Formation of Al and Cr Dual Coatings by Pack Cementation on

127 © 2012 ISIJ

ISIJ International, Vol. 52 (2012), No. 1, pp. 127–133

Formation of Al and Cr Dual Coatings by Pack Cementation on SNCM439 Steel

Jhewn-Kuang CHEN,* Shih-Fan CHEN and Che-Shun HUANG

National Taipei University of Technology, Institute of Materials Science and Engineering, 1, Sec.3, Zhong-Xiao E. Rd., Taipei,10608 Taiwan. E-mail: [email protected]

(Received on May 23, 2011; accepted on September 27, 2011)

Two stage pack cementation processes are developed to form dual Fe–Al–Cr layers on surfaces ofSNCM439 steels. The first 550°C treatment assists to modulate adequate aluminum activity for theformation of iron-rich intermetallics. In the second 750°C treatment stage, simultaneous chromizing andaluminizing treatments are achieved by first forming a FeAl ferritic layer and then a surface layer withhigher Cr content at later time. The current study examines the effects of second stage 750°C holdingtime, activator concentration, and Cr:Al ratio on coating structures. Fe–Al coatings consisting of Fe3Al andFeAl intermetallic phases are observed to form initially. This Fe–Al layer accounts for 25 to 32 μ mthickness of the coatings and show good adherence with the substrates. The coating thickness increasesparabolically with 750°C holding time. With prolonged treatment at 750°C, surface concentration ofaluminum in powder packs drops with treatment time and increasing concentration of activator. A peakconcentration exists at a depth below substrate surface. Aluminum is back diffused from the steel surfaceinto the powder packs. The growth of Fe–Al intermetallics slows down. Surface layer then forms athickness of 6 μm coating with 2–5 wt.% of chromium. Samples treated for longer than 6 h with over 12wt.% of NH4Cl activator concentration or with Cr:Al ratio higher than 90:10 induce earlier chromiuminfusion and lead to porous coating structure due to Kirkendall effect. Eventually chromium carbide formsto cease further growth of the dual coating structures.

KEY WORDS: pack cementation; aluminization; intermetallics; gas-solid reactions; diffusion.

1. Introduction

Steel surface can be modified by forming Fe–Al interme-tallic compounds as hard coatings. Such coating improveswear resistance. The coated steels can also be employed inhigher temperature and corrosive environment.1,2) Literaturereports that the steel surface after hot dippingaluminization3) can withstand temperature up to 1 093°C(2 000°F).4)

Pack cementation process is an ideal alternative processfor hot dipping aluminization. It is essentially an in situchemical vapor deposition (CVD) coating process.5) Fourinter-processes, namely halide activation, gas diffusion,solid deposition, and solid diffusion reactions, take place insequence to form solid coatings by gas-solid reactions.6,7)

The pack cementation process has the advantages of lowcost and applicability to various materials shapes and sizes.The compositions of coating are dependent on processingtemperature, time, substrate composition, and atmo-spheres.8,9)

Pack cementation aluminization of steel surface is nor-mally performed at temperatures as high as 1 050°C.10,11)

Pack aluminization at temperature below 700°C was report-ed by Xiang and Datta12) who observed the formation of Al-rich Fe2Al5 phase. According to Fe–Al phase diagram,13)

aluminum has high solubility in iron. Fe and Al can also

form intermetallic compounds including Fe3Al, FeAl,FeAl2, Fe2Al5, and FeAl3 below 1 000°C.14) Among these,the Fe–rich Fe–Al intermetallic compounds, eg. Fe3Al andFeAl, are favorable due to their less brittle and moredesirable mechanical properties.15,16) Fe3Al is also reportedto resist sulfidation and oxidation at high temperatures.17)

Minor Cr content in the Fe–Al compounds can furtherimprove their resistance to room temperature aqueous cor-rosion and hot corrosion by fused salt deposits.18) Tosimultaneously chromize and aluminize steels, the partialpressures of Cr-halide and Al-halide in the powder packmust be comparable. Since the partial pressures of Al-halides are normally much higher than those of Cr-halides,the coatings of Cr–Al alloys are only possible when theactivity of Al in the pack is 2–3 orders below that of Cr.18–20)

