229-sWELDING JOURNAL

WE

LD

ING

RE

SE

AR

CH

Introduction

The Fe-Cr-C alloy, well known for itshigh hardness and excellent wear and cor-rosion resistance, has been widely appliedin harsh working conditions. Many impor-tant workpieces, such as hammers in min-ing and mineral processing, squeezingrolls in cement production, and abrasion-resistant plates in the manufacturing andmetallurgy industries, are manufacturedfrom Fe-Cr-C alloy (Ref. 1). The excellentabrasive wear resistance results primarilyfrom the type, morphology, amount, di-

mension, and distribution of the carbides,while the toughness of the matrix also con-tributes to the wear resistance (Ref. 2). Fe-Cr-C alloys have been classified into hy-poeutectic, eutectic, and hypereutecticstructures (Refs. 3–5). Compared with thehypoeutectic one, the hypereutectic Fe-Cr-C alloy is regarded as having betterwear resistance, because its microstruc-ture consists of primary M7C3 carbide andeutectic (γ+M7C3) (Ref. 6). While theprimary carbides in the hypereutectic mi-crostructure maintain their forms as

coarser and larger blocks, this in turn de-creases the cast ability (Ref. 7). In general,Fe-Cr-C alloys with hypoeutectic mi-crostructures are applied in engineeringby the casting method.

Workpieces made of Fe-Cr-C alloy failthrough excessive wear over a period oftime. Failed workpieces can be remanu-factured using a hardfacing method. Nor-mally, the hardfacing layers are expectedto be hypereutectic microstructures forobtaining higher hardness and better wearresistance (Ref. 8). Much attention hasbeen focused on improving the wear re-sistance of hypereutectic Fe-Cr-C alloys(Refs. 8–12).

Tungsten carbide (WC), acting as anadvanced ceramic material with wear re-sistance and good thermal shock resist-ance, has been widely used for wear-resistance applications. Kambakas tried touse a double casting technique to producea WC-particle-reinforced high-Cr whitecast iron, and informed that the wear re-sistance of the high-Cr white cast iron withWC particle reinforcement was signifi-cantly better than that without thestrengthening phase. For hardfacing con-sumables, WC particles are not suitablefor reinforcing Fe-Cr-C alloy due to thehigh temperature of the weld pool (Ref.9).

The applications of rare earth (RE) ele-ments have been of much concern recentlybecause of their excellent properties. Byadding RE elements to steel, the crystalgrain can be refined. Hao explored the ef-fect of RE oxides on the morphology of car-bides in hardfacing metal of high-chromiumcast iron. In his studies, the volume fractionand roundness of the carbides were gradu-ally increased, while their area and perime-ter were gradually reduced. The carbideswere refined and spheroidized, with the REoxide additions increasing. Nevertheless,the relationship between the volume frac-tion of carbides and the wear resistance ofthe hardfacing metal was not established in

Effect of Titanium Content on Microstructureand Wear Resistance of Fe-Cr-C

Hardfacing Layers

By adding different amounts of ferrotitanium into flux cored wire, a hardfacinglayer with good performance was obtained, and the M7C3 carbide refinement

mechanism is discussed

BY Y. F. ZHOU, Y. L. YANG, D. LI, J. YANG, Y. W. JIANG, X. J. REN, AND Q. X. YANG

ABSTRACT

Layers of Fe-Cr-C hardfacing material containing various amounts of titaniumwere deposited on ASTM 1045 steel base metal. Optical microscope (OM), field emis-sion scanning electron microscope (FESEM) with energy-dispersive spectrometer(EDS), and X-ray diffraction (XRD) were used to investigate the effect of titaniumcontent on the microstructural characteristics of Fe-Cr-C hardfacing layers. The so-lidification sequence calculation and lattice misfit theory were employed to discuss theM7C3 carbide refinement mechanism. The experimental results show the microstruc-tures of Fe-Cr-C hardfacing layers consist of primary (Cr, Fe)7C3 carbides and the eu-tectic phases (γ-Fe+(Cr, Fe)7C3). In the solidification process, the formation andgrowth of the primary (Cr, Fe)7C3 carbides occur along their long axis, which parallelsthe direction of heat flow. With the increase of titanium content, the primary (Cr,Fe)7C3 carbides are refined. However, it is not proper to increase titanium contentwithout limits. When titanium content reaches 1.17 wt-%, its microstructure changesfrom a hypereutectic form to a hypoeutectic one. The thermodynamic calculationshows MC carbide precipitates prior to M7C3 carbide from Fe-C-Cr-Ti alloy. More-over, the lattice misfit between (110)TiC and (010)Cr7C3 is 9.257%, which indicates thatTiC acting as heterogeneous nuclei of the Cr7C3 is medium effective. Therefore, M7C3carbide can be refined significantly.

