摘要
The effect of deep cryogenic treatment on the formation of reversed austenite （RA） in super martensitic stainless steel was investigated. RA was found to form in steels without （A） and with （B） deep cryogenic treatment. The volume fraction of RA initially increased and then decreased with increasing tempering temperature over 550-- 750 ℃ for the two steels, which were quenched at 1050 ℃. In addition, for both with and without deep cryogenic treatment, the RA content reached a maximum value at 650 ℃ although the RA content in steel B was greater than that in steel A over the entire range of tempering temperatures. Furthermore, the hardness （HRC） of steel B was greater than that of steel A at tempering temperatures of 550--750 ℃. From these results, the basic mechanism for the formation of RA in steels A and B was determined to be Ni diffusion. However, there were more Ni enriched points, a lower degree of enrichment, and a shorter diffusion path in steel B. It needed to be noted that the shapes of the RA consisted of blocks and stripes in both steels. These shapes resulted because the RA redissolved and trans- formed to martensite along the martensitic lath boundaries when the tempering temperature was 650--750 ℃, and a portion of RA in the martensitie lath divided the originally wider martensitic laths into a number of thinner ones. In- terestingly, the RA redissolved more rapidly in steel B and consequently resulted in a stronger refining effect.
The effect of deep cryogenic treatment on the formation of reversed austenite （RA） in super martensitic stainless steel was investigated. RA was found to form in steels without （A） and with （B） deep cryogenic treatment. The volume fraction of RA initially increased and then decreased with increasing tempering temperature over 550-- 750 ℃ for the two steels, which were quenched at 1050 ℃. In addition, for both with and without deep cryogenic treatment, the RA content reached a maximum value at 650 ℃ although the RA content in steel B was greater than that in steel A over the entire range of tempering temperatures. Furthermore, the hardness （HRC） of steel B was greater than that of steel A at tempering temperatures of 550--750 ℃. From these results, the basic mechanism for the formation of RA in steels A and B was determined to be Ni diffusion. However, there were more Ni enriched points, a lower degree of enrichment, and a shorter diffusion path in steel B. It needed to be noted that the shapes of the RA consisted of blocks and stripes in both steels. These shapes resulted because the RA redissolved and trans- formed to martensite along the martensitic lath boundaries when the tempering temperature was 650--750 ℃, and a portion of RA in the martensitie lath divided the originally wider martensitic laths into a number of thinner ones. In- terestingly, the RA redissolved more rapidly in steel B and consequently resulted in a stronger refining effect.
