Solid solutions in sputter-deposits of iron with 0 to 5 Wt Pct C

Supersaturated iron-carbon solid solutions containing 0.06, 0.18, 0.66, 2, 3, and 5 wt pct C were produced by sputter-deposition at 6° to 21°C. Homogeneous deposits of the same carbon composition as the multiphase source materials were produced by high rate (up to 0.0004 in. per hr) sputtering techniques and were 0.005 to 0.027 in. thick. The microhardnesses of the deposits were higher than the hardnesses of martensite with the same carbon content. The hardness increased rapidly from 680 Dph for 0.06 wt pct C content to an unusually high maximum of 1240 Dph at 2 wt pct C and then decreased slowly to 920 Dph at 5 wt pct C. The 0.06, 0.18, and 0.66 wt pct C deposits were bcc, and the 2 and 3 wt pct C deposits were bct. The tetragonality of the 5 wt pct C deposit was detected only after tempering at 150°C. Lattice parameterc/a ratios for the tetragonal deposits were lower than expected from extrapolated martensite data, and they corresponded to the equationc/a=1.06+0.019 (wt pct C). Nevertheless, thec/a ratio of 1.10 for 5 wt pct C deposit was higher than previously observed for martensite. No evidence of a martensitic transformation was found in the microstructures, which typically consisted of 0.5 to 1.0 μ diam columnar grains. Hardness data from tempered eposit samples and the lack of tetragonality of the low carbon deposits indicated that autotempering occurred during sputter-deposition.     

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).     

200, 020, and 002 X-ray peaks in tempered Fe-18 Ni-C martensites

Separate 200, 020, and 002 X-ray peaks were recorded for 0.0, 0.4, and 0.8 wt pct carbon (18 pct Ni) martensites after tempering between 25 and 500°C. The carbon bearing martensites studied here have been tempered initially enough to eliminate the “high tetragonality” 002 peak usually recorded for as-quenched martensite and the present results apply to tempered martensite only. The peak maximum is taken to determine the lattice parameter and the peak shape is recorded. At all carbon levels and after all tempering treatments, the “crd parameter is larger than or equal to the “a” or “b”. The relative enlargement is very small (0.08 pct) for the lowest carbon level and for any carbon level after severe tempering (500°C for 15 min). For the two higher carbon alloys tempered at temperatures below 400°C (for 15 min) the “c” parameter is significantly larger than the “a” and “b” and for the 0.4 wt pct C alloy the “b” is significantly smaller than the“a” whereas in the 0.8 pct C alloy the “b” is slightly larger than the “a”. Within experimental error the mean volume of the unit cell does not change during the tempering studied here and is nearly unaffected by the initial carbon content. This indicates that little (at most 0.1 wt pct) carbon is dissolved in tempered martensite. In the low carbon alloy the peaks are symmetric and sharpen symmetrically during tempering. In the higher carbon alloys the peaks are nearly symmetric and sharp after severe tempering. After less severe tempering the 002 peak is asymmetrically broadened toward lower9 values (higher lattice parameters) whereas the 200 and 020 peaks are asymmetrically broadened toward higher 0 values corresponding to lower lattice parameters. This collection of results is tentatively interpreted as being due to strains in martensite due to transformation induced substructure and precipitated carbides.     

Microstructure-strength relations in a hardenable stainless steel with 16 pct Cr, 1.5 pct Mo,and 5 pct Ni

Metallographic studies have been conducted on a 0.024 pct C-16 pct Cr-1.5 pct Mo-5 pct Ni stainless steel to study the phase reactions associated with heat treatments and investigate the strengthening mechanisms of the steel. In the normalized condition, air cooled from 1010 °C, the microstructure consists of 20 pct ferrite and 80 pct martensite. Tempering in a temperature range between 500 and 600 °C results in a gradual transformation of martensite to a fine mixture of ferrite and austenite. At higher tempering temperatures, between 600 and 800 °C, progressively larger quantities of austenite form and are converted during cooling to proportionally increasing amounts of fresh martensite. The amount of retained austenite in the microstructure is reduced to zero at 800 °C, and the microstructure contains 65 pct re-formed martensite and 35 pct total ferrite. Chromium rich M₂₃C₆ carbides precipitate in the single tempered microstructures. The principal strengthening is produced by the presence of martensite in the microstructure. Additional strengthening is provided by a second tempering treatment at 400 °C due to the precipitation of ultrafine (Cr, Mo) (C,N) particles in the ferrite.     