KEYWORDS

CarbidesHardfacingMicrostructureNucleationTitanium

Y. F. ZHOU (yef.zhou@gmail.com), Y. L.YANG, D. LI, J. YANG, Y. W. JIANG, and Q. X.YANG (qxyang@ysu.edu.cn) are with State KeyLaboratory of Metastable Materials Science &Technology, Yanshan University, Qinhuangdao,China. LI is also with School of Material Scienceand Engineering, Southwest Jiaotong University,Chengdu, China. X. J. REN is with School of En-gineering, Liverpool John Moores University, Liv-erpool, UK.

Zhou Supplement August 2012_Layout 1 7/11/12 3:04 PM Page 229

his work (Ref. 10).Vanadium (V), niobium (Nb), and tita-

nium (Ti) are strong carbide-forming ele-ments, and are of benefit for refining themicrostructure and improving the wear re-sistance of the Fe-Cr-C alloy. Qi investi-gated the effects of vanadium additive onstructural properties and tribological per-formance of high-chromium cast ironhardfacing metal. In his study, V was abenefit element of the Fe-Cr-C alloy. Withthe addition of V, vanadium carbide wasformed as a secondary carbide of Fe-Cr-C-V alloy. The microstructure of the alloywas obviously refined with the increase ofV additive, and the amount of bulk pri-mary carbide was reduced with an increasein refined eutectic carbide (Ref. 8). Whilethe carbides in the Fe-Cr-C alloy appar-ently could also be refined with the Nb ad-dition, the shape of the primary M7C3 car-bides became isotropic. Via the XRD andEDS analyses, NbC carbide was identifiedwhen the Nb element was added into theFe-Cr-C alloy (Ref. 11).

The contribution of Ti to the Fe-Cr-Calloy also can be found in the literature, butthe views contained therein have not yetfully become the consensus. Chung foundthe added titanium in the Fe–25wt-%Cr–4wt-%C alloy did not act as an inoculant torefine primary M7C3 carbides (Ref. 12). In-stead, it was just the reverse, as Zhi ex-plained that the heterogeneous nuclei role

of TiC in the Fe-Cr-Calloy (Refs. 13, 14).Moreover, the wear re-sistance of Fe-Cr-Calloy related to themass fraction of thecarbide with added Tiwas not quantified as

described in previous literature.Based on the above study, the effect of

titanium on hardfacing metal of Fe-Cr-Calloy is reinvestigated in this work. Thevariation of microstructure, phase trans-formations, and wear resistance are ob-served, the carbide refinement under dif-ferent Ti content is describedquantitatively, and the carbide refinementis discussed.

Experimental Procedures

The base metals (100 × 80 ×10 mm)for hardfacing were prepared from ASTM1045 steel plates. Before welding, the basemetals were ground and cleaned with ace-tone. Flux cored wire, which consisted ofan outer steel strip and wrapped powder,was prepared. H08A was selected as thesteel strip due to its good toughness. In ad-dition, the composition of the wrappedpowder was adjusted by adding differentraw materials. The graphite (2 wt-%), fer-rochrome (25 wt-%), ferrosilicon (3 wt-%), ferromanganese (3 wt-%), and fer-rotitanium were uniformly mixed andprepared. Moreover, to investigate the ef-fect of titanium on the microstructures ofthe hardfacing layers, 0, 1, 2, and 4 wt-%ferrotitanium, respectively, were alsoadded to the powder. After the powderwas prepared, the forming roller was usedto roll the steel strip into a U-groove, and

then, before the steel strip was rolled intoa tubular shape, the well-mixed powderwas filled into the U-groove. Furthermore,the required dimension of the flux coredwire was achieved by rolling or wire draw-ing methods. The diagram of flux coredwire fabrication is shown in Fig. 1A, B.

The bead-on-plate technique with fluxcored arc welding (FCAW) was used to de-posit the layers via an automated system inwhich the welding torch was moved backand forth above the base metal at a constantspeed in a multitrack overlapping process.The length of the single track was 50 mm,and the overlap width was 4 mm. To reducethe effect of base metal on the microstruc-ture and property of the hardfacing metal,the hardfacing claddings were welded inthree layers. Table 1 presents the range ofwelding conditions, and the hardfacingequipment and process used in this researchare shown in Fig. 1C, D.

