Fe-C Phase Diagram Research Articles

Efficient estimation of material property curves and surfaces via active learning

The relationship between material properties and independent variables such as temperature, external field or time, is usually represented by a curve or surface in a multi-dimensional space. Determining such a curve or surface requires a series of experiments or calculations which are often time and cost consuming. A general strategy uses an appropriate utility function to sample the space to recommend the next optimal experiment or calculation within an active learning loop. However, knowing what the optimal sampling strategy to use to minimize the number of experiments is an outstanding problem. We compare a number of strategies based on directed exploration on several materials problems of varying complexity using a Kriging based model. These include one dimensional curves such as the fatigue life curve for 304L stainless steel and the Liquidus line of the Fe-C phase diagram, surfaces such as the Hartmann 3 function in 3D space and the fitted intermolecular potential for Ar-SH, and a four dimensional data set of experimental measurements for BaTiO3 based ceramics. We also consider the effects of experimental noise on the Hartmann 3 function. We find that directed exploration guided by maximum variance provides better performance overall, converging faster across several data sets. However, for certain problems, the trade-off methods incorporating exploitation can perform at least as well, if not better than maximum variance. Thus, we discuss how the choice of the utility function depends on the distribution of the data, the model performance and uncertainties, additive noise as well as the budget.

Effect of Solute Elements and Cooling Rate on Strain in Brittle Temperature Range of Continuously Cast Strand

A micro-segregation model of solute elements in mushy zone with δ/γ transformation during solidification was established based on the regular hexagon transverse cross section of dendrite shape proposed by finite difference method under the non-equilibrium solidification condition. The model was used to calculate the non-equilibrium pseudo binary Fe-C phase diagram and the strain of steels induced by variation of temperature in brittle temperature range. On the basis of the phase diagram and the strain, the strain curve in brittle temperature range as a function of carbon content for continuously cast strand was introduced and obtained. Solute elements change the position of the strain curve. And cooling rate changes the position and the shape of the strain curve. The comprehensive formula of the strain as functions of solute elements and cooling rate in brittle temperature range has been obtained by nonlinear fitting program.

Theoretic analysis on the mechanism of particular pattern formed on the ancient Damascus steel blades

What causes the specific pattern on the ancient Damascus steel blades has been a myth for centuries and is still a hot topic for the today’s metallurgist. Although it has been studied and argued for centuries, the disputes on this topic still remain among the present metallurgists, which are summarized in this paper. Besides, histories of Wootz ingots, Damascus blade and Damascus pattern have been also reviewed respectively in details. With the help of modern metallurgy knowledge, we are now able to quantify all the possibly relevant factors by thermodynamic and kinetic calculations, such as influence of S, P and V elements on the Fe-C phase diagram, the dendritic microsegregation of S and P during a slow solidification, the temperature for precipitation of V(CN) in austenite and the diffusion kinetics of V during high-temperature homogenization. All of those factors have been considered as the possible perquisites for the formation of Damascus pattern. The thermodynamic and kinetics simulations indicate that the sulfide and phosphide, which precipitates at the segregated zone due to the slow solidification, may not be completely dissolved at the high-temperature solution treatment and then might act as the nuclei for the subsequent precipitation of cementite particles during hot hammering. In particular, the P-rich liquid, which is formed due to the segregation during solidification, may transform to eutectic product of phosphide and cementite. This is most likely responsible for the spatial distribution of cementite in Damascus blades and the formation of Damascus pattern.

Formation and Structure of Graphene Waves on Fe(110)

A very rich Fe-C phase diagram makes the formation of graphene on iron surfaces a challenging task. Here we demonstrate that the growth of graphene on epitaxial iron films can be realized by chemical vapor deposition at relatively low temperatures, and that the formation of carbides can be avoided in excess of the carbon-containing precursors. The resulting graphene monolayer creates a novel periodically corrugated pattern on Fe(110). Using low-energy electron microscopy and scanning tunneling microscopy, we show that it is modulated in one dimension forming long waves with a period of ∼4 nm parallel to the [001] direction of the substrate, with an additional height modulation along the wave crests. The observed topography of the graphene/Fe superstructure is well reproduced by density functional theory calculations, and found to result from a unique combination of the lattice mismatch and strong interfacial interaction, as probed by core-level photoemission and x-ray absorption spectroscopy.

