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Acta Materialia 52 (2004) 2407–2417
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Discontinuous cellular precipitation in a Cr–Mn–N steelwith niobium and vanadium additions
R.D. Knutsen *, C.I. Lang, J.A. Basson 1
Department of Mechanical Engineering, Centre for Materials Engineering, University of Cape Town,
Private Bag Rondebosch, 7701 Cape Town, South Africa
Received 6 November 2003; received in revised form 22 January 2004; accepted 26 January 2004
Abstract
The influence of niobium and vanadium additions on the precipitation behaviour in a 24Cr–18Mn–1N austenitic stainless steel is
characterised during ageing at temperatures from 800 to 1100 �C. Niobium demonstrates a tendency to stabilise the cubic MX-type
precipitates, whereas vanadium encourages the formation of hexagonal close-packed M2X-type precipitates. Vanadium, further-
more, promotes formation of M2X-type precipitates by the discontinuous cellular precipitation (DCP) reaction. The presence of
sigma phase and a high frequency of austenite twinning occur in association with the DCP reaction during ageing. This behaviour,
together with the influence of niobium and vanadium, is used to understand the driving force for boundary migration during the
DCP reaction.
� 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Keywords: Stainless steels; Ageing; Interface migration; Discontinuous cellular precipitation
1. Introduction
High nitrogen austenitic stainless steels provide im-
proved mechanical properties and corrosion resistance
compared to conventional austenitic grades [1]. As a
result, these high nitrogen steels have found broader
application in engineering, particularly where advantage
is taken of their excellent combination of high strength
and toughness. Modern developments in steel making
have made possible the manufacture of steels withnitrogen levels up to 1 wt% or even more [1,2]. The
processes employed to achieve these high nitrogen levels
can vary: solid-state nitriding processes are easily
adapted to thin sheet and wire stock, whereas thicker
sections are limited to liquid state metallurgy involving
either high or low (atmospheric) pressure melting prac-
tices. For example, the manufacture of high nitrogen
steels with up to 1 wt% nitrogen is made possible at low
* Corresponding author. Tel.: +27-216503172; fax: +27-216897571.
E-mail address: rknutsen@eng.uct.ac.za (R.D. Knutsen).1 Present address: 180 Degrees Engineering Solutions, Cape Town,
South Africa.
1359-6454/$30.00 � 2004 Acta Materialia Inc. Published by Elsevier Ltd. A
doi:10.1016/j.actamat.2004.01.031
pressure by manipulating the solubility of the nitrogen
in the melt with appropriate alloying. Central to thisapproach is the effective use of manganese in controlling
the melt chemistry [3]. Besides influencing the nitrogen
levels in the melt, the manganese additions reinforce the
high work hardening ability associated with nitrogen in
solution, and consequently promote even greater
strength and toughness. In addition to the structural
advantage provided by good mechanical properties, the
combination of high strength and toughness also pro-motes excellent wear resistance in these steels, which
when coupled with good corrosion resistance, provides
attractive tribological properties [4].
Notwithstanding the good wear resistance provided
by the high nitrogen austenitic steels, attempts have
been made to further improve the wear resistance by
promoting nitride dispersions in the austenite matrix. To
this end the addition of strong nitride formers, niobiumand vanadium, to casting alloys has been considered
with a view to promoting desirable precipitate disper-
sions under controlled ageing conditions. However,
second phase particles can also be quite damaging to
mechanical and corrosion properties [5]. For example,
ll rights reserved.
2408 R.D. Knutsen et al. / Acta Materialia 52 (2004) 2407–2417
chromium nitrides (Cr2N) form discontinuously in high
nitrogen stainless steels on ageing in the temperature
range 700–1000 �C [6–8] and their cellular morphology
significantly reduces toughness and corrosion resistance.
During discontinuous cellular precipitation (DCP),Cr2N lamellae form behind a migrating boundary
leaving precipitate cells consisting of alternate layers of
Cr2N precipitate and less nitrogen-saturated austenite
[6,9]. The Cr2N precipitate is of the M2X-type, where M
is usually Cr, but can be a mixture of other compatible
elements. X usually represents nitrogen, but can also be
a mixture of carbon and nitrogen [10]. The addition of
niobium and vanadium, on the other hand, is expectedto result in the formation of MX precipitates that form
in the temperature range 1000–1200 �C [10–12].
This study reports on the influence of niobium and
vanadium additions in the range 0–1 wt% on the pre-
cipitation reactions in high nitrogen steel with an ap-
proximate base composition of 24Cr–18Mn–1N (all
figures in wt%). The alloy compositions were chosen so
that the individual and combined effects of the elementadditions could be determined.
2. Experimental procedure
2.1. Materials
Eight cast ingots of similar base composition, eachweighing approximately 20 kg, were investigated. The
ingots differed from each other with regard to niobium/
vanadium balance as shown in Table 1. The base com-
position is represented by alloy A, which contains low
levels of niobium and vanadium. Alloys B and C contain
approximately 1 wt% niobium and vanadium, respec-
tively, whereas alloys D–H are composed to consider the
influence of different Nb/V ratios on the microstructuralevolution.