The center of the hardfacing layers wasselected as the analytical region. Speci-mens were machined into cuboids (10 ×10 × 18 mm) by a wire cutting machine foranalysis. The chemical composition of thelayers was determined by a SPECTRO-MAXx optical emission spectrum (OES),and the data are listed in Table 2. Both thehorizontal and vertical faces of the speci-mens were treated with rubdown and pol-ishing processes, and then etched with 4%nitric acid. The microstructures of speci-mens were observed through an Axiovert200 MAT optical microscope and a Hi-tachi S4800 field emission scanning elec-tron microscope (FESEM). The morphol-ogy, size, and grade of the primarycarbides, and transfer of matrix structureswere measured with Image-Pro Plus Ver-sion 6.0 software. In addition, 10 OM im-ages were selected randomly from eachlayer in the horizontal direction at 200×magnification to describe the statisticalnature of the maximum diameter and areaof each M7C3 carbide. The inclusion com-positions were analyzed by an EMAX en-ergy-dispersive spectrometer (EDS).D/max-2500/PC X-ray diffraction (XRD)with Cu Kα radiation was used to analyzethe constituent phases of the top surface

230-s

WE

LD

ING

RE

SE

AR

CH

AUGUST 2012, VOL. 91

Fig. 1 — A, B — Diagram of flux cored wire fabrication; C — hardfacingequipment; D — hardfacing process.

Fig. 2 — Schematic diagram of analysis layer.

Table 1 — FCAW Condition

Parameter Wire Voltage Current Travel Speed Welding LayerDiameter Layers Thickness

Value 3.2 mm 22∼24 V 240∼260 A 300 mm⋅min–1 3 8 mm

A B

C D

Zhou Supplement August 2012_Layout 1 7/11/12 3:04 PM Page 230

of the hardfacing layers. A wear resistancetest was conducted on an abrasive belt-type wear testing machine in dry friction.SiC of 80 mesh was selected as the abra-sive material, and the wear velocity of theabrasive belt was 1.8 × 104 mm/min. Elec-tronic balance was used in the wear test toweigh the specimens’ loss of mass perhour. To decrease experimental error andmaintain accuracy, six samples of eachchemical composition were prepared forthe wear resistance test.

The NETZSCH STA 449 C differential

scanning calorimeter(DSC) was used tostudy the phase trans-formation of the hard-facing metal. Theheating and cooling

rates were 40° and 10°C/min, respectively.The Thermo-Calc software was used forphase mass fraction calculation of Fe-Cr-C-Ti alloy with temperature.

Experimental Results andDiscussions

Microstructure and Phase Characteristicsof the Fe-Cr-C Hardfacing Layers

Figure 2 illustrates the three-dimen-sional microstructure schematic of the Fe-

Cr-C hardfacing layer. The XRD results ofsurface layers with and without titaniumare shown in Fig. 3. From the microscopeimage of the vertical face, it can be con-cluded that the specimen can be dividedinto the hardfacing zone, dilution zone,heat-affected zone (HAZ), and substratein sequence. With the aid of X-ray diffrac-tion, the hardfacing microstructure withfree Ti addition is found to consist of twophases: the primary (Cr, Fe)7C3 and theeutectic (γ-Fe+(Cr, Fe)7C3). Besides, afterTi was added into the hardfacing layer, TiCcarbide can also be detected in the hard-facing microstructures. Moreover, the mi-crostructure of the dilution zone can be anadmixture of γ-Fe, (Cr, Fe)7C3 and ferrite.The microstructure of the HAZ consists ofthe coarse grain caused by the heat input

231-sWELDING JOURNAL

WE

LD

ING

RE

SE

AR

CH

Fig. 4 — OM photographs of Fe-Cr-C-Ti layers in horizontal direction withdifferent titanium contents: A — 0 wt-%; B — 0.28 wt-%; C — 0.63 wt-%;D — 1.17 wt-%.

Fig. 5 — Microstructure of Fe-Cr-C-Ti layers with different titanium additions.

Fig. 3 — XRD of hardfacing layer with and without titanium contents: A — 0 wt-% Ti content; B — 0.28 wt-% Ti content.

