Influence of secondary carbides precipitation and transformation on the secondary hardening of laser melted high chromium steel

The influence of secondary carbides precipitation and transformation on the secondary hardening of laser melted high chromium steels was analyzed by means of scanning electron microscopy, transmission electron microscopy, and X-ray diffraction. The microstructure of laser melted high chromium steel is composed of austenite with supersaturated carbon and alloy elements and granular interdendritic carbides of type M₂₃C₆. Secondary hardening of the laser melted layer begins at 450 °C after tempering, and the hardness reaches a peak of 672HV at 560 °C and then decreases gradually. After tempering at 560 °C, a large amount of lamellar martensite was formed in the laser melted layer with a small quantity of thin lamellar M₃C cementite due to the martensitic decomposition. The stripy carbides precipitating at the grain boundaries were determined to be complex hexagonal M₇C₃ carbides and face centered cubic M₂₃C₆ carbides. In addition, the granular M₂₃C₆ carbides and fine rod-like shaped M₇C₃ carbides coexisted within the dendrites. As a result, the combined effects of martensitic transformation, ultrafine carbide precipitations, and dislocation strengthening result in the secondary hardening of the laser melted layer when the samples were tempered at 560 °C.     

Characterization of precipitates in a 7.9Cr–1.65Mo–1.25Si–1.2V steel during tempering

In this paper, the precipitates formed during the tempering after quenching from temperature 1150 °C for 7.90Cr–1.65Mo–1.25Si–1.2V steels are investigated using an analytical transmission electron microscope (A-TEM).The study of this tempering is carried out in isothermal and anisothermal conditions, by comparing the results given by dilatometry and hot hardness.Tempering is performed in the range of 300–700 °C. Coarse primary carbides retained after heat treatment are V-rich MC and Cr–Mo-rich M₇C₃ types. In turn, it gives a significant influence on the precipitation of fine secondary carbides, that is, secondary hardening during tempering. The major secondary carbides are Cr–Mo–V-rich M′C (and/or) Cr–Mo-rich M₂C type. The peak hardness is observed in the tempering range of 450–500 °C. In the end, we observe between 600 and 700 °C, that this impoverished changes the phase. At these high temperatures of tempering, we observe that there is a carbide formation of the types M₆C developing at the expense of the fine M₇C₃ carbides previously formed.     

Microstructural evolution in bainite,martensite, and δ ferrite of low activation Cr–2W ferritic steels

Abstract

The microstructural evolution in (2–15)Cr–2W–0·1C (wt-%) firritic steels after quenching, tempering, and subsequent prolonged aging was investigated, using mainly transmission electron microscopy. The steels examined were low induced radioactivation ferritic steels for fusion reactor structures. With increasing Cr concentration, the matrix phase changed from bainite to martensite and a dual phase of martensite and δ ferrite. During tempering, homogeneous precipitation of fine W₂C rich carbides occurred in bainite and martensite, causing secondary hardening between 673 and 823 K. With increasing tempering temperature, dislocation density decreased and carbides had a tendency to precipitate preferentially along interfaces such as bainite or martensite subgrain boundaries. During aging at high temperature, carbides increased in size and carbide reaction from W₂C and M₆C to stable M₂₃C₆ occurred. No carbide formed in δ ferrite. The precipitation sequence of carbides was analogous to that in conventional Cr–Mo steels.

MST/1049     

Effect of heat treatment on microstructure and mechanical properties of Cr–Ni–Mo–Nb steel

Abstract

The microstructure and mechanical properties of a medium carbon Cr–Ni–Mo–Nb steel in quenched and tempered conditions were investigated using transmission electron microscopy (TEM), X-ray analysis, and tensile and impact tests. Results showed that increasing austenitisation temperature gave rise to an increase in the tensile strength due to more complete dissolution of primary carbides during austenitisation at high temperatures. The austenite grains were fine when the austenitisation temperature was <1373 K owing to the pinning effect of undissolved Nb(C,N) particles. A tensile strength of 1600 MPa was kept at tempering temperatures up to 848 K, while the peak impact toughness was attained at 913 K tempering, as a result of the replacement of coarse Fe rich M₃C carbides by fine Mo rich M₂C carbides. Austenitisation at 1323 K followed by 913 K tempering could result in a combination of high strength and good toughness for the Cr–Ni–Mo–Nb steel.     

