Kinetics of Secondary CarbidesPrecipitation in a High-Chromium White Iron

A Bedolla-Jacuinde L Arias and B Hernaacutendez

(Submitted 26 March 2003)

This work analyzes the rate of secondary carbides precipitation during the destabilization heat treatmentof a 17Cr white iron The experimental iron was characterized in the as-cast conditions to have com-parable parameters with the heat treated samples Destabilization heat treatments were undertaken attemperatures of 900 1000 and 1150 degC for between 5 min and 8 h each sample was water quenchedimmediately after being taken out of the furnace Characterization was carried out by optical and electronmicroscopy image analysis and energy dispersive spectroscopy (EDS) microanalysis hardness and mi-crohardness were also evaluated It was found that most of the secondary carbides that precipitate (be-tween 2-30 of the matrix volume) precipitated in less than 2 h for the lowest destabilization temperature(900 degC) The secondary carbides volume fraction was found to increase for lower destabilization tem-peratures and large soaking times A very low carbide precipitation along with a stabilization of theaustenite phase occurred for heat treatments at 1150 degC The results are discussed in terms of the solubilityof chromium and carbon in the austenite phase at the different treatment temperatures

Keywords cast iron high-chromium secondary carbides

1 Introduction

High-chromium white irons are ferrous alloys containingbetween 11-30Cr and 18-36C It is also common to findsome alloying elements such as molybdenum manganese cop-per and nickel The typical as-cast microstructure of thesealloys consists of primary andor eutectic carbides (M7C3) in ametastable austenitic matrix[ 1 ] The hard eutectic carbides aremainly responsible for the good abrasion resistance of thesealloys Therefore these alloys have been widely used for ap-plications where stability in severe environments is the mainrequirement such as the mineral processing industry cementand paper production and the steel manufacturing industry [ 2 ]

Both carbides and matrix contribute to wear resistance andfracture toughness Eutectic carbides have a hexagonal crystal-line structure and solidify as colonies of plates or bars (eutecticgrains) Once solidified carbide morphology is relatively im-mune to a subsequent modification by heat treatment The as-cast austenitic structure in contrast is readily heat treated fordestabilization and for obtaining small secondary carbides pre-cipitated in a matrix that is a mixture of martensite and retainedaustenite[ 3 ]

The commonly applied heat treatment for maximumstrengthening is denoted destabilization which consists ofheating the alloy at temperatures within 800-1100 degC soakingat these temperatures followed by an air quenching at roomtemperature During soaking carbon and chromium from thematrix react to form small-distributed carbide particles The

new chromium- and particularly carbon-depleted matrixreadily transforms to martensite during the subsequent coolingdown Therefore the final structure after destabilization con-sists of M7C3 eutectic carbides and a martensitic matrix withsecondary carbides distributed in it[1 ] Secondary carbide pre-cipitation and the transformed martensitic matrix promote aneven more brittle alloy However a martensitic matrix is rec-ommended to obtain greater hardness and better wear proper-ties but this increase in hardness affects fracture toughnessSuch a phenomenon brings attention to the importance of op-eration factors of the destabilization treatment to be studied todetermine the appropriate temperature and soaking time forsuch a high level of hardness and fracture toughness

Although relatively little systematic research on the second-ary carbide precipitation phenomena and their effect on the ironmatrix have been developed some works regarding this fieldare relevant[ 4 -9 ]

2 Experimental Procedure

The experimental white iron for the present work was madein a vacuum induction furnace by using high purity raw mate-rials therefore obtaining an impurities-free well-controlledchemical composition alloy The melt was poured at 1450 degCinto a metallic mold placed within the vacuum chamber anda 6 cm diameter and 10 kg bar ingot was obtained The ingotwas sectioned with an alumina abrasive cutting wheel in aDISCOTOM (Buehler Inc Waukegan IL) by copious wateramounts The cutting process manually controlled was asslow as possible to avoid overheating the sample and rectan-gular 12 times 12 times 4 mm samples were obtained

Destabilization heat treatments were undertaken in a fur-nace using electric heating elements at an air atmosphere Acomplete series of 10 samples was placed into the furnace fora heat treatment once the desired temperature was reached thesamples were periodically taken out and water quenched toensure that the secondary carbides had precipitated during the

A Bedolla-Jacuinde L Arias and B Hernaacutendez Instituto de In-vestigaciones Metaluacutergicas Universidad Michoacana de San Nicolaacutesde Hidalgo Morelia Michoacan CP 58000 Meacutexico Contact e-mailabedollazeusumichmx

JMEPEG (2003) 12371-382 copyASM International

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash371

soaking time only Heat treatments were undertaken at 9001000 and 1150 degC for a range of 5 min and 8 h

Materials characterization in both as-cast and heat treatedconditions was undertaken by optical and electronic micros-copy energy dispersive spectroscopy (EDS) microanalysis x-ray diffraction (XRD) and image analysis

Samples were prepared for metallography in the traditionalway grinding on abrasive paper of different mesh sizes (80180 320 400 600 and 1200) Samples were then polished onnylon cloth with 6 mm diamond paste as the abrasive and thenwith 1 mm paste Etching was carried out with two differentsolutions depending on the purpose Villelarsquos etchant (5 mlHCl and 1 g picric acid in 100 ml ethanol) for 30-60 s to revealthe microstructure and a solution of 50 ml FeCl3 plus 20 mlHCl in 930 ml ethanol where the samples were immersed forabout 3 h for a deep etching This latter etching readily re-

moves part of the matrix on the surface allowing the nakedcarbides to be observed

As-cast eutectic carbide volume fraction was measured byimage analysis on digitized micrographs obtained at 250X on aNikon EPHIPHOT 3000 inverted metallurgical microscope(Nikon Melville NY) For this purpose the samples weredeeply etched and 20 micrographs were processed by the Sig-mascan (Jandel Scientific San Rafael CA) V5 software on a PC

XRD studies were also undertaken to identify the differentphases present in the alloy both before and after heat treat-ment Retained austenite quantification was also calculated byXRD data by the technique described by Kim[1 0 ] XRD testswere carried out by using Cu-Ka radiation in a 2u range of30-130deg Further characterization was undertaken on a JEOL(JEOL Peabody MA) 6400 SEM at 20 kV for imaging andmicroanalysis and a PHILIPS TECNAI TEM (Philips Eind-hoven The Netherlands) Secondary carbides volume fractionand size was obtained by point counting measurements onSEM micrographs by using a transparent 250 point grid

Bulk hardness and microhardness of the matrix were under-taken on metallographic samples Twenty tests for each sam-ple were carried out by a diamond indenter and a 50 g load for15 s for Vickers microhardness (HV5 0) Bulk hardness wasundertaken in the Rockwell C scale

3 Results and DiscussionTable 1 shows the chemical composition of the experimen-

tal high-chromium white iron 169Cr and 258C one of

Table 1 Chemical Composition of the ExperimentalWhite Iron wt

CarbonChro-mium

Molyb-denum Nickel

Vana-dium

Sul-phur

Phos-phorus

258 1690 198 180 198 0007 0006

Fig 1 Microstructure of the as-cast experimental iron showing pro-eutectic austenite areas (g) and the austenite-M7C3 eutectic scanningelectron microscopy (SEM) Villelarsquos etching 30 s

Fig 2 XRD pattern of the as-cast alloy showing the presence of austenite g eutectic carbides M7C3 and carbides type M2C

372mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

the most commercial high-chromium iron alloys[ 1 ] Molybde-num and nickel contents are close to 2 enough to ensure agood hardenability Also 2 vanadium was added to this ironfor further hardness Higher as-cast hardness values in thesealloys have been reported when adding 2 vanadium as thiselement partitions almost totally to the eutectic carbide makingthe composition (FeCrVMo)7C3 Vanadium increases thehardness of the eutectic carbide and also the carbide volumefraction which in turn increases the hardness of the alloy[ 1 1 ]

31 As-Cast Structure

Figure 1 shows the typical as-cast structure of the alloywhere the austenitic matrix and eutectic carbides can be ob-served Because the alloy is hypoeutectic austenite dendritesare the first to solidify followed by the eutectic Only eutecticcarbide and matrix phases can be observed from this picturehowever two additional phases are present as shown by XRDtraces from Fig 2 These additional phases present in smallamounts are martensite and a molybdenum rich carbide typeM2C Micrographs at higher magnifications show the presenceof martensite at the matrixcarbide interface in the eutecticzones (Fig 3) It has been widely reported that martensiteforms due to the carbon and chromium depletion by diffusiontowards the carbide during the solidification and the subse-quent cooling to room temperature[ 3 1 2 1 3 ] The martensiteamount as measured by XRD is about 6 The presence ofthe M2C carbide was also detected by EDS in the SEM Figure4 shows the morphology and composition of this carbide beinga molybdenum rich carbide its composition is of the type M2Ctypical for molybdenum carbide The main role of molybde-num in these irons is to improve hardenability however not allthe molybdenum contributes to this goal The presence of car-bon leads to the formation of the molybdenum rich carbide andpart of the molybdenum in these irons forms a eutectic carbidetype M2C at the end of the solidification stage[1 4 -1 9 ]

The elements distribution was analyzed by EDS in a randomarea of the as-cast microstructure Figure 5(a) shows an SEMmicrograph where some microanalyses were undertaken on the

matrix (Fig 5b) on the carbide (Fig 5c) and on a scan alongthe line shown on the micrograph in Fig 5(a) across an aus-tenite dendrite arm (Fig 5d)

Eutectic carbides are of the type M7C3 as revealed by XRD(Fig 2) and as reported for chromium contents of 17[ 7 ] Theelements present in the carbide phase can be seen from Fig5(b) despite being a qualitative analysis the intensity of thepeaks indicates that chromium and iron are present in higheramounts in the carbide The carbide contains also vanadiumand maybe small amounts of molybdenum which could not bedetected However the most important thing to highlight is thepresence of vanadium in the eutectic carbide XRD traces didnot show the presence of vanadium carbide Vanadium is acarbide-forming element of the type VC however the solidi-fication temperature for this carbide is below the solidificationtemperature for the alloy Therefore vanadium carbide forma-

Fig 3 SEM micrograph showing the presence of martensite at thematrixcarbide interface Villelarsquos etching 45 s

Fig 4 (a) SEM micrograph of a deep-etched sample showing amolybdenum rich carbide type M2C formed at the last solidificationstage etching solution 50 ml FeCl3 plus 20 ml HCl in 930 ml ethanolfor about 3 h (b) EDS of the M2C carbide in (a)

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash373

tion in this kind of irons is not likely These vanadium carbidesare believed to form only when the vanadium content exceeds3 as reported by Radulovic et al[ 2 0 ] Therefore the majorproportion of vanadium partitions to the carbide and the restremains dissolved in matrix as shown in the EDS in Fig 5(c)A minor proportion of chromium related to iron can be ob-served from this spectrum contrary to the findings in the car-bide microanalysis where the chromiumiron ratio is close to 1A high amount of molybdenum is also observed in matrix Thiselement remains in matrix in major proportion contributing thisway to improve hardenability From the same spectrum in Fig5(c) to the right of the last iron peak a small peak correspond-ing to nickel can be observed Nickel remains totally in matrixcontributing synergistically with molybdenum to hardenabilityVanadium can also be observed in matrix but in much smalleramounts than in the eutectic carbide The scan on the line inFig 5(a) shows the presence of chromium carbon and vana-dium in matrix a higher proportion of these elements is ob-

served at the edges of the graph These ends include the carbidephase as seen from the line The presence of these elements inmatrix is the source for secondary carbides precipitation duringthe heat treatment of austenite destabilization

Eutectic carbide is a very important phase in the propertiesof these alloys It has been mentioned that eutectic carbide is a3D network or an interconnected skeleton into the materialhowever the level of interconnection is much less than inunalloyed white irons This is due to the differences in prefer-ential growing directions for cementita and the carbide M7C3 M7C3 carbide has a hexagonal structure and a preferentialgrowing direction lt0001gt [ 7 ] therefore the typical morphol-ogy of these carbides is bars or plates (Fig 6) On a polishedsurface these bars could be perpendicular to the surface ap-pearing to be isolated particles under the microscope howeverdeep-etching techniques on the SEM make it possible to ob-serve that many bars may be separated at the surface but con-nected at the base

Fig 5 (a) SEM micrograph of the structure of the as-cast experimental iron (b) EDS corresponding to the eutectic carbide phase (c) EDScorresponding to the proeutectic matrix and (d) vanadium chromium and carbon profile of matrix along the line shown in (a)

374mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

32 Structure After Destabilization Heat Treatment

After destabilization heat treatment some phenomena tookplace which altered the structure of the material Such phe-nomena are a function of temperature and soaking time as inany heat treatment The phases present in the as-cast alloy are

the same as in the heat treated samples but the proportion haschanged The XRD pattern from Fig 7 is a representative proofof what happens after heat treatment The present phases areaustenite martensite M7C3 carbide and M2C carbide Themain difference that can be observed from this XRD comparedwith that from Fig 2 (as-cast material) is the higher intensityfor the martensite peaks (a phase) Therefore the amount ofmartensite is higher in the heat treated samples except in thosetreated at 1150 degC where the austenite phase was over-stabilized (see Section 35) According to the XRD traces sec-ondary carbides precipitated in matrix are of the type M7C3 thesame as the eutectic because no other type was detected Thekind of secondary carbides precipitating depend on the alloycomposition and the destabilization temperature The presenceof secondary carbides of the type M2 3C6 has been observed asfine interconnected bars in irons with very high chromiumamounts (gt25)[ 5 7 ] On the other hand secondary carbides ofthe type M7C3 were observed as agglomerated plates in ironswith chromium contents between 15-20[ 5 ] These resultsagree with the findings in the present work

33 Secondary Carbide Volume Fraction

Figure 8 shows the obtained results for the volume fractionof secondary carbides precipitated as a function of the soakingtime for each of the studied temperatures the horizontal axiswhich corresponds to time is plotted in logarithmic scale At900 degC precipitation was observed to start at about 10 min anda volume fraction of 2 was measured after 8 h of treatmentthe volume fraction was about 27 At 1000 degC the precipi-tation process was faster 2 at 5 min up to 21 at 10 minfrom 10 min to 8 h the volume fraction was about 22 Thetotal volume fraction precipitated within the first 10 min af-terwards the coarsening of some carbides and the dissolution ofothers occurred

At 1150 degC a volume fraction of 15 occurred at only 5min subsequently the volume fraction diminished to 2 at 1 hsoaking Figures 9 10 and 11 show sequences of SEM micro-graphs of some samples at different soaking times and at thetemperatures of 900 1000 and 1150 degC respectively Theseseries of micrographs represent visually what is plotted in Fig8 At 900 degC the diffusion rate of chromium and carbon lead tovolume fraction of secondary carbides of 27 in 8 h Howevermost of the precipitation took place within the first 2 h whichis not clear because the scale of time is plotted logarithmicallyHigher temperatures such as 1000 degC accelerated the precipi-tation process due to the higher diffusion rate of the elementsthat form the carbides chromium and carbon At this tempera-ture almost all of the volume of secondary carbides precipitatedwithin the first 10 min reaching a total volume fraction of22 In contrast at 1150 degC where diffusion rates for chro-mium and carbon are higher a carbide precipitation of 15was observed within the first 5 min but it started to diminishafterwards It is likely that the initial amount of carbides hadprecipitated during heating to the soaking temperature andthese carbides dissolved during soaking when the equilibriumcomposition of chromium and carbon in austenite was reachedAccording to the Fe-C phase diagram austenite can dissolvehigher carbon contents as the temperature increases to the eu-tectic

Fig 6 SEM micrographs with the 3-D structure of eutectic carbidesshowing the real morphology deep-etching with a solution of 50 mlFeCl3 plus 20 ml HCl in 930 ml ethanol for 3 h

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash375

As mentioned at 900 degC a higher secondary carbide volumefraction was observed at lower temperatures the ability ofaustenite to dissolve carbon is lower Therefore there is ahigher amount of carbon to precipitate which in turn is ob-served as a higher carbide volume fraction At higher tempera-tures austenite can retain carbon and chromium in higheramounts so there is very low raw material available to pre-cipitate carbides[ 5 8 9 ] and the final volume fraction is muchlower It is difficult to explain the diminishing of carbide vol-ume fraction observed at 1150 degC it is likely the equilibriumconcentration of carbon in austenite will be higher than theactual amount of this element in austenite so that during soak-ing precipitation is void Perhaps the volume fraction in thefirst sample (15 at 2 min) precipitated during heating to thesoaking temperature and these 2 min at 1150 degC were notenough for austenite to reach its equilibrium composition and

the secondary carbides did not dissolve Samples treated forlonger times (10 and 60 min) show a precipitated volume frac-tion of 2-5 perhaps the equilibrium volume in the alloy is at1150 degC

From micrographs in Fig 9 10 and 11 and particularlyfrom Fig 12 the morphology of secondary carbides can beobserved Secondary carbides are bars connected within thebulk matrix It has been reported[5 ] in white irons of similarcomposition that secondary carbides precipitated after destabi-lization treatments are of the type M7C3 having a hexagonalstructure and a preferred growing direction lt0001gt Figure 13shows some TEM micrographs of secondary carbides and adiffraction pattern of these particles which agree well with thefindings in Ref 5

34 Precipitated Carbides Size and Number

Figure 14 shows the particle size as a function of the soak-ing time these measurements are taken as the diameter of thecarbide bar It is well known that precipitation processes arefavored by temperature because temperature accelerates dif-fusion of the elements dissolved in matrix A higher diffusionrate of elements is observed for precipitation processes under-taken at high temperatures Because precipitation takes placeby nucleation and growth phenomena particles of a certain sizeare formed which depend on the treatment temperature In thecurrent study when the iron was treated at 900 degC the second-ary carbides diameter was about 120 nm within the first 2 hafterwards they coarsened up to 300 nm for 8 h soaking timeAt temperatures of 1000 degC the initial diameter is close to 200nm and gradually coarsens up to 250 nm for 8 h soaking timeAccording to this sequence it would be expected that the di-ameter of secondary carbides formed at 1150 degC to be higherhowever their size is only about 160 nm Because the volumeof material that precipitates in equilibrium as secondary car-

Fig 7 XRD pattern of the simple destabilized at 1000 degC for 1 h showing the detected phases austenite martensite M2C and M7C3

Fig 8 Volume of secondary carbides precipitated as a function ofthe soaking time during destabilization for each of the studied tem-peratures

376mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

bides at this temperature is almost void there is not enoughmaterial to precipitate coarse carbides

What is important to highlight is the precipitation behaviorat 900 and 1000 degC (Fig 15) This figure shows the numberprecipitated particles per mm2 At 900 degC the number of car-bides went from 2 particles per mm2 for 10 min soaking time to14 particles for 90 min Within the next 3 h the number ofcarbides remained at 14 and afterwards decreased to 8 particlesper mm2 for 8 h soaking time This finding indicates that theprecipitation process occurred during the first 90 min duringthe next 3 h there was no new precipitation just coarsening ofthe existing particles Finally during the next 4 h dissolutionof some carbides occurred contributing to the coarsening ofsome others thermodynamically more stable Similar phenom-ena during destabilization heat treatments have been reportedby Kuwano et al[ 8 ] All these phenomena are better observedfrom SEM micrographs (Fig 9 10 and 11) Note that from thecurrent study the secondary carbides precipitation was ob-served homogeneously in matrix According to Powell and

Bee[ 9 ] and Bee et al[ 2 1 ] secondary carbides precipitate at slipbands and sub-grain boundaries within austenite Such slipbands and sub-grain boundaries may form due to stresses gen-erated by differences in the thermal expansion coefficients ofcarbides and matrix On the contrary Kuwano et al[8 ] reportedthat the secondary carbides precipitation progressed from theperiphery adjacent to the eutectic carbides towards the centerof the dendrite in low carbon irons In high carbon alloys thecontrary occurredmdashprecipitation progressed from the center ofthe dendrite towards the periphery They attributed this phe-nomenon to chromium and carbon micro-segregation towardsthe periphery of the dendrites

Just a few works on the kinetics of secondary carbide pre-cipitation have been published Powell and Laird[ 5 ] suggestedthat precipitation occurs during the first 25 min at the soakingtemperature and keeping the alloy for longer periods of timeonly causes coarsening of carbides Kuwano et al[ 8 ] felt thatprecipitation occurs within very short periods of time such as 1min at temperatures between 950-1000 degC which disagree with

Fig 9 Series of SEM micrographs showing secondary carbide precipitation during destabilization heat treatment at 900 degC of the experimentalwhite iron (a) 20 min (b) 1 h (c) 2 h and (d) 8 h Villelarsquos etching 45 s Matrix is composed of martensite retained austenite and secondarycarbides

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash377

the findings from this work What has been observed andwidely accepted by most researchers is that once carbides haveformed longer periods of time at the treatment temperatureincrease the size and volume fraction of them Furthermorestill longer periods of time at the soaking temperature cause thecoarsening of carbides at the expense of othersrsquo dissolution(Ostwald ripening process) reducing the number of carbides[ 8 ]

This effect can be observed in Fig 15

35 Retained Austenite After Destabilization

Figure 16 shows the results of retained austenite as mea-sured by XRD As discussed above secondary carbide precipi-tation depletes the austenitic matrix in chromium and car-bon[ 4 1 0 ] Such depletion produces an increase in MS and it ismore likely to obtain martensitic structures Such an effect isobserved in Fig 16 at 900 degC and long soaking times a re-duction in the retained austenite content is observed from 48at 10 min to just 8 for 8 h soaking time This means an

increase in the martensite volume fraction At this temperatureand long soaking times the volume fraction of secondary car-bides was 27 The chromium- and carbon-depleted matrixtransformed almost totally to martensite during the subsequentcooling Something similar occurred during destabilization at1000 degC residual austenite decreased from 52 for 5 minsoaking time to 15 for 8 h In this case the volume fractionof secondary carbides was less than in the alloy treated at 900degC a clear indication is that chromium and carbon dissolved inmatrix are higher in the alloy treated at 1000 degC and suchelements contribute to lower the MS temperature Thereforethe alloy treated at 1000 degC showed a higher amount of residualaustenite These results are in agreement with the findings ofsome authors[ 5 8 9 ] in high- and low-chromium white irons Onthe contrary retained austenite increased with soaking time forthe iron treated at 1150 degC For a 5 min soaking time retainedaustenite was 25 and for 1 h the whole matrix became aus-tenitic after heat treatment In this case the high temperaturesallowed carbon and chromium to dissolve completely in matrixand inhibited the driving force for secondary carbides precipi-

