Retained austenite and tempered martensite embrittlement in medium carbon steels

Electron microscopy, diffraction and microanalysis, X-ray diffraction, and auger spectroscopy have been used to study quenched and quenched and tempered 0.3 pct carbon low alloy steels. Somein situ fracture studies were also carried out in a high voltage electron microscope. Tempered martensite embrittlement (TME) is shown to arise primarily as a microstructural constraint associated with decomposition of interlath retained austenite into M₃C filM_s upon tempering in the range of 250 °C to 400 °C. In addition, intralath Widmanstätten Fe₃C forms from epsilon carbide. The fracture is transgranular with respect to prior austenite. The sit11Ation is analogous to that in upper bainite. This TME failure is different from temper embrittlement (TE) which o°Curs at higher tempering temperatures (approximately 500 °C), and is not a microstructural effect but rather due to impurity segregation (principally sulfur in the present work) to prior austenite grain boundaries leading to intergranular fracture along those boundaries. Both failures can o°Cur in the same steels, depending on the tempering conditions.     

Mechanisms of tempered martensite embrittlement in low alloy steels

An investigation into the mechanisms of tempered martensite embrittlement (TME), also know as “500°F” or “350°C” or one-step temper embrittlement, has been made in commercial, ultra-high strength 4340 and Si-modified 4340 (300-M) alloy steels, with particular focus given to the role of interlath films of retained austenite. Studies were performed on the variation of i) strength and toughness, and ii) the morphology, volume fraction and thermal and mechanical stability of retained austenite, as a function of tempering temperature, following oil-quenching, isothermal holding, and continuous air cooling from the austenitizing temperature. TME was observed as a decrease in bothK _Ic and Charpy V-notch impact energy after tempering around 300°C in 4340 and 425°C in 300-M, where the mechanisms of fracture were either interlath cleavage or largely transgranular cleavage. The embrittlement was found to be concurrent with the interlath precipitation of cementite during temperingand the consequent mechanical instability of interlath films of retained austenite during subsequent loading. The role of silicon in 300-M was seen to retard these processes and hence retard TME to higher tempering temperatures than for 4340. The magnitude of the embrittlement was found to be significantly greater in microstructures containing increasing volume fractions of retained austenite. Specifically, in 300-M the decrease inK _Ic, due to TME, was a 5 MPa√m in oil quenched structures with less than 4 pct austenite, compared to a massive decrease of 70 MPa√m in slowly (air) cooled structures containing 25 pct austenite. A complete mechanism of tempered martensite embrittlement is proposed involving i) precipitation of interlath cementite due to partial thermal decomposition of interlath films of retained austenite, and ii) subsequent deformation-induced transformation on loading of remaining interlath austenite, destabilized by carbon depletion from carbide precipitation. The deterioration in toughness, associated with TME, is therefore ascribed to the embrittling effect of i) interlath cementite precipitates and ii) an interlath layer of mechanically-transformed austenite,i.e., untempered martensite. The presence of residual impurity elements in prior austenite grain boundaries, having segregated there during austenitization, may accentuate this process by providing an alternative weak path for fracture. The relative importance of these effects is discussed. Formerly with the Lawrence Berkeley Laboratory, University of California.     

Correlation of microstructure and tempered martensite embrittlement in two 4340 steels

This study is concerned with a correlation between the microstructure and fracture behavior of two AISI 4340 steels which were vacuum induction melted and then deoxidized with aluminum and titanium additions. This allowed a comparison between microstructures that underwent large increases in grain size and those that did not. When the steels were tempered at 350°C,K _Ic and Charpy impact energy plots showed troughs which indicated tempered martensite embrittlement (TME). The TME results of plane strain fracture toughness are interpreted using a simple ductile fracture initiation model based on large strain deformation fields ahead of cracks, suggesting thatK _Icscales roughly with the square root of the spacing of cementite particles precipitated during the tempering treatment. The trough in Charpy impact energy is found to coincide well with the amount of intergranular fracture and the effect of segregation of phosphorus on the austenite grain boundaries. In addition, cementite particles are of primary importance in initiating the intergranular cracks and, consequently, reducing the Charpy energy. These findings suggest that TME in the two 4340 steels studied can be explained quantitatively using different fracture models.     