2.2. Heat treatments and metallography
In order systematically to study the precipitation
behaviour of the eight alloys, the as cast alloys were first
solution treated at 1300 �C for 2 h and later aged at
Table 1
Composition of Alloys A–H
Element wt% (at%) A B C D
Fe Balance Balance Balance Ba
Cr 24.1 24.4 23.5 23
Mn 17.5 17.5 18.3 18
N 1.03 (3.89) 1.02 (3.87) 1.17 (4.39) 0.
C 0.08 (0.35) 0.11 (0.49) 0.13 (0.57) 0.
Nb 0.04 (0.02) 1.17 (0.67) 0.04 (0.02) 0.
V 0.20 (0.21) 0.11 (0.11) 1.18 (1.21) 0.
All figures in weight percent except where atomic percent appears in par
temperatures ranging from 800 to 1100 �C for periods of
0.5, 2, 5, 10 and 100 h. All heat treatments were per-
formed under argon partial pressure conditions (102
Torr) and specimens were water quenched after each
heat treatment, including the initial solution treatment.Standard metallographic techniques were used to pre-
pare specimens for microscopic investigation. The final
preparation stages for light microscopy included electro-
polishing in a solution containing 25 g chromium
trioxide, 133 ml acetic acid, and 7 ml distilled water,
followed by electro-etching in the same solution. Both
electrolytic procedures were performed at room tem-
perature; electro-polishing at 20 V and electro-etching at10 V. The electro-etching stage was omitted in situations
where orientation images were generated using electron
backscattered diffraction in the scanning electron mi-
croscope (SEM). The twin-jet polishing technique was
utilised to prepare specimens for transmission electron
microscope (TEM) studies. Suitable foils were prepared
using a solution containing 10% perchloric acid and 20%
glycerol in ethanol. The foils were electro-thinned byapplying a potential difference of 15 V at temperatures
ranging from )10 to 0 �C. Semi-quantitative composi-
tional analysis of the precipitate phases was performed
using energy dispersive X-ray spectroscopy in the SEM.
In most cases, elemental compositions were deduced
from high resolution X-ray maps. The extent of DCP is
reported by reference to the volume fraction of the
precipitate cells as measured by areal analysis performedon a series of light micrographs.
3. Results
3.1. Solution treated condition
The group of alloys showed varying levels of minorphase content in an austenite grain structure after the
standard solution treatment at 1300 �C for 2 h. These
secondary phases, also referred to as particles, were
generally absent in the very low Nb-containing alloys (A
and C), but increased in volume fraction with increasing
Nb/V ratio in the remaining alloys. Although a rea-
sonable level of homogenisation was expected after the
E F G H
lance Balance Balance Balance Balance
.1 23.9 24.7 23.0 23.7
.8 18.8 18.3 18.4 18.4
98 (3.70) 1.02 (3.86) 1.15 (4.32) 1.13 (4.26) 1.13 (4.26)
23 (1.01) 0.11 (0.49) 0.14 (0.61) 0.12 (0.53) 0.14 (0.62)
92 (0.52) 0.62 (0.35) 0.29 (0.16) 0.29 (0.16) 0.40 (0.23)
94 (0.97) 0.45 (0.47) 0.75 (0.78) 0.47 (0.49) 0.22 (0.23)
entheses.
Fig. 1. Occurrence of MX-type particles in alloy E after solution
treatment at 1300 �C.
Fig. 3. [0 1 1] Zone axis diffraction pattern for austenite matrix. Extra
spots correspond to the [1)2 1 0] zone axis for the HCP system.
R.D. Knutsen et al. / Acta Materialia 52 (2004) 2407–2417 2409
solution heat treatment, the distribution of the particles
highlighted the interdendritic regions of the as caststructure, as shown by the distribution pattern in alloy E
in Fig. 1. Compositional analysis has demonstrated that
these particles are rich in niobium and vanadium when
both elements are present in the alloy, but only niobium
is present in the particles in the very low V-containing
alloy (alloy B). Electron diffraction analysis in the TEM
shows that the particles are face-centred cubic with lat-
tice parameters in the range 4.24–4.38 �A, which is con-sistent with the mixed (Nb,V)N phase. Accordingly, the
interdendritic phases are identified as MX-type particles,
where the M component can be a combination of nio-
bium and vanadium, with increasing dominance in
niobium content as the niobium level in the alloy
increases.