Effect of tempering temperature on the microstructure and properties of ultrahigh-strength stainless steel

The microstructure, precipitation and mechanical properties of Ferrium S53 steel, a secondary hardening ultrahigh-strength stainless steel with 10% Cr developed by QuesTek Innovations LLC, upon tempering were studied by scanning electron microscopy (SEM), transmission electron microscopy (TEM), X-ray diffraction (XRD), and tensile and impact tests. Based on these results, the influence of the tempering temperature on the microstructure and properties was discussed. The results show that decomposition occurred when the retained austenite was tempered above 440 °C and that the hardening peak at 482 °C was caused by the joint strengthening of the precipitates and martensite transformation. Due to the high Cr content, the trigonal M₇C₃ carbide precipitated when the steel was tempered at 400 °C, and M₇C₃ and M₂C (5–10 nm in size) coexisted when it was tempered at 482 °C. When the steel was tempered at 630 °C, M₂C and M₂₃C₆ carbides precipitated, and the sizes were greater than 50 nm and 500 nm, respectively, but no M₇C₃ carbide formed. When the tempering temperature was above 540 °C, austenitization and large-size precipitates were the main factors affecting the strength and toughness.     

TEM investigation of the tempering behaviour of the maraging PH 17.4 Mo stainless steel

The PH 17-4 Mo steel (Z6 CND 17.04.02), used in the steam generator of nuclear reactors, was investigated in order to determine the structural evolution occurring during tempering carried out under various conditions of duration and temperature. The formation and growth of different types of carbides such as Mo₂C, M₂₃C₆ and M₇C₃ and of Fe₂Mo intermetallic compound were studied and also of reversed austenite. A small secondary hardening peak was observed for tempering close to 400' C which is related to the Mo₂C carbide precipitation; beyond this temperature, softening occurs.     

Surface electron beam melting and alloying of tool steels

Abstract

Surface melting and alloying of D3 steel using an electron beam has been carried out to improve its surface microstructure and properties. The solution of primary carbides, together with rapid solidification and subsequent cooling, enhance the solubility of alloying elements in the γ Fe phase and thus influence the behaviour of the steel on subsequent tempering. The surface melted zone consists of dendrites without primary carbides, which is also the case for samples alloyed with WC, SiC, or Al₂O₃. When alloyed with TiC or TiB₂, the materials contain TiC or TiB₂ primary phase respectively in addition to the iron rich dendrites. Some unmelted TiB₂ particles are also present. On tempering, both electron beam melting and alloying change the secondary hardening characteristics, increasing the peak hardness and the peak hardness temperature.

MST/1194     

On the microstructure and mechanical properties of high-speed steels

The microstructure of high-speed steels consists of a martensitic matrix with a dispersion of two sets of carbides. These carbides are usually known as primary and secondary carbides. The role of the primary carbides has been reported to be of no importance in strengthening the steels, due to their large size and large interparticle spacing. The present authors have studied the role of the primary carbides on the wear of high-speed steels and found them to be of no importance, and under certain conditions contributing to higher wear rates. It has been shown analytically and experimentally that in quenched and tempered high-speed steels, the precipitation of the secondary hardening carbide (cubic M₂C type) is the main reason for the improved strength and wear resistance. This shows that the secondary hardening phenomenon of high-speed steels is a direct result of the hardening caused by the precipitation of the cubic M_2C-type carbide. The present study has estimated that at peak hardness the volume fraction of secondary hardening carbides is approximately 20%. The measured strength of high-speed steels was found to be lower than the theoretically calculated strength due to non-homogeneous precipitation of the secondary hardening carbides. Areas which were observed to be free from secondary hardening carbides are real and are not artefacts. It has been shown that the strength of high-speed steel in the region of peak hardness depends primarily on the precipitation of the secondary hardening carbide and secondarily on martensitic strengthening.     

