Onset of recrystallization during the tensile deformation of austenitic iron at intermediate strain rates

The onset of recrystallization during the tensile deformation of austenitic iron has been fully documented for the temperature range 950 to 1350°C and strain-rate range 2.8 x 10^-5 to 2.3 x 10^-2 s^-1. Representative materials are zone-refined iron, electrolytic iron, Fe-0.05 C and Fe-5.2 Mn. In general, the strain at the onset of recrystallization decreases with increasing temperature of deformation and decreasing strain rate. The postponement of recrystallization is favored by prior annealing at temperatures above 1200°C and is greatest for the Fe-5.2 Mn alloy; however, for the range of strain rates used, it is difficult to completely eliminate recrystallization. The effects of test conditions on the onset of recrystallization are discussed in terms of a nucleation process that requires a critical amount of stored energy.     

Plastic deformation of austenitic iron at intermediate strain rates

The plastic deformation of austenitic iron, represented by a zone-refined iron, an electrolytic iron, an Fe−0.05 C alloy, and an Fe−5.2 Mn alloy, has been documented for the temperature range 950 to 1350°C (1740 to 2460°F) and the strain-rate range 2.8 × 10⁻⁵ to 2.3 × 10⁻² s⁻¹. The intrusion of recrystallization during deformation restricts the documentation to initial periods of strain usually less than 0.10. The general problem of retaining grain structures representative of polycrystals in specimens annealed at temperatures above 0.95T _m is recognized, and a basis for its solution is presented. Chemical composion appears to influence the plastic-flow behavior of austenitic iron primarily through its effect on the grain structure. Thus, the large-grained zone-refined iron is relatively weak, and the difference in behavior between the Fe−0.05 C alloy and the Fe−5.2 Mn alloy is small. Formerly Associate Scientist, U.S. Steel Corporation.     

Tensile failure of austenitic iron at intermediate strain rates

The basic failure behavior of austenitic iron has been established for the temperature range 950 to 1350°C and the strain-rate range 2.8 x 10^-5 to 2.3 x 1(10^-2 s^-1. Failure in zone-refined iron is determined solely by plastic deformation, leading first to multiple necking, continuing by the exclusive growth of a single neck, and concluding by separation at a point within that neck. With the increasing impurity content of electrolytic iron, Fe-0.05 C and Fe-5.2 Mn, this failure process is interrupted at the lower temperatures by fracture at either second-phase particles or grain boundaries. The regimes of these two fracture modes have been determined as functions of strain rate, deformation temperature, and annealing temperature. Recrystallization is prevalent during the plastic deformation of austenitic iron and influences the necking process to some extent. Recrystallization is more influential as a means of stabilizing arrays of intergranular cracks, thereby allowing the cracks to undergo appreciable plastic deformation during the final stage of failure. The concept of failure diagrams is introduced as a simple means of representing the complex interposition of plastic instability, recrystallization, and fracture during the failure process.     

Tensile failure of austenitic iron at intermediate strain rates

The basic failure behavior of austenitic iron has been established for the temperature range 950 to 1350°C and the strain-rate range 2.8 x 10^-5 to 2.3 x 1(10^-2 s^-1. Failure in zone-refined iron is determined solely by plastic deformation, leading first to multiple necking, continuing by the exclusive growth of a single neck, and concluding by separation at a point within that neck. With the increasing impurity content of electrolytic iron, Fe-0.05 C and Fe-5.2 Mn, this failure process is interrupted at the lower temperatures by fracture at either second-phase particles or grain boundaries. The regimes of these two fracture modes have been determined as functions of strain rate, deformation temperature, and annealing temperature. Recrystallization is prevalent during the plastic deformation of austenitic iron and influences the necking process to some extent. Recrystallization is more influential as a means of stabilizing arrays of intergranular cracks, thereby allowing the cracks to undergo appreciable plastic deformation during the final stage of failure. The concept of failure diagrams is introduced as a simple means of representing the complex interposition of plastic instability, recrystallization, and fracture during the failure process.     

