25Metallurgical Science and TechnologyVol. 29-2 - Ed. 2011

Continuous galvanizing of martensitic and complex phase steels for automotive

anti-intrusion applicationsA. Chakraborty, R. Kavitha, J.R. McDermid – McMaster Steel Research Centre, McMaster University, Hamilton, ON, Canada

B. Voyzelle, E. Essadiqi – Materials Technology Laboratory, CANMET, Ottawa, ON, Canada F.E. Goodwin – InternationalZinc Association, Durham, USA

ABSTRACT

From the perspective of crashworthiness andpassenger safety, martensitic and complexphase Ultra High Strength Steels (UHSS) areideal candidates for automotive anti-intrusioncomponents. However, these steels must beprotected from corrosive environments in orderto maintain the longterm integrity of thestructures involved for which continuousgalvanizing is a cost-effective solution. Severalchallenges have to be overcome in order toprocess the above steels in the continuousgalvanizing line (CGL) while achieving theminimum target tensile strength of 1250 MPa.Steel chemical compositions should be selectedin such a way that maintaining a suitablecooling rate produces martensite or bainite,and also provides a substrate surface withsufficient reactive wetting suitable forgalvanizing. In the present study, steelchemistries were designed around relativelylean compositions based on carbon,manganese and silicon with additionalhardenability being provided by molybdenumor chromium additions. Annealing cycles weredetermined based on the continuous coolingtransformation behaviour of the steels. For bothsteel compositions the target tensile strength of1250 MPa was achieved using austeniticannealing for 120s followed by cooling to roomtemperature at 50°C/s. The steels weresuccessfully reactively wet by the Zn(Al,Fe) bathusing a 95%N2-5%H2, -30°C dew pointprocess atmosphere. From scanning electronmicroscopy, X-ray photoelectron spectroscopyand scanning Auger microscopy it wasdetermined that oxides of manganese, siliconand chromium formed during annealing.However, these oxides did not have an adverseeffect on coatability and both steels formed highquality, adherent coatings.

RIASSUNTO

Gli acciai a resistenza ultra elevata (UHSS),con struttura formata da martensite e da fasicomplesse, sono candidati ideali perrealizzare componenti anti-intrusione atti aincrementare la sicurezza dei passeggerinel caso di scontro frontale di autoveicoli.Per mantenere a lungo la loro integritàstrutturale in ambienti aggressivi questiacciai necessitano di un trattamentoanticorrosione. A questo fine ben si prestala galvanizzazione continua, che presentaanche il vantaggio di un costo limitato. Perconseguire come obiettivo minimo la resistenza a trazione di 1250 MPa una volta ultimato il processo digalvanizzazione, occorre tuttavia superarediverse difficoltà. La composizione chimicadegli acciai va scelta in modo che i) si formimartensite o bainite a seguito di unaragionevole velocità di raffreddamento e ii)la superficie del substrato siasufficientemente bagnabile, reattiva e quindiadatta alla galvanizzazione. Nel presentestudio la chimica dell’acciaio è stata basatasu C, Mn e Si con aggiunta di Molibdenooppure di Cromo per aumentare la durezza.Sono stati definiti cicli di rinvenimento basatisul comportamento dell’acciaio al CCT. Conqueste due composizioni dell’acciaio è statoraggiunto l’obiettivo della resistenza tensiledi 1250 MPa, adottando il rinvenimentoaustenitico di 120s seguito daraffreddamento a velocità di 50°C/s fino atemperatura ambiente. Gli acciai sono statipoi immersi in bagno di Zn(Al,Fe) inatmosfera di 95%N2-5%H2, con punto dirugiada a -30°C. Le indagini effettuatemediante SEM, spettroscopia fotoelettronicaa raggi X e microscopia Auger a scansionehanno messo in evidenza che durante laricottura si formano ossidi di manganese,silicio e cromo. Non avendo questi ossidi

KEYWORDS

Martensitic and complex phase steels;continuous cooling transformation;continuous galvanizing; selective oxidation;reactive wetting.

effetti contrari alla formazione dirivestimenti, su ambedue gli acciai si sonoottenuti rivestimenti aderenti e di elevataqualità.

