STEEL FOR NITROCARBURIZING AND NITROCARBURIZED COMPONENT USING THE STEEL AS MATERIAL
According to the present invention, it is possible to obtain steel for nitrocarburizing having a predetermined chemical composition, a bainite area ratio exceeding 50% and excellent machinability by cutting before nitrocarburizing, and having strength and toughness equivalent to conventional steel, such as SCr420 carburized steel material, and excellent fatigue properties after nitrocarburizing.
This disclosure relates to steel for nitrocarburizing and nitrocarburized components using the steel as material. In particular, the disclosure relates to steel for nitrocarburizing that has excellent fatigue properties after nitrocarburizing and is suitable for use in automobiles and construction equipment and to nitrocarburized components using the steel as a material.
BACKGROUNDSince excellent fatigue properties are desired for machine structural components such as automobile gears, surface hardening is generally performed. Carburizing treatment, induction quench hardening and nitriding treatment are well-known forms of surface hardening.
With carburizing treatment, carbon is caused to infiltrate and diffuse into a high-temperature austenite region, yielding a deep hardening depth. Carburizing treatment is thus useful to improve fatigue strength.
However, since heat treatment distortion occurs, it is difficult to apply carburizing treatment to components that, from the perspective of noise or the like, require high dimensional accuracy.
Induction quench hardening is a process of quenching a surface part by high frequency induction heating and, like carburizing treatment, causes degradation of dimensional accuracy.
Nitriding treatment is a process to harden a surface by causing nitrogen to infiltrate and diffuse into a high-temperature region at or below the Ac1 critical point. The treatment is long, taking 50 to 100 hours, and requires removal of a brittle compound layer on the surface after treatment.
Therefore, nitrocarburizing treatment has been developed for nitriding at approximately the same treatment temperature as nitriding treatment yet in a short time. In recent years, nitrocarburizing treatment has become commonly used on machine structural components and the like. During nitrocarburizing treatment, nitrogen and carbon are simultaneously caused to infiltrate and diffuse into a temperature region at 500° C. to 600° C. to harden the surface, making it possible to reduce the treatment time to half or less that of conventional nitriding treatment.
However, whereas it is possible to increase the core hardness by quench hardening during carburizing treatment, nitrocarburizing treatment is performed at a temperature at or below the critical point of steel, thus causing the core hardness not to increase and yielding nitrocarburized material with poorer fatigue strength than carburized material.
To improve the fatigue strength of nitrocarburized material, quenching and tempering are generally performed before nitrocarburizing to increase the core hardness. The resulting fatigue properties, however, cannot be considered sufficient. Furthermore, this approach increases manufacturing costs and reduces mechanical workability.
To address these problems, it has been proposed to form steel with a chemical composition including Ni, Al, Cr and Ti, to age-harden the core during nitrocarburizing by Ni—Al and Ni—Ti intermetallic compounds or by Cu compounds, and to precipitation-harden nitrides and carbides such as Cr, Al and Ti in a nitrided layer of the surface (JP 5-59488 A, JP 7-138701 A).
JP 2002-69572 A discloses cogging steel that contains 0.5% to 2% of Cu by hot forging and then air cooling the steel to provide a ferrite-based microstructure with solute Cu, precipitating the Cu during nitrocarburizing treatment at 580° C. for 120 minutes and, furthermore, concurrently precipitation-hardening Ti, V and Nb carbonitrides to yield a steel that, after the nitrocarburizing treatment, has excellent bending fatigue properties. JP 2010-163671 A discloses steel for nitrocarburizing having dispersed therein Ti—Mo carbides and carbides including at least one element selected from the group consisting of Nb, V and W.
While the nitrocarburizing steel recited in JP 5-59488 A and JP 7-138701 A improves bending fatigue strength through precipitation-hardening of Cu and the like, the resulting workability cannot be considered sufficient. By requiring the addition of a relatively large amount of Cu, Ti, V and Nb, the nitrocarburizing steel recited in JP 2002-69572 A has a high production cost. The steel for nitrocarburizing recited in JP 2010-163671 A has the problem of high production cost due to the inclusion of a relatively large amount of Ti and Mo.
