M A T E R I A L S C H A R A C T E R I Z A T I O N 5 9 ( 2 0 0 8 ) 1 0 0 7 – 1 0 1 4

Microstructure evolution during hot rolling and heat treatmentof the spray formed Vanadis 4 cold work steel

Fei Yana,⁎, Haisheng Shib, Bingzhong Jinb, Junfei Fanb, Zhou Xua

aKey Laboratory for High Temperature Materials and Tests of Ministry of Education,Shanghai Jiao Tong University, 1954 Hua Shan Road, Shanghai 200030, ChinabShanghai Baosteel Research Institute, Shanghai 201900, China

A R T I C L E D A T A

⁎ Corresponding author. Tel.: +86 21 26649626E-mail address: yanfeisjtu@163.com (F. Yan

1044-5803/$ – see front matter © 2007 Elsevidoi:10.1016/j.matchar.2007.08.012

A B S T R A C T

Article history:Received 25 March 2007Received in revised form3 August 2007Accepted 13 August 2007

A high alloyed Vanadis 4 steel was produced by spray forming, and the microstructureevolution during hot rolling and annealing processes was characterized. It was found thatthe as-sprayed Vanadis 4 steel has a homogeneous and fine microstructure with uniformdispersion of carbides. The hot rolling temperature is the key factor in controlling theevolution of type, morphology and distribution of carbides, as well as the matrixmicrostructure of the as-rolled steels. The optimized processing parameters for the as-sprayed Vanadis 4 steel are rolling at 1050 °C and annealing at 900 °C. The microstructuralevolution mechanisms during hot rolling and annealing are determined according to themicrostructural analysis of the material at different stages.

© 2007 Elsevier Inc. All rights reserved.

Keywords:Spray formingVanadis 4 steelMicrostructure evolution

1. Introduction

Vanadis 4 (V4) steel is a powder metallurgical (PM) highalloyed cold work tool steel containing (wt.%) 1.5 C, 1.0 Si, 0.4Mn, 8.0 Cr, 1.5 Mo, 4.0 V. The steel offers an extremely goodcombination of wear resistance and ductility for high perfor-mance tools. However, high cost and complicated workingprocedures are the main disadvantages of powder metallurgy,which restrains the wide application of the steel. For somenear-net-shape applications, spray forming is expected toreplace powder metallurgy to produce high alloyed steels.Spray forming, also called spray casting or spray deposition, isthe inert gas atomization of a liquid metal stream into varioussized droplets which are then propelled away from the regionof atomization by the fast flowing, atomizing gas [1]. Thedroplets were collected and solidified on a substrate, andfinally deposit into a coherent, near fully dense product. Rapidsolidification effects inherent in the spray forming processdue to high heat exchange rate at the droplet–gas interfaceand also on the deposition surface ensures considerablechemical and microstructural homogeneity of the product

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[2]. As a result, the material microstructures by spray formingdiffer significantly from those of materials by both conven-tional casting and powdermetallurgy. Application of the sprayforming technique in the field of steels has been focused onhigh alloyed steels where the benefits of the techniqueespecially the possibility to produce a fine-grained, segrega-tion-free microstructure are very promising [3–6]. The costbenefit derives from the single-step operation of convertingmolten alloy directly into a semi-finished product. Therefore,spray forming has emerged as a key competitor for existingtechnologies, especially powder metallurgy. But, as reportedinmany papers, an excess solid fraction in the spray generatesa product with some pores due to insufficient liquid phaseavailable to provide bonding of particles during solidificationof the spray formedmaterials [6,9,10]. After spray forming, thebillet need to be forged or hot rolled to eliminate the remainingporosity. Although many efforts have been made by usingspray forming to produce high alloyed steels [3–8], there islittle work published on the microstructural evolution behav-ior during the whole hot working process [4]. Furthermore, toour knowledge, few works have been published on comparing

.

1008 M A T E R I A L S C H A R A C T E R I Z A T I O N 5 9 ( 2 0 0 8 ) 1 0 0 7 – 1 0 1 4

the microstructure of spray formed steel with commercialpowder metallurgical equivalent.

