Morphology and properties of low-carbon bainite

Morphology of low-carbon bainite in commercial-grade high-tensile-strength steels in both isothermal transformation and continuous cooling transformation is lathlike ferrite elongated in the 〈11l〉_b direction. Based on carbide distribution, three types of bainites are classified: Type I, is carbide-free, Type II has fine carbide platelets lying between laths, and Type III has carbides parallel to a specific ferrite plane. At the initial stage of transformation, upper bainitic ferrite forms a subunit elongated in the [−101]f which is nearly parallel to the [lll]_b direction with the cross section a parallelogram shape. Coalescence of the subunit yields the lathlike bainite with the [−101]_f growth direction and the habit plane between (232)_f and (lll)_f. Cementite particles precipitate on the sidewise growth tips of the Type II bainitic ferrite subunit. This results in the cementite platelet aligning parallel to a specific ferrite plane in the laths after coalescence. These morphologies of bainites are the same in various kinds of low-carbon high-strength steels. The lowest brittle-ductile transition temperature and the highest strength were obtained either by Type III bainite or bainite/martensite duplex structure because of the crack path limited by fine unit microstructure. It should also be noted that the tempered duplex structure has higher strength than the tempered martensite in the tempering temperature range between 200 °C and 500 °C. In the case of controlled rolling, the accelerated cooling afterward produces a complex structure comprised of ferrite, cementite, and martensite as well as BI-type bainite. Type I bainite in this structure is refined by controlled rolling and plays a very important role in improving the strength and toughness of low-carbon steels. This paper is based on a presentation made in the symposium “International Conference on Bainite” presented at the 1988 World Materials Congress in Chicago, IL, on September 26 and 27, 1988, under the auspices of the ASM INTERNATIONAL Phase Transformations Committee and the TMS Ferrous Metallurgy Committee.     

Bainite morphologies in a 0.2 C–1.5 Mn steel

The isothermal transformation products of austenite over a wide range of temperatures and times in the bainitic range in a 0.2 wt.% C–1.5 wt.% Mn steel have been studied by transmission electron microscopy in order to characterise the bainitic microstructures in low-carbon low-alloy steels. Widmanstätten ferrite has formed with alternate layers of austenite (martensite) as a transition product at 600 and 500°C that has finally transformed on further isothermal transformation to either pearlite (at 600°C) or upper bainite (at 500°C). This type of transformation product was referred to as BI bainite by earlier investigators, but on the basis of the present investigation it is concluded that such ferrite-austenite (martensite) structures are not bainitic as this is not the final transformation product either at 600 or 500°C. Both upper bainite and lath-type lower bainite are formed at 450°C while the transformation product has been only lath-type lower bainite at 400°C.     

A Study of the Influence of Thermomechanical Controlled Processing on the Microstructure of Bainite in High Strength Plate Steel

Steels with compositions that are hot rolled and cooled to exhibit high strength and good toughness often require a bainitic microstructure. This is especially true for plate steels for linepipe applications where strengths in excess of 690 MPa (100 ksi) are needed in thicknesses between approximately 6 and 30 mm. To ensure adequate strength and toughness, the steels should have adequate hardenability (C. E. >0.50 and Pcm >0.20), and are thermomechanically controlled processed, i.e., controlled rolled, followed by interrupted direct quenching to below the Bs temperature of the pancaked austenite. Bainite formed in this way can be defined as a polyphase mixture comprised a matrix phase of bainitic ferrite plus a higher carbon second phase or micro-constituent which can be martensite, retained austenite, or cementite, depending on circumstances. This second feature is predominately martensite in IDQ steels. Unlike pearlite, where the ferrite and cementite form cooperatively at the same moving interface, the bainitic ferrite and MA form in sequence with falling temperature below the Bs temperature or with increasing isothermal holding time. Several studies have found that the mechanical properties may vary strongly for different types of bainite, i.e., different forms of bainitic ferrite and/or MA. Thermomechanical controlled processing (TMCP) has been shown to be an important way to control the microstructure and mechanical properties in low carbon, high strength steel. This is especially true in the case of bainite formation, where the complexity of the austenite-bainite transformation makes its control through disciplined processing especially important. In this study, a low carbon, high manganese steel containing niobium was investigated to better understand the effects of austenite conditioning and cooling rates on the bainitic phase transformation, i.e., the formation of bainitic ferrite plus MA. Specimens were compared after transformation from recrystallized, equiaxed austenite to deformed, pancaked austenite, which were followed by seven different cooling rates ranging between 0.5 K/s (0.5 °C/s) and 40 K/s (40 °C/s). The CCT curves showed that the transformation behaviors and temperatures varied with starting austenite microstructure and cooling rate, resulting in different final microstructures. The EBSD results and the thermodynamics and kinetics analyses show that in low carbon bainite, the nucleation rate is the key factor that affects the bainitic ferrite morphology, size, and orientation. However, the growth of bainite is also quite important since the bainitic ferrite laths apparently can coalesce or coarsen into larger units with slower cooling rates or longer isothermal holding time, causing a deterioration in toughness. This paper reviews the formation of bainite in this steel and describes and rationalizes the final microstructures observed, both in terms of not only formation but also for the expected influence on mechanical properties.     