Extended Tensile Ductility of a Formable High-Performance High-Manganese Steel

The aim of the contribution is to review the structure/property relationship of a high-performance FeMnN steel composition. In this work the influence of the Mn content ranging from 12 wt.-% to 24 wt.-% on the Fe-C phase diagram was calculated by thermodynamic simulations. The influence of coiling temperature on the microstructure of the FeMnN steel was studied by annealing treatments from 350 °C up to 750 °C. It showed the possible presence of cementite and martensite/pearlite after long annealing treatments at temperatures close to 500 °C. The steel was observed to exhibit outstanding combinations of tensile strength and ductility. The FeMnN steel exhibited at room temperature a remarkably high total elongation of about 100 % with a high tensile strength level of 1100 MPa. This product, proportional to toughness, achieved an outstanding value of about 110000 %MPa. The extended tensile ductility of the high-performance FeMnN steel was attributed to the nano-size twinning induced plasticity (NS-TWIP), which was observed by in-depth transmission electron microscopy study.

Time-Resolved X-Ray Diffraction Study on Solidification of Fe-B and Fe-C Eutectic Alloys

Solidification processes of Fe-B and Fe-C eutectic alloys have been investigated by a time-resolved synchrotron x-ray diffraction under containerless cooling conditions using a conical nozzle levitation technique. To observe relative variations of structure from the undercooled liquid to crystalline phase, we have conducted millisecond order time-resolved x-ray diffraction experiments with a two-dimensional detector. The structural variations observed during the solidification of the Fe83C17 alloy were identified as the phase transformation process expected from the Fe-C phase diagram. As for the Fe83B17 alloy, it was revealed that a metastable phase composed of Fe23B6 compound was precipitated as a primary crystalline phase from the undercooled liquid. In addition, decomposition of the metastable Fe23B6 phase showed dependence on the cooling rate of the sample. At the cooling rate of 30 K/s, the Fe23B6 phase decomposed to bcc-Fe and Fe2B phases with decreasing temperature. On the contrary, at the cooling rate of 180 K/s, the metastable Fe23B6 phase remained in spite of an appearance of the bcc-Fe phase. By comparing the primary crystalline phase between the Fe83C17 and the Fe83B17 alloys, we suggest that the formability of the metastable Cr23C6-type compound is closely related with the glass-forming ability of Fe-metalloid binary alloys.

Synthesis of large-area, few-layer graphene on iron foil by chemical vapor deposition

We demonstrate a simple and controllable way to synthesize large-area, few-layer graphene on iron substrates by an optimized chemical vapor deposition (CVD) method using a mixture of methane and hydrogen. Based on an analysis of the Fe-C phase diagram, a suitable procedure for the successful synthesis of graphene on Fe surfaces was designed. An appropriate temperature and cooling process were found to be very important in the synthesis of highly crystalline few-layer graphene. Graphene-based field-effect transistor (FET) devices were fabricated using the resulting few-layer graphene, and showed good quality with extracted mobilities of 300–1150 cm2/(V·s). Open image in new window

Foreword: Symposium on Austenite Formation and Decomposition IV

The technological importance of ferrous alloys rests to a large extent on their relatively low cost and their ability to achieve desirable strength-ductility combinations. The attainment of superior properties depends to a large degree on the phase transformations to and from austenite that occur during processing. From a scientific standpoint, ferrous alloys have been and continue to be benchmark materials from which solid-solid phase transformations understanding, broadly considered, was first formulated and developed. Since pioneering studies began in the 1880s, significant advances in this understanding have been made in view of developments such as the first Fe-C phase diagram (1900), x-ray diffraction (1920s), T-T-T diagrams (1930s), classical nucleation-and-growth theory (1940s), theories of martensite (1950s), and electron microscopy (1950s).

Application on Statistical Value of Phase Structure Formation Factor to Fe-C Phase Diagram

Based on the empirical electron theory in solids and molecules (EET), the statistical value of phase structure formation factor (S') was calculated by using the calculation methods of statistical value of alloying valence electron structure parameters. The effects of alloying elements upon austenitic zone of Fe-C phase diagram were discussed. The results show that: 1) Mn and Ni can increase the S' value of γ-Fe and austenitic phase, which increase their stability, leading to expand the γ-zone; 2) Cr, Mo and Si can increase the S' value of ferrite and cementite, which make them easier to precipitate, leading to reduce γ-zone. The research results agree well with real situations.

