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ASM Handbook, Volume 4A, Steel Heat Treating Fundamentals and Processes
J. Dossett and G.E. Totten, editors
Fundamentals of Nitriding and
Nitrocarburizing
E.J. Mittemeijer, Max Planck Institute for Intelligent Systems (formerly Max Planck Institute for Metals Research) and
Institute for Materials Science, University of Stuttgart
Introduction
The nitriding process, which involves the
introduction of atomic nitrogen (N) into the surface of a component, has been a most versatile
and efficacious method of surface treatment of
(usually) iron-base materials for many decades.
The advent of nitriding as a technical process in
the early 20th century was the work of Adolph
Machlet (American Gas Company) in the
United States (Ref 1) and Adolph Fry (Krupp
Works) in Germany (Ref 2). The nitriding
process has since become a technology of ever
growing importance, finding an even at present
still widening field of applications. In the
course of time, a great number of process variants have been developed. One principal development is the process of nitrocarburizing,
whereby carbon is introduced simultaneously
with nitrogen.
Different methods exist for introducing
atomic nitrogen, or both atomic nitrogen and
atomic carbon, into the surface of steel. As
described in other articles in this Volume, various nitriding/nitrocarburizing atmospheres can
be indicated:
Gas (NH3-H2) mixtures (see “Gas Nitriding
and Gas Nitrocarburizing of Steels”)
Salt (cyanate-cyanide) baths (see “Liquid
Nitriding of Steels”)
Plasmas (ionized gases, as, for example,
based on N2-H2 gas mixtures) (see “Plasma
(Ion) Nitriding and Nitrocarburizing of
Steels”)
Powder media (based on CaCN2, calciumcyanamide, which delivers NH3 upon reaction
with H2O from an added “activator”) have been
successfully applied in the laboratory (Ref 3)
but have not found technological application.
This article provides an overview of the current understanding of essential aspects of the
thermodynamics and kinetics of the nitriding
and nitrocarburizing of iron-base materials.
The focus of this article is on the scientific
background of nitriding/nitrocarburizing with
gaseous processes. The reason is that only the
gaseous treatments allow for precise control of
the thermodynamic conditions, even if only in
the laboratory. Thus, the chemical potentials
of nitrogen and carbon at the surface of the
component to be nitrided/nitrocarburized (the
nitriding/carburizing “power”) can be varied
over a wide range in a controlled manner (i.e.,
these chemical potentials can be assigned specified values), whereas such prescribed variation
of these chemical potentials is not possible in
the case of the other processes mentioned. This
leaves unimpeded that reproducible results by
salt bath and plasma nitriding/nitrocarburizing
processes can be obtained, but the thermodynamic conditions of these processes are illdefined, thereby obstructing development of
corresponding process variants that can be
tuned in the above sense.
After a brief overview of the early history
and the basics of the generated microstructure
and properties, the interpretation of the “normal” iron-nitrogen phase diagram and the
“potential” Lehrer diagram are discussed. Next,
the essential elements for realizing (in practice)
controlled nitriding and controlled nitrocarburizing are treated to some extent. Special attention is paid to the occurrence of local
equilibria and stationary states. The concept of
the “diffusion path” is highlighted in order to
understand the complicated microstructures that
develop in the iron-carbonitride compound
layer upon nitrocarburizing. The very pronounced interaction of carbon and nitrogen
atoms dissolved on the same sublattice of
interstitial sites is demonstrated by recent diffusion experiments. The role of alloying elements, Me, is characterized by employing
the distinction of strong, intermediate, and
weak Me-N interactions, for the evolution
of the microstructure of both the compound
layer and the diffusion zone. The cardinal
role of the so-called “excess nitrogen” for the
nitriding kinetics is emphasized. Possibilities
to describe the time (and temperature)
dependencies of the nitriding/nitrocarburizing
processes are summarized.
1. Advent of Nitriding
On May 25, 1906, a patent application was
filed by A. Machlet that proposed to avoid oxidation of steel components by replacing the air
atmosphere in the retort with ammonia. This
patent was granted on June 24, 1913 (Patent
1,065,697). Apparently, soon after submitting
this patent application in 1906, Machlet discovered that treating such components in an ammonia atmosphere at elevated temperature led to
a “skin, casing, shell or coating” that was
very hard and difficult “to tarnish, corrode, rust,
or to oxidize.” This invention was presented
as a patent application filed on March 19,
1908. The patent was granted (also on June
24, 1913) as Patent 1,065,379. This patent
(Ref 1) represents the invention of
the nitriding process in the United States!
In a further patent application filed on July 12,
1907, Machlet in fact introduced the
(gaseous) nitrocarburizing process, with a
hydrocarbon gas combined with ammonia as
the treatment atmosphere and including
a subsequent treatment in pure ammonia.
This patent was granted on April 14, 1914
(Patent 1,092,925).