3.2. Ageing treatments at 1100 �C
With the exception of the base alloy (alloy A), ageing
at 1100 �C resulted in either the formation of lamellar
precipitates or fine globular precipitates, depending on
the Nb/V ratio in the steel. No additional second phases
were observed in alloy A after ageing at 1100 �C. Fig. 2shows the lamellar-type precipitates that were observed
in alloys C, D, F and G after ageing for 2 h. Theseprecipitates formed by DCP and indications of the mi-
grating front normally associated with DCP are marked
by arrows in both micrographs. The volume fraction of
Fig. 2. Lamellar precipitates after ageing at 110
lamellar precipitates is greatest in alloy C, where the
total cell volume per grain ranges from 0.2 to 0.3. The
cell volume fraction for alloys D, F and G is estimated
at 0.08, 0.15 and 0.05, respectively. Examination of the
same precipitate distributions in the TEM reveals fur-
ther distinction in that the precipitate lamellae are often
fragmented in alloys F and G compared to the more
continuous nature of the individual lamellae in alloys Cand D. Electron diffraction analysis has shown that the
lamellae are hexagonal close-packed and that they de-
velop a close packed orientation relationship with the
host austenite phase such that (1 1 1)c//(0 0 0 1)M2X and
[)1 1 0]c//[1)2 1 0]M2X. Fig. 3 shows the [0 1 1] diffrac-
tion pattern for precipitate lamellae formed in alloy C.
Although the diffraction analysis is consistent with the
identification of Cr2N type precipitates, compositionalanalysis by EDS has shown that the lamellae are rich in
both chromium and vanadium. Furthermore, the lattice
parameters for the lamellar precipitates vary according
to a ¼ 2:57–2.84 �A and c ¼ 4:57–4.61 �A. Given this
situation, it is reasonable to conclude that the lamellae
are of the mixed M2X-type precipitate of the form
(Cr,V)2N. After extended ageing (up to 100 h) at 1100
�C, the DCP process ceases, and instead the same phaseforms continuously as fine needle-like precipitates
(Fig. 4).
Discontinuous cellular precipitation is notably absent
in alloys B, E and H, but instead fine globular precipi-
tates form at grain boundaries and within the large
austenite grains. This precipitate form is depicted in
a TEM micrograph in Fig. 5 after ageing alloy E at
1100 �C for 2 h. The [1 1 2] zone axis pattern once again
0 �C for 2 h, (a) alloy C and (b) alloy F.
Fig. 6. Lamellar precipitation distributions after ageing
Fig. 7. Extended ageing for 100 h at 1000 �C, (a) continuous precip
Fig. 5. Globular MX-type precipitates in alloy E after ageing at 1100
�C. Inset: [1 1 2] zone axis pattern for precipitate.
Fig. 4. Continuous precipitation in alloy C after extended ageing for
100 h at 1100 �C.
2410 R.D. Knutsen et al. / Acta Materialia 52 (2004) 2407–2417
confirms the MX-type precipitate and the calculated
lattice parameter (4.2–4.4 �A) is consistent with the pre-
cipitate form that was already present after the solution
treatment process. DCP did not occur in alloys B, E and
H even after extended ageing up to 100 h at 1100 �C.
3.3. Ageing treatments at 1000 �C
Discontinuous cellular precipitation occurred in all
but alloy B after ageing at 1000 �C for up to 2 h. The
niobium and vanadium levels influenced the extent of
lamellar precipitation and the greatest volume fraction
of precipitate cells was observed in alloy C, which hasmaximum vanadium level and lowest Nb/V ratio. The
cell volume fraction per grain is nearly 1.0. Alloy A
had the next greatest lamellar precipitate volume (cell
volume¼ 0.15), with the remaining alloys D–H having
much lower precipitate volumes (0.05–0.10). The con-
trast in lamellar precipitate distribution is exhibited in
Fig. 6 where alloys C and F are compared after ageing
at 1000 �C for 2 h. The entire austenite grain structureis transformed to lamellar precipitate cells in alloy C
(Fig. 6(a)). After extended ageing at 1000 �C for up to
100 h, needle-like precipitates formed in the remaining
untransformed austenite matrix in alloys D–H, as
shown in Fig. 7(a). The DCP reaction ceased in alloys
at 1000 �C for 2 h, (a) alloy C and (b) alloy F.
itation in alloy F and (b) discontinuous coarsening in alloy C.
Fig. 8. Parallel twin boundaries in alloy A after ageing at 1000 �C for
10 h.
Fig. 9. Globular MX-type precipitation in alloy B after ageing at
1000 �C for 2 h.
R.D. Knutsen et al. / Acta Materialia 52 (2004) 2407–2417 2411
D, F and G after 5 h, but continued for longer periods
in alloys E and H, although it is obvious that extensive
plate-like precipitates occurred during ageing up to 100
h. Diffraction analysis of the needle-like precipitates
yielded the same result as the lamellar phase. As an-
ticipated, both precipitate forms were found to contain
chromium and vanadium. On the other hand, needle-
like precipitates did not form in alloys A and C afterprolonged ageing, but instead the DCP reaction was
regenerated by the formation of new cells on old
cell boundaries, as shown in Fig. 7(b). The new pre-
cipitate cells have a much greater interlamellar spacing
than the �old� precipitate cells and the reaction has
been described as discontinuous coarsening [13]. A high
Fig. 10. Distribution of DCP volumes after ageing a
incidence of austenite twinning was also detected in
the DCP cells in all the alloys and an example after
ageing at 1000 �C for 10 h is shown for alloy A in
Fig. 8.