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uniform temperature and coating area on the panel of 90mm x 90mm was centred in thelower portion of the steel panel and only material from this area was analyzed. After heat treatment a cross-section of the steel sample was metallographically polishedusing conventional techniques and etchedusing 2% nitric acid in ethanol.Microstructural analysis of the sample wasperformed using a JEOL 7000F fieldemission scanning electron microscope(FEG SEM). An acceleration voltage of10keV was used for all samples. Vickers micro-hardness was measured in a

INTRODUCTIONIn order to lower vehicle weight and reducefuel consumption and undesirablegreenhouse gas emissions withoutdeteriorating crashworthiness andpassenger safety, the automotive industry isadopting the use of thinner gauge ultra-high strength steels (UHSS) for anti-intrusion components such as side impactdoor reinforcement beams, bumper beamsand roof reinforcements [1]. However, toensure long-term stability of thesestructures, these steels must be protectedfrom corrosion, for which continuous hot-dip galvanizing is amongst the most cost-effective solutions. Several challenges haveto be overcome in order to process theabove steels in the continuous galvanizingline (CGL) without compromising the alloytarget tensile strength (>1250 MPa). Thechallenges are as follows: (a) maintaining asuitable cooling rate to prevent thetransformation of austenite to ferrite andpearlite during processing rather than thedesired martensite or bainite and, (b)providing a substrate surface which can bereactively wet by the Zn(Al,Fe) CGL bath.

Challenge (a) can be addressed bychoosing alloy compositions with sufficienthardenability such that martensitic and/orbainitic complex phase microstructures areproduced using cooling rates compatiblewith the CGL. To prevent tempering duringisothermal holding while galvanizing at460°C, the martensite start (Ms)temperature should be lower than 460°C,thereby preventing formation of martensiteprior to galvanizing. It is well known that inorder to avoid the formation of pearlite,alloying elements such as Mn, Al, Si, Cr,Mo and B can be used [2,3]. Theseelements shift the continuous coolingtransformation (CCT) curves to longercooling times, such that the desiredmicrostructure can be obtained using areasonable cooling rate. On the otherhand, many of the cited alloying elementswill selectively oxidize in commonlyemployed CGL annealing atmospheres,which can in turn prevent reactive wettingby the CGL bath and result in poor qualitycoatings. However, this challenge can beaddressed by choosing proper alloying

elements and a process atmosphere pO2

which is both cost effective and willproduce surfaces with sufficient reactivewetting. Some open literature is available on thegalvanizing of martensitic and complexphase UHSS [4-7], most of which arerelated to electro-galvanizing or theQuench and Partition process. The presentstudy deals with developing a processwindow for the galvanizing of a martensiticand/or bainitic complex phase UHSShaving a target tensile strength in excess of1250 MPa. In this investigation two gradesof steels were fabricated in the laboratory.The majority of the hardenability of thesesteels was provided by the C and Mnadditions, where Si, Cr or Mo were addedto increase the alloy strength andhardenability. The selective oxidation of thealloying elements and their effect on thehot dip galvanizability of the twoexperimental steels was systematicallyanalyzed to determine a suitable processwindow for the CGL.

Table 1. Chemical composition of experimental steels (wt.%).

Sample No. C Mn Si Al S P Mo Cr

steel 1 0.16 2.12 0.58 0.01 0.006 0.016 0.11 0.03

steel 2 0.16 2.14 0.57 0.02 0.006 0.017 0.06 0.12

Table 2. Peak annealing temperatures for experimental steels.