The V4 cold work steel was produced by spray forming andits microstructure was studied. Detailed study was focused onthemicrostructural evolution during hot rolling and annealingin order to optimize processing conditions and to examinetheir effects on microstructure and carbides. The objectives ofthe present work were to: (1) select the optimum parametersof hot rolling and annealing for the as-sprayed V4 steel;(2) investigate the microstructural evolution in this steelduring hot working and heat treatment process; and (3)compare the microstructure of the as-sprayed steel after hotrolling and annealing with that of a commercial powder met-allurgical equivalent manufactured by UDDEHOLM, Sweden.

Fig. 1 – (a) Optical micrograph of the as-sprayed V4 steel,(b) TEM morphology of the twined matrensite.

2. Experimental Procedure

The steels were melted in a vacuum induction furnace andthen cast into rods as the feedstock for spray forming. Thefeedstock was heated in an induction-furnace and soaked atabove themelting point for 20min. Themoltenmetal flow ratewas set at approximately 0.1 kg/s, using N2 with a pressure of2.2 MPa as the atomizing gas. The atomized droplets werecooled and driven towards a revolving substrate to form adense product. The distance from the nozzle to the substratewas set as 360 mm and the copper substrate was rotated at aspeed of 10 rpm during atomization and deposition. Thespray-forming process was completed in about 40 s, anda gauss-shaped billet with about 130 mm in diameter and30 mm in height was obtained.

Hot rolling and annealing were carried out based on thecritical temperatures, which were determined by dilatometryusing a “ThermecMastor_Z” thermomechanical simulator. Cy-lindrical specimens 12 mm in height and 8 mm in diameterwere heated to 1100 °C at a rate of 0.2 °C/s, being held for10min, and then cooled firstly to 600 °C at a rate of 0.1 °C/s andfinally to 200 °C at a rate of 1 °C/s. Phase transformations weredetected by monitoring the fractional change in dilatationwith temperature.

The spray formed billet wasmachined into specimens witha thickness of 16 mm. The specimens were then heated at arate of 10 °C/min to the rolling temperatures of 850 °C, 900 °C,950 °C, 1050 °C, 1150 °C, respectively. After holding for 15 min,some specimens were quenched in water in order to holdthe carbides' morphology before hot rolling. Other specimenswere rolled in a single 60% reduction pass, and then cooled toroom temperature in sand.

Three groups of as-rolled specimens were annealed at850 °C, 900 °C, and 950 °C respectively, with each group con-taining five specimens rolled at different temperatures. Afterbeing isothermally held for 2 h, the specimens were cooled to500 °C at a rate of 30 °C/h, and then cooled in air.

Microstructural observations were made optically, with anS-4200 field emission scanning electronmicroscope (SEM) andwith an H-800 transmission electron microscope (TEM). Thinfoils for transmission electron microscopic studies were pre-pared from 3 mm disks, grounded to a thickness of about50 μm and electropolished in an electrolyte containing 5%perchloric and 95% ethanol at −20 °C. Bright field and dark

field imaging and selected area diffraction patterns were usedto identify the carbides.

The specimens for Vicker's hardness testing were per-formed on a Vickers hardness tester (Model: FR-3E) with 30 kgload and each measurement was tested for 8 s. The averagehardness of five measurements was obtained as the finalhardness.

3. Results and Discussion

3.1. Microstructure of the AS-sprayed V4 Steel and PhaseTransformation Temperature

An optical micrograph of the as-sprayed microstructure isshown in Fig. 1a. Spray forming resulted in a substantialreduction in microstructural scale for both grain size andcarbide size compared with conventionally cast equivalent[11]. Fine, uniformly distributed equiaxed grains in diametersranging from 8 to 10 μm are observed. In addition, large pri-mary and interconnected eutectic carbides observed in theconventional cast material are replaced by a more uniformlydistributed spheroidal carbides in the as-sprayed material. Asshown in Fig. 1b, observation by TEM shows that the matrix ofthe microstructure of the as-sprayed material was predomi-nantly twinned martensite. Selected area diffraction patternsconfirmed the presence of V-rich MC and Cr-base M7C3 type

Fig. 2 –Measured dilatation variation with temperature.

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carbides. Most of the MC and some big M7C3 carbides rangingfrom 0.5 to 2 μm in diameter are distributed on the grainboundaries. Other small M7C3 carbides with the diameter ofabout 180 nm distribute in the grains.