Fig 10 Series of SEM micrographs showing secondary carbide precipitation during destabilization heat treatment at 1000 degC of the experimentalwhite iron (a) 5 min (b) 10 min (c) 1 h and (d) 6 h Villelarsquos etching 45 s Matrix is composed of martensite retained austenite and secondarycarbides

378mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

tation Chromium- and carbon-rich austenite is stable underthese conditions and remains stable during cooling to roomtemperature[1 2 2 2 3 ]

36 Microhardness of Matrix and Bulk Hardness of theAlloy After Destabilization

Figure 17 shows the microhardness values obtained for theironrsquos matrix after each heat treatment At 900 degC for 10 minmicrohardness was 800HV5 0 and increased to a maximum of1300HV5 0 after 2-3 h soaking time afterwards it decreased to900HV5 0 for a destabilization of 8 h The increase in micro-hardness of the matrix is influenced by two factors as thesoaking time increased the precipitation of secondary carbidesthat strengthens the matrix by particle dispersion together withthe increase in martensite volume fraction However dimin-ishing microhardness after 3 h soaking time is attributed to thesecondary carbide coarsening At this stage longer times pro-mote a major diffusion of elements favoring coarsening ofsome carbides at the expense of othersrsquo dissolution such aphenomenon leads to a diminishing number of particles in thematrix (Fig 15) The final structure of the matrix is thereforehighly martensitic with relatively few coarse secondary car-bides For the structures of the irons with maximum hardnessa highly martensitic matrix was also observed but it containeda higher number of fine well-dispersed secondary carbides

A similar effect occurred for the iron treated at 1000 degChowever maximum hardness (1159HV5 0 ) was obtained atshorter times (30-60 min) Higher destabilization temperaturesallowed higher diffusion leading to a reduced amount ofcoarse nuclei (Fig 14 and 15) Under this situation thestrengthening due to dispersed particles is lower Furthermoreat these higher temperatures austenite can dissolve higher car-bon contents so that the amount of secondary carbides is alsominor The above phenomena lead to a higher retained austen-ite content in the final structure which reduced matrix hard-ness

In contrast for the iron treated at 1150 degC hardness of thematrix diminished since the beginning For this case a stabi-lized austenitic matrix was obtained Such a matrix was solidsolution strengthened by chromium carbon nickel molybde-num etc but not by particle dispersion Secondary precipita-tion occurred only in the early stages of the heat treatment

Regarding bulk hardness the alloy showed a behavior simi-

Fig 11 Series of SEM micrographs showing secondary carbide pre-cipitation during destabilization heat treatment at 1150 degC of the ex-perimental white iron (a) 2 min matrix is composed of martensiteretained austenite and secondary carbides (b) 20 min matrix com-posed of austenite mainly and (c) 1 h matrix is totally austeniticVillelarsquos etching 45 s

Fig 12 SEM micrograph showing the real morphology of the sec-ondary carbides precipitated during destabilization treatments Deep-etching was with a solution of 50 ml FeCl3 plus 20 ml HCl in 930 mlethanol for 15 min

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash379

lar to that for matrix microhardness Figure 18 shows the re-sults for hardness where maximum values can be observed forthe different temperatures For the iron treated at 900 degC maxi-mum hardness was 68 HRC and was obtained for1-2 h soakingtime For the iron treated at 1000 degC maximum hardness was60 HRC obtained for 30 min Finally for the iron treated at1150 degC hardness remains constant (45 HRC) during the first20 min of treatment and decreases afterwards The strengthen-ing of the matrix is the main factor responsible for the modi-fications in the ironrsquos hardness Microhardness of matrix andthe eutectic carbide volume fraction are the factors that con-

tribute to strengthening of the alloys For this case the carbidevolume fraction was constant (24) therefore the strengthen-ing of matrix caused the changes in bulk hardness

These strengthening phenomena are similar to those foundin aging processes of some alloys and are explained in terms ofthe diffusion of elements that participate in the precipitatingphase Such diffusion in turn is a function of treatment tem-perature and soaking time Kuwano et al[ 8 ] Powell andLaird[ 5 ] and Powell and Bee[ 9 ] have found similar behaviorsfor high- and low-chromium irons in studies of precipitationand hardness

Fig 13 Bright field TEM micrographs showing some secondary carbides (SC) in (a) sample treated at 900 degC for 1 h and (b) sample treated at1000 degC for 1 h The selected area diffraction pattern (SADP) on the particle arrowed in (b) indicates the crystallographic nature of these carbidesto be hexagonal M7C3

Fig 14 Diameter of the precipitated particles as a function of thesoaking time during destabilization heat treatment

Fig 15 Number of precipitated particles in a square micron (mm2) asa function of the soaking time during destabilization heat treatment

380mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

4 Conclusions

The structure of the as-cast alloy was composed of 24eutectic carbides about 5 martensite and 70 austenitesmall amounts of the molybdenum-rich carbide Mo2Cwere also present

Vanadium carbide was not detected in the alloy nonethe-less the iron contained 2 vanadium a carbide-promotingelement Vanadium was observed to partition to the car-bide and in minor amounts to the matrix

According to the studied temperatures the secondary car-bide precipitation level decreased with higher tempera-tures being the highest for the samples treated at 900 degC(27 of matrix)

High temperatures like 1150 degC over-stabilized the aus-tenitic matrix which led to a secondary carbides free aus-tenitic matrix after heat treatment

The amount of precipitated particles showed an increasewith time for the samples treated at 900 and 1000 degCreaching a maximum and then decreased due to dissolutionof some carbides and coarsening of some others

Matrix hardness showed a maximum when plotted againstsoaking time due to these precipitation processes and tothe transformation of austenite to martensite Maximumhardness was obtained for the samples treated for 1-2 h at900 degC where the structure showed a high volume fractionof martensite and a high amount of fine well-dispersedsecondary carbide particles

Acknowledgments

The authors are grateful to the University of Michoacan forfunding this project

References

1 CP Tabrett IR Sare and MR Gomashchi ldquoMicrostructure-Property Relationships in High-Chromium White Iron Alloysrdquo IntMater Rev 1996 41(2) pp 59-82

2 ON Dogan JA Hawk and G Laird II ldquoSolidification Structure andAbrasion Resistance of High Chromium White Ironsrdquo Metall MaterTrans A 1997 28(6) pp 1315-27

3 G Laird II and GLF Powell ldquoSolidification and Solid-State Trans-formation Mechanisms in Si Alloyed High-Chromium White CastIronsrdquo Metall Trans A 1993 24(2) pp 981-88

4 WW Cias ldquoAustenite Transformation Kinetics and Hardenability ofHeat Treated 175Cr White Cast Ironsrdquo Trans AFS 1974 82 pp317-28

5 GLF Powell and G Laird II ldquoStructure Nucleation Growth andMorphology of Secondary Carbides in High Chromium and Cr-NiWhite Cast Ironsrdquo J Mater Sci 1992 27 pp 29-35

6 F Maratray and A Poulalion ldquoAustenite Retention in High-Chromium White Ironsrdquo Trans AFS 1982 90 pp 795-804

7 JTH Pearce ldquoStructure and Wear Performance of Abrasion ResistantChromium White Cast Ironsrdquo Trans AFS 1984 92 pp 599-621

8 M Kuwano K Ogi A Sawamoto and K Matsuda ldquoStudies onPrecipitation Process of Secondary Carbides in High-Chromium CastIronrdquo Trans AFS 1990 98 pp 725-34

9 GLF Powell and JV Bee ldquoSecondary Carbide Precipitation in an 18wtCr-1 wtMo White Ironrdquo J Mater Sci 1996 31 pp 707-11

10 C Kim ldquoX-Ray Method of Measuring Retained Austenite in HeatTreated White Cast Ironsrdquo J Heat Treating 1979 1(2) pp 43-51

11 A Bedolla-Jacuinde ldquoMicrostructure of Vanadium- Niobium- andTitanium-Alloyed High-Chromium White Cast Ironsrdquo Int J CastMetals Res 2001 13(6) pp 343-61

Fig 16 Retained austenite content in iron matrix as a function of thesoaking time during the destabilization heat treatment

Fig 17 Matrix microhardness as a function of the soaking timeduring the destabilization heat treatment

Fig 18 Bulk hardness of the alloy as a function of the soaking timeduring the destabilization heat treatment

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash381

soaking time only Heat treatments were undertaken at 9001000 and 1150 degC for a range of 5 min and 8 h

Materials characterization in both as-cast and heat treatedconditions was undertaken by optical and electronic micros-copy energy dispersive spectroscopy (EDS) microanalysis x-ray diffraction (XRD) and image analysis

Samples were prepared for metallography in the traditionalway grinding on abrasive paper of different mesh sizes (80180 320 400 600 and 1200) Samples were then polished onnylon cloth with 6 mm diamond paste as the abrasive and thenwith 1 mm paste Etching was carried out with two differentsolutions depending on the purpose Villelarsquos etchant (5 mlHCl and 1 g picric acid in 100 ml ethanol) for 30-60 s to revealthe microstructure and a solution of 50 ml FeCl3 plus 20 mlHCl in 930 ml ethanol where the samples were immersed forabout 3 h for a deep etching This latter etching readily re-

moves part of the matrix on the surface allowing the nakedcarbides to be observed

As-cast eutectic carbide volume fraction was measured byimage analysis on digitized micrographs obtained at 250X on aNikon EPHIPHOT 3000 inverted metallurgical microscope(Nikon Melville NY) For this purpose the samples weredeeply etched and 20 micrographs were processed by the Sig-mascan (Jandel Scientific San Rafael CA) V5 software on a PC

XRD studies were also undertaken to identify the differentphases present in the alloy both before and after heat treat-ment Retained austenite quantification was also calculated byXRD data by the technique described by Kim[1 0 ] XRD testswere carried out by using Cu-Ka radiation in a 2u range of30-130deg Further characterization was undertaken on a JEOL(JEOL Peabody MA) 6400 SEM at 20 kV for imaging andmicroanalysis and a PHILIPS TECNAI TEM (Philips Eind-hoven The Netherlands) Secondary carbides volume fractionand size was obtained by point counting measurements onSEM micrographs by using a transparent 250 point grid

Bulk hardness and microhardness of the matrix were under-taken on metallographic samples Twenty tests for each sam-ple were carried out by a diamond indenter and a 50 g load for15 s for Vickers microhardness (HV5 0) Bulk hardness wasundertaken in the Rockwell C scale

3 Results and DiscussionTable 1 shows the chemical composition of the experimen-

tal high-chromium white iron 169Cr and 258C one of

Table 1 Chemical Composition of the ExperimentalWhite Iron wt

CarbonChro-mium

Molyb-denum Nickel

Vana-dium

Sul-phur

Phos-phorus

258 1690 198 180 198 0007 0006

Fig 1 Microstructure of the as-cast experimental iron showing pro-eutectic austenite areas (g) and the austenite-M7C3 eutectic scanningelectron microscopy (SEM) Villelarsquos etching 30 s

Fig 2 XRD pattern of the as-cast alloy showing the presence of austenite g eutectic carbides M7C3 and carbides type M2C

372mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

the most commercial high-chromium iron alloys[ 1 ] Molybde-num and nickel contents are close to 2 enough to ensure agood hardenability Also 2 vanadium was added to this ironfor further hardness Higher as-cast hardness values in thesealloys have been reported when adding 2 vanadium as thiselement partitions almost totally to the eutectic carbide makingthe composition (FeCrVMo)7C3 Vanadium increases thehardness of the eutectic carbide and also the carbide volumefraction which in turn increases the hardness of the alloy[ 1 1 ]

31 As-Cast Structure

Figure 1 shows the typical as-cast structure of the alloywhere the austenitic matrix and eutectic carbides can be ob-served Because the alloy is hypoeutectic austenite dendritesare the first to solidify followed by the eutectic Only eutecticcarbide and matrix phases can be observed from this picturehowever two additional phases are present as shown by XRDtraces from Fig 2 These additional phases present in smallamounts are martensite and a molybdenum rich carbide typeM2C Micrographs at higher magnifications show the presenceof martensite at the matrixcarbide interface in the eutecticzones (Fig 3) It has been widely reported that martensiteforms due to the carbon and chromium depletion by diffusiontowards the carbide during the solidification and the subse-quent cooling to room temperature[ 3 1 2 1 3 ] The martensiteamount as measured by XRD is about 6 The presence ofthe M2C carbide was also detected by EDS in the SEM Figure4 shows the morphology and composition of this carbide beinga molybdenum rich carbide its composition is of the type M2Ctypical for molybdenum carbide The main role of molybde-num in these irons is to improve hardenability however not allthe molybdenum contributes to this goal The presence of car-bon leads to the formation of the molybdenum rich carbide andpart of the molybdenum in these irons forms a eutectic carbidetype M2C at the end of the solidification stage[1 4 -1 9 ]

The elements distribution was analyzed by EDS in a randomarea of the as-cast microstructure Figure 5(a) shows an SEMmicrograph where some microanalyses were undertaken on the

matrix (Fig 5b) on the carbide (Fig 5c) and on a scan alongthe line shown on the micrograph in Fig 5(a) across an aus-tenite dendrite arm (Fig 5d)

Eutectic carbides are of the type M7C3 as revealed by XRD(Fig 2) and as reported for chromium contents of 17[ 7 ] Theelements present in the carbide phase can be seen from Fig5(b) despite being a qualitative analysis the intensity of thepeaks indicates that chromium and iron are present in higheramounts in the carbide The carbide contains also vanadiumand maybe small amounts of molybdenum which could not bedetected However the most important thing to highlight is thepresence of vanadium in the eutectic carbide XRD traces didnot show the presence of vanadium carbide Vanadium is acarbide-forming element of the type VC however the solidi-fication temperature for this carbide is below the solidificationtemperature for the alloy Therefore vanadium carbide forma-

Fig 3 SEM micrograph showing the presence of martensite at thematrixcarbide interface Villelarsquos etching 45 s

Fig 4 (a) SEM micrograph of a deep-etched sample showing amolybdenum rich carbide type M2C formed at the last solidificationstage etching solution 50 ml FeCl3 plus 20 ml HCl in 930 ml ethanolfor about 3 h (b) EDS of the M2C carbide in (a)

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash373

tion in this kind of irons is not likely These vanadium carbidesare believed to form only when the vanadium content exceeds3 as reported by Radulovic et al[ 2 0 ] Therefore the majorproportion of vanadium partitions to the carbide and the restremains dissolved in matrix as shown in the EDS in Fig 5(c)A minor proportion of chromium related to iron can be ob-served from this spectrum contrary to the findings in the car-bide microanalysis where the chromiumiron ratio is close to 1A high amount of molybdenum is also observed in matrix Thiselement remains in matrix in major proportion contributing thisway to improve hardenability From the same spectrum in Fig5(c) to the right of the last iron peak a small peak correspond-ing to nickel can be observed Nickel remains totally in matrixcontributing synergistically with molybdenum to hardenabilityVanadium can also be observed in matrix but in much smalleramounts than in the eutectic carbide The scan on the line inFig 5(a) shows the presence of chromium carbon and vana-dium in matrix a higher proportion of these elements is ob-

served at the edges of the graph These ends include the carbidephase as seen from the line The presence of these elements inmatrix is the source for secondary carbides precipitation duringthe heat treatment of austenite destabilization

Eutectic carbide is a very important phase in the propertiesof these alloys It has been mentioned that eutectic carbide is a3D network or an interconnected skeleton into the materialhowever the level of interconnection is much less than inunalloyed white irons This is due to the differences in prefer-ential growing directions for cementita and the carbide M7C3 M7C3 carbide has a hexagonal structure and a preferentialgrowing direction lt0001gt [ 7 ] therefore the typical morphol-ogy of these carbides is bars or plates (Fig 6) On a polishedsurface these bars could be perpendicular to the surface ap-pearing to be isolated particles under the microscope howeverdeep-etching techniques on the SEM make it possible to ob-serve that many bars may be separated at the surface but con-nected at the base

Fig 5 (a) SEM micrograph of the structure of the as-cast experimental iron (b) EDS corresponding to the eutectic carbide phase (c) EDScorresponding to the proeutectic matrix and (d) vanadium chromium and carbon profile of matrix along the line shown in (a)

374mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

32 Structure After Destabilization Heat Treatment

After destabilization heat treatment some phenomena tookplace which altered the structure of the material Such phe-nomena are a function of temperature and soaking time as inany heat treatment The phases present in the as-cast alloy are

the same as in the heat treated samples but the proportion haschanged The XRD pattern from Fig 7 is a representative proofof what happens after heat treatment The present phases areaustenite martensite M7C3 carbide and M2C carbide Themain difference that can be observed from this XRD comparedwith that from Fig 2 (as-cast material) is the higher intensityfor the martensite peaks (a phase) Therefore the amount ofmartensite is higher in the heat treated samples except in thosetreated at 1150 degC where the austenite phase was over-stabilized (see Section 35) According to the XRD traces sec-ondary carbides precipitated in matrix are of the type M7C3 thesame as the eutectic because no other type was detected Thekind of secondary carbides precipitating depend on the alloycomposition and the destabilization temperature The presenceof secondary carbides of the type M2 3C6 has been observed asfine interconnected bars in irons with very high chromiumamounts (gt25)[ 5 7 ] On the other hand secondary carbides ofthe type M7C3 were observed as agglomerated plates in ironswith chromium contents between 15-20[ 5 ] These resultsagree with the findings in the present work

33 Secondary Carbide Volume Fraction

Figure 8 shows the obtained results for the volume fractionof secondary carbides precipitated as a function of the soakingtime for each of the studied temperatures the horizontal axiswhich corresponds to time is plotted in logarithmic scale At900 degC precipitation was observed to start at about 10 min anda volume fraction of 2 was measured after 8 h of treatmentthe volume fraction was about 27 At 1000 degC the precipi-tation process was faster 2 at 5 min up to 21 at 10 minfrom 10 min to 8 h the volume fraction was about 22 Thetotal volume fraction precipitated within the first 10 min af-terwards the coarsening of some carbides and the dissolution ofothers occurred

At 1150 degC a volume fraction of 15 occurred at only 5min subsequently the volume fraction diminished to 2 at 1 hsoaking Figures 9 10 and 11 show sequences of SEM micro-graphs of some samples at different soaking times and at thetemperatures of 900 1000 and 1150 degC respectively Theseseries of micrographs represent visually what is plotted in Fig8 At 900 degC the diffusion rate of chromium and carbon lead tovolume fraction of secondary carbides of 27 in 8 h Howevermost of the precipitation took place within the first 2 h whichis not clear because the scale of time is plotted logarithmicallyHigher temperatures such as 1000 degC accelerated the precipi-tation process due to the higher diffusion rate of the elementsthat form the carbides chromium and carbon At this tempera-ture almost all of the volume of secondary carbides precipitatedwithin the first 10 min reaching a total volume fraction of22 In contrast at 1150 degC where diffusion rates for chro-mium and carbon are higher a carbide precipitation of 15was observed within the first 5 min but it started to diminishafterwards It is likely that the initial amount of carbides hadprecipitated during heating to the soaking temperature andthese carbides dissolved during soaking when the equilibriumcomposition of chromium and carbon in austenite was reachedAccording to the Fe-C phase diagram austenite can dissolvehigher carbon contents as the temperature increases to the eu-tectic

Fig 6 SEM micrographs with the 3-D structure of eutectic carbidesshowing the real morphology deep-etching with a solution of 50 mlFeCl3 plus 20 ml HCl in 930 ml ethanol for 3 h

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash375

As mentioned at 900 degC a higher secondary carbide volumefraction was observed at lower temperatures the ability ofaustenite to dissolve carbon is lower Therefore there is ahigher amount of carbon to precipitate which in turn is ob-served as a higher carbide volume fraction At higher tempera-tures austenite can retain carbon and chromium in higheramounts so there is very low raw material available to pre-cipitate carbides[ 5 8 9 ] and the final volume fraction is muchlower It is difficult to explain the diminishing of carbide vol-ume fraction observed at 1150 degC it is likely the equilibriumconcentration of carbon in austenite will be higher than theactual amount of this element in austenite so that during soak-ing precipitation is void Perhaps the volume fraction in thefirst sample (15 at 2 min) precipitated during heating to thesoaking temperature and these 2 min at 1150 degC were notenough for austenite to reach its equilibrium composition and

the secondary carbides did not dissolve Samples treated forlonger times (10 and 60 min) show a precipitated volume frac-tion of 2-5 perhaps the equilibrium volume in the alloy is at1150 degC

From micrographs in Fig 9 10 and 11 and particularlyfrom Fig 12 the morphology of secondary carbides can beobserved Secondary carbides are bars connected within thebulk matrix It has been reported[5 ] in white irons of similarcomposition that secondary carbides precipitated after destabi-lization treatments are of the type M7C3 having a hexagonalstructure and a preferred growing direction lt0001gt Figure 13shows some TEM micrographs of secondary carbides and adiffraction pattern of these particles which agree well with thefindings in Ref 5

34 Precipitated Carbides Size and Number

Figure 14 shows the particle size as a function of the soak-ing time these measurements are taken as the diameter of thecarbide bar It is well known that precipitation processes arefavored by temperature because temperature accelerates dif-fusion of the elements dissolved in matrix A higher diffusionrate of elements is observed for precipitation processes under-taken at high temperatures Because precipitation takes placeby nucleation and growth phenomena particles of a certain sizeare formed which depend on the treatment temperature In thecurrent study when the iron was treated at 900 degC the second-ary carbides diameter was about 120 nm within the first 2 hafterwards they coarsened up to 300 nm for 8 h soaking timeAt temperatures of 1000 degC the initial diameter is close to 200nm and gradually coarsens up to 250 nm for 8 h soaking timeAccording to this sequence it would be expected that the di-ameter of secondary carbides formed at 1150 degC to be higherhowever their size is only about 160 nm Because the volumeof material that precipitates in equilibrium as secondary car-

