Discontinuous cellular precipitation in a Cr–Mn–N steel with niobium and vanadium additions

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: [email protected] (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,

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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.


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.


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.


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.


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).


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-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)


(supersaturated)

Grain Boundary


(supersaturated)


(supersaturated)

Fig. 13. Proposed origin of chemical potential (Dl) for boundary

migration.

Fig. 14. Austenite/austenite boundary migrating ahead of lamellar

precipitates.


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


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


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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Discontinuous cellular precipitation in a Cr–Mn–N steel with niobium and vanadium additions

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