Characterization of Transformation Stasis in Low-Carbon Steels Microalloyed with B and Mo

In the present study, bainite transformation kinetics was examined in low C-Mn steels with the addition of small amounts of B and Mo. This addition delays the onset of the bainite transformation. Mo addition causes transformation stasis at temperatures between 873 K and 823 K (600 °C and 550 °C) just below the bainite-start (B _s) temperature, resulting from an incomplete bainite transformation. Post-stasis transformation after a prolonged hold proceeds by the formation of ferrite with a low dislocation density, and in Mo-containing alloys, often the formation of carbides. The volume fraction at which the transformation stops is higher for lower carbon contents and lower transformation temperatures. By contrast, at 773 K (500 °C), the bainite transformation accompanying cementite precipitation occurs regardless of microalloying and is completed after shorter hold times. EDX measurement performed on the Mo-added 0.15 pct C alloy with aberration-corrected STEM revealed that segregation at the bainite/austenite interphase boundary is small for Mn and negligible for Mo in the early stages of stasis, which does not support the incomplete transformation mechanism based on the solute drag theory for the alloys used.     

Identification of Inverse Bainite in Fe-0.84C-1Cr-1Mn Hypereutectoid Low Alloy Steel

A unique dilatation trend is observed for isothermal bainite transformation in Fe-0.84 pct C-1 pct Cr-1 pct Mn steel. The dilatation is found to occur in two stages with volumetric contraction dominating the first stage, followed by volumetric expansion dominating the second stage. Through electron microscopic characterization, bainitic microstructure is identified as inverse bainite with cementite (Fe₃C) nucleating first from supersaturated austenite followed by the transformation of ferrite and secondary carbides (Fe₃C, Fe₂C, and Fe₅C₂) from carbon-depleted austenite.     

The effect of chromium in high carbon bearing steels

Chromium is present in the 52100 bearing steel composition in the range 1.30 to 1.60 wt pct. Chromium has a significant influence on the spheroidization of cementite, finer carbides being formed due to chromium additions. The ferrite to austenite transformation temperatures are increased due to chromium. The microchemistry of the cementite in 52100 changes during heat treatment; the chromium content of secondary carbides is generally lower than that in the spheroidal (FeCr)₃C produced by soft annealing. The rate of carbide dissolution is controlled by the rate of chromium diffusion from the carbide-matrix interface. Also, the chromium content of the residual, spheroidal (FeCr)₃C increases during austenitization. The effect of substitution of chromium in bearing steel compositions is discussed. In view of the beneficial effect of chromium, only substitution by similar strong carbide forming elements should be considered for bearing steels.     

On the bainite structure in Cr-Mo-V rotor steel

The morphology of continuously cooled and isothermally transformed bainite structures formed in a Cr-Mo-V rotor steel has been studied using transmission electron microscopy. The samples were austenitised at 955°C for an hour followed by air cooling to room temperature. The isothermal transformation reaction was carried out at 450°C for up to 100 000 s. The microconstituents observed are predominantly lower bainite with very small amount of upper bainite and martensite (formed from untransformed austenite due to water quenching). Analysis of the selected area diffraction patterns confirm that the carbide in bainite is orthorhombic cementite and the orientation relationship between ferrite and cementite is consistent with that of Bagaryatskii. The carbide particles in isothermally transformed bainite are coarser than those of continuously cooled bainite. Tempering one hour at 670°C of continuously cooled steel samples exhibited formation of fine spheroidal MC type carbides. In addition tempering leads to the enrichment of prior austenite grain boundaries by cementite particles. Tempering ten hours at 670°C exhibited microstructures almost identical to those observed in one hour tempering.     