Onset of recrystallization during the tensile deformation of austenitic iron at intermediate strain rates

The onset of recrystallization during the tensile deformation of austenitic iron has been fully documented for the temperature range 950 to 1350°C and strain-rate range 2.8 × 10^-5 to 2.3 × 10^-2 s^-1. Representative materials are zone-refined iron, electrolytic iron, Fe−0.05 C and Fe−5.2 Mn. In general, the strain at the onset of recrystallization decreases with increasing temperature of deformation and decreasing strain rate. The postponement of recrystallization is favored by prior annealing at temperatures above 1200°C and is greatest for the Fe−5.2 Mn alloy; however, for the range of strain rates used, it is difficult to completely eliminate recrystallization. The effects of test conditions on the onset of recrystallization are discussed in terms of a nucleation process that requires a critical amount of stored energy.     

Plastic deformation of delta-ferritic iron at intermediate strain rates

The plastic deformation of delta-ferritic iron, represented by an electrolytic iron and Fe-0.028 C, Fe-0.044 C and Fe-3.0 Si alloys, has been measured for the temperature range 1200 to 1525‡C and the strain rate range 2.8 x 10^-5 to 2.3 x 10^-2 s^-1. For the bamboo-like tension specimens the plastic flow behavior is approximately nonlinear viscous. Delta-ferrite is more than four times weaker than austenite, and does not exhibit dynamic recrystallization. At the melting point of iron the extrapolated steady-state flow stress increases from 0.31 to 1.86 MN/m² (45 to 370 psi) over the range of strain rate examined.      

Plastic deformation of delta-ferritic iron at intermediate strain rates

The plastic deformation of delta-ferritic iron, represented by an electrolytic iron and Fe-0.028 C, Fe-0.044 C and Fe-3.0 Si alloys, has been measured for the temperature range 1200 to 1525‡C and the strain rate range 2.8 x 10^-5 to 2.3 x 10^-2 s^-1. For the bamboo-like tension specimens the plastic flow behavior is approximately nonlinear viscous. Delta-ferrite is more than four times weaker than austenite, and does not exhibit dynamic recrystallization. At the melting point of iron the extrapolated steady-state flow stress increases from 0.31 to 1.86 MN/m² (45 to 370 psi) over the range of strain rate examined.     

Effect of cooling rate and alloying on the

The pearlitic hardenability of a high-purity Fe-0.8 pct C alloy and zone-refined iron binary alloys containing Mn, Ni, Si, Mo, or Co was studied by means of hot-stage microscopy. The binary alloys were carburized in a gradient furnace to produce eutectoid compositions, thus eliminating proeutectoid phases. A special technique based on hot-stage microscopy was used to study the effect of cooling rate (10°F/min to 25,000°F/min) on the transformation of austenite and provided data for the construction of continuous cooling-transformation diagrams. From these diagrams critical cooling rates were obtained for hardenability calculations. It was found that molybdenum is the most effective element, followed by Si, Ni, Co, and Mn, in suppressing the pearlite transformation,i.e., in increasing the hardenability of the alloys studied. The alloying additions were grouped into two classes according to their effect on hardenability: α-stabilizers (Mo and Si) and γ-stabilizers (Ni, Co, Mn), with the α-stabilizers being the more effective in improving hardenability. This paper is based on a presentation made at a symposium on “Hardenability” held at the Cleveland Meeting of The Metallurgical Society of AIME, October 17, 1972, under the sponsorship of the IMD Heat Treatment Committee.     