Table 2 — Chemical Compositions of the Hardfacing Layers and Base Metal

Composition (wt-%)

Layer C Cr Mn Si Ti Fe

Base metal (1045) 0.43 0.23 0.65 0.21 balSpecimen a 3.82 16.35 2.24 2.09 0Specimen b 3.79 15.97 2.27 2.11 0.28Specimen c 3.85 16.14 2.18 2.14 0.63 balSpecimen d 3.77 16.27 2.22 2.03 1.17

A B

A B

C D

Zhou Supplement August 2012_Layout 1 7/11/12 3:04 PM Page 231

energy, and the substrate contains ferrite.The solidification morphology and thegrowth pattern of the layer are controlledby the thermal conditions in the weld pool.The formation and growth of the primary(Cr, Fe)7C3 carbides occur along their longaxis, which parallels the direction of theheat flow. The primary (Cr, Fe)7C3 is thehexagonal-columniation structures. Dueto different angles, it is an acicular or

blade-like morphology on the verticalfaces, and a hexagonal-shaped morphol-ogy on the horizontal.

Effect of Titanium on Microstructure ofthe Fe-Cr-C Hardfacing Layers

For the hardfacing application, thewear resistance is mainly determined bythe morphology and distribution of pri-

mary (Cr, Fe)7C3 carbides in the horizon-tal direction. The OM photographs of thelayers in the horizontal direction areshown in Fig. 4.

As shown, the microstructure of thehardfacing layers consisted of primary(Cr, Fe)7C3 carbides and eutectic (γ-Fe+(Cr, Fe)7C3), while the carbides and,in particular, the primary carbides, are re-fined gradually with the increase of Ti con-tent. However, in Fig. 4D, γ-Fe dendritecan be observed. Too much carbon is con-sumed with the formation of TiC domainswhen titanium content reaches 1.17 wt-%,resulting in a change in microstructure ofthe alloy from the hypereutectic form to ahypoeutectic one. Therefore, it is notproper to increase the titanium contentwithout limits.

The schematic diagram of the mi-crostructural changes is illustrated in Fig.5. As shown, the morphology of primary(Cr, Fe)7C3 carbides changes from a bulkform to a refined one and the size of pri-mary (Cr, Fe)7C3 carbides becomes muchsmaller. Besides, matrix microstructurestransform from eutectic (γ-Fe + (Cr,

232-s

WE

LD

ING

RE

SE

AR

CH

AUGUST 2012, VOL. 91

Fig. 6 — Quantitative analysis of M7C3 carbide.

Fig. 8 — Differential scanning calorimeter curves of Fe-Cr-C alloy with 1.17Ti content.

Fig. 7 — Mass loss and hardness of Fe-Cr-C layer with different titanium con-tents.

Fig. 9 — Phase mass fraction calculation of Fe-Cr-C-Ti alloy with temperature.

Table 3 — Planar Lattice Misfit between Orthorhombic Cr7C3 and TiC

Matching Interface (110)TiC // (010)Cr C7 3

[uvw]TiC [001] [110] [111]

[uvw]Cr C [001] [100] [101]7 3

Θ 0 0 12.1

dTiC(nm) 0.432 0.610 0.747

dCrC (nm) 0.453 0.701 0.8347 3

δ,% 9.257

- -

Zhou Supplement August 2012_Layout 1 7/11/12 3:04 PM Page 232

Fe)7C3) to dendrite γ-Fe +eutectic (γ-Fe+ (Cr, Fe)7C3).

Because the M7C3 carbide is the majorreinforcing phase in Fe-Cr-C alloy to re-duce friction and wear, the refinement ofthe carbide was analyzed quantitatively,which is shown in Fig. 6. As shown, thearea of M7C3 carbide is gradually reducedwith the increasing Ti content. When theFe-Cr-C alloy is Ti-free, the area of M7C3carbide is nonuniform from 200 to 900μm2. With higher Ti content, the unifor-mity of M7C3 carbide is enhanced. With0.28 wt-% Ti content, a majority of M7C3carbide area are distributed between 200and 600 μm2 and in the case of 0.63 wt-%Ti content, the M7C3 carbide area fell tothe range of 100~300 μm2. As the Ti con-tent increased to 1.17 wt-%, the area ofM7C3 carbide reduced further, which re-mained constant at less than 200 μm2. Be-sides, the maximum diameter of the M7C3carbide declined markedly as the Ti con-tent increased.