Sample No. Temperature (°C)

steel 1 862

steel 2 890

EXPERIMENTAL PROCEDURE The chemical composition of the experimentalsteels is shown in Table 1. Except for Mo insteel 1 and Cr in steel 2, the compositions ofthe steels were quite similar. Experimentalsteels were fabricated at the MaterialsTechnology Laboratory of CANMET. Alloyswere melted in a vacuum induction furnace,cast, hot rolled, pickled, sand blasted andfinally cold rolled to a thickness of ~0.9mm. Itshould be noted that the cooling from the hotrolling step was controlled such that the as-received microstructure for both alloyscomprised ferrite and pearlite. Theexperimental galvanizing heat treatmentsconsisted of austenitic (100% austenite (γ))annealing for 120s followed by cooling at -50°C/s to the 460°C galvanizingtemperature, dipping for 4s and finally coolingto room temperature at 50°C/s. Galvanizingwas carried out in an iron saturated Zn-0.18wt.%Al (dissolved) bath maintained at460°C [8]. The 100% γ formationtemperatures and continuous coolingtransformation (CCT) diagrams of theexperimental steels were determined using aBÄHR-Thermoanalyse quench dilatometer. Thepeak annealing temperatures for eachexperimental alloy are reported in Table 2.

All heat treatments and galvanizing experiments were carried out using the McMasterGalvanizing Simulator. Prior to heat treatment, samples were degreased in an 80°Caqueous 2% NaOH solution, rinsed with de-ionized water, cleaned ultrasonically in iso-propanol and dried with warm air. A final cleaning with acetone was performedimmediately prior to the sample entering the galvanizing simulator. In the present study,all heat treatments were carried out in the quartz lamp infrared furnace under a 95%N2-5%H2 atmosphere with a controlled dew point of -30°C. The sample thermal cyclewas controlled using a type K (0.5mm) thermocouple welded directly to the sample beforethe start of the experiment. Experimental samples comprised 120mm x 200mm panelswith the longitudinal axis of the sample parallel to the rolling and dipping directions. A

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Clemex micro-hardness tester using a200gf load and CMT 5.0 software.Hardness measurements were taken at fivedifferent locations for each sample. Tensiletesting was performed using an Instron100kN electro-mechanical load frame withMerlin software. Tensile samples were cutusing an abrasive water jet and finishmachined using conventional carbide toolsto the dimensions outlined in ASTM E 8/E8M-08 [9] for sub-sized samples with agage length of 25mm. The sample gagewidth was 6mm. A constant crossheadspeed of 1mm/min was used for all tests.An extensometer was used during tensiletesting to measure sample extension. In order to analyze the selective oxidation ofthe steel surface, 50mm x 10mm samples weremirror polished to remove surface roughness.Samples were then heat treated withoutgalvanizing and stored in anhydrous iso-propanol to minimize further oxidation beforeSEM, X-ray Photoelectron Spectroscopy (XPS)and Scanning Auger Microscopy (SAM)examination. The effect of annealing heattreatment parameters on selective oxidationand oxide morphology were determined usingthe JEOL 700F FEG-SEM. An accelerationvoltage of 10keV was used for all samples.Samples were sputter coated with gold toavoid sample charging. Chemical analysis ofthe oxidized steel surfaces was performed by

XPS using a PHI Quantera X-ray photoelectronspectroscope with an Al Kα Xray source(Physical Electronics, Chanhassen, MN). Thespot size used was 100µm and the take offangle was 45°. All spectra were calibratedusing the metallic iron peak position at706.62eV. The binding energy valuesobtained were accurate to within ±0.1eV, andthe accuracy of the chemical compositionmeasurements was ±5% of the measuredvalue in atomic percent for each element. Thebinding energy results shown were collectedafter sputtering with Ar to a depth ofapproximately 5nm to avoid oxide layersformed during sample handling. Elementaldepth profiles were obtained by Ar sputteringfollowed by XPS analysis of the new surface.Sputtering depth measurements wereconsidered to be accurate to within ±10%. The correlation between oxide morphologyand elemental distribution was studiedusing SAM and Auger elemental mapping.SAM data was collected using a JEOLJAMP-9500F field emission Augermicroprobe. The energy of the primaryelectron beam was 15keV for all samples.Samples were tilted at 30° towards theAuger electron analyzer and auto-probetracking was in effect during the collectionof elemental maps to eliminate sample drift.Auger elemental maps were collected after

