An investigation of the effects of austenite strength and austenite stacking fault energy on the morphology of martensite in Fe-Ni-Cr-0.3C alloys

The relative effects of austenite stacking fault energy and austenite yield strength on martensite morphology have been investigated in a series of three Fe-Ni-Cr-C alloys. Carbon content (0.3 wt pct) andM ₆ temperature (− 15°) were held constant within the series. Austenite yield strength atM _s was measured by extrapolating elevated temperature tensile data. Austenite stacking fault energy was measured by the dislocation node technique. Martensite morphologies were characterized by transmission electron microscopy and electron diffraction techniques. A transition from plate to lath martensite occurred with decreasing austenite stacking fault energy. The austenite yield strength atM _s for the low SFE, lath-forming alloy was found to be higher than previously reported for lath-forming alloys. The relative effects of these variables on martensite morphologies in these alloys is discussed.     

The effect of austenite prestrain above the Md temperature on the martensitic transformation in Fe-Ni-Cr-C alloys

The effect of austenite prestrain above theM _d temperature on the structure and transformation kinetics of the martensitic transformation observed on cooling was determined for a series of Fe-Ni-Cr-C alloys. The alloys exhibited a shift in martensite morphology in the nondeformed state from twinned plate to lath while theM _s temperature, carbon content, and austenite grain size were constant. The transformation behavior was observed over the temperature range 0 to -196°C as a function of tensile prestrains performed above theM _d temperature. A range of prestrains from 5 pct to 45 pct was investigated. It is concluded that the response of a given alloy to austenite prestrain above theM _d temperature can be correlated with the morphology of the martensite observed in the nondeformed, as-quenched state. For the range of prestrains investigated, the transformation of austenite to lath martensite is much more susceptible to stabilization by austenite prestrain above theM _d temperature than is the transformation of austenite to plate martensite.     

The effect of quench rate on the properties and morphology of ferrous martensite

The effect of high quench rate on theM _s temperature, percent transformed, martensite morphology and austenite hardness has been studied for several Fe-Ni-C steels. For these steels the quench rate was varied only in the austenite region. TheM _s temperature was found to increase with increased quench rate for both high- and low carbon steels while the percent transformation increased or decreased depending upon the morphology of the steel. No variations in martensite hardness were found in the as-quenched condition, but a difference in tempering rate was found between fast and slow quenched specimens. Austenite hardness decreased slightly with increasing quench rate while the martensite morphology changed from lath to plate. Parallel aligned plate structures were observed which resemble a twinned lath morphology. It was demonstrated that the actual difference between this morphology and a true lath morphology is the self-accommodating nature of the lath structure. The morphology changes were compared to the measured changes in martensite properties in order to identify the mechanism of the morphology shift. It was concluded that for these alloys the morphology was controlled by the austenite shear mode.     

The effect of quench rate on the properties and morphology of ferrous martensite

The effect of high quench rate on theM _s temperature, percent transformed, martensite morphology and austenite hardness has been studied for several Fe-Ni-C steels. For these steels the quench rate was varied only in the austenite region. TheM _s temperature was found to increase with increased quench rate for both high- and low carbon steels while the percent transformation increased or decreased depending upon the morphology of the steel. No variations in martensite hardness were found in the as-quenched condition, but a difference in tempering rate was found between fast and slow quenched specimens. Austenite hardness decreased slightly with increasing quench rate while the martensite morphology changed from lath to plate. Parallel aligned plate structures were observed which resemble a twinned lath morphology. It was demonstrated that the actual difference between this morphology and a true lath morphology is the self-accommodating nature of the lath structure. The morphology changes were compared to the measured changes in martensite properties in order to identify the mechanism of the morphology shift. It was concluded that for these alloys the morphology was controlled by the austenite shear mode. S. J. Donachie was formerly a Graduate Assistant.     