Effect of Titanium on Wear Resistanceand Hardness of the Fe-Cr-C HardfacingLayers

Figure 7 presents the wear resistanceand hardness of the hardfacing layers. Asillustrated, the wear resistance of the Fe-Cr-C alloy increases and then declineswith respect to the amount of added Ti.Meanwhile, the changes in hardness areassociated with the wear resistance be-havior. When the titanium content is 0.63wt-%, the Fe-Cr-C hardfacing layer pres-ents the best wear resistance and thehighest hardness. There are two factorsthat lead to the variation in wear-resis-tance property. TiC carbide, which has ahigh micro-hardness of 3200–3800 HV(Ref. 15), is present in the Fe-Cr-C alloyas added Ti. However, the high micro-hardness of TiC carbide may not work tocontribute to the improvement in wearresistance. When the Ti content is 1.17wt-%, the wear resistance behavior is the

worst and thehardness has re-duced dramati-cally. Therefore,the TiC carbide it-self may not havethe main role inthe variation ofwear-resistancebehavior.

Nevertheless,the morphologyand distribution ofM7C3 carbide ischanged when Ti isadded to the Fe-Cr-C alloy. Partic-ularly when com-pared with thelarge block car-bide, it has beenconfirmed that therefined carbide,which has muchmore contact areawith the matrix,causes a better antistripping ability.Therefore, the wear-resistance is betterwhen the carbide is refined and well dis-tributed. However, the M7C3 carbide isnot only refined but changes its phasestructure when too much Ti is added.When the Fe-Cr-C alloy contains 1.17 wt-% Ti, more carbon is consumed to formTiC carbide. The loss of carbon reducesthe formation of chromium carbide andcauses the hardfacing layer to change fromthe hypereutectic microstructure to a hy-poeutectic one. It is well known that thewear resistance of hypoeutectic Fe-Cr-Calloy without coarse primary carbide is notas good as the hypereutectic alloy, whichwould lead to a decline in hardness and re-duced wear resistance.

Furthermore, the hardness of the hard-facing metal increases from 58 to 61 HRC

with the increasing titanium content from0 to 0.63 wt-%, while the hardness de-creases to 55 HRC when the titanium con-tent reaches 1.17 wt-%. The variation inthe hardness is consistent with that of thewear resistance results.

Effect of Titanium on Phase Transforma-tion of the Fe-Cr-C-Ti Alloy

The differential scanning calorimeter(DSC) results for the Fe-Cr-C hardfacingalloy with 1.17 wt-% Ti content are shownin Fig. 8. There are, respectively, two en-dothermic peaks in the heat process andtwo exothermic peaks in the coolingprocess shown in the curves.

The first peak in cooling curve at1284.6°C is due to the formation of M7C3carbide and γ-phase. And, then, the liquiddisappears. At lower temperature, the sec-ond exothermic peak can be seen at

233-sWELDING JOURNAL

WE

LD

ING

RE

SE

AR

CH

Fig. 10 — Effect of Ti on phase mass fraction calculation of Fe-Cr-C-Tialloy.

Fig. 11 — FESEM morphology and line scanning results of hardfacing layer.

Fig. 12 — Correspondence condition of (110)TiC and (010)Cr7C3.

Zhou Supplement August 2012_Layout 1 7/11/12 3:04 PM Page 233

764.6°C, corresponding to the transforma-tion from γ-phase to α-phase.

Nevertheless, the MC carbide, whosemelting point is higher, cannot be observedin the DSC result. Therefore, the computa-tional thermodynamics method is used toanalyze the phase transformation of the Fe-Cr-C-Ti alloy during the equilibrium state.The calculation result is shown in Fig. 9.

The primary phase precipitates fromliquid at 1570°C is MC carbide. With thetemperature dropping, the MC content isalmost invariable and remains at about0.015 wt-%. At 1280°C, a small quantity ofM7C3 carbide precipitates from the liq-uid, and then austenite is expected to formbelow 1270°C together with some addi-tional precipitation of M7C3. At 1260°C,the liquid disappears and the mass frac-tion of M7C3 still increases with a little re-duction of austenite. Thereafter, M7C3starts to be transformed from the austen-ite. When the temperature falls to 780°C,the transformation of austenite to marten-site occurs and leads to exotherm in thecooling process. The temperature ofexothermic peaks in DSC result (Fig. 8) isclose to the calculation, which verifies theexactness of the calculation model.

The influence of titanium on phasemass fraction is shown in Fig. 10. The massfraction variation of each phase is not ob-vious during the changes in Ti content.With Ti content increasing, the mass frac-tion of MC rises slightly. Meanwhile, themass fraction of M7C3 declines becausesome carbon has been consumed by the Ticontent to form the initial MC carbide.Martensite in Fe-Cr-C-Ti alloy ascendsslightly due to the rising of austenite at ahigh temperature.