In Germany, the nitriding process of steel
for surface hardening was developed by
A. Fry (Ref 2). In particular Fry’s work led to
the technological application of nitriding as
a surface engineering process, especially
by his development of steels (containing aluminum as an alloying element) dedicated to
nitriding.
Starting with these early developments, a
great number of process variants have been
developed with significantly different effects
on the structure and properties of the surface
after nitriding/nitrocarburizing (see, for example, the article “Gas Nitriding and Gas Nitrocarburizing of Steels” in this Volume).
620 / Nitriding and Nitrocarburizing of Steels
2. Nitrided/Nitrocarburized
Microstructure; Thermodynamics,
and Kinetics
be large and their interaction can be strong,
also because of their surrounding strain fields
accommodating the misfit of the sizes of the atom
and the interstitial site.
To understand the nitriding/nitrocarburizing
process, full understanding of its thermodynamics does not suffice at all. The kinetics of the
various process steps (as mass transfer from
The nitriding process typically involves the
introduction of nitrogen into the surfaceadjacent zone of a component, usually at a temperature between 500 and 580 C (773 to
853 K). Depending on the nitriding “power”
of the nitriding atmosphere surrounding the
component, a nitrided zone emerges that, especially in the case of ferritic iron-base alloys or
ferritic steels nitrided at temperatures lower than
590 C (863 K), can be subdivided into (Fig. 1):
A compound layer (of thickness, say, up to
The technological importance of nitriding is
derived from the pronounced increase of the
resistances against fatigue, wear, and corrosion,
which can be achieved by tuned applications of
the nitriding process.
The pronounced property improvement is,
roughly speaking, due to the high hardness, the
internal stresses, and the modified chemistry in
the nitrided zone. It can be crudely said that, in
general, favorable wear and corrosion properties
can be due to specific compound layers and that
favorable fatigue and also wear properties
(if the compound layer has been removed after
nitriding or its formation has been avoided) can
be ascribed to the diffusion zone (Fig. 1).
Nitrocarburizing processes, as compared to
nitriding processes, largely influence the composition and constitution of the compound layer
and thus can be relevant for wear (and corrosion) properties.
To understand what is happening during nitriding/nitrocarburizing and to be able to optimize the
process in view of desired properties, it is imperative to understand the thermodynamics and kinetics of the process. Such basic understanding is less
common than one may expect and, in view of the
experience of approximately one century of nitriding, amazingly has only rather incompletely
been obtained. Hence also to this day, research
on nitriding and nitrocarburizing is timely. From
a scientific point of view, the properties of interstitially dissolved elements as nitrogen and carbon in
iron-base matrices are intriguing because their
mobility on the sublattices of interstitial sites can
Fig. 1
Schematic cross section of nitrided region of an iron-base ferritic specimen/component showing the compound
layer and the diffusion zone with their (possible) constituents
N content, at.%
0
10
20
30
40
1000
900
912°C
(γ-Fe)
Magn.
trans.
800
ε
770°C
Temperature, °C
several 10 mm), largely composed of iron
nitrides, as g0 -Fe4N1x and e-Fe2N1z (see
the iron-nitrogen phase diagram reproduced
in Fig. 2, and see Table 1)
A diffusion zone (of thickness, say, up to
several 100 mm), where, in the case of pure
iron or carbon steel, after nitriding, upon
either slowly cooling or upon aging
subsequent to quenching, the nitrogen dissolved at the nitriding temperature precipitates as iron nitrides in the diffusion zone,
or, in the case of steel containing alloying
elements with affinity for nitrogen, as aluminum and chromium, alloying element
nitrides precipitate during nitriding
700
5.7%,680°C
(α-Fe)
600
0.10
650°C
2.8 4.5 5.7
592°C
2.4
5.7
γ’-Fe4N
500
480°C
508°C
Magnetic
transformation
Fe2N
400
300
0
2
4
6
8
N content, wt%
Fig. 2
Iron-nitrogen phase diagram. Redrawn from Ref 4
10
12
14
Fundamentals of Nitriding and Nitrocarburizing / 621
the nitriding/nitrocarburizing atmosphere to the
substrate, diffusion, and precipitation) are of at
least equal importance, albeit recognizing that
the thermodynamics of the process constrain
the possible kinetics.
Physical materials science has not been able
until now to generally provide a description of
the kinetics of a process once its thermodynamics
have been established; if the Gibbs energies of
the beginning and end stages are known, only seldom it is possible to provide a quantitative
description of the kinetics (i.e., only rarely can
the path taken by the system in the energy landscape of the system be predicted as a function of
the independent, or state, variables) (Ref 14).
This serves to indicate that also in the field of
nitriding/nitrocarburizing, even supposing that
our understanding of thermodynamics is complete
(which is not the case), the complex kinetic phenomena encountered (e.g., in the case of nitrocarburizing; see Section 9.2) have obstructed until
now the presentation of a theoretical treatment
comprehensively integrating the thermodynamics
and kinetics of nitriding/nitrocarburizing.