There are two possible ways to overcome this problem: (1)using a “lean” pack where the metal powders contain higherCr and lower Al concentrations;21,22) or (2) performing dualinstead of single heating processes first by treating at 925°Cand then at 1 150°C.20)

Meanwhile, chromium carbide (Cr23C6) was reported toform at surface due to the rapid outward diffusion of carbonduring simultaneously chromizing and aluminizing. Forma-tion of such carbide layer blocks the inward diffusion of oth-er elements and locally depletes carbon from the steels.Pores are often observed due to vacancy-interstitial interac-

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ISIJ International, Vol. 52 (2012), No. 1

tions.23) Therefore, it is important to carefully control thecombination of chromizing and aluminizing.

In this study, it is our intention to form Fe–Al–Cr inter-metallic coatings on a SNCM439 steel. We intend todevelop a two-stage treatment at low temperatures (550 and750°C) to achieve Cr and Al co-deposition by packcementation. The lower temperature processes are morecost-effective from energy point of view. It is also theobjective of current research to control such that thepreferable Fe-rich instead of Al-rich Fe-Al intermetallics areformed.

2. Experimental Procedures

The substrate used in this study is commercial SNCM439steel. Its nominal composition is listed in Table 1. Thesample size is 30 mm dia × 5 mm t disk cut from a roundbar. These samples are first ground to 800-grit SiC paperand then ultrasonically cleaned in methanol for 10 min.

Pack cementation process places the substrates, orSNCM439 steel in current study, within the powder packs.The powder packs consist of pure metal powders (Al andCr) for diffusion, halide activator (NH4Cl), and inactiveceramic powder (Al2O3) to protect the specimen from beingoxidized. At elevated temperature, halide activator isdecomposed by heating and form metal halides (e.g. AlCl,AlCl3, CrCl2, and CrCl3) with metal powders through gas-eous reactions. The high metal activity at surface of steelcauses metal atoms to diffuse rapidly into the samples bydecomposing metallic halides. An intermetallic compoundlayer is thus formed on steel surface. The decomposed halo-gen gas, e.g. Cl2, then continues to react with metal powdersleft in the powder packs till the activity of metals in powderpack become equilibrium with that at steel surface. There-fore, halide activator plays an important role to control thevapor pressures of metal halides in powder packs21) whichmaintain the activity of diffusing metals at specimen surfaceto sustain the diffusion with aids of gaseous reactions.

The powder packs for current pack cementation processcontain 3–12 wt.% Cr and Al metal powder mixtures, 3–12wt.% NH4Cl activator, and the rest 85 wt.% Al2O3 filler. Themetal powder is a mixture of 75–90 wt.% Cr and rest 10–25 wt.% Al powders. SNCM439 steel specimens arewrapped in powder packs and loaded into a 316L stainlesssteel tube. Ar is supplied to flush the reaction chamber toremove moisture and air for 5 min prior the cementationtreatment. The specimens and powder packs are then treatedusing a two-stage process, first at 550°C for 2 h beforeheating at 10°C/min to 750°C and then hold for 2–8 h.

After treatments, the packs are cooled to roomtemperature by air cooling. The coated samples are removedfrom the pack and ultrasonically cleaned. XRD (X-raydiffraction) analyses are performed on steel surface to identifythe intermetallic phases formed using Rigaku D/max-Bdiffractometer equipped with Cu target. Cross section ofeach sample is cut, mounted, ground and polished for SEM

(scanning electron microscopy, HITACHI S-4700 fieldemission SEM) and EDS (energy dispersive spectroscopy,HORIBA 7200-H) chemical analyses.

3. Results and Discussion

3.1. Effects of Treatment TimeTable 2 lists the compositions of powder packs and

treatment parameters by varying the treatment time at 750°C(sample 1–4). The total coating thickness increases from 25to 43 μm and is proportional to the square root of 750°Ctreatment time with a 0.99 coefficient of linear correlation(R2). The formations of these coatings are evidentlydiffusion controlled. Figure 1 shows that two layers areformed. The thickness of surface layer is approximately 6μm for all conditions listed in Table 2, while the thicknessof sub-surface layer increases with treatment time. In Figs.1(c) and 1(d), pores in the coating layers are observed toincrease with holding time as well.

