Coating Degradation in Hot Press Forming

ISIJ International, Vol. 50 (2010), No. 4, pp. 561–568

1. Introduction

Hot press forming (HPF), also known as hot stamping,press hardening or die quenching, is an alternative technol-ogy to produce ultra high strength steels (UHSS) with atensile strength above 1 GPa for automotive bodies. HPFsteels are mainly used in passenger safety-related anti-in-trusion barrier parts, such as door impact beams, bumpers,pillars, roof rails, and tunnels. These anti-intrusion parts areincreasingly being made of UHSS with a martensitic mi-crostructure obtained by HPF. In the HPF process coldrolled coated steel blanks are heated in a gas fired furnaceto a temperature in the range of 900–950°C for 3–10 min toachieve full austenization. The hot sheet steel is then fed ina forming press equipped with water-cooled dies which si-multaneously deform and quench the part to a temperaturebelow the martensitic transformation finish temperature Mf,which is typically �200°C.1–4)

In earlier HPF technologies bare steel was utilized, and athermal oxide scale was formed due to the high temperatureoxidation during heat treatment. The scale was removed bychromium shot blasting, which left a thin film of chromiumand iron on the surface, thus eliminating the need for re-oil-ing to prevent oxidation and corrosion.2) Coated steel sheetsare currently used in the HPF process instead of bare steel.The coatings prevent surface oxidation and decarburization,and enhance corrosion resistance.

A composite coating has recently been developed forHPF.5) This 6–7 mm thin coating consists of a hybrid inor-ganic–organic matrix filled with aluminum particles andsolid lubricants, which can prevent scale formation duringheat treatment and largely reduce the friction during hotforming. This type of coating must be removed by sand-blasting before further treatment because it cannot be resist-ance spot welded. Composite coatings with higher spot

weldability and corrosion resistance are under develop-ment.

Al based and Zn based metallic coatings are currentlywidely used for HPF. Imai et al.6) has studied the behaviorof galvannealed HPF steels. A galvannealed coating typi-cally contains a thin G intermetallic layer (�1 mm) and athick d intermetallic layer. After the HPF process, the top-most layer consists mainly of Zn oxide. A Zn solid solutionlayer is formed at the coating/steel interface. It contains20–30 wt% Zn. The galvanized coating has been reportedto be utilized for as coating on a 22MnB5 HPF steel.7) Theoxide layer and ZnFe intermetallic layers become thicker asthe heating time increases. The Zn rich oxides must be removed by an abrasive blasting in order to obtain adherentpaint layers. After abrasive blasting, the intermetallic layerremains on the steel. This intermetallic layer causes the for-mation of fine cracks in cyclic stress condition. Thesecracks have been found to stop at the interface betweencoating and substrate steel under certain condition.

In the current contribution, the influences of the thermalcycle and the mechanical deformation on the type 1 alu-minized coating during HPF will be reviewed. Two types ofhot dipped aluminized coatings are in industrial production.Type 1 aluminized coating consists of an Al–Si alloy with anear eutectic composition (7–11 wt% Si). Its melting tem-perature is 577°C.8) The Si addition in type 1 aluminizedcoating results in the formation of an inhibition layerFe2SiAl7 at the interface between coating and substratesteel. This layer suppresses the formation of Fe2Al5 phaseand generally delays the further formation of the inter-metallic phases in service. The coating has excellent resist-ance to both corrosion and elevated temperature oxidation,as required in the application such as for e.g. automotiveexhaust systems, heating boilers and cookers. Typical coat-ing thicknesses range from 15 to 25 mm. Most of the HPF

Coating Degradation in Hot Press Forming

Dong Wei FAN,1) Han Soo KIM,1) Jin-Keun OH,2) Kwang-Geun CHIN2) and B. C. De COOMAN1)

1) Graduate Institute of Ferrous Technology, Pohang University of Science and Technology, 790-784, Pohang, South Korea. E-mail: [email protected] 2) POSCO Technical Research Laboratories 545-090, Gwangyang, South Korea.