Tempered martensite embrittlement and intergranular fracture in an ultra-high strength sulfur doped steel

This paper reports a study of tempered martensite embrittlement in a Ni-Cr steel doped with 0.01 wt pct S. The segregation of sulfur to the grain boundaries and the associated embrittlement of this material are very dependent upon the austenitizing temperature. If the austenitizing temperature is below 1050 °C very little embrittlement and very little intergranular fracture are observed because sulfur remains precipitated as chromium sulfide. At higher austenitizing temperatures the sulfides dissolve and sulfur segregates to the grain boundaries. Because of the high bulk content, the sulfur concentration at the grain boundaries becomes great enough for the sulfides to reprecipitate there. This leads to low energy intergranular ductile fracture. However, some sulfur remains unprecipitated at the boundary and can lower the cohesive strength across the boundary. When plate-like cementite precipitates at the grain boundary during tempering heat treatments at 300 to 400 °C, the combination of the carbides and the unprecipitated sulfur causes intergranular fracture and tempered martensite embrittlement.     

200, 020, and 002 X-ray peaks in tempered Fe-18 Ni-C martensites

Separate 200, 020, and 002 X-ray peaks were recorded for 0.0, 0.4, and 0.8 wt pct carbon (18 pct Ni) martensites after tempering between 25 and 500°C. The carbon bearing martensites studied here have been tempered initially enough to eliminate the “high tetragonality” 002 peak usually recorded for as-quenched martensite and the present results apply to tempered martensite only. The peak maximum is taken to determine the lattice parameter and the peak shape is recorded. At all carbon levels and after all tempering treatments, the “crd parameter is larger than or equal to the “a” or “b”. The relative enlargement is very small (0.08 pct) for the lowest carbon level and for any carbon level after severe tempering (500°C for 15 min). For the two higher carbon alloys tempered at temperatures below 400°C (for 15 min) the “c” parameter is significantly larger than the “a” and “b” and for the 0.4 wt pct C alloy the “b” is significantly smaller than the“a” whereas in the 0.8 pct C alloy the “b” is slightly larger than the “a”. Within experimental error the mean volume of the unit cell does not change during the tempering studied here and is nearly unaffected by the initial carbon content. This indicates that little (at most 0.1 wt pct) carbon is dissolved in tempered martensite. In the low carbon alloy the peaks are symmetric and sharpen symmetrically during tempering. In the higher carbon alloys the peaks are nearly symmetric and sharp after severe tempering. After less severe tempering the 002 peak is asymmetrically broadened toward lower9 values (higher lattice parameters) whereas the 200 and 020 peaks are asymmetrically broadened toward higher 0 values corresponding to lower lattice parameters. This collection of results is tentatively interpreted as being due to strains in martensite due to transformation induced substructure and precipitated carbides.     

Internal friction in ferrous martensites

Internal friction measurements were made at 1 Hz in the temperature range of 25 to 500°C on quenched, tempered and cold-worked Fe-Ni-C martensites. The alloys, which contained 15 to 30 wt pct Ni and 0.1 to 1.0 wt pct C, andM _s temperatures <0°C, transformed to martensite with a twinned (259)_γ habit and exhibited a relaxation peak at ~160°C. These results could be contrasted with those for Fe-C martensites, which form above room temperature, have predominantly dislocated (111)_γ or (225)_γ habits, and exhibit an internal friction peak at about 250°C. The literature on substructures, tempering and internal friction of all ferrous martensites and cold-worked ferrites was utilized in the interpretation of the 160°C peak. Several dissimilarities between the 160°C peak and the 250°C peak in Fe-C martensites or the cold-work peak in ferrite were noted such that models of dislocation-interstitial interaction for these peaks could not explain the 160°C peak. It was concluded that the 160°C peak is associated with the stress-induced motion of twin boundaries containing mobile carbon atoms. Such a mechanism was shown to be consistent with the present experimental observations and all other available data.     

