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.     

Crack growth from internal hydrogen—temperature and microstructural effects in 4340 steel

Internal hydrogen effects on stage II crack growth rates in AISI 4340 steel have been studied as a function of test temperature. A model is developed that is physically based in that classical thermodynamics relates to solubility and trapping and Fick’s second law controls hydrogen transport. Both of these are microstructurally related to how trapping affects both the crack initiation site and diffusion to it. For two tempered conditions of 4340 steel, it is shown that there is a test temperature,T ₀, for stage II crack growth, above which the crack does not grow. The fractography associated with test temperatures approachingT ₀ tends toward 100 pct intergranular for both 1340 MPa and 1620 MPa strength levels. At lower test temperatures, there is as much as 50 pct microvoid coalescence or 30 pct quasi-cleavage. In the lower strength condition, hydrogen traps at oxysulfide particles with a binding energy near 75 kJ/mol. Where these intersect the prior austenite grain boundaries, this promotes fingers of intergranular fracture which later triggers tearing of 100 μm size ligaments by microvoid coalescence. For the higher strength material, it is proposed that hydrogen traps along martensite lath intersections with prior austenite grain boundaries, the binding energy being near 27 kJ/mol. This promotes 1 μm size striations along intergranular facets. In both cases the fractography is consistent with a proposed model of stress field concentration of hydrogen, further concentration along trap sites, fracture nucleation at trap sites, and local, discontinuous fracture instabilities.     

Effect of microstructure on the temper embrittlement of Cr-Mo-V steels

The effect of transformation product on the temper embrittlement susceptibility of a Cr-Mo-V steel doped with P and Sn has been investigated at different strength levels. Results show that at low strength levels (< 10R _C ) embrittlement susceptibilities of tempered bainite and ferrite-pearlite structures are comparable to each other, but lower relative to tempered martensite. The lower susceptibility of tempered bainite relative to tempered martensite obtains up to about 40R _C , above which the susceptibilities for the two structures are similar. Variation of embrittlement susceptibility with microstructure is completely consistent with the degree of grain boundary segregation of P and Sn. The segregation is smaller in bainite than in martensite at a given strength level and increases with increasing strength level for a given structure. When compared at the same degree of embrittlement (i.e., same shift in FATT), the amount of grain boundary segregate and the extent of intergranular fracture are lower in tempered bainite compared to martensite indicating that embrittlement of interfaces other than prior austenite boundaries might be an important factor in the embrittlement of bainite structures.     

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).     

The microstructure and fracture of a carburized steel

The case microstructure and fracture of a coarse-grained 8620 steel carburized to 1 pet surface carbon are quite sensitive to austenitizing conditions. Reheating martensitic speci-mens below theA _cm produces in the case a refined austenitic grain size, a very fine mar-tensite, spherical carbide particles and a minimum of retained austenite and microcrack-ing. Overload fracture through the latter microstructure is transgranular and scanning electron microscopy shows both microvoid coalescence around thecarbide particles and an apparent fine cleavage in other areas. As-carburized specimens and specimens re-austenitized above theA _cm developed a case microstructure characterized by a coarse austenitic grain structure in which plate martensite with microcracks developed on cool-ing within a large amount of retained austenite. The overload fracture through this mi-crostructure followed a predominately intergranular path and effectively by-passed the retained austenite and microcracked martensite. Auger electron analysis showed that C and P were present on the intergranular fracture surfaces at concentrations above bulk, an observation consistent with literature reports of P segregation during austenitizing. This paper is based on a presentation made at a symposium on “Carburizing and Nitriding: Fundamentals, Processes and Properties” held at the Cincinnati Meeting of The Metallurgical Society of AIME, November 11 and 12, 1975 under the sponsorship of the Heat Treatment Committee.     

The influence of Mo and Co on the embrittlement of an Fe-Ni-Mn alloy

In the “as rolled” condition an Fe-6 Ni-5 Mn maraging type alloy was found to be brittle exhibiting intergranular fractures. The addition of 2.5 pct Mo and 5.0 pct Mo increased the impact toughness of the “as rolled” material and changed the mode of brittle fracture to transgranular cleavage. The addition of 9 pct Co embrittled the alloy. On aging Mo and Co raised the peak hardness of the base Fe-6 Ni-5 Mn alloy, however, aging led to rapid embrittlement. The base alloy and an alloy containing 2.5 pct Mo showed brittle intergranular fractures on aging. The addition of 5 pct Mo gave rise to brittle transgranular cleavage fractures on aging at 450°C, but at temperatures less than 450°C there was always up to 20 pct intergranular fracture present in brittle fractures. At temperatures greater than 475°C brittle intergranular failure occurred in the 5 pct Mo alloy due to a grain boundary film of M₆C and Fe₂Mo. This paper is based upon a thesis submitted by D. R. Squires in partial fulfilment for a higher degree of CNAA at Sheffield Polytechnic.     