Fig 7 XRD pattern of the simple destabilized at 1000 degC for 1 h showing the detected phases austenite martensite M2C and M7C3

Fig 8 Volume of secondary carbides precipitated as a function ofthe soaking time during destabilization for each of the studied tem-peratures

376mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

bides at this temperature is almost void there is not enoughmaterial to precipitate coarse carbides

What is important to highlight is the precipitation behaviorat 900 and 1000 degC (Fig 15) This figure shows the numberprecipitated particles per mm2 At 900 degC the number of car-bides went from 2 particles per mm2 for 10 min soaking time to14 particles for 90 min Within the next 3 h the number ofcarbides remained at 14 and afterwards decreased to 8 particlesper mm2 for 8 h soaking time This finding indicates that theprecipitation process occurred during the first 90 min duringthe next 3 h there was no new precipitation just coarsening ofthe existing particles Finally during the next 4 h dissolutionof some carbides occurred contributing to the coarsening ofsome others thermodynamically more stable Similar phenom-ena during destabilization heat treatments have been reportedby Kuwano et al[ 8 ] All these phenomena are better observedfrom SEM micrographs (Fig 9 10 and 11) Note that from thecurrent study the secondary carbides precipitation was ob-served homogeneously in matrix According to Powell and

Bee[ 9 ] and Bee et al[ 2 1 ] secondary carbides precipitate at slipbands and sub-grain boundaries within austenite Such slipbands and sub-grain boundaries may form due to stresses gen-erated by differences in the thermal expansion coefficients ofcarbides and matrix On the contrary Kuwano et al[8 ] reportedthat the secondary carbides precipitation progressed from theperiphery adjacent to the eutectic carbides towards the centerof the dendrite in low carbon irons In high carbon alloys thecontrary occurredmdashprecipitation progressed from the center ofthe dendrite towards the periphery They attributed this phe-nomenon to chromium and carbon micro-segregation towardsthe periphery of the dendrites

Just a few works on the kinetics of secondary carbide pre-cipitation have been published Powell and Laird[ 5 ] suggestedthat precipitation occurs during the first 25 min at the soakingtemperature and keeping the alloy for longer periods of timeonly causes coarsening of carbides Kuwano et al[ 8 ] felt thatprecipitation occurs within very short periods of time such as 1min at temperatures between 950-1000 degC which disagree with

Fig 9 Series of SEM micrographs showing secondary carbide precipitation during destabilization heat treatment at 900 degC of the experimentalwhite iron (a) 20 min (b) 1 h (c) 2 h and (d) 8 h Villelarsquos etching 45 s Matrix is composed of martensite retained austenite and secondarycarbides

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash377

the findings from this work What has been observed andwidely accepted by most researchers is that once carbides haveformed longer periods of time at the treatment temperatureincrease the size and volume fraction of them Furthermorestill longer periods of time at the soaking temperature cause thecoarsening of carbides at the expense of othersrsquo dissolution(Ostwald ripening process) reducing the number of carbides[ 8 ]

This effect can be observed in Fig 15

35 Retained Austenite After Destabilization

Figure 16 shows the results of retained austenite as mea-sured by XRD As discussed above secondary carbide precipi-tation depletes the austenitic matrix in chromium and car-bon[ 4 1 0 ] Such depletion produces an increase in MS and it ismore likely to obtain martensitic structures Such an effect isobserved in Fig 16 at 900 degC and long soaking times a re-duction in the retained austenite content is observed from 48at 10 min to just 8 for 8 h soaking time This means an

increase in the martensite volume fraction At this temperatureand long soaking times the volume fraction of secondary car-bides was 27 The chromium- and carbon-depleted matrixtransformed almost totally to martensite during the subsequentcooling Something similar occurred during destabilization at1000 degC residual austenite decreased from 52 for 5 minsoaking time to 15 for 8 h In this case the volume fractionof secondary carbides was less than in the alloy treated at 900degC a clear indication is that chromium and carbon dissolved inmatrix are higher in the alloy treated at 1000 degC and suchelements contribute to lower the MS temperature Thereforethe alloy treated at 1000 degC showed a higher amount of residualaustenite These results are in agreement with the findings ofsome authors[ 5 8 9 ] in high- and low-chromium white irons Onthe contrary retained austenite increased with soaking time forthe iron treated at 1150 degC For a 5 min soaking time retainedaustenite was 25 and for 1 h the whole matrix became aus-tenitic after heat treatment In this case the high temperaturesallowed carbon and chromium to dissolve completely in matrixand inhibited the driving force for secondary carbides precipi-

Fig 10 Series of SEM micrographs showing secondary carbide precipitation during destabilization heat treatment at 1000 degC of the experimentalwhite iron (a) 5 min (b) 10 min (c) 1 h and (d) 6 h Villelarsquos etching 45 s Matrix is composed of martensite retained austenite and secondarycarbides

378mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

tation Chromium- and carbon-rich austenite is stable underthese conditions and remains stable during cooling to roomtemperature[1 2 2 2 3 ]

36 Microhardness of Matrix and Bulk Hardness of theAlloy After Destabilization

Figure 17 shows the microhardness values obtained for theironrsquos matrix after each heat treatment At 900 degC for 10 minmicrohardness was 800HV5 0 and increased to a maximum of1300HV5 0 after 2-3 h soaking time afterwards it decreased to900HV5 0 for a destabilization of 8 h The increase in micro-hardness of the matrix is influenced by two factors as thesoaking time increased the precipitation of secondary carbidesthat strengthens the matrix by particle dispersion together withthe increase in martensite volume fraction However dimin-ishing microhardness after 3 h soaking time is attributed to thesecondary carbide coarsening At this stage longer times pro-mote a major diffusion of elements favoring coarsening ofsome carbides at the expense of othersrsquo dissolution such aphenomenon leads to a diminishing number of particles in thematrix (Fig 15) The final structure of the matrix is thereforehighly martensitic with relatively few coarse secondary car-bides For the structures of the irons with maximum hardnessa highly martensitic matrix was also observed but it containeda higher number of fine well-dispersed secondary carbides

A similar effect occurred for the iron treated at 1000 degChowever maximum hardness (1159HV5 0 ) was obtained atshorter times (30-60 min) Higher destabilization temperaturesallowed higher diffusion leading to a reduced amount ofcoarse nuclei (Fig 14 and 15) Under this situation thestrengthening due to dispersed particles is lower Furthermoreat these higher temperatures austenite can dissolve higher car-bon contents so that the amount of secondary carbides is alsominor The above phenomena lead to a higher retained austen-ite content in the final structure which reduced matrix hard-ness

In contrast for the iron treated at 1150 degC hardness of thematrix diminished since the beginning For this case a stabi-lized austenitic matrix was obtained Such a matrix was solidsolution strengthened by chromium carbon nickel molybde-num etc but not by particle dispersion Secondary precipita-tion occurred only in the early stages of the heat treatment

Regarding bulk hardness the alloy showed a behavior simi-

Fig 11 Series of SEM micrographs showing secondary carbide pre-cipitation during destabilization heat treatment at 1150 degC of the ex-perimental white iron (a) 2 min matrix is composed of martensiteretained austenite and secondary carbides (b) 20 min matrix com-posed of austenite mainly and (c) 1 h matrix is totally austeniticVillelarsquos etching 45 s

Fig 12 SEM micrograph showing the real morphology of the sec-ondary carbides precipitated during destabilization treatments Deep-etching was with a solution of 50 ml FeCl3 plus 20 ml HCl in 930 mlethanol for 15 min

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash379

lar to that for matrix microhardness Figure 18 shows the re-sults for hardness where maximum values can be observed forthe different temperatures For the iron treated at 900 degC maxi-mum hardness was 68 HRC and was obtained for1-2 h soakingtime For the iron treated at 1000 degC maximum hardness was60 HRC obtained for 30 min Finally for the iron treated at1150 degC hardness remains constant (45 HRC) during the first20 min of treatment and decreases afterwards The strengthen-ing of the matrix is the main factor responsible for the modi-fications in the ironrsquos hardness Microhardness of matrix andthe eutectic carbide volume fraction are the factors that con-

tribute to strengthening of the alloys For this case the carbidevolume fraction was constant (24) therefore the strengthen-ing of matrix caused the changes in bulk hardness

These strengthening phenomena are similar to those foundin aging processes of some alloys and are explained in terms ofthe diffusion of elements that participate in the precipitatingphase Such diffusion in turn is a function of treatment tem-perature and soaking time Kuwano et al[ 8 ] Powell andLaird[ 5 ] and Powell and Bee[ 9 ] have found similar behaviorsfor high- and low-chromium irons in studies of precipitationand hardness

Fig 13 Bright field TEM micrographs showing some secondary carbides (SC) in (a) sample treated at 900 degC for 1 h and (b) sample treated at1000 degC for 1 h The selected area diffraction pattern (SADP) on the particle arrowed in (b) indicates the crystallographic nature of these carbidesto be hexagonal M7C3

Fig 14 Diameter of the precipitated particles as a function of thesoaking time during destabilization heat treatment

Fig 15 Number of precipitated particles in a square micron (mm2) asa function of the soaking time during destabilization heat treatment

380mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

4 Conclusions

The structure of the as-cast alloy was composed of 24eutectic carbides about 5 martensite and 70 austenitesmall amounts of the molybdenum-rich carbide Mo2Cwere also present

Vanadium carbide was not detected in the alloy nonethe-less the iron contained 2 vanadium a carbide-promotingelement Vanadium was observed to partition to the car-bide and in minor amounts to the matrix

According to the studied temperatures the secondary car-bide precipitation level decreased with higher tempera-tures being the highest for the samples treated at 900 degC(27 of matrix)

High temperatures like 1150 degC over-stabilized the aus-tenitic matrix which led to a secondary carbides free aus-tenitic matrix after heat treatment

The amount of precipitated particles showed an increasewith time for the samples treated at 900 and 1000 degCreaching a maximum and then decreased due to dissolutionof some carbides and coarsening of some others

Matrix hardness showed a maximum when plotted againstsoaking time due to these precipitation processes and tothe transformation of austenite to martensite Maximumhardness was obtained for the samples treated for 1-2 h at900 degC where the structure showed a high volume fractionof martensite and a high amount of fine well-dispersedsecondary carbide particles

Acknowledgments

The authors are grateful to the University of Michoacan forfunding this project

References

1 CP Tabrett IR Sare and MR Gomashchi ldquoMicrostructure-Property Relationships in High-Chromium White Iron Alloysrdquo IntMater Rev 1996 41(2) pp 59-82

2 ON Dogan JA Hawk and G Laird II ldquoSolidification Structure andAbrasion Resistance of High Chromium White Ironsrdquo Metall MaterTrans A 1997 28(6) pp 1315-27

3 G Laird II and GLF Powell ldquoSolidification and Solid-State Trans-formation Mechanisms in Si Alloyed High-Chromium White CastIronsrdquo Metall Trans A 1993 24(2) pp 981-88

4 WW Cias ldquoAustenite Transformation Kinetics and Hardenability ofHeat Treated 175Cr White Cast Ironsrdquo Trans AFS 1974 82 pp317-28

5 GLF Powell and G Laird II ldquoStructure Nucleation Growth andMorphology of Secondary Carbides in High Chromium and Cr-NiWhite Cast Ironsrdquo J Mater Sci 1992 27 pp 29-35

6 F Maratray and A Poulalion ldquoAustenite Retention in High-Chromium White Ironsrdquo Trans AFS 1982 90 pp 795-804

7 JTH Pearce ldquoStructure and Wear Performance of Abrasion ResistantChromium White Cast Ironsrdquo Trans AFS 1984 92 pp 599-621

8 M Kuwano K Ogi A Sawamoto and K Matsuda ldquoStudies onPrecipitation Process of Secondary Carbides in High-Chromium CastIronrdquo Trans AFS 1990 98 pp 725-34

9 GLF Powell and JV Bee ldquoSecondary Carbide Precipitation in an 18wtCr-1 wtMo White Ironrdquo J Mater Sci 1996 31 pp 707-11

10 C Kim ldquoX-Ray Method of Measuring Retained Austenite in HeatTreated White Cast Ironsrdquo J Heat Treating 1979 1(2) pp 43-51

11 A Bedolla-Jacuinde ldquoMicrostructure of Vanadium- Niobium- andTitanium-Alloyed High-Chromium White Cast Ironsrdquo Int J CastMetals Res 2001 13(6) pp 343-61

Fig 16 Retained austenite content in iron matrix as a function of thesoaking time during the destabilization heat treatment

Fig 17 Matrix microhardness as a function of the soaking timeduring the destabilization heat treatment

Fig 18 Bulk hardness of the alloy as a function of the soaking timeduring the destabilization heat treatment

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash381

the most commercial high-chromium iron alloys[ 1 ] Molybde-num and nickel contents are close to 2 enough to ensure agood hardenability Also 2 vanadium was added to this ironfor further hardness Higher as-cast hardness values in thesealloys have been reported when adding 2 vanadium as thiselement partitions almost totally to the eutectic carbide makingthe composition (FeCrVMo)7C3 Vanadium increases thehardness of the eutectic carbide and also the carbide volumefraction which in turn increases the hardness of the alloy[ 1 1 ]

31 As-Cast Structure

Figure 1 shows the typical as-cast structure of the alloywhere the austenitic matrix and eutectic carbides can be ob-served Because the alloy is hypoeutectic austenite dendritesare the first to solidify followed by the eutectic Only eutecticcarbide and matrix phases can be observed from this picturehowever two additional phases are present as shown by XRDtraces from Fig 2 These additional phases present in smallamounts are martensite and a molybdenum rich carbide typeM2C Micrographs at higher magnifications show the presenceof martensite at the matrixcarbide interface in the eutecticzones (Fig 3) It has been widely reported that martensiteforms due to the carbon and chromium depletion by diffusiontowards the carbide during the solidification and the subse-quent cooling to room temperature[ 3 1 2 1 3 ] The martensiteamount as measured by XRD is about 6 The presence ofthe M2C carbide was also detected by EDS in the SEM Figure4 shows the morphology and composition of this carbide beinga molybdenum rich carbide its composition is of the type M2Ctypical for molybdenum carbide The main role of molybde-num in these irons is to improve hardenability however not allthe molybdenum contributes to this goal The presence of car-bon leads to the formation of the molybdenum rich carbide andpart of the molybdenum in these irons forms a eutectic carbidetype M2C at the end of the solidification stage[1 4 -1 9 ]

The elements distribution was analyzed by EDS in a randomarea of the as-cast microstructure Figure 5(a) shows an SEMmicrograph where some microanalyses were undertaken on the

matrix (Fig 5b) on the carbide (Fig 5c) and on a scan alongthe line shown on the micrograph in Fig 5(a) across an aus-tenite dendrite arm (Fig 5d)

Eutectic carbides are of the type M7C3 as revealed by XRD(Fig 2) and as reported for chromium contents of 17[ 7 ] Theelements present in the carbide phase can be seen from Fig5(b) despite being a qualitative analysis the intensity of thepeaks indicates that chromium and iron are present in higheramounts in the carbide The carbide contains also vanadiumand maybe small amounts of molybdenum which could not bedetected However the most important thing to highlight is thepresence of vanadium in the eutectic carbide XRD traces didnot show the presence of vanadium carbide Vanadium is acarbide-forming element of the type VC however the solidi-fication temperature for this carbide is below the solidificationtemperature for the alloy Therefore vanadium carbide forma-

Fig 3 SEM micrograph showing the presence of martensite at thematrixcarbide interface Villelarsquos etching 45 s

Fig 4 (a) SEM micrograph of a deep-etched sample showing amolybdenum rich carbide type M2C formed at the last solidificationstage etching solution 50 ml FeCl3 plus 20 ml HCl in 930 ml ethanolfor about 3 h (b) EDS of the M2C carbide in (a)

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash373

tion in this kind of irons is not likely These vanadium carbidesare believed to form only when the vanadium content exceeds3 as reported by Radulovic et al[ 2 0 ] Therefore the majorproportion of vanadium partitions to the carbide and the restremains dissolved in matrix as shown in the EDS in Fig 5(c)A minor proportion of chromium related to iron can be ob-served from this spectrum contrary to the findings in the car-bide microanalysis where the chromiumiron ratio is close to 1A high amount of molybdenum is also observed in matrix Thiselement remains in matrix in major proportion contributing thisway to improve hardenability From the same spectrum in Fig5(c) to the right of the last iron peak a small peak correspond-ing to nickel can be observed Nickel remains totally in matrixcontributing synergistically with molybdenum to hardenabilityVanadium can also be observed in matrix but in much smalleramounts than in the eutectic carbide The scan on the line inFig 5(a) shows the presence of chromium carbon and vana-dium in matrix a higher proportion of these elements is ob-

served at the edges of the graph These ends include the carbidephase as seen from the line The presence of these elements inmatrix is the source for secondary carbides precipitation duringthe heat treatment of austenite destabilization

Eutectic carbide is a very important phase in the propertiesof these alloys It has been mentioned that eutectic carbide is a3D network or an interconnected skeleton into the materialhowever the level of interconnection is much less than inunalloyed white irons This is due to the differences in prefer-ential growing directions for cementita and the carbide M7C3 M7C3 carbide has a hexagonal structure and a preferentialgrowing direction lt0001gt [ 7 ] therefore the typical morphol-ogy of these carbides is bars or plates (Fig 6) On a polishedsurface these bars could be perpendicular to the surface ap-pearing to be isolated particles under the microscope howeverdeep-etching techniques on the SEM make it possible to ob-serve that many bars may be separated at the surface but con-nected at the base

Fig 5 (a) SEM micrograph of the structure of the as-cast experimental iron (b) EDS corresponding to the eutectic carbide phase (c) EDScorresponding to the proeutectic matrix and (d) vanadium chromium and carbon profile of matrix along the line shown in (a)

374mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

32 Structure After Destabilization Heat Treatment

After destabilization heat treatment some phenomena tookplace which altered the structure of the material Such phe-nomena are a function of temperature and soaking time as inany heat treatment The phases present in the as-cast alloy are

the same as in the heat treated samples but the proportion haschanged The XRD pattern from Fig 7 is a representative proofof what happens after heat treatment The present phases areaustenite martensite M7C3 carbide and M2C carbide Themain difference that can be observed from this XRD comparedwith that from Fig 2 (as-cast material) is the higher intensityfor the martensite peaks (a phase) Therefore the amount ofmartensite is higher in the heat treated samples except in thosetreated at 1150 degC where the austenite phase was over-stabilized (see Section 35) According to the XRD traces sec-ondary carbides precipitated in matrix are of the type M7C3 thesame as the eutectic because no other type was detected Thekind of secondary carbides precipitating depend on the alloycomposition and the destabilization temperature The presenceof secondary carbides of the type M2 3C6 has been observed asfine interconnected bars in irons with very high chromiumamounts (gt25)[ 5 7 ] On the other hand secondary carbides ofthe type M7C3 were observed as agglomerated plates in ironswith chromium contents between 15-20[ 5 ] These resultsagree with the findings in the present work

33 Secondary Carbide Volume Fraction

Figure 8 shows the obtained results for the volume fractionof secondary carbides precipitated as a function of the soakingtime for each of the studied temperatures the horizontal axiswhich corresponds to time is plotted in logarithmic scale At900 degC precipitation was observed to start at about 10 min anda volume fraction of 2 was measured after 8 h of treatmentthe volume fraction was about 27 At 1000 degC the precipi-tation process was faster 2 at 5 min up to 21 at 10 minfrom 10 min to 8 h the volume fraction was about 22 Thetotal volume fraction precipitated within the first 10 min af-terwards the coarsening of some carbides and the dissolution ofothers occurred

At 1150 degC a volume fraction of 15 occurred at only 5min subsequently the volume fraction diminished to 2 at 1 hsoaking Figures 9 10 and 11 show sequences of SEM micro-graphs of some samples at different soaking times and at thetemperatures of 900 1000 and 1150 degC respectively Theseseries of micrographs represent visually what is plotted in Fig8 At 900 degC the diffusion rate of chromium and carbon lead tovolume fraction of secondary carbides of 27 in 8 h Howevermost of the precipitation took place within the first 2 h whichis not clear because the scale of time is plotted logarithmicallyHigher temperatures such as 1000 degC accelerated the precipi-tation process due to the higher diffusion rate of the elementsthat form the carbides chromium and carbon At this tempera-ture almost all of the volume of secondary carbides precipitatedwithin the first 10 min reaching a total volume fraction of22 In contrast at 1150 degC where diffusion rates for chro-mium and carbon are higher a carbide precipitation of 15was observed within the first 5 min but it started to diminishafterwards It is likely that the initial amount of carbides hadprecipitated during heating to the soaking temperature andthese carbides dissolved during soaking when the equilibriumcomposition of chromium and carbon in austenite was reachedAccording to the Fe-C phase diagram austenite can dissolvehigher carbon contents as the temperature increases to the eu-tectic

Fig 6 SEM micrographs with the 3-D structure of eutectic carbidesshowing the real morphology deep-etching with a solution of 50 mlFeCl3 plus 20 ml HCl in 930 ml ethanol for 3 h