The relationship between austenite strength and the transformation to martensite in Fe-10 pct Ni-0.6 pct C alloys

The effect of austenite yield strength on the transformation to martensite was investigated in Fe-10 pct Ni-0.6 pct C alloys. The strength of the austenite was varied by 1) additions of yttrium oxide particles to the base alloy and 2) changing the austenitizing temperature. The austenite strength was measured at three temperatures above theM _s temperature and the data extrapolated to the experimentally determinedM _s temperature. It is shown that the austenite yield strength is determined primarily by the austenite grain size and that the yttrium oxide additions influence the effect of austenitizing temperature on grain size. As the austenite yield strength increases, both theM _s temperature and the amount of transformation product at room temperature decrease. The effect of austenitizing temperature on the transformation is to determine the austenite grain size. The results are consistent with the proposal¹ that the energy required to overcome the resistance of the austenite to plastic deformation is a substantial portion of the non-chemical free energy or restraining force opposing the transformation to martensite.     

The influence of austenite strength upon the austenite-martensite transformation in alloy steels

The resistance of austenite to plastic deformation (austenite flow stress) was measured using a high temperature tensile apparatus. The flow stress was then correlated with the M_s temperature as determined magnetically during subsequent cooling. In one part of the study, the flow stress of the austenite was varied only by work hardening the austenite, allowing the austenite composition, which is known to affect M_s, to be held constant. A decrease in M_s temperature with increasing austenite flow stress was observed. This observation was supported by the observation of a decrease in the amount of austenite transformed at 25°C. In the other part of the study, a series of alloy steels of different chemical compositions was tested. A decrease in M_s temperature with increasing austenite flow stress was again observed. Strengthening of austenite by plastic deformation was shown not to change the chemical driving force for transformation. The effect of deformation on M_s temperature thus results from its influence on either the nucleation or the growth process. While the effect of austenite deformation on martensite nucleation is uncertain, specific nucleation models can account for only approximately one-third of the nonchemical free energy change which accompanies transformation. A proposal, consistent with the observations, was made that the energy expended for the deformation of austenite during martensite plate growth could reasonably account for a substantial part of the nonchemical free energy change.     

The effect of quench rate on the martensitic transformation in Fe−C alloys

The effect of very high quench rates on the transformation kinetics of a series of Fe?C and low-alloy steels and the morphology of an Fe?14Ni-0.76C alloy was investigated. TheM _S temperatures of the Fe?C and Fe?C?X alloys increased between 90° and 122°C in a sigmoidal fashion over a quench rate range from 2,750° to 24,800°C per sec. The sensitivity of theM _s temperature to the quench rate from the austenitizing temperature to 315°C was shown to be related to the influence of the third alloying element on the diffusivity of carbon in austenite. Transmission electron microscopy and optical metallography showed that the morphology of an Fe?14Ni?0.76C martensite is changed from a lath structure in slow quenched samples to a plate structure in fast quenched samples. The substructure of the untransformed austenite adjacent to the martensite plates changed from planar dislocation arrays to dislocation tangles with increased quench rate. These results were explained using a model for ferrous martensite strengthening based upon the extent of carbon segregation to imperfections in the austenite during cooling.     

Substructure of ausformed martensite in Fe-Ni and Fe-Ni-C alloys

The martensite substructure after ausforming has been studied for two different martensite morphologies: partially twinned, lenticular martensite (Fe-33 pct Ni, M_s =-105?C) and completely twinned “thin plate” martensite (Fe-31 pct Ni-0.23 pct C, M_s = -170?C), and in both cases ausforming produces a dislocation cell structure in the austenite which is inherited, without modification, by the martensite. In the Fe-Ni alloy, the dislocation cell structure is found in both the twinned (near the midrib) and untwinned (near the interface) regions, the latter also containing a regular dislocation network generated by the transformation itself and which is unaltered by the austenite dislocation cell structure. Similarly, in the Fe-Ni-C alloy, the transformation twins are unimpeded by the prior cell structure. These observations show that carbide precipitation during ausforming is not necessarily required to pin the austenite cell structure and that the martensite-austenite interface, backed by either twins or dislocations, does not exhibit a ”sweeping” effect. Although the martensite transformation twins are not inhibited by the ausforming cell structure, they do undergo a refinement with increased ausforming, and it is indicated that the transformation twin width in martensite depends on the austenite hardness. However, the relative twin widths remain unchanged, as expected from the crystallographic theory.     

