Thermally Stable Ni-rich Austenite Formed Utilizing Multistep Intercritical Heat Treatment in a Low-Carbon 10 Wt Pct Ni Martensitic Steel

Austenite reversion and its thermal stability attained during the transformation is key to enhanced toughness and blast resistance in transformation-induced-plasticity martensitic steels. We demonstrate that the thermal stability of Ni-stabilized austenite and kinetics of the transformation can be controlled by forming Ni-rich regions in proximity of pre-existing (retained) austenite. Atom probe tomography (APT) in conjunction with thermodynamic and kinetic modeling elucidates the role of Ni-rich regions in enhancing growth kinetics of thermally stable austenite, formed utilizing a multistep intercritical (Quench-Lamellarization-Tempering (QLT)-type) heat treatment for a low-carbon 10 wt pct Ni steel. Direct evidence of austenite formation is provided by dilatometry, and the volume fraction is quantified by synchrotron X-ray diffraction. The results indicate the growth of nm-thick austenite layers during the second intercritical tempering treatment (T-step) at 863 K (590 °C), with austenite retained from first intercritical treatment (L-step) at 923 K (650 °C) acting as a nucleation template. For the first time, the thermal stability of austenite is quantified with respect to its compositional evolution during the multistep intercritical treatment of these steels. Austenite compositions measured by APT are used in combination with the thermodynamic and kinetic approach formulated by Ghosh and Olson to assess thermal stability and predict the martensite-start temperature. This approach is particularly useful as empirical relations cannot be extrapolated for the highly Ni-enriched austenite investigated in the present study.     

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.     

On the Stability of Reversely Formed Austenite and Related Mechanism of Transformation in an Fe-Ni-Mn Martensitic Steel Aided by Electron Backscattering Diffraction and Atom Probe Tomography

The stability of reversely formed austenite and related mechanism of transformation were investigated against temperature and time in an Fe-9.6Ni-7.1Mn (at. pct) martensitic steel during intercritical annealing at a dual-phase (α + γ) region. Dilatometry, electron backscattering diffraction (EBSD), atom probe tomography (APT), and X-ray diffraction (XRD) were used to characterize the mechanism of reverse transformation. It was found that under intercritical annealing at 853 K (580 °C), when the heating rate is 20 K/s (20 °C/s), reverse transformation takes place through a mixed diffusion control mechanism, i.e., controlled by bulk diffusion and diffusion along the interface, where Ni controls the diffusion as its diffusivity is lower than that of Mn in the martensite and austenite. Increasing the intercritical annealing to 873 K (600 °C) at an identical heating rate of 20 K/s (20 °C/s) showed that reverse transformation occurs through a sequential combination of both martensitic and diffusional mechanisms. The transition temperature from diffusional to martensitic transformation was obtained close to 858 K (585 °C). Experimental results revealed that the austenite formed by the diffusional mechanism at 853 K (580 °C) mainly remains untransformed after cooling to ambient temperature due to the enrichment with Ni and Mn. It was also found that the stability of the reversely formed austenite by martensitic mechanism at 873 K (600 °C) is related to grain refinement.     