(Received on November 2, 2009; accepted on January 20, 2010)

The thermo-mechanically induced microstructure changes occurring in a type 1 aluminized coating on hotpress forming (HPF) steel were studied in detail. The formation of intermetallic phases at the soaking tem-perature prior to die quenching revealed that the coating matrix consists mainly of FeAl2 intermetallic phaseby the time the press forming carried out. Kirkendall void formation was observed to take place. The ther-mal oxidation of aluminized coating during HPF was found to be limited, with the coating acting as an effec-tive barrier for the oxygen during heating. The deterioration caused by the high temperature plastic defor-mation was shown to lead to coating cracking without loss of adhesion. The steel surface was oxidizedwhere it was exposed between the coating segments.

KEY WORDS: hot press forming; aluminized; intermetallic; diffusion; Kirkendall voids.

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steels are coated with a 25 mm thick Al–10wt%Si type 1coating. Type 2 coating consists of pure aluminum, and isused mainly for high reflectivity surface appearance, suchas building cladding panels and ventilation systems. Thetypical thickness of type 2 coating is 38–60 mm.9)

2. Experimental

The material used in the present study was cold rolledand continuous hot dipped Al–10wt%Si aluminized steelsheet of 1.6 mm thickness supplied by the POSCO Techni-cal Research Laboratories. The steel chemical compositionis given in Table 1. The steel is a CMn steel with additionsmainly of Cr and B. Carbon, an austenite stabilizer, is themost effective strengthening elements for martensite via themechanism of interstitial solute strengthening. Weldabilityrequirements limit the carbon equivalent to a maximum of about 0.5. A carbon content of 0.2 wt% C ensures the decomposition of austenite into a high strength lath mar-tensite microstructure when relatively low cooling rates are used, providing B-additions are made. Mn is an effec-tive hardening addition as it retards most austenite decom-position reactions. Cr is another strong hardenability agent,which effectively suppresses the bainite transformation. B is added to HPF steels to increase the hardenability by etarding the heterogeneous nucleation of ferrite at the austenite grain boundaries. Additions of Al and Ti are often used to avoid the formation of BN. AlN or TiN pre-cipitates are formed as instead. TiN also suppresses graingrowth.10,11)

Rectangular specimens with dimension of 40�150 mmwere cut. Their weight was approximately 100 g. The speci-mens had large enough surfaces to minimize the effect ofthe increased weight caused by the oxidation of the cutedges during high temperature tests. The specimens weretransferred to a box furnace with a stationary air atmos-phere and kept at 930°C for 2–120 min. The specimenswere then removed from the furnace and cooled in air toroom temperature. The specimens were weighted beforeand after the heat treatment. The specimens were grindedand polished with a 1 mm diamond suspension. The crosssectional specimens were not etched in order to avoid themeasurement errors introduced by the specimen topogra-phy. The cross sections were observed in a Zeiss Ultra 55FE-SEM and analyzed by means of Energy DispersiveSpectroscopy (EDS) and Wavelength Dispersive Spec-troscopy (WDS). The specimens were also analyzed byGlow Discharge Optical Emission Spectrometry (GDOES)to obtain elemental depth profiles. The GDOES analysiswas carried out on a 4 mm diameter area. This gave an aver-aged elemental depth distribution rather than a more precisemicroanalysis as in the case of EDS or WDS. The alu-minized steel sheets were also deformed at 800°C in aGleeble 3500 thermo-mechanical simulator, and the deteri-oration of coating due to a uniaxial plastic deformation wascompared to the deformation-free degradation.

3. Results

3.1. Determination of Stoichiometry of the Intermetal-lic Phases

Figure 1 shows cross sections of the aluminized coatingbefore and after holding at 930°C for different times. Fig-ure 1(a) is the aluminized specimen prior to the heat treat-ment, while (c), (e) and (f) are the specimens heated at930°C for 2, 5 and 8 min, respectively. The specimen crosssections were observed in the FE-SEM and analyzed byEDS. The composition of the phases indicated by the num-bers in each figure is listed in Fig. 1. The solid state reac-tion between the aluminized coating and the steel substrateabove the liquidus temperature of the alloy coating hasrarely been studied in the past. In the present study, the for-

ISIJ International, Vol. 50 (2010), No. 4

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Table 1. Chemical composition of the HPF steel (wt%).