Environmentally assisted cracking of two-phase Fe-Mn-Al alloys in NaCl solution

Three two-phase Fe-Mn-Al alloys with nominal compositions, Fe-24Mn-9Al, Fe-27Mn-9Al-3Cr,. and Fe-27Mn-9Al-6Cr, were prepared in the solution-treated and cold-rolled conditions. The fractions of ferrite in the solution-treated condition were controlled at 46 to 60 pct, mainly by adjusting the carbon content and the relative amounts of Mn and Al. The ferrite fractions were reduced to 30 to 37 pct after 75 pct deformation by cold-rolling. Specimens were tensile tested at open circuit in aerated 3.5 pct NaCl solution at slow strain rates ranging from 4 × 10^-7 to 4 × 10^-5 s^-1 at room temperature. All of the alloys were quite susceptible to environmentally assisted cracking (EAC). The deformed specimens showed less susceptibility, presumably because the plasticity was already too limited. The EAC appeared to occur at or after the onset of plastic deformation. In this alloy system, the ferritic phase was less resistant to EAC than the austenitic phase, in contrast to the Fe-Cr-Ni stainless steels. The crack propagated preferentially through the ferrite grains or along the ferrite/austenite grain boundaries. The addition of up to 6 pct Cr did not improve the EAC resistance. Formerly Graduate Student, Department of Materials Science and Engineering, National Tsing Hua University     

Flow stress and microstructural evolution during hot working of alloy 22cr-13ni-5mn-0.3n austenitic stainless steel

The stress-strain behavior and the development of microstructure between 850 °C and 1150 °C in an austenitic stainless steel, 22Cr-13Ni-5Mn-0.3N, were investigated by uniaxial compression of cylindrical specimens at strain rates between 0.01 and 1 s^-1 up to a strain of one. The measured (anisothermal) and corrected (isothermal) flow curves were distinctly different. The flow stress at moderate hot working temperatures, compared to a number of other austenitic alloys, was second only to that of alloy 718. Both static and dynamic recrystallization were observed. Recrystallization was sluggish in comparison to alloy 304L, apparently due to the presence of a fine Cr- and Nb-rich second-phase dispersion, identified as Z phase, which tended to pin the high-angle grain boundaries even at a high temperature of 1113 °C. Recrystallization may also be retarded by preferential res-toration through the competitive process of recovery, which is consistent with the relatively high stacking-fault energy for this alloy. It is concluded that this alloy must be hot worked at temperatures higher than usual for austenitic stainless steels in order to minimize flow stress and refine grain size.     

Tracer diffusion of63Ni in Fe-17 wt pct Cr-12 wt pct Ni

The tracer diffusion of⁶³Ni in Fe-17 Cr-12 Ni by both volume and grain boundary transport has been studied from 600° to 1250°C. The use of an RF sputtering technique for serial sectioning allowed the determination of very small volume diffusion coefficients at the lower temperatures. Volume diffusion of nickel in this alloy was observed to be much slower than in pure iron or austenitic stainless steel at comparable temperatures. The volume diffusion coefficient is described byD _v =8.8 exp (−60,000/RT) cm²/s and grain boundary diffusion is described by σD _gb =3.7×10⁻⁹ exp (−32,000/RT) cm³/s. R. A. PERKINS, formerly Presidential Intern, Metals and Ceramics Division, Oak Ridge National Laboratory, Oak, Ridge, Tenn. 37830, is     

Microstructure Evolution and Age-Hardening Behavior of Microalloyed Austenitic Fe-30Mn-9Al-0.9C Light-Weight Steels

The aging behavior and mechanical properties of microalloyed austenitic Fe-30Mn-9Al-0.9C light-weight steels were investigated through transmission electron microscopy analysis, electron backscatter diffraction analysis, tensile tests, and Vickers hardness tests. The base steels were aged at 823 K (550 °C) for up to 10,000 minutes. The true stress–strain responses of solution-treated samples before aging showed that the addition of Nb and/or V improved the strength by grain refinement and precipitation hardening. During the process of tensile deformation, the strain-hardening rate of Fe-30Mn-9Al-0.9C steel steadily increased due to the microband-induced plasticity (MBIP) from the onset of plastic deformation to ε = 25 pct, while such behavior was weakened and not observed in Nb- and/or V-added steels despite MBIP. In the early stage of aging, the Vickers hardness gradually increased with an increase in the aging time due to the precipitation of κ-carbide of (Fe,Mn)₃AlC and remained stagnant between the aging times of 1000 and 3000 minutes. The hardness increased again after 3000 minutes due to the formation of ferrite and brittle β-Mn.     