Ar sputtering to a depth of 20nm in orderto eliminate the effect of oxidation arisingfrom sample handling. The accuracy of thesputtering depth was ±10%. Steel reactive wetting by the liquid Zn bathwas studied by measuring the bare spotarea in the galvanized coating.Measurements of any bare spots present onthe galvanized panels were performedusing a Nikon AZ100M Stereoscope on the90mm x 90mm uniform coating area.Images were captured using 10xmagnification. The bare spot area wasmeasured by manually selecting the bare spotperimeter. The reactive wetting behaviour wasalso investigated by analyzing the steel/coatinginterface via SEM. The interfacial layer on thegalvanized panels were exposed for SEManalysis by stripping the zinc overlay using: (i) aninhibited 10% H2SO4 solution in water, whichleaves any Fe–Zn intermetallics as well as theinhibition layer (Fe2Al5-xZnx 0<x<1) intact and (ii)fuming nitric acid (HNO3), which removes anypure Zn and Fe–Zn intermetallics, leaving only theFe-Al interfacial layer intact. Adherence of thecoating to the substrate steel was determined usingthe 180° ‘U’ bend test per ASTM A 653/A653M-09 [10]. Flaking of the coating wasanalyzed using a Nikon AZ100M stereoscope atthe top of the ‘U’ bend sample.

RESULTS AND DISCUSSION Steel composition plays a vital role in theproduction of galvanized martensitic andcomplex phase steel. Alloying elements suchas Mn, Si, Mo and Cr improve steel quench-hardenability by concentrating the carbon inaustenite, which increases its stability, and aidsin the formation of martensite or bainite. Onthe other hand, except for Mo, these alloyingelements will form stable oxides whenexposed to typical CGL annealing furnaceatmospheres and can prevent reactive wettingby the liquid zinc bath during galvanizing. Forthese reasons, the steel composition should bechosen in such a way that it produces sufficientquench-hardenability to be compatible withthe cooling capabilities of the CGL along withallowing the production of a surface suitablefor reactive wetting by the CGL bath. Figure 1 shows the experimentally determinedCCT diagrams for the present steels. Steelmicrohardness values as a function of coolingrates are appended to the CCT diagram. From

this figure it is observed that for both steels theMs temperatures are well below thegalvanizing bath temperatures (460°C). Sincethe martensitic transformation takes place aftergalvanizing there is no possibility of temperingthe martensite during galvanizing. Figure 1also shows that in order to achieve theminimum target strength of 1250 MPa, aminimum cooling rate of 50°C/s was requiredfor both steels. As shown by the highlightedcooling curves in Figure 1, the predictedmicrostructures at this cooling rate for bothsteels are dominated by bainite and martensitewith minor amounts of ferrite. Austenitic annealing was carried out at 862°C

and 890°C for steel 1 and steel 2,respectively. Figure 2 shows the steelmicrostructures along with their microhardnessafter completing the specified annealingcycles. In the case of steel 1, the microstructurecontains mostly bainite and lath martensitealong with minor amounts of ferrite, whereasbainite and lath martensite along with aconsiderable amount of ferrite was observedin steel 2. As expected, steel 1 showed higherhardness due to the presence of significantlyless ferrite in the microstructure. Representative engineering stress-straindata for the steels are shown in Figure 3. Asummary of the tensile results (yield

Table 3. Summary of steel mechanical properties after annealing.