Structure and mechanical properties of Fe-Cr-Mo-C alloys with and without boron

A study of the structure and mechanical properties of Fe-Cr-Mo-C martensitic steels with and without boron addition has been carried out. Nonconventional heat treatments have subsequently been designed to improve the mechanical properties of these steels. Boron has been known to be a very potent element in increasing the hardenability of steel, but its effect on structure and mechanical properties of quenched and tempered martensitic steels has not been clear. The present results show that the as-quenched structures of both steels consist mainly of dislocated martensite. In the boron-free steel, there are more lath boundary retained austenite films. The boron-treated steel shows higher strengths at all tempering temperatures but with lower Charpy V-notch impact energies. Both steels show tempered martensite embrittlement when tempered at 350 °C for 1 h. The properties above 500 °C tempering are significantly different in the two steels. While the boron-free steel shows a continuous increase in toughness when tempered above 500 °C, the boron-treated steel suffers a second drop in toughness at 600 °C tempering. Transmission electron microscopy studies show that in the 600 °C tempered boron-treated steel large, more or less continuous cementite films are present at the lath boundaries, which are probably responsible for the embrittlement. The differences in mechanical properties at tempering temperatures above 500 °C are rationalized in terms of the effect of boron-vacancy interactions on the recovery and recrystallization behavior of these steels. Although boron seems to impair room temperature impact toughness at low strength levels, it does not affect this property at high strength levels. By simple nonconventinal heat treatments of the present alloys, martensitic steels may be produced with quite good strength-toughness properties which are much superior to those of existing commercial ultra-high strength steels. It is also shown that very good combinations of strength and toughness can be obtained with as-quenched martensitic steels.     

Reversible temper embrittlement of rotor steels

The susceptibility to temper embrittlement of eight different rotor steels has been studied in terms of the effects of composition, of cooling rate from tempering temperature, of isothermal aging, of steel-making practice and of strength level and tempering temperature. The Ni Cr Mo V steels tested showed increasing susceptibility to temper embrittlement with increasing nickel content. The normally marked susceptibility of a high phosphorus 3 pct Cr Mo steel was eliminated by the removal of manganese. Embrittlement in a 3 pct Ni Cr Mo V steel was caused by the equilibrium segregation of solute atoms to the prior austenite grain boundaries. Two Cr Mo V steels tested were not susceptible to temper embrittlement. Electroslag remelting and refining had very little effect on the susceptibility of the steels tested. Strength level and tempering temperature had no effect on the degree of embrittlement of the 3 pct Ni Cr Mo V disc steel. The possibilities of remedial action include an adjustment of the post tempering cooling rate, to optimize the conflicting interests of minimum temper embrittlement and adequate stress relief, and the production of very low manganese rotor steels.     

Tempering behaviour of a dual-phase low-alloy steel

The tempering behaviour of a dual phase steel of 0.08% C, 1.21% Mn, 1.00% Si, 0.42% Cr, and 0.41% Mo composition with two different martensite contents of 30 and 52%. (obtained by intercritical treatments at 820 and 860°C, respectively) has been studied. The ultimate tensile strength decreased and percentage elongation increased continuously with increasing tempering temperature up to 600°C for both intercritical treatments. The yield strength has, however, increased up to 300°C, beyond which it decreased for the steel with 30% martensite. In contrast it remained almost constant for 52% martensite up to 300°C, beyond which it decreased. The martensite of dual-phase steel for both the intercritical treatments has undergone microstructural changes on tempering that are akin to those of fully martensitic low carbon steels. The SEM fractographs from the as-quenched specimens indicate that the tensile specimens failed by microvoid coalescence with the martensite areas appearing facetted and featureless while those for 600°C tempered condition by the formation of equiaxed dimples.     

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.     

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).     

Nanoscale Analyses of High-Nickel Concentration Martensitic High-Strength Steels

Austenite reversion in martensitic steels is known to improve fracture toughness. This research focuses on characterizing mechanical properties and the microstructure of low-carbon, high-nickel steels containing 4.5 and 10 wt pct Ni after a QLT-type austenite reversion heat treatment: first, martensite is formed by quenching (Q) from a temperature in the single-phase austenite field, then austenite is precipitated by annealing in the upper part of the intercritical region in a lamellarization step (L), followed by a tempering (T) step at lower temperatures. For the 10 wt pct Ni steel, the tensile strength after the QLT heat treatment is 910 MPa (132 ksi) at 293 K (20 °C), and the Charpy V-notch impact toughness is 144 J (106 ft-lb) at 188.8 K (?84.4 °C, ?120 °F). For the 4.5 wt pct Ni steel, the tensile strength is 731 MPa (106 ksi) at 293 K (20 °C) and the impact toughness is 209 J (154 ft-lb) at 188.8 K (?84.4 °C, ?120 °F). Light optical microscopy, scanning electron and transmission electron microscopies, synchrotron X-ray diffraction, and local-electrode atom-probe tomography (APT) are utilized to determine the morphologies, volume fractions, and local chemical compositions of the precipitated phases with sub-nanometer spatial resolution. The austenite lamellae are up to 200 nm in thickness, and up to several micrometers in length. In addition to the expected partitioning of Ni to austenite, APT reveals a substantial segregation of Ni at the austenite/martensite interface with concentration maxima of 10 and 23 wt pct Ni for the austenite lamellae in the 4.5 and 10 wt pct Ni steels, respectively. Copper-rich and M₂C-type metal carbide precipitates were detected both at the austenite/martensite interface and within the bulk of the austenite lamellae. Thermodynamic phase stability, equilibrium compositions, and volume fractions are discussed in the context of Thermo-Calc calculations.     