The high Nb-containing alloy B appeared much thesame after ageing at 1000 �C compared to the ageing
treatments at 1100 �C, even for periods of up to 100 h at
temperature. DCP was notably absent and instead TEM
studies identified prolific globular precipitate popula-
tions on the austenite grain boundaries. An example of
this precipitate form after ageing for 2 h is exhibited in
Fig. 9. Diffraction and compositional analysis indicates
that these precipitates are the same MX-type precipi-tates that were identified at 1100 �C, although not sur-
prisingly, they are much smaller in size than those
depicted in Fig. 5.
3.4. Ageing treatments at 800 �C
Ageing treatments at 800 �C resulted in the formation
of lamellar precipitates in all of the alloys including alloyB. The volume fraction of lamellar precipitation after
ageing for 2 h was greatest in alloy C, followed by alloy
A and then the alloys with significant niobium addi-
tions. The volume fraction of lamellar precipitation in
alloy B was the lowest after the same ageing time.
Fig. 10(a) and (b) show the lamellar precipitation in
alloys A and E, respectively. Alloy E is representative of
the microstructures that evolved in Nb-containing al-loys. Of particular importance is the observation that
the interlamellar spacing of the precipitates within the
precipitate cells is much finer than those that formed
after ageing at 1000 �C and that extensive austenite
twinning continued to occur in the cell structures. Al-
though twinning was also observed in the DCP cell
structures at the higher temperatures (Fig. 8), the fre-
quency of twinning is much greater at 800 �C. Diffrac-tion analysis in the TEM once again showed that these
are M2X-type precipitates with the same orientation
relationship with the austenite as those that form at
1000 �C. The host austenite in alloys A and C were
t 800 �C for 2 h, (a) alloy A and (b) alloy E.
Fig. 11. Occurrence of r-phase at DCP cell boundaries (alloy A aged
at 800 �C).
2412 R.D. Knutsen et al. / Acta Materialia 52 (2004) 2407–2417
completely transformed to DCP cells after ageing at
800 �C for 100 h.Sigma phase (r-FeCr) also appears after ageing at
800 �C and is shown in bright contrast between the
DCP cells in alloy A in Fig. 11. The body-centred
tetragonal r-phase is known to occur in the Fe–Cr
system at temperatures below 820 �C, although the
reaction is quite sluggish [14]. Nevertheless, r-phasecan dissolve large amounts of other elements, including
vanadium and manganese [14–16], and it has beenshown that manganese in particular increases the sta-
bility of the r-phase [15,16]. The high manganese
levels, as well as the additions of vanadium, have
clearly assisted in the formation of r-phase in these
alloys.
4. Discussion
The precipitation behaviour in the present range of
alloys during ageing between 800 and 1100 �C has been
shown to be quite variable. Two predominant precipi-
tate types occur, namely MX and M2X. The substitu-
tional element composition of these precipitates is
complex, although niobium dominates the MX-type,
whereas chromium dominates the M2X-type. Notwith-standing the variation in total alloy concentration over
the range of alloys, the influence of niobium and vana-
dium on the morphology and mechanisms of precipita-
tion over this temperature range is significant and
distinctions can be made even in the solution treated
cast structures. Although it is unfortunate that the ele-
ment additions, excluding niobium and vanadium, are
not highly consistent, there is sufficient excess interstitialelement to encourage the formation of the respective
precipitates, and thus greater influence on the progress
of precipitation is partitioned to the substitutional nio-
bium and vanadium additions. The sequence of precip-
itation events will be explained in terms of the influence
of niobium and vanadium on precipitate stability as well
as the influence that these elements have on controlling
the progress of DCP.
4.1. Influence of niobium and vanadium during solution
treatment
The role of niobium in promoting nitride formation
in the interdendritic regions in the cast structure is es-tablished by the abundance of MX-type particles that
are still visible after the solution treatment at 1300 �C,which is a mere 50–70 �C below the solidus temperature
in these alloys. Although vanadium also occurs in solid
solution in these particles, there is no evidence to suggest
that vanadium promotes their formation. The formation
of large, blocky particles of the MX-type and face-cen-
tred cubic structure has been reported in mild steels [17],high strength low alloy steels [18], stainless steels [19] as
well as high nitrogen steels [20]. Further, it has been
shown that niobium additions strongly favour this re-
action, especially in the presence of nitrogen [17,21,22].