In view of the foregoing, it could be helpful to provide steel for nitrocarburizing and a nitrocarburized component using the steel as material, the steel having a low hardness and excellent mechanical workability before nitrocarburizing while allowing for an increase in core hardness via nitrocarburizing treatment and allowing for relatively inexpensive manufacture of nitrocarburized components with excellent fatigue properties.
SUMMARYWe intensely studied the effects of the microstructure and composition of steel on the fatigue properties after nitrocarburizing of steel. As a result, we discovered that with a steel material provided with a specific amount of V and Nb in the steel composition and a bainite-based microstructure before nitrocarburizing, excellent fatigue properties are obtained after nitrocarburizing by performing nitrocarburizing treatment on the steel material while utilizing the rise in temperature to increase the core hardness by age precipitating fine precipitates in the core structure other than the nitrocarburized surface part.
We thus provide:
[1] A steel for nitrocarburizing comprising, in, mass %, C: 0.01% or more and less than 0.10%, Si: 1.0% or less, Mn: 0.5% to 3.0%, Cr: 0.30% to 3.0%, Mo: 0.005% to 0.4%, V: 0.02% to 0.5%, Nb: 0.003% to 0.15%, Al: 0.005% to 0.2%, S: 0.06% or less, P: 0.02% or less, B: 0.0003% to 0.01%, and the balance being Fe and incidental impurities, and including a microstructure with a bainite area ratio exceeding 50% before nitrocarburizing.
[2] The steel for nitrocarburizing according to [1], wherein after nitrocarburizing, precipitates including V and Nb are dispersed in a bainite phase.
[3] A nitrocarburized component using the steel for nitrocarburizing according to [1] or [2] as material.
It is thus possible to obtain steel for nitrocarburizing, and nitrocarburized components using the steel as material, that has excellent machinability by cutting before nitrocarburizing, and that after nitrocarburizing has strength and toughness equivalent to conventional steel, such as SCr420 carburized steel material, and excellent fatigue properties, thus proving extremely useful in industrial terms.
Our steels will be further described below with reference to the accompanying drawings, wherein:
The microstructure, chemical composition and manufacturing conditions of the steel for nitrocarburizing will be described.
1. Microstructure
The microstructure before nitrocarburizing has a bainite area ratio exceeding 50%, and the microstructure after nitrocarburizing has V and Nb precipitates dispersed in a bainite phase. When a matrix phase before nitrocarburizing is a bainite-based microstructure with a bainite area ratio exceeding 50%, formation of V and Nb precipitates in the matrix phase is drastically inhibited compared to a ferrite-pearlite microstructure. As a result, formation of the V and Nb precipitates before nitrocarburizing and consequent increased hardness of the steel can be prevented, thereby improving workability of cutting generally performed before nitrocarburizing. Furthermore, applying nitrocarburizing treatment to the steel causes the surface part to be nitrided and simultaneously age precipitates the V and Nb precipitates in the core bainite phase other than the nitrided surface part, thereby increasing the core hardness. Both the fatigue properties and the strength after nitrocarburizing therefore dramatically improve.
Note that the “microstructure with a bainite area ratio exceeding 50%” contemplated herein refers to the area ratio of the bainite microstructure (phase) exceeding 50% under cross-sectional microstructure observation (microstructure observation with a 200× optical microscope). The area ratio of the bainite phase preferably exceeds 60% and even more preferably exceeds 80%. Moreover, the V and Nb precipitates in the bainite phase are preferably a dispersion of fine precipitates having a grain size of less than 10 nm. Furthermore, for sufficient strengthening by precipitation, 500 or more of the V and Nb precipitates with the grain size of less than 10 nm preferably exist per 1 μm2.
2. Chemical Composition
Reasons for the limitations of the chemical composition in the steel for nitrocarburizing will now be described. The fraction of each steel component represents mass %.