The previous text should not suggest ignorance of any scientific background, and, indeed,
a lot of research has been done in recent years
that has provided accumulated fundamental
knowledge of specific aspects of nitriding/nitrocarburizing. The following attempts to provide
an overview of the current basis of our understanding, including results of recent research.
References 15 and 16 are well-known, previous review papers but are more limited and of
rather different character and scope.
phases at normal temperatures and pressure
for the iron-nitrogen system. This is the origin
of the well-known porosity in the compound
(nitride) layer (Ref 18–20) that already develops during the nitriding process, because the
nitride not in contact with the outer atmosphere
is simply the nitride at elevated temperature at
approximately 1 atm pressure and thus is prone
to decomposition in Fe and N2 (see Fig. 3; see
also Fig. 8 and 9 and their discussion in
Section 9). Even nitrogen ferrite (a-Fe[N]),
as produced by a nitriding treatment, is super-
Table 1 Crystal structures and composition ranges of Fe-N-C phases
Phase
a-Fe[N,C]
g-Fe[N,C]
a00 -Fe16N2
g0 -Fe4N1x
e-Fe3(N,C)1+y
Y-Fe3C
Crystal structure
Nitrogen content, at.%
Carbon
content, at.%
Lattice parameter
references
Body-centered cubic Fe
N and C disordered in octahedral
interstices
Face-centered cubic Fe
N and C disordered in octahedral
interstices
Body-centered tetragonal Fe sublattice
N ordered in octahedral interstices
<0.4
<0.1
Ref 5–8
<10.3
<9.1
Ref 5–8
...
Ref 10
<0.7
Ref 11
15–33
<8
Ref 12 (only for
e-Fe3N1+y)
0
25
Ref 13
Face-centered cubic Fe sublattice
N ordered in octahedral interstices
Hexagonal close-packed Fe sublattice
N and C more or less ordered in
octahedral interstices
Orthorhombic Fe sublattice
C in bicapped trigonal prisms
12.5
(structural N vacancies can
occur)(Ref 9)
19.4–20
3. The Iron-Nitrogen Phase Diagram
Interpretation and The Development
of Porosity
There is a lot of misunderstanding among heat
treatment practitioners and also in the literature
regarding the interpretation of the stability, metastability, and instability of nitride and carbide
(and carbonitride and nitrocarbide) phases in
iron-base materials. This already is manifested
upon discussing the iron-nitrogen phase diagram.
Phase-stability regions for the iron-nitrogen
system are usually presented in a “normal”
phase diagram: temperature versus nitrogen
content (Fig. 2; for temperatures below
300 C, or 573 K, see Ref 17). “Normal” phase
diagrams, as presented in the compilation of,
for example, Ref 4, present stability regions
for phases as a function of temperature and
composition at 1 atm pressure. If this would
hold for the iron-nitrogen phase diagram presented in Fig. 2, then all phases shown as, for
example, the solid solutions a-Fe[N] (nitrogen
ferrite) and g-Fe[N] (nitrogen austenite) and
the nitrides g0 -Fe4N1x and e-Fe2N1z are
unstable. Indeed, if the kinetics allow (i.e., the
temperature is high enough and the needed time
is of practical length), these phases decompose
in Fe (solid) and N2 (gas), which are the stable
Porosity in the e sublayer, as developed during nitriding of an a-Fe substrate at 570 C (843 K) for 7 h at rN =
1.93 atm1/2, by the association of dissolved nitrogen as N2 gas molecules at grain boundaries in the sublayer
but also within the grains of the sublayer. The porosity is most pronounced in the surface-adjacent part of the sublayer,
because this is the oldest part of the sublayer and this part of the sublayer has the largest dissolved nitrogen content,
implying a larger driving force for N2 gas formation close to the surface than at larger depths. The pressure in the pores
is so large that local distortion at the surface leads to local bulging out of the nitrided material. The developing pores at
the grain boundaries can coalesce in advanced stages of nitriding, leading to “open” grain boundaries in contact with the
outer nitriding atmosphere (see also Fig. 9b and its discussion in Section 9). Light optical micrograph of cross section;
etched in 1 vol% Nital; oblique illumination, green light. Courtesy of M.A.J. Somers
Fig. 3
622 / Nitriding and Nitrocarburizing of Steels
saturated with dissolved nitrogen and can
decompose into practically nitrogen-less iron
and N2, that is, show pore development, after
nitriding for a long time, as shown in dedicated
experiments, although this is not generally
recognized.