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash375

As mentioned at 900 degC a higher secondary carbide volumefraction was observed at lower temperatures the ability ofaustenite to dissolve carbon is lower Therefore there is ahigher amount of carbon to precipitate which in turn is ob-served as a higher carbide volume fraction At higher tempera-tures austenite can retain carbon and chromium in higheramounts so there is very low raw material available to pre-cipitate carbides[ 5 8 9 ] and the final volume fraction is muchlower It is difficult to explain the diminishing of carbide vol-ume fraction observed at 1150 degC it is likely the equilibriumconcentration of carbon in austenite will be higher than theactual amount of this element in austenite so that during soak-ing precipitation is void Perhaps the volume fraction in thefirst sample (15 at 2 min) precipitated during heating to thesoaking temperature and these 2 min at 1150 degC were notenough for austenite to reach its equilibrium composition and

the secondary carbides did not dissolve Samples treated forlonger times (10 and 60 min) show a precipitated volume frac-tion of 2-5 perhaps the equilibrium volume in the alloy is at1150 degC

From micrographs in Fig 9 10 and 11 and particularlyfrom Fig 12 the morphology of secondary carbides can beobserved Secondary carbides are bars connected within thebulk matrix It has been reported[5 ] in white irons of similarcomposition that secondary carbides precipitated after destabi-lization treatments are of the type M7C3 having a hexagonalstructure and a preferred growing direction lt0001gt Figure 13shows some TEM micrographs of secondary carbides and adiffraction pattern of these particles which agree well with thefindings in Ref 5

34 Precipitated Carbides Size and Number

Figure 14 shows the particle size as a function of the soak-ing time these measurements are taken as the diameter of thecarbide bar It is well known that precipitation processes arefavored by temperature because temperature accelerates dif-fusion of the elements dissolved in matrix A higher diffusionrate of elements is observed for precipitation processes under-taken at high temperatures Because precipitation takes placeby nucleation and growth phenomena particles of a certain sizeare formed which depend on the treatment temperature In thecurrent study when the iron was treated at 900 degC the second-ary carbides diameter was about 120 nm within the first 2 hafterwards they coarsened up to 300 nm for 8 h soaking timeAt temperatures of 1000 degC the initial diameter is close to 200nm and gradually coarsens up to 250 nm for 8 h soaking timeAccording to this sequence it would be expected that the di-ameter of secondary carbides formed at 1150 degC to be higherhowever their size is only about 160 nm Because the volumeof material that precipitates in equilibrium as secondary car-

Fig 7 XRD pattern of the simple destabilized at 1000 degC for 1 h showing the detected phases austenite martensite M2C and M7C3

Fig 8 Volume of secondary carbides precipitated as a function ofthe soaking time during destabilization for each of the studied tem-peratures

376mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

bides at this temperature is almost void there is not enoughmaterial to precipitate coarse carbides

What is important to highlight is the precipitation behaviorat 900 and 1000 degC (Fig 15) This figure shows the numberprecipitated particles per mm2 At 900 degC the number of car-bides went from 2 particles per mm2 for 10 min soaking time to14 particles for 90 min Within the next 3 h the number ofcarbides remained at 14 and afterwards decreased to 8 particlesper mm2 for 8 h soaking time This finding indicates that theprecipitation process occurred during the first 90 min duringthe next 3 h there was no new precipitation just coarsening ofthe existing particles Finally during the next 4 h dissolutionof some carbides occurred contributing to the coarsening ofsome others thermodynamically more stable Similar phenom-ena during destabilization heat treatments have been reportedby Kuwano et al[ 8 ] All these phenomena are better observedfrom SEM micrographs (Fig 9 10 and 11) Note that from thecurrent study the secondary carbides precipitation was ob-served homogeneously in matrix According to Powell and

Bee[ 9 ] and Bee et al[ 2 1 ] secondary carbides precipitate at slipbands and sub-grain boundaries within austenite Such slipbands and sub-grain boundaries may form due to stresses gen-erated by differences in the thermal expansion coefficients ofcarbides and matrix On the contrary Kuwano et al[8 ] reportedthat the secondary carbides precipitation progressed from theperiphery adjacent to the eutectic carbides towards the centerof the dendrite in low carbon irons In high carbon alloys thecontrary occurredmdashprecipitation progressed from the center ofthe dendrite towards the periphery They attributed this phe-nomenon to chromium and carbon micro-segregation towardsthe periphery of the dendrites

Just a few works on the kinetics of secondary carbide pre-cipitation have been published Powell and Laird[ 5 ] suggestedthat precipitation occurs during the first 25 min at the soakingtemperature and keeping the alloy for longer periods of timeonly causes coarsening of carbides Kuwano et al[ 8 ] felt thatprecipitation occurs within very short periods of time such as 1min at temperatures between 950-1000 degC which disagree with

Fig 9 Series of SEM micrographs showing secondary carbide precipitation during destabilization heat treatment at 900 degC of the experimentalwhite iron (a) 20 min (b) 1 h (c) 2 h and (d) 8 h Villelarsquos etching 45 s Matrix is composed of martensite retained austenite and secondarycarbides

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash377

the findings from this work What has been observed andwidely accepted by most researchers is that once carbides haveformed longer periods of time at the treatment temperatureincrease the size and volume fraction of them Furthermorestill longer periods of time at the soaking temperature cause thecoarsening of carbides at the expense of othersrsquo dissolution(Ostwald ripening process) reducing the number of carbides[ 8 ]

This effect can be observed in Fig 15

35 Retained Austenite After Destabilization

Figure 16 shows the results of retained austenite as mea-sured by XRD As discussed above secondary carbide precipi-tation depletes the austenitic matrix in chromium and car-bon[ 4 1 0 ] Such depletion produces an increase in MS and it ismore likely to obtain martensitic structures Such an effect isobserved in Fig 16 at 900 degC and long soaking times a re-duction in the retained austenite content is observed from 48at 10 min to just 8 for 8 h soaking time This means an

increase in the martensite volume fraction At this temperatureand long soaking times the volume fraction of secondary car-bides was 27 The chromium- and carbon-depleted matrixtransformed almost totally to martensite during the subsequentcooling Something similar occurred during destabilization at1000 degC residual austenite decreased from 52 for 5 minsoaking time to 15 for 8 h In this case the volume fractionof secondary carbides was less than in the alloy treated at 900degC a clear indication is that chromium and carbon dissolved inmatrix are higher in the alloy treated at 1000 degC and suchelements contribute to lower the MS temperature Thereforethe alloy treated at 1000 degC showed a higher amount of residualaustenite These results are in agreement with the findings ofsome authors[ 5 8 9 ] in high- and low-chromium white irons Onthe contrary retained austenite increased with soaking time forthe iron treated at 1150 degC For a 5 min soaking time retainedaustenite was 25 and for 1 h the whole matrix became aus-tenitic after heat treatment In this case the high temperaturesallowed carbon and chromium to dissolve completely in matrixand inhibited the driving force for secondary carbides precipi-

Fig 10 Series of SEM micrographs showing secondary carbide precipitation during destabilization heat treatment at 1000 degC of the experimentalwhite iron (a) 5 min (b) 10 min (c) 1 h and (d) 6 h Villelarsquos etching 45 s Matrix is composed of martensite retained austenite and secondarycarbides

378mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

tation Chromium- and carbon-rich austenite is stable underthese conditions and remains stable during cooling to roomtemperature[1 2 2 2 3 ]

36 Microhardness of Matrix and Bulk Hardness of theAlloy After Destabilization

Figure 17 shows the microhardness values obtained for theironrsquos matrix after each heat treatment At 900 degC for 10 minmicrohardness was 800HV5 0 and increased to a maximum of1300HV5 0 after 2-3 h soaking time afterwards it decreased to900HV5 0 for a destabilization of 8 h The increase in micro-hardness of the matrix is influenced by two factors as thesoaking time increased the precipitation of secondary carbidesthat strengthens the matrix by particle dispersion together withthe increase in martensite volume fraction However dimin-ishing microhardness after 3 h soaking time is attributed to thesecondary carbide coarsening At this stage longer times pro-mote a major diffusion of elements favoring coarsening ofsome carbides at the expense of othersrsquo dissolution such aphenomenon leads to a diminishing number of particles in thematrix (Fig 15) The final structure of the matrix is thereforehighly martensitic with relatively few coarse secondary car-bides For the structures of the irons with maximum hardnessa highly martensitic matrix was also observed but it containeda higher number of fine well-dispersed secondary carbides

A similar effect occurred for the iron treated at 1000 degChowever maximum hardness (1159HV5 0 ) was obtained atshorter times (30-60 min) Higher destabilization temperaturesallowed higher diffusion leading to a reduced amount ofcoarse nuclei (Fig 14 and 15) Under this situation thestrengthening due to dispersed particles is lower Furthermoreat these higher temperatures austenite can dissolve higher car-bon contents so that the amount of secondary carbides is alsominor The above phenomena lead to a higher retained austen-ite content in the final structure which reduced matrix hard-ness

In contrast for the iron treated at 1150 degC hardness of thematrix diminished since the beginning For this case a stabi-lized austenitic matrix was obtained Such a matrix was solidsolution strengthened by chromium carbon nickel molybde-num etc but not by particle dispersion Secondary precipita-tion occurred only in the early stages of the heat treatment

Regarding bulk hardness the alloy showed a behavior simi-

Fig 11 Series of SEM micrographs showing secondary carbide pre-cipitation during destabilization heat treatment at 1150 degC of the ex-perimental white iron (a) 2 min matrix is composed of martensiteretained austenite and secondary carbides (b) 20 min matrix com-posed of austenite mainly and (c) 1 h matrix is totally austeniticVillelarsquos etching 45 s

Fig 12 SEM micrograph showing the real morphology of the sec-ondary carbides precipitated during destabilization treatments Deep-etching was with a solution of 50 ml FeCl3 plus 20 ml HCl in 930 mlethanol for 15 min

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash379

lar to that for matrix microhardness Figure 18 shows the re-sults for hardness where maximum values can be observed forthe different temperatures For the iron treated at 900 degC maxi-mum hardness was 68 HRC and was obtained for1-2 h soakingtime For the iron treated at 1000 degC maximum hardness was60 HRC obtained for 30 min Finally for the iron treated at1150 degC hardness remains constant (45 HRC) during the first20 min of treatment and decreases afterwards The strengthen-ing of the matrix is the main factor responsible for the modi-fications in the ironrsquos hardness Microhardness of matrix andthe eutectic carbide volume fraction are the factors that con-

tribute to strengthening of the alloys For this case the carbidevolume fraction was constant (24) therefore the strengthen-ing of matrix caused the changes in bulk hardness

These strengthening phenomena are similar to those foundin aging processes of some alloys and are explained in terms ofthe diffusion of elements that participate in the precipitatingphase Such diffusion in turn is a function of treatment tem-perature and soaking time Kuwano et al[ 8 ] Powell andLaird[ 5 ] and Powell and Bee[ 9 ] have found similar behaviorsfor high- and low-chromium irons in studies of precipitationand hardness

Fig 13 Bright field TEM micrographs showing some secondary carbides (SC) in (a) sample treated at 900 degC for 1 h and (b) sample treated at1000 degC for 1 h The selected area diffraction pattern (SADP) on the particle arrowed in (b) indicates the crystallographic nature of these carbidesto be hexagonal M7C3

Fig 14 Diameter of the precipitated particles as a function of thesoaking time during destabilization heat treatment

Fig 15 Number of precipitated particles in a square micron (mm2) asa function of the soaking time during destabilization heat treatment

380mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

4 Conclusions

The structure of the as-cast alloy was composed of 24eutectic carbides about 5 martensite and 70 austenitesmall amounts of the molybdenum-rich carbide Mo2Cwere also present

Vanadium carbide was not detected in the alloy nonethe-less the iron contained 2 vanadium a carbide-promotingelement Vanadium was observed to partition to the car-bide and in minor amounts to the matrix

According to the studied temperatures the secondary car-bide precipitation level decreased with higher tempera-tures being the highest for the samples treated at 900 degC(27 of matrix)

High temperatures like 1150 degC over-stabilized the aus-tenitic matrix which led to a secondary carbides free aus-tenitic matrix after heat treatment

The amount of precipitated particles showed an increasewith time for the samples treated at 900 and 1000 degCreaching a maximum and then decreased due to dissolutionof some carbides and coarsening of some others

Matrix hardness showed a maximum when plotted againstsoaking time due to these precipitation processes and tothe transformation of austenite to martensite Maximumhardness was obtained for the samples treated for 1-2 h at900 degC where the structure showed a high volume fractionof martensite and a high amount of fine well-dispersedsecondary carbide particles

Acknowledgments

The authors are grateful to the University of Michoacan forfunding this project

References

1 CP Tabrett IR Sare and MR Gomashchi ldquoMicrostructure-Property Relationships in High-Chromium White Iron Alloysrdquo IntMater Rev 1996 41(2) pp 59-82

2 ON Dogan JA Hawk and G Laird II ldquoSolidification Structure andAbrasion Resistance of High Chromium White Ironsrdquo Metall MaterTrans A 1997 28(6) pp 1315-27

3 G Laird II and GLF Powell ldquoSolidification and Solid-State Trans-formation Mechanisms in Si Alloyed High-Chromium White CastIronsrdquo Metall Trans A 1993 24(2) pp 981-88

4 WW Cias ldquoAustenite Transformation Kinetics and Hardenability ofHeat Treated 175Cr White Cast Ironsrdquo Trans AFS 1974 82 pp317-28

5 GLF Powell and G Laird II ldquoStructure Nucleation Growth andMorphology of Secondary Carbides in High Chromium and Cr-NiWhite Cast Ironsrdquo J Mater Sci 1992 27 pp 29-35

6 F Maratray and A Poulalion ldquoAustenite Retention in High-Chromium White Ironsrdquo Trans AFS 1982 90 pp 795-804

7 JTH Pearce ldquoStructure and Wear Performance of Abrasion ResistantChromium White Cast Ironsrdquo Trans AFS 1984 92 pp 599-621

8 M Kuwano K Ogi A Sawamoto and K Matsuda ldquoStudies onPrecipitation Process of Secondary Carbides in High-Chromium CastIronrdquo Trans AFS 1990 98 pp 725-34

9 GLF Powell and JV Bee ldquoSecondary Carbide Precipitation in an 18wtCr-1 wtMo White Ironrdquo J Mater Sci 1996 31 pp 707-11

10 C Kim ldquoX-Ray Method of Measuring Retained Austenite in HeatTreated White Cast Ironsrdquo J Heat Treating 1979 1(2) pp 43-51

11 A Bedolla-Jacuinde ldquoMicrostructure of Vanadium- Niobium- andTitanium-Alloyed High-Chromium White Cast Ironsrdquo Int J CastMetals Res 2001 13(6) pp 343-61

Fig 16 Retained austenite content in iron matrix as a function of thesoaking time during the destabilization heat treatment

Fig 17 Matrix microhardness as a function of the soaking timeduring the destabilization heat treatment

Fig 18 Bulk hardness of the alloy as a function of the soaking timeduring the destabilization heat treatment

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash381

tion in this kind of irons is not likely These vanadium carbidesare believed to form only when the vanadium content exceeds3 as reported by Radulovic et al[ 2 0 ] Therefore the majorproportion of vanadium partitions to the carbide and the restremains dissolved in matrix as shown in the EDS in Fig 5(c)A minor proportion of chromium related to iron can be ob-served from this spectrum contrary to the findings in the car-bide microanalysis where the chromiumiron ratio is close to 1A high amount of molybdenum is also observed in matrix Thiselement remains in matrix in major proportion contributing thisway to improve hardenability From the same spectrum in Fig5(c) to the right of the last iron peak a small peak correspond-ing to nickel can be observed Nickel remains totally in matrixcontributing synergistically with molybdenum to hardenabilityVanadium can also be observed in matrix but in much smalleramounts than in the eutectic carbide The scan on the line inFig 5(a) shows the presence of chromium carbon and vana-dium in matrix a higher proportion of these elements is ob-

served at the edges of the graph These ends include the carbidephase as seen from the line The presence of these elements inmatrix is the source for secondary carbides precipitation duringthe heat treatment of austenite destabilization

Eutectic carbide is a very important phase in the propertiesof these alloys It has been mentioned that eutectic carbide is a3D network or an interconnected skeleton into the materialhowever the level of interconnection is much less than inunalloyed white irons This is due to the differences in prefer-ential growing directions for cementita and the carbide M7C3 M7C3 carbide has a hexagonal structure and a preferentialgrowing direction lt0001gt [ 7 ] therefore the typical morphol-ogy of these carbides is bars or plates (Fig 6) On a polishedsurface these bars could be perpendicular to the surface ap-pearing to be isolated particles under the microscope howeverdeep-etching techniques on the SEM make it possible to ob-serve that many bars may be separated at the surface but con-nected at the base

Fig 5 (a) SEM micrograph of the structure of the as-cast experimental iron (b) EDS corresponding to the eutectic carbide phase (c) EDScorresponding to the proeutectic matrix and (d) vanadium chromium and carbon profile of matrix along the line shown in (a)

374mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

32 Structure After Destabilization Heat Treatment

After destabilization heat treatment some phenomena tookplace which altered the structure of the material Such phe-nomena are a function of temperature and soaking time as inany heat treatment The phases present in the as-cast alloy are

the same as in the heat treated samples but the proportion haschanged The XRD pattern from Fig 7 is a representative proofof what happens after heat treatment The present phases areaustenite martensite M7C3 carbide and M2C carbide Themain difference that can be observed from this XRD comparedwith that from Fig 2 (as-cast material) is the higher intensityfor the martensite peaks (a phase) Therefore the amount ofmartensite is higher in the heat treated samples except in thosetreated at 1150 degC where the austenite phase was over-stabilized (see Section 35) According to the XRD traces sec-ondary carbides precipitated in matrix are of the type M7C3 thesame as the eutectic because no other type was detected Thekind of secondary carbides precipitating depend on the alloycomposition and the destabilization temperature The presenceof secondary carbides of the type M2 3C6 has been observed asfine interconnected bars in irons with very high chromiumamounts (gt25)[ 5 7 ] On the other hand secondary carbides ofthe type M7C3 were observed as agglomerated plates in ironswith chromium contents between 15-20[ 5 ] These resultsagree with the findings in the present work

33 Secondary Carbide Volume Fraction

Figure 8 shows the obtained results for the volume fractionof secondary carbides precipitated as a function of the soakingtime for each of the studied temperatures the horizontal axiswhich corresponds to time is plotted in logarithmic scale At900 degC precipitation was observed to start at about 10 min anda volume fraction of 2 was measured after 8 h of treatmentthe volume fraction was about 27 At 1000 degC the precipi-tation process was faster 2 at 5 min up to 21 at 10 minfrom 10 min to 8 h the volume fraction was about 22 Thetotal volume fraction precipitated within the first 10 min af-terwards the coarsening of some carbides and the dissolution ofothers occurred

At 1150 degC a volume fraction of 15 occurred at only 5min subsequently the volume fraction diminished to 2 at 1 hsoaking Figures 9 10 and 11 show sequences of SEM micro-graphs of some samples at different soaking times and at thetemperatures of 900 1000 and 1150 degC respectively Theseseries of micrographs represent visually what is plotted in Fig8 At 900 degC the diffusion rate of chromium and carbon lead tovolume fraction of secondary carbides of 27 in 8 h Howevermost of the precipitation took place within the first 2 h whichis not clear because the scale of time is plotted logarithmicallyHigher temperatures such as 1000 degC accelerated the precipi-tation process due to the higher diffusion rate of the elementsthat form the carbides chromium and carbon At this tempera-ture almost all of the volume of secondary carbides precipitatedwithin the first 10 min reaching a total volume fraction of22 In contrast at 1150 degC where diffusion rates for chro-mium and carbon are higher a carbide precipitation of 15was observed within the first 5 min but it started to diminishafterwards It is likely that the initial amount of carbides hadprecipitated during heating to the soaking temperature andthese carbides dissolved during soaking when the equilibriumcomposition of chromium and carbon in austenite was reachedAccording to the Fe-C phase diagram austenite can dissolvehigher carbon contents as the temperature increases to the eu-tectic

Fig 6 SEM micrographs with the 3-D structure of eutectic carbidesshowing the real morphology deep-etching with a solution of 50 mlFeCl3 plus 20 ml HCl in 930 ml ethanol for 3 h

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash375

As mentioned at 900 degC a higher secondary carbide volumefraction was observed at lower temperatures the ability ofaustenite to dissolve carbon is lower Therefore there is ahigher amount of carbon to precipitate which in turn is ob-served as a higher carbide volume fraction At higher tempera-tures austenite can retain carbon and chromium in higheramounts so there is very low raw material available to pre-cipitate carbides[ 5 8 9 ] and the final volume fraction is muchlower It is difficult to explain the diminishing of carbide vol-ume fraction observed at 1150 degC it is likely the equilibriumconcentration of carbon in austenite will be higher than theactual amount of this element in austenite so that during soak-ing precipitation is void Perhaps the volume fraction in thefirst sample (15 at 2 min) precipitated during heating to thesoaking temperature and these 2 min at 1150 degC were notenough for austenite to reach its equilibrium composition and

the secondary carbides did not dissolve Samples treated forlonger times (10 and 60 min) show a precipitated volume frac-tion of 2-5 perhaps the equilibrium volume in the alloy is at1150 degC

From micrographs in Fig 9 10 and 11 and particularlyfrom Fig 12 the morphology of secondary carbides can beobserved Secondary carbides are bars connected within thebulk matrix It has been reported[5 ] in white irons of similarcomposition that secondary carbides precipitated after destabi-lization treatments are of the type M7C3 having a hexagonalstructure and a preferred growing direction lt0001gt Figure 13shows some TEM micrographs of secondary carbides and adiffraction pattern of these particles which agree well with thefindings in Ref 5

34 Precipitated Carbides Size and Number

Figure 14 shows the particle size as a function of the soak-ing time these measurements are taken as the diameter of thecarbide bar It is well known that precipitation processes arefavored by temperature because temperature accelerates dif-fusion of the elements dissolved in matrix A higher diffusionrate of elements is observed for precipitation processes under-taken at high temperatures Because precipitation takes placeby nucleation and growth phenomena particles of a certain sizeare formed which depend on the treatment temperature In thecurrent study when the iron was treated at 900 degC the second-ary carbides diameter was about 120 nm within the first 2 hafterwards they coarsened up to 300 nm for 8 h soaking timeAt temperatures of 1000 degC the initial diameter is close to 200nm and gradually coarsens up to 250 nm for 8 h soaking timeAccording to this sequence it would be expected that the di-ameter of secondary carbides formed at 1150 degC to be higherhowever their size is only about 160 nm Because the volumeof material that precipitates in equilibrium as secondary car-