M7C3 Carbide Refinement Mechanism

The results described previously sug-gest that the improvement in wear resist-ance is mainly dependent on the refine-ment of the carbide in Fe-Cr-C-Ti alloy byadded Ti content. As an effective element,Ti is widely used in metallurgy of iron andsteel for partition to the matrix as well asmodification of the carbides (Ref. 14).How it works for refining the Cr7C3 car-bide is discussed in this section.

Figure 11 shows the field emissionscanning electron morphology of thehardfacing layers. From Fig. 11A, it can beseen that square particles are surroundedby (Cr, Fe)7C3 carbides. According to theEDS analysis and line scanning results ofthe square particle shown in Fig. 11B andD, the main compositions are titaniumand carbon, which indicates that thesquare particle is TiC carbide. Besides, themore clear morphology of the TiC carbidecan be seen in Fig. 11C.

Therefore, it can be said that the re-fined primary (Cr,Fe)7C3 carbides are re-

lated to TiC carbides. During the hard-facing solidification process, the fastercooling rate results in smaller dimensionsand greater number of nuclei. The resist-ance of heterogeneous nucleation mainlydepends on the interfacial energy be-tween nucleation basement and crys-talline phase. And the interfacial energyis constituted by its chemistry item andstructural one. The chemistry item in-cludes bond strength, bond energy, andbond types between atoms, and the struc-tural item is mainly decided by lattice dis-tortion energy, which is caused by theatomic misfit. Misfit is the major factor ofthe interfacial energy in higher lattice dis-tortion energy.

The value of the two-dimensional lat-tice misfit is used to estimate whethersome inclusions can act as the heteroge-neous nuclei. A mathematical model ofthe two-dimensional lattice misfit is as fol-lows (Ref. 16):

Where(hkl)s is a low-index plane of the matrix;[uvw]s is a low-index direction in (hkl)s;(hkl)n is a low-index plane in the nucleatedsolid;[uvw]n is a low-index direction in (hkl)n;d[uvw]n is the interatomic spacing along[uvw]n;d[uvw]s is the interatomic spacing along[uvw]s;θ is the angle between the [uvw]s and[uvw]n (θ≤90 deg).

Bramfitt (Ref. 16) proposed a theoryregarding the heterogeneous nucleationprocess. The nuclei with δ<6% is themost effective, and that with δ between 6and 12% is medium effective, while thatwith δ>12% is ineffective.

The crystal lattice of TiC is face-cen-tered cubic, and its lattice parameter is a= 0.432 (nm). Orthorhombic Cr7C3 isone mode of (Cr, Fe)7C3, and its latticeparameters are a = 0.701 (nm), b = 1.214(nm), and c = 0.453 (nm) (Refs. 17, 18).The atom correspondence condition ofthose two planes is shown in Fig. 12. Table3 lists the calculated result of the latticemisfit δ between (110)TiC and (001)Cr7C3.It can be seen that the lattice misfit be-tween (110)TiC and (001)Cr7C3 is 9.257%.According to Bramfitt’s two-dimensionallattice misfit theory, TiC acting as het-erogeneous nuclei of the Cr7C3 is middleeffective and the primary Cr7C3 carbidesare refined. These results are appropriatesupplements that some works (Ref. 13)also point out that titanium and/or nio-bium can refine the microstructure of theFe-Cr-C alloy.

Conclusions

A series of Fe-Cr-C hardfacing layerswith varying amounts of titanium was de-posited by the FCAW process. The mi-crostructure and wear resistance of the Fe-Cr-C hardfacing layers were determinedand correlated to the varying titaniumcontents. Meanwhile, the carbide refine-ment mechanism and the phase precipita-tion rule were discussed. Following are themajor conclusions that can be drawn fromthis work:• Microstructures of the hardfacing layers

consisted of the primary (Cr, Fe)7C3and the eutectic (γ-Fe+(Cr, Fe)7C3).The existence of M7C3-type carbidemaintains a high hardness and goodwear resistance of the Fe-Cr-C alloy.

• Primary (Cr, Fe)7C3 carbides are refinedgradually with the increase in titaniumcontent. The morphology changes froma bulk form to a refined one. Mean-while, the increase in hardness and wearresistance improve until the titaniumcontent is increased to 0.63 wt-%. Whenthe titanium content is 1.17 wt-%, toomuch carbon is consumed by titaniumto form TiC carbide. This leads the mi-crostructure of the Fe-Cr-C alloy tochange from a hypereutectic form to ahypoeutectic one. In addition, the hard-ness decreases and wear resistance be-comes worse. Therefore, it is not properto increase the titanium content unlim-itedly, and the x%Fe-16%Cr-3.8%Calloy with 0.63 wt-% titanium content ismore appropriate.

