Volume 6 Preprint 99
Multi-Zone Internal Precipitation Reactions
Keywords: Oxidation, carburisation, nitridation, high temperature, alloys, diffusion
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Volume 6 Paper H056
Multi-Zone Internal Precipitation Reactions
School of Materials Science & Engineering, University of New South
Wales, Sydney NSW 2052, Australia, email@example.com
The kinetics and morphological development of internal precipitation
zones are reported for Fe-Ni-Cr alloys subjected to carburisingoxidising and carburising-nitriding gases and for Ni-Cr-Al alloys
exposed to oxidising-nitriding gases, all at 1000oC. In all cases,
successive zones of different precipitates formed in a sequence
reflecting thermodynamic stabilities, providing that the oxidant
forming the less stable precipitate diffused faster than the other
reactant. Thus carbides formed beneath subsurface oxides, as did
nitrides, but nitrides did not form beneath near-surface carbides. The
Ni-Cr-Al alloys formed more complex sequences : Al2O3 + Cr2O3 near
the surface, then Al2O3, then AlN, then AlN + Cr2N and finally AlN,
reflecting the relative stabilities of chromium and aluminium
compounds together with solute thermodynamic interactions. It is
shown that prior formation of a precipitate followed by conversion to a
more stable one can alter both the shape and the volume fraction of
the final reaction product, substantially affecting the severity of the
Keywords: Oxidation, carburisation, nitridation, high temperature,
If an alloy fails to form a stable protective scale, then a gaseous
corrodent species can dissolve in the alloy, diffuse inwards and react
preferentially with solute metals to form internal precipitates. This
process of “internal oxidation” can occur if the oxidant involved has an
activity too low to react with the alloy solvent component. Internal
oxidation kinetics have been successfully modelled [#ref 1,2] for the
simple case of a single corrodent producing small, discrete
precipitates of a very stable reaction product.
If diffusion of the solute metal can be neglected, the rate at which the
precipitation zone deepens is described by:
X 2 = 2k p t
Here, case X is the precipitation depth formed in time, t , Do and No
are the diffusion coefficient and surface concentration (mole fraction)
of oxidant, N m is the initial alloy reactive metal solute content, ε is a
labyrinth factor reflecting diffusional blocking by precipitates, and ν is
the stoichiometric factor for the precipitate MO ν . Internal precipitation
resulting from reaction with a single oxidant has been extensively
studied and the success of the above description well established.
Heat resisting alloys are commonly used in service environments
containing a number of oxidants. For example, combustion of various
fossil fuels produces atmospheres which can be oxidising, carburising,
sulphurising and/or nitriding depending on fuel/air stoichiometry.
Also important is the petro chemical industry in which strongly
carburising and mildly oxidising gases are handled at high
Relatively little information is available for the situation of practical
interest, where internal attack by more than one oxidant occurs. A
number of qualitative observations have been made, but kinetic data is
scant. On the other hand, a theoretical prediction of multi-zone
internal precipitation kinetics was provided some time ago by
Meijering [#ref 3]. Assuming no interaction between the solutes in the
alloy, and no intersolubility between the precipitate compounds, the
theory predicts formation of discrete precipitate zones, each of which
widens parabolically. For the dual oxidant case, the rate of advance of
the inner precipitation front is given by the expression:
ε 1 N1 D 1
ε N D
+ 2 2 2
ν 1 Nm
ν 2 Nm
where the subscripts 1 and 2 refer to the two reactant species. The
separate precipitated zones, are predicted by the theory to occur in
order of thermodynamic stability, with the more stable phase at the
surface, and the less stable at greater depth. The reason for this
ordering is as follows. At the surface the solute metal activity will be
low and the activities of the oxidants high. Consequently the
precipitate formed will be the more stable one. Beneath the surface,
this precipitate will continue to form in preference to the other
possibility for so long as the reactant activity is high enough. If the
second reactant has significant permeability in the depleted subsurface alloy matrix, it will penetrate the first precipitate zone, and
result in the activity of the second reactant exceeding that of the first,
beneath this zone. In this event, the second phase will precipitate
ahead of the first precipitation front.
