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