Clearly the level of 1 wt% niobium is more than suffi-
cient to promote the strong presence of Nb-rich nitrides
in the interdendritic liquid during solidification, whereas
these particles are absent in the two low Nb-containingalloys (alloys A and C). The formation of the stable
MX-type particles during solidification will reduce the
nitrogen level in solution at 1300 �C, although the im-
pact is quite small. Despite the fact that the volume
fractions of the MX-type particles have not been
quantified, even complete consumption of the 1 wt% Nb
in alloy B in the form of NbN particles will still leave
0.89 wt% nitrogen in solution. Since it is also likely thatsome niobium will still remain in solution, the level of
nitrogen in solution in the austenite phase after the
solution treatment at 1300 �C will be above at least
0.9 wt%.
4.2. Discontinuous cellular precipitation in the base alloy
In an attempt to explain the respective roles of theniobium and vanadium additions in affecting precipita-
tion behaviour at the different ageing temperatures, it is
appropriate to first discuss the precipitation behaviour
in the base alloy (alloy A). No precipitation was ob-
served at 1100 �C; precipitation only occurred in earnest
during the 1000 �C ageing treatment. At this tempera-
ture profuse DCP occurred to produce Cr-rich nitride
(M2X) lamellae behind migrating austenite grainboundaries. Similar precipitate reactions have been re-
ported by Vanderschaeve et al. [7] and Kikuchi et al. [6]
in high nitrogen stainless steels. The DCP reaction
continued until the entire austenite matrix was trans-
formed to the cellular structure, and upon continued
ageing after the saturation point, the DCP reaction was
regenerated and discontinuous coarsening occurred. At
800 �C, the same DCP reaction occurs, although theinterlamellar spacing is much finer. In addition to the
precipitation of MX2 lamellae at 800 �C, r-phase also
formed at the DCP cell boundaries. The formation of
Fig. 12. Parent austenite orientation reversal at grain boundary due to
DCP, (a) cells migrating either side of austenite grain boundary and
(b) grey-scale austenite orientation map of same area as (a).
R.D. Knutsen et al. / Acta Materialia 52 (2004) 2407–2417 2413
r-phase is not unexpected and has been predicted by
thermodynamic analysis to occur in these steels [23].
However, the distribution of r-phase in relation to the
DCP cells is interesting (Fig. 11) and suggests that a
point is reached during the migration of the cellboundaries when the chromium content becomes suffi-
ciently enriched at the migrating interface to allow
r-phase to form. After a certain volume of r-phase hasformed, the DCP process is again initiated and the
boundary migration continues. Therefore, the sequence
of events portrayed by the morphological arrangement
of the DCP cells and r-phase indicates that chromium
becomes enriched at the migrating grain boundary, in-stead of being depleted if one places emphasis on the
growth of the Cr-rich lamellae. This deduction is at odds
with the conclusions reached by Vanderschaeve et al. [7]
who argue that the DCP reaction is prematurely termi-
nated in their steels due to the development of chro-
mium impoverished zones ahead of the migrating
interface. Their support for the incomplete reaction is
based on the fact that the untransformed austeniteahead of the migrating DCP cell still contains sufficient
nitrogen (about 75% of the original level) to cause su-
persaturation in the host, and therefore drive the for-
mation of Cr2N. Although they have elegantly
illustrated that the DCP reaction in high nitrogen au-
stenitic stainless steels does not occur under steady state
conditions, they have based their argument on the
comparison of the measured growth rate of the precip-itate cells with the theoretical growth rate of lamellar
structures predicted by grain boundary diffusion and
volume diffusion of chromium, respectively. They con-
cluded that, whilst the chemical driving force for the
lamellar precipitation is governed by the level of nitro-
gen supersaturation, the rate of cell advancement is
controlled by the supply of chromium to the growing
lamellae. In the early stages of the DCP reaction theyreport that the growth rate is consistent with grain
boundary diffusion of chromium, but that the growth
rate steadily retards thereafter towards the slower bulk
diffusion rate and then prematurely terminates as men-
tioned above. In our deliberations, we wish to place
different emphasis on the issues affecting the non-steady
state growth conditions of the DCP cells, particularly
with regard to the migration of the austenite/austenitegrain boundary.