Equations (2) and (3) refer to the situation where the reacting metal
solute is essentially immobile. If, however, the solute has a significant
mobility, it will diffuse towards the surface and enrich within the
precipitation zone. This effect has been accounted for by Gesmundo
and Niu [#ref 4].
The qualitative success of Meijering’s model has been demonstrated
[#ref 5] using commercial, heat resisting alloys. However, those
experiments involved simultaneous scale growth, and the results are
therefore difficult to interpret. The purpose of the present work was
to investigate internal precipitation by dual oxidants in the absence of
A commercial stainless steel grade 304, and model iron- nickelchromium and nickel- chromium- aluminium alloys, were used.
Composition are shown in Table 1.
Table 1 Alloy compositions (wt.%)
The stainless steel was cut into coupons, which were mechanically
ground to an 800 grit surface finish and subsequently chemically
etched to remove the work hardened surface region. The ternary
alloys were prepared by agon are melting using non-consumable
electrodes. Alloyed buttons were remelted several times and
subsequently annealed at 1000ºC for 120 hours under a flowing
Ar-10%H2 mixture. Samples of these materials were ground to a 1200
grit finish and ultrasonically degreased in acetone immediately prior to
Oxidation-carburization of the commercial materials was carried out
at T = 700ºC in flowing CO/CO2 gas mixture. This mixture was
adjusted so as to yield a supersaturated carbon activity of seven and
oxygen partial pressure of 1 x10–23 atm. Carbon was not deposited on
inert surfaces within the reaction apparatus, demonstrating that the
gases did not equalibrate, but remained supersatured. The oxygen
potential was sufficient to oxidise Cr, but not Fe or Ni. Reaction was
carried out under temperature cycling conditions (45 min at reaction
temperature and 15 min at ambient), in order to promote protective
Oxidation-nitridation of the Ni-Cr-Al alloy, was carried out in still
laboratory air at 1100ºC using temperature cycling conditions.
Carburization-nitridation experiments were carried out at 1000ºC in a
gas mixture of 89% N2, 10% H2, and 1% C3H6. Unit carbon activity was
achieved as demonstrated by carbon deposition within the reactor. It
was necessary also to carry out carburization alone, in a gas of 10% H2
89% Cr, 1% C3H6, and nitridation in 90% N2, 10% H2 followed by
carburization. These reactions were carried out isothermally, as no
surface scale was formed.
Results and Discussion
Selective oxidation of chromium led to the growth on SS304 of an
external scale consisting mainly of Cr2O3 with some spinel. Repeated
thermal cycles led to scale breakage and spallation as shown by
periodic weight losses. Rehealing of the protective scale led to
depletion of chromium from the alloy surface. At 700ºC, metal
diffusion within the alloy was slow, and removal of chromium from the
alloy subsurface region could not be compensated by diffusion from
the alloy interior. Repeated scale damage events led to further
lowering of the chromium concentration until not enough remained to
reform an external Cr203 scale. At that point the alloy was attacked
Metallographic cross sections of reacted SS304 revealed that it had
sustained internal attack. An example is shown in Fig.1. The internal
precipitates were identified as oxides in the subsurface region ( spinel
near the surface and Cr203 at greater depths), and carbides in the
deeper zone. The carburization process was rapid, reaching a depth
after 528 cycles of 520 µm. Similar reaction morpnologies were
observed at 650 and 750ºC [#ref 6].
Oxygen and carbon can both dissolve in SS304 and diffuse within it
much more rapidly than alloy metal components. In the absence of a
protective scale, gas access to the alloy led to inwards diffusion of
Fig. 1. Type 304 stainless steel after cyclic oxidation-carburisation for
oxygen and carbon, both of which reacted with alloy solute chromium,
to form internal precipitates:
2Cr + 3O = Cr2O3
23Cr = 6C = Cr23C6
7Cr23C6 + 27C = 23C7C3
The very rapid penetration rates clearly reflect high permabilities of
the alloy for oxygen and carbon. Unfortunately, the necessary data is
not available with which to test the applicability of equation 3.