Fig 7 XRD pattern of the simple destabilized at 1000 degC for 1 h showing the detected phases austenite martensite M2C and M7C3

Fig 8 Volume of secondary carbides precipitated as a function ofthe soaking time during destabilization for each of the studied tem-peratures

376mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

bides at this temperature is almost void there is not enoughmaterial to precipitate coarse carbides

What is important to highlight is the precipitation behaviorat 900 and 1000 degC (Fig 15) This figure shows the numberprecipitated particles per mm2 At 900 degC the number of car-bides went from 2 particles per mm2 for 10 min soaking time to14 particles for 90 min Within the next 3 h the number ofcarbides remained at 14 and afterwards decreased to 8 particlesper mm2 for 8 h soaking time This finding indicates that theprecipitation process occurred during the first 90 min duringthe next 3 h there was no new precipitation just coarsening ofthe existing particles Finally during the next 4 h dissolutionof some carbides occurred contributing to the coarsening ofsome others thermodynamically more stable Similar phenom-ena during destabilization heat treatments have been reportedby Kuwano et al[ 8 ] All these phenomena are better observedfrom SEM micrographs (Fig 9 10 and 11) Note that from thecurrent study the secondary carbides precipitation was ob-served homogeneously in matrix According to Powell and

Bee[ 9 ] and Bee et al[ 2 1 ] secondary carbides precipitate at slipbands and sub-grain boundaries within austenite Such slipbands and sub-grain boundaries may form due to stresses gen-erated by differences in the thermal expansion coefficients ofcarbides and matrix On the contrary Kuwano et al[8 ] reportedthat the secondary carbides precipitation progressed from theperiphery adjacent to the eutectic carbides towards the centerof the dendrite in low carbon irons In high carbon alloys thecontrary occurredmdashprecipitation progressed from the center ofthe dendrite towards the periphery They attributed this phe-nomenon to chromium and carbon micro-segregation towardsthe periphery of the dendrites

Just a few works on the kinetics of secondary carbide pre-cipitation have been published Powell and Laird[ 5 ] suggestedthat precipitation occurs during the first 25 min at the soakingtemperature and keeping the alloy for longer periods of timeonly causes coarsening of carbides Kuwano et al[ 8 ] felt thatprecipitation occurs within very short periods of time such as 1min at temperatures between 950-1000 degC which disagree with

Fig 9 Series of SEM micrographs showing secondary carbide precipitation during destabilization heat treatment at 900 degC of the experimentalwhite iron (a) 20 min (b) 1 h (c) 2 h and (d) 8 h Villelarsquos etching 45 s Matrix is composed of martensite retained austenite and secondarycarbides

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash377

the findings from this work What has been observed andwidely accepted by most researchers is that once carbides haveformed longer periods of time at the treatment temperatureincrease the size and volume fraction of them Furthermorestill longer periods of time at the soaking temperature cause thecoarsening of carbides at the expense of othersrsquo dissolution(Ostwald ripening process) reducing the number of carbides[ 8 ]

This effect can be observed in Fig 15

35 Retained Austenite After Destabilization

Figure 16 shows the results of retained austenite as mea-sured by XRD As discussed above secondary carbide precipi-tation depletes the austenitic matrix in chromium and car-bon[ 4 1 0 ] Such depletion produces an increase in MS and it ismore likely to obtain martensitic structures Such an effect isobserved in Fig 16 at 900 degC and long soaking times a re-duction in the retained austenite content is observed from 48at 10 min to just 8 for 8 h soaking time This means an

increase in the martensite volume fraction At this temperatureand long soaking times the volume fraction of secondary car-bides was 27 The chromium- and carbon-depleted matrixtransformed almost totally to martensite during the subsequentcooling Something similar occurred during destabilization at1000 degC residual austenite decreased from 52 for 5 minsoaking time to 15 for 8 h In this case the volume fractionof secondary carbides was less than in the alloy treated at 900degC a clear indication is that chromium and carbon dissolved inmatrix are higher in the alloy treated at 1000 degC and suchelements contribute to lower the MS temperature Thereforethe alloy treated at 1000 degC showed a higher amount of residualaustenite These results are in agreement with the findings ofsome authors[ 5 8 9 ] in high- and low-chromium white irons Onthe contrary retained austenite increased with soaking time forthe iron treated at 1150 degC For a 5 min soaking time retainedaustenite was 25 and for 1 h the whole matrix became aus-tenitic after heat treatment In this case the high temperaturesallowed carbon and chromium to dissolve completely in matrixand inhibited the driving force for secondary carbides precipi-

Fig 10 Series of SEM micrographs showing secondary carbide precipitation during destabilization heat treatment at 1000 degC of the experimentalwhite iron (a) 5 min (b) 10 min (c) 1 h and (d) 6 h Villelarsquos etching 45 s Matrix is composed of martensite retained austenite and secondarycarbides

378mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

tation Chromium- and carbon-rich austenite is stable underthese conditions and remains stable during cooling to roomtemperature[1 2 2 2 3 ]

36 Microhardness of Matrix and Bulk Hardness of theAlloy After Destabilization

Figure 17 shows the microhardness values obtained for theironrsquos matrix after each heat treatment At 900 degC for 10 minmicrohardness was 800HV5 0 and increased to a maximum of1300HV5 0 after 2-3 h soaking time afterwards it decreased to900HV5 0 for a destabilization of 8 h The increase in micro-hardness of the matrix is influenced by two factors as thesoaking time increased the precipitation of secondary carbidesthat strengthens the matrix by particle dispersion together withthe increase in martensite volume fraction However dimin-ishing microhardness after 3 h soaking time is attributed to thesecondary carbide coarsening At this stage longer times pro-mote a major diffusion of elements favoring coarsening ofsome carbides at the expense of othersrsquo dissolution such aphenomenon leads to a diminishing number of particles in thematrix (Fig 15) The final structure of the matrix is thereforehighly martensitic with relatively few coarse secondary car-bides For the structures of the irons with maximum hardnessa highly martensitic matrix was also observed but it containeda higher number of fine well-dispersed secondary carbides

A similar effect occurred for the iron treated at 1000 degChowever maximum hardness (1159HV5 0 ) was obtained atshorter times (30-60 min) Higher destabilization temperaturesallowed higher diffusion leading to a reduced amount ofcoarse nuclei (Fig 14 and 15) Under this situation thestrengthening due to dispersed particles is lower Furthermoreat these higher temperatures austenite can dissolve higher car-bon contents so that the amount of secondary carbides is alsominor The above phenomena lead to a higher retained austen-ite content in the final structure which reduced matrix hard-ness

In contrast for the iron treated at 1150 degC hardness of thematrix diminished since the beginning For this case a stabi-lized austenitic matrix was obtained Such a matrix was solidsolution strengthened by chromium carbon nickel molybde-num etc but not by particle dispersion Secondary precipita-tion occurred only in the early stages of the heat treatment

Regarding bulk hardness the alloy showed a behavior simi-

Fig 11 Series of SEM micrographs showing secondary carbide pre-cipitation during destabilization heat treatment at 1150 degC of the ex-perimental white iron (a) 2 min matrix is composed of martensiteretained austenite and secondary carbides (b) 20 min matrix com-posed of austenite mainly and (c) 1 h matrix is totally austeniticVillelarsquos etching 45 s

Fig 12 SEM micrograph showing the real morphology of the sec-ondary carbides precipitated during destabilization treatments Deep-etching was with a solution of 50 ml FeCl3 plus 20 ml HCl in 930 mlethanol for 15 min

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash379

lar to that for matrix microhardness Figure 18 shows the re-sults for hardness where maximum values can be observed forthe different temperatures For the iron treated at 900 degC maxi-mum hardness was 68 HRC and was obtained for1-2 h soakingtime For the iron treated at 1000 degC maximum hardness was60 HRC obtained for 30 min Finally for the iron treated at1150 degC hardness remains constant (45 HRC) during the first20 min of treatment and decreases afterwards The strengthen-ing of the matrix is the main factor responsible for the modi-fications in the ironrsquos hardness Microhardness of matrix andthe eutectic carbide volume fraction are the factors that con-

tribute to strengthening of the alloys For this case the carbidevolume fraction was constant (24) therefore the strengthen-ing of matrix caused the changes in bulk hardness

These strengthening phenomena are similar to those foundin aging processes of some alloys and are explained in terms ofthe diffusion of elements that participate in the precipitatingphase Such diffusion in turn is a function of treatment tem-perature and soaking time Kuwano et al[ 8 ] Powell andLaird[ 5 ] and Powell and Bee[ 9 ] have found similar behaviorsfor high- and low-chromium irons in studies of precipitationand hardness

Fig 13 Bright field TEM micrographs showing some secondary carbides (SC) in (a) sample treated at 900 degC for 1 h and (b) sample treated at1000 degC for 1 h The selected area diffraction pattern (SADP) on the particle arrowed in (b) indicates the crystallographic nature of these carbidesto be hexagonal M7C3

Fig 14 Diameter of the precipitated particles as a function of thesoaking time during destabilization heat treatment

Fig 15 Number of precipitated particles in a square micron (mm2) asa function of the soaking time during destabilization heat treatment

380mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

4 Conclusions

The structure of the as-cast alloy was composed of 24eutectic carbides about 5 martensite and 70 austenitesmall amounts of the molybdenum-rich carbide Mo2Cwere also present

Vanadium carbide was not detected in the alloy nonethe-less the iron contained 2 vanadium a carbide-promotingelement Vanadium was observed to partition to the car-bide and in minor amounts to the matrix

According to the studied temperatures the secondary car-bide precipitation level decreased with higher tempera-tures being the highest for the samples treated at 900 degC(27 of matrix)

High temperatures like 1150 degC over-stabilized the aus-tenitic matrix which led to a secondary carbides free aus-tenitic matrix after heat treatment

The amount of precipitated particles showed an increasewith time for the samples treated at 900 and 1000 degCreaching a maximum and then decreased due to dissolutionof some carbides and coarsening of some others

Matrix hardness showed a maximum when plotted againstsoaking time due to these precipitation processes and tothe transformation of austenite to martensite Maximumhardness was obtained for the samples treated for 1-2 h at900 degC where the structure showed a high volume fractionof martensite and a high amount of fine well-dispersedsecondary carbide particles

Acknowledgments

The authors are grateful to the University of Michoacan forfunding this project

References

1 CP Tabrett IR Sare and MR Gomashchi ldquoMicrostructure-Property Relationships in High-Chromium White Iron Alloysrdquo IntMater Rev 1996 41(2) pp 59-82

2 ON Dogan JA Hawk and G Laird II ldquoSolidification Structure andAbrasion Resistance of High Chromium White Ironsrdquo Metall MaterTrans A 1997 28(6) pp 1315-27

3 G Laird II and GLF Powell ldquoSolidification and Solid-State Trans-formation Mechanisms in Si Alloyed High-Chromium White CastIronsrdquo Metall Trans A 1993 24(2) pp 981-88

4 WW Cias ldquoAustenite Transformation Kinetics and Hardenability ofHeat Treated 175Cr White Cast Ironsrdquo Trans AFS 1974 82 pp317-28

5 GLF Powell and G Laird II ldquoStructure Nucleation Growth andMorphology of Secondary Carbides in High Chromium and Cr-NiWhite Cast Ironsrdquo J Mater Sci 1992 27 pp 29-35

6 F Maratray and A Poulalion ldquoAustenite Retention in High-Chromium White Ironsrdquo Trans AFS 1982 90 pp 795-804

7 JTH Pearce ldquoStructure and Wear Performance of Abrasion ResistantChromium White Cast Ironsrdquo Trans AFS 1984 92 pp 599-621

8 M Kuwano K Ogi A Sawamoto and K Matsuda ldquoStudies onPrecipitation Process of Secondary Carbides in High-Chromium CastIronrdquo Trans AFS 1990 98 pp 725-34

9 GLF Powell and JV Bee ldquoSecondary Carbide Precipitation in an 18wtCr-1 wtMo White Ironrdquo J Mater Sci 1996 31 pp 707-11

10 C Kim ldquoX-Ray Method of Measuring Retained Austenite in HeatTreated White Cast Ironsrdquo J Heat Treating 1979 1(2) pp 43-51

11 A Bedolla-Jacuinde ldquoMicrostructure of Vanadium- Niobium- andTitanium-Alloyed High-Chromium White Cast Ironsrdquo Int J CastMetals Res 2001 13(6) pp 343-61

Fig 16 Retained austenite content in iron matrix as a function of thesoaking time during the destabilization heat treatment

Fig 17 Matrix microhardness as a function of the soaking timeduring the destabilization heat treatment

Fig 18 Bulk hardness of the alloy as a function of the soaking timeduring the destabilization heat treatment

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash381

32 Structure After Destabilization Heat Treatment

After destabilization heat treatment some phenomena tookplace which altered the structure of the material Such phe-nomena are a function of temperature and soaking time as inany heat treatment The phases present in the as-cast alloy are

the same as in the heat treated samples but the proportion haschanged The XRD pattern from Fig 7 is a representative proofof what happens after heat treatment The present phases areaustenite martensite M7C3 carbide and M2C carbide Themain difference that can be observed from this XRD comparedwith that from Fig 2 (as-cast material) is the higher intensityfor the martensite peaks (a phase) Therefore the amount ofmartensite is higher in the heat treated samples except in thosetreated at 1150 degC where the austenite phase was over-stabilized (see Section 35) According to the XRD traces sec-ondary carbides precipitated in matrix are of the type M7C3 thesame as the eutectic because no other type was detected Thekind of secondary carbides precipitating depend on the alloycomposition and the destabilization temperature The presenceof secondary carbides of the type M2 3C6 has been observed asfine interconnected bars in irons with very high chromiumamounts (gt25)[ 5 7 ] On the other hand secondary carbides ofthe type M7C3 were observed as agglomerated plates in ironswith chromium contents between 15-20[ 5 ] These resultsagree with the findings in the present work

33 Secondary Carbide Volume Fraction

Figure 8 shows the obtained results for the volume fractionof secondary carbides precipitated as a function of the soakingtime for each of the studied temperatures the horizontal axiswhich corresponds to time is plotted in logarithmic scale At900 degC precipitation was observed to start at about 10 min anda volume fraction of 2 was measured after 8 h of treatmentthe volume fraction was about 27 At 1000 degC the precipi-tation process was faster 2 at 5 min up to 21 at 10 minfrom 10 min to 8 h the volume fraction was about 22 Thetotal volume fraction precipitated within the first 10 min af-terwards the coarsening of some carbides and the dissolution ofothers occurred

At 1150 degC a volume fraction of 15 occurred at only 5min subsequently the volume fraction diminished to 2 at 1 hsoaking Figures 9 10 and 11 show sequences of SEM micro-graphs of some samples at different soaking times and at thetemperatures of 900 1000 and 1150 degC respectively Theseseries of micrographs represent visually what is plotted in Fig8 At 900 degC the diffusion rate of chromium and carbon lead tovolume fraction of secondary carbides of 27 in 8 h Howevermost of the precipitation took place within the first 2 h whichis not clear because the scale of time is plotted logarithmicallyHigher temperatures such as 1000 degC accelerated the precipi-tation process due to the higher diffusion rate of the elementsthat form the carbides chromium and carbon At this tempera-ture almost all of the volume of secondary carbides precipitatedwithin the first 10 min reaching a total volume fraction of22 In contrast at 1150 degC where diffusion rates for chro-mium and carbon are higher a carbide precipitation of 15was observed within the first 5 min but it started to diminishafterwards It is likely that the initial amount of carbides hadprecipitated during heating to the soaking temperature andthese carbides dissolved during soaking when the equilibriumcomposition of chromium and carbon in austenite was reachedAccording to the Fe-C phase diagram austenite can dissolvehigher carbon contents as the temperature increases to the eu-tectic

Fig 6 SEM micrographs with the 3-D structure of eutectic carbidesshowing the real morphology deep-etching with a solution of 50 mlFeCl3 plus 20 ml HCl in 930 ml ethanol for 3 h

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash375

As mentioned at 900 degC a higher secondary carbide volumefraction was observed at lower temperatures the ability ofaustenite to dissolve carbon is lower Therefore there is ahigher amount of carbon to precipitate which in turn is ob-served as a higher carbide volume fraction At higher tempera-tures austenite can retain carbon and chromium in higheramounts so there is very low raw material available to pre-cipitate carbides[ 5 8 9 ] and the final volume fraction is muchlower It is difficult to explain the diminishing of carbide vol-ume fraction observed at 1150 degC it is likely the equilibriumconcentration of carbon in austenite will be higher than theactual amount of this element in austenite so that during soak-ing precipitation is void Perhaps the volume fraction in thefirst sample (15 at 2 min) precipitated during heating to thesoaking temperature and these 2 min at 1150 degC were notenough for austenite to reach its equilibrium composition and

the secondary carbides did not dissolve Samples treated forlonger times (10 and 60 min) show a precipitated volume frac-tion of 2-5 perhaps the equilibrium volume in the alloy is at1150 degC

From micrographs in Fig 9 10 and 11 and particularlyfrom Fig 12 the morphology of secondary carbides can beobserved Secondary carbides are bars connected within thebulk matrix It has been reported[5 ] in white irons of similarcomposition that secondary carbides precipitated after destabi-lization treatments are of the type M7C3 having a hexagonalstructure and a preferred growing direction lt0001gt Figure 13shows some TEM micrographs of secondary carbides and adiffraction pattern of these particles which agree well with thefindings in Ref 5

34 Precipitated Carbides Size and Number

Figure 14 shows the particle size as a function of the soak-ing time these measurements are taken as the diameter of thecarbide bar It is well known that precipitation processes arefavored by temperature because temperature accelerates dif-fusion of the elements dissolved in matrix A higher diffusionrate of elements is observed for precipitation processes under-taken at high temperatures Because precipitation takes placeby nucleation and growth phenomena particles of a certain sizeare formed which depend on the treatment temperature In thecurrent study when the iron was treated at 900 degC the second-ary carbides diameter was about 120 nm within the first 2 hafterwards they coarsened up to 300 nm for 8 h soaking timeAt temperatures of 1000 degC the initial diameter is close to 200nm and gradually coarsens up to 250 nm for 8 h soaking timeAccording to this sequence it would be expected that the di-ameter of secondary carbides formed at 1150 degC to be higherhowever their size is only about 160 nm Because the volumeof material that precipitates in equilibrium as secondary car-

Fig 7 XRD pattern of the simple destabilized at 1000 degC for 1 h showing the detected phases austenite martensite M2C and M7C3

Fig 8 Volume of secondary carbides precipitated as a function ofthe soaking time during destabilization for each of the studied tem-peratures

376mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

bides at this temperature is almost void there is not enoughmaterial to precipitate coarse carbides

What is important to highlight is the precipitation behaviorat 900 and 1000 degC (Fig 15) This figure shows the numberprecipitated particles per mm2 At 900 degC the number of car-bides went from 2 particles per mm2 for 10 min soaking time to14 particles for 90 min Within the next 3 h the number ofcarbides remained at 14 and afterwards decreased to 8 particlesper mm2 for 8 h soaking time This finding indicates that theprecipitation process occurred during the first 90 min duringthe next 3 h there was no new precipitation just coarsening ofthe existing particles Finally during the next 4 h dissolutionof some carbides occurred contributing to the coarsening ofsome others thermodynamically more stable Similar phenom-ena during destabilization heat treatments have been reportedby Kuwano et al[ 8 ] All these phenomena are better observedfrom SEM micrographs (Fig 9 10 and 11) Note that from thecurrent study the secondary carbides precipitation was ob-served homogeneously in matrix According to Powell and

Bee[ 9 ] and Bee et al[ 2 1 ] secondary carbides precipitate at slipbands and sub-grain boundaries within austenite Such slipbands and sub-grain boundaries may form due to stresses gen-erated by differences in the thermal expansion coefficients ofcarbides and matrix On the contrary Kuwano et al[8 ] reportedthat the secondary carbides precipitation progressed from theperiphery adjacent to the eutectic carbides towards the centerof the dendrite in low carbon irons In high carbon alloys thecontrary occurredmdashprecipitation progressed from the center ofthe dendrite towards the periphery They attributed this phe-nomenon to chromium and carbon micro-segregation towardsthe periphery of the dendrites

Just a few works on the kinetics of secondary carbide pre-cipitation have been published Powell and Laird[ 5 ] suggestedthat precipitation occurs during the first 25 min at the soakingtemperature and keeping the alloy for longer periods of timeonly causes coarsening of carbides Kuwano et al[ 8 ] felt thatprecipitation occurs within very short periods of time such as 1min at temperatures between 950-1000 degC which disagree with

Fig 9 Series of SEM micrographs showing secondary carbide precipitation during destabilization heat treatment at 900 degC of the experimentalwhite iron (a) 20 min (b) 1 h (c) 2 h and (d) 8 h Villelarsquos etching 45 s Matrix is composed of martensite retained austenite and secondarycarbides

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash377

the findings from this work What has been observed andwidely accepted by most researchers is that once carbides haveformed longer periods of time at the treatment temperatureincrease the size and volume fraction of them Furthermorestill longer periods of time at the soaking temperature cause thecoarsening of carbides at the expense of othersrsquo dissolution(Ostwald ripening process) reducing the number of carbides[ 8 ]

