HIGH-TOUGHNESS STEEL AND METHOD FOR MANUFACTURING THE SAME
A method for manufacturing high-toughness steel includes the following steps. A steel billet is provided. The steel billet is subjected to a heating step, a hot rolling step, a spheroidizing step and a quenching and tempering heat treatment step to obtain the high-toughness steel. The steel billet includes a specific proportion of composition. After the steel billet undergoes the heating step, the hot rolling step and the spheroidizing step, the obtained spheroidized steel has good processability and formability. The high-toughness steel has good hardness and impact toughness.
This application claims priority to Taiwan Application Serial Number 113116828, filed May 7, 2024, which is herein incorporated by reference in its entirety.
BACKGROUND Field of InventionThe present invention relates to a high-toughness steel and a method for manufacturing high-toughness steel.
Description of Related ArtIt is known that a steel containing a specific proportion of nickel, copper, vanadium and a method for manufacturing the same, in which the steel is subjected to a rolling step under high temperature and air cooling step, followed by a quenching step under high temperature and a tempering step under high temperature, so as to obtain the steel including a ferrite microstructure of 35% to 40% and a balanced amount of bainite microstructure. However, addition of the high-priced elements such as nickel, copper and vanadium may increase the alloy cost, and the above-mentioned steel with ferrite microstructure and bainite microstructure has poor cold-worked properties and cannot be subjected to subsequent cold-working step for large-scale reduction requirements.
At present, it is also known that a steel including a specific proportion of nickel, copper and molybdenum and a method for manufacturing the same. After the steel is rolled, it is rapidly cooled at a cooling rate of 50° C./second and then subjected to a tempering step under high temperature. However, the addition of high-priced elements such as nickel, copper and molybdenum may increase the alloy cost, and the steel produced by the above rapidly cooling step has poor cold-worked properties and cannot be subjected to subsequent cold-working step for large-scale reduction requirements.
At present, it is also known that a steel containing a specific proportion of nickel, molybdenum and vanadium and a method for manufacturing the same. The rolled microstructure is not controlled during the hot rolling temperature process. The cold-worked properties of the produced are poor and cannot be subjected to subsequent cold-working step for large-scale reduction requirements. In addition, the addition of high-priced elements such as nickel, molybdenum and vanadium may increase the alloy cost.
At present, it is also known that a steel containing a specific proportion of nickel and molybdenum and a method for manufacturing the same, which uses a specific forging method, and then performs a quenching step under high temperature and a tempering step under high temperature. However, the hot forged materials have the poor cold-worked properties and cannot be subjected to subsequent cold-working step for large-scale reduction requirements. In addition, the addition of high-priced elements such as nickel and molybdenum may increase the alloy cost.
At present, it is also known that a steel including a specific proportion of nickel and copper and a method for manufacturing the same, which is rapidly cooled at a cooling rate of 15° C./second to 20° C./second after the hot rolled step, thereby obtaining a steel with a ferrite microstructure and a bainite microstructure. However, the addition of high-priced elements such as nickel and copper may increase the alloy cost, and the above mentioned steel with ferrite microstructure and a bainite microstructure has poor cold-worked properties and cannot be subjected to subsequent cold-working step for large-scale reduction requirements. In addition, the above mentioned steel has not undergone a quenching and tempering heat treatment step. Therefore, the resulting hot rolled steel microstructure has poor distribution uniformity and large variation in properties.
At present, it is also known that a steel including a specific proportion of nickel, molybdenum and vanadium and a method for manufacturing the same. After the hot rolling step, two stage of rapid cooling are used. In the first stage, the cooling rate is greater than 60° C./second, and in the second stage, the rapidly cooling rate is greater than 50° C./second to a certain temperature, thereby obtaining a steel plate with tempered martensite microstructure and bainite microstructure. However, the addition of high-priced elements such as nickel, molybdenum and vanadium may increase the alloy cost, and the steel produced by the above rapid cooling step has poor cold-worked properties and cannot be subjected to subsequent cold-working step for large-scale reduction requirements. In addition, the above mentioned steel has not undergone a quenching and tempering heat treatment step. Therefore, the resulting hot rolled steel microstructure has poor distribution uniformity and large variation in properties.
Accordingly, there is an urgent need to develop a steel and a method for manufacturing the same to improve the above problems.
SUMMARYThe invention provides a steel billet including specific content of niobium in the high-toughness steel and a method for manufacturing the same, which uses high heating temperature, low rolled temperature and slow cooling to improve the precipitation of niobium compounds, so that the rolled steel has a fine-grained microstructure with low hardness and easy spheroidization. Next, after the spheroidizing step, the steel is subjected to a quenching and tempering heat treatment step of low temperature tempering to obtain a high-toughness steel of the present invention. The steel billet of the present invention excludes nickel, copper, molybdenum, vanadium and aluminum, so the alloy cost can be reduced.
The present invention of the high-toughness steel is obtained by the above-mentioned method. The high-toughness steel has good hardness and impact toughness after the above quenching and tempering heat treatment step.