The location of the oxide zone, immediately beneath the alloy surface,
with the carbides at greater depth, is as expected from the fact that
Cr2O3 is much more stable than the carbides. It is a consequence of
the conversion of carbide to oxide
2Cr7C3 + 21O = 7Cr2O3 + 6C
which is thermodynamically favoured at high dissolved oxygen
concentrations. These existed near the surface, but not at greater
depths. The very finely dispersed internal oxide reflects the in situ
conversion of prior carbides, which were themselve finely dispersed,
according to reaction (7).
The penetration of carbon past the internal oxidation front can only
occur if DcNc > DoNo . Measurements of these quantities have not
been carried out at such a low temperature. However, extrapolation of
high temperature data for carbon [#ref 9,10] and oxygen [#ref 11,12]
leads to the estimates DcNc ≈ 10-11 cm2 s-1 and DoNo ≈ 10-15 cm2 s-1
ac = 1 and pO2 = 10-23 atm. The development of a deep
carburisation zone is thereby explained.
The almost continuous oxide bands formed near the surface were a
result of the large precipitate volume fractions. Approximating that all
alloy chromium was precipitated, and using mole volumes of 29.1,
45.0, 57.8, and 181.9 cm3 mol-1 for Cr2O3, FeCr2 O 4, Cr7C3 and Cr23C6
respectively, the volume fractions precipitated within SS304 were
calculated and the results are shown in Table 2. With the usual
criterion of fv ≥ 0.3 for formation of a continuous product, it is
concluded that dispersions of carbide precipitates are to be expected,
dispersed Cr2O3 is possible if not all chromium is precipitated, and
that continuous regions of subsurface spinel are explained simply by
the increase in oxide volume accompanying conversion of Cr2O3 to
Table 2 Calculated precipitate volume fractions and
volume change (SS304)
Reaction of SS304 with the CO/CO2 gas mixture, led not only to
internal precipitation but to roughening and pitting of the surface.
This surface disintegration cannot be attributed to the metal dusting
process, [#ref 7,8] which involves deposition of carbon on the surface.
No carbon was ever observed on the reacted metal surfaces. Instead,
the disintegration is attributed to the unusual mode of oxidation and
consequent volume expansion.
Normally, oxidation of SS304 leads to formation of a Cr2O3 external
scale because of the large alloy chromium content. After a sufficient
number of thermal cycles, that process became impossible because
the alloy surface had been depleted of chromium. From that point on,
internal attack by carbon and oxygen occurred. Carbon rapidly
diffused into the steel, precipitating internal carbide. Further
subsequent oxidation of the remaining chromium and the chromium
carbides caused substantial volume expansion.
From the mole volumes of the precipitate phases and assuming still
that all of the chromium in SS304 is precipitated, the expected volume
expansions were calculated, and the results are shown in Table 2.
They show why large volume fractions of internal carbides can
precipitate within austenite. Subsequent oxidation led to insitu
conversion of those precipitates and formation of a very large internal
oxide volume fraction. The calculated volume expansions are large,
implying the development of a high mechanical stress. This stress
would have been accentuated by thermal cycling, leading to
mechanical failure, and the observed surface damage.
Cyclic oxidation kinetics [#ref 13] showed that the alloy lost
substantial weight over the first hundred and fifty cycles, and then
experienced a slow weight increase. Examination of the oxidised
surface confirmed that scale spallation was extensive. X-ray
diffraction analysis of the surface showed that γ -nickel was the major
component along with small amounts of Al203, Cr2 O 3 and NiCr2O4.