Internal friction in ferrous martensites

Internal friction measurements were made at 1 Hz in the temperature range of 25 to 500°C on quenched, tempered and cold-worked Fe-Ni-C martensites. The alloys, which contained 15 to 30 wt pct Ni and 0.1 to 1.0 wt pct C, andM _s temperatures <0°C, transformed to martensite with a twinned (259)_γ habit and exhibited a relaxation peak at ~160°C. These results could be contrasted with those for Fe-C martensites, which form above room temperature, have predominantly dislocated (111)_γ or (225)_γ habits, and exhibit an internal friction peak at about 250°C. The literature on substructures, tempering and internal friction of all ferrous martensites and cold-worked ferrites was utilized in the interpretation of the 160°C peak. Several dissimilarities between the 160°C peak and the 250°C peak in Fe-C martensites or the cold-work peak in ferrite were noted such that models of dislocation-interstitial interaction for these peaks could not explain the 160°C peak. It was concluded that the 160°C peak is associated with the stress-induced motion of twin boundaries containing mobile carbon atoms. Such a mechanism was shown to be consistent with the present experimental observations and all other available data.     

The mechanical stability of austenite and cryogenic toughness of ferritic Fe-Mn-AI Alloys

In an attempt to understand the role of retained austenite on the cryogenic toughness of a ferritic Fe-Mn-AI steel, the mechanical stability of austenite during cold rolling at room temperature and tensile deformation at ambient and liquid nitrogen temperature was investigated, and the microstructure of strain-induced transformation products was observed by transmission electron microscopy (TEM). The volume fraction of austenite increased with increasing tempering time and reached 54 pct after 650 °C, 1-hour tempering and 36 pct after 550 °C, 16-hour tempering. Saturation Charpy impact values at liquid nitrogen temperature were increased with decreasing tempering temperature, from 105 J after 650 °C tempering to 220 J after 550 °C tempering. The room-temperature stability of austenite varied significantly according to the(α + γ) region tempering temperature;i.e., in 650 °C tempered specimens, 80 to 90 pct of austenite were transformed to lath martensite, while in 550 °C tempered specimens, austenite remained untransformed after 50 pct cold reductions. After tensile fracture (35 pct tensile strain) at -196 °C, no retained austenite was observed in 650 °C tempered specimens, while 16 pct of austenite and 6 pct of e-martensite were observed in 550 °C tempered specimens. Considering the high volume fractions and high mechanical stability of austenite, the crack blunting model seems highly applicable for improved cryogenic toughness in 550 °C tempered steel. Other possible toughening mechanisms were also discussed. Formerly Graduate Student, Seoul National University.     

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. Some in 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 films 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 situation is analogous to that in upper bainite. This TME failure is different from temper embrittlement (TE) which occurs 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 occur in the same steels, depending on the tempering conditions.     

An approach to developing an alternative hot work die steel

This paper evaluates an approach to developing a steel which combines resistance to softening on tempering with an economical use of alloy elements. To obtain the desired tempering behavior advantage has been taken of the ability of non-carbide forming elements to enhance the secondary hardening response and the tendency of vanadium additions to reduce the coarsening rate of Mo₂C. Five alloys were investigated; these were a base secondary hardening steel, that base steel modified by the addition of 2 wt pct silicon and by the combined addition of 1 wt pct silicon and 1 wt pct aluminum. To these two modified alloys were made additions of 0.4 wt pct vanadium. It was found that both types of additions without vanadium enhance the secondary hardening response to the same significant degree. Both of these alloys soften rapidly as the tempering temperature is increased above 600 °C. However, with the addition of vanadium, hardnesses over R_gC 50 are obtained after tempering at 650 °C. While silicon additions appear indispensable to this resistance to softening on tempering, silicon also favors the retention of primary carbides after austenitizing and, if present in sufficient amounts, can cause brittle intergranular fracture after tempering at high temperatures.     