Precipitate design for creep strengthening of 9% Cr tempered martensitic steel for ultra-supercritical power plants

Abstract

It is crucial for the carbon concentration of 9% Cr steel to be reduced to a very low level, so as to promote the formation of MX nitrides rich in vanadium as very fine and thermally stable particles to enable prolonged periods of exposure at elevated temperatures and also to eliminate Cr-rich carbides M₂₃C₆. Sub-boundary hardening, which is inversely proportional to the width of laths and blocks, is shown to be the most important strengthening mechanism for creep and is enhanced by the fine dispersion of precipitates along boundaries. The suppression of particle coarsening during creep and the maintenance of a homogeneous distribution of M₂₃C₆ carbides near prior austenite grain boundaries, which precipitate during tempering and are less fine, are effective for preventing the long-term degradation of creep strength and for improving long-term creep strength. This can be achieved by the addition of boron. The steels considered in this paper exhibit higher creep strength at 650 °C than existing high-strength steels used for thick section boiler components.     

Secondary carbide precipitation in an 18 wt%Cr-1 wt% Mo white iron

High chromium (18%) white irons solidify with a substantially austenitic matrix supersaturated with chromium and carbon. The austenite is destabilized by a hightemperature heat treatment which precipitates chromium-rich secondary carbides. In the as-cast condition the eutectic M₇Ca₃ carbides are surrounded by a thin layer of martensite and in some instances an adjacent thicker layer of lath martensite. The initial secondary carbide precipitation occurs on sub-grain boundaries during cooling of the as-cast alloy. After a short time (0.25 h) at the destabilization temperature of 1273 K, cuboidal M₂₃C₆ precipitates within the austenite matrix with the cube-cube orientation relationship. After the normal period of 4 h at 1273 K, there is a mixture of M₂₃C₆ and M₇C₃ secondary carbides and the austenite is sufficiently depleted in chromium and carbon to transform substantially to martensite on cooling to room temperature.     

Influence of grain boundary carbide density on impact behaviour of C–Mn–Nb–Al steels

Abstract

The influence of grain boundary carbide density on impact behaviour has been examined for C–Mn–Nb–Al steels by (i) normalizing at increasing temperatures above the Ac₃, and (ii) tempering for long times below the AC₁. Low normalizing temperatures (i) resulted in a large number of grain boundary carbides possibly because incomplete homogenization on austenitizing produces a high concentration of carbon at the boundaries. Raising the normalizing temperature reduced the number of grain boundary carbides as well as refining their size, but the expected improvement in impact behaviour was not realized because grain size also increased. Tempering at 680°C raised the grain boundary carbide density considerably and completely destroyed the pearlite colonies; tempering at 600°C (ii) gave a lower increase in carbide density and destroyed fewer pearlite colonies. Only small changes in grain size and grain boundary carbide thickness were noted so that the deterioration in impact behaviour obtained on tempering at 680°C could be ascribed mainly to this increase in grain boundary carbide density. Analysis of all the results suggests that an increase in grain boundary carbide density by 20 mm^?1 at constant grain size results in an increase in transition temperature of ~30 K. This agrees with the 20–30 K rise in transition temperature reported in a previously published paper which relied on a linear regression approach.

MST/424     

Einfluß des Anlassens auf Gefüge und Härte warmfester 9%Cr-Stähle

Influence of tempering on microstructure and hardness of high-temperature 9%Cr-steels The influences of temperature and duration of tempering on hardness and microstructure were investigated at high-temperature martensitic and low-carbon steels with 9% chrome and the further alloying elements molybdenium, vanadium, niobium and partially tungsten. After austenitizing and subsequent air cooling the steels were tempered at temperatures below, at and above Ac_1b for different times and finally a hardness test was performed. Making use of the temperature dependence of the hardness tempering diagrams were constructed and the Hollomon-Jaffe-Parameter on the three steels was determined within its application limits. Micrographs of the structure shows the formation of the carbides and the martensite. At tempering temperatures below Ac_1b a decrease of hardness occurs, above Ac_1b, a hardness rise due to the partial austenitizing was obtained. While hardening below Ac_1b, the tempering quality increases from P 91, NF 616 to E 911.     