Fig. 1. Cross sectional FE-SEM micrographs for Al–10wt%Sicoated HPF steel, (a) as aluminized, (b) detail of the coat-ing/steel interface, (c) after 2 min at 930°C, (d) detail ofthe specimen after 2 min at 930°C; (e) after 5 min at930°C, (f) after 8 min at 930°C. The table lists the com-position of the phases identified by the numbers in themicrographs a–f.

mation of intermetallic phases was studied and their stoi-chiometry was determined by EDS measurement. TheFe2O3, Al2O3 and SiO2 were selected as standard samplesfor Fe, Al and Si analysis respectively.

The thickness of the aluminized coating (Fig. 1(a)) wasapproximately 25 mm. The coating consisted mainly of analuminum matrix and pure Si formed in the eutectic reac-tion during the cooling of aluminizing process. The peaksin the local Si content (Fig. 2(a)) indicate the presence ofpure Si in the aluminum coating matrix. Phases with anelongated shape (point 1 in Fig. 1(a)) and bright contrastwere observed. Between the coating and the steel substratea layer with bright contrast (point 2) was clearly visible. Itsthickness was about 5 mm. Both the phases with the elon-gated shape and this layer were found to be Fe2SiAl7. At theinterface between the intermetallic phase Fe2SiAl7 and thesubstrate steel (point 3 in Fig. 1(b)), a thin layer of FeAl3

was also detected.In Fig. 1(c), the coating thickness is approximately

38 mm. In the coating, five distinct regions with differentcontrast were detected. The point 4 in the outermost part of

coating, which had the darkest contrast and contained thehighest Al content (Fig. 2(b)), was identified as pure Al. Inthe magnified figure (Fig. 1(d)) point 5, which had a darkercontrast, was identified as the intermetallic phase Fe2SiAl7.From Fig. 1(c), it can be seen that the surface consisted ofpure Al and Fe2SiAl7. In Fig. 1(d) the contrast at points 6and 7 was brighter than at point 5. The Fe/Al averageatomic ratio of point 6 obtained by EDS was close to 0.4corresponding to the stoichiometry of the intermetallicphase Fe2Al5. The Fe/Al average atomic ratio at point 7 wasclose to 0.5 corresponding to the composition of the inter-metallic phase FeAl2. The composition profiles of Fe2Al5

and FeAl2 intermetallic phases can be seen in Fig. 2(b). InFig. 1(c), the regions with the brightest contrast (e.g. point8) inside the FeAl2 intermetallic were identified as Fe2SiAl2.

Figure 1(e) shows the specimen after a 5 min holding at930°C. This condition corresponds to a typical industrialaustenizing time prior to HPF. Compared to Fig. 1(c), thecoating had a more uniform contrast, as a result of the ho-mogenization of Fe, Al and Si contents in the coating. Mostof the coating consisted of the intermetallic phase FeAl2

(point 11) with approximately 3 at% Si in solid solution(Fig. 2(c)). Some isolated intermetallic phases were stillpresent in the coating, as indicated at the points 9 and 10.The intermetallic phase at point 9, close to the top surfaceand with a relatively darker contrast, was identified asFe2SiAl2. The intermetallic phase at point 10, close to thesubstrate and with a brighter contrast, was identified asFe5SiAl4.

Some coating features with a non-uniform contrast wereobserved. They contained both Fe2SiAl2 and Fe5SiAl4 asshown in Fig. 1(e). Images taken at a higher magnificationare shown in Fig. 3. Comparing Fig. 3(a) with (b), it can beseen that the Fe2SiAl2 phase grew in size with time. Com-paring Fig. 3(b) with (c), it is clear that the Fe2SiAl2 phase,which had a darker contrast, gradually transformed toFe5SiAl4, which had a brighter contrast by Fe enrichment.Figure 2(c) shows the composition profiles of both theFe5SiAl4 and Fe2SiAl2 phases. Due to the gradual change ofFe content it was not possible to precisely separate the dif-fusion region from the substrate steel in SEM images (Fig.1). The diffusion region will be discussed in the followingparagraph. From Fig. 2(c) it is clear that the thickness of thecoating had increased to approximately 42 mm after a 5 minholding at 930°C.