Genuine equilibrium for the phases mentioned
at normal temperatures and pressures can only be
realized at the surface of the specimens by contact with a nitriding medium of fixed chemical
potential of nitrogen. Hence, the usual ironnitrogen phase diagram (Fig. 2) represents the
equilibria between an iron-nitrogen phase, or
two or (maximally) three iron-nitrogen phases,
and a medium of largely variable chemical
potential of nitrogen. Such a medium is an
NH3-H2 gas atmosphere of constant composition
that can be varied over a wide range. Evidently, then, whereas at a certain temperature a
single-phase region in the iron-nitrogen phase
diagram can represent equilibria with a range of
nitriding atmospheres of variable value of the
chemical potential of nitrogen, a two-phase
region in the iron-nitrogen phase diagram at a
certain temperature presents the equilibrium
with a nitriding atmosphere of a single, specific
value of the chemical potential of nitrogen. This
knowledge is needed for understanding of the
Lehrer diagram dealt with in the next section.
The discussion in the preceding paragraph
also implies that the whole nitrided zone of a
nitrided component, even if the component
would be homogeneously (through) nitrided,
and apart from (possibly) the very surface of
the component is not in (far from) thermodynamic equilibrium.
Within this context, it should moreover
be realized that the possibility to nitride in
NH3-H2 gas mixtures is a fortunate consequence
of slow ammonia-dissociation kinetics: at the
usual nitriding temperatures and pressures ammonia gas should thermodynamically be practically
fully decomposed in nitrogen gas and hydrogen
gas. However, this is a slow process, and thus,
for example, by maintaining a sufficiently
high flow rate in the nitriding furnace, such dissociation can be more or less avoided (see Sections 5
and 8). Only in this way can nature be prompted to
impose a high chemical potential on the surface of
the component to be nitrided.
The preceding two paragraphs illustrate that
much of materials science involves dealing
with systems in states far from equilibrium.
A discussion similar to that for the ironnitrogen phase diagram can be given for the
iron-carbon phase diagram, as published in compilations as given by Ref 4. Indeed, cementite is
an unstable phase at normal temperatures and
pressures and prone to decomposition in iron
and carbon (graphite), which is indeed happening
(kinetically possible) at elevated temperature
within reasonable times. Upon inspection, there
are similarities between the iron-nitrogen and
iron-carbon phase diagrams. This implies that
graphite and molecular nitrogen gas play similar
roles for the iron-carbon and iron-nitrogen systems, respectively. This has led one German
researcher (Ref 21) to propose the name “Molnit”
(“molnite” in English) for this molecular nitrogen gas as the pendant for “Graphit” (“graphite”).
Then it is obtained:
aN;s ¼
4. Nitriding Potential and the
Lehrer Diagram
1
=2 N2 $ ½N
(Eq 1a)
NH3 $1=2 N2 þ3=2 H2
(Eq 1b)
giving
(Eq 2)
where [N] represents nitrogen dissolved in the
solid substrate M. Establishment of the equilibrium (Eq 2) implies the occurrence of local
equilibrium at (only) the surface of the substrate M (see Section 8).
Equation 1(a) implies that:
=2 mN2 ;g ¼ mN;s
1
(Eq 3)
where mN2 ;g and mN;s represent the chemical
potentials of nitrogen in the gas atmosphere
and the solid substrate, respectively. Assuming
ideal gases or, at least, adopting a constant
fugacity coefficient,* it follows from Eq 3:
=2 m0N2 ;g þ1=2 RT ln
1
pN2
¼ m0N;s þ RT ln aN;s
p0
(Eq 5a)
Using the equilibrium constants, K, of Eq 1(b)
and 2:
The nitriding of a solid M in an NH3-H2 gas
mixture can formally be imagined as the result of
bringing N2 gas into contact with M under a certain pressure. This statement is a consequence of
the Gibbs energy (and thus the chemical potential)
being a state variable. Therefore, the route followed to reach a certain (final) state is irrelevant
for the value of the Gibbs energy of that (final)
state. Thus, nitriding in an NH3-H2 gas mixture
can be conceived as the sum of the following
(hypothetical) reactions (the subsequent treatment
and also that of Section 6, “Carburizing Potential
and Controlled Carburizing,” is largely derived
from the one presented in Ref 20):
NH3 $ ½N þ3=2 H2
rffiffiffiffiffiffiffi
pN2
p0
(Eq 4)
where m0i is the chemical potential of component i (i = N2,g or Ns) in the reference state
(temperature dependent at the selected pressure
of the reference state; see the following), R is
the gas constant, T is the absolute temperature,
pN2 is the partial pressure of the (hypothetical)
nitrogen gas in Eq 1(a) and (b), p0 is the pressure of N2 in the reference state,** and aN;s is
the activity of nitrogen in the substrate. Now
m0N;s is selected such that:
1= m0 ¼ m0
N;s
2 N2 ;g
* In this latter case, the fugacity coefficient is thought to be
incorporated in the reference chemical potential 0N2 ;g (cf. Eq 4).