The variation in specimen dimensions with temperatureduring heating and cooling process by the dilatometry ex-periment is shown in Fig. 2. The transformation temperatures,845 °C (Ac1s) and 890 °C (Ac1f), are defined as the points wherethe curve leaves the tangent to the linear portion. It is in-

Fig. 3 –SEM micrographs of the as-sprayed steels quench

teresting to find that an obvious flexure appears duringheating between temperatures ranging from 650 °C to 760 °C,indicating the decomposition of martensite. The cooling curveindicates that the high alloyed austenite has a strong stability,as no obvious transformations could be observed other thanthe bainite reaction which occurred only at very slow coolingrates.

3.2. Microstructure Evolution During Hot Rolling

Fig. 3 shows the SEM micrographs of the microstructuresbefore hot rolling at different temperatures. It can be seen thatprecipitation occurred in all of the cases and obvious dif-ferences lie in the amount andmorphology of the precipitates.Fig. 3 indicates that the amount of the small carbidesdecreases with increasing temperature. As shown in Fig. 3a,a larger number of small carbides appear at 850 °C. TEMobservation shows that they are M3C or M7C3, however, thebigger, irregular phases, as shown by arrows, are MC (VC)carbides. When the temperature is elevated to 950 °C or evenhigher, only MC and M7C3 can be found. Thus, reheating tem-perature leads to major microstructural changes involvingphase transformation, dissolution and spheroidization of car-bides formed in the as-sprayed steels.

Fig. 4 shows the SEM observation on the microstructures ofthe as-rolled steels which were rolled at 850 °C, 950 °C, 1050 °C,and 1150 °C, respectively. The stable MC carbide particles

ed at: (a) 850 °C; (b) 950 °C; (c) 1050 °C and (d) 1150 °C.

Fig. 4 –SEM micrographs of the as-sprayed steel rolled at: (a) 850 °C; (b) 950 °C; (c) 1050 °C and (d) 1150 °C.

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show little tendency to coarsen. The M7C3 carbide, unlike MC,is very sensitive to the rolling temperature, i.e., both the sizeand morphology of the M7C3 carbides vary with temperature.For example, TEM observation suggests that many coarsecarbides shown in Fig. 4a are of M7C3 type. But their size andamount are all decreased with increased rolling temperature.Fig. 5a is the TEM image of those small carbides shown inFig. 4a. It was confirmed by electron diffraction patterns thatmost of the small carbides are of M7C3 type, and some M3Ccarbides are also found. Electron diffraction patterns of M3Cand M7C3 carbides are shown in Fig. 5b and c, respectively.

The characteristics shown in Fig. 3a and Fig. 4a suggest thatthe microstructural evolution process at 850 °C differs sig-nificantly from those at other temperatures. This can be at-tributed to the difference of the diffusion coefficient of thealloy elements.

Nucleation accompanied by growth and followed by coars-ening constitutes the processes during the hot rolling and sub-sequent cooling process. The nucleation rate as the derivative ofthe precipitate density N vs time t can be expressed as [12]:

dNdt

¼ N0Zb⁎ exp �DG⁎kt

� �exp � s

t

� �and b⁎ ¼ 4kR⁎2DXC0

a4: ð1Þ

Where N0 is the number of nucleation sites per unitvolume, Z is the Zeldovich factor, τ is the incubation time,ΔG⁎ is the nucleation driving force, R⁎ is the nucleation radius,

D is the diffusion coefficient of solute atoms in thematrix, andXC0

is the initial solute mole fraction.The coarsening of second phase particles as a result of a

reduction in interfacial energy was theoretically treated byLifshitz and Slyozov [13] and Wagner where the result wasexpressed as:

r̄3 � r̄30 ¼ 8gDceV2m

9RTt: ð2Þ

Where r ̄ is the average particle radius at time t, r ̄0 is theaverage particle radius at the onset of coarsening, γ is theinterfacial free-energy of the particle–matrix interface, D isthe diffusion coefficient of solute atoms in the matrix, ce isthe concentration of solute in the matrix in equilibrium witha particle of infinite size, Vm is the molar volume of theparticle, R is the gas constant, and T is the temperature.

The nucleation rate and coarsening of the carbides are allgoverned by the diffusion coefficient D of solute atoms, andtherefore, themagnitude of the diffusion coefficients providesa useful comparison basis. The diffusion coefficient D can bepresented in terms of an Arrhenius equation:

D ¼ D0 exp � QRT

� �: ð3Þ

Where D0 is the pre-experimental factor for the diffusioncoefficient, and Q is the activation energy for diffusion.