Analytical Modeling of Austenite Growth and Phase Evolution during Reverse Transformation from Pearlite in High Carbon Steels

Based on an analytical one-dimensional model, austenite growth into pearlite lamella and the corresponding phase evolution during isothermal reverse transformation to austenite at 1000-1183 K in Fe-C fully pearlitic steels containing 0.6-1.0 mass% C (in the austenite single phase field of Fe-C phase diagram) were simulated. It was found that the rate of austenite growth into ferrite increases faster with increasing reversion temperature than into cementite. Three types of phase evolution dependent on reversion temperature and carbon content were classified: 1) cementite rather than ferrite disappears first; 2) ferrite and cementite simultaneously disappear; 3) ferrite rather than cementite disappears first. The type of phase evolution in a hypoeutectoid steel heated above its Ae3 temperature possibly changes in the order of 1), 2) and 3) as the reversion temperature increases. For eutectoid and hypereutectoid steels, the phase evolution during isothermal reversion always obeys the type 3).

On the iron-carbon phase diagram

Experimental results that are obtained by electron microscopy, X-ray diffraction, and carbide analysis and indicate the precipitation of carbon atoms clusters in a hypereutectoid steel during its annealing above the eutectoid temperature are presented. These results are compared to the reported data in order to construct a new Fe-C phase diagram, where cementite forms below the eutectoid temperature due to the tendency of the Fe-C system toward ordering and carbon unbound to iron precipitates above this temperature in the form of clusters or graphite particles due to the tendency of this system toward phase separation.

Paraequilibria in the Fe-Si-C system and their relation to the bainite transformation in steels

The thermodynamic properties of silicon-containing cementite and e carbide are estimated, and the results obtained are used to perform the thermodynamic calculation of paraequilibria in the Fe-Si-C system. Even a relatively low silicon content is shown to substantially change the positions of virtually all boundaries of phase fields in the Fe-C phase diagram. The relation between the paraequilibria and bainite transformation was analyzed. The maximum supersaturation of retained austenite is found to be controlled by a thermodynamic factor, namely, the solubility of the paraequilibrium cementite. The thermodynamic specific features also cause the differences in the mechanisms of carbide precipitation from the α and γ phases in the absence of silicon redistribution: only cementite can precipitate from austenite, whereas both cementite and e carbide can precipitate from bainitic ferrite.

C—Fe—U (Carbon‐Iron‐Uranium)

AbstractFor Abstract see ChemInform Abstract in Full Text.

C-Fe-U (Carbon-Iron-Uranium)

The Fe-C phase diagram [1992Oka] is in the form of a double diagram, corresponding to the metastable equilibrium with cementite (Fe3C) or the stable equilibrium with graphite. The peritectic formation of the solid solution (austenite) based on the face-centered cubic (fcc) Fe is followed by the eutectic reaction, which yields austenite and cementite (or graphite). The eutectoidal decomposition of austenite yields body-centered cubic (bcc) (ferrite) and cementite (or graphite). In the C-U phase diagram [Massalski2], there are three intermediate phases. The monocarbide UC covers a wide composition range from 47.5-66 at.% C and has the NaCl-type cubic structure. Below the congruent formation temperature of this phase, there is a miscibility gap, where this phase splits into two phases with different C contents. The sesquicarbide U2C3 forms at 1823 °C through a peritectoid reaction between the above two phases. It has the Pu2C3 type cubic structure. UC2 forms through a peritectoid reaction at 1793 °C between the C-rich monocarbide and U2C3 and decomposes eutectoidally at 1516 °C. It has the CaC2 type tetragonal structure. The Fe-U phase diagram [1993Oka] has two intermediate phases of stoichiometric composition: Fe2U (MgCu2-type cubic) and FeU6 (MnU6 type tetragonal).

Simulation of segregation in continuous casting of high-carbon Fe-C alloys

In this paper, we propose a mathematical model devised to simulate the segregation profile on a cross-section of a high-carbon steel bloom produced by continuous casting, as a function of the operation conditions of the plant (withdrawal speed, superheat and coolant flow). The model consists of two coupled routines simulating the thermal and segregation behaviour of the system. The segregation routine is innovative, because (i) it simulates the segregation process accurately, and (ii) it can be used for on-line simulation of the behaviour of an industrial continuous casting process. In order to reduce the computational problem to a manageable degree of complexity, the physics of the process can be modelled with two time-dependent parameters describing (i) convective effects on the distribution of the solute (homogenisation parameter Co); (ii)displacements from thermodynamic equilibrium brought about by mass transport at the solid-liquid interface (effective partition coefficient oPe), a function of the total liquid fraction Ci,tot through the local thermal gradient ΔT). The numerical model is 2-D, but 3-D information is conveyed by the parameters Co); and oPe. The mathematical model includes the flowing balances energy, mass, species and momentum; the overall system can be simplified by including convective effects through the time-and space-dependent homogenisation parameter Co. The number of equations is matched to the number of unknowns by correlating the liquid phase composition with the temperature through the linearised equilibrium Fe-C phase diagram. The results of the numerical simulation of the segregation profile are in good agreement with the experimental data available in the literature; the results of the thermal simulation are in excellent agreement with experimental data obtained in this work.

