JP4008391B2 - High strength steel with excellent hydrogen embrittlement resistance and method for producing the same - Google Patents

Description

  The present invention relates to a high-strength steel having excellent hydrogen embrittlement resistance against hydrogen embrittlement, which is one of the causes of fatigue life and delayed fracture, and a method for producing the same.

  Fatigue strength and delayed fracture resistance are properties required for high-strength springs and high-strength bolts used in automobiles, aircraft, various industrial machines, and the like. These properties are significantly degraded, especially when used in an environment where hydrogen penetrates into the material. Such deterioration of characteristics is called hydrogen embrittlement. Hydrogen embrittlement is a phenomenon in which hydrogen that has penetrated into steel gathers at a stress-concentrated part of the member macroscopically and at a location of a metal structure that tends to be the starting point of fracture such as a grain boundary microscopically. It is. For this reason, it is important to be excellent in hydrogen embrittlement resistance as a material characteristic of high-strength steels used as those materials.

  In order to improve hydrogen embrittlement resistance, there has been a conventional idea of capturing (trapping) hydrogen that has penetrated into steel and preventing diffusion of hydrogen in the steel. For example, the following techniques are known.

  In Patent Document 1 (Japanese Patent Laid-Open No. 2001-288539), as high-strength spring steel, V, Mo, which has activation energy for a predetermined hydrogen desorption and has a predetermined amount of hydrogen trap capacity. Steel having hydrogen trap sites made of oxides, carbides, nitrides, or composite precipitates containing one or more of Ti, Nb, and Zr, or a temperature range of 150 ° C. to 600 ° C. in temperature rising hydrogen analysis Describes a steel that has a release peak and has a predetermined amount of released hydrogen.

Patent Document 2 (Japanese Patent Laid-Open No. 10-17985) discloses that any one of Mo-based compounds, Ti-based compounds, V-based compounds, and composite compounds of 50 nm or less in the steel is 20 / (500 nm). A high-strength steel excellent in hydrogen embrittlement resistance, characterized by having ^two or more, has been proposed.

  As related technologies, Patent Document 3 (Japanese Patent Laid-Open No. 5-214484) includes C: 0.3 to 0.5%, Si: 1.0 to 3.0%, and Mn: 0.5 to 1. 0.5%, P: 0.025 or less, S: 0.03% or less, Ni: 0.1-2.0%, Cr: 0.5-1.0%, Mo: 0.1-0.5 %, V: 0.1 to 0.5%, Al as selective element: 0.01 to 0.03%, balance Fe and impurities, and high strength spring steel in which the number of coarse oxide inclusions is regulated Is disclosed. This technique is intended to improve the fatigue strength by regulating the amount of large oxide that becomes the starting point of fatigue fracture (paragraph 0030), and is not a technique for improving hydrogen embrittlement resistance.

JP 2001-288539 A (Claims) Japanese Patent Laid-Open No. 10-17985 (Claims) JP-A-5-214484 (Claims)

  According to the above patent document, by precipitating oxides, carbides, nitrides, etc., acting as hydrogen trap sites in the steel, hydrogen (free hydrogen) that has penetrated into the steel can be fixed. The hydrogen embrittlement resistance could be improved to some extent.

  However, in recent years, there has been an increasing demand for weight reduction and downsizing of members, and there has been a further demand for improved strength of steel materials. Along with this, further improvements in the fatigue life and delayed fracture resistance of members are further required. It is getting to be. For this reason, as for the hydrogen embrittlement resistance, there is a strong demand for characteristics superior to those of the prior art.

  The present invention has been made in view of such demands, and an object of the present invention is to provide a high-strength steel having hydrogen embrittlement resistance superior to that of the prior art and a method for producing the same.

  In order to improve the hydrogen embrittlement resistance, the present inventors diligently studied the kind and particle size of precipitates acting as hydrogen trap sites, and some precipitates are not necessarily effective as hydrogen trap sites. Moreover, the particle size is not necessarily fine, and it has been found that there are trap sites with weak trapping and strong trap sites depending on the composition (type) and particle size. It was also found that a synergistic effect can be expected by optimizing the combination of trap sites having different compositions and particle sizes, and is extremely effective for improving the hydrogen embrittlement resistance. In addition, the hydrogen trapping ability of the precipitate functioning as a trap site is based on the peak of the amount of hydrogen released at low and high temperatures in the hydrogen release behavior obtained by temperature analysis of steel charged with hydrogen. Confirmed to have. The present invention has been made based on such knowledge.

That is, the high-strength steel excellent in hydrogen embrittlement resistance of the present invention is mass%, C: 0.3 to 0.7%, Si: 0.1 to 3.0%, Mn: 0.01 to 1.8%, P: 0.01% or less, S: 0.01% or less, Cr: 0.2 to 1.8%, Ti: 0.004 to 0.20%, N: 0.0010 to 0 .0080%, V: 0.03 to 1.0%, Mo: 0.01 to 1.0% of one or two of them, the balance being Fe and impurities, A high strength steel including carbide, nitride, carbonitride as a product and having a tensile strength of 1300 MPa or more, and using a test piece of 10 mm × 10 mm × 2 mm in thickness as a cathode, 0.1 M NaOH at room temperature and 0 in a mixed aqueous solution of thiourea .05M, conducted cathode charge to be energized for 24 hours at a current density of 0.1 mA / cm ² Immediately after that, with respect to the amount of hydrogen obtained by sequentially measuring the amount of hydrogen released while heating the test piece at a heating rate of 12 ° C./min, the low-temperature region released hydrogen amount H1 released from room temperature to 350 ° C. is 1.0 to 10 mass ppm, and the amount of hydrogen H2 released in the high temperature range from 350 ° C. to 800 ° C. is 0.3 to 5 mass ppm. Measured after the ratio H3 / H1 of the low-temperature region released hydrogen amount H3 and the H1 measured after standing for a period of time is 0.05 or more, and after standing in the atmosphere at 80 ° C. for 24 hours after the cathodic charging. The ratio H4 / H2 between the H2 released hydrogen amount H4 and the H2 is 0.8 or more.