The micro-mechanisms of tempered martensite embrittlement in 4340-type steels

A study of the micro-mechanisms of tempered martensite embrittlement was made on a series of 4340-type steels in which the contents of manganese, silicon, and trace impurities, especially phosphorus and sulfur, were varied. One plain-carbon steel was also examined. The study employed Charpy impact tests and four-point slow-bend tests coupled with an elastic-plastic stress analysis, as well as scanning electron fractography, Auger electron spectroscopy, transmission electron microscopy of extraction replicas, and magnetic measurements of the transformation of retained austenite. The results indicate that in these steels the TME phenomenon is an intergranular embrittlement problem caused by carbide precipitation on prior austenite grain boundaries which are already weakened by segregated phosphorus and sulfur. The transformation of intragranular retained austenite is concluded not to be of primary significance in the TME in these steels, although it may contribute to the magnitude of the TME toughness trough.     

Effect of molybdenum on tempered martensite embrittlement of medium carbon Si-Mn steels

This paper reports a study of the effects of molybdenum on tempered martensite embrittlement (TME) of medium carbon Si-Mn steels. The study employed standard U-notch impact tests, scanning electron fractography, transmission electron microscopy, Auger electron spectroscopy, and dilatometric measurements. It is shown that the addition of molybdenum to Si-Mn steels does not eliminate TME, but molybdenum acts to decrease the ductile-brittle transition temperature (DBTT), thus making the impact test temperature for revealing TME lowered. Furthermore, the deferring role of molybdenum on TME is observed. In the molybdenum doped steels the TME embrittlement trough is displaced to a higher tempering temperature. The embrittlement is found to be concurrent with the replacement of ?-carbide by cementite during tempering. In the molybdenum doped steels the displacement of TME to a higher tempering temperature is seen to be associated with the role of molybdenum retarding cementite precipitation.     

The evaluation of tempered martensite embrittlement in 4130 steel by instrumented charpy V-notch testing

Tempered martensite embrittlement (TME) was studied in vacuum-melted 4130 steel with either 0.002 or 0.02 wt pct P. TME was observed as a severe decrease in Charpy V-notch impact energy, from 46 ft-lb. at 200 °C to 35 ft-lb. at 300 °C in the low P alloy. The impact energy of the high P alloy was consistently lower than that of the low P alloy in all tempered conditions. Fracture was transgranular for all specimens; therefore, segregation of P to the prior austenitic grain boundaries was not a factor in the o°Currence of TME. Analysis of load-time curves obtained by instrumented Charpy testing revealed that the embrittlement is associated with a drop in the pre-maximum-load and post-unstable-fracture energies. In specimens tempered at 400 °C the deleterious effect of phosphorus on impact energy became pronounced, a result more consistent with classical temper embrittlement rather than TME. A constant decrease in pre-maximum-load energy due to phosphorus content was observed. The pre-maximum-load energy decreases with increasing tempering temperature in the range of 200 °C to 400 °C, a result explained by the change in work hardening rate. Carbon extraction replicas of polished and etched as-quenched specimens revealed the presence of Fe₂MoC and/or Fe₃C carbides retained after austenitizing. Ductile crack extension close to the notch root was related to the formation of fine micro voids at the retained carbides. This paper is based on a presentation made at the “pcter G. Winchell Symposium on Tempering of Steel” held at the Louisville Meeting of The Metallurgical Society of AIME, October 12-13, 1981, under the sponsorship of the TMS-AIME Ferrous Metallurgy and Heat Treatment Committees.     

The relation between the fracture behavior of 4340 bend specimens and the observation of tempered martensite embrittlement

V-notched and fatigue precracked Charpy specimens of various sizes, tested in impact and slow bend, were used to study tempered martensite embrittlement in a 4340 steel. When plotted as a function of tempering temperature, the results showed that the magnitude of the toughness decrease caused by embrittlement varied with the type and size of the specimens. Embrittlement was always detected using thin samples but its detection in thick specimens depended on whether or not they contained a precrack. In particular, no embrittlement-associated fall in toughness was observed using standard size precracked samples tested in slow bend. Separation of the shear and flat fracture components of the absorbed energies showed that the variation of shear energy is a major factor contributing to embrittlement. The results are interpreted as indicating that intergranular fracture occurs more as the result of inhibition of plastic flow within the grains rather than directly as the result of the appearance of a low resistance crack path at the grain boundaries.     