Strength (MPa) Elongation (%)

Sample No. YS UTS Uniform Total

steel 1 975 1379 4.50 6.07

steel 2 948 1295 5.29 7.22

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Fig. 1: Continuous cooling transformation diagram of (a) steel 1 and (b) steel 2. (FÜFerrite, PÜPearlite,BÜBainite and MÜMartensite).

a

b

strength (YS), ultimate tensile strength(UTS), uniform and total elongation) for theaverage of five samples are provided inTable 3. From Figure 3 and Table 3, it canbe seen that both steels met the minimumUTS target of 1250 MPa. It was alsoobserved that in steel 1, due to the highervolume fraction of bainite and martensite,there was a significant increase in YS andUTS without a significant deterioration inelongation versus steel 2. From the tensileresults, it can also be observed that, whenprocessed using similar cooling rates, theuse of Mo in the steel compositionproduced higher quench hardenabilitycompared to Cr. Figure 4 shows the SEM microstructures ofoxides present on the steel surfacesimmediately before galvanizing. For bothsteels, nodular shaped oxides weredistributed over the bulk grain surfaces aswell as on the grain boundaries. Figure 5 shows the chemical composition ofthe steel surfaces determined via XPS depthprofiling. Table 4 lists the binding energiesof the elements of interest and the oxidespecies identified on the steel surfaces. Forboth steels, considerable surfaceenrichment of Mn, Si, O and Cr (only forsteel 2) was observed, from which it can beconcluded that the annealing temperaturesignificantly affected the surfacecomposition as well as the oxides present. Itshould be noted that no significant surfacesegregation of Mo was observed for steel1 (Figure 5(a)). As shown in Figure 5(b),the surface of steel 2, showed a slightlyhigher level of Mn (~25at.%) compared tosteel 1 (~23at.%, Figure 5(a)), whereas aslightly higher amount of Si was observedon steel 1 (~13at.%) compared to steel 2(~10at.%). On the steel 2 surface, asshown in Figure 5(b), the amount of Crincreased with increasing depth to 30nm,after which it decreased to bulkconcentration levels. This could be due tothe internal oxidation of Cr. SAM analysis shows the link between oxidemorphology and selective oxidation of thesteel alloying elements. Figures 6(a) and6(b) show the Auger maps for Mn, Si, Mo,Fe and O for steel 1 and Mn, Si, Cr, Fe andO for steel 2, respectively. From Figure 6 itcan be observed that during annealingprior to galvanizing, considerableenrichment of Mn and Si (both steels) andCr (for steel 2) in the form of oxides waspresent on the steel surfaces. From Figure 6

Fig 2: SEM micrographs of steel microstructures after annealing: (a) steel 1, (b) steel 2.

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it can be observed that nodular MnO andmixed (Mn, Si) oxides were distributedover the bulk grain surfaces as well asalong the grain boundaries for both steels.Consistent with the XPS observations forsteel 1 (Figure 5(a)), there was no surfacesegregation of Mo observed during theSAM analysis, confirming that under thepresent annealing process atmosphere Mo-oxides were not formed. On the otherhand, on the surface of steel 2, thesegregation of Cr-oxides was prominentprimarily along the grain boundaries. Asthere is some correspondence between theMn, Cr and O maps for steel 2, it is possiblethat some nodules comprised mixed(Mn,Cr) oxides.The ability of the zinc bath to reactively wetthe steel surface was determined byassessing the population and size of barespots on the coating and by observing theFe–Zn interface to determine if anyreaction products were present. The areapercentages of bare spots and averagebare spot size within the uniform coatingarea are summarized in Table 5. Bothsteels showed excellent coatability despitethe presence of oxides on the surfacebefore galvanizing (Figures 4, 5, 6 andTable 4). This may be the result of in-situaluminothermic reduction of the surfaceoxides by bath dissolved aluminumexposing the underlying Fe, therebyallowing reactive wetting to occur [11,12]. Areas of the galvanized panels thatFig 4: SEM micrographs of the oxides formed on the steel surface during annealing: (a) steel 1 and (b) steel 2.

Fig 3: Engineering stress-strain plots for the experimental steels after annealing.

Fig 5: XPS depth profile of oxides formed on surface during annealing of (a) steel 1 and (b) steel 2.

