Formation of Al and Cr Dual Coatings by Pack Cementation

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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 of

SNCM439 steels. The first 550°C treatment assists to modulate adequate aluminum activity for the

formation of iron-rich intermetallics. In the second 750°C treatment stage, simultaneous chromizing and

aluminizing treatments are achieved by first forming a FeAl ferritic layer and then a surface layer with

higher Cr content at later time. The current study examines the effects of second stage 750°C holding

time, activator concentration, and Cr:Al ratio on coating structures. Fe–Al coatings consisting of Fe 3Al andFeAl intermetallic phases are observed to form initially. This Fe–Al layer accounts for 25 to 32 μ m

thickness of the coatings and show good adherence with the substrates. The coating thickness increases

parabolically with 750°C holding time. With prolonged treatment at 750°C, surface concentration of

aluminum in powder packs drops with treatment time and increasing concentration of activator. A peak

concentration exists at a depth below substrate surface. Aluminum is back diffused from the steel surface

into the powder packs. The growth of Fe–Al intermetallics slows down. Surface layer then forms a

thickness of 6 μ m coating with 2–5 wt.% of chromium. Samples treated for longer than 6 h with over 12

wt.% of NH4Cl activator concentration or with Cr:Al ratio higher than 90:10 induce earlier chromium

infusion and lead to porous coating structure due to Kirkendall effect. Eventually chromium carbide forms

to 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 improves

wear resistance. The coated steels can also be employed in

higher temperature and corrosive environment.1,2) Literature

reports that the steel surface after hot dipping

aluminization3) can withstand temperature up to 1093°C

(2000°F).4)

Pack cementation process is an ideal alternative processfor hot dipping aluminization. It is essentially an in situ

chemical vapor deposition (CVD) coating process.5) Four

inter-processes, namely halide activation, gas diffusion,

solid deposition, and solid diffusion reactions, take place in

sequence to form solid coatings by gas-solid reactions.6,7)

The pack cementation process has the advantages of low

cost and applicability to various materials shapes and sizes.

The compositions of coating are dependent on processing

temperature, 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 and

FeAl, are favorable due to their less brittle and more

desirable mechanical properties.15,16) Fe3Al is also reported

to resist sulfidation and oxidation at high temperatures.17)

Minor Cr content in the Fe–Al compounds can further

improve their resistance to room temperature aqueous cor-

rosion and hot corrosion by fused salt deposits.18) To

simultaneously chromize and aluminize steels, the partial pressures of Cr-halide and Al-halide in the powder pack

must 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 the

activity 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 higher

Cr and lower Al concentrations;21,22) or (2) performing dual

instead of single heating processes first by treating at 925°C

and then at 1150°C.20)

Meanwhile, chromium carbide (Cr 23C6) was reported to

form 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-



© 2012 ISIJ 128

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

tions.23) Therefore, it is important to carefully control the

combination 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 to

develop a two-stage treatment at low temperatures (550 and

750°C) to achieve Cr and Al co-deposition by pack

cementation. The lower temperature processes are more

cost-effective from energy point of view. It is also theobjective of current research to control such that the

preferable Fe-rich instead of Al-rich Fe-Al intermetallics are

formed.

2. Experimental Procedures

The substrate used in this study is commercial SNCM439

steel. Its nominal composition is listed in Table 1. The

sample size is 30 mm dia × 5 mm t disk cut from a round

bar. These samples are first ground to 800-grit SiC paper

and then ultrasonically cleaned in methanol for 10 min.

Pack cementation process places the substrates, or SNCM439 steel in current study, within the powder packs.

The powder packs consist of pure metal powders (Al and

Cr) for diffusion, halide activator (NH4Cl), and inactive

ceramic powder (Al2O3) to protect the specimen from being

oxidized. At elevated temperature, halide activator is

decomposed 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 steel

causes metal atoms to diffuse rapidly into the samples by

decomposing metallic halides. An intermetallic compound

layer is thus formed on steel surface. The decomposed halo-

gen gas, e.g . Cl2, then continues to react with metal powders

left in the powder packs till the activity of metals in powder

pack become equilibrium with that at steel surface. There-

fore, halide activator plays an important role to control the

vapor pressures of metal halides in powder packs21) which

maintain the activity of diffusing metals at specimen surface

to sustain the diffusion with aids of gaseous reactions.