High temperature plastic-flow behavior of mixtures of austenite,cementite, ferrite,and pearlite in plain-carbon steels

The plastic-flow behavior of ferrite + pearlite, pearlite + cementite, and austenite + cementite mixtures in plain carbon steels has been examined over the temperature range 500 to 1050 °C, strain-rate range 6 x l0⁻⁶ to 2 x l0⁻² s⁻¹, and carbon range 0.005C to 1.89C. Up to the eutectoid temperature the strength of the ferrite + pearlite mixture more than doubles as the carbon content increases from 0.005C to 0.7C, so that whereas in low-carbon steels the ferrite is weaker than the higher temperature austenite phase, in eutectoid steels the fully pearlitic structure is stronger than the fully austenitic structure. Manganese and silicon strengthen ferrite more effectively than they do austenite. A 0.17 pct phosphorus addition strengthens the ferrite + pearlite mixture across the range of microstructures from fully ferritic to fully pearlitic. Beyond the eutectoid composition, the amount of proeutectoid cementite does not significantly affect the strength of the pearlite, but above the eutectoid temperature it appreciably strengthens the austenite and cementite mixture at the strain rate of 2 X 10^-2 s^-1.     

Interrelations of cooling rate,microstructure, and mechanical properties in four HSLA steels

The present study was carried out on four steels containing 0.1 pct C-1.5 pct Mn-0.003 pct B* in common, with additions of 1 pct Cr, 0.5 pct Mo, 0.25 pct Mo + 1 pct Cr, 0.2 pct Ti + 1 pct Cr. They were designated, accordingly, as Cr, Mo, Mo-Cr, and Cr-Ti steels. All the steels exhibited a complete lath martensite microstructure with thin interlaths of retained austenite (≈0.05 pct) in the quenched condition. The normalized microstructures, granular bainite, contained massive areas of ferrite and granules of bainite laths. Both microconstituents contained a fine dispersion of cementite particles (size ≈50 Å) together with high dislocation densities. A mechanism explaining their for-mation has been given. The Cr steel, due to its low hardenability, showed in addition polygonal ferrite in the neighborhood of the so-called M-A constituent (twinned martensite and/or austenite). The annealed microstructure (using a cooling rate of 0.033 °C s^?1) of the Cr steel consisted of coarse ferrite-pearlite. Addition of 0.2 pct Ti to the Cr steel markedly refined the structure, whereas an addition of 0.25 pct Mo altered the microstructure to ferrite-lower bainite. In the 0.5 pct Mo steel, polygonal ferrite was found to be completely missing. The mechanical properties of the four steels after quenching, normalizing, and annealing were investigatedvia hardness and tensile test mea-surements. An empirical equation, relating the ultimate tensile strength to the steel composition, for steels that had granular bainite microstructures in the normalized condition, was proposed. The fracture surfaces exhibited cleavage and variable-size dimples depending on the microstructure and steel composition.     

Metallography of bainitic transformation in silicon containing steels

The formation of carbide in lower bainite was studied in two silicon containing carbon steels by transmission electron microscopy and diffraction techniques. Epsilon carbide was identified in the low temperature isothermally transformed bainite structure. The crystallographic relationship between epsilon carbide and bainitic ferrite was found to follow the Jack orientation relationship,viz, (0001)ε l l(011)α, (101l)ε l 1(101)α. The cementite observed in lower bainite was in the shape of small platelets and obeyed the Isaichev orientation relationship with the bainitic ferrite,viz, (010) cl 1(1-11)α, (103) cl 1 (011)α. Direct evidence showing the sequence of carbide formation from aus-tenite in bainite has also been obtained. Based on the observations and all the crystallo-graphical features, it is strongly suggested that in silicon containing steels the bainitic carbide precipitated directly from austenite instead of from ferrite at the austenite/fer-rite interface as has been proposed by Kinsman and Aaronson (Ref. 1). The uniformity of the carbide distribution is thus envisaged to be the outcome of precipitation at the aus-tenite-ferrite interphase boundary. DER-HUNG HUANG, formerly with the Department of Materials Science and Mineral Engineering, University of California     