At least one embodiment of the present invention provides a method for manufacturing high-toughness steel, which includes the following steps. A steel billet is provided, in which a total weight of the steel billet is as 100 weight percent. The steel billet includes 0.3 weight percent to 0.35 weight percent of carbon, 0.01 weight percent to 0.1 weight percent of silicon, 0.8 weight percent to 1.4 weight percent of manganese, 0.6 weight percent to 0.8 weight percent of chromium, 0.025 weight percent to 0.045 weight percent of niobium, 0.02 weight percent to 0.03 weight percent of titanium, 0.0015 weight percent to 0.0025 weight percent of boron, the balance being iron and inevitable impurities, in which the steel billet excludes nickel, copper, molybdenum, vanadium and aluminum. The steel billet is subjected to a heating step to obtain a heated steel billet. The heated steel billet is subjected to a hot rolling step to obtain rolled steel. The rolled steel is subjected to a spheroidizing step to obtain spheroidized steel. The spheroidized steel is subjected to a quenching and tempering heat treatment step to obtain a high toughness steel.
At least one embodiment of the present invention, the heating temperature of the heating step is 1150° C. to 1250° C., and the heating temperature holding time of the heating step is 1.0 hour to 2.5 hours.
At least one embodiment of the present invention, the rolled temperature of the hot rolling step is 800° C. to 900° C., and the post-rolling cooling rate of the hot rolling step is 0.3° C./second to 0.5° C./second.
At least one embodiment of the present invention, the rolled steel consists of a ferrite microstructure with a volume fraction of 40% to 50% and a balanced amount of pearlite microstructure.
At least one embodiment of the present invention, the spheroidizing temperature of the spheroidizing step is 765° C. to 775° C., and the spheroidizing temperature holding time of the spheroidizing step is 2 hours to 4 hours.
At least one embodiment of the present invention, the post-spheroidizing cooling rate of the spheroidizing step is 10° C./hour to 12° C./hour.
At least one embodiment of the present invention, the spheroidization rate of spheroidized steel is classified as grade 1, and the hardness of spheroidized steel is HRB70 to HRB73 according to Rockwell hardness B scale.
At least one embodiment of the present invention, the quenching and tempering heat treatment step includes an austenitization treatment and a tempering treatment after the austenitization treatment. The austenitization temperature of the austenitization treatment is 840° C. to 870° C., and an austenitization temperature holding time of the austenitization treatment is 1 hour to 2 hours. The tempering temperature of the tempering treatment is 200° C. to 250° C., and the tempering temperature holding time of the tempering treatment is 1 hour to 2 hours.
At least one embodiment of the present invention, the above mentioned method for manufacturing high-toughness steel further includes the step of selectively performing a wire drawing step on the rolled steel before performing the spheroidizing step on the rolled steel, and the step of selectively performing a forming step on the spheroidized steel before performing the quenching and tempering heat treatment step on the spheroidized steel.
Another embodiment of the present invention provides a method for manufacturing high-toughness steel, which includes the following steps. A steel billet is provided, in which a total weight of the steel billet is as 100 weight percent. The steel billet includes 0.3 weight percent to 0.35 weight percent of carbon, 0.01 weight percent to 0.1 weight percent of silicon, 0.8 weight percent to 1.4 weight percent of manganese, 0.6 weight percent to 0.8 weight percent of chromium, 0.025 weight percent to 0.045 weight percent of niobium, 0.02 weight percent to 0.03 weight percent of titanium, 0.0015 weight percent to 0.0025 weight percent of boron, the balance being iron and inevitable impurities, in which the steel billet excludes nickel, copper, molybdenum, vanadium and aluminum. The steel billet is subjected to a hot rolling step to obtain rolled steel. The rolled steel is subjected to a wire drawing step to obtain drawn rolled steel. The drawn rolled steel is subjected to a spheroidizing step to obtain spheroidized steel. The spheroidized steel is subjected to a quenching and tempering heat treatment step to obtain a high toughness steel. The quenching and tempering heat treatment step includes an austenitization treatment and a tempering treatment after the austenitization treatment. The austenitization temperature of the austenitization treatment is 840° C. to 870° C., and an austenitization temperature holding time of the austenitization treatment is 1 hour to 2 hours. The tempering temperature of the tempering treatment is 200° C. to 250° C., and the tempering temperature holding time of the tempering treatment is 1 hour to 2 hours.
At least one embodiment of the present invention, the rolled temperature of the hot rolling step is 800° C. to 900° C., and the post-rolling cooling rate of the hot rolling step is 0.3° C./second to 0.5° C./second.
At least one embodiment of the present invention, the rolled steel consists of a ferrite microstructure with a volume fraction of 40% to 50% and a balanced amount of pearlite microstructure.
At least one embodiment of the present invention, the spheroidizing temperature of the spheroidizing step is 765° C. to 775° C., and the spheroidizing temperature holding time of the spheroidizing step is 2 hours to 4 hours.
At least one embodiment of the present invention, the post-spheroidizing cooling rate of the spheroidizing step is 10° C./hour to 12° C./hour.
At least one embodiment of the present invention, the spheroidization rate of spheroidized steel is classified as grade 1, and the hardness of spheroidized steel is HRB70 to HRB73 according to Rockwell hardness B scale.
At least one embodiment of the present invention, the above mentioned method for manufacturing high-toughness steel further includes the step of selectively performing a forming step on the spheroidized steel before performing the quenching and tempering heat treatment step on the spheroidized steel.
At least one embodiment of the present invention, the high-toughness steel made by the above method, the hardness of the high-toughness steel is HRC47 to HRC49, and the impact toughness of the high-toughness steel is 55.7J to 60.3J.
The invention can be more fully understood by reading the following detailed description of the embodiment, with reference made to the accompanying drawings as follows:
The single figure is a schematic flow chart of a method for manufacturing high-toughness steel according to some embodiments of the present invention.