The powder packs for current pack cementation process

contain 3–12 wt.% Cr and Al metal powder mixtures, 3–12

wt.% NH4Cl activator, and the rest 85 wt.% Al2O3 filler. The

metal powder is a mixture of 75–90 wt.% Cr and rest 10–

25 wt.% Al powders. SNCM439 steel specimens are

wrapped in powder packs and loaded into a 316L stainless

steel tube. Ar is supplied to flush the reaction chamber toremove moisture and air for 5 min prior the cementation

treatment. The specimens and powder packs are then treated

using a two-stage process, first at 550°C for 2 h before

heating at 10°C/min to 750°C and then hold for 2–8 h.

After treatments, the packs are cooled to room

temperature by air cooling. The coated samples are removed

from the pack and ultrasonically cleaned. XRD (X-ray

diffraction) analyses are performed on steel surface to identify

the intermetallic phases formed using Rigaku D/max-B

diffractometer equipped with Cu target. Cross section of

each sample is cut, mounted, ground and polished for SEM

(scanning electron microscopy, HITACHI S-4700 field

emission SEM) and EDS (energy dispersive spectroscopy,

HORIBA 7200-H) chemical analyses.

3. Results and Discussion

3.1. Effects of Treatment Time

Table 2 lists the compositions of powder packs andtreatment parameters by varying the treatment time at 750°C

(sample 1–4). The total coating thickness increases from 25

to 43 μ m and is proportional to the square root of 750°C

treatment time with a 0.99 coefficient of linear correlation

(R 2). The formations of these coatings are evidently

diffusion controlled. Figure 1 shows that two layers are

formed. The thickness of surface layer is approximately 6

μ m for all conditions listed in Table 2, while the thickness

of sub-surface layer increases with treatment time. In Figs.

1(c) and 1(d), pores in the coating layers are observed to

increase with holding time as well.

The surface coating phase is analyzed by XRD to consistof mainly FeAl solid solution in 2 h-treated sample, Fe3Al

in 4 h-treated samples, Cr 23C6 and Cr 7C3 in 6 h treated

samples, and Cr 7C3 in 8 h treated samples (Fig. 2). These

results are very different from that reported recently24) by a

single step process at 700°C using Cr-Al alloy metal

powders which forms an aluminum-rich brittle Fe2Al5 phase

in contrast to iron-rich Fe3Al phase formed by dual stage

process 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) gaseous

aluminum halides apparently provide fairly high activity of

aluminum at the steel surface to form such aluminum-rich

Fe2Al5 coatings. However, Fe2Al5 coatings are not favored

due to its brittle characteristics. To avoid the formation of

brittle Fe2Al5 phase, activity of aluminum must be reduced.

In current study, the first stage 550°C treatment is designed

to reduce the starting activity of aluminum. Therefore, a

FeAl solid solution is first formed during the second stage

750°C treatment instead of the less favored Al-rich Fe2Al5.

In Fig. 3, Surface aluminum concentration is shown to

attain 42 at.% after 2 h-treatment at 750°C. The concentra-

tion of aluminum drops with diffusion depth till 24 μ m into

the steel. This layer corresponds to ferritic FeAl solid solu-

tion and is consistent with XRD analyses. When 750°C

treatment 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 to

Fe3Al 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 at

750°C stage.

Sampleno.

750°C holdingtime (h)

NH4Cl(wt.%)

Cr:Alin pack

Coating thickness( μ m)

0 0 3 85:15 0

1 2 3 85:15 242 4 3 85:15 32

3 6 3 85:15 38

4 8 3 85:15 43




129 © 2012 ISIJ

It is also important to note that a highest aluminum con-

centration appears at depth of ~6 μ m below surface after 4

h 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 peak

aluminum concentration at a depth below steel surface

requires the aluminum accumulation in steel to diffuse in

two directions, one diffusing further into the steel and the

other diffusing back to surface. Diffusions of aluminum in

two 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 12

and 18 μ m below steel surface according to Fig. 3, respec-

tively. The locations of maximum aluminum concentration

move toward into the steel with time showing some of thealuminum in-take during the earlier treatment time diffuses

further into steels.