Fig. 5 – (a) TEM micrograph of the carbides in the materialrolled at 850 °C, electron diffraction patterns in [100] zoneaxis of M3C particle (b), and in [431] zone of M7C3 particle (c).

Fig. 6 – (a) TEM micrograph of the polygonal carbide, (b) itselectron diffraction pattern in [312] zone axis.

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According to Eq. (3), the diffusion coefficients of carbon (withD0=0.738 cm2 · s−1 and Q=158.98 kJ · mol−1 [14]) and chromium(with D0=4.08 cm2 · s−1 and Q=286.8 kJ · mol−1 [14]) at 850 °Care 2.97×10−8 cm2 · s−1 and 1.86×10−13cm2 · s−1, respectively.The diffusion coefficients between carbon and chromium areof great difference.

Sincemartensite is not an equilibrium phase, when steel isheated below the eutectoid temperature, the thermodynam-ically stable phases α and M3C start to precipitate. At tem-peratures ranging from 650 °C to 760 °C, precipitation of M3Ccarbides can be written as: α′⇒α+M3C. This reaction is similar

to the tempering of martensite around 700 °C [15]. In the earlystages of the decomposition, the reaction is controlled bycarbon diffusion and is thus very rapid. The change of thematerial's volume occurs because a large number of M3Ccarbides precipitate from matrix. This is why an obviousflexure appears in dilatability–temperature curve (Fig. 2)between the temperatures of 650 °C and 760 °C. The amountof the precipitates below the A1 temperature was tremendous,as shown in Fig. 3a. A metastable supersaturated phase willnot only generate nuclei but also cause deposition on thegenerated nuclei and consequent growth of the clusters. Thisdecreases the supersaturation, and correspondingly increasesthe critical nucleus size, and therefore promotes the coarsen-ing process [16]. As the result of this, the Gibbs–Thomsoneffect becomes so significant that smaller clusters dissolve,transfer their mass to larger growing clusters and coarsen thesize distribution [17]. When small clusters shrink to theircritical nucleus size, they become thermodynamically unsta-ble, and spontaneously disintegrate and eventually vanish[18]. This is why a large number of clusters nucleate and thengrow, Fig. 3a, and many coarse carbides appear, Fig. 4a.Occasionally coarse carbides, which are favorably situated atthe grain corners, appear to have grown from those smallcarbides close to them. This conclusion is supported by theobservation that such particles are frequently polygonal inshape, and there are comparatively fewer small carbides that

Fig. 8 –TEM micrograph of martensite in the as-rolled steel.

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exist in the vicinity of the coarse particles. On the other hand,inmany cases, the firstly formed carbides are non-equilibriumphase and they are replaced bymore stable forms gradually inthe precipitation sequence [19]. After the precipitation of M3Ccarbides, the kinetics of the diffusion mechanism changesfrom carbon diffusion control to chromium diffusion control.Therefore, the observed small M7C3 carbides, as shown in Fig.5, may be attributed to the transformation from M3C to M7C3

carbides. The possibility of M7C3 carbide formation from M3Ccarbides was also proved by Homolová [20] and Liu [21,22].

During the heat up process, the carbide precipitation anddissolution is a competition process. At relatively lower tem-perature the precipitation dominates the processwhile at hightemperature dissolution dominates. Therefore, the decreaseof the amount of the small carbides can be attributed to thedissolution with the increase of temperature, and the satura-tion of the alloy elements in the matrix will increase also. Themicrostructure of the steel which rolled at 1050 °C shows thebest uniform distribution of the carbides with ideal averagesize, Fig. 4c. Some big polygonal carbides appear at the triple-junctions of grain boundaries in thematerial which was rolledat 1150 °C, Fig. 4d. Those polygonal carbides were not found inthe materials rolled at other temperatures. TEM morphologyof this kind of carbide and its corresponding electrondiffraction pattern indicate that this carbide is of M7C3, Fig 6.

After hot rolling, the supersaturated alloying elements willprecipitate during cooling. The precipitation of the carbides isdetermined by two basic factors. One is that the alloying ele-ments in the matrix should be in supersaturation state, andthe other that the atoms have enough ability to diffuse. Both ofthese factors closely correlate with the temperature. It is thusreasonable to infer that the supersaturation and the diffusionabilities of the alloy atoms are at their highest at 1150 °C incomparison with those at other temperatures.