1=2
Kð1bÞ ¼
3=2
pN2 pH2
pNH3 p0
3=2
Kð2Þ ¼
aN;s pH2
pffiffiffiffiffi
pNH3 p0
it follows:
aN;s ¼ Kð1bÞ pffiffiffiffiffi
pffiffiffiffiffi
p0 rN ¼ Kð2Þ p0 rN
(Eq 5b)
with the so-called nitriding potential (rN)
defined by:
rN pNH3
3=2
(Eq 6)
pH2
where pNH3 and pH2 are the partial pressures of
the gas components NH3 and H2.
The pressure of the reference state for the gas
components is selected as one pressure unit (usually 1 atm), requiring that the partial pressures of
all gas components in the equations must be
expressed in the same unit as the pressure of the
reference state. As a consequence of this step,
the numerical value of the nitrogen activity,
aN;s , can be interpreted as the square root of
the pressure (in pressure units) of the hypothetical nitrogen gas occurring in Eq 1(a) and (b).
The phase-stability regions for the ironnitrogen system can obviously be visualized in a
diagram showing temperature versus chemical
potential of nitrogen of the atmosphere surrounding the specimen, provided “local equilibrium”
with the surface of the specimen has been realized. This “potential diagram” should not be
confused with a “normal” phase diagram presenting the phase-stability regions in a diagram
showing temperature versus nitrogen content of
the iron-nitrogen alloy.
The described given treatment makes clear that
there is a one-to-one relationship between the
chemical potential of nitrogen, mN,s, and the nitrogen activity, aN,s (Eq 3, 4), and the nitriding potential, rN (Eq 5b). So, instead of the previously
described potential diagram, corresponding diagrams are obtained upon plotting temperature versus nitrogen activity (activity diagram) or upon
plotting temperature versus nitriding potential.
The latter diagram for the iron-nitrogen system
has, for the first time, been presented by Lehrer
(Ref 22) and is called the “Lehrer diagram.” The
activity diagram and the Lehrer diagram for the
iron-nitrogen system are shown in Fig. 4(a) and
(b), respectively.
** The pressure of the reference state is taken equal for all gas
components, and thus, pN0 2 is replaced by p0. Normally, p0 is
taken equal to 1 atm (see the following).
Fundamentals of Nitriding and Nitrocarburizing / 623
Fig. 4
(a) Temperature vs. nitrogen activity (Eq 5b). The corresponding equilibrium pressure of the hypothetical N2 gas occurring in Eq 1(a) and (b), as calculated using Eq 5(a) with the
reference pressure, p0, taken as 1 atm, has been indicated on the abscissa shown at the top of the figure. This (very high) equilibrium pressure actually must be read as a fugacity
(see text following Eq 3) (taken from Ref 20). (b) Lehrer diagram of temperature vs. nitriding potential (cf. Eq 6). The ammonia content in an NH3-H2 gas mixture at 1 atm total pressure
corresponding with a given nitriding potential has been indicated on the abscissa shown at the top of the figure (taken from Ref 20). (c) Part of the Lehrer diagram showing the original a/g0
and a/g phase boundaries as presented by Lehrer (Ref 22) and the corresponding newer data pertaining to the experimentally accessible stationary states and the (hypothetical) (local)
equilibria derived as described in section 8.1 and, in particular, Ref 23. Source: Ref 23
It should be realized that a two-phase field as
in the normal iron-nitrogen phase diagram does
not occur in the potential diagram, the activity
diagram, or the Lehrer diagram, because at each
temperature, the chemical potentials of nitrogen
(the nitrogen activities, the nitriding potentials)
are the same for the two iron-nitrogen phases in
equilibrium with each other and the surrounding nitriding atmosphere; that is, the two-phase
stability “region” in the Lehrer diagram is given
by a line. Similarly, the eutectoid, where three
solid phases, a-Fe[N], g-Fe[N], and g0 -Fe4N1x,
are in equilibrium with each other and the surrounding nitriding atmosphere, is represented
by a point in the Lehrer diagram.
It is impressive to establish how high the
quality of the experimental work performed by
Lehrer in 1930 was; up until now, only minor
modifications of the experimental data in the
Lehrer diagram have been necessary. However,
as discussed in the Section 8.1, the a/g0 and
a/g phase-boundary lines in the experimental
Lehrer diagram above approximately 580 C
(853 K) may actually represent stationary states
rather than (local) equilibrium. The transition of
local equilibrium (at lower temperatures) to stationary state (at higher temperatures) occurs at
decreasing temperature for increasing nitriding
potential, and thus, in the case of iron nitrides
present at the surface of the specimen, the transition from local equilibrium to stationary state
can occur at a temperature distinctly lower than
for nitrogen ferrite (see Section 8.1).