Effect of Continuous Galvanizing Heat Treatments on the Microstructure and Mechanical Properties of High Al-Low Si Transformation Induced Plasticity Steels

Heat treatments were performed using an isothermal bainitic transformation (IBT) temperature compatible with continuous hot-dip galvanizing on two high Al–low Si transformation induced plasticity (TRIP)-assisted steels. Both steels had 0.2 wt pct C and 1.5 wt pct Mn; one had 1.5 wt pct Al and the other had 1 wt pct Al and 0.5 wt pct Si. Two different intercritical annealing (IA) temperatures were used, resulting in intercritical microstructures of 50 pct ferrite (α)-50 pct austenite (γ) and 65 pct α-35 pct γ. Using the IBT temperature of 465 °C, five IBT times were tested: 4, 30, 60, 90, and 120 seconds. Increasing the IBT time resulted in a decrease in the ultimate tensile strength (UTS) and an increase in the uniform elongation, yield strength, and yield point elongation. The uniform elongation was higher when using the 50 pct α-50 pct γ IA temperature when compared to the 65 pct α-35 pct γ IA temperature. The best combinations of strength and ductility and their corresponding heat treatments were as follows: a tensile strength of 895 MPa and uniform elongation of 0.26 for the 1.5 pct Al TRIP steel at the 50 pct γ IA temperature and 90-second IBT time; a tensile strength of 880 MPa and uniform elongation of 0.27 for the 1.5 pct Al TRIP steel at the 50 pct γ IA temperature and 120-second IBT time; and a tensile strength of 1009 MPa and uniform elongation of 0.22 for the 1 pct Al-0.5 pct Si TRIP steel at the 50 pct γ IA temperature and 120-second IBT time.     

Role of Tungsten in the Tempered Martensite Embrittlement of a Modified 9 Pct Cr Steel

The effect of tempering on the mechanical properties and fracture behavior of two 3 pct Co-modified 9 pct Cr steels with 2 and 3 wt pct W was examined. Both steels were ductile in tension tests and tough under impact tests in high-temperature tempered conditions. At T ≤ 923 K (650 °C), the addition of 1 wt pct W led to low toughness and pronounced embrittlement. The 9Cr2W steel was tough after low-temperature tempering up to 723 K (450 °C). At 798 K (525 °C), the decomposition of retained austenite induced the formation of discontinuous and continuous films of M₂₃C₆ carbides along boundaries in the 9Cr2W and the 9Cr3W steels, respectively, which led to tempered martensite embrittlement (TME). In the 9Cr2W steel, the discontinuous boundary films played a role of crack initiation sites, and the absorption energy was 24 J cm^?2. In the 9Cr3W steel, continuous films provided a fracture path along the boundaries of prior austenite grains (PAG) and interlath boundaries in addition that caused the drop of impact energy to 6 J cm^?2. Tempering at 1023 K (750 °C) completely eliminated TME by spheroidization and the growth of M₂₃C₆ carbides, and both steels exhibited high values of adsorbed energy of ≥230 J cm^?2. The addition of 1 wt pct W extended the temperature domain of TME up to 923 K (650 °C) through the formation of W segregations at boundaries that hindered the spheroidization of M₂₃C₆ carbides.     