Moreover, the other high strength steel of this invention is mass%, C: 0.3-0.7%, Si: 0.1-3.0%, Mn: 0.01-1.8%, P : 0.01% or less, S: 0.01% or less, Cr: 0.2 to 1.8%, Ti: 0.004 to 0.20%, N: 0.0010 to 0.0080% In addition, V: 0.03 to 1.0%, Mo: contain one or two of 0.01 to 1.0%, the balance consists of Fe and impurities, and carbides and nitrides as precipitates , A high-strength steel containing carbonitride and having a tensile strength of 1300 MPa or more, and a particle size of 10 nm among carbides, nitrides, and carbonitrides excluding Fe-based carbides and Fe-Cr-based carbides from the precipitates. the following: containing 20 / (500 nm) ² or more, more than a particle size 30nm containing 10 / (500 nm) ² or more, or Average content of Ti in the precipitates or grain size 30nm is not less than 30 mass%, in which the average content of Ti in the precipitates following Tsubukei 10nm is less than 30 mass%.

As the alloy components of the steel of the present invention , in addition to the basic components, (1) Zr: 0.03 to 1.0%, Nb: 0.03 to 1.0%, Al: 0.001 to 0 .1% or more of 1%, (2) Cu: 0.02 to 0.5%, Ni: 0.02 to 2.0% of one or two elements of each group alone Or in combination.

  Further, in the method for producing high-strength steel of the present invention, the steel having the above components is melted, the melted steel is cast, the obtained cast piece is hot-rolled, and then quenched into a hot-rolled material, A method for producing a high-strength steel to be tempered, in the casting, a temperature range of 1500 ° C. to 1400 ° C. during solidification is 20 ° C./min or more, and a temperature range of 1400 ° C. to 1200 ° C. is 5 ° C. / Min., Cooled at an average cooling rate of 30 ° C./min or less, and during the hot rolling, the obtained cast piece was heated to 1000 ° C. to 1250 ° C., and hot-rolled at a rolling end temperature of 900 ° C. or higher, The temperature range from the rolling end temperature to 700 ° C is cooled at an average cooling rate of 100 ° C / min or higher, and the temperature range from 700 ° C to 400 ° C is limited to an average cooling rate of 50 ° C / min or lower to be cooled. In the case of the quenching and tempering, from room temperature to 600 ° C / Heat to a temperature of 840 ° C. to 1100 ° C. at the above heating rate, rapidly quench so that the time to be heated to a temperature of 840 ° C. or higher is 2 to 120 seconds, and further a heating rate of 600 ° C./min or higher. Is heated to a temperature of 400 ° C. to 600 ° C. and rapidly cooled so that the time of heating to 400 ° C. or higher is 2 seconds to 120 seconds.

  Another method for producing the high-strength steel of the present invention is to melt a steel having the above components, cast the melted steel, hot-roll the obtained cast piece, and then bak it on a hot-rolled material. A method for producing high-strength steel that is tempered and tempered. In the casting, a temperature range of 1500 ° C. to 1400 ° C. during solidification is 20 ° C./min or more, and a temperature range of 1400 ° C. to 1200 ° C. is 5 In the quenching and tempering, the quenching temperature Tq is 840 ° C <Tq <1300 ° C and Tq> 9500 / {6.72-log ( [% V] [% C])}-400 ° C., quenched from Tq, and then tempered by heating to a temperature of 500 ° C. or higher for 5 minutes or longer.

According to the high-strength steel of the present invention, in a temperature rising analysis immediately after hydrogen charging and after leaving at 80 ° C. for 24 hours, an effective combination of a trap site with a weak hydrogen trapping action and a strong trap site under a predetermined component . Defines the amount of low-temperature region released hydrogen, the amount of high-temperature region released hydrogen, and a predetermined ratio thereof, and defines the size and number of precipitates that become trap sites, and the average Ti content contained in these precipitates. Therefore, it has excellent hydrogen embrittlement resistance while having high strength. Moreover, according to the manufacturing method of this invention, the said high strength steel can be manufactured industrially easily .

  In the prior art (for example, Patent Document 1), when hydrogen thermal desorption analysis is performed, the hydrogen embrittlement resistance characteristics are focused only on that the amount of released hydrogen has a peak in a temperature range of 150 ° C. to 350 ° C. In the present invention, depending on the precipitate that captures hydrogen, the amount of released hydrogen has a peak not only in the low temperature range from room temperature to 350 ° C but also in the high temperature range from 350 ° C to 800 ° C. Attention was also paid to the fact that the peak of the released hydrogen amount in this high temperature region has an important role in suppressing hydrogen embrittlement. In other words, the amount of released hydrogen having a peak in the low temperature range is due to hydrogen desorbed from a relatively weak trapped state in the precipitate, while the amount of released hydrogen having a peak in the high temperature range is trapped relatively strongly. This is due to hydrogen desorbed from the state. By combining the two trap sites so that the weak trap site and the strong trap site compensate for each other while taking advantage of each other, excellent hydrogen embrittlement resistance can be exhibited.