The morphology,crystallography, and mechanism of carbide precipitation in an Fe-0.12 Pct C-3.28 Pct Ni alloy

Carbide precipitation during the eutectoid decomposition of austenite has been studied in an Fe-0.12 pct C-3.28 pct Ni alloy by transmission electron microscopy (TEM) supplemented by optical microscopy. Nodular bainite which forms during the latter stages of austenite decomposition at 550 °C exhibits two types of carbide arrangement: (a) banded interphase boundary carbides with particle diameters of about 20 to 90 nm and mean band spacings between 180 and 390 nm and (b) more randomly distributed (“nonbanded”) elongated particles exhibiting a wide range of lengths between 33 and 2500 nm, thicknesses of approximately 11 to 50 nm, and mean intercarbide spacings of approximately 140 to 275 nm. Electron diffraction analysis indicated that in both cases, the carbides are cementite, obeying the Pitsch orientation relationship with respect to the bainitic ferrite. The intercarbide spacings of both morphologies are significantly larger than those previously reported for similar microstructures in steels containing alloy carbides other than cementite (e.g., VC, TiC). Both curved and straight cementite bands were observed; in the latter case, the average plane of the interphase boundary precipitate sheets was near {110}_α//{011}_c consistent with cementite precipitation on low-energy {110}_α//{111}_γ ledge terrace planes (where α,β, andc refer to ferrite, austenite, and cementite, respectively). The results also suggest that the first stage in the formation of the nonbanded form of nodular bainite is often the precipitation of cementite rods, or laths, in austenite at the α:γ interfaces of proeutectoid ferrite secondary sideplates formed earlier. Although these cementite rods frequently resemble the “fibrous” microstructures observed by previous investigators in carbide-forming alloy steels, they are typically much shorter than fibrous alloy carbides. The bainitic microstructures observed here are analyzed in terms of a previously developed model centered about the roles of the relative nucleation and growth rates of the product phases in controlling the evolution of eutectoid microstructures.     

Influence of carbon concentration and reaction temperature upon bainite morphology in Fe-C-2 Pct Mn alloys

Isothermal transformation of austenite to bainite was studied by optical, replication, and thin foil transmission electron microscopy (TEM) in both hypo- and hypereutectoid Fe-C-2 wt pct Mn alloys (and a single 3 pct Mn alloy) containing from 0.1 to 1.37 wt pct C in order to characterize and explain the transitions which occur in bainite morphology as a function of carbon content and reaction temperature. A “morphology map” was constructed showing the temperaturecarbon composition(T-x) regions in which four different bainite morphologies predominate: upper bainite, lower bainite, nodular bainite, and inverse bainite. Calculations of the volume free energy changes associated with the nucleation of ferrite, ?G _v ^α , and cementite, ?G _v ^c , and the parabolic rate constant for growth of ferrite, α_α, and cementite, α_c, were performed in order to explain the observed morphological transitions. In hypoeutectoid alloys, where |?G _v ^α | ? |ΔG _v ^c | and α_α ? α _c , the Widmanstätten ferrite-dominated morphologies of upper and lower bainite predominate. The upper-to-lower bainite transition appears to be associated with the emergence of edge-to-face sympathetic nucleation at high values of |ΔG _v ^α |. In hypereutectoid alloys, the ratios ΔG _v ^α /ΔG _v ^c and α_α/α_c are considerably smaller; hence, cementite can compete much more readily with ferrite during both nucleation and growth, resulting in the formation of nodular bainite. With decreasing temperature in the hypereutectoid regime,ΔG _v ^α /ΔG _v ^c , and especially α_ga/α_c, increase, resulting in the replacement of nodular bainite by lower bainite at temperatures below about 250 °C to 275°C.     