This effect can be observed in Fig 15

35 Retained Austenite After Destabilization

Figure 16 shows the results of retained austenite as mea-sured by XRD As discussed above secondary carbide precipi-tation depletes the austenitic matrix in chromium and car-bon[ 4 1 0 ] Such depletion produces an increase in MS and it ismore likely to obtain martensitic structures Such an effect isobserved in Fig 16 at 900 degC and long soaking times a re-duction in the retained austenite content is observed from 48at 10 min to just 8 for 8 h soaking time This means an

increase in the martensite volume fraction At this temperatureand long soaking times the volume fraction of secondary car-bides was 27 The chromium- and carbon-depleted matrixtransformed almost totally to martensite during the subsequentcooling Something similar occurred during destabilization at1000 degC residual austenite decreased from 52 for 5 minsoaking time to 15 for 8 h In this case the volume fractionof secondary carbides was less than in the alloy treated at 900degC a clear indication is that chromium and carbon dissolved inmatrix are higher in the alloy treated at 1000 degC and suchelements contribute to lower the MS temperature Thereforethe alloy treated at 1000 degC showed a higher amount of residualaustenite These results are in agreement with the findings ofsome authors[ 5 8 9 ] in high- and low-chromium white irons Onthe contrary retained austenite increased with soaking time forthe iron treated at 1150 degC For a 5 min soaking time retainedaustenite was 25 and for 1 h the whole matrix became aus-tenitic after heat treatment In this case the high temperaturesallowed carbon and chromium to dissolve completely in matrixand inhibited the driving force for secondary carbides precipi-

Fig 10 Series of SEM micrographs showing secondary carbide precipitation during destabilization heat treatment at 1000 degC of the experimentalwhite iron (a) 5 min (b) 10 min (c) 1 h and (d) 6 h Villelarsquos etching 45 s Matrix is composed of martensite retained austenite and secondarycarbides

378mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

tation Chromium- and carbon-rich austenite is stable underthese conditions and remains stable during cooling to roomtemperature[1 2 2 2 3 ]

36 Microhardness of Matrix and Bulk Hardness of theAlloy After Destabilization

Figure 17 shows the microhardness values obtained for theironrsquos matrix after each heat treatment At 900 degC for 10 minmicrohardness was 800HV5 0 and increased to a maximum of1300HV5 0 after 2-3 h soaking time afterwards it decreased to900HV5 0 for a destabilization of 8 h The increase in micro-hardness of the matrix is influenced by two factors as thesoaking time increased the precipitation of secondary carbidesthat strengthens the matrix by particle dispersion together withthe increase in martensite volume fraction However dimin-ishing microhardness after 3 h soaking time is attributed to thesecondary carbide coarsening At this stage longer times pro-mote a major diffusion of elements favoring coarsening ofsome carbides at the expense of othersrsquo dissolution such aphenomenon leads to a diminishing number of particles in thematrix (Fig 15) The final structure of the matrix is thereforehighly martensitic with relatively few coarse secondary car-bides For the structures of the irons with maximum hardnessa highly martensitic matrix was also observed but it containeda higher number of fine well-dispersed secondary carbides

A similar effect occurred for the iron treated at 1000 degChowever maximum hardness (1159HV5 0 ) was obtained atshorter times (30-60 min) Higher destabilization temperaturesallowed higher diffusion leading to a reduced amount ofcoarse nuclei (Fig 14 and 15) Under this situation thestrengthening due to dispersed particles is lower Furthermoreat these higher temperatures austenite can dissolve higher car-bon contents so that the amount of secondary carbides is alsominor The above phenomena lead to a higher retained austen-ite content in the final structure which reduced matrix hard-ness

In contrast for the iron treated at 1150 degC hardness of thematrix diminished since the beginning For this case a stabi-lized austenitic matrix was obtained Such a matrix was solidsolution strengthened by chromium carbon nickel molybde-num etc but not by particle dispersion Secondary precipita-tion occurred only in the early stages of the heat treatment

Regarding bulk hardness the alloy showed a behavior simi-

Fig 11 Series of SEM micrographs showing secondary carbide pre-cipitation during destabilization heat treatment at 1150 degC of the ex-perimental white iron (a) 2 min matrix is composed of martensiteretained austenite and secondary carbides (b) 20 min matrix com-posed of austenite mainly and (c) 1 h matrix is totally austeniticVillelarsquos etching 45 s

Fig 12 SEM micrograph showing the real morphology of the sec-ondary carbides precipitated during destabilization treatments Deep-etching was with a solution of 50 ml FeCl3 plus 20 ml HCl in 930 mlethanol for 15 min

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash379

lar to that for matrix microhardness Figure 18 shows the re-sults for hardness where maximum values can be observed forthe different temperatures For the iron treated at 900 degC maxi-mum hardness was 68 HRC and was obtained for1-2 h soakingtime For the iron treated at 1000 degC maximum hardness was60 HRC obtained for 30 min Finally for the iron treated at1150 degC hardness remains constant (45 HRC) during the first20 min of treatment and decreases afterwards The strengthen-ing of the matrix is the main factor responsible for the modi-fications in the ironrsquos hardness Microhardness of matrix andthe eutectic carbide volume fraction are the factors that con-

tribute to strengthening of the alloys For this case the carbidevolume fraction was constant (24) therefore the strengthen-ing of matrix caused the changes in bulk hardness

These strengthening phenomena are similar to those foundin aging processes of some alloys and are explained in terms ofthe diffusion of elements that participate in the precipitatingphase Such diffusion in turn is a function of treatment tem-perature and soaking time Kuwano et al[ 8 ] Powell andLaird[ 5 ] and Powell and Bee[ 9 ] have found similar behaviorsfor high- and low-chromium irons in studies of precipitationand hardness

Fig 13 Bright field TEM micrographs showing some secondary carbides (SC) in (a) sample treated at 900 degC for 1 h and (b) sample treated at1000 degC for 1 h The selected area diffraction pattern (SADP) on the particle arrowed in (b) indicates the crystallographic nature of these carbidesto be hexagonal M7C3

Fig 14 Diameter of the precipitated particles as a function of thesoaking time during destabilization heat treatment

Fig 15 Number of precipitated particles in a square micron (mm2) asa function of the soaking time during destabilization heat treatment

380mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

4 Conclusions

The structure of the as-cast alloy was composed of 24eutectic carbides about 5 martensite and 70 austenitesmall amounts of the molybdenum-rich carbide Mo2Cwere also present

Vanadium carbide was not detected in the alloy nonethe-less the iron contained 2 vanadium a carbide-promotingelement Vanadium was observed to partition to the car-bide and in minor amounts to the matrix

According to the studied temperatures the secondary car-bide precipitation level decreased with higher tempera-tures being the highest for the samples treated at 900 degC(27 of matrix)

High temperatures like 1150 degC over-stabilized the aus-tenitic matrix which led to a secondary carbides free aus-tenitic matrix after heat treatment

The amount of precipitated particles showed an increasewith time for the samples treated at 900 and 1000 degCreaching a maximum and then decreased due to dissolutionof some carbides and coarsening of some others

Matrix hardness showed a maximum when plotted againstsoaking time due to these precipitation processes and tothe transformation of austenite to martensite Maximumhardness was obtained for the samples treated for 1-2 h at900 degC where the structure showed a high volume fractionof martensite and a high amount of fine well-dispersedsecondary carbide particles

Acknowledgments

The authors are grateful to the University of Michoacan forfunding this project

References

1 CP Tabrett IR Sare and MR Gomashchi ldquoMicrostructure-Property Relationships in High-Chromium White Iron Alloysrdquo IntMater Rev 1996 41(2) pp 59-82

2 ON Dogan JA Hawk and G Laird II ldquoSolidification Structure andAbrasion Resistance of High Chromium White Ironsrdquo Metall MaterTrans A 1997 28(6) pp 1315-27

3 G Laird II and GLF Powell ldquoSolidification and Solid-State Trans-formation Mechanisms in Si Alloyed High-Chromium White CastIronsrdquo Metall Trans A 1993 24(2) pp 981-88

4 WW Cias ldquoAustenite Transformation Kinetics and Hardenability ofHeat Treated 175Cr White Cast Ironsrdquo Trans AFS 1974 82 pp317-28

5 GLF Powell and G Laird II ldquoStructure Nucleation Growth andMorphology of Secondary Carbides in High Chromium and Cr-NiWhite Cast Ironsrdquo J Mater Sci 1992 27 pp 29-35

6 F Maratray and A Poulalion ldquoAustenite Retention in High-Chromium White Ironsrdquo Trans AFS 1982 90 pp 795-804

7 JTH Pearce ldquoStructure and Wear Performance of Abrasion ResistantChromium White Cast Ironsrdquo Trans AFS 1984 92 pp 599-621

8 M Kuwano K Ogi A Sawamoto and K Matsuda ldquoStudies onPrecipitation Process of Secondary Carbides in High-Chromium CastIronrdquo Trans AFS 1990 98 pp 725-34

9 GLF Powell and JV Bee ldquoSecondary Carbide Precipitation in an 18wtCr-1 wtMo White Ironrdquo J Mater Sci 1996 31 pp 707-11

10 C Kim ldquoX-Ray Method of Measuring Retained Austenite in HeatTreated White Cast Ironsrdquo J Heat Treating 1979 1(2) pp 43-51

11 A Bedolla-Jacuinde ldquoMicrostructure of Vanadium- Niobium- andTitanium-Alloyed High-Chromium White Cast Ironsrdquo Int J CastMetals Res 2001 13(6) pp 343-61

Fig 16 Retained austenite content in iron matrix as a function of thesoaking time during the destabilization heat treatment

Fig 17 Matrix microhardness as a function of the soaking timeduring the destabilization heat treatment

Fig 18 Bulk hardness of the alloy as a function of the soaking timeduring the destabilization heat treatment

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash381

As mentioned at 900 degC a higher secondary carbide volumefraction was observed at lower temperatures the ability ofaustenite to dissolve carbon is lower Therefore there is ahigher amount of carbon to precipitate which in turn is ob-served as a higher carbide volume fraction At higher tempera-tures austenite can retain carbon and chromium in higheramounts so there is very low raw material available to pre-cipitate carbides[ 5 8 9 ] and the final volume fraction is muchlower It is difficult to explain the diminishing of carbide vol-ume fraction observed at 1150 degC it is likely the equilibriumconcentration of carbon in austenite will be higher than theactual amount of this element in austenite so that during soak-ing precipitation is void Perhaps the volume fraction in thefirst sample (15 at 2 min) precipitated during heating to thesoaking temperature and these 2 min at 1150 degC were notenough for austenite to reach its equilibrium composition and

the secondary carbides did not dissolve Samples treated forlonger times (10 and 60 min) show a precipitated volume frac-tion of 2-5 perhaps the equilibrium volume in the alloy is at1150 degC

From micrographs in Fig 9 10 and 11 and particularlyfrom Fig 12 the morphology of secondary carbides can beobserved Secondary carbides are bars connected within thebulk matrix It has been reported[5 ] in white irons of similarcomposition that secondary carbides precipitated after destabi-lization treatments are of the type M7C3 having a hexagonalstructure and a preferred growing direction lt0001gt Figure 13shows some TEM micrographs of secondary carbides and adiffraction pattern of these particles which agree well with thefindings in Ref 5

34 Precipitated Carbides Size and Number

Figure 14 shows the particle size as a function of the soak-ing time these measurements are taken as the diameter of thecarbide bar It is well known that precipitation processes arefavored by temperature because temperature accelerates dif-fusion of the elements dissolved in matrix A higher diffusionrate of elements is observed for precipitation processes under-taken at high temperatures Because precipitation takes placeby nucleation and growth phenomena particles of a certain sizeare formed which depend on the treatment temperature In thecurrent study when the iron was treated at 900 degC the second-ary carbides diameter was about 120 nm within the first 2 hafterwards they coarsened up to 300 nm for 8 h soaking timeAt temperatures of 1000 degC the initial diameter is close to 200nm and gradually coarsens up to 250 nm for 8 h soaking timeAccording to this sequence it would be expected that the di-ameter of secondary carbides formed at 1150 degC to be higherhowever their size is only about 160 nm Because the volumeof material that precipitates in equilibrium as secondary car-

Fig 7 XRD pattern of the simple destabilized at 1000 degC for 1 h showing the detected phases austenite martensite M2C and M7C3

Fig 8 Volume of secondary carbides precipitated as a function ofthe soaking time during destabilization for each of the studied tem-peratures

376mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

bides at this temperature is almost void there is not enoughmaterial to precipitate coarse carbides

What is important to highlight is the precipitation behaviorat 900 and 1000 degC (Fig 15) This figure shows the numberprecipitated particles per mm2 At 900 degC the number of car-bides went from 2 particles per mm2 for 10 min soaking time to14 particles for 90 min Within the next 3 h the number ofcarbides remained at 14 and afterwards decreased to 8 particlesper mm2 for 8 h soaking time This finding indicates that theprecipitation process occurred during the first 90 min duringthe next 3 h there was no new precipitation just coarsening ofthe existing particles Finally during the next 4 h dissolutionof some carbides occurred contributing to the coarsening ofsome others thermodynamically more stable Similar phenom-ena during destabilization heat treatments have been reportedby Kuwano et al[ 8 ] All these phenomena are better observedfrom SEM micrographs (Fig 9 10 and 11) Note that from thecurrent study the secondary carbides precipitation was ob-served homogeneously in matrix According to Powell and

Bee[ 9 ] and Bee et al[ 2 1 ] secondary carbides precipitate at slipbands and sub-grain boundaries within austenite Such slipbands and sub-grain boundaries may form due to stresses gen-erated by differences in the thermal expansion coefficients ofcarbides and matrix On the contrary Kuwano et al[8 ] reportedthat the secondary carbides precipitation progressed from theperiphery adjacent to the eutectic carbides towards the centerof the dendrite in low carbon irons In high carbon alloys thecontrary occurredmdashprecipitation progressed from the center ofthe dendrite towards the periphery They attributed this phe-nomenon to chromium and carbon micro-segregation towardsthe periphery of the dendrites

Just a few works on the kinetics of secondary carbide pre-cipitation have been published Powell and Laird[ 5 ] suggestedthat precipitation occurs during the first 25 min at the soakingtemperature and keeping the alloy for longer periods of timeonly causes coarsening of carbides Kuwano et al[ 8 ] felt thatprecipitation occurs within very short periods of time such as 1min at temperatures between 950-1000 degC which disagree with

Fig 9 Series of SEM micrographs showing secondary carbide precipitation during destabilization heat treatment at 900 degC of the experimentalwhite iron (a) 20 min (b) 1 h (c) 2 h and (d) 8 h Villelarsquos etching 45 s Matrix is composed of martensite retained austenite and secondarycarbides

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash377

the findings from this work What has been observed andwidely accepted by most researchers is that once carbides haveformed longer periods of time at the treatment temperatureincrease the size and volume fraction of them Furthermorestill longer periods of time at the soaking temperature cause thecoarsening of carbides at the expense of othersrsquo dissolution(Ostwald ripening process) reducing the number of carbides[ 8 ]

This effect can be observed in Fig 15

35 Retained Austenite After Destabilization

Figure 16 shows the results of retained austenite as mea-sured by XRD As discussed above secondary carbide precipi-tation depletes the austenitic matrix in chromium and car-bon[ 4 1 0 ] Such depletion produces an increase in MS and it ismore likely to obtain martensitic structures Such an effect isobserved in Fig 16 at 900 degC and long soaking times a re-duction in the retained austenite content is observed from 48at 10 min to just 8 for 8 h soaking time This means an

increase in the martensite volume fraction At this temperatureand long soaking times the volume fraction of secondary car-bides was 27 The chromium- and carbon-depleted matrixtransformed almost totally to martensite during the subsequentcooling Something similar occurred during destabilization at1000 degC residual austenite decreased from 52 for 5 minsoaking time to 15 for 8 h In this case the volume fractionof secondary carbides was less than in the alloy treated at 900degC a clear indication is that chromium and carbon dissolved inmatrix are higher in the alloy treated at 1000 degC and suchelements contribute to lower the MS temperature Thereforethe alloy treated at 1000 degC showed a higher amount of residualaustenite These results are in agreement with the findings ofsome authors[ 5 8 9 ] in high- and low-chromium white irons Onthe contrary retained austenite increased with soaking time forthe iron treated at 1150 degC For a 5 min soaking time retainedaustenite was 25 and for 1 h the whole matrix became aus-tenitic after heat treatment In this case the high temperaturesallowed carbon and chromium to dissolve completely in matrixand inhibited the driving force for secondary carbides precipi-

Fig 10 Series of SEM micrographs showing secondary carbide precipitation during destabilization heat treatment at 1000 degC of the experimentalwhite iron (a) 5 min (b) 10 min (c) 1 h and (d) 6 h Villelarsquos etching 45 s Matrix is composed of martensite retained austenite and secondarycarbides

378mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

tation Chromium- and carbon-rich austenite is stable underthese conditions and remains stable during cooling to roomtemperature[1 2 2 2 3 ]

36 Microhardness of Matrix and Bulk Hardness of theAlloy After Destabilization

Figure 17 shows the microhardness values obtained for theironrsquos matrix after each heat treatment At 900 degC for 10 minmicrohardness was 800HV5 0 and increased to a maximum of1300HV5 0 after 2-3 h soaking time afterwards it decreased to900HV5 0 for a destabilization of 8 h The increase in micro-hardness of the matrix is influenced by two factors as thesoaking time increased the precipitation of secondary carbidesthat strengthens the matrix by particle dispersion together withthe increase in martensite volume fraction However dimin-ishing microhardness after 3 h soaking time is attributed to thesecondary carbide coarsening At this stage longer times pro-mote a major diffusion of elements favoring coarsening ofsome carbides at the expense of othersrsquo dissolution such aphenomenon leads to a diminishing number of particles in thematrix (Fig 15) The final structure of the matrix is thereforehighly martensitic with relatively few coarse secondary car-bides For the structures of the irons with maximum hardnessa highly martensitic matrix was also observed but it containeda higher number of fine well-dispersed secondary carbides

A similar effect occurred for the iron treated at 1000 degChowever maximum hardness (1159HV5 0 ) was obtained atshorter times (30-60 min) Higher destabilization temperaturesallowed higher diffusion leading to a reduced amount ofcoarse nuclei (Fig 14 and 15) Under this situation thestrengthening due to dispersed particles is lower Furthermoreat these higher temperatures austenite can dissolve higher car-bon contents so that the amount of secondary carbides is alsominor The above phenomena lead to a higher retained austen-ite content in the final structure which reduced matrix hard-ness

In contrast for the iron treated at 1150 degC hardness of thematrix diminished since the beginning For this case a stabi-lized austenitic matrix was obtained Such a matrix was solidsolution strengthened by chromium carbon nickel molybde-num etc but not by particle dispersion Secondary precipita-tion occurred only in the early stages of the heat treatment

Regarding bulk hardness the alloy showed a behavior simi-

Fig 11 Series of SEM micrographs showing secondary carbide pre-cipitation during destabilization heat treatment at 1150 degC of the ex-perimental white iron (a) 2 min matrix is composed of martensiteretained austenite and secondary carbides (b) 20 min matrix com-posed of austenite mainly and (c) 1 h matrix is totally austeniticVillelarsquos etching 45 s

Fig 12 SEM micrograph showing the real morphology of the sec-ondary carbides precipitated during destabilization treatments Deep-etching was with a solution of 50 ml FeCl3 plus 20 ml HCl in 930 mlethanol for 15 min

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash379

lar to that for matrix microhardness Figure 18 shows the re-sults for hardness where maximum values can be observed forthe different temperatures For the iron treated at 900 degC maxi-mum hardness was 68 HRC and was obtained for1-2 h soakingtime For the iron treated at 1000 degC maximum hardness was60 HRC obtained for 30 min Finally for the iron treated at1150 degC hardness remains constant (45 HRC) during the first20 min of treatment and decreases afterwards The strengthen-ing of the matrix is the main factor responsible for the modi-fications in the ironrsquos hardness Microhardness of matrix andthe eutectic carbide volume fraction are the factors that con-

tribute to strengthening of the alloys For this case the carbidevolume fraction was constant (24) therefore the strengthen-ing of matrix caused the changes in bulk hardness

These strengthening phenomena are similar to those foundin aging processes of some alloys and are explained in terms ofthe diffusion of elements that participate in the precipitatingphase Such diffusion in turn is a function of treatment tem-perature and soaking time Kuwano et al[ 8 ] Powell andLaird[ 5 ] and Powell and Bee[ 9 ] have found similar behaviorsfor high- and low-chromium irons in studies of precipitationand hardness

Fig 13 Bright field TEM micrographs showing some secondary carbides (SC) in (a) sample treated at 900 degC for 1 h and (b) sample treated at1000 degC for 1 h The selected area diffraction pattern (SADP) on the particle arrowed in (b) indicates the crystallographic nature of these carbidesto be hexagonal M7C3

Fig 14 Diameter of the precipitated particles as a function of thesoaking time during destabilization heat treatment

Fig 15 Number of precipitated particles in a square micron (mm2) asa function of the soaking time during destabilization heat treatment

380mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

4 Conclusions

The structure of the as-cast alloy was composed of 24eutectic carbides about 5 martensite and 70 austenitesmall amounts of the molybdenum-rich carbide Mo2Cwere also present

Vanadium carbide was not detected in the alloy nonethe-less the iron contained 2 vanadium a carbide-promotingelement Vanadium was observed to partition to the car-bide and in minor amounts to the matrix

According to the studied temperatures the secondary car-bide precipitation level decreased with higher tempera-tures being the highest for the samples treated at 900 degC(27 of matrix)

High temperatures like 1150 degC over-stabilized the aus-tenitic matrix which led to a secondary carbides free aus-tenitic matrix after heat treatment

The amount of precipitated particles showed an increasewith time for the samples treated at 900 and 1000 degCreaching a maximum and then decreased due to dissolutionof some carbides and coarsening of some others

Matrix hardness showed a maximum when plotted againstsoaking time due to these precipitation processes and tothe transformation of austenite to martensite Maximumhardness was obtained for the samples treated for 1-2 h at900 degC where the structure showed a high volume fractionof martensite and a high amount of fine well-dispersedsecondary carbide particles

Acknowledgments

The authors are grateful to the University of Michoacan forfunding this project

References

1 CP Tabrett IR Sare and MR Gomashchi ldquoMicrostructure-Property Relationships in High-Chromium White Iron Alloysrdquo IntMater Rev 1996 41(2) pp 59-82