• The M7C3 carbide refinement is relatedto the complex metallurgical reactions.According to the thermodynamic calcu-lations, the MC carbide is found to pre-cipitate prior to the M7C3 carbide. Thisprovides the MC carbide with thechance to act as the heterogeneous nu-clei of M7C3 carbide. Moreover, the lat-tice misfit between (110)TiC and(010)Cr7C3 is 9.257%, which indicatesthat TiC acting as heterogeneous nucleiof the Cr7C3 is medium effective due toBramfitt’s theory. Therefore, the M7C3carbide can be refined.

Acknowledgments

The authors would like to express theirgratitude for projects supported by Pro-gram for 100 excellent talents of HebeiProvince of China (SPRC 021) and keyproject of science and technology of HebeiProvince (09215106D).

References

1. Jacuinde, A. B., Correa, A. R., andQuezada, J. G. 2005. Effect of titanium on theas-cast microstructure of a 16% chromiumwhite iron. Materials Science and Engineering A398(1-2): 297–308.

δ θhkl

hkld ds

n

uvw si

uvw ni( )

( )= −⎡

⎣⎢⎤⎦⎥

⎡⎣⎢

⎤⎦⎥

cos // /duvw n

i

i

⎡⎣⎢

⎤⎦⎥

=

⎛

⎝⎜⎜

⎞

⎠⎟⎟

⎡

⎣⎢⎢

⎤

⎦⎥⎥×∑ 3 100 (1

1

3))

AUGUST 2012, VOL. 91234-s

WE

LD

ING

RE

SE

AR

CH

Zhou Supplement August 2012_Layout 1 7/11/12 3:04 PM Page 234

2. Llewellyn, R. J., Yick, S. K., and Dolman,K. F. 2004. Scouring erosion resistance of metal-lic materials used in slurry pump service. Wear256(6): 592–599.

3. Menon, R., and Wallin, J. 2008. Specialtycored wires for wear and corrosion applica-tions. Welding Journal 87(2): 31–36.

4. Menon, R. 2002. Recent advances incored wires for hardfacing. Welding Journal81(11): 53–58.

5. Wiengmoon, A., Chairuangsri, T., Brown,A., and Pearce, J. T. H. 2005. Microstructuraland crystallographical study of carbides in 30wt.% Cr cast irons. Acta Materialia. 53(15):4143–4154.

6. Asensio, J., Pero-Sanz, J. A., and Verdej,J. I. 2003. Microstructure selection criteria forcast irons with more than 10 wt.% chromium forwear applications. Materials Characterization49(2): 83–93.

7. Liu, H. N., Sakamoto, M., and Nomura,M. 2001. Abrasion resistance of high Cr castirons at an elevated temperature. Wear 250(1-12): 71–75.

8. Qi, X. W., Jia, Z. N.,and Yang, Q. X. 2011.Effects of vanadium additive on structure prop-erty and tribological performance of highchromium cast iron hardfacing metal. Surfaceand Coatings Technology 205(23-24):5510–5514.

9. Kambakas, K., and Tsakiropoulos, P.2005. Solidification of high-Cr white cast iron-WC particle reinforced composites. MaterialsScience and Engineering A 413-414: 538–544.

10. Hao, F. F., Li, D., and Yang, Q. X. 2011.Effect of rare earth oxides on the morphologyof carbides in hardfacing metal of highchromium cast iron. Journal of Rare Earths29(2): 168–172.

11. Zhi, X. H., Xing, J. D., and Fu, H. G.2008. Effect of niobium on the as-cast mi-crostructure of hypereutectic high chromiumcast iron. Materials Letters 62(6-7): 857–860.

12. Chung, R. J., Tang, X., and Li, D. Y.2009. Effects of titanium addition on mi-crostructure and wear resistance of hypereutec-tic high chromium cast iron Fe-25wt.%Cr-4wt.%C. Wear 267(1-4): 356–361.

13. Zhi, X. H., Xing, J. D., and Fu, H. G.2008. Effect of titanium on the as-cast mi-crostructure of hypereutectic high chromiumcast iron. Materials Characterization 59(9):1221–1226.