The surface coating phase is analyzed by XRD to consistof mainly FeAl solid solution in 2 h-treated sample, Fe3Alin 4 h-treated samples, Cr23C6 and Cr7C3 in 6 h treatedsamples, and Cr7C3 in 8 h treated samples (Fig. 2). Theseresults are very different from that reported recently24) by asingle step process at 700°C using Cr-Al alloy metalpowders which forms an aluminum-rich brittle Fe2Al5 phasein contrast to iron-rich Fe3Al phase formed by dual stageprocess in current research.

In the case that Fe2Al5 phase containing 71 at.% of alu-minum is formed at temperature as low as 700°C,24) gaseousaluminum halides apparently provide fairly high activity ofaluminum at the steel surface to form such aluminum-richFe2Al5 coatings. However, Fe2Al5 coatings are not favoreddue to its brittle characteristics. To avoid the formation ofbrittle Fe2Al5 phase, activity of aluminum must be reduced.In current study, the first stage 550°C treatment is designedto reduce the starting activity of aluminum. Therefore, aFeAl solid solution is first formed during the second stage750°C treatment instead of the less favored Al-rich Fe2Al5.

In Fig. 3, Surface aluminum concentration is shown toattain 42 at.% after 2 h-treatment at 750°C. The concentra-tion of aluminum drops with diffusion depth till 24 μm intothe steel. This layer corresponds to ferritic FeAl solid solu-tion and is consistent with XRD analyses. When 750°Ctreatment further prolongs to 4 h, the surface aluminum con-centration is reduced instead of increasing. The surface alu-minum concentration reduces to 25 at.% corresponding toFe3Al phase (Fig. 2) which is the favorable Fe-rich interme-tallic compound coating.

Table 1. Nominal composition (wt.%) of SNCM439 steel.

Element C Si Mn Ni Cr Mo Fe

wt.% 0.36 –0.43 0.15 –0.35 0.6 –0.9 1.6 –2.0 0.6 –1.0 0.15 –0.3 Bal.

Table 2. Thickness of coatings treated for different holding time at750°C stage.

Sampleno.

750°C holdingtime (h)

NH4Cl(wt.%)

Cr:Alin pack

Coating thickness(μm)

0 0 3 85:15 0

1 2 3 85:15 24

2 4 3 85:15 32

3 6 3 85:15 38

4 8 3 85:15 43


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It is also important to note that a highest aluminum con-centration appears at depth of ~6 μm below surface after 4h treatment at 750°C rather than at steel surface (Fig. 3).Apparently, the reduced aluminum concentration at steelsurface indicates that aluminum in powder packs is exhaust-ed after long treatment time at 750°C. The presence of peakaluminum concentration at a depth below steel surfacerequires the aluminum accumulation in steel to diffuse intwo directions, one diffusing further into the steel and theother diffusing back to surface. Diffusions of aluminum intwo directions are also observed in the concentration gradi-ents of 6 h- and 8 h-treated samples. The 6 h- and 8 h- treat-ed samples demonstrate peak concentrations at depth of 12and 18 μ m below steel surface according to Fig. 3, respec-tively. The locations of maximum aluminum concentrationmove toward into the steel with time showing some of thealuminum in-take during the earlier treatment time diffusesfurther into steels.

On the other hand, the surface aluminum concentrationreduces from 42 to 25, 7, and 5 at.% for 2, 4, 6, and 8 h-

treated samples, respectively. These are compared to thepeak concentrations of 42, 28, 12, and 9 at.% for 2, 4, 6, and8 h-treated samples, respectively. It is obvious that the sur-face aluminum concentration is lower than the peak concen-trations for specimens treated longer than 2 h. The concen-tration gradient indicates that part of aluminum in steelcoatings diffuse outward back into the powder packs. Thisdual diffusing activity provides an important mechanismcontrolling the formation of dual coatings in current study.