2 ON Dogan JA Hawk and G Laird II ldquoSolidification Structure andAbrasion Resistance of High Chromium White Ironsrdquo Metall MaterTrans A 1997 28(6) pp 1315-27

3 G Laird II and GLF Powell ldquoSolidification and Solid-State Trans-formation Mechanisms in Si Alloyed High-Chromium White CastIronsrdquo Metall Trans A 1993 24(2) pp 981-88

4 WW Cias ldquoAustenite Transformation Kinetics and Hardenability ofHeat Treated 175Cr White Cast Ironsrdquo Trans AFS 1974 82 pp317-28

5 GLF Powell and G Laird II ldquoStructure Nucleation Growth andMorphology of Secondary Carbides in High Chromium and Cr-NiWhite Cast Ironsrdquo J Mater Sci 1992 27 pp 29-35

6 F Maratray and A Poulalion ldquoAustenite Retention in High-Chromium White Ironsrdquo Trans AFS 1982 90 pp 795-804

7 JTH Pearce ldquoStructure and Wear Performance of Abrasion ResistantChromium White Cast Ironsrdquo Trans AFS 1984 92 pp 599-621

8 M Kuwano K Ogi A Sawamoto and K Matsuda ldquoStudies onPrecipitation Process of Secondary Carbides in High-Chromium CastIronrdquo Trans AFS 1990 98 pp 725-34

9 GLF Powell and JV Bee ldquoSecondary Carbide Precipitation in an 18wtCr-1 wtMo White Ironrdquo J Mater Sci 1996 31 pp 707-11

10 C Kim ldquoX-Ray Method of Measuring Retained Austenite in HeatTreated White Cast Ironsrdquo J Heat Treating 1979 1(2) pp 43-51

11 A Bedolla-Jacuinde ldquoMicrostructure of Vanadium- Niobium- andTitanium-Alloyed High-Chromium White Cast Ironsrdquo Int J CastMetals Res 2001 13(6) pp 343-61

Fig 16 Retained austenite content in iron matrix as a function of thesoaking time during the destabilization heat treatment

Fig 17 Matrix microhardness as a function of the soaking timeduring the destabilization heat treatment

Fig 18 Bulk hardness of the alloy as a function of the soaking timeduring the destabilization heat treatment

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash381

bides at this temperature is almost void there is not enoughmaterial to precipitate coarse carbides

What is important to highlight is the precipitation behaviorat 900 and 1000 degC (Fig 15) This figure shows the numberprecipitated particles per mm2 At 900 degC the number of car-bides went from 2 particles per mm2 for 10 min soaking time to14 particles for 90 min Within the next 3 h the number ofcarbides remained at 14 and afterwards decreased to 8 particlesper mm2 for 8 h soaking time This finding indicates that theprecipitation process occurred during the first 90 min duringthe next 3 h there was no new precipitation just coarsening ofthe existing particles Finally during the next 4 h dissolutionof some carbides occurred contributing to the coarsening ofsome others thermodynamically more stable Similar phenom-ena during destabilization heat treatments have been reportedby Kuwano et al[ 8 ] All these phenomena are better observedfrom SEM micrographs (Fig 9 10 and 11) Note that from thecurrent study the secondary carbides precipitation was ob-served homogeneously in matrix According to Powell and

Bee[ 9 ] and Bee et al[ 2 1 ] secondary carbides precipitate at slipbands and sub-grain boundaries within austenite Such slipbands and sub-grain boundaries may form due to stresses gen-erated by differences in the thermal expansion coefficients ofcarbides and matrix On the contrary Kuwano et al[8 ] reportedthat the secondary carbides precipitation progressed from theperiphery adjacent to the eutectic carbides towards the centerof the dendrite in low carbon irons In high carbon alloys thecontrary occurredmdashprecipitation progressed from the center ofthe dendrite towards the periphery They attributed this phe-nomenon to chromium and carbon micro-segregation towardsthe periphery of the dendrites

Just a few works on the kinetics of secondary carbide pre-cipitation have been published Powell and Laird[ 5 ] suggestedthat precipitation occurs during the first 25 min at the soakingtemperature and keeping the alloy for longer periods of timeonly causes coarsening of carbides Kuwano et al[ 8 ] felt thatprecipitation occurs within very short periods of time such as 1min at temperatures between 950-1000 degC which disagree with

Fig 9 Series of SEM micrographs showing secondary carbide precipitation during destabilization heat treatment at 900 degC of the experimentalwhite iron (a) 20 min (b) 1 h (c) 2 h and (d) 8 h Villelarsquos etching 45 s Matrix is composed of martensite retained austenite and secondarycarbides

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash377

the findings from this work What has been observed andwidely accepted by most researchers is that once carbides haveformed longer periods of time at the treatment temperatureincrease the size and volume fraction of them Furthermorestill longer periods of time at the soaking temperature cause thecoarsening of carbides at the expense of othersrsquo dissolution(Ostwald ripening process) reducing the number of carbides[ 8 ]

This effect can be observed in Fig 15

35 Retained Austenite After Destabilization

Figure 16 shows the results of retained austenite as mea-sured by XRD As discussed above secondary carbide precipi-tation depletes the austenitic matrix in chromium and car-bon[ 4 1 0 ] Such depletion produces an increase in MS and it ismore likely to obtain martensitic structures Such an effect isobserved in Fig 16 at 900 degC and long soaking times a re-duction in the retained austenite content is observed from 48at 10 min to just 8 for 8 h soaking time This means an

increase in the martensite volume fraction At this temperatureand long soaking times the volume fraction of secondary car-bides was 27 The chromium- and carbon-depleted matrixtransformed almost totally to martensite during the subsequentcooling Something similar occurred during destabilization at1000 degC residual austenite decreased from 52 for 5 minsoaking time to 15 for 8 h In this case the volume fractionof secondary carbides was less than in the alloy treated at 900degC a clear indication is that chromium and carbon dissolved inmatrix are higher in the alloy treated at 1000 degC and suchelements contribute to lower the MS temperature Thereforethe alloy treated at 1000 degC showed a higher amount of residualaustenite These results are in agreement with the findings ofsome authors[ 5 8 9 ] in high- and low-chromium white irons Onthe contrary retained austenite increased with soaking time forthe iron treated at 1150 degC For a 5 min soaking time retainedaustenite was 25 and for 1 h the whole matrix became aus-tenitic after heat treatment In this case the high temperaturesallowed carbon and chromium to dissolve completely in matrixand inhibited the driving force for secondary carbides precipi-

Fig 10 Series of SEM micrographs showing secondary carbide precipitation during destabilization heat treatment at 1000 degC of the experimentalwhite iron (a) 5 min (b) 10 min (c) 1 h and (d) 6 h Villelarsquos etching 45 s Matrix is composed of martensite retained austenite and secondarycarbides

378mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

tation Chromium- and carbon-rich austenite is stable underthese conditions and remains stable during cooling to roomtemperature[1 2 2 2 3 ]

36 Microhardness of Matrix and Bulk Hardness of theAlloy After Destabilization

Figure 17 shows the microhardness values obtained for theironrsquos matrix after each heat treatment At 900 degC for 10 minmicrohardness was 800HV5 0 and increased to a maximum of1300HV5 0 after 2-3 h soaking time afterwards it decreased to900HV5 0 for a destabilization of 8 h The increase in micro-hardness of the matrix is influenced by two factors as thesoaking time increased the precipitation of secondary carbidesthat strengthens the matrix by particle dispersion together withthe increase in martensite volume fraction However dimin-ishing microhardness after 3 h soaking time is attributed to thesecondary carbide coarsening At this stage longer times pro-mote a major diffusion of elements favoring coarsening ofsome carbides at the expense of othersrsquo dissolution such aphenomenon leads to a diminishing number of particles in thematrix (Fig 15) The final structure of the matrix is thereforehighly martensitic with relatively few coarse secondary car-bides For the structures of the irons with maximum hardnessa highly martensitic matrix was also observed but it containeda higher number of fine well-dispersed secondary carbides

A similar effect occurred for the iron treated at 1000 degChowever maximum hardness (1159HV5 0 ) was obtained atshorter times (30-60 min) Higher destabilization temperaturesallowed higher diffusion leading to a reduced amount ofcoarse nuclei (Fig 14 and 15) Under this situation thestrengthening due to dispersed particles is lower Furthermoreat these higher temperatures austenite can dissolve higher car-bon contents so that the amount of secondary carbides is alsominor The above phenomena lead to a higher retained austen-ite content in the final structure which reduced matrix hard-ness

In contrast for the iron treated at 1150 degC hardness of thematrix diminished since the beginning For this case a stabi-lized austenitic matrix was obtained Such a matrix was solidsolution strengthened by chromium carbon nickel molybde-num etc but not by particle dispersion Secondary precipita-tion occurred only in the early stages of the heat treatment

Regarding bulk hardness the alloy showed a behavior simi-

Fig 11 Series of SEM micrographs showing secondary carbide pre-cipitation during destabilization heat treatment at 1150 degC of the ex-perimental white iron (a) 2 min matrix is composed of martensiteretained austenite and secondary carbides (b) 20 min matrix com-posed of austenite mainly and (c) 1 h matrix is totally austeniticVillelarsquos etching 45 s

Fig 12 SEM micrograph showing the real morphology of the sec-ondary carbides precipitated during destabilization treatments Deep-etching was with a solution of 50 ml FeCl3 plus 20 ml HCl in 930 mlethanol for 15 min

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash379

lar to that for matrix microhardness Figure 18 shows the re-sults for hardness where maximum values can be observed forthe different temperatures For the iron treated at 900 degC maxi-mum hardness was 68 HRC and was obtained for1-2 h soakingtime For the iron treated at 1000 degC maximum hardness was60 HRC obtained for 30 min Finally for the iron treated at1150 degC hardness remains constant (45 HRC) during the first20 min of treatment and decreases afterwards The strengthen-ing of the matrix is the main factor responsible for the modi-fications in the ironrsquos hardness Microhardness of matrix andthe eutectic carbide volume fraction are the factors that con-

tribute to strengthening of the alloys For this case the carbidevolume fraction was constant (24) therefore the strengthen-ing of matrix caused the changes in bulk hardness

These strengthening phenomena are similar to those foundin aging processes of some alloys and are explained in terms ofthe diffusion of elements that participate in the precipitatingphase Such diffusion in turn is a function of treatment tem-perature and soaking time Kuwano et al[ 8 ] Powell andLaird[ 5 ] and Powell and Bee[ 9 ] have found similar behaviorsfor high- and low-chromium irons in studies of precipitationand hardness

Fig 13 Bright field TEM micrographs showing some secondary carbides (SC) in (a) sample treated at 900 degC for 1 h and (b) sample treated at1000 degC for 1 h The selected area diffraction pattern (SADP) on the particle arrowed in (b) indicates the crystallographic nature of these carbidesto be hexagonal M7C3

Fig 14 Diameter of the precipitated particles as a function of thesoaking time during destabilization heat treatment

Fig 15 Number of precipitated particles in a square micron (mm2) asa function of the soaking time during destabilization heat treatment

380mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

4 Conclusions

The structure of the as-cast alloy was composed of 24eutectic carbides about 5 martensite and 70 austenitesmall amounts of the molybdenum-rich carbide Mo2Cwere also present

Vanadium carbide was not detected in the alloy nonethe-less the iron contained 2 vanadium a carbide-promotingelement Vanadium was observed to partition to the car-bide and in minor amounts to the matrix

According to the studied temperatures the secondary car-bide precipitation level decreased with higher tempera-tures being the highest for the samples treated at 900 degC(27 of matrix)

High temperatures like 1150 degC over-stabilized the aus-tenitic matrix which led to a secondary carbides free aus-tenitic matrix after heat treatment

The amount of precipitated particles showed an increasewith time for the samples treated at 900 and 1000 degCreaching a maximum and then decreased due to dissolutionof some carbides and coarsening of some others

Matrix hardness showed a maximum when plotted againstsoaking time due to these precipitation processes and tothe transformation of austenite to martensite Maximumhardness was obtained for the samples treated for 1-2 h at900 degC where the structure showed a high volume fractionof martensite and a high amount of fine well-dispersedsecondary carbide particles

Acknowledgments

The authors are grateful to the University of Michoacan forfunding this project

References

1 CP Tabrett IR Sare and MR Gomashchi ldquoMicrostructure-Property Relationships in High-Chromium White Iron Alloysrdquo IntMater Rev 1996 41(2) pp 59-82

2 ON Dogan JA Hawk and G Laird II ldquoSolidification Structure andAbrasion Resistance of High Chromium White Ironsrdquo Metall MaterTrans A 1997 28(6) pp 1315-27

3 G Laird II and GLF Powell ldquoSolidification and Solid-State Trans-formation Mechanisms in Si Alloyed High-Chromium White CastIronsrdquo Metall Trans A 1993 24(2) pp 981-88

4 WW Cias ldquoAustenite Transformation Kinetics and Hardenability ofHeat Treated 175Cr White Cast Ironsrdquo Trans AFS 1974 82 pp317-28

5 GLF Powell and G Laird II ldquoStructure Nucleation Growth andMorphology of Secondary Carbides in High Chromium and Cr-NiWhite Cast Ironsrdquo J Mater Sci 1992 27 pp 29-35

6 F Maratray and A Poulalion ldquoAustenite Retention in High-Chromium White Ironsrdquo Trans AFS 1982 90 pp 795-804

7 JTH Pearce ldquoStructure and Wear Performance of Abrasion ResistantChromium White Cast Ironsrdquo Trans AFS 1984 92 pp 599-621

8 M Kuwano K Ogi A Sawamoto and K Matsuda ldquoStudies onPrecipitation Process of Secondary Carbides in High-Chromium CastIronrdquo Trans AFS 1990 98 pp 725-34

9 GLF Powell and JV Bee ldquoSecondary Carbide Precipitation in an 18wtCr-1 wtMo White Ironrdquo J Mater Sci 1996 31 pp 707-11

10 C Kim ldquoX-Ray Method of Measuring Retained Austenite in HeatTreated White Cast Ironsrdquo J Heat Treating 1979 1(2) pp 43-51

11 A Bedolla-Jacuinde ldquoMicrostructure of Vanadium- Niobium- andTitanium-Alloyed High-Chromium White Cast Ironsrdquo Int J CastMetals Res 2001 13(6) pp 343-61

Fig 16 Retained austenite content in iron matrix as a function of thesoaking time during the destabilization heat treatment

Fig 17 Matrix microhardness as a function of the soaking timeduring the destabilization heat treatment

Fig 18 Bulk hardness of the alloy as a function of the soaking timeduring the destabilization heat treatment

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash381

the findings from this work What has been observed andwidely accepted by most researchers is that once carbides haveformed longer periods of time at the treatment temperatureincrease the size and volume fraction of them Furthermorestill longer periods of time at the soaking temperature cause thecoarsening of carbides at the expense of othersrsquo dissolution(Ostwald ripening process) reducing the number of carbides[ 8 ]

This effect can be observed in Fig 15

35 Retained Austenite After Destabilization

Figure 16 shows the results of retained austenite as mea-sured by XRD As discussed above secondary carbide precipi-tation depletes the austenitic matrix in chromium and car-bon[ 4 1 0 ] Such depletion produces an increase in MS and it ismore likely to obtain martensitic structures Such an effect isobserved in Fig 16 at 900 degC and long soaking times a re-duction in the retained austenite content is observed from 48at 10 min to just 8 for 8 h soaking time This means an

increase in the martensite volume fraction At this temperatureand long soaking times the volume fraction of secondary car-bides was 27 The chromium- and carbon-depleted matrixtransformed almost totally to martensite during the subsequentcooling Something similar occurred during destabilization at1000 degC residual austenite decreased from 52 for 5 minsoaking time to 15 for 8 h In this case the volume fractionof secondary carbides was less than in the alloy treated at 900degC a clear indication is that chromium and carbon dissolved inmatrix are higher in the alloy treated at 1000 degC and suchelements contribute to lower the MS temperature Thereforethe alloy treated at 1000 degC showed a higher amount of residualaustenite These results are in agreement with the findings ofsome authors[ 5 8 9 ] in high- and low-chromium white irons Onthe contrary retained austenite increased with soaking time forthe iron treated at 1150 degC For a 5 min soaking time retainedaustenite was 25 and for 1 h the whole matrix became aus-tenitic after heat treatment In this case the high temperaturesallowed carbon and chromium to dissolve completely in matrixand inhibited the driving force for secondary carbides precipi-

Fig 10 Series of SEM micrographs showing secondary carbide precipitation during destabilization heat treatment at 1000 degC of the experimentalwhite iron (a) 5 min (b) 10 min (c) 1 h and (d) 6 h Villelarsquos etching 45 s Matrix is composed of martensite retained austenite and secondarycarbides

378mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

tation Chromium- and carbon-rich austenite is stable underthese conditions and remains stable during cooling to roomtemperature[1 2 2 2 3 ]

36 Microhardness of Matrix and Bulk Hardness of theAlloy After Destabilization

Figure 17 shows the microhardness values obtained for theironrsquos matrix after each heat treatment At 900 degC for 10 minmicrohardness was 800HV5 0 and increased to a maximum of1300HV5 0 after 2-3 h soaking time afterwards it decreased to900HV5 0 for a destabilization of 8 h The increase in micro-hardness of the matrix is influenced by two factors as thesoaking time increased the precipitation of secondary carbidesthat strengthens the matrix by particle dispersion together withthe increase in martensite volume fraction However dimin-ishing microhardness after 3 h soaking time is attributed to thesecondary carbide coarsening At this stage longer times pro-mote a major diffusion of elements favoring coarsening ofsome carbides at the expense of othersrsquo dissolution such aphenomenon leads to a diminishing number of particles in thematrix (Fig 15) The final structure of the matrix is thereforehighly martensitic with relatively few coarse secondary car-bides For the structures of the irons with maximum hardnessa highly martensitic matrix was also observed but it containeda higher number of fine well-dispersed secondary carbides

A similar effect occurred for the iron treated at 1000 degChowever maximum hardness (1159HV5 0 ) was obtained atshorter times (30-60 min) Higher destabilization temperaturesallowed higher diffusion leading to a reduced amount ofcoarse nuclei (Fig 14 and 15) Under this situation thestrengthening due to dispersed particles is lower Furthermoreat these higher temperatures austenite can dissolve higher car-bon contents so that the amount of secondary carbides is alsominor The above phenomena lead to a higher retained austen-ite content in the final structure which reduced matrix hard-ness

In contrast for the iron treated at 1150 degC hardness of thematrix diminished since the beginning For this case a stabi-lized austenitic matrix was obtained Such a matrix was solidsolution strengthened by chromium carbon nickel molybde-num etc but not by particle dispersion Secondary precipita-tion occurred only in the early stages of the heat treatment

Regarding bulk hardness the alloy showed a behavior simi-

Fig 11 Series of SEM micrographs showing secondary carbide pre-cipitation during destabilization heat treatment at 1150 degC of the ex-perimental white iron (a) 2 min matrix is composed of martensiteretained austenite and secondary carbides (b) 20 min matrix com-posed of austenite mainly and (c) 1 h matrix is totally austeniticVillelarsquos etching 45 s

Fig 12 SEM micrograph showing the real morphology of the sec-ondary carbides precipitated during destabilization treatments Deep-etching was with a solution of 50 ml FeCl3 plus 20 ml HCl in 930 mlethanol for 15 min

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash379

lar to that for matrix microhardness Figure 18 shows the re-sults for hardness where maximum values can be observed forthe different temperatures For the iron treated at 900 degC maxi-mum hardness was 68 HRC and was obtained for1-2 h soakingtime For the iron treated at 1000 degC maximum hardness was60 HRC obtained for 30 min Finally for the iron treated at1150 degC hardness remains constant (45 HRC) during the first20 min of treatment and decreases afterwards The strengthen-ing of the matrix is the main factor responsible for the modi-fications in the ironrsquos hardness Microhardness of matrix andthe eutectic carbide volume fraction are the factors that con-

tribute to strengthening of the alloys For this case the carbidevolume fraction was constant (24) therefore the strengthen-ing of matrix caused the changes in bulk hardness

These strengthening phenomena are similar to those foundin aging processes of some alloys and are explained in terms ofthe diffusion of elements that participate in the precipitatingphase Such diffusion in turn is a function of treatment tem-perature and soaking time Kuwano et al[ 8 ] Powell andLaird[ 5 ] and Powell and Bee[ 9 ] have found similar behaviorsfor high- and low-chromium irons in studies of precipitationand hardness

Fig 13 Bright field TEM micrographs showing some secondary carbides (SC) in (a) sample treated at 900 degC for 1 h and (b) sample treated at1000 degC for 1 h The selected area diffraction pattern (SADP) on the particle arrowed in (b) indicates the crystallographic nature of these carbidesto be hexagonal M7C3

Fig 14 Diameter of the precipitated particles as a function of thesoaking time during destabilization heat treatment

Fig 15 Number of precipitated particles in a square micron (mm2) asa function of the soaking time during destabilization heat treatment

380mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

4 Conclusions

The structure of the as-cast alloy was composed of 24eutectic carbides about 5 martensite and 70 austenitesmall amounts of the molybdenum-rich carbide Mo2Cwere also present

Vanadium carbide was not detected in the alloy nonethe-less the iron contained 2 vanadium a carbide-promotingelement Vanadium was observed to partition to the car-bide and in minor amounts to the matrix

According to the studied temperatures the secondary car-bide precipitation level decreased with higher tempera-tures being the highest for the samples treated at 900 degC(27 of matrix)

High temperatures like 1150 degC over-stabilized the aus-tenitic matrix which led to a secondary carbides free aus-tenitic matrix after heat treatment

The amount of precipitated particles showed an increasewith time for the samples treated at 900 and 1000 degCreaching a maximum and then decreased due to dissolutionof some carbides and coarsening of some others

