scopic plane of maximum shear, i.e. 45° to the loading direc- tion. These shear bands measured between tens of microns to several hundred microns in length. Shear band nucleation and crack formation appears to be promoted by porosity in prefer- ence to any other microstructural feature. When nucleated at the specimen edge, shear band and subsequent crack formation followed an interdendritic path for about 50 µm before merging with the macroscopic 45° band (see Figure 12).
deformation mechanism in Fe-Mn-Al-C steels. Frommeyer and Brux5
mm2
than 3 wt.% suppress twinning. Frommeyer and Brux have shown that a Fe-28Mn-12Al-1C alloy with a γSFE
γSFE greater than 50 mJ/mm2 of 110 mJ/
adiabatic shear under high strain rates. The calculated γSFE for the 1 and 1.4% silicon containing alloys are 100 mJ/mm2
deformed by plastic flow in quasistatic testing and by .
and observations showing planar slip are in agreement with the observations by Frommeyer and Brux.5
The silicon difference did not have any significant impact on a change in the calculated γSFE
. The calculated γSFE
An interesting comparison can be made between the cast Fe- Mn-Al-C alloys studied here and the wrought alloy typically used to benchmark steel armor, i.e. rolled homogeneous armor. The specific compressive strength of the 1% silicon containing alloy is greater than rolled homogenous armor (RHA); both steels were tested at 3000 s-1
strain rate. Figure
10 shows excellent work hardening prior to fracture for the silicon containing Fe-Mn-Al-C steels with higher ultimate strength than RHA. Factoring in the 13% density reduc- tion for the Fe-Mn-Al-C alloy and the specific strengths are greater still. Table 3 compares specific high strain rate com- pression strengths of the solution treated and aged 1% sili- con modified alloy to that of RHA, which has been quenched and tempered to similar hardness values. The 1% alloy ex- ceeded the specific compressive tensile strength of RHA by 58 MPa/ρ (287 ksi/ρ) in the solution treated condition and 54 MPa/ρ (220 ksi/ρ) for the aged condition.
Observation of work hardening at high strain rates, in both the solution treated and aged Fe-Mn-Al-C-Si alloys, indi- cates an inherent resistance to the formation of adiabatic shear bands. However, the final failure was still by shear band formation, but at higher stresses than RHA. Casting porosity appears to be the main microstructural feature that nucleates the shear band.
Investigations are ongoing to increase Fe-Mn-Al-C alloy strength and notch toughness through improved foundry prac- tices and a physics based, first principles modeling of alloy additions. Foundry practice studies are focused on grain re- finement, porosity minimization, mold design and fluid mod- eling, reduction of phosphorous, and improving the cleanli- ness by reducing the number of nonmetallic inclusions.
International Journal of Metalcasting/Winter 10 energy
Thermodynamic calculations for the stacking fault energy (γSFE
) based on Olsen’s26 model have been used to predict the
stated that twinning is suppressed in systems with . Aluminum additions greater
Conclusions
Lightweight Fe-Mn-Al-C steels can be solution treated and aged 10 hours at 530°C (986°F) to meet hardness require- ments specified in MIL-PRF-32269. When tested at high strain rates, a combination of lower weight and higher com- pressive strength for the Fe-30Mn-9Al-1Si-0.5Mo alloy re- sults in a specific strength that is 28% greater than rolled homogeneous armor with equivalent hardness of 352 BHN. More importantly, the Fe-30Mn-9Al-1Si-0.5Mo alloy ex- hibits work hardening during high strain rate testing, which indicates an inherent resistance to the formation of adiabatic shear bands. However, casting porosity was observed to pro- mote adiabatic shear bands and crack formation.
Increasing silicon content from 1 to 1.4% in a Fe-30Mn-9Al- XSi-0.9C-0.5Mo alloy increased tensile strength and hard- ness, but reduced ductility and the work hardening exponent in aged materials. The decrease in tensile ductility appears to be related to an increase in non-metallic inclusions for the 1.4% Si alloy since the additional silicon did not appear to change the tensile fracture or microstructural deformation mechanisms. Improved properties are expected by reducing the phosphorous content. This will require a steelmaking process using charge materials and furnace refractories that are low in phosphorous. Lower phosphorous content should also lower the amount of ferrite in the solution treated and aged conditions.
REFERENCES
1. Ham J. L., Cairns R. E., “Manganese Joins Aluminum to Give Strong Stainless,” Product Engineering, Dec, pp 51-52 (1958).
2. Banerji S. K., “An Update on Fe-Mn-Al Steels,” pre- sented at Vanderbilt University’s workshop on “Con- servation and Substitution Technology for Critical Material,” June 15-17, (1981).
3. Brady G. S., Clauser H. R., “Manganese-Aluminum Steel,” Materials Handbook, 11th
Ed., p. 497 (1977).
4. Acselrad O., Pereira L. C., Amaral M. R., “Processing Condition, Microstructure and Strength of an Austen- itic FeMnAlC Alloy,” Proceedings of Proc. and Prop. of Mats., pp 829-834 (1992).
5. Frommeyer G., Brux U., “Microstructures and Me- chanical Properties of High-Strength Fe-Mn-Al-C Light-Weight TRIPLEX Steels,” Steel Research Int., vol. 77, pp 627-633 (2006).
6. Goretskii G. P., Gorev K. V., “Phase Equilibria in Fe- Mn-Al-C Alloys,” Russian Met., vol. 2, pp 217-221 (1990).
7. Han K. H., Choo W. K., “Phase Decomposition of Rapidly Solidified Fe-Mn-Al-C Austenitic Alloy,” Met. Trans. A., vol. 20A, pp 205-214 (1989).
8. Ishida K., Ohtani H., Naoya S., Kainuma R., Nishizawa,T., “Phase Equilibria in Fe-Mn-Al-C Al- loys,” ISIJ, vol. 30, pp 680-686 (1990).
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