The most recent evaluation of the a/g0 and
a/g phase boundaries in the Lehrer diagram is
shown in Fig. 4(c), together with the original
Lehrer phase-boundary lines. The experimentally determined coordinates of the triple-point
a/g0 /g (where the eutectoid reaction g ! a + g0
runs) are T = 593 C (866 K) and rN =
0.139 atm1/2. However, these data intrinsically
refer to the in-reality-occurring stationary state
at the surface; the coordinates of the triple-point
a/g0 /g in the (hypothetical above approximately
580 C, or 853K) Lehrer diagram pertaining to
(local) equilibria (at the surface of the specimen/component) are T = 593 C (866 K) and
rN = 0.135 atm1/2; for full discussion, see
Section 8 and Ref 23. The solubility of nitrogen
in a-Fe at this point is 0.441 at.% (Fig. 2), which
is the maximum equilibrium solubility of nitrogen in ferrite. This last result follows from substitution of rN in the absorption function of ferrite
describing the equilibrium solubility of ferrite
as a function of rN and T (Ref 23).
5. Controlled Nitriding
To be able to steer the outcome of the nitriding process, the chemical potential of nitrogen
in the nitriding atmosphere must be controlled;
that is, it should be set and maintained at a certain selected value, or it should be varied in a
prescribed way. In the context of the treatment
in Section 4, it then follows that control of the
624 / Nitriding and Nitrocarburizing of Steels
nitriding potential in an NH3-H2 gas mixture is
imperative. This can be achieved by nitriding in
a flow of NH3-H2 gas mixture that is of a fixed
composition (corresponding to the desired
nitriding potential).
A stationary gas atmosphere is inappropriate,
because NH3 will tend to decompose (catalytically activated by the surface of the iron-base
component and/or the furnace walls) to realize
equilibrium with respect to its thermal decomposition in the gas atmosphere, which should
be avoided (if controlled nitriding is the goal).
This implies that a certain minimum (linear)
gas flow rate is required (see Fig. 4 in Ref 24).
A simple way to check this is to determine the
gas composition at the entrance of the furnace
and at the outlet; these should be identical.
One could imagine that the gas from the outlet
is reused as entrant gas at the inlet.
To apply the aforementioned approach, additional kinetic conditions must be met:
Equilibrium (Eq 2) is established instanta-
neously (in any case, fast as compared to
the nitriding time). This is not fully the case;
at the start of nitriding, some time is needed
to attain the activity/concentration of nitrogen in the solid at the surface corresponding
with local equilibrium at the surface of the
solid with the gas atmosphere (i.e.,
corresponding to the fixed nitriding potential, Eq 5b). See the discussion in Section 8.
Other gas components potentially present in
the furnace should behave as inert gases.
For example, this holds for the presence of
N2, which may have been added to the gas
atmosphere; it can be taken for certainty that
equilibrium (Eq 1a) is not established.
The iron-nitrogen phases of the solid as
developed in the nitrided zone (not at the
very surface) are not in thermodynamic
equilibrium, and therefore, these iron-nitrogen phases (underneath the surface) in principle can decompose, causing N2 gas
development and thus the formation of
pores. Ideally, the kinetics of this process
should be infinitely slow at the nitriding
temperature. Unfortunately, that is not the
case, and thus, appreciable porosity can
occur during nitriding, especially close to
the surface, which is the oldest part of the
nitrided zone and which is the part of the
nitrided zone where the largest nitrogen concentration/activity occurs and thus where the
thermodynamic driving force for the decomposition is largest. Indeed, especially the surface-adjacent e-Fe2N1z phase is known to
show porosity. As a consequence, transport of
nitrogen through the nitrided zone (induced
by the loss of dissolved nitrogen by pore formation) is nonnegligible, and a local equilibrium cannot be established at the surface.
However, the situation could have been worse:
if nature would have given much faster kinetics to these decomposition processes in the
solid iron-nitrogen phases, no nitriding process
of iron-base components would be possible.
Epilogue. The phenomenon of severe pore formation is not restricted to only the e-Fe2N1z
phase. Nitrogen austenite (g-Fe[N]) can accommodate much more nitrogen than nitrogen ferrite
(a-Fe[N]), up to approximately 10.3 at.% N
versus approximately 0.4 at.% N (Table 1), and
the driving force for its decomposition is larger
compared to the case of nitrogen ferrite. Also,
because of higher temperatures necessary
for austenitization by nitriding (cf. Fig.2),
the kinetics of its decomposition can be
relatively fast.
Thus, only a thin layer of austenite at the surface can be maintained upon continued nitriding of a pure iron foil in the austenite-phase
field (Fig. 2) beyond the time necessary for
complete transformation of the entire, originally pure iron foil to nitrogen austenite. The
experiment discussed here (Ref 25) involved
nitriding a pure iron foil with a thickness of
1 mm (0.04 in.) at 810 C (1083 K) in a 5%
NH3-95%H2 flowing gas mixture for times up to
93 h. (The entire foil had transformed to nitrogen
austenite after approximately 24 h.) Underneath
the nitrogen-austenite layer at the surface, which
is stabilized by nitrogen uptake at the surface
from the nitriding atmosphere, the originally produced nitrogen austenite at these depths has
decomposed into a-Fe and N2, leading to severe
pore development. These pores develop preferentially at the grain boundaries in the nitrogenaustenitic foil, where grain boundaries run
approximately perpendicular to the surface. By
subsequent coalescence of the pores, channels
at these grain boundaries are formed, which are
in contact with the outer atmosphere.