If the formation of the lamellar precipitates is ignored
for the moment, then it is possible to develop an argu-
ment for the migration of the austenite/austenite grain
boundaries based on the existence of a pulling force for
the migrating boundary. To establish this, it is first
necessary to consider the widely accepted conditions for
normal grain growth in a single phase polycrystallinesolid. In the case of two adjacent grains separated by a
curved boundary, the effect of the pressure difference
caused by the curved boundary is to create a difference
in free energy (DG) that drives the atoms across the
boundary, and the boundary migrates towards the
centre of curvature. Consequently, the free energy dif-
ference is thought of as a force pulling the grain
boundary towards the grain with the higher energy. Thispulling force per unit area of boundary is given by
F ¼ DG=Vm; ð1Þwhere Vm is the molar volume. The above situation not
only applies to grain growth, but also to recrystallisation
for example, where the boundaries between the new
strain-free grains and the original deformed grains are
acted on by a force DG=Vm, where in this case, DG is due
to the difference in dislocation strain energy between the
two grains. In the present case, it is proposed that the
energy difference is provided by the difference in nitro-gen saturation across the moving interface, and hence
DG can be referred to as the chemical potential differ-
ence, Dl. Accordingly, the transformed austenite grain
will migrate into the untransformed, and still supersat-
urated, neighbouring grain. However, just as the mi-
gration of a recrystallisation front is made possible by
the annihilation of dislocations as the boundary sweeps
into the deformed grain, so is the present situation madepossible by the consumption of nitrogen by the ad-
vancement of the lamellae behind the migrating grain
boundary. The similarity to the process of recrystalli-
sation is further suggested by grain orientation mea-
surements depicted in Fig. 12. In this case (alloy A),
precipitate cells have formed at either side of the original
grain boundary separating the austenite grains 1 and 2
(Fig. 12(a)). The grey-scale orientation image for aus-tenite in Fig. 12(b) demonstrates that the cell volume (1)
migrating into grain 2 has the same orientation as grain
1, whereas the cell volume (2) migrating into grain 1 has
the same orientation as grain 2. The same orientation
reversal can arise during recrystallisation when large
subgrains on one side of a grain boundary can grow into
the adjacent grain and vice versa, and the process is
often referred to as strain induced grain boundary
Boundary
migrationdirectionTransformed
austeniteregion
(low nitrogen)
Grain Boundary
untransformedaustenite
(supersaturated)
untransformedaustenite
(supersaturated)
Grain Boundary
untransformedaustenite
(supersaturated)
untransformedaustenite
(supersaturated)
Fig. 13. Proposed origin of chemical potential (Dl) for boundary
migration.
Fig. 14. Austenite/austenite boundary migrating ahead of lamellar
precipitates.
2414 R.D. Knutsen et al. / Acta Materialia 52 (2004) 2407–2417
migration. In this way it is proposed that the boundary
only indirectly migrates to seek solute for the precipitate
growth, and that the principal force behind boundary
migration is the difference in chemical potential associ-
ated with the transformed and untransformed austenite.In order for the process of DCP to progress at some
given rate, the mobility of the grain boundaries must be
considered as well as the driving force. According to the
discussion in the preceding paragraph, the rate of DCP
is influenced by the rate of austenite/austenite boundary
migration and hence it is common to write the progress,
or velocity, as
V ¼ MDl=Vm; ð2Þwhere M is the mobility of the migrating grain bound-
ary, and Vm in this case is the molar volume of thetransformed austenite phase. The mobility of the grain
boundary will be determined by the mechanism of
boundary migration and whether the migration is dif-
fusion or interface controlled, as well as the influence of
solute drag that is expected to occur in such a complex
alloy.
Returning to the precipitation of the M2X-type la-
mellae, it is widely observed that the precursor to theDCP reaction is that precipitates first form at a grain
boundary and develop an orientation relationship with
one of the grains (the host grain) to reduce the interfa-
cial energy [9]. After this initial stage, the host grain
migrates into the neighbouring grain and the lamellar
precipitates advance in the direction of the migrating
grain boundary. But, if the grain boundary did not
migrate, the process of precipitation would still continueand would occur by growth of the lamellae into the in-
terior of the host grain to maintain the favourable ori-
entation relationship; in other words, opposite to the
direction in which the grain boundary would normally
migrate. The fact that the grain boundary does migrate
provides a much faster supply of solute for precipitate
growth and, consequently, the precipitate growth fol-
lows closely behind the migrating grain boundary. Thisis the preferred growth direction for the lamellae since it
may be argued that the solute is provided by grain
boundary diffusion along the advancing grain boundary
rather than by volume diffusion. In fact, the precipita-
tion of r-phase at the cell boundaries suggests that thereis an oversupply of at least chromium at the migrating
grain boundary.