Superplastic behavior of two ultrahigh boron steels

The high-temperature deformation behavior of two ultrahigh boron steels containing 2.2 pct and 4.9 pct B was investigated. Both alloys were processedvia powder metallurgy involving gas atomization and hot isostatic pressing (hipping) at various temperatures. After hipping at 700 °C, the Fe-2.2 pct B alloy showed a fine microstructure consisting of l-μm grains and small elongated borides (less than 1μm) . At 1100 °C, a coarser microstructure with rounded borides was formed. This alloy was superplastic at 850 °C with stress exponents of about two and tensile elongations as high as 435 pct. The microstructure of the Fe-4.9 pct B alloy was similar to that of the Fe-2.2 pct B alloy showing, in addition, coarse borides. This alloy also showed low stress exponent values but lacked high tensile elongation (less than 65 pct), which was attributed to the presence of stress accumulation at the interface between the matrix and the large borides. A change in the activation energy value at theα-γ transformation temperature was seen in the Fe-2.2 pct B alloy. The plastic flow data were in agreement with grain boundary sliding and slip creep models. J.A. JIMéNEZ, Postdoctoral Fellow, formerly with Centro Nacional de Investigaciones Metalurgicas, C.S.I.C.     

Effect of thermomechanical treatments on the room-temperature mechanical behavior of iron aluminide Fe3AI

The room-temperature hydrogen embrittlement (HE) problem in iron aluminides has restricted their use as high-temperature structural materials. The role of thermomechanical treatments (TMT),i.e., rolling at 500 °C, 800 °C, and 1000 °C, and post-TMT heat treatments,i.e., recrystallization at 750 °C and ordering at 500 °C, in affecting the room-temperature mechanical properties of Fe-25A1 intermetallic alloy has been studied from a processing-structure-properties correlation viewpoint. It was found that when this alloy is rolled at higher temperature, it exhibits a higher fracture strength. This has been attributed to fine subgrain size (28/μ) due to dynamic recrystallization occurring at the higher rolling temperature of 1000 °C. However, when this alloy is rolled at 1000 °C and then recrystallized, it shows the highest ductility but poor fracture strength. This behavior has been ascribed to the partially recrystallized microstructure, which prevents hydrogen ingress through grain boundaries and minimizes hydrogen embrittlement. When the alloy is rolled at 1000 °C and then ordered at 500 °C for 100 hours, it shows the highest fracture strength, due to its finer grain size. The alloy rolled at 500 °C and then ordered undergoes grain growth. Hence, it exhibits a lower fracture strength of 360 MPa. Fracture morphologies of the alloy were found to be typical of brittle fracture,i.e., cleavage-type fracture in all the cases.     

Microstructural hot corrosion behaviour of austenitic Fe-29.7Mn-8.7Al-1.04C alloy in fused Na2SO4 and Na2SO4-NaCl mixture

A short-time hot corrosion test was performed on the austenitic Fe-29.7Mn-8.7Al-1.04C alloy in sodium sulphate at 900°C. The corrosion scales formed on the alloy investigated were studied by scanning electron microscopy and X-ray mapping techniques. The hot-corrosion scales morphology of the austenitic Fe-29.7Mn-8.7Al-1.04C alloy were characterized by the formation of Al₂S₃ and α-MnS sulphides at and beneath the internal alumina scale. No fluxing of the scales was observed. The effect of the addition of sodium chloride to the fused sodium sulphate on the hot corrosion scales morphology was also investigated.     