A study of the hydrogen-induced degradation of two steels differing in sulfur content

Two steels with different sulfur contents: 0.003 and 0.024 wt pct, were cathodically charged under three different conditions and brought to fracture in tension immediately after charging or after aging at room temperature. All hydrogen charged specimens showed embrittlement, with a little higher loss of ductility in the high sulfur steel. The hydrogen embrittlement was reversible in both steels when specimens were charged in arsenic-free sulfuric acid solution at room temperature but was irreversible when charged in arsenic-containing acid at the same temperature. After charging in molten salts at 200 °C, some of the low sulfur steel specimens exhibited irreversible hydrogen damage with the appearance of quasicleavage fractures, while all high sulfur steel specimens were restored to the uncharged ductility by aging at room temperature. These results are interpreted by assuming that an increased sulfur content in steel increases the density of trapping sites for hydrogen at the sulfide/matrix interfaces. These traps are inactive above 150 °C and become operative after cooling. Therefore, at the same hydrogen content in steel after cooling, the greater content of sulfur results in a decreased activity of the lattice dissolved hydrogen, hence in reduced embrittlement.     

Strength of initially virgin martensites at - 196 °C after aging and tempering

The compressive strength at -196°C of martensites in Fe-0.26 pct C-24 pct Ni, Fe-0.4 pct C-21 pct Ni, and Fe-0.4 pct C-18 pct Ni-3 pct Mo alloys, all with subzero M temperatures, has been determined in the virgin condition and after one hour at temperatures from -80 to +400 °C. The effects of ausforming (20 pct reduction in area of the austenite by swaging at room temperature prior to the martensitic transformation) were also investigated. For the unausformed martensites, aging at temperatures up to 0 °C results in relatively small increases in strength. Above 0 °C, the age hardening increment increases rapidly, reaching a maximum at 100 °C. Above 100 °C, the strength decreases continuously with increasing tempering temperature except for the molybdenum-containing alloy, which exhibits secondary hardening on tempering at 400 °C. For the ausformed martensites, the response to aging at subzero temperatures is greater than for unausformed material. Strength again passes through a maximum on aging at 100 °C. However, on tempering just above 100 °C, the ausformed materials show a slower rate of softening than the unausformed martensites. The strengthening produced by the ausforming treatment is largest for the Fe-0.4 pct C-18 pct Ni-3 pct Mo alloy, but there is no evidence of carbide precipitation in the deformed austenite to a°Count for this effect of molybdenum.     

Design of strong tough Fe/Mo/C martensitic steels and the effects of cobalt

The microstructures and mechanical properties of a series of vacuum melted Fe/(2 to 4) Mo/(0.2 to 0.4) C steels with and without cobalt have been investigated in the as-quenched fully martensitic condition and after quenching and tempering for 1 h at 673 K (400°C) and 873 K (600°C); austenitizing was done at 1473 K (1200°C) in argon. Very good strength and toughness properties were obtained with the Fe/2 Mo/0.4 C alloy in the as-quenched martensitic condition and this is attributed mainly to the absence of internal twinning. The slightly inferior toughness properties compared to Fe/Cr/C steels is attributed to the absence of interlath retained austenite. The two 0.4 pct carbon steels having low Mo contents had approximately one-half the amount of transformation twinning associated with the two 0.4 pct carbon steels having high Mo contents. The plane strain fracture toughness of the steels with less twinning was markedly superior to the toughness of those steels with similar alloy chemistry which had more heavily twinned microstructures. Experiments showed that additions of Co to a given Fe/Mo/C steel raised M_s but did not decrease twinning nor improve toughness. Molybdenum carbide particles were found in all specimens tempered at 673 K (400°C). The Fe/Mo/C system exhibits secondary hardening after tempering at 873 K (600°C). The precipitate is probably Mo₂C. This secondary hardening is associated with a reduction in toughness. Additions of Co to Fe/Mo/C steels inhibited or eliminated the secondary hardening effect normally observed. Toughness, however, did not improve and in fact decreased with Co additions.     