After holding for 8 min at 930°C, the coating consistedmainly of FeAl2 (point 14). Some intermetallic Fe5SiAl4

was still present in the FeAl2 matrix (points 12 and 13). Theintermetallic phases grew larger and tended to form twoseparate layers parallel to the original coating/steel inter-face. The thickness of the coating was approximately45 mm (Fig. 2(d)).

3.2. Kirkendall Void Formation

Figure 4 shows the detail of the interface between thecoating and the steel substrate. In Fig. 1(b) and Fig. 2(a),the original interface between the coating and the steel sub-strate is sharp. As the specimen was heated for longer timethe coating/steel interface became more diffuse (Figs. 4(a),(b) and (c)). From Figs. 2(b), (c) and (d) it can be seen thatthe diffusion zone became thicker with increasing holding


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Fig. 2. Cross sectional EDS depth profiles for Al, Si and Fe ofthe Al–10wt%Si coating, (a) as aluminized, (b) after2 min at 930°C, (c) after 5 min at 930°C, and (d) after8 min at 930°C.

time at 930°C. In Fig. 4(a) the diffusion zone is not clearlyvisible because the thickness of the diffusion zone is onlyabout 1 mm (Fig. 2(b)). As time increases, this diffusionzone becomes thicker and more visible (Figs. 4(b) and (c)).In the concentration profile of 2 min heated specimen (Fig.2(b)), there is a 2 mm thin layer between the coating matrixFeAl2 and the diffusion zone, which was identified as inter-metallic phases Fe2SiAl2. Similar situation was observed inthe specimens heated for 5 and 8 min (Figs. 2(c) and (d)).The thin layers of intermetallic phase Fe5SiAl4 were foundat the interface between the coating matrix FeAl2 and thediffusion zone. Due to the thin thickness and close contrastwith diffusion zone, these intermetallic layers cannot beclearly seen in the FE-SEM micrographs (Figs. 4(a), (b)and (c)). Kirkendall voids were observed in the diffusionzone. The Kirkendall voids became larger and merged atlonger holding time.

3.3. Surface Porosity and Thermal Oxidation

In Fig. 1(a), the original surface of coating can be seen tobe flat without any pores. In Figs. 1(e) and (f), surfacepores can clearly be observed at the surfaces of the coating.These surface pores became larger as the holding time increased at 930°C.

Figure 5 shows the increased sample weight during hold-ing at 930°C caused by thermal oxidation. It is apparentthat the weight increment was very limited. This small in-crement of coating weight is very likely the result of a verylimited coating oxidization, due to the fact that the oxygendid not diffuse deeply into the coating and the substratesteel was not oxidized, which can be seen from Fig. 6 thatin 120 min oxygen only reached about 5 mm which is muchless than the coating thickness (�25 mm). It can also beseen from Fig. 5 that initially the specimen weight in-creased slowly, while after 8 min the weight increased more


564© 2010 ISIJ

Fig. 3. FE-SEM micrographs of the intermetallic phase transfor-mation form Fe2SiAl2 to Fe5SiAl4 after (a) 2 min, (b)5 min and (c) 8 min at 930°C.

Fig. 4. FE-SEM micrographs showing the formation and growthof Kirkendall voids at the coating/steel interface after (a)2 min, (b) 5 min and (c) 8 min at 930°C.

Fig. 5. Specimen weight changes as a function of the square rootof the holding time at 930°C.

Fig. 6. GDOES oxygen depth profiles after different holdingtime at 930°C.

rapidly with time. From that point onward, a linear relation-ship with the square root of heating time was found. Thepresence of Al2O3 in the coating was detected by means ofX-Ray Diffraction (XRD). The experimental results clearlyindicate that a high temperature (930°C) holding time inexcess of 8 min was needed for a pronounced coating oxi-dation.