In many diffusion phase transformation and precipitationprocess, nucleation of the product phase occurs heteroge-neously at some preferential nucleation sites in the matrixsuch as the grain boundaries, dislocations and second phases.Grain boundaries often contain a high density of defects,which can be increased by plastic deformation [23]. Therefore,

Fig. 7 –Hardness of the as-sprayed steels rolled at differenttemperatures.

carbide precipitation on grain boundaries will be more easily,and precipitates which are growing on boundaries will gain abigger proportion of their material from within the grains.Although the dissolution of chromium carbides in matrix hasbeen reported by researchers as a more rapid process than thedissolution of other carbides, the complete dissolution of M7C3

carbides is very difficult even at very high austenitizingtemperature [24]. When the as-sprayed material was heatedto 1150 °C, many M7C3 carbides on the grain boundary still notdissolved completely. They act as preferential nucleation sitesand grow up quickly. Thus, big polygonal M7C3 carbidesformed and the various shapes of grain boundary M7C3

carbides may possibly be explained by variation of highangle boundaries. There is evidence from Fig. 4d that themost frequently occurring form of grain boundary precipitateis discontinuous. The limited boundary movement observedcan be explained in the manner that the precipitation of thecarbides in the grain boundaries act as effective obstacles ofthe growth of the grains, which ensures fine grains remainedin elevated temperature.

Results of the hardness as a function of the rollingtemperature are shown in Fig. 7. It can be seen that thehardness increases with increased rolling temperature. Espe-cially there is a sharp increase from 900 °C to 950 °C. Theorigination of this effect can be ascribed to the completelydifferent matrix microstructure after hot rolling. TEM obser-vation shows the matrix of the as rolled steels rolled at 850 °Cand 900 °C consists of ferrite, as shown in Fig. 5a. Whereas thematrix of the as rolled steels which rolled at or higher than950 °C consists of martensite, and the representative mor-phology of the matrix is shown in Fig. 8 (by the steel rolled at1050 °C).

The as-rolled microstructure was controlled by the phasetransformation during hot rolling and cooling process, and therolling temperature is actually a key in this process. When thesteel is heated to 900 °C (slightly above A1), little carbidedissolves and the contents of both carbon and alloyingelements in austenite are lower. The unstable austenitecould transform into ferrite during cooling. But when thetemperature is elevated to 950 °C or higher, more alloyelements dissolve in the austenite and improve the hard-enability of the steel. Martensite can be therefore form even

Table 1 – Hardness of the as-rolled steels annealed atdifferent temperatures (HV)

Rollingtemperature

(°C)

Annealing temperature

850 °C 900 °C 950 °C

850 °C

900 °C 288 243 231268 240 230

950 °C 278 237 228289 246 238

1050 °C 280 240 219

1150 °C

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when the steel is cooled in sand. As a result, the completelydifferent microstructure, ferrite and martensite, cause thesharp increase of the hardness from 900 °C to 950 °C.

3.3. Microstructure Evolution During Annealing

All the as-rolled steels were annealed at different tempera-tures in order to find the optimum combination of hot rollingand annealing parameters. The bestmicrostructure, with idealaverage size and uniform distribution of the carbides, isshown in Fig. 9a. It was obtained by the as-sprayed steel rolledat 1050 °C and then annealed at 900 °C, whereas all other as-annealed steels showed inhomogeneous distribution of car-bide size to some extent. It can be seen from Fig. 9 that theaverage carbide size in the spray formed steel is even smallerthan the equivalent in the commercial PM V4 steel. TEMobservation shows that the as-annealed steel consists offerrite and homogeneously distributed granular MC and M7C3

carbides. Table 1 compares Vickers hardness of all the as-annealed steels. The hardness of those steels which annealedat 900 °C is close to that of the commercial PM V4 steel (HV:243). The results give strong evidence that the high alloyed V4steel can be produced by spray forming.