Fracture toughness of aisi M2 high-speed steel and corresponding matrix tool steel

The influence of microstructural variations on the fracture toughness of two tool steels with compositions 6 pct W-5 pct Mo-4 pct Cr-2 pct V-0.8 pct C (AISI M2 high-speed steel) and 2 pct W-2.75 pct Mo-4.5 pct Cr-1 pct V-0.5 pct C (VASCO-MA) was investigated. In the as-hardened condition, the M2 steel has a higher fracture toughness than the MA steel, although the latter steel is softer. In the tempered condition, MA is softer and has a higher fracture toughness than M2. When the hardening temperature is below 1095 °C (2000 °F), tempering of both steels causes embrittlement,i.e., a reduction of fracture toughness as well as hardness. The fracture toughness of both steels was enhanced by increasing the grain size. The steel samples with intercept grain size of 5 (average grain diameter of 30 microns) or coarser exhibit 2 to 3 MPa√m (2 to 3 ksi√in.) higher fracture toughness than samples with intercept grain size of 10 (average grain diameter of 15 microns) or finer. Tempering temperature has no effect on the fracture toughness of M2 and MA steels as long as the final tempered hardness of the steels is constant. Retained austenite has no influence on the fracture toughness of as-hardened MA steel, but a high content of retained austenite appears to raise the fracture toughness of as-hardened M2 steel. There is a temperature of austenitization for each tool steel at which the retained austenite content in the as-quenched samples is a maximum. The above described results were explained through changes in the microstructure and the fracture modes. CHONGMIN KIM, formerly with Climax Molybdenum Company of Michigan, Ann Arbor, MI.     

Strength and toughness of Fe-10ni alloys containing C,Cr, Mo,and Co

The effects of C (0.10 to 0.20 pct), Cr (0 to 3 pct), Mo (0 to 2 pct), and Co (0 to 8 pct) on the yield strength, toughness (Charpy shelf energy), and tempering behavior of martensitic lONiCr-Mo-Co steels have been investigated. Variations in the carbon content between 0.10 and 0.20 pct result in yield strengths between 160 and 210 ksi (1.1 and 1.45 GN/m²) when these steels are tempered at 900° to 1000°F (480° to 540°C) for times of 1 to 100 h. These steels exhibit a secondary-hardening peak at 900° to 1000° F (480° to 540°C) where coarse Fe₃C carbides are gradually replaced by a fine, dislocation-nucleated dispersion of (Mo, Cr)₂C carbides. Maximum toughness at a given yield strength in these steels is only obtained when they are tempered for sufficiently long times so that the coarse Fe₃C carbides are completely dissolved. Molybdenum is primarily responsible for the secondary-hardening peak observed in these steels. However, chromium additions do result in lower secondaryhardening temperatures and promote coarsening of the secondary-hardening carbide. Best combinations of strength and toughness are obtained with steels containing 2 pct Cr and 1 pct Mo. Cobalt increases the yield strength of these steels over the entire tempering range and results in a higher secondary-hardening peak. This effect of cobalt is attributed to 1) a retardation in the rate of recovery of the dislocation substructure of the martensite, 2) the formation of a finer dispersion of secondary-hardening carbides, and 3) solid-solution strengthening. The finer dispersion of secondary-hardening carbides in steels containing cobalt is favored by the finer dislocation substructure in these steels since the (Mo, Cr)₂C carbide is dislocation-nucleated. This fine dispersion of (Mo, Cr)₂C carbide combined with the high nickel content accounts for the excellent combination of strength and toughness exhibited by these steels.     