  According to the inventor's research, the characteristics of the trap site with weak hydrogen capture and the strong trap site are as follows. Relatively weak trap sites are fine precipitates that are coherently precipitated, and therefore can be easily distributed in many steels, quickly trapping invading hydrogen and suppressing hydrogen embrittlement. However, since the trap is weak, hydrogen is diffused by using a hydrogen concentration gradient or temperature rise as a driving force, and this hydrogen may cause delayed fracture depending on the amount of hydrogen in the member and the stress state. On the other hand, strong trap sites are relatively large size Ti-based precipitates that are inconsistently precipitated, and it is difficult to distribute a large number of them in the steel. Since hydrogen is difficult to diffuse in a normal use state of the member, hydrogen embrittlement can be more stably prevented.

  In the present invention, hydrogen embrittlement is effectively suppressed by combining these trap sites with weak hydrogen capture and strong trap sites. In order to know the state of the optimal combination, how hydrogen actually enters the steel and trapped in the weak trap site and the strong trap site changes (releases) over time. Examined. In order for hydrogen trapped at weak trap sites not to adversely affect hydrogen embrittlement, it is desirable to slowly diffuse and decrease, while hydrogen trapped at strong trap sites should not change. Therefore, it is possible to achieve an effect by defining how much the amount of low-temperature released hydrogen with a peak of released hydrogen in a low-temperature region and how much the amount of high-temperature released hydrogen with a peak of released hydrogen in a high-temperature region changes over time. In particular, hydrogen embrittlement can be suppressed.

Specifically, using a test piece of 10 mm × 10 mm × 2 mm in thickness as a cathode, in a mixed aqueous solution of 0.1 M NaOH and 0.05 M thiourea at room temperature at a current density of 0.1 mA / cm ^2. Immediately after carrying out the cathode charging for 24 hours, the amount of hydrogen released while heating the test piece is sequentially measured, and the amount of released hydrogen H1 released from room temperature to 350 ° C., from 350 ° C. to 800 ° C. The amount of hydrogen H2 released in the high temperature range was calculated. Further, in order to know the state of change with time, the amount of hydrogen released at low temperature H3 and the amount of hydrogen released at high temperature H4 after standing for 24 hours in the atmosphere at 80 ° C. after the cathode charge were obtained, and H1, H2 and H3 The optimum values of / H1 and H4 / H2 were determined from the viewpoint of improving hydrogen embrittlement resistance. The low temperature range is from room temperature to 350 ° C., and the high temperature range is from 350 ° C. to 800 ° C. As is clear from the examples described later, a peak is observed in the amount of released hydrogen at each boundary from 350 ° C. Because it is done.

  H1 indicates the amount of hydrogen trapped relatively weakly, and the trap sites contributing to H1 are precipitates of, for example, Mo and V that are coherently precipitated in the base material, so that it is easy to disperse the precipitates in large quantities. Therefore, securing a certain amount of H1 is effective in improving the hydrogen embrittlement resistance. On the other hand, if the amount is too large, the hydrogen at the crystal grain boundary, which is the starting point of hydrogen cracking, increases, so that the hydrogen embrittlement characteristics deteriorate. For this reason, H1 shall be 1.0-10 mass ppm.

  The H2 indicates the amount of hydrogen trapped relatively strongly, and if it is a certain amount or more, it means that there are sufficient strong trap sites and contributes to prevention of hydrogen embrittlement. However, too much is not good because the hydrogen at the grain boundaries also increases. For this reason, H2 shall be 0.3-5 mass ppm.

  H3 / H1 is an index indicating that charged hydrogen is diffused and released in steel even at a low temperature. When H3 is reduced to less than 0.05 × H1, it is clear from the examples described later. Thus, the hydrogen trapping effect by the weak trap site is too small, and hydrogen embrittlement becomes apparent as a whole. For this reason, H3 / H1 is specified to be 0.05 or more. Although the upper limit is not particularly defined, it is usually about 0.6 or less in many cases.

  Further, H4 / H2 indicates that the closer to 1, the stronger the trap site does not change over time, and as is clear from the examples, if it is less than 0.80, the release of hydrogen once trapped becomes excessive. As a whole, hydrogen embrittlement becomes apparent. For this reason, H4 / H2 is specified to be 0.80 or more.

According to the inventor's research, it has been found that, in addition to oxides, Fe-based carbides and Fe-Cr-based carbides are not effective as hydrogen traps or are too small to contribute to the hydrogen release characteristics. In addition, it was found that the trap site with weak hydrogen trapping is a fine precipitate having a particle size of 10 nm or less, which is generated by, for example, co-precipitation of V or Mo carbide with the base material. For this reason, it is possible to disperse these precipitates in large quantities. On the other hand, strong trap sites cannot be obtained by precipitation of carbides of V and Mo, but are limited to precipitates containing Ti, and fine ones that are aligned with the base material are weak trap sites. It was found that this was a Ti-based precipitate which was enlarged to some extent with a particle size of 30 nm or more and precipitated in inconsistent precipitation.
And in order to implement | achieve the discharge | release characteristic of the said hydrogen, as evident from the below-mentioned Example, among carbide | carbonized_materials, nitride, and carbonitride except Fe type carbide and Fe-Cr type carbide, the particle size is 10 nm or less. It is necessary to contain 20 / (500 nm) ² or more, and 10 / (500 nm) ² or more of particles having a particle size of 30 nm or more. Furthermore, since the precipitate having a particle size of 10 nm or less is preferably a non-Ti precipitate such as V or Mo, the average content of Ti in the precipitate is less than 30 mass%, while the particle size is 30 nm or more. Therefore, the average content of Ti in these precipitates is 30 mass% or more. The method for measuring the size (particle size) and number of the precipitates and the method for calculating the average content of Ti will be described in Examples.