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Fig 7: SEM micrographs of interfacial layer and Fe-Zn intermetallic phase after the Zn overlay was stripped with10% H2SO4 in water (a) steel 1 and (b) steel 2.

demonstrated good wettability werestripped of their zinc overlay with 10%H2SO4 in water leaving any interfacialFe–Zn and Fe–Al–Zn intermetallics intact.SEM micrographs of these samples areshown in Figure 7. Fe–Zn intermetallicswere identified at the interfaces of bothsteels. Fe-Zn intermetallics should not beformed when galvanizing in a 0.18%dissolved aluminium containing zinc bath.Fe-Zn intermetallic phases typically formdue to breakdown of the inhibition layer orcan precipitate on the Fe-Al layer due tolocal Al depletion. Inhibition breakdownmay have arisen from the heat generatedduring the austenite to bainite or martensitephase transformation, as elevatedtemperatures are known to promote thisprocess [13]. The precipitation of Fe-Znintermetallics may also have beenpromoted by aluminothermic reduction ofsurface oxides by the dissolved aluminiumof the bath, causing local depletion ofaluminium and promoting precipitation ofFe-Zn phases on the Fe-Al interfacial layer. In order to investigate the Fe2Al5-xZnx

(0<x<1) interfacial layer for both steels, awell wetted portion of the galvanizedcoating area was stripped with fumingHNO3 leaving only the Fe-Zn-Al inhibitionlayer intact. Figure 8 shows the SEMmicrographs of these samples. In general, acontinuous Fe2Al5-xZnx (0<x<1) layer waspresent for all samples, indicating completereactive wetting and that the Fe-Znintermetallics observed in Figure 7precipitated on the surface of the Fe-Allayer due to local depletion of dissolved Alin the melt. From the above SEMmicrographs it was also observed that theinhibition layer comprised two sub-layers:the lower sub-layer – next to the substratesurface – consisted of small, equiaxedcrystals and the upper layer of coarse,elongated crystals [13,14]. Coating adherence with the steel substratewas determined by means of a 180° ‘U’bend test per ASTM A 653/A 653M-09[10]. The top surfaces of the ‘U’ bend testsamples were analyzed using an opticalstereoscope. Figure 9 shows the stereo-micrographs of the galvanized samples. Itcan be observed that aside from minorcracking of the galvanized coating, noflaking of the coating was observed andthat, overall, the coating showed goodadherence with the substrate steel.

Table 4. Identification of oxides on the steel surface using XPS .

Measured Binding Energies (eV) of Element State (Compound)

Mn 2p3/2 Mn 2p1/2 Si 2p Cr 2p3/2 Fe 2p3/2 O 1s

MnO, SiO2,Mn2SiO4,

Cr2O3 and Fe (metallic)

641.6 653.4 103.4 576.5 706.6 530.4

Table 5. Bare spot analysis of galvanized panels(total area analyzed: 8100mm2).

Sample No. Average Bare Total Bare Bare SpotSpot Size (mm2) Area (mm2) (%)

steel 1 0.06 9.72 0.12

steel 2 0.10 13.77 0.17

Fig 6: Elemental Auger maps of oxides formed on the surface during annealing of (a) steel 1 and (b) steel 2.

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CONCLUSIONS The hot dip galvanizability of twolaboratory-produced ultra high strengthsteels has been systematically analyzed todetermine a suitable process window forCGL production. Microstructuralcharacterization along with substratemechanical properties as a function ofthermal cycles based on their continuouscooling transformation behaviour was alsodetermined. In the present case, the steelschemistries were designed aroundrelatively lean compositions based on C,

Fig 8: SEM micrographs of inhibition layer after the Zn overlay was stripped with fuming HNO3 (a) steel 1 and (b) steel 2.

Fig 9: Stereo-micrographs of the top surface of 180° ‘U’ bend galvanized samples: (a) steel 1 and (b) steel 2.