C: 0.01% or More and Less Than 0.10%
Carbon (C) is added for bainite phase formation and to ensure strength. When the amount of C added is less than 0.01%, the amount of bainite formed decreases, as does the amount of V and Nb precipitates, thus making it difficult to ensure strength. On the other hand, when 0.10% or greater of C is added, the bainite phase becomes harder, thereby reducing the mechanical workability. Accordingly, the amount of C added is 0.01% or more and less than 0.10%. C preferably 0.03% or more and less than 0.10%.
Si: 1.0% or Less
Silicon (Si) is added for its usefulness in deoxidizing and bainite phase formation. Adding an amount of Si exceeding 1.0%, however, deteriorates mechanical workability and cold-rolling workability due to solid solution hardening of ferrite and bainite phases. Accordingly, the amount of Si added is 1.0% or less. The amount is preferably 0.5% or less and more preferably 0.3% or less. Note that for Si to contribute effectively to deoxidation, the amount of Si added is preferably 0.01% or more.
Mn: 0.5% to 3.0%
Manganese (Mn) is added for its usefulness in bainite phase formation and in increasing strength. When the amount of Mn added is less than 0.5%, the amount of bainite phase formed decreases, and V and Nb precipitates are formed, causing the hardness before nitrocarburizing to increase and the amount of V and Nb precipitates formed after nitrocarburizing treatment to decrease. In turn, this lowers the hardness after nitrocarburizing and makes it difficult to ensure strength. On the other hand, adding an amount of Mn exceeding 3.0% deteriorates mechanical workability and cold-rolling workability. Accordingly, the amount of Mn added is 0.5% to 3.0%. The amount is preferably 0.5% or more and 2.5% or less, and more preferably 0.6% or more and 2.0% or less.
Cr: 0.30% to 3.0%
Chromium (Cr) is added for its usefulness in bainite phase formation. When the amount of Cr added is less than 0.30%, the amount of bainite phase formed decreases, and V and Nb precipitates are formed, causing the hardness before nitrocarburizing to increase and the amount of V and Nb precipitates formed after nitrocarburizing treatment to decrease. In turn, this lowers the hardness after nitrocarburizing and makes it difficult to ensure strength. On the other hand, adding an amount of Cr exceeding 3.0% deteriorates mechanical workability and cold-rolling workability. Accordingly, the amount of Cr added is 0.30% to 3.0%. The amount is preferably 0.5% or more and 2.0% or less, and more preferably 0.5% or more and 1.5% or less.
V: 0.02% to 0.5%
Vanadium (V) forms fine precipitates along with Nb due to the rise in temperature during nitrocarburizing and is therefore an important element to increase core hardness and improve strength. An added amount of V less than 0.02% does not satisfactorily achieve these effects. On the other hand, adding an amount of V exceeding 0.5% causes the precipitates to coarsen. Accordingly, the amount of V added is 0.02% to 0.5%. The amount is preferably 0.03% or more and 0.3% or less, and more preferably 0.03% or more and 0.25% or less.
Nb: 0.003% to 0.15%
Niobium (Nb) forms fine precipitates along with V due to the rise in temperature during nitrocarburizing and is therefore an extremely effective element to increase core hardness and improve fatigue strength. An added amount of Nb less than 0.003% does not satisfactorily achieve these effects. On the other hand, adding an amount of Nb exceeding 0.15% causes the precipitates to coarsen. Accordingly, the amount of Nb added is 0.003% to 0.15%. The amount is preferably 0.02% or more and 0.12% or less.
Mo: 0.005% to 0.4%
Molybdenum (Mo) causes fine V and Nb precipitates to form and is effective in improving the strength of the nitrocarburized material. Mo is therefore an important element. Mo is also useful for bainite phase formation. To improve strength, 0.005% or more is added, but since Mo is an expensive element, adding more than 0.4% leads to increased component cost. Accordingly, the amount of Mo added is 0.005% to 0.4%. The amount is preferably 0.01% to 0.3% and more preferably 0.04% to 0.2%.