Metallographic cross-sections of the reacted alloy revealed large
quantities of internally precipitated corrosion products, as seen in
Figure 2. Large irregular shaped oxide precipitates were located
immediately beneath the alloy surface, and faceted and needle-like
nitride precipitates were located at greater depths. Microprobe
Fig. 2. Ni-Cr-Al alloy after oxidation-nitridation for 260 cycles.
analysis of the precipitation zone showed [#ref 13] that precipitate
phases were arranged in a complex way: Al2O3 and Cr2O3 beneath
the surface, followed by Al2O3 then AlN, then AlN + Cr2N and finally
AlN alone at the reaction front. The metal matrix in the precipitation
zone was single phase γ -nickel and alloy transformation fronts near
the precipitation front are clearly visible in Figure 2. Aluminium and
chromium were both depleted from the γ -matrix, aluminium almost
completely and chromium less severely.
Clearly, scale formation and spallation led to weight losses, which were
eventually offset by the weight gains associated with massive internal
precipitation. The selective formation of Al2O3 and Cr2O3 led to
depletion of these elements from the alloy subsurface zone. This
depletion led eventually to formation of the γ -phase. This phase
possesses a high oxygen permeability [#ref 11], which together with
the loss of protective scale, led to rapid oxygen entry and internal
The distribution of precipitates within the alloy, is generally in
agreement with the predictions of Meijering. After an initial period
during which aluminium and chromium were depleted from the alloy
surface, nickel oxide formed, and controlled the oxygen potential at
the oxide-alloy interface. Frequent damage to this scale would have
rendered it permeable to nitrogen, establishing an effective surface
nitrogen partial pressure of 0.8 atmosphere. Under these conditions,
the alumina and chromium are both much more stable than the
nitrides of either aluminium or chromium, and the oxides formed in
preference to the nitride at the alloys surface. The activity of inwardly
diffusion oxygen was maximum at the surface and decreased with
depth into the alloy. The pattern of oxide precipitation reflects this,
with both Al2O3 and Cr2O3 precipitating together near the surface,
but only Al2O3 at greater depths, where the oxygen activity was low.
The interface between oxide and nitride precipitation zones
corresponds to the reaction
2 AlN + 3O = Al2O3 + 2N
It was calculated [#ref 13] that this equilibrium is sustained by
NN > 3.6 x 10-11, a value sufficient to precipitate AlN, but not the
less stable Cr2N. This explains the appearance of a zone containing
only AlN immediately beneath the oxide precipitation front. At greater
depths, where Ncr was somewhat higher, Cr2N precipitated together
with AlN. An important factor leading to the stabilisation of Cr2N was
the strong thermodynamic interaction between dissolved chromium
and nitrogen [#ref 14] this leads to apparent “uphill” diffusion of
nitrogen from a low concentration in the pure nickel metal at the
surface, to higher concentrations within the alloy, driven by the
chromium gradient. This same phenomenon has been observed [#ref
15] in the reaction between Ni-Cr alloys and nitrogen alone. At still
greater depths, the concentration of inwardly diffusion nitrogen was
further decreased, and Cr2N could no longer precipitate, but AlN was
still stable. Thus, the entire sequence of precipitation zones reflects
local thermodynamic equilibrium with the two inwardly diffusion
oxidants, just as predicted by Meijering.
Internal precipitation kinetics have been measured [#ref 13]. Both
oxide and nitride precipitation zones, were found to widen according
to parabolic kinetics during the first 25h, prior to mechanical damage
to the reaction zone. The principal qualitative prediction of the
Meijering theory is thereby confirmed. In this case, it is expected that
equation (3) should provide an estimate of the rate constant. Using
measured values, [#ref 11,12] and taking the surface oxygen potential
to be defined by the Ni/NiO equilibrium, the oxygen permeability
DoNo in nickel, is estimated at 1 x 10-11 cm2 s-1. A recent estimate
of DNNN [#ref 14] yields a value of 10-10 cm2 s-1 in Ni-20Cr. The
condition for nitrogen penetration is seen to be met for most of the
internal precipitation zone width.
Applying equation (3) to the aluminium precipitation case with the
above permeabilities, a value of kp = 7 x 10-10 cm2 s-1 was
calculated. This is in close agreement with the measured quantity, 6.5
x 10-10 cm2 s-1.