Martensitic Phase Transformation from Non-isothermally Deformed Austenite in High Strength Steel 22SiMn2TiB

Non-isothermal compressive deformation was performed on high strength steel 22SiMn2TiB for the study of martensitic phase transformation from deformed austenite. The transformation start temperature M _s decreased with the increase of deformation from 0 to 50 pct, and the variation of deformation rate (0.1 and 10 s^?1) and the appearance of deformation-induced ferrite and bainite showed no influence on the change of M _s temperature. The deformation at both the rates increased the volume fraction of retained austenite; however, the carbon content of retained austenite decreased at 10 s^?1 and remained basically unchanged at 0.1 s^?1. The yield strength of austenite at M _s temperature and the stored energy in deformed austenite were experimentally obtained, with which the relationships between the change of M _s temperature and the thermodynamic driving force for martensitic phase transformation from deformed austenite were established by the use of the Fisher-ADP–Hsu model. And finally, the transformation kinetics was analyzed by the Magee–Koistinen–Marhurger equation.     

Thermodynamically Based Prediction of the Martensite Start Temperature for Commercial Steels

A thermodynamic method for predicting the martensite start temperature of commercial steels is developed. It is based mainly on information on M _s from binary Fe-X systems obtained from experiments with very rapid cooling, and M _s values for lath and plate martensite are treated separately. Comparison with the experimental M _s of several sets of commercial steels indicates that the predictive ability is comparable to models based on experimental information of M _s from commercial steels.     

Influence of austenite strength on martensite start temperature Ms

It is known that austenite strength determines the morphology of the new phase during martensitic transformation. As the strength of austenite influences the growth of a martensite crystal, i.e. the movement of the austenite/martensite interface, a correlation between strength of the parent phase and M_s has to exist. M_s depends on thermodynamical and mechanical properties of the alloys. To distinguish the individual variables, austenite strength was changed by different hardening mechanisms: solid solution hardening, plastic deformation or both.     

Experiments on directional martensitic transformation in steel

This paper describes some preliminary experiments on the feasibility of producing steels with an aligned martensitic microstructure. The parent austenite is initially textured, following which the martensite habit plane “activity” is regulated by transforming under strain. An Fe-20 pct Ni-5 pct Mn alloy exhibiting a lowM _s temperature and a lath mar-tensite morphology was found suitable for producing aligned martensite. On heavy defor-mation this alloy developed a typical (123) [412] fcc texture. Subsequent annealing then produced a strong (001) [100] cube texture. Nearly 80 pct of the martensite formed was aligned when such a cube-textured specimen was strained 2.5 pet at 45 deg to the rolling direction followed by cooling under strain in liquid nitrogen to produce martensite. M. Khobaib and R. Quattrone, Formerly with the Construc-tion Engineering Research Laboratory, U.S. Army Corps of Engineers, Champaign, Illinois     

The role of the constituent phases in determining the low temperature toughness of 5.5Ni cryogenic steel

Ferritic Fe-Ni steels that are intended for service at low temperature are usually given an intercritical temper as the final step in their heat treatment. The temper dramatically decreases the ductile-brittle transition temperature, T_B. Its metallurgical effect is to temper the lath martensite matrix and precipitate a distribution of fine austenite particles along the lath boundaries. Prior research suggests that the low value of T_B is a consequence of the small effective grain size of the ferrite-austenite composite. The present research was done to test this suggestion against the counter-hypothesis that the low T_B is due to the inherent toughness of the constituent phases. The approximate compositions of the tempered martensite and precipitated austenite phases in the composite microstructure of tempered 5.5Ni steel are known from STEM analysis. Bulk alloys were cast with these two compositions. Their mechanical properties were measured after heat treatment and compared to those of the parent alloy in the toughened ‘QLT’ condition. Both of the constituent phases are brittle at low temperature. It follows that the outstanding low-temperature toughness of the tempered alloy cannot be attributed to the inherent properties of the constituent phases, but must reflect their cooperative behavior in the composite microstructure. The austenitic bulk alloy was also used to investigate the stability of the precipitated austenite phase. The thermomechanical stability of the bulk alloy approximates that of the precipitated austenite within tempered 5.5Ni steel. This result is consistent with previous data, and supports the conclusion that the stability of the precipitated austenite is determined mainly by its chemical composition.     