Volume Fractions of Proeutectoid Ferrite/Pearlite and Their Dependence on Prior Austenite Grain Size in Hypoeutectoid Fe-Mn-C Alloys

It has been generally believed that pearlite transformation in hypoeutectoid steels starts when the average carbon concentration in untransformed austenite reaches the A_cm line after the formation of proeutectoid ferrite. To test this concept experimentally, volume fractions of proeutectoid ferrite/pearlite and carbon contents in the austenite being transformed into pearlite were measured for the Fe-2Mn-0.3C alloy isothermally transformed in the temperature range 848 K to 898 K (575 °C to 625 °C). It was found that lamellar pearlite can form even when the average carbon content in untransformed austenite is much lower than the A_cm line. This peculiar observation is probably due to the two-dimensional diffusion of carbon, i.e., parallel to and normal to the austenite/pearlite interface, which enables lamellar cementite to grow continuously by supplying carbon atoms to its growth front. This results in proeutectoid ferrite fractions with respect to pearlite being much lower than those predicted by the lever rule. With decreasing prior austenite grain size, proeutectoid ferrite fractions with respect to pearlite were found to increase, but the thickness of proeutectoid ferrite was constant within the range of grain size investigated. This is due to the existence of the critical α/γ interface velocity only below which pearlite (actually cementite) can be nucleated at the migrating α/γ interface. Furthermore, the upper limit temperatures for pearlite formation in the Fe-1Mn-0.33C and Fe-2Mn-0.3C alloys were found to be well between the PLE/NPLE and PE A_e1 temperatures.     

Microstructural evolution during thermomechanical processing of a Ti-Nb interstitial-free steel just below the Ar 3 temperature

Laboratory thermomechanical processing (TMP) experiments have been carried out to study the austenite transformation characteristics, precipitation behavior, and recrystallization of deformed ferrite for an interstitial-free (IF) steel in the temperature range just below Ar ₃. For cooling rates in the range 0.1 °C s⁻¹ to 130 °C s⁻¹, austenite transforms to either polygonal ferrite (PF) or massive ferrite (MF). The transformation temperatures vary systematically with cooling rate and austenite condition. There is indirect evidence that the transformation rates for both PF and MF are decreased by the presence of substitutional solute atoms and precipitate particles. When unstable austenite is deformed at 850 °C, it transforms to an extremely fine strain-induced MF. Under conditions of high supersaturation of Ti, Nb, and S, (Ti,Nb) _x S _y precipitates form at 850 °C as coprecipitates on pre-existing (Ti,Nb)N particles and as discrete precipitates within PF grains. Pre-existing intragranular (Ti,Nb) _x S _y precipitates retard recrystallization and grain coarsening of PF deformed at 850 °C and result in a stable, recovered subgrain structure. The results are relevant to the design of TMP schedules for warm rolling of IF steels.     

Bainite formation in low carbon Cr-Ni steels

A low carbon Cr-Ni steel has been used to investigate the formation of upper bainite. Experimental results indicate that the start temperatures of the three morphologies of upper bainite in this steel,i.e., carbide-free bainite, bainite with carbide between and within ferrite laths, are about 600°, 500δ, and 425 °C, respectively; the habit plane of bainitic ferrite in this steel is close to (1 7 11)_α, which is 13.3 deg away from (0 ll)_α; and the orientation relationship between cementite and ferrite is consistent with Bagaryatskii’s. By means of the superelement approach, a thermodynamic treatment which applies to Fe-C alloys is extended into that suitable for low alloy steels, and calculation shows that the driving force for bainite formation at B_S temperatures is insufficient to compensate for shear strain energy. Formerly Graduate Student, Department of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai, People’s Republic of China.     

Isothermal martensite transformation in a 1.80 Wt Pct C steel

The linear products formed isothermally at 373 K in a 1.80 wt pct C steel (Ms = 346 K) were examined by means of transmission electron microscopy. They were first reported as “black line products” by Greninger and Troiano. The isothermal product was of a thin plate with about 0.5 μm width, and it contained {112}_b transformation twins and revealed a habit plane of {3 15 10}_f. The orientation relationship between austenite and product was close to the Nishiyama relationship. These crystallographic data were in good agreement with those calculated by the phenomenological theory of martensitic transformation. Consequently, the product was determined not to be lower bainite, but isothermal martensite. The black color of isothermal martensites resulted from the fact that they were easily etched by the precipitates of epsilon carbide formed during the isothermal holding.     