The stress corrosion resistance and fatigue crack growth rate of a high strength martensitic stainless steel,AFC 77

The plane-strain stress corrosion thresholdK _Iscc and fatigue crack growth rate have been determined for a high strength martensitic stainless steel, AFC 77, in both conventionally processed and strain-aged conditions. TheK _Iscc (in 3.5 pct sodium chloride solution) is markedly affected by both the tempering temperature and the degree of strain aging. The highestK _Iscc of 105 ksi \(\sqrt {in} \) . was obtained by tempering at 500°F and the lowestK _Iscc of 10 ksi \(\sqrt {in} \) . by tempering at 1100°F. Retained austenite raisedK _Iscc at tempering temperatures up to 1000°F, which was the highest tempering temperature at which austenite could be maintained. Fatigue crack growth rates in both dry air (<10pct relative humidity) and 3.5 pct sodium chloride solution were at a maximum for material tempered at 700°F. Over the range of stress intensity studied, retained austenite reduced fatigue crack growth rate in salt solution but increased it in dry air.     

Characterization of the Carbon and Retained Austenite Distributions in Martensitic Medium Carbon, High Silicon Steel

The retained austenite content and carbon distribution in martensite were determined as a function of cooling rate and temper temperature in steel that contained 1.31 at. pct C, 3.2 at. pct Si, and 3.2 at. pct noniron metallic elements. Mössbauer spectroscopy, transmission electron microscopy (TEM), transmission synchrotron X-ray diffraction (XRD), and atom probe tomography were used for the microstructural analyses. The retained austenite content was an inverse, linear function of cooling rate between 25 and 560 K/s. The elevated Si content of 3.2 at. pct did not shift the start of austenite decomposition to higher tempering temperatures relative to SAE 4130 steel. The minimum tempering temperature for complete austenite decomposition was significantly higher (>650 °C) than for SAE 4130 steel (～300 °C). The tempering temperatures for the precipitation of transition carbides and cementite were significantly higher (>400 °C) than for carbon steels (100 °C to 200 °C and 200 °C to 350 °C), respectively. Approximately 90 pct of the carbon atoms were trapped in Cottrell atmospheres in the vicinity of the dislocation cores in dislocation tangles in the martensite matrix after cooling at 560 K/s and aging at 22 °C. The 3.2 at. pct Si content increased the upper temperature limit for stable carbon clusters to above 215 °C. Significant autotempering occurred during cooling at 25 K/s. The proportion of total carbon that segregated to the interlath austenite films decreased from 34 to 8 pct as the cooling rate increased from 25 to 560 K/s. Developing a model for the transfer of carbon from martensite to austenite during quenching should provide a means for calculating the retained austenite. The maximum carbon content in the austenite films was 6 to 7 at. pct, both in specimens cooled at 560 K/s and at 25 K/s. Approximately 6 to 7 at. pct carbon was sufficient to arrest the transformation of austenite to martensite. The chemical potential of carbon is the same in martensite that contains 0.5 to 1.0 at. pct carbon and in austenite that contains 6 to 7 at. pct carbon. There was no segregation of any substitutional elements.     

Aging embrittlement and grain boundary

“Clean” 3.5NiCrMoV steels with limited contents in trace elements (P, Sn, As, Sb) are commonly provided for manufacturing big rotor shafts. The possible increase in temperature in future steam turbines has promoted the development of “superclean” steels characterized by an extra drastic decrease of manganese and silicon contents. Their higher cost in comparison to “clean” steels leads to concern above which temperature they must be considered as mandatory for resisting aging embrittlement in operation. 3.5NiCrMoV “clean” steel samples (Mn = 0.30 pct; Si = 0.10 pct) were aged at 300 °C, 350 °C, and 400 °C for 10,000 hours up to 30,000 hours. No embrittlement results from aging at 300 °C and 350 °C, but holding at 400 °C is highly detrimental. Auger spectroscopy confirms that, when aging at 400 °C, phosphorus is the main embrittling trace element. It is suggested that grain boundary embrittlement is associated with the building of a layer that contains, on the one hand, Ni and P and, on the other hand, Mo and Cr. Head of the Testing and Head of the Testing and Head of the Testing and