Thermal-induced evolution of secondary phases in Cr–Mo–V low alloy steels

Four low alloy steels with different contents of molybdenum and vanadium were investigated. The steels were annealed at 773, 793, 853, 873, 933, 973, and 993 K for 500, 1000, 3000, and 10000 h. Techniques of transmission electron microscopy and thermodynamic calculations (ThermoCalc) were used to characterise influence of the steel bulk composition and the annealing conditions on evolution of carbides M₃C, M₂C, M₇C₃, M₂₃C₆, M₆C, and MC (M=metallic element). Changes in structure types and metal compositions of the carbides were characterised in detail. The work was done with the intention to obtain more information about the secondary phase evolution in low alloy steels used in energy industries.     

Microstructural characterisation of vacuum sintered T42 powder metallurgy high-speed steel after heat treatments

High-speed steel powders (T42 grade) have been uniaxially cold-pressed and vacuum sintered to full density. Subsequently, the material was heat treated following an austenitising + quenching + multitempering route or alternatively austenitising + isothermal annealing. The isothermal annealing route was designed in order to attain a hardness value of ～50 Rockwell C (HRC) (adequate for structural applications) while the multitempering parameters were selected to obtain this value and also the maximum hardening of the material (～66 HRC). Microstructural characterisation has been carried out by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). The microstructure consists of a ferrous (martensitic or ferritic) matrix with a distribution of second phase particles corresponding to nanometric and submicrometric secondary carbides precipitated during heat treatment together with primary carbides. The identification of those secondary precipitates (mainly M₃C, M₆C and M₂₃C₆ carbides) has allowed understanding the microstructural evolution of T42 high-speed steel under different processing conditions.     

Influence of nickel on secondary hardening of a modified AISI H13 hot work die steel

This paper focuses on the effects of nickel on secondary hardening of a modified H13 hot work die steel. Both the non‐nickel steel and the nickel‐added steel get a secondary hardening peak at 520 °C, and the secondary hardening peak trends to increase in the nickel‐added steel. On the basis of scanning electron microscope and transmission electron microscope observation, the rise of the secondary hardening peak is in connection with the precipitation of M₃C type carbides. More strip‐shaped and needle‐shaped M₃C type carbides precipitated from matrix. By means of internal friction, the result suggests that nickel does not affect the position of the Snoek‐Kê‐Köster peak, but the height of Snoek‐Kê‐Köster peak of the nickel‐added steel is higher, which indicates nickel enhances the interaction between dislocations and interstitial atoms, promoting the precipitation of carbides.     

Precipitation behaviour of a sensitized AISI type 316 austenitic stainless steel in hydrogen

The purpose of this study was to characterize the precipitation behaviour of AISI type 316 steel in hydrogen. The different precipitates (M₂₃C₆, M₆C), the intermetallicχ-phase and the martensitic phase (α′,ε) were determined by using transmission electron microscopy (TEM) and X-ray diffraction techniques. All the specimens were sensitized at 650? C for 24 h. Some samples were carburized up to 2 wt% C. Additions of carbon content decrease the time required for sensitization. Short-term (24 h) exposure of this steel to sensitization temperature results in a complex precipitation reaction of various carbides and intermetallic phases. Hydrogen was introduced by severe cathodic charging at room temperature. This study indicates that by conventional X-ray techniques it is possible to detect those precipitates and their behaviour in a hydrogen environment. The zero shift as observed by X-ray diffraction from the carbides (M₂₃C₆, M₆C) and the intermetallicχ-phase, indicates that those phases absorb far less hydrogen than the austenitic matrix. TEM studies reveal that hydrogen inducesα′ martensite at chromium-depleted grain-boundary zones, near the formation of the carbides.     