Effect of cerium content on the deformation behavior of rapidly solidified Al-Fe-Ce alloys

The deformation behavior of a rapidly solidified Al-8.9Fe-6.9Ce (wt pct) alloy was studied in the temperature range of 250 °C to 350 °C and stress range of 20 to 175 MPa. The stress exponents and activation energies suggest that the alloy exhibits a pronounced diffusional creep regime with a transition to power law creep behavior at stresses beyond 60 MPa. Comparing these data with those obtained earlier for an Al-8.8Fe-3.7Ce alloy, it was found that in the diffusional creep regime, the Ce content had no effect on the creep rate. However, in the power law creep regime, a strong dependence on the precipitate spacing, as predicted by the structureinvariant creep law,^[5] was observed. The higher volume fraction of precipitates in the Al-8.9Fe6.9Ce alloy causes a decrease in the power law creep rates by a factor of 5. Formerly Graduate Student. Formerly Assistant Professor, University of Illinois at Urbana-Champaign.     

The improvement of cryogenic mechanical properties of Fe-12 Mn and Fe-8 Mn alloy steels through thermal/mechanical treatments

An investigation has been made to improve the low temperature mechanical properties of Fe-8Mn and Fe-12Mn-0.2 Ti alloy steels. A reversion annealing heat treatment in the two-phase (α+ γ) region following cold working has been identified as an effective treatment. In an Fe-12Mn-0.2Ti alloy a promising combination of low temperature (-196°C) fracture toughness and yield strength was obtained by this method. The improvement of properties was attributed to the refinement of grain size and to the introduction of a uniform distribution of retained austenite (γ). It was also shown that an Fe-8Mn steel could be grain-refined by a purely thermal treatment because of its dislocated α martensitic structure and absence of ε martensite. As a result, a significant reduction of ductile to brittle transition temperature was obtained. formerly with the Lawrence Berkeley Laboratory, University of California.     

Mechanical Behavior and Microstructural Change of a High Nitrogen CrMn Austenitic Stainless Steel during Hot Deformation

The hot deformation behavior of a high nitrogen CrMn austenitic stainless steel in the temperature range 1173 to 1473 K (900 to 1200 °C) and strain rate range 0.01 to 10 s⁻¹ was investigated using optical microscopy, stress-strain curve analysis, processing maps, etc. The results showed that the work hardening rate and flow stress decreased with increasing deformation temperature and decreasing strain rate in 18Mn18Cr0.5N steel. The dynamic recrystallization (DRX) grain size decreased with increasing Z value; however, deformation heating has an effect on the DRX grain size under high strain rate conditions. In the processing maps, flow instability was observed at higher strain rate regions (1 to 10 s⁻¹) and manifested as flow localization near the grain boundary. Early in the deformation, the flow instability region was at higher temperatures, and then the extent of this unstable region decreased with increasing strain and was restricted to lower temperatures. The hot deformation equation as well as the quantitative dependence of the critical stress for DRX and DRX grain size on Z value was obtained.     

New aspects on the superplasticity of fine-grained 7475 aluminum alloys

Thermomechanical processes were developed which give fine grain sizes of 6 and 8 μm in the 7475 Al alloy. Superplastic properties of this material were evaluated in the temperature range of 400 °C to 545 °C over the strain-rate range of 2.8 x 10^-4 to 2.8 X 10^-2 s^-1. The maximum ductility exhibited by the alloy was approximately 2000 pct, and optimum superplasticity was achieved at a strain rate of 2.8 X 10^-3 s^-1 which is higher by an order of magnitude than other 7475 Al alloys. This result is attributed to the presence of fine dispersoids which maintain the fine grain size at high homologous temperatures. The flow stress and strain-rate sensitivity strongly depend on the grain size. The superplastic 7475 Al alloy has strain-rate sensitivities of 0.67 (6 μm) and 0.5 (13 μm) and an activation energy which is similar to the one for grain boundary diffusion of aluminum. Microstructural investigation after superplastic tests revealed zones free of dispersoid particles at grain boundaries primarily normal to the tensile direction. These dispersoidfree zones (DFZs) appear even after 100 pct elongation and are occasionally as large as 5 μm across. This result demonstrates the importance of diffusional flow in superplastic deformation of the fine-grained 7475 Al alloy especially at low elongations.