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.     

Chi-carbide in tempered high carbon martensite

Martensite in an Fe-1.22C alloy was tempered at 523, 573, and 623 K and examined by transmission electron microscopy (TEM) and Mössbauer effect spectroscopy (MES) to identify the morphology and type of carbide formed at the beginning of the third stage of tempering. Carbides formed in three morphologies: on twins within the martensite plates, in the matrix of twin-free areas of the martensite plates, and along the interfaces of the martensite plates. Chi-carbide(x), as identified by selected area diffraction (SAD), was associated with each carbide morphology in specimens tempered at 573 K. Cementite (0) together with chi-carbide was observed in specimens tempered at 623 K. Small amounts (about 2 pct) of retained austenite were observed by MES of specimens tempered at 523 K. The transformation of the 25 pct retained austenite in as-quenched specimens was related to theX-carbide formed at the martensite plate interfaces during tempering. The MES results also show the presence of κ-carbide in the specimen tempered at 523 K and yields parameters indicative of a mixture of κ and θ carbides for the specimens tempered at 573 K and 623 K. MES measurements of the magnetic transition temperatures of the carbides show diffuse transitions but suggest that κ is the dominant carbide in the tempering temperature range examined.     

Tempered martensite embrittlement in Fe-Mo-C and Fe-W-C steel

A study on the phenomenon of tempered martensite embrittlement (TME) has been made in experimental Fe-Mo-C and Fe-W-C steel. Charpy impact testing was conducted to evaluate the impact toughness, sensitive to TME. Retained austenite was observed by an analytical transmission electron microscopy in both steels. Both steels represented TME. TME was correlated with the formation of the interlath cementite, resulting from the decomposition of interlath retained austenite. TME occurred in a limited range of test temperatures where the interlath cementite could act as a source of embrittling cracks. Therefore, both the interlath cementite resulting from the decomposition of the interlath retained austenite, and the level of matrix toughness, enabling the interlath cementite to act as an effective embrittler, are necessary to produce TME.     

Tempered martensite embrittlement in SAE 4340 steel

The toughness of SAE 4340 steel with low (0.003 wt pct) and high (0.03 wt pct) phosphorus has been evaluated by Charpy V notch (CVN) impact and compact tension plane strain fracture toughness (K _1c) tests of specimens quenched and tempered up to 673 K (400°C). Both the high and low P steel showed the characteristic tempered martensite embrittlement (TME) plateau or trough in room temperature CVN impact toughness after tempering at temperatures between 473 K (200°C) and 673 K (400°C). The CVN energy absorbed by low P specimens after tempering at any temperature was always about 10 J higher than that of the high P specimens given the same heat treatment. Interlath carbide initiated cleavage across the martensite laths was identified as the mechanism of TME in the low P 4340 steel, while intergranular fracture, apparently due to a combination of P segregation and carbide formation at prior austenite grain boundaries, was associated with TME in the high P steel.K _IC values reflected TME in the high P steels but did not show TME in the low P steel, a result explained by the formation of a narrow zone of ductile fracture adjacent to the fatigue precrack during fracture toughness testing. The ductile fracture zone was attributed to the low rate of work hardening characteristic of martensitic steels tempered above 473 K (200°C).     

Tempered martensite embrittlement in phosphorus doped steels

In this paper the effect of phosphorus on tempered martensite embrittlement of Ni−Cr steels is reported. It is shown that the measured degree of embrittlement depends on the phosphorus concentration, test temperature, grain size, and austenitizing temperature. Although reducing the prior austenite grain size tends to reduce the observed embrittlement, this can be offset by the fact that the low austenitizing temperatures used to produce the fine grain size cause an increased amount of impurity segregation. It is further shown that bulk phosphorus concentrations below 100 wppm may be required to avoid embrittlement of this type in ultra-high strength steels.     