The manufacture and uses of embodiments of the invention are discussed in detail below. It will be understood that the embodiments provide many applicable inventive concepts that can be practiced in a wide variety of specific contexts. The specific embodiments discussed are illustrative only and are not intended to limit the scope of the invention.
The range of “one value to another value” recited herein is a summary expression that avoids enumerating all the values in the range one by one in the specification. Therefore, the description of a specific numerical range covers any numerical value within the numerical range and the smaller numerical range defined by any numerical value within the numerical range, as if the arbitrary numerical value and the smaller numerical range are expressly written in the description.
Referring to the figure, which is a schematic flow chart of a method 100 for manufacturing high-toughness steel according to some embodiments of the present invention. A steel billet is provided, as shown in step 110 of the figure. Based on a total weight of the steel billet as 100 weight percent, the steel billet includes 0.3 weight percent to 0.35 weight percent of carbon, 0.01 weight percent to 0.1 weight percent of silicon, 0.8 weight percent to 1.4 weight percent of manganese, 0.6 weight percent to 0.8 weight percent of chromium, 0.025 weight percent to 0.045 weight percent of niobium, 0.02 weight percent to 0.03 weight percent of titanium, 0.0015 weight percent to 0.0025 weight percent of boron, the balance being iron and inevitable impurities. It should be noted that the “inevitable impurities” mentioned herein can be trace elements such as phosphorus, sulfur, copper, nitrogen or aluminum without additional addition of the above trace elements.
It is worth noting that the steel billet of the present invention excludes nickel, copper, molybdenum, vanadium and aluminum. In other words, the steel billet of the present invention does not add additional nickel, copper, molybdenum, vanadium and aluminum. Compared with the conventional ones, the steel billet of the present invention excludes high-priced elements such as nickel, copper, molybdenum, vanadium and aluminum. Therefore, the steel billet used in the present invention can reduce alloy cost.
In some embodiments, the steel billet includes 0.31, 0.32, 0.33, or 0.34 weight percent of carbon. If the carbon content was less than 0.3 weight percent, the strength and/or impact toughness of the steel obtained by the quenching and tempering heat treatment step would be insufficient. If the carbon content was greater than 0.35 weight percent, the rolled steel obtained after the hot rolling step would be too much pearlite microstructure (i.e., the volume fraction is greater than 60%), making the steel difficult to process and form, thus affecting the hardness and impact toughness of the steel obtained after quenching and tempering heat treatment step.
In some embodiments, the steel billet includes 0.02, 0.04, 0.06 or 0.08 weight percent of silicon. If the silicon content was less than 0.01 weight percent, the effect of the inhibiting carbide precipitation would not be obtained and would not be conducive to the spheroidizing step. If the silicon content was greater than 0.1 weight percent, the hardening rate of the formed steel would be increased, which would not be conducive to the processability and formability of the steel.
In some embodiments, the steel billet includes 0.85, 0.9, 1, 1.1, 1.2 or 1.3 weight percent of manganese. Manganese can optimize the hardenability of steel so that the steel will not have non-uniform distribution of hardenability due to size effect. If the manganese content was less than 0.8 weight percent, the hardenability of the steel obtained through the subsequent quenching and tempering heat treatment step would be insufficient, so that the steel with specific hardness and impact toughness of the present invention would not be obtained. If the manganese content was greater than 1.4 weight percent, the ferrite microstructure and pearlite microstructure of the rolled steel obtained by the hot rolling step would be non-uniform distributed, thereby affecting the processability and formability of the steel, and thus the steel with specific hardness and impact toughness of the present invention would not be obtained.
In some embodiments, the steel billet includes 0.65, 0.7 or 0.75 weight percent of chromium. Chromium can also optimize the hardenability of steel so that the steel will not have non-uniform distribution of hardenability due to size effect. In addition, adding an appropriate amount of chromium can make the steel have a good spheroidization rate. If the chromium content was less than 0.6 weight percent, the spheroidization rate of the spheroidized steel obtained by the spheroidizing step would not be good, and the hardenability of the steel obtained by the quenching and tempering heat treatment step would be insufficient, so that the steel with the specific hardness and impact toughness of the present invention would not be obtained. If the chromium content was greater than 0.8 weight percent, the ferrite microstructure and pearlite microstructure of the rolled steel obtained by the hot rolling step would be non-uniform distributed, thereby affecting the processability and formability of the steel, and thus the steel with specific hardness and impact toughness of the present invention would not be obtained.
In some embodiments, the steel billet includes 0.03, 0.032, 0.035, 0.038 or 0.04 weight percent of niobium. Niobium has the effect of refining grains. If the niobium content was less than 0.025 weight percent, the effect of refining grains would be insufficient, so that the steel with the specific impact toughness of the present invention would not be obtained. If the niobium content was greater than 0.045 weight percent, it would not only increase the alloy cost, but the excess niobium would not be completely dissolved and would be residual crystals in the heated steel billet, and thus the steel with specific impact toughness of the present invention would not be obtained.
In some embodiments, the steel billet includes 0.022, 0.025 or 0.028 weight percent of titanium. Titanium also has the effect of refining grains. Specifically, TiC produced by titanium and carbon has the effect of refining grains. In addition, titanium can be formed TiN with the nitrogen in the steel billet, so that the nitrogen in the steel billet cannot be formed BN with boron, thus achieving the effect of adding boron. If the titanium content was less than 0.02 weight percent, it would not be the effect of refining grains, and the boron in the steel billet would not be any effect. If the titanium content was greater than 0.03 weight percent, there would be no significant effect on the refining grains.