The relationship between toughness and microstructure in Fe-high Mn binary alloys

The influence of Mn content on the ductile-brittle transition in 16 to 36 wt pct Mn steels was investigated and interpreted in light of the evolving microstructure. It was found that when hcp ε martensite is present in the as-quenched condition or forms during deformation, it lowers the toughness. In 25Mn steel, the stress concentrations at e plate intersections result in the formation of planar void sheets along the {111}_γ planes. The deformation-induced α’ martensite in 16 to 20 pct Mn alloys enhances the toughness, but leads to a ductile-to-brittle transition at low temperatures that is due to the intrusion of an intergranular fracture mode. Binary alloys with greater than 31 pct Mn also fracture in an intergranular mode at 77 K although the impact energy remains quite high. Auger spectroscopy of the fracture surfaces shows no evidence of significant impurity segregation, which suggests the importance of slip heterogeneity in controlling intergranular fracture in these alloys.     

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.     

The tensile properties of an Fe-4 pct Mo-0.2 pct C martensite tempered under applied stress

In Fe-4 pct Mo-0.2 pct C martensite which is a typical secondary hardening steel, premature failure o°Curred in tensile test at 600 °C to 700°C where solute atoms could diffuse easily. To clarify this phenomenon, the quenched specimens were tempered under applied stress and tensile-tested at room temperature. The following results were obtained: (1) Typical intergranular fracture was observed in specimens tempered in a temperature range of 600 °C to 650 °C with tempering times of five minutes to 10 minutes and applied stress (70 MPa to 140 MPa). (2) Based on Auger analysis, this phenomenon was considered to be caused by segregation of P, S, and Mo on prior austenite grain boundaries due to applied stress. (3) The direction of applied stress was found to be very significant. Namely, when the tensile direction was parallel to the applied stress during tempering, the specimen was more brittle, and when tensile direction was normal to the applied stress, the specimen was not so brittle. (4) To reduce this embrittlement, solution treatment temperature was adjusted, and it was found that the embrittlement was considerably reduced both in specimens with fine prior austenite grains and with some ferrite phase on prior austenite grain boundaries. TAKATOSHI OGAWA, formerly with Kyoto University. YOSHIFUMI OHMURA, formerly with Kyoto University. 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.     

Improving the weldability and service performance of nickel-and iron-based superalloys by grain boundary engineering

The principal limitation of today’s Ni- and Fe-based superalloys continues to be their susceptibility to intergranular degradation arising from creep, hot corrosion, and fatigue. Many precipitation-strengthened superalloys are also difficult to weld, owing to the formation of heat-affected zone (HAZ) cracks during postweld heat treatments (PWHTs). The present work highlights significant improvements in high-temperature intergranular degradation susceptibility and weldability arising from increasing the relative proportion of crystallographically “special” low-Σ CSL grain boundaries in the microstructure. Susceptibility to intergranular degradation phenomena is reduced by between 30 and 90 pct and is accompanied by decreases in the extent and length of PWHT cracking of up to 50-fold, with virtually no compromise in mechanical (tensile) properties upon which the functionality of these specialty materials depends. Collectively, the data presented suggest that “engineering” the crystallographic structure of grain boundaries offers the possibility to extend superalloy lifetimes and reliability, while minimizing the need for specialized welding techniques which can negatively impact manufacturing costs and throughput.     

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.     

Granular product in a high strength,low alloy containing molybdenum and niobium

The formation and microstructure of the granular product and its effect on the mechanical properties of a high-strength, low alloy steel containing molybdenum and niobium have been investigated. It was found that the granular product “islands” are composed of both twinned martensite and dislocated martensite. The effect of the granular “islands” on the strength at room temperature and at 400 °C has been determined. The results showed that the strength increased and both the impact and fracture toughness decreased as the volume fraction of granular “islands” was increased.In situ fracture studies indicated that the three stages of the microfracture process of the specimen containing granular “islands” are the initiation of voids at interfaces between the granular “islands” and the bainitic ferrite matrix, followed by void growth and finally, coalescence by shear.     