Transformation of austenite into bainitic ferrite and martensite

The stress induced martensitic transformation in the upper metastable intermediate state of γ-α transformation in ferrous materials, structured as ferritic bainite, is discussed. The fibrous structured ferritic bainite consists of retained austenite and ferrite platelets growing in the [111]α//[101]γ direction. The ferrite growth Induces carbon enrichment of the adjacent austenite at the phase boundaries. Strengthening at high stress levels up to the yield point causes dislocation tangles in the ferrite fibre and the formation of shear bands crossing each other in the retained austenite. At lower carbon contents of the austenite, lath martensite precipitates at the shear band intersections and at high shear band densities martensite blocks are observed. In carbon enriched austenite martensite lenses formed by shear processes have been observed. At alternating loading conditions, exceeding the stress level for athermic martensite formation, various shear planes are activated forming characteristic patterns of plate martensite.     

Hardness of tempered martensite in carbon and low-alloy steels

This paper presents the results of a systematic study of the effect of carbon, manganese, phosphorus, silicon, nickel, chromium, molybdenum, and vanadium on the hardness of martensite in low to medium carbon steels tempered for one hour at 100°F (56°C) intervals in the range 400 to 1300°F (204 to 704°C). Results show that the as-quenched hardness depends solely on carbon content. On tempering, the effect of carbon on hardness decreases markedly with increasing tempering temperature. Studies of carbon-0.5 manganese steels showed that the incremental increase in hardness from 0.5 pct manganese after a given tempering treatment was independent of carbon content. Based on this result, studies of the effects of the other alloying elements were made using a 0.2 or 0.3 pct carbon, 0.3 to 0.5 pct manganese steel base composition. The hardness of the resulting tempered martensite was assumed to be due to a given alloy addition, and when two or more alloying elements were added, their effects were assumed to be additive. Each of the seven alloying elements increased the hardness of tempered martensite by varying amounts, the increase being greater as more of each element was present. Nickel and phosphorus have substantially the same effect at all tempering temperatures. Manganese has essentially the same hardening effect at any temperature in the range 700 (371°C) to 1300°F; silicon is most effective at 600°F (316°C), chromium at 800°F (427°C), molybdenum at 1000 to 1100°F (538 to 592°C), and vanadium at 1200°F (649°C). Using the data obtained, a procedure is established for calculating the hardness of tempered martensite for carbon and alloy steel compositions in the range studied and for any combination of tempering time and temperature. R. A. GRANGE was formerly with U. S. Steel Corporation (retired)     

Mössbauer effect studies of tempered martensite

The⁵⁷Fe Mössbauer effect has been widely used to trace the tempering of Fe-C and Fe-N martensitic systems, and has provided much detailed information on the interstitial atom configurations formed during these processes. Carbon and nitrogen atoms are found predominantly to o°Cupy octahedral interstitial positions in the virgin martensite structure. During aging at room temperature, the interstitial atoms tend to agglomerate to regions of high carbon or nitrogen content. Carbon atoms are believed to cluster by jumping from tetrahedral interstitial sites and/or forming regions of ordered Fe₄C. Nitrogen atoms agglomerate to regions of ordered Fe₁₆N₂. On tempering above room temperature, the sequential precipitation of the ε or η, ? and θ carbides is observed in Fe-C systems. The α (Fe₁₆N₂) and γ′ (Fe₄N) nitrides are observed during the decomposition of Fe-N martensite. Mössbauer hyperfine spectra associated with each of these structures are described.     

Tempered martensite embrittlement in a high purity steel

This paper reports a study of tempered martensite embrittlement in a low alloy steel. Even though this material would ordinarily be considered high purity (0.004 wt pct S and 0.004 wt pct P) the 40 wppm S are sufficient to cause the embrittlement. The results show that this embrittlement will only occur if sulfur is present on the grain boundaries in its elemental form. It it is precipitated along the grain boundaries as chromium sulfides, no embrittlement trough is observed. However, the fracture energy of all samples is reduced. This is because microvoid coalescence occurs at these grain boundary precipitates in the same manner as is observed in overheated steels.