3.4. Coating Cracking

Large cracks were observed in the coating (Figs. 1(c), (e)and (f)) after holding at 930°C, prior to deformation. Thecracks were found to initiate from the surface of the coatingand stop in the diffusion zones (Figs. 1(e) and (f)). Thecracks did not reach the steel substrate. In HPF, the coatedsteel goes through a high temperature deformation processafter annealing. During deformation, the brittle intermetal-lic phases will not deform plastically and break. Figures7(a) and (b) show a comparison of the Al–10wt%Si coatingfor undeformed and deformed specimens. Figure 7(b) is aspecimen deformed to an engineering strain of about 30%at a strain rate of 0.5/s. Compared with the undeformedsample (Fig. 7(a)) the coating on the deformed sample wasseen to be broken into segments. Between the broken seg-ments of the coating bare steel was exposed. The averagelength of coating segments was 280 mm. The average lengthof the exposed steel gaps between coating segments was160 mm. The lengths of coating segments and gaps betweenthe segments were shorter when the strain rate increased,i.e. the coating fracture pattern became finer at higher strainrates (Fig. 7(c)).

In Fig. 7(d) the edge of a broken coating is seen to peeloff from the steel surface and there is a thin layer of thermaloxide (1 mm) between the coating and steel substrate, whichis highlighted by a white box in Fig. 7(d). A small gap canbe observed between this layer and substrate steel, whichimplies that this layer might easily peel off. EDS and WDSanalysis revealed that this layer was FeOx.

4. Discussions

The composition of the intermetallic phases observed inthe present work is shown in Fig. 8. Based on the analysispresented in the previous paragraphs, the original Al–Sialloy coating is changed to a coating consisting mostly ofintermetallic phases. The changes are the following. In the initial stage, the coating composition is Al–9.6at%Si(Al–10wt%Si), and an intermetallic phase Fe2SiAl7 with athin FeAl3 layer is present at the coating/steel interface.During the next stage, the steel is transferred to the furnaceand the temperature increases to 930°C. This temperature is higher than the melting point of the type 1 aluminizedcoating (577°C) and higher than the melting point ofFe2SiAl7.

8,12) The coating melts, and the Si and the Fe2SiAl7

intermetallic phase dissolves into liquid aluminum. AFe–Si–Al liquid solution is formed. Fe diffuses into thecoating and reacts with the liquid alloy coating, and thenthe coating gradually becomes rich of Fe. Fe2SiAl7 (Fig. 8)can be formed during cooling at this stage, and the Si con-tent changes slightly from 9.6 at% in the coating to 10 at%in the intermetallic phase. As holding time increases, moreFe diffuses into the coating. Fe2Al5 phase forms and theFe2Al5 transforms further into the FeAl2 phase due to theincrease of its Fe content. Both Fe2Al5 and FeAl2 are solidbecause their melting points are higher than 930°C.9) Theupper part of coating with less Fe content can become toFe2SiAl7 intermetallic phase during cooling. As holdingtime increases, the coating will be full of FeAl2 due to moreFe diffusion. The composition of intermetallic phaseschanges form 70 at% Al (Fe2SiAl7) to 71.4 at% Al (Fe2Al5),then to 66.7 at% Al (FeAl2), and Si is involved into the formation of Fe2SiAl2. In the following stage, as time in-creases Fe continues to diffuse into the coating, theFe2SiAl2 phase gradually grows larger. The Fe2SiAl2 phaseeventually transforms to Fe5SiAl4. During this transforma-tion the Al content is constant at 40 at% (Fig. 8). In the final


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Fig. 7. Cross sectional light optical micrographs of the heattreated coating (a) before and (b) after uniaxial strainingat 800°C. (c) Comparison of segmentation of the coatingafter deformation at different strain rates. (d) Cross sec-tional FE-SEM micrograph showing the delamination ofthe edges of the coating segment at the coating/steel in-terface and the formation of a thin surface oxide film.

Fig. 8. Fe–Al–Si ternary phase diagram showing the composi-tional evolution of the intermetallic phases observed inthe present study.

stage, the Fe5SiAl4 phase continues to grow larger with time(Fig. 3). It is clear that the changes in the coating are essen-tially due to the formation of intermetallics with an increas-ingly higher Fe content. This process is driven by the diffu-sion of Fe.