Al: 0.005% to 0.2%
Aluminum (Al) is a useful element to improve surface hardness and effective hardened case depth after nitrocarburizing and is therefore intentionally added. Al also yields a finer microstructure by inhibiting the growth of austenite grains during hot forging and is thus a useful element to improve toughness. Therefore, 0.005% or more is added. On the other hand, including over 0.2% does not increase this effect, but rather causes the disadvantage of higher component cost. Accordingly, the amount of Al added is 0.005% to 0.2%. The amount is preferably over 0.020% and 0.1% or less, and more preferably over 0.020% and 0.040% or less.
S: 0.06% or Less
Sulfur (S) forms MnS in the steel and is a useful element to improve the machinability by cutting. Including over 0.06%, however, lessens toughness. Accordingly, the amount of S added is 0.06% or less. The amount is preferably 0.04% or less. Note that for S to achieve the effect of improving machinability by cutting, the amount of S added is preferably 0.002% or more.
P: 0.02% or Less
Phosphorus (P) exists in a segregated manner at austenite grain boundaries and lowers the grain boundary strength, thereby lowering strength and toughness. Accordingly, the P content is preferably kept as low as possible, but a content of up to 0.02% is tolerable. The P content is therefore 0,02% or less. Note that setting the content of P to less than 0.001% requires a high cost. Therefore, it suffices in industrial terms to reduce the content of P to 0.001%.
B: 0.0003% to 0.01%
Boron (B) effectively promotes bainite phase formation. An added amount of B less than 0.0003% does not satisfactorily achieve this effect. On the other hand, adding over 0.01% does not increase this effect and only leads to higher component cost. Accordingly, the amount of B added is 0.0003% to 0.01%. The amount is preferably 0.0010% or more and 0.01% or less.
Note that to achieve the effect of promoting bainite phase formation, it is preferable that B be present in the steel as a solute. When solute N is present in the steel, however, the B in the steel is consumed by formation of BN. B does not contribute to improved quench hardenability when existing in the steel as BN. Accordingly, when solute N exists in the steel, B is preferably added in an amount greater than that consumed by formation of BN, and the amounts of B (% B) and of N (% N) in the steel preferably satisfy formula (1) below.
% B≧% N/14×10.8+0.0003. (1)
In the steel for nitrocarburizing, after subjection to forging or when improving machinability by cutting of the nitrocarburized material, one or more selected from the group of Pb≦0.2% and Bi≦0.02% may be added. Note that the desired effects achieved are not diminished regardless of whether these elements are added and regardless of their content.
Furthermore, in the steel for nitrocarburizing, the balance other than the above added elements consists of Fe and incidental impurities. In particular, however, Ti not only adversely affects strengthening by precipitation of V and Nb, but also lowers the core hardness and therefore is not to be included insofar as possible. The amount of Ti is preferably less than 0.010% and more preferably less than 0.005%.
3. Manufacturing Conditions
In
Specifically, in the steel bar manufacturing process (S1), a steel ingot is hot rolled into a steel bar and shipped after quality inspection. After shipping, the steel bar is transported (S2), and during the process (S3) of finishing the product (nitrocarburized component), the steel bar is cut to predetermined dimensions and subjected to hot forging or cold forging. After cutting the steel bar into a predetermined shape by drill boring, lathe turning or the like as necessary, nitrocarburizing treatment is performed, yielding the final product.
Alternatively, hot rolling material may be directly cut into a predetermined shape by lathe turning, drill boring or the like, with nitrocarburizing treatment then being performed to yield the final product. In the case of hot forging, cold straightening may be performed afterwards. Coating treatment such as painting or plating, may also be applied to the final product. Preferable manufacturing conditions will now be described.
Rolling Heating Temperature
The rolling heating temperature is preferably 950° C. to 1250° C. This range is adopted to cause carbides remaining after melting to be present as a solute during hot rolling, so as not to diminish forgeability due to formation of fine precipitates in the rolling material (the steel bar which is the material for the hot forging component).