On the other hand, the surface aluminum concentration

reduces from 42 to 25, 7, and 5 at.% for 2, 4, 6, and 8 h-

treated samples, respectively. These are compared to the

peak 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 steel

coatings diffuse outward back into the powder packs. This

dual diffusing activity provides an important mechanism

controlling the formation of dual coatings in current study.

Chromium concentrations in this series of experiments all

remain at 2–5 wt.% in the steel which represent slow but

steady diffusion (Fig. 3). It has been reported18) that

chromium activity in powder packs is well below that of

aluminum, even though chromium powder content isdesigned to be greater than that of aluminum in powder

packs. Furthermore, for chromium diffusion to occur

requires higher temperature. Although, the 750°C treatment

allows 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, and

4 in Table 2).



© 2012 ISIJ 130


diffuses at a comparably slower rate than aluminum

diffusing outward to the powder packs. The difference in

diffusion directions and rates of chromium and aluminum

thus causes Kirkendal effects.

Geib and Rapp8) report that, for simultaneous chromizing

and aluminizing, porous coating starts to form due to back

diffusion of chromium and aluminum late in the process

when Cr- and Al-depletion zones form in the powder packs.Because iron has a much higher solubility for aluminum

than chromium, aluminum is the main substitutional atoms

diffusing into the steel in the initial treatment stage. When

aluminum in-take reaches an equilibrium activity on steel

surface with powder packs, aluminum diffusion slows

down. 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 on

steel surface when chromium infuses more pronouncedly

into the steel due to chromium’s higher activity in the pow-

der packs. When chromium diffuses into the steels by

decomposing 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 back

from 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 to

keep the pack cementation reactions in equilibrium with

reduced chromium content in powder packs. Chromium and

aluminum thus demonstrates a positive interaction parame-

ter in both powder packs and steels.

Although both aluminum and chromium of the powder

packs tend to diffuse into the steel substrate at the beginning

of reactions, the sequence of faster aluminum diffusion and

slower chromium diffusion causes dual layers to form in

competition. The competition of chromium with aluminum

diffusion into steel is less pronounced in the beginning of

cementation reactions. But after 4 h of diffusion time at

750°C, aluminum activity attains equilibrium between steel

and powder pack, and chromium becomes the main ele-

ments diffusing into the steel. Further chromium diffusion

drives the earlier diffused aluminum to diffuse back into the

powder pack. Back diffusion is expected, because when

chromium is reduced from the powder pack, the powder

pack then has a high chlorine activity and tends to react with

aluminum. The aluminum cementation reactions essentially proceed 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 the

surface 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 Concentration

Table 3 lists a series of pack cementation samples by

varying NH4Cl activator concentrations (sample 5–8) while

keeping the metal powder content and 750°C processingtime constant for 2 h. The observations of dual layer coating

thickness does not change much with the concentration of

activator for samples treated using 3–9 wt.% of NH4Cl

activator as shown in Fig. 4 and stay at a thickness of 23–

25 μ m. In 12 wt.% NH4Cl treated sample, only a thin layer

of 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:AlAl2O3

(wt.%)Coating thickness

( μ m)

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) NH 4Cl concentration in the powder packs.




131 © 2012 ISIJ

surface coating phases are analyzed using XRD as shown in

Fig. 5. In sample 5 and 6, surface coatings correspond to

FeAl, while in sample 7 and 8, only carbides and Fe or

substrate are present. The presence of Fe phase on surface

analyses indicates that only small amount of aluminum are

diffused into the steel when NH4Cl activator concentration

is greater than 9 wt.%. It is also interesting to note that many

pores 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 aluminum

concentration 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 outward

directions from 7 μ m depth inside the steel. Only 5 at.% alu-

minum concentration is observed at surface. It again shows

that aluminum diffuses outward to the powder packs as sug-

gested in Section 3.1. This is a confirmation of aluminum

depletion in powder packs in using higher activator concen-

tration. On the other hand, chromium continues to diffuse

into the steel substrate at a slower rate. The great amount of pores 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 12

wt.%, only carbide layer is present (Figs 4(d) and 5) on the

steel surface, and no distinguished subsurface Fe–Al solid

solution is formed. Almost none porosity is observed in the

coatings as well as shown in Fig. 4(d). This indicates that

the interdiffusion of chromium and aluminum is limited and

thus no Kirkendall effect occurs. Apparently, aluminum dif-

fusion is restricted when activator concentration is higher

than a critical value.