In some embodiments, the steel billet includes 0.0018, 0.002 or 0.0022 weight percent of boron. Boron affects the hardenability of steel. If the boron content was less than 0.0015 weight percent, the hardenability of the steel obtained by the quenching and tempering heat treatment step would be insufficient. If the boron content was greater than 0.0025 weight percent, it would not be significantly helpful for the high-toughness steel of the present invention.
Afterwards, the steel billet is subjected to a heating step to obtain a heated steel billet, as shown in step 120 of the figure. In some embodiments, the heating temperature of the heating step is 1150° C. to 1250° C., such as 1200° C. or 1220° C. In some embodiments, the heating temperature holding time of the heating step is 1.0 hour to 2.5 hours. In the hot rolling step, fine niobium will be strained and induced to precipitate. Therefore, in order to promote the precipitation of large amounts of niobium precipitates, the solid solution amount of niobium is increased by increasing the heating temperature of the steel billet, thereby increasing the amount of niobium precipitate during the hot rolling step. If the heating temperature was less than 1150° C. and/or the heating holding time was less than 1.0 hour, because there was too little niobium dissolved into the steel billet, the effect of refining grains in the steel would not be achieved, and therefore the specific impact toughness of the present invention would not be obtained. If the heating temperature was greater than 1250° C. and/or the heating holding time was greater than 2.5 hours, it would not meet the cost considerations of the process and would not be conducive to hot rolling step.
In other words, the steel billet including 0.025 to 0.045 weight percent of niobium can be completely dissolved into the steel billet at 1150° C. to 1250° C.
Then, the heated steel billet is subjected to a hot rolling step to obtain rolled steel, as shown in step 130 of the figure. In some embodiments, the rolled temperature of the hot rolling step is 800° C. to 900° C., such as 820° C. or 850° C. If the rolled temperature was less than 800° C., the grain size of the obtained rolled steel would be too small and the strength would be high, which would not be conducive to the wire drawing step and/or spheroidizing step. If the rolled temperature was greater than 900° C., the grain size of the obtained rolled steel would be large and the strength would be insufficient.
In some embodiments, the post-rolling cooling rate of the hot rolling step is 0.3° C./second to 0.5° C./second, such as 0.4° C./second.
Although the present invention has been described in considerable detail with reference to certain embodiments thereof, other embodiments are possible. Therefore, the spirit and scope of the appended claims should not be limited to the description of the embodiments contained herein. It can be understood that the “post-rolling cooling rate” in the present invention refers to the steel billet to be heated to 800° C. to 900° C., and the heated steel billet is cooled at a cooling rate of 0.3° C./second to 0.5° C./second to obtain rolled steel. If the post-rolling cooling rate was less than 0.3° C./second, the production efficiency of the steel would be poor, the cost would be high, and the grain size of the obtained rolled steel would be too large, resulting in poor processability and formability. If the post-rolling cooling rate was greater than 0.5° C./second, the hardness of the obtained rolled steel would be high, which would not be conducive to the wire drawing step.
In other words, if the rolled temperature is 800° C. to 900° C. and the post-rolling cooling rate is 0.3° C./second to 0.5° C./second, a large amount of niobium compounds can be precipitated and the grains can be refined, making the CCT curve left shift, and use a large amount of precipitated niobium compounds as nucleation points of the ferrite microstructure to further obtain more soft-phase ferrite microstructure. Furthermore, the effect of refining grains is also used to break up the originally pearlite microstructure, so that the distribution of ferrite microstructure and pearlite microstructure is more uniform can improve the uniformity of carbides dispersion after the spheroidizing step, thereby facilitating forming step.
In some embodiments, the rolled steel is consisted of a ferrite microstructure with a volume fraction of 40% to 50% and a balanced amount of pearlite microstructure. In other words, the rolled steel is consisted of a ferrite microstructure with a volume fraction of 40% to 50% and a pearlite microstructure with a volume fraction of 50% to 60%. Specifically, the rolled steel of the present invention has no other phase microstructure except the ferrite microstructure and pearlite microstructure, for example, there is no martensite microstructure and/or bainite microstructure.
In some embodiments, after the hot rolling step described above (i.e., step 130 of the figure), the spheroidizing step can optionally be subjected to a wire drawing step. For example, the rolled steel with a large diameter is drawn into smaller diameter steel.
Next, the rolled steel is subjected to a spheroidizing step to obtain spheroidized steel, as shown in step 140 of the figure. The purpose of the spheroidizing step is to allow the flake-shaped pearlite microstructure to re-melt into the steel, and the remaining pearlite microstructure that has not been re-melted into the steel can become the nucleation point for carbides after the spheroidizing step. In some embodiments, the spheroidizing temperature of the spheroidizing step is 765° C. to 775° C., such as 770° C. In some embodiments, the spheroidizing temperature holding time in of the spheroidizing step is 2 hours to 4 hours, such as 3 hours. If the spheroidizing temperature was less than 765° C. and/or the spheroidizing temperature holding time was less than 2 hours, an insufficient amount of the pearlite microstructure would be re-melted into the steel, thus affecting the spheroidization rate of the spherical carbides, so the steel with specific impact toughness of the present invention would not be obtained. If the spheroidizing temperature was greater than 775° C. and/or the spheroidizing temperature holding time was greater than 4 hours, all the flake-shaped pearlite microstructure would be re-melted into the steel without any nucleation point for spherical carbides. Therefore, the steel with specific impact toughness of the present invention would not be obtained.
