The main aim of this work is to study the mechanisms that control the austenitisation process in steels with different initial microstructures. The compiled knowledge in literature regarding the isothermal formation of austenite from different initial microstructures (pure and mixed microstructures), has been used in this work to develop a model for non-isothermal austeniteformation in steels with initial microstructure consisting of ferrite and/or pearlite. The microstructural parameters that affect the nucleation and growth kinetics of austenite, and the influence of the heating rate have been considered in the modelling. Moreover, since dilatometric analysis is a technique very often employed to study phase transformations in steels, a second model to describe the dilatometric behaviour of the steel and calculate the relative change in length which occurs during the austeniteformation has been developed. Both kinetics and dilatometric models have been validated. Experimental kinetic transformation, critical temperatures as well as the magnitude of the overall contraction due to austeniteformation are in good agreement with calculations.
temperature is almost insensitive to heating rate over the range studied. Experimental results in steels with a pearlite and ferrite-pearlite initial microstructures also show that the elevation of the critical temperatures with heating rate is quite sensitive to the morphology of pearlite. It seems that the higher the heating rate is, the stronger the influence of morphology on the critical temperatures are. This experimental study and the knowledge of the mechanisms that control the austeniteformation process have allowed to establish the variables that most directly influence this reaction in steels with pearlite and ferrite-pearlite initial microstructures. Those are the heating rate and the two parameters that characterise the morphology of pearlite, the mean true interlamellar spacing and the edge length of the pearlite colonies interface in pearlitic steels, together with the volume fraction of pearlite and the mean free distance of pearlite in ferrite plus pearlite initial microstructures. Likewise, two equations have been proposed for the determination of the start (Ac 1 ) and (Ac 3 ) finish temperatures of austeniteformation as a
The microstructure from which austenite may form can be infinitely varied (ferrite, martensite, carbide and various morphologies and aggregates of each, e.g. pearlite, bainite). Many variables are therefore needed to describe the kinetics of austeniteformation. Factors such as particle size, the distribution and chemistry of individual phases, homogeneity, the presence of non-metallic inclusions, should all be important. [2-5] Thus, in the case of formation of austenite from pearlite one of the most relevant structural factor to be considered is the interlamellar spacing of pearlite. 
Samajdar et al. (5) have characterized the transformations during intercritical annealing in a Fe- 0.11C-1.53Mn-1.5Si steel, and they reported that the initial formation of austenite was observed inside pearlite colonies. They claimed that the transformation may be divided in two stages: spheroidization of cementite and formation of austenite from the spheroidized cementite particles. Below 1 min annealing, they observed the austenite structure to be very fine. When the intercritical soaking time increased, a general coarsening of the austenite particles was observed while the volume fraction of the austenitic phase leveled off. Also, the Ms temperature determined by dilatometry was found to increase with longer soaking time. Even after a few seconds annealing, about 5% of retained austenite could be detected, and the measured carbon content of the retained austenite steadily dropped with longer annealing times. The rate of austeniteformation was much slower than that of ferrite recrystallization, the completion of the latter being observed within about 5 seconds at 750°C. Pure austenite grain coarsening was observed for prolonged holding, beyond 4 minutes at 750°C. Austenite, in general, was less stable at long intercritical annealing times. This was in part interpreted by the influence of grain size on the stability of austenite. For a given chemical composition, larger austenite grains are expected to be less stable. Also, the high stability of intercritical austenite in the very beginning of the soaking treatment was a direct consequence of its formation process from spheroidized cementite, i.e., high C austenite. By this, the new austenite particles resulted with a high carbon content leading to a considerable stability.
Background: High strength low carbon steels used in the automobile industry are often annealed at the intercritical temperature, and recrystallization of ferrite and formation of austenite can occur at the same time during the intercritical annealing. However, detailed studies on the competition between recrystallization of ferrite and formation of austenite at the early stage of annealing have not been carried. Therefore, ferrite recrystallization and austeniteformation in cold-rolled low carbon steels at an early stage of annealing have been observed. Methods: The chemical composition of the steel (in mass %) was 0.1C-2.0Mn. After hot and cold rolling, the specimens were heated from room temperature to 1003, 1013 and 1023 K at a rate of 30 K s -1 , and then gas-quenched to room temperature at a rate of 50 K s -1 . We observed the microstructure of the annealed specimens.