In this study, when activator concentration reaches 12wt.%, the limited amount of aluminum is only enough to

react with decomposed chlorine from the activator. Not

much aluminum is available in the powder packs for diffu-

sion. Therefore, diffusion of aluminum into steel is limited

and 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 alumiunum

depletion and makes chromium diffusing at earlier time in

comparison with the samples discussed in Section 3.1. Once

the chromium reaches ~4 at.% at steel surface, chromium

carbides are formed. The carbide layer serves as surface bar-rier and forbids further growth of coatings as observed in

Figs. 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 Packs

Table 4 lists a set of samples coated using varied Cr:Al

ratio in the metal powders while NH4Cl concentration is

fixed at 3 wt.% and 750°C treatment time is fixed at 2 h.

The coating thickness all remains at 24–26 μ m for Cr:Al

ratio below 85:15. In the highest chromium containing

powder pack where Cr:Al is 90:10 (sample 10), the

thickness of coatings reduces sharply to 18 μ m as shown in

Fig. 7(d). The coatings in Fig. 7(d) also consist of more

pores than other specimen in this series of experiments.

According to XRD analyses shown in Fig. 8, the surface

coating composition is FeAl for samples treated using Cr:Al

ratio below 85:15, while Fe3Al and (Cr, Fe)7C3 appear in

sample 10 (Cr:Al=90:10).

By observing aluminum concentration profiles in Fig. 9,

it is noted that, for samples treated with Cr:Al ratio below

85:15, aluminum concentration all attain similar level of 39–

42 at.% at the steel surface. Aluminum concentrations then

drop with depth till approximately 25 μ m below steel surface

which correspond to the thickness of FeAl phase as shown in

Figs. 7(a)–7(c) and 8. The similar concentration profile sug-gests that increased aluminum content in powder packs does

not necessarily increase surface aluminum concentration. A

saturated aluminum level is attained by thermodynamic equi-

librium between activator and the steel substrates. Activator

concentration 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 samples

treated in powder packs containing different Cr:Al ratio.

Sampleno

holdingtime(h)

NH4Cl(wt.%)

Cr:Al(weight ratio)

Al2O3

(wt.%)Coating

thickness ( μ m)

9 2 3 75:25 85 265 2 3 80:20 85 25

1 2 3 85:15 85 24

10 2 3 90:10 85 18



© 2012 ISIJ 132


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 obviously

lower than those obtained in samples treated using more alu-

minum. The lesser amount of aluminum diffusion is due to

reduced supply of aluminum in powder packs besides those

reacting with activator. Higher chromium in this sample also

causes brittle chromium carbide to form on steel surface as

shown in Fig. 8. Pores are thus visible within 5 μ m below

surface.

According to the above observations, a critical amount of

aluminum is required to react with the activator in powder

packs. In current study, at least 15% of aluminum in metal

powders is needed to achieve a saturated level of aluminumat the surface. When aluminum is below this level, the

thickness and aluminum concentration of Fe–Al layer is

reduced. Chromium carbide reaction can then takes place

earlier 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 μ m Cr

containing 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. The

FeAl intermetallic layer then starts to form during the 750°C

treatments. The saturated content of surface aluminum is

controlled to attain 40 at.% which permits the formation of

favorable Fe-rich FeAl intermetallics on steel surface.

The coating thickness and pores increases with the second

stage 750°C holding time. When aluminum in-take is

completed, aluminum starts to deplete in the powder pack.

The surface Al concentration then back diffuses into the

powder 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 using

different 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 and

causes metal powder to deplete at earlier time. Back

diffusion is then induced and porous coating structures are

formed by Kirkendal effect on the surface due to chromium

and aluminum interdiffusion. Maximum NH4Cl activator

concentration of 6% should be used.The increased Cr content in metal powder decreases

coating thickness, since aluminum concentration is

relatively lower and forms a thinner Fe–Al layer. The lesser

aluminum content incurs premature Cr deposition to form

chromium carbides and inhibit further growth of coatings.

Ideal Cr:Al ratio in the powder packs is below 85:15.

Acknowledgements

The financial support of National Science Council of

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

project is acknowledged.

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Documents

Formation of Al and Cr Dual Coatings by Pack Cementation