In other words, when the rolling heating temperature is less than 950° C., it becomes difficult for the carbides remaining after melting to form a solute. On the other hand, a temperature exceeding 1250° C. facilitates coarsening of the crystal grains, thus reducing forgeability. Accordingly, the rolling heating temperature is preferably 950° C. to 1250° C.
Rolling Finishing Temperature
The rolling finishing temperature is preferably 800° C. or more. This temperature is adopted because at a rolling finishing temperature of less than 800° C., a ferrite phase forms. Particularly when the next process is nitrocarburizing after cold forging or cutting, such a ferrite phase is disadvantageous to obtain a bainite phase with an area ratio exceeding 50% of the matrix phase after nitrocarburizing. Moreover, at a rolling finishing temperature of less than 800° C., the rolling load increases, which degrades the out-of-roundness of the rolling material. Accordingly, the rolling finishing temperature is preferably 800° C. or more.
Cooling Rate
To prevent fine precipitates from forming before forging, thereby reducing forgeability, it is preferable to specify the cooling rate after rolling. In the precipitation temperature range of fine precipitates of 700° C. to 550° C., it is preferable to cool the steel bar faster than the critical cooling rate at which fine precipitates are produced (0.5° C./s).
Nitrocarburizing Treatment (Precipitation Treatment)
The resulting steel bar is then used as material that is forged and shaped into components by cutting and the like. Nitrocarburizing treatment is then performed. The temperature for nitrocarburizing treatment is preferably 550° C. to 700° C. to yield fine precipitates including V and Nb, and the treatment time is preferably 10 minutes or more. This range is adopted because at less than 550° C., insufficient precipitates are obtained, whereas over 700° C., the temperature enters the austenite region, making nitrocarburizing difficult. A more preferable range is 550° C. to 630° C. Furthermore, the treatment time is 10 minutes or more to obtain a sufficient amount of V and Nb precipitates.
Note that when hot forging is used, the hot forging is preferably performed with the heating temperature during hot forging at 950° C. to 1250° C., with the forging finishing temperature at 800° C. or more and the cooling rate after forging exceeding 0.5° C./s for the bainite phase to exceed 50% in area ratio of the matrix phase after nitrocarburizing and in order to prevent formation of fine precipitates from the standpoints of cold straightening and workability of cutting after hot forging.
EXAMPLESNext, our steels are further described by examples.
Steel samples with the composition shown in Table 1 (steel samples No. 1 to 17) were obtained by steelmaking in a 150 kg vacuum melting furnace, then rolling by heating at 1150° C., finishing at 970° C., and subsequently cooling to room temperature at a cooling rate of 0.9° C./s to prepare steel bars with o 50 mm. No. 17 is a conventional material, JIS SCr420. Note that P was not intentionally added to any of the steel samples in Table 1. Accordingly, the content of P in Table 1 indicates the amount mixed in as an incidental impurity. Furthermore, Ti was added to steel samples No. 14 and No. 15 but not intentionally added to steel samples No. 1 to 13 and No. 16 to 17 in Table 1. Accordingly, the content of Ti in steel samples No. 1 to 13 and No. 16 to 17 in Table 1 indicates the amount mixed in as an incidental impurity.
These materials were then heated to 1200° C. and subsequently hot forged at 1100° C. to a size of o 30 mm. The materials were cooled to room temperature at a cooling rate of 0.8° C./s, with a portion being cooled at 0.1° C./s for the sake of comparison.
The microstructure of the above materials was observed, hardness was measured, and machinability by cutting was tested. During microstructure observation, a cross-section was observed under an optical microscope, and the core microstructure was identified. For samples in which a bainite phase was present in the core, the area fraction of the bainite phase in the core was calculated. Machinability by cutting was assessed by a drill cutting test. Specifically, hot forging material was sliced to yield 20 mm thick pieces of test material in which through holes were bored in five locations per cross section using a JIS high-speed tool steel SKH51 straight drill with a 6 mm, under the following conditions: feed rate, 0.15 mm/rev; revolution speed, 795 rpm. Machinability by cutting was assessed by the total number of holes before the drill could no longer cut.