Substructure of ausformed martensite in Fe-Ni and Fe-Ni-C alloys

The martensite substructure after ausforming has been studied for two different martensite morphologies: partially twinned, lenticular martensite (Fe-33 pct Ni, M_s =-105‡C) and completely twinned “thin plate” martensite (Fe-31 pct Ni-0.23 pct C, M_s = -170‡C), and in both cases ausforming produces a dislocation cell structure in the austenite which is inherited, without modification, by the martensite. In the Fe-Ni alloy, the dislocation cell structure is found in both the twinned (near the midrib) and untwinned (near the interface) regions, the latter also containing a regular dislocation network generated by the transformation itself and which is unaltered by the austenite dislocation cell structure. Similarly, in the Fe-Ni-C alloy, the transformation twins are unimpeded by the prior cell structure. These observations show that carbide precipitation during ausforming is not necessarily required to pin the austenite cell structure and that the martensite-austenite interface, backed by either twins or dislocations, does not exhibit a ”sweeping” effect. Although the martensite transformation twins are not inhibited by the ausforming cell structure, they do undergo a refinement with increased ausforming, and it is indicated that the transformation twin width in martensite depends on the austenite hardness. However, the relative twin widths remain unchanged, as expected from the crystallographic theory. T. MAKI, Formerly with the University of Illinois     

Austenite ferromagnetism and martensite morphology

The Curie temperature of the austenite, the martensite-start temperature, and martensite morphology have been determined in a series of nil-carbon Fe?Ni and Fe?Ni?Co alloys. For these alloys, austenite ferromagnetism aboveM _s is a necessary, but not sufficient, condition for the formation of lenticular rather than packet martensite. In contrast to Fe?Ni alloys where lenticular martensite only forms below ≈O°C, some of the Fe?Ni?Co alloys transform to this structure at temperatures up to ≈200°C. The results support the hypothesis that the resistance of austenite to plastic deformation affects the habit plane and thus morphology of the martensite which forms.     

Influence of austenite and martensite strength on martensite morphology

The martensite morphology and austenite flow strength have been determined in a variety of ferrous alloys chosen so that the austenites were paramagnetic, ferromagnetic, substitutional strengthened, and interstitial strengthened. It is demonstrated that two of the most important variables in determining the habit plane (and thus morphology) of martensite in a given alloy are the resistances to dislocation motion in austenite and in ferrite (i. e., martensite). In the wide variety of alloys where martensite with a {259}_γ habit plane was observed, the austenite flow strength atM _s is greater than 30,000 psi. At lower austenite strengths, either {225}_γ or {111}_γ habit planes are found depending on the resistance to dislocation motion in ferrite. Thus, {225} martensites are not always found as part of the spectrum between {111} and {259} martensites but only in the cases (e. g., interstitial strengthening) where ferrite is preferentially strengthened relative to austenite. All of the observations are consistent with the idea that the habit plane observed in a given alloy is the one involving the minimum plastic work for the lattice invariant shear.     