Continued nitriding then merely implies
“pumping” of nitrogen through the system
without the gain of nitrogen being realized by
the specimen. Nitrogen enters the specimen
through the surface of the specimen, where the
thin layer of nitrogen austenite is located, and
leaves the specimen in the ferritic regions
within the specimen at the channel walls,
according to 2[N]a-Fe ! N2↑ (Ref 25).
Carburizing in a CO-CO2 gas mixture can be
conceived as the sum of the following (hypothetical) reactions:
CGr $ ½C
(Eq 7a)
2CO $ CGr þ CO2
(Eq 7b)
giving:
2CO $ ½C þ CO2
where CGr denotes a hypothetical graphite,
playing a role as the hypothetical N2 gas in
Eq 1(a) and (b), and [C] represents carbon dissolved in the solid substrate M. Establishment
of the equilibrium (Eq 8a), which is called the
Boudouard reaction, implies the occurrence of
local equilibrium at (only) the surface of the
substrate M (cf. Sections 4 and 8).
Equation 7(a) implies that:
mCGr ;s ¼ mC;s
(Eq 9)
where CGr ;s and mC,s represent the chemical
potentials of carbon of the (hypothetical) graphite and the solid substrate, respectively. It follows from Eq 9:
m0CGr ;s þ RT ln aCGr ;s ¼ m0C;s þ RT ln aC;s
(Eq 10)
where m0i is the chemical potential of component i (i = CGr,s or Cs) in the reference state
(temperature dependent at the selected pressure
of the reference state; see the following), R is
the gas constant, T is the absolute temperature,
aCGr ;s is the activity of the (hypothetical) graphite in Eq 7(a) and (b), and aC,s is the activity of
carbon in the substrate. Now, m0C;s is selected
such that m0CGr ;s ¼ m0C;s . Then it is obtained:
aC;s ¼ aCGr
(Eq 11a)
Using the equilibrium constants, K, of (Eq 7b)
and (Eq 8a):
Kð7bÞ ¼
6. Carburizing Potential and
Controlled Carburizing
Carburizing of steel can be realized, in a COCO2 gas atmosphere for example. The definition
of a carburizing potential is possible in terms of
just the CO and CO2 gas components alone, as
long as no other gas components are present that
can react with CO and CO2. However, the desired
nitriding effect in a nitrocarburizing gas atmosphere involves the presence of H2 (see Section
4), which introduces additional reactions involving CO and CO2 that obstruct a unique definition
of a carburizing potential in the case of nitrocarburizing (see the next section). To prepare for a
discussion of nitrocarburizing, first a treatment
for the case of carburizing in a CO-CO2 gas
mixture is given that parallels the one given for
nitriding in Section 4.
(Eq 8a)
aCGr ;s pCO2 p0
p2CO
Kð8aÞ ¼
aC;s pCO2 p0
p2CO
it follows:
aC;s ¼
Kð7bÞ rC Kð8aÞ rC
¼
p0
p0
(Eq 11b)
with a so-called carburizing potential, rC,a,
defined by:
rC;a p2CO
pCO2
(Eq 12a)
where pCO and pCO2 are the partial pressures of
the really present gas components CO and CO2.
The reference state for the carbon dissolved
in M is selected as graphite at 1 atm. As a
Fundamentals of Nitriding and Nitrocarburizing / 625
consequence of this step, the numerical value of
the carbon activity, aCGr ;s , of the hypothetical
graphite appearing in Eq 7(a) and (b) can have
a value deviating considerably from 1.
The above presented treatment makes clear
that there is a one-to-one relationship between
the chemical potential of carbon, mC,s, and the
carbon activity, aC,s (Eq 9, 10), and the carburizing potential, rC (Eq 11b). So, the carburizing
potential can be used as a direct measure for the
chemical potential of carbon in the solid at the
surface, provided that “local equilibrium” at the
surface prevails (see Section 8). Thus, as long as
pure carburizing is considered, the carburizing
potential, as defined by Eq 12(a) for the CO-CO2
gas mixture as carburizing atmosphere, can be
applied as a carburizing controlling parameter,
similar to the nitriding potential in the case of
nitriding.