Hardness was measured by testing the hardness of the core using a Vickers hardness tester, with a test force of 100 g.
For steel samples No. 1 to 16, gas nitrocarburizing treatment was further applied to the hot forging material, and for steel sample No. 17, gas carburizing treatment was applied to the hot forging material. The gas nitrocarburizing treatment was performed by heating to 570° C. to 620° C. and retaining for 3.5 h under an atmosphere of NH3:N2:CO2=50:45:5. The gas carburizing treatment was performed by carburizing at 930° C. for 3 h, then oil quenching after retaining at 850° C. for 40 minutes and, furthermore, tempering at 170° C. for 1 h.
The microstructure of these heat treatment materials was observed, hardness measured, precipitates observed, and impact properties and fatigue properties tested.
During microstructure observation, a cross-section was observed under an optical microscope, and the core microstructure was identified. For samples in which a bainite phase was present in the core, the area fraction of the bainite phase was calculated.
To measure the hardness of the nitrocarburized material and the carburized material, the core hardness and surface hardness were measured. The surface hardness was measured at a position 0.02 mm from the surface, and the effective hardened case depth was measured as the depth from the surface at a hardness of HV 400. Samples for transmission electron microscopy observation were created from the cores of the nitrocarburized material and the carburized material by Twin-jet electropolishing. Precipitates were observed in the resulting samples using a transmission electron microscope with an acceleration voltage of 200 kV. Furthermore, the composition of the observed precipitates was calculated with an energy-dispersive X-ray spectrometer (EDX).
The assessment of impact properties was made by performing a Charpy impact test and calculating the impact value (J/cm2). Notched test pieces (R: 10 mm, depth: 2 mm) were used as test pieces. The notched test pieces were collected from the hot forging material, and after performing the above-described nitrocarburizing treatment or carburizing treatment, the collected test pieces were used in the Charpy impact test.
The assessment of fatigue properties was made by an Ono-type rotary bending fatigue test, and the fatigue limit was calculated. Notched test pieces (notch R: 1.0 mm; notch diameter: 8 mm; stress concentration factor: 1.8) were used as test pieces. The test pieces were collected from the hot forging material and, after the above-described nitrocarburizing treatment or carburizing treatment, were used in the fatigue test.
Table 2 shows the test results. No. 1 to 6 are our examples, No. 7 to 17 are comparative examples, and No. 18 is a conventional example provided by JIS SCr420 steel.
As is clear from Table 2, nitrocarburized materials No. 1 to 6 have better fatigue strength than the material resulting from carburizing, quenching, and tempering the conventional example (No. 18). As for workability of drill cutting, the material before nitrocarburizing treatment in No. 1 to 6 (hot forging material) has a level equivalent to or greater than the conventional material in practical terms. Furthermore, the results of transmission electron microscopy observation and of testing the precipitate composition by EDX confirm that the nitrocarburized materials No. 1 to 6 contain 500 or more fine precipitates, including V and Nb, with a grain size of less than 10 nm dispersed per 1 μm2 in the bainite phase. Based on these results, it can be concluded that our nitrocarburized material exhibits a high fatigue strength due to strengthening by precipitation based on the above fine precipitates.
By contrast, comparative examples No. 7 to 17 have a chemical composition or a resulting microstructure that are outside of our scope and thus have worse fatigue strength or drill workability.
In particular, No. 7 has low fatigue strength as compared to our examples due to the slow cooling rate after hot forging. For No. 7, the results of transmission electron microscopy observation showed no dispersion of fine precipitates with a grain size of less than 10 nm, whereas course precipitates with a grain size greatly exceeding 10 nm were observed. Based on these results, the coarseness of such resulting precipitates can be considered the cause of the reduction in fatigue strength. In other words, we believe that if the cooling rate after hot forging is slow and the desired bainite phase is not obtained, course precipitates are formed before nitrocarburizing. The amount of fine precipitates that form after nitrocarburizing treatment then decreases, resulting in insufficient strengthening by precipitation.