A new empirical formula for the bainite upper temperature limit of steel

The definition of the practical upper temperature limit of the bainite reaction in steels is discussed. Because the theoretical upper temperature limit of bainite reaction, B 0, can neither be obtained directly from experimental measurements, nor from calculations, then, different models related to the practical upper temperature limit of bainite reaction, B S, are reviewed and analyzed first in order to define the B 0 temperature. A new physical significance of the B S and B 0 temperatures in steels is proposed and analyzed. It is found that the B 0 temperature of the bainite reaction in steels can be defined by the point of intersection between the thermodynamic equilibrium curve for the austenite→ferrite transformation by coherent growth (curve Z $$\gamma \to \overrightarrow \alpha $$ ) and the extrapolated thermodynamic equilibrium curve for the austenite→cementite transformation (curve ES in the Fe-C phase diagram). The B S temperature for the bainite reaction is about 50–55 °C lower than the B 0 temperature in steels. Using this method, the B 0 and B S temperatures for plain carbon steels were found to be 680 °C and 630 °C, respectively. The bainite reaction can only be observed below 500 °C because it is obscured by the pearlite reaction which occurs prior to the bainite reaction in plain carbon steels. A new formula, B S(°C) =, 630-45Mn-40V-35Si-30Cr-25Mo-20Ni-15W, is proposed to predict the B S temperature of steel. The effect of steel composition on the B S temperature is discussed. It is shown that B S is mainly affected by alloying elements other than carbon, which had been found in previous investigations. The new formula gives a better agreement with experimental results than for 3 other empirical formulae when data from 82 low alloy steels from were examined. For more than 70% of these low alloy steels, the B S temperatures can be predicted by this new formula within ±25°C. It is believed that the new equation will have more extensive applicability than existing equations since it is based on data for a wide range of steel compositions (7 alloying elements).

Application of a Nodular Iron Microstructure Model for an Industrial Casting

The objective of this study was to determine the applicability of using a microstructural model to predict the heat flow, microstructure and hardness distribution in an industrial casting. A mathematical model has been developed to describe the microstructure evolution in nodular cast iron during solidification of primary and eutectic phases, and the subsequent solid-state transformations during the eutectoid temperature range. Microstructural features such as volume fraction of constituent phases are calculated utilising the constraint conditions imposed by the Fe-C phase diagram and kinetic equations The model's predictive ability has been experimentally tested using the directional solidification technique on a eutectic composition nodular iron as well as being compared with a commercial microstructure simulation program. The hardness distribution of the nodular iron casting was calculated from the microstructure and the amount of alloying elements present using multiple regression equations available in the literature.

Effect of silicon content on transformation kinetics of austempered ductile iron

The present study investigated the effect of austenitising temperature (850, 900, and 950°C) and austempering time (0–7 h) on the volume of retained austenite of a 0.3 wt-%Mn ductile iron containing two different levels of silicon, namely 2.02 wt-% and 3.31 wt-%, and austempered at 360°C. The volume fraction of retained austenite and austenite carbon content results were then correlated with microstructural changes and impact toughness results. It is shown that the austenite stability is of great significance with respect to impact toughness, and that ferrite, when present in acicular form, can increase the mechanical stability of the austenite. It is observed that decreasing the austenitising temperature increases the driving force for the stage I transformation reaction in which mother austenite transforms to high carbon austenite plus acicular ferrite. However, the austenitising temperature has only a small effect on the kinetics of the stage II reaction in which high carbon austenite transforms to bainiticferrite plus carbides. In the low silicon iron, austenitising at 950°C results in a continuous network of intercellular and low carbon austenite which reduces impact properties. Intercellular austenite is attributed to the segregation of manganese and high austenitising temperature which decrease the carbon diffusion rate and delay ferrite nucleation and growth. Decreasing the austenitising temperature to 850°C increases the rate of transformation which results in a more uniform microstructure, stable high carbon austenite, and higher impact toughness. Silicon has the effect of modifying the Fe-C phase diagram such that a higher solution treatment temperature is needed to fully austenitise the iron. Furthermore, a three phase region of austenite-ferrite-graphite is introduced in the Fe-C-Si phase diagram. Consequently, austenitising at low solution treatment temperatures produces structures containing proeutectoidferrite. Increasing the austenitising temperature to 950°C leads to a more uniform acicular microstructure of stable high carbon retained austenite andferrite and results in optimum impact properties. Following short austempering times in irons containing 2.02% and 3.31% silicon, the carbon content of the retained austenite is low and on subsequent cooling to room temperature it transforms to martensite, resulting in low impact values. Optimum properties are obtained at intermediate austempering periods when both the amount of retained austenite and the austenite carbon content are maximum. Extending the austempering time causes the high carbon austenite to decompose to ferrite plus carbides, the stage II reaction, leading to a reduction in impact toughness.