In some embodiments, the post-spheroidizing cooling rate of the spheroidizing step is 10° C./hour to 12° C./hour, such as 11° C./hour. It can be understood that the “post-spheroidizing cooling rate” in the present invention refers to the rolled steel to be heated to 765° C. to 775° C. and maintained 2 hours to 4 hours, the rolled steel billet is cooled at a cooling rate of 10° C./hour to 12° C./hour to obtain spheroidized steel. If the cooling rate after spheroidization was less than 10° C./hour, the process cost would be increased and the carbide size of the obtained spheroidized steel would be large, resulting in poor processability and formability. Therefore, the steel with the specific impact toughness of the present invention would not be obtained. If the cooling rate after spheroidization was greater than 12° C./hour, the carbide size of the obtained spheroidized steel would be too small and non-uniform distributed. In addition, there would also be the formation of regenerated pearlite microstructure.
In some embodiments, the spheroidization rate of the spheroidized steel is classified as grade 1. In some embodiments, the hardness of the spheroidized steel is HRB70 to HRB73 according to Rockwell hardness B scale, such as HRB71 or HRB72. It can be understood that the spheroidization rate and the hardness of the spheroidized steel can be used to determine the properties of processability (for example, cold working). If the scale of the spheroidization rate is lower, it means that the distribution of spherical carbides is more uniform and the processing properties are good. On the contrary, if the scale of the spheroidization rate is higher, it means that the distribution of spherical carbides is non-uniform and the processing properties are poor. If the distribution of spherical carbides is non-uniform, the uniformity of forming will be worse, for example, the steel after processing is east to breakage or the size does not meet the specifications.
In some embodiments, after the above-described spheroidizing step (i.e., step 140 of the figure), the spheroidized steel can optionally be subjected to a forming step. The forming step can be, for example, cold working (i.e., cold forging). In the forming step, the spheroidized steel is formed into a specific shape, for example, a nut, a bolt, or a sleeve. It should be noted that the “forming step” of the present invention involves subjecting spheroidized steel to a severe deformation process, with a reduction rate of at least 35% to 50%.
Afterwards, the spheroidized steel is subjected to a quenching and tempering heat treatment step to obtain the high-toughness steel, as shown in steps 150 and 160 of the figure. In some embodiments, the quenching and tempering heat treatment step includes a tempering treatment and a tempering treatment after the austenitization treatment. It can be understood that the “austenitization treatment” of the present invention can also be called quenching treatment.
The purpose of the austenitization treatment is to re-melt the spherical carbides into the steel. In some embodiments, the austenitization temperature of the austenitization treatment is 840° C. to 870° C., such as 850° C. or 860° C. In some embodiments, the austenitization temperature holding time of the austenitization treatment is 1 hour to 2 hours. If the austenitization temperature was less than 840° C. and/or the austenitization temperature holding time was less than 1 hour, the carbides would not be completely re-melted into the steel. If the austenitization temperature was greater than 870° C., it would not be conducive to cooling (for example, cold quenching) step because of equipment limitations and high process costs. If the austenitization temperature holding time was greater than 2 hours, the grain size of the steel obtained would be too large and the toughness would be poor.
After the above-mentioned austenitization treatment and cooling steps, the steel has residual stress and strain energy. The purpose of the tempering treatment is relieving residual stress and strain energy in the steel. In some embodiments, the tempering temperature of the tempering treatment is 200° C. to 250° C., such as 210° C. or 230° C. In some embodiments, the tempering temperature holding time of the tempering treatment is 1 hour to 2 hours. If the tempering temperature is between 200° C. and 250° C., the martensite microstructure in the steel is tempered into the tempered martensite microstructure, and the carbide can be precipitated finer and evenly distributed. If the tempering temperature was less than 200° C. and/or the tempering temperature holding time was less than 1 hour, the carbide would not easily precipitate out of the tempered martensite microstructure. If the tempering temperature was greater than 250° C. and/or the tempering temperature holding time was greater than 2 hours, the grain size of the steel obtained would be larger and the toughness would be poor.
In some embodiments, the high-toughness steel made by the above method, in which the hardness of the high-toughness steel is HRC47 to HRC49, such as HRC48 according to Rockwell hardness C scale. In some embodiments, the impact toughness of the high-toughness steel is 55.7J (Joules) to 60.3J, such as 57J or 59J.
Briefly, the high-toughness steel of the present invention is processed by the types and contents of the aforementioned elements and after specific processes (for example, high-temperature heating step, low-temperature rolling temperature, slow cooling post-cooling rates and low-temperature tempering), the high-toughness steel formed has specific hardness and impact toughness. In addition, the steel billet of the present invention excludes nickel, copper, molybdenum, vanadium and aluminum, so the alloy cost can be reduced.
It will be apparent to those skilled in the art that various modifications and variations can be made to the structure of the present invention without departing from the scope or spirit of the invention. In view of the foregoing, it is intended that the present invention cover modifications and variations of this invention provided they fall within the scope of the following claims.
Example 1In Example 1, the steel billet included 0.3 weight percent of carbon, 0.01 weight percent of silicon, 0.8 weight percent of manganese, 0.8 weight percent of chromium, 0.025 weight percent of niobium, 0.025 weight percent of titanium, 0.0020 weight percent of boron, and the balance being iron and inevitable impurities. The compounds and contents of the steel billet in Example 1 were shown in Table 1 below.