The Tensile Properties of an Fe-4 Pct Mo-0.2 Pct C Martensite Tempered under Applied Stress

In Fe-4 pct Mo-0.2 pct C martensite which is a typical secondary hardening steel, premature failure occurred in tensile test at 600 °C to 700 °C where solute atoms could diffuse easily. To clarify this phenomenon, the quenched specimens were tempered under applied stress and tensile-tested at room temperature. The following results were obtained: (1) Typical intergranular fracture was observed in specimens tempered in a temperature range of 600 °C to 650 °C with tempering times of five minutes to 10 minutes and applied stress (70 MPa to 140 MPa). (2) Based on Auger analysis, this phenomenon was considered to be caused by segregation of P, S, and Mo on prior austenite grain boundaries due to applied stress. (3) The direction of applied stress was found to be very significant. Namely, when the tensile direction was parallel to the applied stress during tempering, the specimen was more brittle, and when tensile direction was normal to the applied stress, the specimen was not so brittle. (4) To reduce this embrittlement, solution treatment temperature was adjusted, and it was found that the embrittlement was considerably reduced both in specimens with fine prior austenite grains and with some ferrite phase on prior austenite grain boundaries. Formerly with Kyoto University Formerly with Kyoto University This paper is based on a presentation made at the “Peter 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.     

Plastic straining effects on the microstructure of a Ti-rich NiTi shape memory alloy

A Ni-52 at. pct Ti shape memory alloy, cold drawn to 30 pct, was annealed at 1173 K for 1 hour, water quenched, and then subjected to differential scanning calorimetry (DSC). No evidence of the premartensiticR transformation was found during either the forward or the reverse transformation. Microstructurally, it was found that the alloy possessed a relatively large volume fraction (∼0.05) of coarse second-phase brittle particles. These precipitates acted as preferential sites for martensite plate nucleation and gave rise to a “starlike” morphology. The tensile and compressive properties of the alloy in the as-received condition were also investigated. The alloy exhibited relatively good ductility (fracture strain=0.28), which was attributed to its inherent ability to relieve or delay the development of plastic instabilities through rapid strain hardening. In addition, X-ray diffraction (XRD) of deformed specimens indicated the presence of an extraintensity peak corresponding to the B2 phase (110)_B2 when the alloy was plastically deformed in compression. Accordingly, it is suggested that plastic deformation induces the reverse transformation to the B2 phase in highly stressed local regions. Transmission electron microscopy (TEM) of deformed martensite structures showed slip lines probably due to dislocation slip, as well as variant interpenetration. Besides, optical and scanning microscopy of regions adjacent to the fractured surfaces indicated that fine martensite plates and/or “apparent” new grains develop at regions of prior stress intensification (former crack-tip regions) during crack propagation.     

Plastic straining effects on the microstructure of a Ti-rich NiTi shape memory alloy

A Ni-52 at. pct Ti shape memory alloy, cold drawn to 30 pct, was annealed at 1173 K for 1 hour, water quenched, and then subjected to differential scanning calorimetry (DSC). No evidence of the premartensitic R transformation was found during either the forward or the reverse transformation. Microstructurally, it was found that the alloy possessed a relatively large volume fraction (∼0.05) of coarse second-phase brittle particles. These precipitates acted as preferential sites for martensite plate nucleation and gave rise to a “starlike” morphology. The tensile and compressive properties of the alloy in the as-received condition were also investigated. The alloy exhibited relatively good ductility (fracture strain = 0.28), which was attributed to its inherent ability to relieve or delay the development of plastic instabilities through rapid strain hardening. In addition, X-ray diffraction (XRD) of deformed specimens indicated the presence of an extraintensity peak corresponding to the B2 phase (110)_B2 when the alloy was plastically deformed in compression. Accordingly, it is suggested that plastic deformation induces the reverse transformation to the B2 phase in highly stressed local regions. Transmission electron microscopy (TEM) of deformed martensite structures showed slip lines probably due to dislocation slip, as well as variant interpenetration. Besides, optical and scanning microscopy of regions adjacent to the fractured surfaces indicated that fine martensite plates and/or “apparent” new grains develop at regions of prior stress intensification (former crack-tip regions) during crack propagation.     

The influence of the lath morphology on the yield stress and transition temperature of martensitic- bainitic steels

A decrease in the packet size or lath width of bainitic-martensitic steels produces a simultaneous increase of toughness and yield stress. Specific interpretations are necessary to give the actual relationships since classical theories of the grain size effect cannot be directly applied to both of these microstructural parameters. To obtain these specific interpretations, a detailed analysis of lath orientations inside a packet is necessary. This analysis reveals that assumptions such as all laths in any one packet are similarly oriented are unfounded and that in fact a packet contains many high angle lath boundaries which are given by laths adopting different Kurjumov-Sachs orientation variants during the γ → α transformation. The yield stress then depends on the average lath “diameter” which is a function of lath width and length, the latter dimension being related to the packet diameter. A Petch agreement is not found, rather the yield stress is found to be related to the reciprocal of the average lath diameter. A theoretical analysis shows that for very fine grain sizes, as encountered in bainites and martensites, macroscopically heterogeneous deformation (a necessary condition leading to the Petch formulation) tends to vanish and that for macroscopically homogeneous deformation the yield stress is expected to be related to the reciprocal of the grain diameter. The fracture transition temperature is determined by the particular fracture characteristics of these steels. It is possible to explain the fracture of bainite and lath martensite without recourse to concepts such as “effective grain size” or “covariant packet size”. In spite of the different lath orientations inside a packet, a brittle crack may adopt an average, approximately straight fracture direction across a packet by following a particular group of different fracture planes that are separated by low angle boundaries. At a packet boundary the crack must find another group of fracture planes, which will impose an important deviation of the crack. At the transition temperature, the controlling event in the fracture sequence is this crack deviation, which imposes an energy requirement for the crack to undergo high angle deviations across the first laths that are adjacent to the packet boundary, until the new average fracture direction is found. Using this model the transition temperature can be related to a logarithmic function of the product of the packet diameter and lath width. Formerly at the Centre des Matériaux de l’Ecole Nationale Supérieure des Mines de Paris.     