  The composition of the high-strength steel excellent in hydrogen embrittlement resistance of the present invention is appropriately adjusted so that the tensile strength after quenching and tempering is 1300 MPa or more, preferably 1500 MPa or more. The steel components that can provide tensile strength and can easily generate carbides, nitrides, and carbonitrides that are effective as hydrogen release characteristics and trap sites, and the reasons for their limitation (unit mass%) will be described in detail below.

C: 0.3 to 0.7%
C forms carbides and carbonitrides together with Fe, Ti, V, Mo, Zr, etc., and has an important function to determine the strength and hydrogen embrittlement resistance as a high-strength steel, but less than 0.3% When the content exceeds 0.7%, the carbides become coarse and the toughness also decreases. Therefore, the lower limit of the C content is set to 0.3% and the upper limit is set to 0.7%.

Si: 0.1-3.0%
In addition to acting as a deoxidizing element, Si has an effect of improving the strength by solid solution in Fe, but if it is less than 0.1%, no effect appears, and if it exceeds 3.0%, the workability decreases. For this reason, the lower limit of Si amount was set to 0.1%, and the upper limit was set to 3.0%.

Mn: 0.01 to 1.8%
Mn combines with S, which causes embrittlement of the steel to make S harmless, and also increases the hardenability of the steel and contributes to high strength. However, if it is less than 0.01%, these effects do not appear. If it exceeds 8%, the toughness is conversely reduced, so the lower limit of the amount of Mn is set to 0.01% and the upper limit is set to 1.8%.

P: 0.01% or less P is more preferable as it causes embrittlement of the steel, and high strength steel has a large effect. In particular, even in a high-strength steel of 1200 MPa or more, the range of 0.01% or less was set as a range where no practical adverse effect occurs.

S: 0.01% or less Since S is also an embrittlement element like P, the smaller the content, the better.

Cr: 0.2 to 1.8%
Cr is an essential element for improving the strength of the steel. However, if it is less than 0.2%, its effect is insufficient, and if it exceeds 1.8%, the toughness decreases. The upper limit was 2% and 1.8%.

Ti: 0.004 to 0.20 %
Ti combines with C and N to form carbonitrides to form hydrogen trap sites, and exerts the effect of improving hydrogen embrittlement resistance if the size and the like are controlled. However, if it is less than 0.004 %, the effect cannot be obtained, and if it is 0.20 % or more, too coarse carbonitride is formed to lower toughness. For this reason, the lower limit of the Ti amount is 0.004 %, preferably 0.02%, and the upper limit is 0.20 %, preferably 0.1%.

N: 0.0010 to 0.0080%
N forms carbonitrides and nitrides together with Ti, V, Al, etc., and functions as pinning particles for hydrogen trap sites and crystal grain size refinement. However, if it is less than 0.0010%, the effect is too small. On the other hand, if it exceeds 0.0080%, coarse nitrides are generated, and the number of precipitates having a desired hydrogen trap effect is reduced. The lower limit of the N amount is 0.0010%, and the upper limit is 0.0080%.

V: 0.03 to 1.0%, Mo: one or two of 0.01 to 1.0% V forms fine carbides and acts as a hydrogen trap site, but less than 0.03% However, the effect does not appear, and if it exceeds 1.0%, coarse carbides are easily formed and the effect becomes insufficient. For this reason, the lower limit of the V amount is 0.03%, more preferably 0.1%, and the upper limit is 1.0%, more preferably 0.5%.

  Mo also forms fine carbides and works as a hydrogen trap site. However, if it is less than 0.03%, the effect does not appear, and if it exceeds 1.0%, coarse carbides are easily formed and the effect becomes insufficient. For this reason, the lower limit of the Mo amount is 0.03%, more preferably 0.1%, and the upper limit is 1.0%, more preferably 0.6%.

  In addition to the basic components described above, the steel of the present invention is formed by the remaining Fe and impurities, but does not impair the action and effect of the basic components and does not hinder the addition of elements that improve the characteristics. For example, (1) any one or more of Zr, Nb, and Al in the following ranges, and / or (2) any one or two of Cu, Ni in the following ranges can be added.

Zr: 0.03 to 1.0%
Zr forms fine carbides and acts as a hydrogen trap site. However, if it is less than 0.03%, the effect does not appear, and if it exceeds 1.0%, the effect is saturated and the cost is increased. The lower limit is 0.03%, and the upper limit is 1.0%.

Nb: 0.03-1.0%
Nb forms carbonitrides and has a crystal grain refining effect. However, if it is less than 0.03%, no effect appears, and if it exceeds 1.0%, coarse carbides increase and the effect is saturated. For this reason, the lower limit of the Nb amount is 0.03%, and the upper limit is 1.0%.

Al: 0.001 to 0.1%
Al acts as a deoxidizing element and also forms a precipitate by combining with N, and may be added because it is effective in refining the crystal grain size of steel. However, if it is less than 0.001%, no effect is seen. If it exceeds 0.1%, a large amount of alumina-based oxides are formed and the fatigue characteristics are deteriorated. For this reason, the lower limit of the Al amount is 0.001%, and the upper limit is 0.1%.

Cu: 0.02 to 0.5%
Since Cu has an effect of improving the corrosion resistance, it may be added. However, if less than 0.02%, no effect is seen, and if over 0.5%, hot workability deteriorates and a problem occurs during rolling. The lower limit of Cu content is 0.02%, and the upper limit is 0.5%.