The steel billet of Example 1 was subjected to a heating step to obtain a heated steel billet, in which the heating temperature of the heating step was 1150° C. Then, a hot rolling step was performed on the hot steel billet to obtain a rolled steel, in which the rolled temperature in the hot rolling step was 800° C. and the post-rolling cooling rate of the hot rolling step was 0.3° C./second. The obtained rolled steel consisted of a ferrite microstructure with a volume fraction of 50%. Next, the rolled steel was subjected to a spheroidizing step to obtain spheroidized steel, in which the spheroidizing temperature was 765° C., the spheroidizing temperature holding time was 2 hours to 4 hours, and the post-spheroidizing cooling rate was 10° C./hour to 12° C./hour. The spheroidization rate of the spheroidized steel was classified as grade 1, and the hardness of the spheroidized steel was HRB72. Afterwards, the spheroidized steel was subjected to a quenching and tempering heat treatment step to obtain the high-toughness steel, in which the austenitization temperature (also known as quenching temperature) of the austenitization treatment (i.e., quenching treatment) was 840° C., the tempering temperature of the tempering treatment was 200° C. The hardness of the obtained high-toughness steel was HRC47, and the impact toughness was 57.8J. The process conditions and evaluation results of Example 1 were shown in Table 2 below.
Example 2 to Example 4 and Comparative Example 1 to Comparative Example 4.
Please refer to Table 1 above for the components and contents of the steel billets in Example 2 to 4 and Comparative Examples 1 to 4. The process conditions and evaluation results of Example 2 to 4 and Comparative Examples 1 to 4 were shown in Table 2 above.
Evaluation Method 1. Measurement of Metallographic MicrostructureThe measurement of the metallographic microstructure of the present invention was to use a scanning electron microscope to capture images of the steel of each example and comparative example, and measured the area percentage occupied by the ferrite microstructure in the image, and the average value of the 5 obtained area percentage was used as the volume fraction, and unit was %. The measurement results of the metallographic microstructure of the steel of the examples and comparative examples were shown in Table 2 above.
2. Measurement of Spheroidization RateThe spheroidization rate of the present invention was measured using the JIS G3509-2 spheroidization classification standard method as the evaluation result. The measurement results of the spheroidization rate of the steel of Examples 1 to 4 and Comparative Examples 1 to 4 were shown in Table 2 above.
3. Test of Post-Spheroidized HardnessThe post-spheroidized hardness referred to here in the present invention was measured by the standard detection method according to Rockwell hardness B scale, in which a Rockwell hardness tester was used as the indenter, and the load was 100 kgf. The post-spheroidized hardness of the steel of Examples 1 to 4 and Comparative Examples 1 to 4 was measured by this method, and the measurement results were shown in Table 2 above.
4. Test of Hardness after Quenching and Tempering
The hardness after quenching and tempering referred to here in the present invention was measured by the standard detection method according to Rockwell hardness C scale, in which a Rockwell hardness tester was used as the indenter, and the load was 150 kgf. The hardness after quenching and tempering of the steel of Examples 1 to 4 and Comparative Examples 1 to 4 was measured by this method, and the measurement results were shown in Table 2 above.
5. Test on Impact Toughness after Quenching and Tempering
The impact toughness after quenching and tempering referred to here in the present invention was tested using the ISO 148-1 impact test standard method. The impact toughness after quenching and tempering of the steel of Examples 1 to 4 and Comparative Examples 1 to 4 was measured by this method, and the unit was J. The measurement results were shown in Table 2 above.
Please refer to Table 2. In Comparative Examples 1 to 4, the heating temperature were 1050° C. and 1100° C., the rolled temperature were 950° C. and 1000° C., the post-rolling cooling rate were 1.5° C./second and 2.0° C./second, and the resulting rolled steel included a ferrite microstructure with a volume fraction of around 16% to 25%.
In Comparative Examples 1 to 4, it was found in the metallographic microstructure of the spheroidized steel after spheroidizing step that the spherical carbides were mixed with a large amount of recycled pearlite microstructure, so the measured spheroidization rate was poor (the spheroidization rate of Comparative Examples 1 to 4 was classified as grade 2 to grade 3), and the hardness was relatively high (the post-spheroidized hardness of Comparative Examples 1 to 4 was HRB76 to HRB79). The main reason for these results was that the rolled temperature used in Comparative Examples 1 to 4 was relatively high, and the grain size of the rolled steel obtained was larger, so a large amount of pearlite microstructure could be found in the rolled microstructure, resulting in a relatively non-uniform distribution of ferrite microstructure and pearlite microstructure. Therefore, the dispersion of the spherical carbides after spheroidization was poor and the post-spheroidized hardness was higher. Therefore, the cold-worked processability and formability of Comparative Examples 1 to 4 was poor. Specifically, in addition to the higher hardness that easily caused mold abrasion and increased process costs, the spheroidized steel of Comparative Examples 1 to 4 could not be cold-formed with large reductions (for example, at least a 35% or 50% reduction rate). If Comparative Examples 1 to 4 wanted to increase the deformation reduction of spheroidized steel, the spheroidizing temperature holding time needed to be extended (for example, more than 4 hours) or additional spheroidizing steps were required to soften the spheroidized steel.