The shape memory effect and pseudoelasticity in polycrystalline Cu-Zn alloys

The shape memory effect associated with the reverse transformation of deformed martensite, pseudoelastic behavior involved in stress-induced martensite formation and the reversion of strained martensite after an applied stress is relaxed aboveA _f have been studied. Grain size and specimen geometry effects have been related to the above phenomena. Although recoverable strains as high as 10.85 pct were observed in coarse-grained (“bamboo” type) specimens, the shape memory effect is restricted in fine-grained specimens because of permanent grain boundary deformation and intergranular fracture which occurs at relatively low strains. A fine grain size also acts to suppress pseudoelastic behavior because permanent, localized deformation is generated concurrent with the formation of stress-induced martensite which inhibits reversion of the latter upon release of stress. The apparent plastic deformation of martensite belowM _f can be restored by transforming back to the original parent phase by heating toA _f (shape memory) or alternatively, can be recovered belowM _f by applying a small stress of opposite sign. Martensite deformed belowM _f with the same stress maintained while heating persists aboveA _f, but reverts to the parent phase in a pseudoelastic manner when the stress is relieved. The athermal thermoelastic martensite, which forms in groups composed of four martensite plate variants, undergoes several morphology changes under deformation. One of the variants within a plate group cluster may grow with respect to the others, and eventually form a single crystalline martensitic region. At a later stage pink colored deformation bands form in the same area and join up with increasing stress, resulting in thermally irreversible kinks. The clusters of plate groups may expand like grain growth or contract as a whole during deformation, or act as immobile “subgrains” which lead to permanent deformation at their boundaries. Stress-induced martensite usually forms as one variant of parallel plates which join up with increasing stress to form single crystalline regions. Further stress leads to pink colored deformation bands, similar to those in the deformed athermal martensite. Other similarities and differences between the stress-induced and athermal martensite have been investigated and are discussed. Formerly with the University of Illinois at Urbana-Champaign     

Effects of intercritical treatment and tempering on fracture behavior in a medium-carbon 2Si-3Ni steel

In the 2Si-3Ni steel intercritically treated in the range of 720 ‡C to 790 ‡C, the fracture behavior under the impact testing has been analyzed and the post-tempering effect has also been investigated. The transgranular fracture occurred in the specimens treated below and at 730 ‡C (SN73 specimen) in relatively low intercritical temperature range, but the intergranular fracture occurred in the specimens treated at 750 ‡C and 770 ‡C (SN75 and SN77 specimens) in relatively high intercritical temperature range. In the SN73 specimen, there was little coarse martensite at the prior austenite grain boundaries, whereas there was continuous, coarse martensite at those boundaries in the SN75 and SN77 specimens. The fracture behavior was mainly discussed in terms of the microstructural differences. In addition, no or a little increase in impact toughness, in spite of great decreases in hardness, in the SN75 and SN77 specimens tempered at 600 ‡C is correlated with the easy occurrence of intergranular fracture, which is caused by the carbide aggregates formed in the continuous, coarse martensite at the grain boundaries.     

Effects of microstructure on the fracture toughness of Ti3Al-based titanium aluminides

The influence of microstructure on the fracture toughness of Ti-23A1-9Nb-2Mo-1Zr-1.2Si (at. pct) and Ti-23A1-11Nb-0.9Si (at. pct) Ti₃Al-based alloys has been investigated. Basket-weave microstructures comprising different volume fractions of α ₂ and retained β phases were produced by systematic heat treatments. Besides the volume fraction of the retained β phase, the average size of the β laths has also been used to characterize these microstructures. The toughness of both alloys was examined at room temperature, and the brittle transgranular fracture modes were found to be controlled by microstructure. However, the toughness is not determined solely by the volume fraction of the retained β phase, and a linear relationship has been obtained between the fracture toughness and the average size of the retained β laths. It appears therefore that the toughness of Ti₃Al-based alloys at room temperature is controlled primarily by the width of retained β laths rather than by the retained β volume fraction.     

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.