Annealing of the as-rolled steels modifies the microstruc-ture in three ways: (1) recrystallization of the matrix;(2) transform the matrix into soft ferrite; and (3) coarseningand further spheroidization of carbides. Microstructural

Fig. 9 –Comparison of V4 annealed microstructures: (a) sprayforming, (b) powder metallurgy.

evolution processes during annealing varied according to thevariations of the as rolled steels' matrix. As to thosematerials,which were rolled at or higher than 950 °C, their matrix iscomposed of martensite. Similar to the reheating process ofthe as-sprayed material, when they were heated to elevatedtemperature the martensite will decompose into ferrite andM3C carbides quickly. But this does not happen during thereheating process to the steels rolled at 850 °C and 900 °C astheir matrix is ferrite. In chromium steel the M3C carbidesusually have stronger tendency to dissolve in matrix thanM7C3 carbides because the later is thermodynamically morestable [24]. Upon heating up to austenitizing temperature,there is a likelihood of in-situ transformation from undis-solvedM3C carbides toM7C3 carbides. The transformationmaybegin with the dissolution of M3C followed by in situnucleation and growth of M7C3 carbides [24], or may proceedby the eutectoid reaction of M3C⇒γ+M7C3 [22].

It is well known that a distribution of immobile particles ina solid matrix tends to lower its interfacial free energy bytransport of matter from the small to the larger members,thereby diminishing the total particles number but increas-ing the average particle size. This coarsening process also isknown as Ostwald ripening [25,26]. As to the MC and M7C3

carbides, V-rich MC carbides have little tendency to coarsenbut Cr-base M7C3 carbides are prone to grow up at elevatedtemperature. In addition, the growth of dispersed particles in asaturated solution by a diffusion mechanism depends onmany variables. As has reported by Greenwood [27], the mainfactors promoting stability (slow growth-rate) are particles oflarge mean radius, high density, and low molecular weight,with low values of the solute diffusion coefficient. Thecarbides' size and distribution of the as-annealed steels areinfluenced by the microstructures of the as-rolled steels. Ingeneral, attractive annealed microstructure can be obtainedfrom the as-rolled steels with fine, homogeneously distributedcarbides. Two representative as-annealed steels, which wereobtained by rolling at 850 °C and 1050 °C respectively and thenall annealed at 900 °C, could be taken to explain the difference.

As shown in Fig. 4a, the radii of the M7C3 carbides variedgreatly. During the isothermal holding process during anneal-ing, the crucial effect of interfacial curvature on energy, andhence on particle size (Gibbs–Thomson effect), causes largerM7C3 carbides to grow at the expense of the smaller parti-cles, so that smaller particles dissolve and eventually vanish.Thus, it is hard to obtain ideal microstructure by the initial

1014 M A T E R I A L S C H A R A C T E R I Z A T I O N 5 9 ( 2 0 0 8 ) 1 0 0 7 – 1 0 1 4

incongruous size and distribution of the carbides. On thecontrary, as to the specimens rolled at 1050 °C, the carbidesare highly dispersed and have a uniform size distribution.Then they have higher stability and grow slowly according toGreenwood's conclusion [27]. No polarizing coarsening behav-ior as discussed above will appear and it is more likely toobtain ideal equal carbides' size and distribution.

4. Conclusions

1. Spray forming resulted in a significant refinement of bothgrain size and carbide size of the high alloyed V4 cold worksteel. The optimized processing parameters for the as-sprayedV4 steel are rolling at 1050 °C and then annealing at900 °C. No interconnected eutectic carbide structure wasobserved and fine spheroidal MC and M7C3 carbides wereuniformly distributed in the matrix.

2. The hot rolling temperature is the key factor in controllingthe evolution of type, morphology and distribution ofcarbides, as well as in controlling the matrix microstruc-ture of the as-rolled steel. The matrix consists of ferritewhen the materials rolled at 850 °C and 900 °C. But whenthe steels were rolled at or higher than 950 °C, martensitewill be obtained because of enhanced hardenability bydissolving of alloy elements in austenite.

3. M3C is the only transient carbide precipitate below the A1

temperature during the reheating of the as-sprayed steel.Irregular M7C3 carbides are preferentially precipitated atthe grain boundaries during cooling after the steel wasrolled at 1150 °C.

4. The size and distribution of carbides in the as-annealedsteels are influenced by themicrostructures of the as-rolledsteels. Attractive annealed microstructure can be obtainedfrom the as-rolled steels with fine, homogeneously distrib-uted carbides.

Acknowledgements

This research was financially supported by Shanghai BaosteelTechnology Center and Key Laboratory for High TemperatureMaterials and Tests of Ministry of Education, Shanghai JiaoTong University. The author would like to thank DrWei Wangfor many useful suggestions and comments.

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