No. 8 includes a high amount of C, outside of our range. The hardness of the bainite phase therefore increases, reducing drill workability.
No. 9 includes high amounts of Si and Mn, outside of our range. The hardness of the hot forging material is therefore high, reducing the drill workability to approximately ⅕ that of conventional material.
No. 10 includes a low amount of Mn, outside of our range. A ferrite-pearlite microstructure thus forms before nitrocarburizing (after hot forging), lowering the area ratio of the bainite phase and forming V and Nb precipitates in the microstructure. The hardness before nitrocarburizing thus increases, reducing the drill workability.
No. 11 includes a low amount of Cr, outside of our range. A ferrite-pearlite microstructure thus forms before nitrocarburizing (after hot forging), lowering the area ratio of the bainite phase and forming V and Nb precipitates in the microstructure. The hardness before nitrocarburizing thus increases, reducing the drill workability.
No. 12 includes a low amount of Mo, outside of our range. Therefore, few fine precipitates exist after the nitrocarburizing treatment, and the resulting core hardness is insufficient. The fatigue strength is therefore lower than the conventional example.
No. 13 includes low amounts of V and Nb, outside of our range. Therefore, few fine precipitates exist after the nitrocarburizing treatment, and the resulting core hardness is insufficient. The fatigue strength is therefore lower than the conventional material.
No. 14 includes a low amount of Nb, outside of our range. Therefore, few fine precipitates exist after the nitrocarburizing treatment, and the resulting core hardness is insufficient. The fatigue strength is therefore lower than the conventional material.
Ti was added to No. 15 and No. 16, thus yielding few precipitates including V and Nb after the nitrocarburizing treatment. The resulting core hardness is therefore insufficient, and the fatigue strength is lower than the conventional material. Furthermore, the impact value is low.
No. 17 includes a low amount of Al, outside of our range. The surface hardness after the nitrocarburizing treatment and the effective hardened case depth are therefore insufficient, resulting in a lower fatigue strength than the conventional material.
Claims
1.-3. (canceled)
4. A steel for nitrocarburizing comprising, in mass %,
- C: 0.01% or more and less than 0.10%,
- Si: 1.0% or less,
- Mn: 0.5% to 3.0%,
- Cr: 0.30% to 3.0%,
- Mo: 0.005% to 0.4%,
- V: 0.02% to 0.5%,
- Nb: 0.003% to 0.15%,
- Al: 0.005% to 0.2%,
- S: 0.06% or less,
- P: 0.02% or less,
- B: 0.0003% to 0.01%, and
- the balance being Fe and incidental impurities, and
- including a microstructure with a bainite area ratio exceeding 50% before nitrocarburizing.
5. The steel according to claim 4, wherein after nitrocarburizing, precipitates including V and Nb are dispersed in a bainite phase.
6. A nitrocarburized component comprising a steel produced by nitrocarburizing the steel according to claim 4.
7. A nitrocarburized component comprising a steel produced by nitrocarburizing the steel according to claim 5.
Type: Application
Filed: Feb 15, 2013
Publication Date: Jan 22, 2015
Inventors: Takashi Iwamoto (Sendai), Keisuke Ando (Sendai), Kunikazu Tomita (Sendai), Yasuhiro Omori (Tokyo), Kiyoshi Uwai (Tokyo), Shinji Mitao (Tokyo)
Application Number: 14/378,553
International Classification: C23C 8/32 (20060101); C22C 38/32 (20060101); C22C 38/28 (20060101); C22C 38/26 (20060101); C22C 38/00 (20060101); C22C 38/22 (20060101); C22C 38/06 (20060101); C22C 38/04 (20060101); C22C 38/02 (20060101); C22C 38/38 (20060101); C22C 38/24 (20060101);