Stress-Assisted and strain-induced martensites in FE-NI-C alloys

A metallographic study was made of the martensite formed during plastic straining of metastable, austenitic Fe-Ni-C alloys withM _s temperatures below 0°C. A comparison was made between this martensite and that formed during the deformation of two TRIP steels. In the Fe-Ni-C alloys two distinctly different types of martensite formed concurrently with plastic deformation. The large differences in morphology, distribution, temperature dependence, and other characteristics indicate that the two martensites form by different transformation mechanisms. The first type, stress-assisted martensite, is simply the same plate martensite that forms spontaneously belowM _s except that it is somewhat finer and less regularly shaped than that formed by a temperature drop alone. This difference is due to the stress-assisted martensite forming from cold-worked austenite. The second type, strain-induced martensite, formed along the slip bands of the austenite as sheaves of fine parallel laths less than 0.5μm wide strung out on the {111}_γ planes of the austenite. Electron diffraction indicated a Kurdjumov-Sachs orientation for the strain-induced martensite relative to the parent austenite. No stress-assisted, plate martensite formed in the TRIP steels; all of the martensite caused by deformation of the TRIP steels appeared identical to the strain-induced martensite of the Fe-Ni-C alloys. It is concluded that the transformation-induced ductility of the TRIP steels is a consequence of the formation of strain-induced martensite. Formerly a graduate student at Stanford University     

Stress- and strain-induced formation of martensite and its effects on strength and ductility of metastable austenitic stainless steels

The effects of deformation-induced formation of martensite have been studied in metastable austenitic stainless steels. The stability of the austenite, being the critical factor in the formation of martensite, was controlled principally by varying the amounts of carbon and manganese. The formation of martensite was also affected by different test and rolling temperatures, rolling time, and various reductions in thickness. The terms “stress-induced” and “strain-induced” formation of martensite are defined. Experimental results show that low austenite stability resulted in stress-induced formation of martensite, high work-hardening rates, high tensile strengths, low “yield strengths,” and low elongation values. When the austenite was stable, plastic deformation was initiated by slip, and the work-hardening rate was too low to prevent early necking. A specific amount of strain-induced martensite led to an “optimum” work-hardening rate, resulting in high strengthand high ductility. For best results processing should be carried out aboveM _d and testing betweenM _d andM _s. Mechanical working aboveM _d had a negligible effect on the yield strength betweenM _d andM _s when the austenite stability was low, but its effect increased as the austenite became, more stable. Serrations appeared in the stress-strain curve when martensite was strain induced.     

Stabilization of retained austenite due to partial martensitic transformations

The stabilization effect of retained austenite has been studied using FeNiC alloys with M_s temperatures below 0°C via a two-step cooling procedure, i.e. the samples were first cooled to a temperature (T_a) below M_s temperature and then heated to room temperature (RT), after being held at RT for a while, the samples were recooled to low temperatures (23 or 82 K) and then heated to RT. It was found that, during the second step of cooling, the martensitic transformation occurred at a temperature of M_s′ which was lower than T_a. With increasing the amount of martensite formed during the first cooling, the difference in the martensitic transformation starting temperatures, ΔM_s = M_s − M_s′, increased. The mechanism of the stabilization of retained austenite during the second step of cooling is proposed to be mainly due to the inhibition effect produced by the previously formed martensite. The aging processes, which retard the growth of the previously formed martensite plates and reduce the number of the available nucleation sites, are the necessary conditions for the above mechanism to operate. By simplifying the internal resisting stress acting on the retained austenite due to the existence of martensite phase as a hydrostatic compressive stress, which increases with increasing the amount of martensite, the change in ΔM_s is discussed from a thermodynamic point of view.     

Reverse Martensitic Transformation and Resulting Microstructure in a Cold Rolled Metastable Austenitic Stainless Steel

The reverse martensitic transformation in cold‐rolled metastable austenitic stainless steel has been investigated via heat treatments performed for various temperatures and times. The microstructural evolution was evaluated by differential scanning calorimetry, X‐ray diffraction and microscopy. Upon heat treatment, both diffusionless and diffusion‐controlled mechanisms determine the final microstructure. The diffusion reversion from α′‐martensite to austenite was found to be activated at about 450°C and the shear reversion is activated at higher temperatures with A_f′ ~600°C. The resulting microstructure for isothermal heat treatment at 650°C was austenitic, which inherits the α′‐martensite lath morphology and is highly faulted. For isothermal heat treatments at temperatures above 700°C the faulted austenite was able to recrystallize and new austenite grains with a low defect density were formed. In addition, carbo‐nitride precipitation was observed for samples heat treated at these temperatures, which leads to an increasing M_s‐temperature and new α′‐martensite formation upon cooling.