In order to form a sufficiently uniform refined struc- ture, the small austenite grains inherent in DRx should cover a high enough volume fraction. However, even at T > 1050 °C, this is possible only if single strains exceed 0.9 %, which is unrealistic in industrial rolling. There- fore, austeniteformation during DRx remains incom- plete, and hence, an unwanted non-uniform structure forms (coexistence of fine and coarse grains) that re- duces the ductility of the steel. As evident in Fig. 1, this characteristic previously known with the traditional steel F620 alloyed with V (Zisman et al. 2012) remains valid for its modification alloyed with Nb. Therefore, the growth of threshold εp due to Nb proves to have an add- itional positive effect because it more reliably excludes in- complete DRx and related grain size variation. The grain size was measured according to ISO 643: 2003 standard.
Abstract: Most high-hardness tool steels comprising forming dies require expensive finish machining operations to compensate for the dimensional distortion and surface oxidation caused by the die heat treatment. Precipitation-hardening (PH) tool steels allow for soft finish machining followed by an aging treatment without major deformation or oxidation in the die, but exhibit poor wear performance owing to the lack of carbides in their structure. This drawback can be overcome by combining laser cladding technology, austenite retention, and cryogenic treatments. Hence, an alternative die manufacturing route based on laser cladding was explored. The forming surface of a modified chemistry tool steel die was subjected to cladding. The martensite finish (M f ) temperature of the steel was tuned to enhance austenite retention at room temperature. The cladded surface was then machined in a reduced-hardness condition resulting from retained austeniteformation. Subsequent deep cryogenic treatment of the die favoured the retained-austenite-to-martensite transformation, thereby increasing the die hardness without major distortion or oxidation. This process combined the advantages of high-carbide-bearing tool steels and PH steels, allowing for a die with hardness exceeding 58 HRC to be finish machined at <52 HRC. Controlling the occurrence of retained austenite represents an effective strategy for achieving new manufacturing scenarios.
In the case that the sample is austenitized subsequent to rapid cooling, austenite will transform from a bainite and/or martensite structure. Heating to the austenite temperature will allow the diffusion of carbon from the supersaturated martensite. The carbon bonds with iron forming cementite precipitates. At the austenitizing temperature austenite grain nucleation will primarily take place along the carbide and PAGB’s. Austenitizing steels from bainite or tempered martensite produces acicular and globular austenite grain morphologies. Acicular austeniteformation is a product of an acicular starting microstructure that is heated at a low austenitizing temperature with slow heating rates. Acicular nuclei also quickly develop to a globular morphology when heated above the upper critical limit. Initial globular austenite grain formation is amplified by utilizing elevated austenitizing
It is well known that hot rolled microalloyed steels can be strengthened by a combination of grain reﬁnement and precipitation strengthening, using microalloying elements niobium (Nb), titanium (Ti) and vanadium (V), individually or in combination. 10 12) Previous experiments prove that the addition of lanthanum (La) suppresses the precipitation of NbC in austenite, and the ﬁrst-principle calculations indicate that the solubility of niobium and carbon in fcc Fe are increased, and the chemical potential of both are decreased in the presence of La. 13) For the case of NbC precipitation in ferrite steels, our di ﬀ usion couple experiment reveals that Nb di ﬀ uses slightly faster in the presence of La, and consequently leads to the faster precipitation kinetics of NbC. 14)
The validity and the efficiency of techniques to reveal austenite grain boundaries in steels are uncertain, because they depend on the chemical composition, heat treatment, and other not well-identified factors. Therefore, the revealing of austenite grain boundaries could be a difficult task, especially in medium-carbon microalloyed steels, which showed low sensibility to chemical etching [1,2]. Recent works in medium-carbon microalloyed steels [3,4] demonstrated that procedures based on the combination of heat treatment and chemical etching are unable to reveal the austenite grain boundaries for certain austenitisation conditions in a given steel and for a particular condition in steels with similar chemical compositions without an apparent reason. However, the thermal etching method gave excellent results in all the tested steels at every austenitisation condition, even at those in which no other method was effective for revealing the austenite grain boundaries. The thermal etching was the most adequate method for revealing the prior austenite grain boundary of this kind of steels.
The aim of this project was to combine the use of cyclic partial phase transformations (CPPT) with the direct observation of individual interfaces using in situ hot stage transmission electron microscopy (TEM). Although hot stage TEM has been used previously to examine phase transformations in steels, the use of hot stage TEM to investigate interface behaviour during CPPT experiments has not been reported in the literature before. The primary purpose of this combination was to determine how the interface behaved during the stagnant stage and in the period immediately before and after it. This includes understanding whether the interface migrated sluggishly or remained completely static as well as any effect on the nature of the interface itself. More general insight could be gained using this technique into the behaviour of the interface during the austenite to ferrite and ferrite to austenite transformation, including how the interface interacts with microstructural features.