Ni: 0.02 to 2.0%
Ni also has the effect of improving the corrosion resistance, so it may be added, but if 0.02% or less, no effect is seen, and if it exceeds 2.0%, the amount of retained austenite generated increases when tempered and tempered. Since the homogeneity of the structure is lowered, the lower limit of Ni content is 0.02%, and the upper limit is 2.0%.

Next, the manufacturing method of the high strength steel rolling material which has the said component as this invention steel is demonstrated.
The steel of the present invention is produced by melting the steel of the above components, casting the molten steel, hot rolling the obtained cast piece, and subsequently quenching and tempering the hot rolled material. . The rolling may be plate rolling or wire rolling, and the rolling shape can be appropriately determined according to the cross-sectional shape of the target steel material.

In casting the steel, the temperature range of 1500 ° C. to 1400 ° C. during solidification is 20 ° C./min or more, preferably 25 ° C./min or more, and the temperature range of 1400 ° C. to 1200 ° C. is 5 ° C./min or more, 30 It is cooled at an average cooling rate of not more than 10 ° C / min, preferably not more than 20 ° C / min, more preferably not more than 10 ° C / min or less than 10 ° C / min. In steels containing Ti, most of Ti carbides, nitrides and carbonitrides (hereinafter sometimes simply referred to as carbonitrides) are produced during cooling from 1500 ° C. to 1200 ° C. during solidification. However, the size of Ti carbonitride varies depending on the temperature at which it is produced. In the temperature range of 1500 ° C. to 1400 ° C., coarse Ti carbonitride exceeding 100 nm is likely to be formed. Therefore, this temperature range is cooled at a fast cooling rate of 20 ° C./min or more to suppress the formation of coarse carbonitride. In addition, by cooling the temperature range of 1400 ° C. to 1200 ° C. at a moderate average rate of 30 ° C./min or less (preferably an average cooling rate lower than the average cooling rate in the temperature range of 1500 ° C. to 1400 ° C.), 30 nm to A large amount of Ti carbonitride of about 100 nm, preferably about 30 nm to 50 nm is produced. That is, when a temperature range of 1500 ° C. to 1400 ° C. is cooled at a slow cooling rate of less than 20 ° C./min, coarse carbonitrides are generated and the number is drastically reduced, and the number of carbonitrides of 30 nm or more is 10/500 nm ^2. It becomes difficult to do it above. Further, even when the temperature range of 1400 ° C. to 1200 ° C. is cooled at a cooling rate slower than 5 ° C./min, coarse carbonitride exceeding 100 nm is generated, and the number of carbonitrides of 30 nm or more is reduced. On the other hand, when the temperature range of 1400 ° C. to 1200 ° C. is cooled at a cooling rate faster than 30 ° C./min, the amount of excessively fine carbonitride increases or the cooling is completed before carbonitride is formed. End up. For this reason, as described above, the temperature range of 1500 ° C. to 1400 ° C. and the temperature range of 1400 ° C. to 1200 ° C. are cooled at a predetermined average cooling rate. Conventionally, the cooling rate after casting of this type of high-strength steel is about 4 to 8 ° C./min on average from 1500 ° C. to 1200 ° C., and the cooling rate is not particularly controlled by the temperature range.

  After casting, the cast piece is heated from 1000 ° C. to 1250 ° C. in order to dissolve carbonitride-forming elements during hot rolling. Thereafter, hot rolling is performed, but the cooling rate is controlled as follows in order to precipitate fine carbonitrides of V and Mo during cooling during rolling. When the average cooling rate from the rolling finish temperature to 700 ° C. is 100 ° C./min or less, carbonitride precipitates at a high temperature up to 700 ° C., and therefore fine precipitates of 10 nm or less become too small. Cool quickly as above. Carbonitride tends to be a fine precipitate when it is deposited in a temperature range of 700 ° C to 400 ° C. In order to ensure the amount of fine precipitates of 10 nm or less, the average cooling rate in this temperature range is regulated to 50 ° C./min or less, desirably 30 ° C./min or less. The rolling end temperature (finish rolling end temperature) is 900 ° C. or higher because rolling cracks are likely to occur when the temperature is less than 900 ° C. The upper limit of the rolling end temperature is not particularly limited, but is naturally determined by the heating temperature, and is usually about 1200 ° C. or lower.

  After hot rolling, the hot-rolled material is quenched and tempered, but in quenching, it is quenched unless heated so that the quenching temperature Tq (gammaization temperature) is 840 ° C. or more and the holding time is 2 seconds or more. Is insufficient, and a martensite-based organization cannot be obtained. Further, when the quenching temperature exceeds 1100 ° C. or the heating time exceeds 120 seconds, fine precipitates of 20 nm or less among the precipitates whose particle size is controlled are either re-dissolved or coarsened. Therefore, it is heated to a temperature range of 840 ° C. or higher and 1100 ° C. or lower, and quenched rapidly so that the time for heating to a temperature of 840 ° C. or higher is 2 seconds or longer and 120 seconds or shorter. In addition, if the rate of temperature rise to the quenching temperature is slow, reprecipitation and coarsening of precipitates occur in the high temperature region, so the rate of temperature rise is set to 600 ° C./min or more.