Chromium concentrations in this series of experiments allremain at 2–5 wt.% in the steel which represent slow butsteady diffusion (Fig. 3). It has been reported18) thatchromium activity in powder packs is well below that ofaluminum, even though chromium powder content isdesigned to be greater than that of aluminum in powderpacks. Furthermore, for chromium diffusion to occurrequires higher temperature. Although, the 750°C treatmentallows chromium to diffuse into the steel, chromium

(a) (b)

(c) (d)Fig. 1. Cross section SEM microstructures of surface coatings with different second stage 750°C holding time: (a) 2 h, (b)

4 h, (c) 6 h, and (d) 8 h.

Fig. 2. XRD spectra of sample 1, 2, 3, and 4 in Table 2. Fig. 3. Al and Cr concentration profiles in coating layers of sam-ples treated for different time at 750°C (samples 1, 2, 3, and4 in Table 2).

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diffuses at a comparably slower rate than aluminumdiffusing outward to the powder packs. The difference indiffusion directions and rates of chromium and aluminumthus causes Kirkendal effects.

Geib and Rapp8) report that, for simultaneous chromizingand aluminizing, porous coating starts to form due to backdiffusion of chromium and aluminum late in the processwhen Cr- and Al-depletion zones form in the powder packs.Because iron has a much higher solubility for aluminumthan chromium, aluminum is the main substitutional atomsdiffusing into the steel in the initial treatment stage. Whenaluminum in-take reaches an equilibrium activity on steelsurface with powder packs, aluminum diffusion slowsdown. The powder packs then start to be depleted of alumi-num and would compete with steel for aluminum.

The chromium atoms continue to substitute aluminum onsteel surface when chromium infuses more pronouncedlyinto the steel due to chromium’s higher activity in the pow-der packs. When chromium diffuses into the steels bydecomposing chromium chlorides in powder packs, activityof chlorine is increased. Therefore, a driving force is pro-duced for chlorine to react with the high aluminum concen-tration at the steel surface. Aluminum is thus migrating backfrom the steel surface into the powder packs while chromi-um diffuses into the steel. The outward diffusion of alumi-num proceeds to react with chlorine in powder packs and tokeep the pack cementation reactions in equilibrium withreduced chromium content in powder packs. Chromium andaluminum thus demonstrates a positive interaction parame-ter in both powder packs and steels.

Although both aluminum and chromium of the powderpacks tend to diffuse into the steel substrate at the beginningof reactions, the sequence of faster aluminum diffusion andslower chromium diffusion causes dual layers to form incompetition. The competition of chromium with aluminumdiffusion into steel is less pronounced in the beginning ofcementation reactions. But after 4 h of diffusion time at

750°C, aluminum activity attains equilibrium between steeland powder pack, and chromium becomes the main ele-ments diffusing into the steel. Further chromium diffusiondrives the earlier diffused aluminum to diffuse back into thepowder pack. Back diffusion is expected, because whenchromium is reduced from the powder pack, the powderpack then has a high chlorine activity and tends to react withaluminum. The aluminum cementation reactions essentiallyproceed in reverse direction when metallic activity is deplet-ed in powder packs.

Eventually, porous coating is formed due to the Kirken-dall effect by aluminum and chromium interdiffusion in thesurface region as 750°C treatment extends longer than 6 h.When surface chromium concentration reaches ~4 at.%,chromium carbides are formed as shown in Fig. 2.

3.2. Effects of Activator ConcentrationTable 3 lists a series of pack cementation samples by

varying NH4Cl activator concentrations (sample 5–8) whilekeeping the metal powder content and 750°C processingtime constant for 2 h. The observations of dual layer coatingthickness does not change much with the concentration ofactivator for samples treated using 3–9 wt.% of NH4Clactivator as shown in Fig. 4 and stay at a thickness of 23–25 μm. In 12 wt.% NH4Cl treated sample, only a thin layerof coating lower than 5 μ m is observed in Fig. 4(d). The

Table 3. Thickness of coatings treating using different NH4Cl con-centrations and their process parameters.