Effect of thermal exposure on the stability of carbides in Ni-Cr-W based superalloy

The microstructure evolutions of Ni-Cr-W based superalloy during thermal exposure have been investigated systematically. M₆C carbides in the alloy decompose into M₂₃C₆ carbides at temperatures from 650 to 1000 °C due to its high content of Cr. The M₆C carbides decompose dramatically from 800 to 900 °C. At temperatures up to 1000 °C, a few M₂₃C₆ carbides form on the surface of M₆C carbides. The decomposition behavior of primary M₆C can be explained by the following reaction: M₆C → M₂₃C₆ + Me (W, Ni, Cr, Mo). At temperatures below 900 °C, coarse lamellar M₂₃C₆ carbides precipitate at the grain boundaries. The carbide lamellae line almost perpendicular to the grain boundaries. While the temperature is above 1000 °C, discrete M₂₃C₆ carbides precipitate at the grain boundaries. Moreover, there are lots of small M₂₃C₆ particles precipitated around M₆C carbides from 650 to 1000 °C.     

Phase transformations in Fe–Ni–Cr heat-resistant alloys for reformer tube applications

Fe–35Ni–25Cr–0.4C alloys with different compositions are aged between 750 and 1150°C up to ～10,000?h. As-cast microstructure contains interdendritic carbides of type M₇C₃ (‘Cr₇C₃’) and MC (‘NbC’). At service temperatures, M₇C₃ transform into M₂₃C₆ (‘Cr₂₃C₆’) within hours. Then, a hardening precipitation of secondary intragranular M₂₃C₆ occurs over hundreds of hours, the nose of the ‘temperature-time-hardening’ curve being around 1000°C. G phase forms after long aging; its solvus temperature and formation kinetics depend on silicon content. Z phase is observed after long aging at 950°C or above. G and Z phases form at the expense of MC. Very long aging causes nitridation under air, with first a transformation of M₂₃C₆ into chromium-rich M₂X carbonitrides (X?=?C,N), then of MC into chromium-rich MX carbonitrides.     

Microstructures and high temperature properties of spray formed niobium‐containing M3 high speed steel

The billets of M3 high speed steel (HSS) with or without niobium addition were prepared via spray forming and forging, and the corresponding microstructures, properties were characterized and analysed. Finer and uniformly‐distributed grains without macrosegregation appear in the as‐deposited high speed steel that are different to the as‐cast high speed steel, and the primary austenite grain size can be decreased with 2% niobium addition. Niobium appears in primary MC‐type carbides to form Nb₆C₅ in MN2 high speed steel, whereas it contributes less to the creation of eutectic M₆C‐type carbides. With same treatments to forged MN2 high speed steel and M3 high speed steel, it is found that the peak hardness of these two steels are almost the same, but the temper‐softening resistance of the former is better. With higher high‐temperature hardness of the forged MN2 high speed steel, its temper softening above 600 °C tends to slow down, which is related to the precipitation of the secondary carbides after tempering. A satisfactory solid solubility of Vanadium and Molybdenum can be obtained by Nb substitution, precipitation strengthening induced by larger numbers of nano‐scaled MC and M₂C secondary carbides accounts for the primary role of determining higher hardness of MN2 high speed steel. The results of the wear tests show that the abrasive and adhesive wear resistance of MN2 high speed steel can be improved by the grain refinement, existence of harder niobium‐containing MC carbides, as well as solute strengthening by more solute atoms. The oxidational wear behavior of MN2 high speed steel can be markedly influenced by the presence of the high hardness and stabilization of primary niobium‐containing MC‐type carbides embedded in the matrix tested at 500 °C or increased loads. The primary MC carbides with much finer sizes and uniform distribution induced by the combined effects of niobium addition and atomization/deposition would be greatly responsible for the good friction performance of the forged MN2 high speed steel.     

Reversed austenite in 0Cr13Ni4Mo martensitic stainless steels

The austenite reversion process and the distribution of carbon and other alloying elements during tempering in 0Cr13Ni4Mo martensitic stainless steel have been investigated by in-situ high temperature X-ray diffraction (XRD) and scanning transmission electron microscopy (STEM). The microstructure of the reversed austenite was characterized using transmission electron microscopy (TEM). The results revealed that the amount of the reversed austenite formed at high temperature increased with the holding time. Direct experimental evidence supported carbon partitioning to carbides and Ni to the reversed austenite. The reversed austenite almost always nucleated in contact with lath boundary M₂₃C₆ carbides during tempering and the diffusion of Ni promoted its growth. The Ni enrichment and the ultrafine size of the reversed austenite were considered to be the main factors that accounted for the stability of the reversed austenite.