The intermetallic phase identification in the present studywas compared with references and the results are listed inTable 2. The intermetallic phase Fe2SiAl7 (points 1, 2, 5 inFig. 1) has been reported as the t5 phase in previous publi-cations.9,13) The purpose of Si addition in type 1 aluminizedcoating is to form a Fe2SiAl7 inhibition layer at the inter-face between coating and substrate steel. The presence ofthis inhibition layer can suppress the formation and growthof intermetallic compounds FeAl3 and Fe2Al5 which impaircoating adherence and formability.14) A thin FeAl3 layerwas observed at the interface between Fe2SiAl7 inhibitionlayer and steel substrate. This is in agreement with previouswork.9)

The coexistence of Fe2Al5 and FeAl2 was observed in thespecimen after holding at 930°C for 2 min (Figs. 1(c) and(d)). This coexistence has also been reported by Suehiro et al.15) who studied the thermal degradation of an Al–10wt%Si coating on HPF steel at 950°C. They indicated the intermetallic phases on a calculated ternary phase dia-gram with the composition obtained from EDS. They re-ported that after 30 s heat treatment the coating consisted offive separate sub-layers, from surface to steel: Fe2Al5, an or-dered BCC phase, FeAl2, a second ordered BCC phase, anda disordered BCC phase. Whether pure Al was present atthe coating surface after a short 30 s heating was not re-ported. The presence of Fe2Al5 was also reported by Jenneret al.16) who studied solid phase reaction of a type 1 coatingheat treated at 925°C for 9 min. The presence of FeAl2 wasnot reported. Wang17) reported the presence of both FeAl2

and Fe2Al5 as two separate phases in a type 2 aluminizedcoating heated at 750°C for 2 h. The same authors reportedthe absence of Fe2Al5 after heating at 950°C for 2 h. In thework of Chang et al.,18) the heating of an Al–7wt%Si coat-ing at 850°C for 10 min was reported to result in the forma-tion of FeAl2 and Fe2Al5 two phases. In the present workFeAl2 and Fe2Al5 were found as two separate phases, andFe2Al5 was transformed to FeAl2. None of the references

mentioned above have reported the evolution of Fe2Al5 toFeAl2 as a result of progressive Fe enrichment.

The presence of Fe2SiAl2 was observed in the specimenafter holding at 930°C for 5 min (Fig. 1(e)). A similar ob-servation was made previously by Jenner et al.,16) who didnot report that this phase eventually disappeared when theholding time was increased as was observed in the courseof the present work. The stoichiometry of this t1b phasehas also been reported in an alternative manner by Li et al.as Fe40Si20Al40.

13) Chang et al.18) has indentified this phaseas t1 phase with composition of 30 at% Fe, 22 at% Si and42 at% Al. In the present study, both Fe2SiAl2 and Fe5SiAl4

intermetallic phases were present in the FeAl2 matrix afterholding at 930°C for 5 min (Fig. 1(e)). In the work of Sue-hiro et al.15) the coating was reported to change gradually totwo sub-layers after heating for 5 min at 950°C: an orderedBCC phase and a disordered BCC phase. The stoichiometryof these two phases was not given.

The intermetallic phase Fe5SiAl4 was observed to appearin the specimen after holding at 930°C for 5 min (Figs. 1(e)and (f)). This intermetallic phase is very likely the same asthe one reported previously by Kubalova et al.,19) who re-ported it as the FeAl1�xSix phase containing 25 at% Si.

The coating/steel interface is sharp in aluminized speci-men (Fig. 1(b)). The interface became more diffuse (Figs.4(a), (b) and (c)) as the holding time increased. This is dueto the formation of an interfacial diffusion zone (Figs. 2(b),(c) and (d)). The concentration profiles of Fe, Al and Si inthe diffusion zone are shown in Figs. 9(a), (b) and (c). Withincreasing the holding time at 930°C, more Fe, Al and Sidiffused into the diffusion zone and this diffusion zone be-came thicker. In Fe–Al binary alloys,20) Al addition closes


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Table 2. Comparison of the intermetallic phases identified inthe present study and reported in previous work.