  Further, if the tempering temperature is less than 400 ° C, the steel is too hard and the toughness is lowered, and if it exceeds 600 ° C, it becomes too soft and sufficient strength cannot be obtained, so the heating temperature (tempering temperature) is 400 ° C. The temperature is 600 ° C. or lower. Further, if the heating time is too short, the toughness is lowered, and if it is too long, sufficient strength cannot be obtained. Therefore, the time for heating to 400 ° C. or higher is set to 2 seconds or more and 120 seconds or less, and then cooled quickly. Further, if the rate of temperature increase to the tempering temperature is slow, the residence time in the high temperature region becomes long, so that the strength is also lowered. For this reason, the temperature rising rate is set to 600 ° C./min or more. The cooling rate after tempering is not particularly limited, but the cooling rate is preferably 300 ° C./min or more.

  The production conditions described above are intended to control the precipitation of fine carbonitrides in the hot rolling process, but the production of fine precipitates can be controlled in the quenching and tempering processes as described below. In this case, the hot rolling before the quenching process and the subsequent cooling conditions are not particularly limited, and may be performed in accordance with an ordinary method. Moreover, you may give an appropriate plastic processing after a hot rolling, for example, wire drawing. In addition, 30-100 nm Ti type carbide | carbonized_material is ensured by cooling the 1500 degreeC-1400 degreeC and the 1400 degreeC-1200 degreeC temperature range at the time of solidification at the predetermined | prescribed average cooling rate as above-mentioned.

  In order to control the formation of precipitates in the quenching and tempering steps, the conditions may be selected so that the alloy carbides are re-dissolved and reprecipitated during quenching and tempering to obtain fine precipitates. That is, quenching is performed under conditions where the carbides of V and Mo are re-dissolved to some extent, and tempering is performed at a temperature at which fine carbides are precipitated. Since it is known that Mo carbide is re-dissolved at a relatively low temperature, the quenching temperature Tq is at least 840 ° C. or higher. On the other hand, if Tq exceeds 1300 ° C., Ti-based precipitates having a large particle size will be re-dissolved, so it is necessary to keep the temperature below 1300 ° C.

Also, it is necessary to heat and quench the carbides of V and Zr to a temperature at which they can be dissolved to some extent. The temperature at which the carbide of V can be dissolved is the solubility product equation (9500 = T {6.72−log ([% V] [%]) reported by Narita and Oyama (Kobe Steel Technical Report, 67 (1966), 179). C])}, T: absolute temperature) can be calculated, but it is 127 from the temperature calculated by the solubility product formula (temperature at which the solid solution is completely dissolved). It was found that a sufficient amount of carbide was dissolved by setting the temperature to be lower than the temperature. For this reason, the quenching temperature Tq is set on the premise of 840 ° C. <Tq <1300 ° C. and so as to satisfy the following equation. In addition, since the temperature calculated | required by the formula of a solubility product is absolute temperature, in the following formula, 127 + 273 = 400 is subtracted and converted into ° C.
Tq> 9500 / {6.72-log ([% V] [% C])}-400 ° C.

  Tempering is performed in order to ensure toughness and to precipitate carbonitrides of alloy elements that have been dissolved during quenching. When the tempering temperature is less than 500 ° C. and less than 5 minutes, fine carbides do not precipitate, so the temperature is set to 500 ° C. or more and 5 minutes or more. The upper limit may be set as appropriate within a range that does not cause a decrease in strength, but it is usually about 600 ° C. or less and about 10 minutes or less depending on the desired tensile strength. Moreover, the cooling rate after heating and holding at the tempering temperature is not particularly limited.

  EXAMPLES Next, although an Example is given and this invention is demonstrated more concretely, this invention is not limitedly interpreted by this Example.

  Steels having the components shown in Table 1 were melted using a high-frequency induction vacuum melting furnace to produce a cast piece (ingot) of about 30 kg. Table 2 shows the cooling rate after casting. The cooling rate was controlled by combining the shape of the mold, heat retention and the like. Thereafter, the cast piece was heated to 1200 ° C. and hot-rolled to a thickness of 6 mm. Tables 2 and 3 also show the rolling end temperature in hot rolling and the cooling rate after rolling (from each rolling end temperature to 700 ° C. and each cooling rate from 700 ° C. to 400 ° C.). These cooling rates were controlled by combining air cooling, heat retention, and the like.

  A plate having a thickness of 2.2 mm, a width of 11 mm, and a length of 120 mm was cut out from the rolled material, and, according to the quenching conditions shown in Tables 2 and 3, heated and held for a predetermined time by energization heating, and then rapidly cooled by gas. Then, after heating and holding at the tempering temperature shown in the same table for a predetermined time, it was air-cooled. The hardness of the steel plate after tempering was about 500 to 650 in terms of Vickers hardness. Moreover, when the tensile strength of the steel plate after tempering was measured by a tensile test, it was about 1300 to 2000 MPa.

  From the steel plate after quenching and tempering, a hydrogen analysis test piece of 10 mm × 10 mm × 2 mm thickness and a 4-point bending cathode charge test piece of 1.5 mm thickness × 10 mm width × 65 mm length were cut out for hydrogen analysis. The hydrogen embrittlement resistance was investigated.

The hydrogen analysis was performed as follows. Immediately after the cathode charge was conducted by energizing the specimen for hydrogen analysis for 48 hours at a current density of 0.1 mA / cm ² in a mixed aqueous solution of 0.1 M NaOH and 0.05 M thiourea at room temperature. The amount of hydrogen released while heating the specimen from room temperature to 800 ° C. at 12 ° C./min using API-MS (atmospheric pressure ionization mass spectrometer, manufactured by East Japan Semiconductor Technology, Inc., UG-240APN) (mass ppm) Were measured sequentially, and the amount of hydrogen H1 released from the room temperature to 350 ° C. and the amount of hydrogen H2 released from 350 ° C. to 800 ° C. were determined. In addition, after hydrogen was charged under the same conditions and left in the atmosphere at 80 ° C. for 24 hours, the hydrogen analysis was performed in the same manner, and the amount of released hydrogen H3 was high temperature hydrogen released from 350 ° C. to 800 ° C. The amount H4 was determined, and the values of H3 / H1 and H4 / H2 were determined. An example of hydrogen release characteristics (sample No. 8) measured by API-MS is shown in FIG. In FIG. 1, the horizontal axis indicates the temperature, but the elapsed time is 0.2 seconds per 1 ° C. (60 seconds / 12 ° C. = 0.2 seconds / ° C.), and the amount of released hydrogen is in the low temperature range and high temperature range. The area enclosed by the temperature axis and the curve can be obtained.