To summarise, the fact that the DCP reaction occursat all is a result of the migration of the austenite/aus-
tenite grain boundary and not that the boundary mi-
grates due to the DCP reaction. The initial driving force
for the grain boundary migration is provided by the
chemical potential difference (Dl) between the nitrogen
depleted volume around the newly formed precipitates
and the nitrogen supersaturated volume in the adjacent
grain, which is shown schematically in Fig. 13. As
mentioned previously, this is analogous to the process of
recrystallisation where the high angle boundary migratesin the direction of higher stored energy. In the case of
recrystallisation, the dislocations are annihilated at the
sweeping grain boundary and consequently the energy
difference is maintained across the boundary. In the
present situation, the advancement of the grain bound-
ary into the supersaturated grain collects nitrogen at the
boundary, which is then consumed by the growing M2X
lamellae, and thus Dl is more or less maintained be-tween the transformed austenite volume and the un-
transformed austenite grain. It is understandable that
the M2X lamellae will follow close behind the migrating
boundary and the lamellae are often observed to be at-
tached to the boundary. However, this is not necessarily
always the case and an example is presented in Fig. 14
where the migrating boundary has moved ahead of the
precipitate lamellae. In the same way that a boundarywould not be able to migrate during recrystallisation if
the dislocations did not dissolve at the migrating
boundary, it is imperative that the solute is removed by
the advancing lamellae. Furthermore, it follows that the
R.D. Knutsen et al. / Acta Materialia 52 (2004) 2407–2417 2415
velocity of the migrating boundary will retard as Dl is
reduced by the overall diminishing supersaturation of
nitrogen in the untransformed austenite due to the
combined effects of grain boundary and volume diffu-
sion as the reaction progresses.It must be recognised that DCP does not occur for
every diffusion controlled precipitation reaction in-
volving a supersaturated matrix and hence the mobility
of the migrating grain boundary is critical. Some in-
dication of the mechanism of boundary migration is
provided by the incidence of twinning that is associ-
ated with the DCP cells [24]. Usually, extensive twin-
ning of the nature exhibited in Fig. 8 is associated withthe formation of annealing twins during recrystallisa-
tion, but in this case the requisite prior deformation is
absent. Nevertheless, it has been demonstrated that the
special boundary migration conditions that arise dur-
ing DCP support the development of twins in the same
way that they arise during recrystallisation [24]. In
essence, the orientation of the {1 1 1} austenite planes
in the precipitate cells and the bowing out of theboundary ahead of the lamellae provides the necessary
conditions for Shockley partial dislocations to be
emitted from the advancing boundary and thereby
form the incoherent and coherent twin boundaries,
where the latter is parallel to the long axis of the la-
mellae. This argument is based on the pop-out model
of Meyers and Murr [25], who suggest that the initi-
ation of twins takes place at grain boundary ledges,and is supported by the more recent analysis by
Mahajan et al. [26] who show that the partial dislo-
cations are generated by growth accidents as the ledges
migrate. Consequently, it is proposed that the austen-
ite/austenite boundary must migrate by the ledge
mechanism in order to satisfy the requirements for
twin formation. The occurrence of twin boundaries in
the precipitate cells depicted in the orientation imagein Fig. 12(b) supports this proposal. The proposal is
further reinforced by the fact that the frequency of
twinning increases with decreasing temperature, which
Table 2
Summary of precipitation after ageing between 800 and 1100 �C
Alloy Solution treated Ageing temperature (�C)
800 1000
A No MX M2X (DCP)+Sigma M2X) (DCP)
coarsening
B MX (interdendritic) M2X (DCP) MX (Globular
C No MX M2X (DCP)+Sigma M2X) (DCP)
coarsening
D MX (interdendritic) M2X (DCP) M2X) (DCP)
E MX (interdendritic) M2X (DCP) M2X) (DCP)
F MX (interdendritic) M2X (DCP) M2X) (DCP)
G MX (interdendritic) M2X (DCP) M2X) (DCP)
H MX (interdendritic) M2X (DCP) M2X) (DCP)
Reference to the solution treated condition is included for comparison.
is what one would expect given that the number of
growth accidents would increase with increasing Dl at
the lower temperature. The fact that the precipitate
lamellae develop an orientation relationship with the
host austenite by aligning parallel to the {1 1 1} planesallows these planes to protrude into the boundary to
create the steps for further boundary migration, and at
the same time provide the boundary migration direc-
tion that is favourable for the advancing lamellae to
follow.
4.3. Influence of niobium and vanadium during ageing
The addition of vanadium expands the stability of the
M2X phase field up to 1100 �C and also increases the
M2X precipitation kinetics at 1000 and 800 �C, which is
illustrated by the ageing results of alloy C in comparison
to the base alloy. The M2X reaction in alloy B is sup-
pressed to temperatures below 1000 �C and additional
MX-type precipitation occurs on ageing at 1000 and
1100 �C. Furthermore, the kinetics of the DCP in allalloys containing niobium additions is significantly
slower when compared to the base alloy, alloy A. Dis-
continuous coarsening occurs in alloy C in the same way
that it is evident in alloy A, but it does not occur in
alloys D–H, which all contain combinations of niobium
and vanadium. Instead, after extended ageing at 1000 �Cand in some cases at 1100 �C, the DCP reaction stops
and the M2X precipitates form continuously as needle-like precipitates (Fig. 7(a)). The results are briefly sum-
marised in Table 2.