Fig. 9. The EDS concentration profiles for (a) Fe, (b) Al and (c)Si in the diffusion zone after different holding time at930°C.

the g-loop at a concentration of 2 at%. At this Al content aFe–Al alloy is fully ferritic and there is g to a phase trans-formation when the Al content increases in the diffusionzone and form a Fe–Al alloy with an Al content in excessof 2 at%. In the Al diffusion profiles of Fig. 9(b), most ofAl content is higher than 2 at%, and Si is also present in thediffusion zone (Fig. 9(c)). As Si is also a strong ferrite sta-bilizer, the diffusion zone is expected to be fully ferritic. Inthe present study the diffusion can be considered to occurbetween the steel substrate in the FCC austenitic state andthe coating matrix FeAl2 (Monoclinic) with, in betweenboth phases, a thin BCC ferritic alloy enriched in Al and Si.The relation between the concentration and the diffusiondepth can be expressed as follows.

..................(1)21)

Where M is the number of atoms per unit area, D is the dif-fusivity at a constant temperature, t is the diffusion time,and C is the concentration at position x. For the diffusionwithin a certain time, D · t will be a constant and hence thefirst term in Eq. (1) will be constant. This leads to a linearrelationship between ln C and x2. The diffusivity can there-fore be calculated from the slope of the diffusion profile inan ln C versus x2 plot by least square fitting. Figure 10shows the calculation of the diffusivity of Fe, Al and Si. Inthe experiment the specimens were manually transferredinto the furnace. The furnace temperature dropped to800°C when the furnace door was open. The temperaturerose back to 930°C in 30 s after closing the door. The tem-perature drop would decrease the diffusion in the coating.This effect would be neglected when the holding time wasmuch longer than this 30 s. The diffusivities were calculatedfrom the data of the longest holding time. The values of dif-fusivity at 930°C are DFe�8.0�10�10 cm2/s and DSi�2.8�10�10 cm2/s. In Fig. 10(b), it was apparent that the slope ofspecimen held for 8 min varied. The diffusivities in the twosegments could be calculated as DAl1�5.8�10�11 cm2/s andDAl2�4.2�10�10 cm2/s, respectively. Hence the diffusivityof Fe (DFe) is almost 14 times larger than Al (DAl1), whichmeans there will be a large mass imbalance during the dif-fusion. Due to this imbalance, vacancies will be formed andaccumulated in the diffusion zone. These accumulated va-cancies coalesce and form the observed Kirkendall voids(Figs. 4(a), (b) and (c)). As time increased, the mass imbal-ance became more severe, and the Kirkendall voids thus be-came larger, eventually merging to form larger cavities. Thediffusivity changing position was approximately 3.3 mm tothe top of the diffusion zone (11 mm2 in Fig. 10(b)). The po-sition of the observed Kirkendall voids is in good agree-ment with this value. The reason why the diffusivity of Alis different respecting to different Al content needs furtherinvestigation. It is noteworthy that in the present study thediffusivity of Fe was found to be larger than the diffusivityof Al. In previous reports focusing on the interdiffusion ofFe and Al in the Fe–Al binary system,22,23) it is usually re-ported that the Al diffusivity is larger than that of Fe. Theunusual observation made in the course of the present workis not yet fully understood and it is very likely related to thefact that Si affects the diffusivities of Fe and Al.

The increased surface roughness and formation of sur-face pores in Fig. 1 were already mentioned in previousparagraphs. The surface of aluminized specimen was ini-tially flat and without any surface pores. The surface flat-ness of the specimen held at 930°C for 2 min changed bythe formation of intermetallic phase Fe2SiAl7 (Fig. 1(c)).For the specimens held for 5 and 8 min at 930°C, the sur-face roughness was mainly due to the formation of surfacepores. As the holding time increased the surface pores be-came larger and surface roughness increased considerably(Figs. 1(e) and (f)).

It was reported that at high temperature the intermetallicphase FeAl reacted with moisture to form Al2O3 and H2

(Eq. (2)). The porosity was formed due to the evolution ofH2 gas.24,25)

2Al�3H2O�Al2O3�3H2 ......................(2)

In the present study a similar surface porosity phenome-non was observed. It is not unlikely that the increasedroughness might be related to the presence of H2 generatingfrom the reaction between intermetallic phase and watervapor. The formation of surface pores needs to be furtherinvestigated in more detail.