  The hydrogen embrittlement resistance was evaluated by conducting a four-point bending cathode charge life test in the following manner, measuring the time until fracture, and measuring the time. As shown in FIG. 2, the test piece S is set in a protruding shape at both ends via four fulcrums 3 arranged symmetrically between the support frame 1 and the push-up stand 2, and the push-up stand 2 is placed on the support frame 1 side. In a state where a maximum stress of 1400 MPa was applied to the test piece, the test piece was used as a cathode in a mixed solution of 0.5 mol / L sulfuric acid and 0.01 mol / L KSCN (potassium thiocyanate). Using a platinum electrode, a cathode potential of −700 mV was applied, hydrogen was added to the test piece, and the time until breaking after applying the potential was measured.

  In addition, specimens for microstructure observation were collected from the steel plates after quenching and tempering, and 5 fields of 500 nm × 500 nm were observed at 150,000 times each using a TEM (transmission electron microscope) by the extraction replica method. , Count the number of carbides (excluding Fe-based carbides and Fe-Cr-based carbides), nitrides and carbonitride-generated products of a given size (crystallized products and precipitates) and average Asked. Specifically, since the product compounds observed in each field of view of 150,000 times are almost spherical, all those having a diameter of 10 nm or less and 30 nm or more were analyzed using an EDX (energy dispersive detection device). The average composition including the average content of Ti is determined, whether or not the carbon is a predetermined carbide, nitride, or carbonitride is determined, and the average number of these carbonitrides having a diameter of 10 nm or less and 30 nm or more is determined. It was. These measurement results are also shown in Table 4. In the table, “the number of precipitates” includes the number of crystallized substances.

  From Table 4, Sample Nos. 1-4, 7-10, 13, and 14 of the invention examples all satisfy the basic components according to the present invention, and the production conditions are also appropriate. The conditions of the present invention are satisfied, and the results of the 4-point bending cathode charge life test are all 1000 seconds or more, indicating excellent hydrogen embrittlement resistance. On the other hand, although the manufacturing conditions are appropriate, Sample No. 5 has a small amount of Ti, No. 6 has a large amount of Ti, No. 11 has a small amount of V, and No. 12 has a large amount of V. No. 15 has too much Mo, so the hydrogen analysis result does not satisfy the conditions of the present invention, and the precipitate size is also out of the scope of the present invention, resulting in poor hydrogen embrittlement resistance. Sample Nos. 16 to 21 (invention examples) are obtained by adding appropriate amounts of Al, Zr, Nb, Cu, and Ni in addition to V and Mo, and all show excellent hydrogen embrittlement resistance. Yes.

  On the other hand, Samples Nos. 22 to 33 all have the inventive components, but Nos. 22 to 25 are manufactured under appropriate manufacturing conditions, especially the cooling rate during solidification is changed. It is. No. 22 (Comparative Example) has a slow cooling rate in the temperature range of 1500 ° C. to 1400 ° C., and a cooling rate in the temperature range of 1400 ° C. to 1200 ° C. is too fast. The number of precipitates of 30 nm or more is reduced, and the hydrogen analysis results show that the amount of hydrogen in H2 is insufficient and the hydrogen embrittlement resistance is deteriorated. Samples Nos. 26 to 28 are obtained by changing the cooling rate of 700 ° C. to 400 ° C. after rolling. In No. 28 (comparative example), the cooling rate is too fast, so that the precipitation of fine carbides of 10 nm or less is insufficient, and the hydrogen analysis results show that the amount of hydrogen in H1 is too small, so the hydrogen embrittlement resistance is significantly deteriorated. . Sample No. 29 has a particularly slow cooling rate up to 700 ° C. after the end of rolling, so that carbides containing a large amount of V are coarsely deposited, resulting in insufficient amount of precipitates of 10 nm or less, and the hydrogen analysis results are also H1. Since the amount of hydrogen became too small, the hydrogen embrittlement resistance deteriorated. Sample No. 30 had a heating time during quenching that was too long and fine carbides were re-dissolved, so the amount of precipitates of 10 nm or less was insufficient, and the hydrogen analysis results showed that the amount of hydrogen in H1 was too small. Therefore, the hydrogen embrittlement resistance is deteriorated.

  Samples Nos. 31 to 33 are those in which the formation of fine precipitates is controlled in the quenching and tempering processes. Samples Nos. 31 and 32 (invention examples) have fine rolling conditions in the hot rolling process. Although the conditions for generating precipitates are not satisfied, but the subsequent heat treatment conditions are appropriate, the size and number of precipitated carbides satisfy the conditions of the present invention, and the hydrogen analysis results also satisfy the conditions of the present invention and are excellent. Furthermore, hydrogen embrittlement resistance is obtained. On the other hand, Sample No. 33 has a low tempering temperature and a short tempering time, so that the precipitation of fine carbides is insufficient and the hydrogen embrittlement resistance is poor.