Martensite lattice changes during tempering

An X-ray diffractometer study of martensite formed in an 18 wt pct nickel, 0.98 wt pct carbon austenite single crystal yields the shapes, positions, and integrated intensities of 200, 020, and 002 peaks. Martensite, which forms below ? 60 °C, was tempered at successively higher temperatures from ?45 to 450 °C. The results show that after subambient aging, during which C atoms in c-oriented octahedral sites have clustered, carbide precipitation starts and small regions (～30Å in the [001]) with negative tetragonality appear. Upon subsequent tempering these are augmented by larger regions which have small positive tetragonality. In this process the “c” lattice parameter changes markedly but the “a” and “b” lattice parameter increase very little. These results indicate the formation of carbon depleted martensite which is coherently strained by the carbide particles. At and just above 100 °C the 200, 020, and 002 peaks all become doublets as the martensite matrix discontinuously breaks free of coherency and becomes highly imperfect ferrite. This change also occurs during the so-called “first stage of tempering.” Further tempering decreases the defect content of this ferrite. The lattice of the martensite is extensively reoriented during tempering just above room temperature. These reorientations probably accommodate the lattice parameter changes described above and may be carried out by movement of twin boundaries.     

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)     

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.     

Creep Rupture Strength and Microstructure of Low C-10Cr-2Mo Heat-Resisting Steels with V and Nb

The new ferritic heat-resisting steels of 0.05C-10Cr-2Mo-0.10V-0.05Nb (Cb) composition with high creep rupture strength and good ductility have already been reported. The optimum amounts of V and Nb that can be added to the 0.05C-10Cr-2Mo steels and their effects on the creep rupture strength and microstructure of the steels have been studied in this experiment. The optimum amounts of V and Nb are about 0.10 pct V and 0.05 pct Nb at 600 °C for 10,000 h, but shift to 0.18 pct V and 0.05 pct Nb at 650 °C. Nb-bearing steels are preferred to other grades on the short-time side, because NbC precipitation during initial tempering stages delays recovery of martensite. On the long-time side, however, V-bearing steels have higher creep rupture strength. By adding V to the steels, electron microscopic examination reveals a stable microstructure, retardation during creep of the softening of tempered martensite, fine and uniform distribution of precipitates, and promotion of the precipitation of Fe₂Mo.     

Response of Phase Transformation Inducing Heat Treatments on Microstructure and Mechanical Properties of Reduced Activation Ferritic-Martensitic Steels of Varying Tungsten Contents

Microstructure and mechanical properties of 9Cr-W-0.06Ta Reduced Activation Ferritic-Martensitic (RAFM) steels having various tungsten contents ranging from 1 to 2 wt pct have been investigated on subjecting the steels to isothermal heat treatments for 5 minutes at temperatures ranging from 973 K to 1473 K (700 °C to 1200 °C) (below Ac₁ to above Ac₃) followed by oil quenching and tempering at 1033 K (760 °C) for 60 minutes. The steels possessed tempered martensite structure at all the heat-treated conditions. Prior-austenitic grain size of the steels was found to decrease on heating in the intercritical temperature range (between Ac₁ and Ac₃) and at temperatures just above the Ac₃ followed by increase at higher heating temperatures. All the steels suffered significant reduction in hardness, tensile, and creep strength on heating in the intercritical temperature range, and the reduction was less for steel having higher tungsten content. Strength of the steels increased on heating above Ac₃ and was higher for higher tungsten content. Transmission Electron Microscopy (TEM) investigations of the steels revealed coarsening of martensitic substructure and precipitates on heating in the intercritical temperature range, and the coarsening was relatively less for higher tungsten content steel, resulting in less reduction in tensile and creep strength on intercritical heating. Tensile and creep strengths of the steels at different microstructural conditions have been rationalized based on the estimated inter-barrier spacing to dislocation motion. The study revealed the uniqueness of inter-barrier spacing to dislocation motion in determining the strength of tempered martensitic steels subjected to different heat treatments.     

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 occurrence 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 microvoids at the retained carbides.     