In Comparative Examples 1 to 4, the cold forming step with a deformation reduction was 35%, an austenitization temperature (i.e., quenching temperature) was 880° C. and a tempering temperature was 350° C. or 400° C. The steel obtained had poor hardness (the hardness after quenching and tempering of Comparative Examples 1 to 4 was HRC41 to HRC45) and poor toughness (the impact toughness after quenching and tempering of Comparative Examples 1 to 4 was 25.3J to 38.6J). These results were mainly due to the insufficient number of precipitates in the base phase that hinder the movement of grain boundaries during the re-heating process of spheroidized steel at the austenitization temperature of 880° C., and there were not enough precipitate particles to fix the grain boundaries, therefore grain boundary movement (grain growth) could not be effectively inhibited, so that the grain size after quenching and tempering was too large, resulting in poor toughness of the steel.
Please refer to Table 2. In Examples 1 to 4, the heating temperature was 1150° C. to 1250° C., the rolled temperature was 800° C. to 900° C., the post-rolling cooling rate was 0.3° C./second to 0.5° C./second, and the resulting rolled steel included a ferrite microstructure with a volume fraction of around 40% to 50%.
In Examples 1 to 4, the spheroidization rate of the spheroidized steel after the spheroidizing step was good (the spheroidization rate of Examples 1 to 4 were all classified as grade 1), and the hardness was low (the post-spheroidized hardness of Examples 1 to 4 was HRB70 to HRB73). Therefore, the spheroidizing and softening effects of Examples 1 to 4 were significant and could meet the cold forming requirements for larger deformation reduction amounts (for example, at least 35% or 50% reduction rate).
In Examples 1 to 4, the cold forming step with a deformation reduction was 35%, an austenitization temperature (i.e., quenching temperature) was 840° C. or 870° C. and a tempering temperature was 200° C. or 250° C. The steel obtained had good hardness (the hardness after quenching and tempering of Examples 1 to 4 was HRC47 to HRC49) and good toughness (the impact toughness after quenching and tempering of Examples 1 to 4 was 55.7J to 60.3J). These results were mainly due to the presence of a large number of fine niobium precipitate particles in the base phase during the re-heating process of spheroidized steel at the austenitization temperature of 840° C. to 870° C., which could effectively hinder the movement of grain boundaries and fixed the grain boundaries, thus inhibiting the growth of grains. After quenching and low-temperature tempering, it could maintain the original fine-grained microstructure to enhance the fine-grained toughness effect. Therefore, after quenching and tempering, alloy steel with excellent strength and toughness could be obtained.
Comprehensive comparison of Comparative Examples 1 to 4 and Examples 1 to 4, the volume fraction 16% to 25% of ferrite microstructure increased to 40% to 50%, grade 2 to grade 3 of the spheroidization rate increased to grade 1, and HRB76 to HRB79 of the post-spheroidized hardness reduced to HRB70 to HRB73. Therefore, the cold-worked formability of Examples 1 to 4 was good. After the quenching and tempering heat treatment step, HRC41 to HRC45 of the hardness after quenching and tempering increased to HRC47 to HRC49, and 25.3J to 38.6J of the impact toughness after quenching and tempering increased to 55.7J to 60.3J. Therefore, the steel of Examples 1 to 4 had good mechanical properties.
The steel billet in the method for manufacturing high-toughness steel of the present invention includes a specific content of niobium, which uses high heating temperature, low rolled temperature and slow cooling to increase the precipitation of niobium compounds, so that the rolled steel has a fine-grained microstructure with low hardness and easy spheroidization. Therefore, the cold-worked formability of the rolled steel is good. Next, after the steel is spheroidized, a low-temperature tempering of quenching and tempering heat treatment step is used to obtain the steel with high hardness and high toughness of the present invention. The steel billet of the present invention excludes high-priced elements such as nickel, copper, molybdenum, vanadium and aluminum, so the alloy cost can be reduced.
It can be understood that a specific composition, a specific metallographic microstructure, a specific manufacturing method, and a specific evaluation method are used as examples in the present invention to illustrate the high-toughness steel and a method for manufacturing the same.
Although the present invention has been described in considerable detail with reference to certain embodiments thereof, other embodiments are possible. Therefore, the spirit and scope of the appended claims should not be limited to the description of the embodiments contained herein.
It will be apparent to those skilled in the art that various modifications and variations can be made to the structure of the present invention without departing from the scope or spirit of the invention. In view of the foregoing, it is intended that the present invention cover modifications and variations of this invention provided they fall within the scope of the following claims.
Claims
1. A method for manufacturing high-toughness steel, comprising:
- providing a steel billet, wherein a total weight of the steel billet is as 100 weight percent, the steel billet comprises: 0.3 weight percent to 0.35 weight percent of carbon; 0.01 weight percent to 0.1 weight percent of silicon; 0.8 weight percent to 1.4 weight percent of manganese; 0.6 weight percent to 0.8 weight percent of chromium; 0.025 weight percent to 0.045 weight percent of niobium; 0.02 weight percent to 0.03 weight percent of titanium; 0.0015 weight percent to 0.0025 weight percent of boron; and a balance being iron and inevitable impurities, wherein the steel billet excludes nickel, copper, molybdenum, vanadium and aluminum;
- performing a heating step on the steel billet, so as to obtain a heated steel billet;
- performing a hot rolling step on the heated steel billet, so as to obtain a rolled steel;
- performing a spheroidizing step on the rolled steel, so as to obtain a spheroidized steel; and
- performing a quenching and tempering heat treatment step on the spheroidized steel, so as to obtain the high-toughness steel.
