Fracture Surface Examination of the Solution Treated Impact Specimens

The fracture surface of alloy K1525 (Fe-8Mn) showed completely

intergranular fracture under impact testing at room temperature after o being water quenched from the solution treatment temperature of 1000 C. The scanning electron micrograph showing this fracture surface is in Fig. 57. The ductile mode of fracture at temperatures above the ductile to brittle transition temperature was ductile dimpling.

The room temperature fracture mode of alloy K1526 (Fe-8Mn-2.5Mo) after the same solution treatment was a brittle cleavage/ductile dimpling mixture indicating tougher room temperature impact properties than the base alloy. The fracture mode below the ductile to brittle transition temperature is seen to be completely brittle cleavage and that above the ductile to brittle transition temperature is seen to be ductile dimpling. The appearance of these fracture surfaces is illustrated in the scanning electron micrographs in Fig. 68.

After the other three solution treatments of 1300°C for one hour then water quenching, 1300°C for one hour plus water quenching followed by a further hour at 1000°C followed by water quenching, and 1300°C for one hour plus water quenching followed by a further hour at 850°C plus water quenching, the brittle fracture mode in all cases was seen to be completely transgranular cleavage and the fracture surface of the alloy above the

ductile to brittle transition temperature was seen to be by ductile dimpling. These fractures may be seen in Figs. 69 - 71. The room temperature fracture mode of alloy K1527 (Fe-8Mn-5Mo) after one hour at 1000°C followed by water quenching was again seen to be of a mixed brittle cleavage/ductile

dimpling type indicating again a shift in the ductile to brittle transition temperature to a lower value by the addition of 5% Mo to the base Fe-8Mn alloy. This room temperature fracture was in the transition region

between fully brittle and fully ductile. Below the ductile to brittle transition temperature the mode of fracture was seen to be by brittle transgranular cleavage and above this temperature was by ductile dimpling

as shown in Fig. 73.

After the other three solution treatments, described earlier, all brittle fractures were by transgranular cleavage with some evidence of faceting on the cleavage planes. The tougher mode of fracture above the ductile to brittle transition temperature was, in all cases, by ductile dimpling. These fracture surfaces are shown in Figs. 74 - 76.

In summary, both 2.5% and 5% additions of molybdenum are seen to increase the intergranular cohesive strength above that of the matrix brittle fracture strength (Fig. 54) such that the brittle fracture mode changed from intergranular to transgranular cleavage.

4.2.5 Hardness of Alloys on Isothermal Aging at 450°C and 525°C

8

Nazim reported only a slight age hardening of alloy K1525 (Fe-8Mn) after isothermally holding at 450°C. The hardness first increased by 12Hv30 in ten minutes from 258Hv30 to 270Hv30 but the alloy then started to soften, the hardness dropping to 235Hv30 after 100 hours. His results of the variation of hardness with aging time, after the alloy was solution treated at 1000°C for one hour and then water quenched,are shown in Fig. 125 for aging temperatures 450°C, 350°C and 400°C.

The variation of hardness on isothermally holding alloy K1526 (Fe-8Mn- 2.5Mo) at 450°C is shown in Fig.83. The prior austenitizing treatment was one hour at 1000°C followed by water quenching. The error bars shown are the 90% confidence limits. There is possibly an initial small hardening response after three minutes but a straight line is assumed as it falls within the errors of the hardness measurement. The alloy is seen to commence softening after 330 hours and the hardness value falls from 348

Hv30 to ~ 300Hv30 after 9000 hours.

The variation of hardness with aging time at 450°C for alloy K1527

(Fe-8Mn-5Mo), after being solution treated for one hour at 1000°C followed

by water quenching, is shown in Fig. 83. An initial incubation period of

ten hours occurred with no variation in hardness from ~ 280Hv30. After

ten hours the hardness increased and reached 480Hv30 after 3000 hours,

an increase of 190Hv30. With a further holding time at 450°C, softening

of the alloy occurred, the hardness level being reduced to ^ 410Hv30

after a total isothermal age of 9000 hours.

The variation of hardness with aging time at 525°C for alloy K1526 (Fe-8Mn-2.5Mo), after being solution treated for one hour at 1000°C followed by water quenching, is shown in Fig. 128. The hardness is seen to fall much quicker than the equivalent alloy aged at 450°C. The

hardness level dropped immediately on aging unlike the same alloy at o

450 C when the hardness was seen to remain constant for 330 hours. At

525°C it fell from 340Hv30 to 318Hv30 after 100 minutes. A levelling of hardness is seen to occur after 17 hours with a value of 255Hv30 being maintained after 30 hours. After an equivalent holding time at 450°C the same alloy had a hardness of 330Hv30.

The variation of hardness on isothermally aging alloy K1527 (Fe- 8Mn- 5Mo) at 525°C, after a solution treatment of 1000°C for one hour followed by water quenching, is shown in Fig. 128. An initial drop in hardness was observed from 280Hv30 to 264Hv30 after 68 minutes before the hardness increased steadily to 310Hv30 after a holding time of 28 hours. No

incubation period was observed as with the 450°C age for the same alloy. The increase in hardness was not as dramatic as noted with the aging at 450°C.