A cross-sectional view of the Fe-Ni-Cr alloy, after exposure to the
carburising-nitriding gas is shown in Figure 3. The precipitation zone
grew according to rapid parabolic kinetics [#ref 16] and consisted of
two zones in which the precipitates where identified by electron
diffraction as chromium-rich M7C3 near the surface, and chromiumrich M23C6 at greater depths. No nitride precipitates where found.
The microstructure and its rate of development were in fact exactly the
same as produced by carburisation alone [#ref 17]. This is to be
expected, given that the chromium rich carbides are more stable than
Cr2N, and that the nitrogen permeability DNNN ≈ 8 x 10-11 cm2 s-1
whereas, DcNc ≈ 9 x 10-10 cm2 s-1 . The nitrogen cannot affect the
carburisation process because it cannot penetrate the carburised zone.
A nitrogen effect on carburisation can be brought about if the alloy is
first nitrided and then carburised. The results of nitriding the alloy for
M 7 C3
Fig. 3. FIB image of sample carbonitrided for 22.5 min.
10 minutes, followed by carburising for 30 minutes, are shown in
Figure 4. Again the reaction products consist of chromium rich M7C3
adjacent to the alloy surface, and chromium rich M23C6 at greater
depths. However, the morphology of the M23C6 precipitates is now
quite different. They are now of lamellar form, orientated
approximately normal to the alloy surface. Furthermore, at the
reaction front, a high angle boundary had developed between the as
yet unreacted austenite, and the chromium-depleted austenite in the
precipitation zone. A transmission electron microscopy investigation,
[#ref 18] established that the M23C6 precipitates had a “cube-oncube” orientation relationship with the austenitic matrix.
Growth of the M23C6 platelets was constrained in direction by the
availability of reactant carbon and chromium. An average growth
direction normal to the sample surface minimised oxidant diffusion
distances to the relatively immobile chromium. This essentially unidirectional growth, coupled with the formation of orientated
precipitate/matrix interfaces, meant that a recrystallisation process in
the austenite at the precipitation front was necessary. The resulting
interface between parent and product austenite grains was in-
coherent, allowing the grain boundary to advance rapidly and to
deliver chromium to the growing carbide platelet tips.
The grain boundary was formed during the initial nitridation stage,
which also developed a platelet precipitate methodology. This
boundary continued to function during subsequent carburisation.
Carburisation alone, did not produce such boundary, reflecting a
different result in the competition between growth of existing
Fig. 4. FIB image of sample nitrided 10 min, then carburised 30 min.
Markers spaced at 30 µm intervals.
precipitates and nucleation of new ones. The greater stability of the
carbide is thought to be an important factor influencing the preference
for new precipitate nucleation over platelet growth. This preference is,
however, overcome at a grain boundary, such as that produced by
Three examples of multi-zone internal precipitation reactions,
involving dual oxidants have been studied. In all cases, the different
precipitated zones were arranged in a sequence corresponding to the
relative thermodynamic stabilities of the reaction products, and the
gradients in oxidant activity. The less stable precipitates were formed
only if the permeability of the less reactive oxidant exceeded that of
the more reactant one.
The separate zones of different precipitates, all grew according to
parabolic kinetics, confirming the qualitative prediction of the
Meijering theory. In the one case, where quantitative comparision was
possible (oxidation-nitridation), the prediction was found to be
Conversion of a less stable precipitate into a more stable compound,
by reaction with a different oxidant, was found to affect the reaction
morphology significantly. Thus, carburisation of a high chromium
content stainless steel, led to internal oxidation of that material,
whereas, oxidation alone would produce external scaling. Prior
nitridation of Ni-Cr-Al alloys, led to subsequent formation of
spheroidal Al2O3 precipitates, rather than the more common acicular
particles. Prior nitridation of Fe-Ni-Cr, led to development of a grain
boundary at the reaction front, which subsequently supported the
growth of a carbide precipitate array.
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