A study of the early stages of tempering of iron-carbon martensites by atom probe field ion microscopy

The redistribution of carbon atoms during the early stages of ageing and tempering of iron-carbon martensites has previously been studied only by indirect methods. The computer-controlled atom probe field ion microscope permits the direct, quantitative determination of carbon concentrations at the atomic level, and thus all the stages of the martensite decomposition process become amenable to direct study. Analyses of a low-carbon martensite, Fe-1.0 at. pct C, (Fe-0.21 wt pct C), water quenched and tempered for 10 min at 150 °C, showed a matrix carbon content of only 0.14 at. pct. Analysis of a 2 nm diam area centered on a lath boundary showed a local concentration of 2.01 at. pct C. There is some evidence that this carbon level is associated with the presence of a thin film of retained austenite at the boundary. In the case of a higher carbon martensite, Fe-0.64 at. pct Mn, 3.47 at. pct C, (Fe-0.65 wt pct Mn-0.78 wt pct C) water quenched and aged for approximately 24 h at room temperature, analysis of twinned regions showed a matrix carbon level of 2.7 at. pct and a concentration enrichment to 6.9 at. pct in a region 2 nm diam, centered on the coherent twin interface. Assuming the segregated carbon to be located in a single atomic layer at the twin interface, this result indicates that a carbon concentration of 24 at. pct exists locally at the boundary. These results appear to be the first direct demonstration of the segregation of carbon atoms to lattice defects in carbon martensites. Tempering of the higher carbon martensite for 1 h at 160 °C produced further segregation of carbon to the region of twin interfaces. The matrix carbon content fell to 1.5 at. pct and the average carbon content over a 2 nm diam region at the interface rose to 8.7 at. pct. The width of the carbon segregated regions also increased, which seems to imply that incipient carbide precipitation in the plane of the twin boundaries is occurring at this stage of the tempering process. Formerly with the Department of Metallurgy and Science of Materials, University of Oxford     

The effects of heat treatment on fracture toughness and fatigue crack growth Rates in 440C and BG42 steels

The fatigue crack growth rates,da/dN, and the fracture toughness, K_Ic have been measured in two high-carbon martensitic stainless steels, 440C and BG42. Variations in the retained austenite contents were achieved by using combinations of austenitizing temperatures, refrigeration cycles, and tempering temperatures. In nonrefrigerated 440C tempered at 150 °C, about 10 vol pct retained austenite was transformed to martensite at the fracture surfaces duringK _Ic testing, and this strain-induced transformation contributed significantly to the fracture toughness. The strain-induced transformation was progressively less as the tempering temperature was raised to 450 °C, and at the secondary hardening peak, 500 °C, strain-induced transformation was not observed. In nonrefrigerated 440C austenitized at 1065 °C,K _Ic had a peak value of 30 MPa m^1/2 on tempering at 150 °C and a minimum of 18 MPa m^1/2 on tempering at 500 °C. Refrigerated 440C retained about 5 pct austenite, and did not exhibit strain-induced transformation at the fracture surfaces for any tempering temperature. TheK _Ic values for corresponding tempering temperatures up to the secondary peak in refrigerated steels were consistently lower than in nonrefrigerated steels. All of the BG42 specimens were refrigerated and double or quadruple tempered in the secondary hardening region; theK _Ic values were 16 to 18 MPa m^1/2 at the secondary peak. Tempered martensite embrittlement (TME) was observed in both refrigerated and nonrefrigerated 440C, and it was shown that austenite transformation does not play a role in the TME mechanism in this steel. Fatigue crack propagation rates in 440C in the power law regime were the same for refrigerated and nonrefrigerated steels and were relatively insensitive to tempering temperatures up to 500 °C. Above the secondary peak, however, the fatigue crack growth rates exhibited consistently lower values, and this was a consequence of the tempering of the martensite and the lower hardness. Nonrefrigerated steels showed slightly higher threshold values, ΔK_th, and this was ascribed to the development of compressive residual stresses and increased surface roughening in steels which exhibit a strain-induced martensitic transformation.