The reason that the addition of vanadium promotes
the formation of M2X is understood from the influence
that vanadium has on raising the thermodynamic sta-
bility of M2X in these alloys [23], and so it is not
surprising that M2X should form at higher tempera-
tures in alloys containing significant vanadium levels.What is interesting though, is the change in M2X
precipitation mechanism from DCP to continuous
precipitation in the same alloy. As already discussed,
Nb, V content
(at%)1100
+discontinuous No precipitation Base alloy
) MX (Globular) 0.67 Nb
+discontinuous M2X) (DCP)+ continuous 1.21 V
+ continuous M2X) (DCP)+ continuous Nb=V ¼ 0:54
+ continuous MX (Globular) Nb=V ¼ 0:74
+ continuous M2X)DCP+continuous Nb=V ¼ 0:21
+ continuous M2X) (DCP)+ continuous Nb=V ¼ 0:33
+ continuous MX (Globular) Nb=V ¼ 1:00
2416 R.D. Knutsen et al. / Acta Materialia 52 (2004) 2407–2417
DCP is energetically more favourable on condition that
the austenite/austenite grain boundary can migrate.
The fact that the DCP reaction terminates and yet
abundant continuous precipitation of M2X still occurs
must point towards a dramatic decrease in the velocityof the migrating austenite/austenite grain boundaries.
As explained by Eq. (2), the velocity is affected by the
mobility and the driving force for boundary migration.
The decrease in driving force at the higher ageing
temperatures (1000 and 1100 �C) can be explained by
the greater level of nitrogen solubility, and therefore it
is feasible that boundary migration should slow or stop
before the M2X precipitation is complete. Given thatvolume diffusion is relatively high at these tempera-
tures, the M2X precipitation can still form continu-
ously and the regular pattern of the precipitate arrays
suggest that favourable orientation relationships are
maintained. However, it is not immediately clear as to
why the DCP process should give way to continuous
precipitation in certain alloys at 1000 �C, yet continuesuntil the entire austenite volume is transformed in thecase of alloys A and C. The difference in the case of
alloy C is that it contains the highest vanadium level,
but on the other hand, alloy A has relatively low va-
nadium level compared to the remaining alloys. The
reasons for this behaviour are provided by considering
the possible influence of niobium on not only lowering
the stability of the M2X-type precipitates, but also on
affecting the mobility of the austenite/austenite grainboundaries.
All the alloys containing niobium, with the exception
of alloy B (1 wt% Nb), formM2X at 1000 �C; initially byDCP and later by continuous precipitation. The two
alloys that are transformed entirely by DCP contain
very minor niobium levels. As discussed above, the latter
situation cannot be entirely explained by the role of
vanadium. Consequently, it is suggested that niobiuminfluences the mobility of the austenite/austenite
boundary migration and therefore terminates the DCP
reaction, whereas DCP is able to continue unabated in
the base alloy. Since niobium does not enter into the
M2X phase in the same way that vanadium does, it is
plausible that niobium might collect at the advancing
austenite boundaries. Furthermore, there is a tendency
for solute to segregate to grain boundaries anyway be-cause in most cases it leads to a reduction in grain
boundary energy. As a result, the condensed niobium
atoms migrate along with the boundary and exert a drag
that reduces the boundary mobility. If we now combine
this reduced mobility with lower chemical potential (Dl)at the higher ageing temperatures, then the cell bound-
ary velocity will decrease according to Eq. (2), and cause
the DCP reaction to cease despite the fact that M2X canstill form continuously. This same situation does not
arise in alloys A and C since niobium is not available in
the same way.
5. Conclusions
1. The addition of niobium and vanadium to the high
nitrogen austenitic stainless steel leads to an increased
stability of the MX-type precipitates in the case of theniobium additions, and an increase in the stability of
the M2X-type precipitates in the case of the vana-
dium additions. More particularly, it was noted that:
• Niobium promoted the formation of MX-type
precipitates during solidification and on ageing.
Vanadium does participate in these reactions but
there are no indications that these reactions are
promoted by the vanadium additions.• Vanadium promoted the stability of M2X-type
precipitates up to 1100 �C and notably increased
the kinetics of the precipitation reaction at temper-
atures between 1000 and 800 �C.• Niobium suppressed the stability and kinetics of
the M2X precipitation reaction.
2. The mode of precipitation of the M2X-type phase is
also influenced by the niobium and vanadium levels,whereby vanadium tends to promote DCP, whereas
niobium suppresses DCP. This behaviour is explained
by the solute drag effect that niobium exerts on the
austenite/austenite boundary migration and thereby
retards the DCP reaction.
3. Finally, it is concluded that DCP occurs as a result of
the propensity for austenite/austenite boundary mi-
gration, which in turn is influenced by the chemicalpotential difference provided across the boundary
by the respective nitrogen levels in the transformed
and untransformed austenite.
Acknowledgements
The authors gratefully acknowledge the financial
sponsorship provided by the National Research Foun-
dation (Pretoria, RSA), Columbus Stainless (Pty) Ltd
(Middelburg, RSA) and the University of Cape Town
Research Committee. The authors are also appreciative
of the facilities and assistance provided by the ElectronMicroscope Unit at the University of Cape Town.
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