The coil having the thickness of about 0.23 mm was slit to have a width of about 8 mm required for forming a spring. Thereafter, a bur removing process, a shape-forming process, a winding process, and a strain aging process were performed to form the spring.
Patenting (or austempering) annealing conditions, quenching annealing conditions, and reduction ratio required for forming a high carbon steel sheet superior in fatigue life for a spring are disclosed in [TABLE 1].
A fatigue test was performed to measure fatigue life of a spring. In the fatigue test, the spring was installed in a fatigue measuring device wherein rotation and reverse rotation were performed. In order to allow the spring to operate in a plastic deformation region, the rotation and the reverse rotation were repeated between a 2nd rotation region to a 22th rotation region until the spring was broken. The fatigue measuring device repeatedly measured the fatigue life, and the results are disclosed in [TABLE 1].
Referring to Table 1, in the case that the temperature of the quenching process was below about 800° C. (Comparative Example 1), a reverse transformation occurs in a matrix. However, the spheroidized cementite may not be completely dissolved. The cementite structure that is not dissolved may reside after the austempering annealing process. Thus, a concentration of carbon measured in the matrix may be insufficient so that the transformation curve may move forward. As a result, pearlite may be formed before injection to the solder pot even though the patenting (or austempering) annealing process is performed at the same cooling speed and some of the austenite may be transformed into fine pearlite while the rest of the austenite is maintained in the solder spot. In this case, the amount of the spheroidized cementite that is not dissolved is larger than that of strained laminar carbide. Thus, a fatigue crack is not effectively prevented and yield strength may be low. As a result, the fatigue life may be relatively short.
In the case that the quenching annealing temperature is over about 1100° C. (Comparative Example 2), the transformation curve may move back because of an increase in a grain size of austenite. Thus, a small amount of martensite may be formed from the austenite remaining in a cooling step after passing the solder spot. In addition, ferrite due to decarburization may be found at a surface portion. The martensite may damage the fatigue life in the plastic transformation region even though the amount of the martensite is small. In addition, in the case that surface decarburization occurs, the fatigue strength may be further reduced due to a decrease in surface strength.
When the time maintained in the solder spot is short (Comparative Example 5), remaining austenite may reside or the fatigue life may reduced due to the martensite formed in a cooling step after passing the solder spot.
In the case that the maintained temperature of the patenting (or austempering) annealing temperature is changed into a low temperature (Comparative Example 3), the yield strength measured after the annealing may increase.
The fatigue life is relatively long in an elastic deformation region. However, in the case that the plastic deformation region is included in an operation region, the fatigue life may be short. This is because bainite carbide is not effective to prevent a growth of the crack.
In the case that the plastic deformation region is included in the spring operation region (i.e., a stress applied to a surface acts larger than the yield strength), a crack may be easily generated. Thus, the fatigue life may be largely determined by the growth of the crack. Referring to
In the case that the maintained patenting (or austempering) temperature is above 570° C. (Comparative Example 4), a pearlite fine structure may be generated. Carbon may be largely strained due to a relatively high reduction ratio. However, the amount of laminar carbide per volume is relatively small as opposed to in the example of the present invention. Thus, the growth of the fatigue crack may not be efficiently prevented as opposed to in the example of the present invention.
In the case that a cold reduction ratio is not sufficient (i.e., Comparative Example 6) in spite of the quenching annealing temperature, and the temperature of the patenting (or austempering) and time are included in proper ranges, the fatigue life is smaller those of the examples of the present invention. However, the fatigue life is larger than those of other comparative examples.
As shown in Table 1, the fatigue life according to the examples of the present invention is superior to the comparative examples. This is because the high carbon steel sheet according to the example of the present invention has the large amount of strained laminar carbide having a relatively large ratio of length to width of the fine structure. The laminar carbide may prevent growth of the fatigue crack. In the case that the plastic deformation region is included in the spring operation region (i.e., a stress applied to a surface acts larger than the yield strength), a crack may be easily generated. Thus, the fatigue life may be mainly determined by the growth of the crack. As a result, the fatigue life may be varied by the fine structure that is capable of preventing the growth of the crack and the length of a path where the fatigue crack grows.
The high carbon steel sheet according to the examples of the present invention includes a fine structure in which a large amount of fine pearlite that is strained by a rolling process per volume is arranged. The fine structure is advantageous to increase the fatigue life.
A high carbon steel sheet according to the embodiments of the present invention has the large amount of laminar carbide having a relatively large ratio of length to width of over about 10:1. Thus, growth of a crack may be efficiently prevented.
In addition, fatigue life of the high carbon steel sheet may be improved because the growth of the crack is prevented.
Further, in the case that a spring is formed using the high carbon steel sheet according to the embodiments of the present invention, a generated crack may not easily grow. Thus, the spring may operate for a relatively long time even though the spring operates in a plastic deformation region.
| Number | Date | Country | Kind |
|---|---|---|---|
| 10-2007-0055199 | Jun 2007 | KR | national |
| Filing Document | Filing Date | Country | Kind | 371c Date |
|---|---|---|---|---|
| PCT/KR2007/006863 | 12/27/2007 | WO | 00 | 11/13/2009 |