Bainite transformation temperatures in high-silicon steels

The bainite transformation temperatures of eight high-silicon steels were determined metallographically. The bainite start (B _s ) temperatures, which define the highest temperature at which bainite can form, all lay below the T ₀ loci, where ferrite and austenite of the same chemical compositions have identical free energy. The established method of calculating B _s temperatures gave reasonable agreement with the experimental results. Careful study of the isothermally reacted samples revealed that Widmanstätten ferrite and bainite could both be observed, even at the beginning of the transformation, at around the B _s temperature. On the other hand, the lower bainite start (LB _s ) temperatures of these steels were found to be very close to the martensite start (M _s ) temperatures. Silicon is considered to be responsible for depressing the LB _s temperature by retarding the formation of cementite. The coformation of upper and lower bainite near the LB _s temperature is also confirmed. The results indicate that the displacive formation mechanism of bainite is sustainable.     

Transformation Characteristics of Ferrite/Carbide Aggregate in Continuously Cooled, Low Carbon-Manganese Steels

Transformation characteristics and morphological features of ferrite/carbide aggregate (FCA) in low carbon-manganese steels have been investigated. Work shows that FCA has neither the lamellae structure of pearlite nor the lath structure of bainite and martensite. It consists of a fine dispersion of cementite particles in a smooth ferrite matrix. Carbide morphologies range from arrays of globular particles or short fibers to extended, branched, and densely interconnected fibers. Work demonstrates that FCA forms over similar cooling rate ranges to Widmanstätten ferrite. Rapid transformation of both phases occurs at temperatures between 798 K and 973 K (525 °C and 700 °C). FCA reaction is not simultaneous with Widmanstätten ferrite but occurs at temperatures intermediate between Widmanstätten ferrite and bainite. Austenite carbon content calculations verify that cementite precipitation is thermodynamically possible at FCA reaction temperatures without bainite formation. The pattern of precipitation is confirmed to be discontinuous. CCT diagrams have been constructed that incorporate FCA. At low steel manganese content, Widmanstätten ferrite and bainite bay sizes are significantly reduced so that large amounts of FCA are formed over a wide range of cooling rates.     

Second Stage of Upper Bainite in a 0.3 Mass Pct C Steel

Upper bainite forms in at least two stages, the formation of parallel plates of ferrite and the transformation of the interspaces to a mixture of cementite and ferrite. The first stage was examined in a preceding metallographic study of the formation of ferrite in hypoeutectoid steels and the second stage, which is initiated by the occurrence of cementite in the interspaces, is the subject of the present study. The alloy from the preceding study will also be used here. The band of austenite in the interspaces between parallel plates is generally transformed by a degenerate eutectoid transformation when this band is thin. When it is thicker, it will transform by a more cooperative growth mechanism and result in a eutectoid colony, often with cementite platelets. A series of sketches are presented which illustrate in detail how the second stage of upper bainite progresses according to the present observations. The cooperative manner did not increase as the temperature was lowered because the tendency to form plates of ferrite was still increasing at lower temperatures, making the interspaces too narrow for the cooperative reaction to dominate over the formation of fine plates of ferrite.

     

Mechanism of Secondary Hardening in Rapid Tempering of Dual-Phase Steel

Dual-phase steel with ferrite-martensite-bainite microstructure exhibited secondary hardening in the subcritical heat affected zone during fiber laser welding. Rapid isothermal tempering conducted in a Gleeble simulator also indicated occurrence of secondary hardening at 773 K (500 °C), as confirmed by plotting the tempered hardness against the Holloman–Jaffe parameter. Isothermally tempered specimens were characterized by analytic transmission electron microscopy and high-angle annular dark-field imaging. The cementite (Fe₃C) and TiC located in the bainite phase of DP steel decomposed upon rapid tempering to form needle-shaped Mo₂C (aspect ratio ranging from 10 to 25) and plate-shaped M₄C₃ carbides giving rise to secondary hardening. Precipitation of these thermodynamically stable and coherent carbides promoted the hardening phenomenon. However, complex carbides were only seen in the tempered bainite and were not detected in the tempered martensite. The martensite phase decomposed into ferrite and spherical Fe₃C, and interlath-retained austenite decomposed into ferrite and elongated carbide.