2. The method for manufacturing high-toughness steel according to claim 1, wherein a heating temperature in the heating step is 1150° C. to 1250° C., and a heating temperature holding time in the heating step is 1.0 hour to 2.5 hours.
3. The method for manufacturing high-toughness steel according to claim 1, wherein a rolled temperature in the hot rolling step is 800° C. to 900° C., and a post-rolling cooling rate in the hot rolling step is 0.3° C./second to 0.5° C./second.
4. The method for manufacturing high-toughness steel according to claim 1, wherein the rolled steel consists of a ferrite microstructure with a volume fraction of 40% to 50% and a balanced amount of pearlite microstructure.
5. The method for manufacturing high-toughness steel according to claim 1, wherein a spheroidizing temperature in the spheroidizing step is 765° C. to 775° C., and a spheroidizing temperature holding time in the spheroidizing step is 2 hours to 4 hours.
6. The method for manufacturing high-toughness steel according to claim 5, wherein a post-spheroidizing cooling rate in the spheroidizing step is 10° C./hour to 12° C./hour.
7. The method for manufacturing high-toughness steel according to claim 1, wherein a spheroidization rate of the spheroidized steel is classified as grade 1, and a hardness of the spheroidized steel is HRB70 to HRB73 according to Rockwell hardness B scale.
8. The method for manufacturing high-toughness steel according to claim 1, wherein the quenching and tempering heat treatment step comprises:
- an austenitization treatment, wherein an austenitization temperature of the austenitization treatment is 840° C. to 870° C., and an austenitization temperature holding time of the austenitization treatment is 1 hour to 2 hours; and a tempering treatment after the austenitization treatment, wherein a tempering temperature of the tempering treatment is 200° C. to 250° C., and a tempering temperature holding time of the tempering treatment is 1 hour to 2 hours.
9. The method for manufacturing high-toughness steel according to claim 1, further comprising:
- before performing the spheroidizing step on the rolled steel, a wire drawing step is performed on the rolled steel; and
- before performing the quenching and tempering heat treatment step on the spheroidized steel, a forming step is performed on the spheroidized steel.
10. A method for manufacturing high-toughness steel, comprising:
- providing a steel billet, wherein a total weight of the steel billet is as 100 weight percent, the steel billet further comprises: 0.3 weight percent to 0.35 weight percent of carbon; 0.01 weight percent to 0.1 weight percent of silicon; 0.8 weight percent to 1.4 weight percent of manganese; 0.6 weight percent to 0.8 weight percent of chromium; 0.025 weight percent to 0.045 weight percent of niobium; 0.02 weight percent to 0.03 weight percent of titanium; 0.0015 weight percent to 0.0025 weight percent of boron; and a balance being iron and inevitable impurities, wherein the steel billet excludes nickel, copper, molybdenum, vanadium and aluminum;
- performing a hot rolling step, so as to obtain a rolled steel;
- performing a wire drawing step on the rolled steel, so as to obtain a drawn rolled steel;
- performing a spheroidizing step on the drawn rolled steel, so as to obtain a spheroidized steel; and
- performing a quenching and tempering heat treatment step on the spheroidized steel, so as to obtain the high-toughness steel, wherein the quenching and tempering heat treatment step comprises:
- an austenitization treatment, wherein an austenitization temperature of the austenitization treatment is 840° C. to 870° C., and an austenitization temperature holding time of the austenitization treatment is 1 hour to 2 hours; and a tempering treatment after the austenitization treatment, wherein a tempering temperature of the tempering treatment is 200° C. to 250° C., and a tempering temperature holding time of the tempering treatment is 1 hour to 2 hours.
11. The method for manufacturing high-toughness steel according to claim 10, wherein a rolled temperature in the hot rolling step is 800° C. to 900° C., and a post-rolling cooling rate in the hot rolling step is 0.3° C./second to 0.5° C./second.
12. The method for manufacturing high-toughness steel according to claim 10, wherein the rolled steel consists of a ferrite microstructure with a volume fraction of 40% to 50% and a balanced amount of pearlite microstructure.
13. The method for manufacturing high-toughness steel according to claim 10, wherein a spheroidizing temperature in the spheroidizing step is 765° C. to 775° C., and the spheroidizing temperature holding time in the spheroidizing step is 2 hours to 4 hours.
14. The method for manufacturing high-toughness steel according to claim 13, wherein a post-spheroidizing cooling rate in the spheroidizing step is 10° C./hour to 12° C./hour.
15. The method for manufacturing high-toughness steel according to claim 10, wherein a spheroidization rate of the spheroidized steel is classified as grade 1, and a hardness of the spheroidized steel is HRB70 to HRB73 according to Rockwell hardness B scale.
16. The method for manufacturing high-toughness steel according to claim 10, further comprising:
- before the quenching and tempering heat treatment step on the spheroidized steel, a forming step is performed on the spheroidized steel.
17. A high-toughness steel made by the method for manufacturing high-toughness steel according to claim 1, wherein a hardness of the high-toughness steel is HRC47 to HRC49, and an impact toughness of the high-toughness steel is 55.7J to 60.3J.
Type: Application
Filed: Apr 8, 2025
Publication Date: Nov 13, 2025
Inventors: Shih-Ming KUO (KAOHSIUNG), Herng-Shuoh JANG (Kaohsiung City), Ying-Hao WANG (Kaohsiung City), Wen-Fang CHIU (Pingtung County)
Application Number: 19/172,690