Sampleno.

holdingtime (h)

NH4Cl(wt.%) Cr:Al Al2O3

(wt.%)Coating thickness

(μm)

5 2 3 80:20 85 25

6 2 6 80:20 85 25

7 2 9 80:20 85 23

8 2 12 80:20 85 ––

(a) (b)

(c) (d)Fig. 4. Cross section SEM microstructures of coatings using (a) 3 wt.% (sample 5), (b) 6 wt.% (sample 6), (c) 9 wt.%

(sample 7), and (d) 12 wt.% (sample 8) NH4Cl concentration in the powder packs.


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surface coating phases are analyzed using XRD as shown inFig. 5. In sample 5 and 6, surface coatings correspond toFeAl, while in sample 7 and 8, only carbides and Fe orsubstrate are present. The presence of Fe phase on surfaceanalyses indicates that only small amount of aluminum arediffused into the steel when NH4Cl activator concentrationis greater than 9 wt.%. It is also interesting to note that manypores are observed only on the surface of sample 7 (Fig.4(c)) while there is no pore in sample 8 (Fig. 4(d)).

In sample 7 treated using 9 wt.% NH4Cl, the aluminumconcentration gradient in Fig. 6 demonstrates a peak con-centration of 14 at.% at 7 μ m below steel surface. The alu-minum concentration reduces in both inward and outwarddirections from 7 μ m depth inside the steel. Only 5 at.% alu-minum concentration is observed at surface. It again showsthat aluminum diffuses outward to the powder packs as sug-gested in Section 3.1. This is a confirmation of aluminumdepletion in powder packs in using higher activator concen-tration. On the other hand, chromium continues to diffuseinto the steel substrate at a slower rate. The great amount ofpores are thus formed by Kirkendall effects due to chromi-um and aluminum interdiffusion as explained in Section 3.1.

When activator concentration increases further to 12wt.%, only carbide layer is present (Figs 4(d) and 5) on the

steel surface, and no distinguished subsurface Fe–Al solidsolution is formed. Almost none porosity is observed in thecoatings as well as shown in Fig. 4(d). This indicates thatthe interdiffusion of chromium and aluminum is limited andthus no Kirkendall effect occurs. Apparently, aluminum dif-fusion is restricted when activator concentration is higherthan a critical value.

In this study, when activator concentration reaches 12wt.%, the limited amount of aluminum is only enough toreact with decomposed chlorine from the activator. Notmuch aluminum is available in the powder packs for diffu-sion. Therefore, diffusion of aluminum into steel is limitedand only higher-containing chromium diffusion can pro-ceed. Pores are thus significantly less in sample 8 than sam-ple 7 as shown in Figs. 4(c) and 4(d).

The activator overdose apparently accelerates alumiunumdepletion and makes chromium diffusing at earlier time incomparison with the samples discussed in Section 3.1. Oncethe chromium reaches ~4 at.% at steel surface, chromiumcarbides are formed. The carbide layer serves as surface bar-rier and forbids further growth of coatings as observed inFigs. 4(c) and 4(d). In other word, formation of surface car-bides marks the end of growth for surface coatings.

3.3. Effects of Cr:Al Ratio in PacksTable 4 lists a set of samples coated using varied Cr:Al

ratio in the metal powders while NH4Cl concentration isfixed at 3 wt.% and 750°C treatment time is fixed at 2 h.The coating thickness all remains at 24–26 μm for Cr:Alratio below 85:15. In the highest chromium containingpowder pack where Cr:Al is 90:10 (sample 10), thethickness of coatings reduces sharply to 18 μm as shown inFig. 7(d). The coatings in Fig. 7(d) also consist of morepores than other specimen in this series of experiments.According to XRD analyses shown in Fig. 8, the surfacecoating composition is FeAl for samples treated using Cr:Alratio below 85:15, while Fe3Al and (Cr, Fe)7C3 appear insample 10 (Cr:Al=90:10).

By observing aluminum concentration profiles in Fig. 9,it is noted that, for samples treated with Cr:Al ratio below85:15, aluminum concentration all attain similar level of 39–42 at.% at the steel surface. Aluminum concentrations thendrop with depth till approximately 25 μm below steel surfacewhich correspond to the thickness of FeAl phase as shown inFigs. 7(a)–7(c) and 8. The similar concentration profile sug-gests that increased aluminum content in powder packs doesnot necessarily increase surface aluminum concentration. Asaturated aluminum level is attained by thermodynamic equi-librium between activator and the steel substrates. Activatorconcentration and diffusion time also modulate the extent of

Fig. 5. XRD spectra of samples 5, 6, 7, and 8 in Table 3 with dif-ferent NH4Cl concentrations (in wt.%).