Samples for metallographic observation were cut on the rolling direction (RD)– normal direction (ND) plane and prepared carefully. A 2% nital solution was used to show the transformed microstructure and a saturated aqueous picric acid solution was used to reveal the prior-austenite grain boundaries (PAGBs). PAGSs were measured optically by the linear intercept method. Optical microscopy (OM) and scanning electron microscopy (SEM) observations were carried out on Nikon Eclipse LV150 and FEI InspectF, respectively.
It can be seen from Figure 2.18 that during the heating process, the surface is kept in compression until it reaches 1000°C. At this temperature, the surface becomes plastic. During the cooling process, the work piece has tensile stress until the surface comes to the martensite formation stage. The level of stresses depends on the thermal coefficient of expansion and elastic modulus of the material. Thus, the amount of heat transfer and rate of heating determines the thermal stresses. Residual stresses on heating will have very little effect except when tensile stresses are high on the inside and the outside compressive stresses are relaxed through subsequent plastic deformation. In this case, the inside remaining tensile stress are large enough to cause the part to crack. This means that induction heating could cause cracks when heating cold-drawn bar stock, when the bar is not stress relieved uniformly, and there are still large tensile stresses inside the core.
transform immediately after the initial maximum n and transformation proceeds until the instability criterion is reached, while in alloy 2 it starts to transform later, since the austenite is mechanically more stable. It should be noted that although both alloys have the same austenite fraction, the total elongation of alloy 2 is 10% smaller than that of alloy 1. It is suggested that this is due to the fact that the chemical composition of retained austenite in alloy 2 makes it mechanically too stable. High fractions of very stable austenite are present at 300 C microstructures for both
Super duplex stainless steel (SDSS) is considered as a composite formed from a microstructure of an approximately equal mixture of two primary constituents (γ-austenite and α-ferrite phases) and the secondary precipitates (sigma, chi, alpha-prime, etc.). While the formation of these phases affects the properties of SDSS, however there are no rules that govern the relationship. In this work, the relationship between toughness as well as corrosion behavior of SDSS (UNS 32760) and the microstructure constituents has been experimentally investigated, and analyzed in view of the composite principles. Another two stainless steels namely; fully austenitic SASS (UNS N08367) and fully ferritic FSS (UNS S42900) are considered to simulate the constituent’s primary components in the composite which are austenite γ and ferrite α phases respectively. Samples of the composite and constituent’s steels are first subjected to solution annealing, where the composite steel has a microstructure of γ austenite and α ferrite grains. They were then subjected to similar different isothermal heat treatment cycles, for the formation of secondary phase precipitations within the transformation temperature ranges of each of γ and α primary grains. Impact toughness and cor- rosion (specific weight loss) tests were conducted on the annealed and isothermally treated sam- ples. The composite rule of the mixtures (ROM) is used to analyze the relationship between the toughness and corrosion properties in the composite SDSS and the SASS and FSS constituent’s steels. The analysis indicates that in case of toughness, ROM applies well on the composite and constituents’ steels in the solution annealed and in isothermal treatment conditions, where better matching between experimental and calculated results is observed. When applying ROM for cor- rosion weight loss, a great difference is found between the experimental and calculated results, which is much reduced for solution treated samples ferritic and austenitic temperature ranges of 480˚C - 500˚C and 700˚C - 750˚C as for ferrite and austenite respectively.
Previous works on the martensitic steels showed that the carbon from the supersaturated martensite can migrate to the untransformed austenite during partitioning stage 8-11 . The carbon enrich retained austenite becomes more stable as a result of this relatively short process (usually ≤0.5 hrs) 12 . Therefore, it appears that for the Q&P treatment, the microstructure can be refined and the transformation time can be considerably shortened, compared to the low temperature micro/nano-structured bainitic transformation. In the present work, conventional micro/nano-structured bainitic transformation combined with partitioning, and quenching process was conducted in order to develop a micro/nano-structured bainitic steel with even finer and more stable retained austenite grains. The heat treatment sequence consisted of a low temperature bainitic transformation followed by a partitioning process at an intermediate temperature, and finally water quenching to room temperature (B+P+Q). The effect of such multi-stage heat treatment on the microstructure of a micro/nano-structured bainitic steel was investigated by using optical, scanning and transmission electron microscopy methods.
OR with each martensite in the prior austenite grain. Hence, the variant selection mechanism is suggested, 15) where a single variant during the reverse transformation is selected according to the relaxation of the local residual stresses induced by martensitic transformation. 16) As no residual stress existed in the ferrite matrix, the austenite could be precipitated by satisfying the K S OR with the ferrite matrix by reducing the interfacial energy between ferrite / austenite. The twenty-four variants of reverted austenite satisfying the KS OR with the ferrite matrix revealed the same interfacial energy between them. Therefore, the retained austenite in the ferrite matrix was chosen to be any variant such that austenite grains with diﬀerent crystal orientations were dispersed in the ferrite grain.