14. Wu, X. J., Xing, J. D., and Fu, H. G. 2007.Effect of titanium on the morphology of pri-mary M7C3 carbides in hypereutectic highchromium white iron. Materials Science and En-gineering A 457(1-2): 180–185.

15. Suzuki, A. 1999. Effect of multiplycharged ions on the Vickers hardness of TiCfilms. Japanese Journal of Applied Physics 38:881–885.

16. Bramfitt, B. L. 1970. The effect of car-bide and nitride additions on the heteroge-neous nucleation behavior of liquid iron. Met-allurgical and Materials Transactions B 1(7):1987–1995.

17. Dennis, W. H., and William, V. G. 2008.Crystallography and metallography of carbidesin high alloy steels. Materials Characterization59(7): 825–841.

18. Shtansky, D. V., Nakai, K., and Ohmori,Y. 1999. Crystallography and interface bound-ary structure of pearlite with M7C3 carbidelamellae. Acta Materialia 47(4): 1105–1115.

235-sWELDING JOURNAL

WE

LD

ING

RE

SE

AR

CH

Do You Have a Story to Tell?If you have a welding-

related issue you’d like to write about, please send it our way.Submissions from members are being taken for the American

Welding Society’s new blog. Send your article torcorona@aws.org and it could get posted!

Authors: Submit Research Papers Online

Peer review of research papers is now managed through anonline system using Editorial Manager software. Papers can besubmitted into the system directly from the Welding Journal pageon the AWS Web site (www.aws.org) by clicking on “submitpapers.” You can also access the new site directly atwww.editorialmanager.com/wj/. Follow the instructions to registeror log in. This online system streamlines the review process, andmakes it easier to submit papers and track their progress. Bypublishing in the Welding Journal, more than 68,000 memberswill receive the results of your research.

Additionally, your full paper is posted on the American WeldingSociety Web site for FREE access around the globe. There are nopage charges, and articles are published in full color. By far, themost people, at the least cost, will recognize your research whenyou publish in the world-respected Welding Journal.

Want to be a Welding Journal

Advertiser?

For information, contact Rob Saltzstein at (800) 443-9353, ext. 243,

or via e-mail at salty@aws.org.

Zhou Supplement August 2012_Layout 1 7/11/12 3:04 PM Page 235

and the IIW. His many awards include theAWS Honorary Membership, NationalMeritorious, Safety and Health, and theRobotic and Automatic Arc Welding.Hinrichs is survived by his wife Patriciaand family members. The AWS Milwau-kee Section has established the John Hin-richs Memorial Endowment Scholarship.To contribute to to this memorial, contactVicki Pinsky, vpinsky@aws.org, (800/305)443-9353, ext. 212.

James M. Sawhill Jr.

James M. Sawhill Jr., 71, an AWS Fel-low, died suddenly June 1 in James CityCounty, Pa. A native of Baltimore, Md.,

he was a Marylandstate championwrestler in 1959 andran the New YorkCity Marathon in atime of 3:30. He re-ceived his degree inmaterials sciencefrom North CarolinaState University inRaleigh where hewas named Out-

standing Engineering Senior with a 4.0average. He received his master’s in 1966from Lehigh University and PhD in ma-terials engineering from Rensselaer Poly-technic Institute in 1972. Sawhill, activewith AWS and ASM International, madenumerous contributions to the weldingand metallurgical fields. He received onepatent and published more than 20 tech-nical papers. Because of his passion forthe children he befriended while workingthrough Rotary International with theRefugio de los Sueños in Quito, Ecuador,the family requests memorials be made tobenefit the Jim Sawhill Memorial Project,Yorktown Rotary Foundation, PO Box142, Yorktown,VA 23690. Sawhill is sur-vived by his wife, Mary, two daughters,two sisters, and a grandson.

Frank D. Pigage

Frank D. Pigage, 78, died April 6 inFort Myers, Fla. Heserved on the execu-tive board of theAWS PhiladelphiaSection for manyyears. Pigage was awelding distributorsales person for morethan 20 years with E.R. Joseph Co., inNorristown, Pa. Healso worked for the

Fisher Tank Co. in Chester, Pa., and BOCAirco in Reading, Pa. Pigage is survivedby his wife, Phyllis, three sons, four grand-children, and a sister.◆

James Sawhill Jr.

Frank Pigage

— continued from page 94

sales@fosterprinting.comCall 866.879.9144 or

sales@fosterprinting.comCall 866.879.9144 or

AUGUST 2012236-s

Personnel August_Layout 1 7/13/12 8:38 AM Page 236