Mn and Si with additional hardenabilitybeing provided by Mo and Cr additions forsteel 1 and steel 2, respectively. For both steel compositions the targettensile strength of 1250 MPa (minimum)was achieved using austenitic annealing for120s followed by cooling to roomtemperature at 50°C/s. It was alsoobserved that when using similar coolingrates the Mo-containing alloy producedhigher quench hardenability as comparedto Cr alloy. For both the steels, surfaceswhich could be successfully reactively wet

by the Zn(Al,Fe) bath were produced usinga 95%N2-5%H2, -30°C dew point processatmosphere. From SEM, XPS and SAManalysis it was determined that oxides ofMn, Si and Cr formed during annealingand that these were distributed on the bulkgrain surface as well as grain boundaries.There was no segregation and oxidation ofMo observed. However, these oxides didnot have an adverse effect on thecoatability, with both steels forming highquality, adherent coatings and a welldeveloped Fe-Al interfacial layer. For both

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ACKNOWLEDGEMENTS The authors would like to thank the members of the International Zinc AssociationGalvanized Autobody Partnership members, ArcelorMittal and the Natural Sciences andEngineering Research Council of Canada (NSERC) for their financial support of thisresearch and the Materials Technology Laboratory, CANMET for the provision of theexperimental materials through the RIEM program. The authors would also like to thankDr. Li. Sun of ArcelorMittal (Dofasco) for XPS analysis, Shihong Xu of the Alberta Centrefor Surface Engineering and Science for the Auger analysis and John Thomson, MarianaBudiman, Doug Culley and Steve Koprich of McMaster University for their technicalsupport.

REFERENCES [1] C. SCOTT, N. GUELTON, S. ALLAIN and

M. FARAL, Mat. Sci. and Tech. Conf.Proc., Pittsburgh, USA, (2005), p 127.

[2] J. MAHIEU, S. CLAESSENS, B.C. DECOOMAN and F.E. GOODWIN,Galvatech’04, Conf. Proc., Chicago,USA, (2004), p 529.

[3] E.DE BRUYCKER, B.C. DE COOMANand M. DE MEYER, Steel Res. Int. 75, 2,(2004), p 147.

[4] K.I. MATSUMURA, N. FUJITA and T.NONAKA, Galvatech’07, Conf. Proc.,Osaka, Japan, (2007), p 392.

[5] E. C. OREN and F.E. GOODWIN,Galvatech’04, Conf. Proc., Chicago,USA, (2004), p. 737.

[6] M. BABBIT, Steel Res. Int., 77, 9-10,(2006), p 620.

[7] K.I. SUGIMOTO and B. YU, ISIJ Int., 45,8, (2005), p 1194.

[8] J.R. MCDERMID, M.H. KAYE and W.T.THOMPSON, Met. Mat. Trans. B, 38B,(2007), p 215.

[9] Standard Test Method for TensionTesting of Metallic Materials, ASTMStandard, Designation: E 8/E 8M – 08

[10] Standard Specification for Steel Sheet,Zinc-Coated (Galvanized) or Zinc-IronAlloy Coated (Galvannealed) by theHot-Dip Process, ASTM Standard,Designation: A 653/A 653M – 09.

[11] E.M. BELLHOUSE and J.R. MCDERMID,Met. Mat. Trans. A, 41A, (2010), p 1539.

[12] R. KHONDKER, A. MERTENS and J.R.MCDERMID, Mat. Sci. Eng. A, 463,(2007), p 157.

[13] L. CHEN, R. FOURMENTIN and J.R.MCDERMID, Met. Mat. Trans. A, 39A,(2008), p 2128.

[14] E. BARIL and G.L’ESPERANCE, Met.Mat. Trans. A, 30A, (1999), p 68

steels, Fe-Zn intermetallics were alsoobserved to have precipitated on thesurface of the Fe-Al interfacial layer andmay be due to the depletion of aluminiumfrom the aluminothermic reduction ofoxides on the steel surface duringgalvanizing. It can be concluded that forthe above steel compositions, a suitableprocess window has been developed toproduce galvanized ultra high strength(>1250MPa) martensitic and complexphase steels.