Effect of silicon content on transformation kinetics of austempered ductile iron

The present study investigated the effect of austenitising temperature (850, 900, and 950°C) and austempering time (0–7 h) on the volume of retained austenite of a 0.3 wt-%Mn ductile iron containing two different levels of silicon, namely 2.02 wt-% and 3.31 wt-%, and austempered at 360°C. The volume fraction of retained austenite and austenite carbon content results were then correlated with microstructural changes and impact toughness results. It is shown that the austenite stability is of great significance with respect to impact toughness, and that ferrite, when present in acicular form, can increase the mechanical stability of the austenite. It is observed that decreasing the austenitising temperature increases the driving force for the stage I transformation reaction in which mother austenite transforms to high carbon austenite plus acicular ferrite. However, the austenitising temperature has only a small effect on the kinetics of the stage II reaction in which high carbon austenite transforms to bainiticferrite plus carbides. In the low silicon iron, austenitising at 950°C results in a continuous network of intercellular and low carbon austenite which reduces impact properties. Intercellular austenite is attributed to the segregation of manganese and high austenitising temperature which decrease the carbon diffusion rate and delay ferrite nucleation and growth. Decreasing the austenitising temperature to 850°C increases the rate of transformation which results in a more uniform microstructure, stable high carbon austenite, and higher impact toughness. Silicon has the effect of modifying the Fe-C phase diagram such that a higher solution treatment temperature is needed to fully austenitise the iron. Furthermore, a three phase region of austenite-ferrite-graphite is introduced in the Fe-C-Si phase diagram. Consequently, austenitising at low solution treatment temperatures produces structures containing proeutectoidferrite. Increasing the austenitising temperature to 950°C leads to a more uniform acicular microstructure of stable high carbon retained austenite andferrite and results in optimum impact properties. Following short austempering times in irons containing 2.02% and 3.31% silicon, the carbon content of the retained austenite is low and on subsequent cooling to room temperature it transforms to martensite, resulting in low impact values. Optimum properties are obtained at intermediate austempering periods when both the amount of retained austenite and the austenite carbon content are maximum. Extending the austempering time causes the high carbon austenite to decompose to ferrite plus carbides, the stage II reaction, leading to a reduction in impact toughness.

Effect of silicon content on impact properties of austempered ductile iron

This study investigates the effect of austenitising temperature (850°C, 900°C, 950°C) and austempering time (0 to 7 h) on the impact energy values of three irons containing increasing silicon contents (2.02%, 2.65%, 3.31%) and around 0.3% manganese, austempered at 360°C. It was shown that increasing the silicon content modifies the Fe-C phase diagram such that (a) a higher solution treatment temperature is required to fully austenitise the iron, and (b) the resulting austenite dissolves less carbon. Increasing the austenitising temperature simultaneously with the silicon content restores a high carbon austenite, a fully ausferritic structure and high impact properties. On the other hand, the low silicon iron austenitised at 950°C results in a microstructure containing a continuous network of intercellular austenite. This is attributed to the manganese whose segregation is more evident in the iron containing 2.02% silicon. For a given silicon content, increasing the austenitising temperature increases the carbon content dissolved in the austenite which in turn reduces the free energy controlling the rate at which the austempering reaction proceeds. Either short or extensive holding in the austempering bath leads to undesirable structures and low mechanical properties. Intermediate austempering periods resulted in structures containing ferrite needles, a large volume of high carbon stable austenite and a better combination of mechanical properties.