A clear difference of the surface oxidation before andafter an 8 min holding at 930°C (Fig. 5) was observed in thepresent study. The reason for this could be that at the startpure Al reacted with O2 to form Al2O3. As time increasedthe liquid Al alloy coating gradually changed to a solid in-termetallic phase, and the intermetallic reacted with watervapor to form Al2O3 and H2. Surface pores were therebyformed, which gradually became larger. The reaction areawas thus gradually increased, which would enhance the re-action between the intermetallic phase and moisture. When

ln lnCM

Dt

x

Dt� �

2 4

2

π


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Fig. 10. The ln C versus x2 plot for (a) Fe, (b) Al and (c) Si.

this effect became more pronounced the oxidation was en-hanced. This may explain why there is a plateau in theweight change data in Fig. 5 and that after 8 min the speci-men weight increased more rapidly. The O depth profilesshown in Fig. 6 reveal very high O concentration in the out-ermost parts of the coating. This is very likely an error re-sulting from the surface roughness. The calculation of theO diffusivity at 930°C is shown in Fig. 11, Do�6.7�10�12

cm2/s.The cracks in the undeformed specimens (Fig. 1) were

very likely caused by the different thermal expansion coef-ficients of the intermetallic compounds formed in the coat-ing and steel substrate. The crack tips were found to stop inthe diffusion zones (Figs. 1(e) and (f)), which might be be-cause that the thermal expansion coefficients of the diffu-sion zone is close to the substrate steel. The segmentationof the deformed specimens was due to the difference of me-chanical properties between the intermetallic coating andthe substrate steel. Mckamey et al.25) studied the influenceof atmosphere on the mechanical properties of intermetallicphase FeAl at room temperature. He reported that the totalelongation of intermetallic phase FeAl in moisture atmos-phere is about 2.2%, but when tested in an O2 atmospherethe total elongation was 11.3%. This phenomenon was ex-plained as hydrogen embitterment due to the reaction ofH2O with Al which generated H atoms. Whether the envi-ronment moisture has any effect on the mechanical proper-ties of the coatings on HPF steels needs further investiga-tion.

In the segmented coating (Fig. 7(d)) it can be seen thatthe edge of the coating segment broke at the interface be-tween substrate steel and the diffusion zone, which meansthat the diffusion zone is more coherent with the coatingmatrix than the substrate steel. It can be seen that except itsedges the segmented coating was still had a strong adhesionto the substrate steel.

5. Conclusions

During the HPF process, the original Al–10wt%Si coat-ing reacts with the steel substrate and forms complex inter-metallic phases. These original intermetallic phases evolveto the phases with a higher Fe content as the holding timeincreases. Meanwhile, the coating thickness increases,Kirkendall voids are formed at the coating surface and thesurface roughness increases. In addition Kirkendall voidsappear in the diffusion zone and cracks are formed in thecoating. The intermetallic coating can still prevent the sub-strate steel from oxidation if it is not fractured. Plastic de-

formation caused the segmentation of the coating, and athin FeOx oxide layer is formed at the exposed steel surface.

In summary there are two challenges facing the use of intype 1 coatings on HPF steels. The coatings become brittleand cannot resist a tensile deformation. The coating adhe-sion to the substrate steel remains excellent but the coatingsegmentation exposes the bare steel and results in the ther-mal oxidation of the steel. It is therefore imperative to re-move the intermetallic coating after HPF for the reasons ofweldability and phosphating. The reason is that the pres-ence of Al2O3 and surface Fe oxides make the resistancespot welding difficult and lead to poor phosphating per-formance.

Acknowledgements

The authors gratefully acknowledge POSCO TechnicalResearch Laboratories for providing the materials. TheTechnology Innovation Center for Metals and Materials(TICM) is also acknowledged for providing all the analysisequipments in the present work.

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Fig. 11. The ln C versus x2 plot for oxygen.

Coating Degradation in Hot Press Forming

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