Matrix hardness showed a maximum when plotted againstsoaking time due to these precipitation processes and tothe transformation of austenite to martensite Maximumhardness was obtained for the samples treated for 1-2 h at900 degC where the structure showed a high volume fractionof martensite and a high amount of fine well-dispersedsecondary carbide particles

Acknowledgments

The authors are grateful to the University of Michoacan forfunding this project

References

1 CP Tabrett IR Sare and MR Gomashchi ldquoMicrostructure-Property Relationships in High-Chromium White Iron Alloysrdquo IntMater Rev 1996 41(2) pp 59-82

2 ON Dogan JA Hawk and G Laird II ldquoSolidification Structure andAbrasion Resistance of High Chromium White Ironsrdquo Metall MaterTrans A 1997 28(6) pp 1315-27

3 G Laird II and GLF Powell ldquoSolidification and Solid-State Trans-formation Mechanisms in Si Alloyed High-Chromium White CastIronsrdquo Metall Trans A 1993 24(2) pp 981-88

4 WW Cias ldquoAustenite Transformation Kinetics and Hardenability ofHeat Treated 175Cr White Cast Ironsrdquo Trans AFS 1974 82 pp317-28

5 GLF Powell and G Laird II ldquoStructure Nucleation Growth andMorphology of Secondary Carbides in High Chromium and Cr-NiWhite Cast Ironsrdquo J Mater Sci 1992 27 pp 29-35

6 F Maratray and A Poulalion ldquoAustenite Retention in High-Chromium White Ironsrdquo Trans AFS 1982 90 pp 795-804

7 JTH Pearce ldquoStructure and Wear Performance of Abrasion ResistantChromium White Cast Ironsrdquo Trans AFS 1984 92 pp 599-621

8 M Kuwano K Ogi A Sawamoto and K Matsuda ldquoStudies onPrecipitation Process of Secondary Carbides in High-Chromium CastIronrdquo Trans AFS 1990 98 pp 725-34

9 GLF Powell and JV Bee ldquoSecondary Carbide Precipitation in an 18wtCr-1 wtMo White Ironrdquo J Mater Sci 1996 31 pp 707-11

10 C Kim ldquoX-Ray Method of Measuring Retained Austenite in HeatTreated White Cast Ironsrdquo J Heat Treating 1979 1(2) pp 43-51

11 A Bedolla-Jacuinde ldquoMicrostructure of Vanadium- Niobium- andTitanium-Alloyed High-Chromium White Cast Ironsrdquo Int J CastMetals Res 2001 13(6) pp 343-61

Fig 16 Retained austenite content in iron matrix as a function of thesoaking time during the destabilization heat treatment

Fig 17 Matrix microhardness as a function of the soaking timeduring the destabilization heat treatment

Fig 18 Bulk hardness of the alloy as a function of the soaking timeduring the destabilization heat treatment

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash381

tation Chromium- and carbon-rich austenite is stable underthese conditions and remains stable during cooling to roomtemperature[1 2 2 2 3 ]

36 Microhardness of Matrix and Bulk Hardness of theAlloy After Destabilization

Figure 17 shows the microhardness values obtained for theironrsquos matrix after each heat treatment At 900 degC for 10 minmicrohardness was 800HV5 0 and increased to a maximum of1300HV5 0 after 2-3 h soaking time afterwards it decreased to900HV5 0 for a destabilization of 8 h The increase in micro-hardness of the matrix is influenced by two factors as thesoaking time increased the precipitation of secondary carbidesthat strengthens the matrix by particle dispersion together withthe increase in martensite volume fraction However dimin-ishing microhardness after 3 h soaking time is attributed to thesecondary carbide coarsening At this stage longer times pro-mote a major diffusion of elements favoring coarsening ofsome carbides at the expense of othersrsquo dissolution such aphenomenon leads to a diminishing number of particles in thematrix (Fig 15) The final structure of the matrix is thereforehighly martensitic with relatively few coarse secondary car-bides For the structures of the irons with maximum hardnessa highly martensitic matrix was also observed but it containeda higher number of fine well-dispersed secondary carbides

A similar effect occurred for the iron treated at 1000 degChowever maximum hardness (1159HV5 0 ) was obtained atshorter times (30-60 min) Higher destabilization temperaturesallowed higher diffusion leading to a reduced amount ofcoarse nuclei (Fig 14 and 15) Under this situation thestrengthening due to dispersed particles is lower Furthermoreat these higher temperatures austenite can dissolve higher car-bon contents so that the amount of secondary carbides is alsominor The above phenomena lead to a higher retained austen-ite content in the final structure which reduced matrix hard-ness

In contrast for the iron treated at 1150 degC hardness of thematrix diminished since the beginning For this case a stabi-lized austenitic matrix was obtained Such a matrix was solidsolution strengthened by chromium carbon nickel molybde-num etc but not by particle dispersion Secondary precipita-tion occurred only in the early stages of the heat treatment

Regarding bulk hardness the alloy showed a behavior simi-

Fig 11 Series of SEM micrographs showing secondary carbide pre-cipitation during destabilization heat treatment at 1150 degC of the ex-perimental white iron (a) 2 min matrix is composed of martensiteretained austenite and secondary carbides (b) 20 min matrix com-posed of austenite mainly and (c) 1 h matrix is totally austeniticVillelarsquos etching 45 s

Fig 12 SEM micrograph showing the real morphology of the sec-ondary carbides precipitated during destabilization treatments Deep-etching was with a solution of 50 ml FeCl3 plus 20 ml HCl in 930 mlethanol for 15 min

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash379

lar to that for matrix microhardness Figure 18 shows the re-sults for hardness where maximum values can be observed forthe different temperatures For the iron treated at 900 degC maxi-mum hardness was 68 HRC and was obtained for1-2 h soakingtime For the iron treated at 1000 degC maximum hardness was60 HRC obtained for 30 min Finally for the iron treated at1150 degC hardness remains constant (45 HRC) during the first20 min of treatment and decreases afterwards The strengthen-ing of the matrix is the main factor responsible for the modi-fications in the ironrsquos hardness Microhardness of matrix andthe eutectic carbide volume fraction are the factors that con-

tribute to strengthening of the alloys For this case the carbidevolume fraction was constant (24) therefore the strengthen-ing of matrix caused the changes in bulk hardness

These strengthening phenomena are similar to those foundin aging processes of some alloys and are explained in terms ofthe diffusion of elements that participate in the precipitatingphase Such diffusion in turn is a function of treatment tem-perature and soaking time Kuwano et al[ 8 ] Powell andLaird[ 5 ] and Powell and Bee[ 9 ] have found similar behaviorsfor high- and low-chromium irons in studies of precipitationand hardness

Fig 13 Bright field TEM micrographs showing some secondary carbides (SC) in (a) sample treated at 900 degC for 1 h and (b) sample treated at1000 degC for 1 h The selected area diffraction pattern (SADP) on the particle arrowed in (b) indicates the crystallographic nature of these carbidesto be hexagonal M7C3

Fig 14 Diameter of the precipitated particles as a function of thesoaking time during destabilization heat treatment

Fig 15 Number of precipitated particles in a square micron (mm2) asa function of the soaking time during destabilization heat treatment

380mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

4 Conclusions

The structure of the as-cast alloy was composed of 24eutectic carbides about 5 martensite and 70 austenitesmall amounts of the molybdenum-rich carbide Mo2Cwere also present

Vanadium carbide was not detected in the alloy nonethe-less the iron contained 2 vanadium a carbide-promotingelement Vanadium was observed to partition to the car-bide and in minor amounts to the matrix

According to the studied temperatures the secondary car-bide precipitation level decreased with higher tempera-tures being the highest for the samples treated at 900 degC(27 of matrix)

High temperatures like 1150 degC over-stabilized the aus-tenitic matrix which led to a secondary carbides free aus-tenitic matrix after heat treatment

The amount of precipitated particles showed an increasewith time for the samples treated at 900 and 1000 degCreaching a maximum and then decreased due to dissolutionof some carbides and coarsening of some others

Matrix hardness showed a maximum when plotted againstsoaking time due to these precipitation processes and tothe transformation of austenite to martensite Maximumhardness was obtained for the samples treated for 1-2 h at900 degC where the structure showed a high volume fractionof martensite and a high amount of fine well-dispersedsecondary carbide particles

Acknowledgments

The authors are grateful to the University of Michoacan forfunding this project

References

1 CP Tabrett IR Sare and MR Gomashchi ldquoMicrostructure-Property Relationships in High-Chromium White Iron Alloysrdquo IntMater Rev 1996 41(2) pp 59-82

2 ON Dogan JA Hawk and G Laird II ldquoSolidification Structure andAbrasion Resistance of High Chromium White Ironsrdquo Metall MaterTrans A 1997 28(6) pp 1315-27

3 G Laird II and GLF Powell ldquoSolidification and Solid-State Trans-formation Mechanisms in Si Alloyed High-Chromium White CastIronsrdquo Metall Trans A 1993 24(2) pp 981-88

4 WW Cias ldquoAustenite Transformation Kinetics and Hardenability ofHeat Treated 175Cr White Cast Ironsrdquo Trans AFS 1974 82 pp317-28

5 GLF Powell and G Laird II ldquoStructure Nucleation Growth andMorphology of Secondary Carbides in High Chromium and Cr-NiWhite Cast Ironsrdquo J Mater Sci 1992 27 pp 29-35

6 F Maratray and A Poulalion ldquoAustenite Retention in High-Chromium White Ironsrdquo Trans AFS 1982 90 pp 795-804

7 JTH Pearce ldquoStructure and Wear Performance of Abrasion ResistantChromium White Cast Ironsrdquo Trans AFS 1984 92 pp 599-621

8 M Kuwano K Ogi A Sawamoto and K Matsuda ldquoStudies onPrecipitation Process of Secondary Carbides in High-Chromium CastIronrdquo Trans AFS 1990 98 pp 725-34

9 GLF Powell and JV Bee ldquoSecondary Carbide Precipitation in an 18wtCr-1 wtMo White Ironrdquo J Mater Sci 1996 31 pp 707-11

10 C Kim ldquoX-Ray Method of Measuring Retained Austenite in HeatTreated White Cast Ironsrdquo J Heat Treating 1979 1(2) pp 43-51

11 A Bedolla-Jacuinde ldquoMicrostructure of Vanadium- Niobium- andTitanium-Alloyed High-Chromium White Cast Ironsrdquo Int J CastMetals Res 2001 13(6) pp 343-61

Fig 16 Retained austenite content in iron matrix as a function of thesoaking time during the destabilization heat treatment

Fig 17 Matrix microhardness as a function of the soaking timeduring the destabilization heat treatment

Fig 18 Bulk hardness of the alloy as a function of the soaking timeduring the destabilization heat treatment

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash381

lar to that for matrix microhardness Figure 18 shows the re-sults for hardness where maximum values can be observed forthe different temperatures For the iron treated at 900 degC maxi-mum hardness was 68 HRC and was obtained for1-2 h soakingtime For the iron treated at 1000 degC maximum hardness was60 HRC obtained for 30 min Finally for the iron treated at1150 degC hardness remains constant (45 HRC) during the first20 min of treatment and decreases afterwards The strengthen-ing of the matrix is the main factor responsible for the modi-fications in the ironrsquos hardness Microhardness of matrix andthe eutectic carbide volume fraction are the factors that con-

tribute to strengthening of the alloys For this case the carbidevolume fraction was constant (24) therefore the strengthen-ing of matrix caused the changes in bulk hardness

These strengthening phenomena are similar to those foundin aging processes of some alloys and are explained in terms ofthe diffusion of elements that participate in the precipitatingphase Such diffusion in turn is a function of treatment tem-perature and soaking time Kuwano et al[ 8 ] Powell andLaird[ 5 ] and Powell and Bee[ 9 ] have found similar behaviorsfor high- and low-chromium irons in studies of precipitationand hardness

Fig 13 Bright field TEM micrographs showing some secondary carbides (SC) in (a) sample treated at 900 degC for 1 h and (b) sample treated at1000 degC for 1 h The selected area diffraction pattern (SADP) on the particle arrowed in (b) indicates the crystallographic nature of these carbidesto be hexagonal M7C3

Fig 14 Diameter of the precipitated particles as a function of thesoaking time during destabilization heat treatment

Fig 15 Number of precipitated particles in a square micron (mm2) asa function of the soaking time during destabilization heat treatment

380mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance

4 Conclusions

The structure of the as-cast alloy was composed of 24eutectic carbides about 5 martensite and 70 austenitesmall amounts of the molybdenum-rich carbide Mo2Cwere also present

Vanadium carbide was not detected in the alloy nonethe-less the iron contained 2 vanadium a carbide-promotingelement Vanadium was observed to partition to the car-bide and in minor amounts to the matrix

According to the studied temperatures the secondary car-bide precipitation level decreased with higher tempera-tures being the highest for the samples treated at 900 degC(27 of matrix)

High temperatures like 1150 degC over-stabilized the aus-tenitic matrix which led to a secondary carbides free aus-tenitic matrix after heat treatment

The amount of precipitated particles showed an increasewith time for the samples treated at 900 and 1000 degCreaching a maximum and then decreased due to dissolutionof some carbides and coarsening of some others

Matrix hardness showed a maximum when plotted againstsoaking time due to these precipitation processes and tothe transformation of austenite to martensite Maximumhardness was obtained for the samples treated for 1-2 h at900 degC where the structure showed a high volume fractionof martensite and a high amount of fine well-dispersedsecondary carbide particles

Acknowledgments

The authors are grateful to the University of Michoacan forfunding this project

References

1 CP Tabrett IR Sare and MR Gomashchi ldquoMicrostructure-Property Relationships in High-Chromium White Iron Alloysrdquo IntMater Rev 1996 41(2) pp 59-82

2 ON Dogan JA Hawk and G Laird II ldquoSolidification Structure andAbrasion Resistance of High Chromium White Ironsrdquo Metall MaterTrans A 1997 28(6) pp 1315-27

3 G Laird II and GLF Powell ldquoSolidification and Solid-State Trans-formation Mechanisms in Si Alloyed High-Chromium White CastIronsrdquo Metall Trans A 1993 24(2) pp 981-88

4 WW Cias ldquoAustenite Transformation Kinetics and Hardenability ofHeat Treated 175Cr White Cast Ironsrdquo Trans AFS 1974 82 pp317-28

5 GLF Powell and G Laird II ldquoStructure Nucleation Growth andMorphology of Secondary Carbides in High Chromium and Cr-NiWhite Cast Ironsrdquo J Mater Sci 1992 27 pp 29-35

6 F Maratray and A Poulalion ldquoAustenite Retention in High-Chromium White Ironsrdquo Trans AFS 1982 90 pp 795-804

7 JTH Pearce ldquoStructure and Wear Performance of Abrasion ResistantChromium White Cast Ironsrdquo Trans AFS 1984 92 pp 599-621

8 M Kuwano K Ogi A Sawamoto and K Matsuda ldquoStudies onPrecipitation Process of Secondary Carbides in High-Chromium CastIronrdquo Trans AFS 1990 98 pp 725-34

9 GLF Powell and JV Bee ldquoSecondary Carbide Precipitation in an 18wtCr-1 wtMo White Ironrdquo J Mater Sci 1996 31 pp 707-11

10 C Kim ldquoX-Ray Method of Measuring Retained Austenite in HeatTreated White Cast Ironsrdquo J Heat Treating 1979 1(2) pp 43-51

11 A Bedolla-Jacuinde ldquoMicrostructure of Vanadium- Niobium- andTitanium-Alloyed High-Chromium White Cast Ironsrdquo Int J CastMetals Res 2001 13(6) pp 343-61

Fig 16 Retained austenite content in iron matrix as a function of thesoaking time during the destabilization heat treatment

Fig 17 Matrix microhardness as a function of the soaking timeduring the destabilization heat treatment

Fig 18 Bulk hardness of the alloy as a function of the soaking timeduring the destabilization heat treatment

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash381

4 Conclusions

The structure of the as-cast alloy was composed of 24eutectic carbides about 5 martensite and 70 austenitesmall amounts of the molybdenum-rich carbide Mo2Cwere also present

Vanadium carbide was not detected in the alloy nonethe-less the iron contained 2 vanadium a carbide-promotingelement Vanadium was observed to partition to the car-bide and in minor amounts to the matrix

According to the studied temperatures the secondary car-bide precipitation level decreased with higher tempera-tures being the highest for the samples treated at 900 degC(27 of matrix)

High temperatures like 1150 degC over-stabilized the aus-tenitic matrix which led to a secondary carbides free aus-tenitic matrix after heat treatment

The amount of precipitated particles showed an increasewith time for the samples treated at 900 and 1000 degCreaching a maximum and then decreased due to dissolutionof some carbides and coarsening of some others

Matrix hardness showed a maximum when plotted againstsoaking time due to these precipitation processes and tothe transformation of austenite to martensite Maximumhardness was obtained for the samples treated for 1-2 h at900 degC where the structure showed a high volume fractionof martensite and a high amount of fine well-dispersedsecondary carbide particles

Acknowledgments

The authors are grateful to the University of Michoacan forfunding this project

References

1 CP Tabrett IR Sare and MR Gomashchi ldquoMicrostructure-Property Relationships in High-Chromium White Iron Alloysrdquo IntMater Rev 1996 41(2) pp 59-82

2 ON Dogan JA Hawk and G Laird II ldquoSolidification Structure andAbrasion Resistance of High Chromium White Ironsrdquo Metall MaterTrans A 1997 28(6) pp 1315-27

3 G Laird II and GLF Powell ldquoSolidification and Solid-State Trans-formation Mechanisms in Si Alloyed High-Chromium White CastIronsrdquo Metall Trans A 1993 24(2) pp 981-88

4 WW Cias ldquoAustenite Transformation Kinetics and Hardenability ofHeat Treated 175Cr White Cast Ironsrdquo Trans AFS 1974 82 pp317-28

5 GLF Powell and G Laird II ldquoStructure Nucleation Growth andMorphology of Secondary Carbides in High Chromium and Cr-NiWhite Cast Ironsrdquo J Mater Sci 1992 27 pp 29-35

6 F Maratray and A Poulalion ldquoAustenite Retention in High-Chromium White Ironsrdquo Trans AFS 1982 90 pp 795-804

7 JTH Pearce ldquoStructure and Wear Performance of Abrasion ResistantChromium White Cast Ironsrdquo Trans AFS 1984 92 pp 599-621

8 M Kuwano K Ogi A Sawamoto and K Matsuda ldquoStudies onPrecipitation Process of Secondary Carbides in High-Chromium CastIronrdquo Trans AFS 1990 98 pp 725-34

9 GLF Powell and JV Bee ldquoSecondary Carbide Precipitation in an 18wtCr-1 wtMo White Ironrdquo J Mater Sci 1996 31 pp 707-11

10 C Kim ldquoX-Ray Method of Measuring Retained Austenite in HeatTreated White Cast Ironsrdquo J Heat Treating 1979 1(2) pp 43-51

11 A Bedolla-Jacuinde ldquoMicrostructure of Vanadium- Niobium- andTitanium-Alloyed High-Chromium White Cast Ironsrdquo Int J CastMetals Res 2001 13(6) pp 343-61

Fig 16 Retained austenite content in iron matrix as a function of thesoaking time during the destabilization heat treatment

Fig 17 Matrix microhardness as a function of the soaking timeduring the destabilization heat treatment

Fig 18 Bulk hardness of the alloy as a function of the soaking timeduring the destabilization heat treatment

Journal of Materials Engineering and Performance Volume 12(4) August 2003mdash381

12 G Laird II ldquoMicrostructures of Ni-Hard I Ni-Hard IV and High-CrWhite Cast Ironsrdquo Trans AFS 1991 99 pp 339-57

13 P Duppin J Saverna and JM Schissler ldquoA Structural Study ofChromium White Cast Ironrdquo Trans AFS 1982 90 pp 711-18

14 F Maratray ldquoChoice of Appropriate Compositions for Chromium-Molybdenum White Ironsrdquo Trans AFS 1971 79 pp 121-24

15 JDB DeMello M Duran-Charre and S Hamar-Thibualt ldquoSolidifi-cation and Solid State Transformations During Cooling of Chromium-Molybdenum White Cast Ironsrdquo Metall Trans A 1983 14(9) pp1793-801

16 JW Choi and SK Chang ldquoEffects of Molybdenum and CopperAdditions on Microstructure of High Chromium Cast Iron Rollsrdquo ISIJInt 1992 32(11) pp 1170-76

17 CR Loper Jr and HK Baik ldquoInfluence of Molybdenum and Tita-nium on the Microstructures of Fe-C-Cr-Nb White Cast Ironsrdquo TransAFS 1989 97 pp 1001-08

18 CP Tong T Suzuki and T Umeda ldquoEutectic Solidification of HighChromium Cast Ironsrdquo in Proc IV Int Symposium on the Physical

Metallurgy Ed Materials Research Society Tokyo Japan 1989pp 403-10

19 M Ikeda T Umeda CP Tong T Suzuki N Niwa and O KatoldquoEffect of Molybdenum Addition on Solidification Structure Me-chanical Properties and Wear Resistivity of High Chromium CastIronrdquo ISIJ Int 1992 32(11) pp 1157-62

20 M Radulovic M Fiset K Peev and M Tomovic ldquoThe Influence ofVanadium on Fracture and Abrasion Resistance in High ChromiumWhite Cast Ironsrdquo J Mater Sci 1994 29 pp 5085-94

21 JV Bee GLF Powell and B Bednarz ldquoA Substructure Within theAustenitic Matrix of High-Chromium White Cast Ironsrdquo Scripta Met-all Mater 1994 31 pp 1735-36

22 ON Dogan and JA Hawk ldquoEffect of Retained Austenite on Abra-sion Resistance of High-Cr White Cast Ironsrdquo Trans AFS 1997 105pp 167-74

23 DK Subramanyam ldquoRetained Austenite Measurements in HighChromium White Ironsrdquo Trans AFS 1985 93 pp 763-68

382mdashVolume 12(4) August 2003 Journal of Materials Engineering and Performance