It is a graph which shows an example of the hydrogen temperature rising analysis result in an Example. It is explanatory drawing of the implementation point of the 4-point bending cathode charge life test in an Example.

Claims (6)

mass%
C: 0.3-0.7%
Si: 0.1 to 3.0%,
Mn: 0.01 to 1.8%,
P: 0.01% or less,
S: 0.01% or less,
Cr: 0.2 to 1.8%,
Ti: 0.004 to 0.20%,
N: containing 0.0010 to 0.0080%,
V: 0.03 to 1.0%, Mo: contain one or two of 0.01 to 1.0%, the balance consists of Fe and impurities, and precipitates are carbide, nitride, charcoal A high strength steel containing nitride and having a tensile strength of 1300 MPa or more,
Cathode energized at a current density of 0.1 mA / cm ² for 24 hours in a mixed aqueous solution of 0.1 M NaOH and 0.05 M thiourea at room temperature using a test piece of 10 mm × 10 mm × 2 mm thickness as a cathode Immediately after the charging, the hydrogen amount obtained by sequentially measuring the amount of hydrogen released while heating the test piece at a heating rate of 12 ° C./min. Is released in a low temperature range from room temperature to 350 ° C. The amount of hydrogen H1 is 1.0 to 10 mass ppm, the amount of hydrogen H2 released in the high temperature range from 350 ° C. to 800 ° C. is 0.3 to 5 mass ppm,
The ratio H3 / H1 of the amount H3 of the low-temperature region released hydrogen and the H1 measured after standing for 24 hours in the atmosphere at 80 ° C. after the cathode charge is 0.05 or more, and the cathode charge is performed. Further, the hydrogen embrittlement resistance is characterized in that the ratio H4 / H2 of the amount H4 of hydrogen released in the high temperature range and H2 measured after standing for 24 hours in the atmosphere at 80 ° C. is 0.8 or more. High strength steel. mass%
C: 0.3-0.7%
Si: 0.1 to 3.0%,
Mn: 0.01 to 1.8%,
P: 0.01% or less,
S: 0.01% or less,
Cr: 0.2 to 1.8%,
Ti: 0.004 to 0.20%,
N: containing 0.0010 to 0.0080%,
V: 0.03 to 1.0%, Mo: contain one or two of 0.01 to 1.0%, the balance consists of Fe and impurities, and precipitates are carbide, nitride, charcoal A high strength steel containing nitride and having a tensile strength of 1300 MPa or more,
Of the precipitates, carbides, nitrides, and carbonitrides excluding Fe-based carbides and Fe-Cr-based carbides, containing 20 particles / (500 nm) ² or more having a particle size of 10 nm or less, and having a particle size of 30 nm or more 10 / (500 nm) ² or more, and the average content of Ti in the precipitate having a particle size of 30 nm or more is 30 mass% or more, and the average content of Ti in the precipitate having a particle size of 10 nm or less Is a high-strength steel excellent in hydrogen embrittlement resistance, characterized by being less than 30 mass%. Furthermore, at mass%,
Zr: 0.03-1.0%,
Nb: 0.03-1.0%,
Al: 0.001 to 0.1%
The high-strength steel according to claim 1 or 2 , containing one or more of them. Furthermore, at mass%,
Cu: 0.02 to 0.5%,
Ni: 0.02 to 2.0%
High strength steel given in any 1 paragraph of Claims 1-3 containing 1 type or 2 types among these. A steel having the components described in any one of claims 1 to 4 is melted, the melted steel is cast, the obtained cast piece is hot-rolled, and then quenched into a hot-rolled material. A method for producing high strength steel to which reversion is performed,
During the casting, the temperature range of 1500 ° C. to 1400 ° C. during solidification is 20 ° C./min or more, and the temperature range of 1400 ° C. to 1200 ° C. is cooled at an average cooling rate of 5 ° C./min to 30 ° C./min. And
At the time of the hot rolling, the obtained cast piece is heated to 1000 ° C. to 1250 ° C., hot rolled at a rolling end temperature of 900 ° C. or higher, and a temperature range from the rolling end temperature to 700 ° C. is 100 ° C./min or higher. The cooling is performed at an average cooling rate of 700 ° C. to 400 ° C. and limited to an average cooling rate of 50 ° C./min or less.
During the quenching and tempering, heating is performed from room temperature to a temperature of 840 ° C. to 1100 ° C. at a heating rate of 600 ° C./min or more, and the heating time to a temperature of 840 ° C. or more is quickly 2 seconds to 120 seconds. And further cooled to a temperature of 400 ° C. to 600 ° C. at a heating rate of 600 ° C./min or more and rapidly cooled so that the time of heating to 400 ° C. or more is 2 seconds to 120 seconds. that, the method of producing a high strength steel excellent in hydrogen embrittlement resistance. A steel having the components described in any one of claims 1 to 4 is melted, the melted steel is cast, the obtained cast piece is hot-rolled, and then quenched into a hot-rolled material. A method for producing high strength steel to which reversion is performed,
During the casting, the temperature range of 1500 ° C. to 1400 ° C. during solidification is 20 ° C./min or more, and the temperature range of 1400 ° C. to 1200 ° C. is cooled at an average cooling rate of 5 ° C./min to 30 ° C./min. And
In the quenching and tempering, the quenching temperature Tq is set to 840 ° C. <Tq <1300 ° C. and Tq> 9500 / {6.72−log ([% V] [% C])} − 400 ° C. And then tempering by heating to a temperature of 500 ° C. or higher for 5 minutes or longer , and a method for producing high-strength steel excellent in hydrogen embrittlement resistance.

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