Fig. 6. Al and Cr concentration profiles of coatings formed in pow-der packs with different NH4Cl concentrations (in wt.%)(samples 5, 6, and 7 in Table 3).

Table 4. Coating thickness and process parameters of samplestreated in powder packs containing different Cr:Al ratio.

Sampleno

holdingtime(h)

NH4Cl(wt.%)

Cr:Al(weight ratio)

Al2O3(wt.%)

Coatingthickness (μm)

9 2 3 75:25 85 26

5 2 3 80:20 85 25

1 2 3 85:15 85 24

10 2 3 90:10 85 18

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aluminum back diffusion in pack cementation.For sample 10 treated using Cr:Al=90:10, surface concen-

tration reaches only 29 at.% corresponding to Fe3Al phase.The entire concentration profile of sample 10 is obviouslylower than those obtained in samples treated using more alu-minum. The lesser amount of aluminum diffusion is due toreduced supply of aluminum in powder packs besides thosereacting with activator. Higher chromium in this sample alsocauses brittle chromium carbide to form on steel surface asshown in Fig. 8. Pores are thus visible within 5 μm belowsurface.

According to the above observations, a critical amount ofaluminum is required to react with the activator in powderpacks. In current study, at least 15% of aluminum in metalpowders is needed to achieve a saturated level of aluminumat the surface. When aluminum is below this level, thethickness and aluminum concentration of Fe–Al layer isreduced. Chromium carbide reaction can then takes placeearlier in the pack cementation process as shown in Fig. 7(d).

4. Conclusions

The Cr/Al dual layer coatings are formed by combininga 550°C and 750°C two-stage pack cementation treatment.The Fe-rich Fe-Al intermetallics, including FeAl and Fe3Al,are first formed on SNCM439 surface and a 6 μm Crcontaining layer is formed on top of the FeAl solid solution.

The low temperature 550°C treatment has a role in mod-ulating the initial aluminum activity in powder packs. TheFeAl intermetallic layer then starts to form during the 750°Ctreatments. The saturated content of surface aluminum iscontrolled to attain 40 at.% which permits the formation offavorable Fe-rich FeAl intermetallics on steel surface.

The coating thickness and pores increases with the secondstage 750°C holding time. When aluminum in-take iscompleted, aluminum starts to deplete in the powder pack.The surface Al concentration then back diffuses into thepowder packs and becomes lower than that in the sub-surface layer. Therefore, there exists an optimum holding

(a) (b)

(c) (d)Fig. 7. Cross section SEM micrographs of samples treated using powder packs of different Cr:Al ratios: (a) 75:25, (b)

80:20, (c) 85:15, and (d) 90:10.

Fig. 8. XRD spectra of coatings treated using different Cr:Al ratios(samples 9, 5, 1, and 10 in Table 4).

Fig. 9. Al and Cr concentration profiles in the coatings usingdifferent Cr:Al ratios as in Table 4.


133 © 2012 ISIJ

time for the 750°C treatment stage.On the effects of NH4Cl activator concentrations, too

much activator can accelerate the powder-gas reactions andcauses metal powder to deplete at earlier time. Backdiffusion is then induced and porous coating structures areformed by Kirkendal effect on the surface due to chromiumand aluminum interdiffusion. Maximum NH4Cl activatorconcentration of 6% should be used.

The increased Cr content in metal powder decreasescoating thickness, since aluminum concentration isrelatively lower and forms a thinner Fe–Al layer. The lesseraluminum content incurs premature Cr deposition to formchromium carbides and inhibit further growth of coatings.Ideal Cr:Al ratio in the powder packs is below 85:15.

AcknowledgementsThe financial support of National Science Council of

Taiwan, R.O.C. through grant #NSC 97-2221-E-027-006project is acknowledged.

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Formation of Al and Cr Dual Coatings by Pack Cementation on