For carburizing on the described basis,
kinetic constraints must be satisfied that are
more or less parallel with those formulated for
nitriding in the previous section:
A stationary gas atmosphere is inappropri-
ate. As follows from the thermodynamics
of equilibrium (Eq 7b) at 1 atm, graphite formation (sooting) could, in principle, occur to
a significant extent at 1 atm at temperatures
below approximately 700 C (973 K). For
the CO-CO2 gas mixture, such sooting can
kinetically be avoided by applying a gas
flow rate high enough to avoid the formation
of graphite (soot) due to the relatively slow
kinetics of equilibrium (Eq 7b).
Equilibrium (Eq 8a) should be established instantaneously (cf. Section 8). This is not the case, and
such effects are nowadays taken into account in
procedures for controlled carburizing.
The iron-carbon phases of the solid as developed in the carburized zone (not at the very surface) are not in thermodynamic equilibrium,
and therefore, these iron-carbon phases (underneath the surface) in principle can decompose,
causing graphite development. Such graphite
formation can lead to metal dusting, i.e. disintegration of iron carbides, produced in the carburizing process, into iron and graphite.
As an interesting point, it is noted that such
sooting and metal dusting phenomena appear
to be suppressed if NH3 is added to the carburizing gas mixture, that is, nitrocarburizing (Ref
26, 27). Moreover it is shown in Ref 26 and 27
that appropriate choice of the gas composition
in CO-H2-N2-NH3 gas mixtures allows the
growth of massive cementite layers on ferrite
(a-Fe).
which is called the heterogeneous water gas
reaction. By a treatment similar to that given
for the Boudouard reaction (Eq 8a) in a pure
CO-CO2 gas mixture, a carburizing potential
can be defined as a direct measure for the
chemical potential of carbon in the solid at the
surface, provided that local equilibrium at the
surface prevails (see Section 8). This carburizing potential for a pure CO-H2-H2O gas mixture is given by (following a treatment
completely analogous to the one in the previous
section):
rC;b pCO pH2
pH2 O
Carburizing can also be achieved by putting
specimens/components in a CH4-H2 atmosphere according to:
CH4 $ ½C þ 2H2
Carburizing can also be achieved by putting
specimens/components in a CO-H2-H2O atmosphere according to:
CO þ H2 $ ½C þ H2 O
(Eq 8b)
(Eq 8c)
Again, like the Boudouard reaction in a pure
CO-CO2 atmosphere (Eq 8a) and the heterogeneous water gas reaction (Eq 8b), a carburizing
potential can be defined as a direct measure for
the chemical potential of carbon in the solid at
the surface, provided that local equilibrium
at the surface prevails. This carburizing potential for a pure CH4-H2 gas mixture is:
rC;c pCH4
p2H2
(Eq 12c)
Obviously, the carburizing potential for realizing the same chemical potential of carbon in the
gas atmosphere can have completely different
values and dimensions (units) for different carburizing atmospheres (Eq 12a–c). This is trivial.
The relevance of the previous discussion is to
demonstrate that in the presence of a (carburizing) gas atmosphere comprising the components
CO, CO2, H2, H2O, and CH4 (Eq 8a–c), it is generally impossible to characterize the chemical
potential of carbon in the gas atmosphere by
either rC,a, rC,b, or rC,c (Eq 12a–c). This recognition has the following unpleasant consequence
for gaseous nitrocarburizing.
A gas atmosphere (initially) containing NH3
and H2 gas components, to establish a nitriding
reaction, and CO and CO2 gas components, to
establish a carburizing reaction, is subject to
side reactions between these gas components:
CO2 þ H2 $ CO þ H2 O
(Eq 13)
which is called the homogeneous water-gas
reaction, and:
CO þ 3H2 $ CH4 þ H2 O
7. Controlled Nitrocarburizing
(Eq 12b)
(Eq 14)
Thus, the simultaneous presence of the gas
components CO, CO2, H2, H2O, and CH4 in a
nitrocarburizing gas atmosphere must be considered. One may hope that of the various carburizing equilibria (Eq 8a–c), possibly only
one is established fast. It has been found
that equilibrium (Eq 8b), the heterogeneous
water-gas reaction, is established much faster
than the equilibria (Eq 8a and c) in the case of
carburizing pure iron (in the ferritic or austenitic state) (Ref 28). However, whether this
statement also holds if nitrides, carbides, carbonitrides, or nitrocarbides are at the surface of
the specimen (developed in the course of the
treatment) is unknown. Moreover, the simultaneous occurrence of the side reactions (Eq 13,
14) will complicate any calculation/prediction
and control of the nitrocarburizing reaction;
application of a nitriding potential and a carburizing potential on the basis of the gas inlet composition appears senseless.
An approach is possible that avoids the ambiguity of the definition of the chemical potential
of carbon according to either of the equilibria
(Eq 8a–c) and that, at the same time, accounts
for the occurrence of the side reactions (Eq 13
and 14), as follows. It can be shown that if the
gas composition is chosen and controlled such
that equilibria (Eq 13 and 14) are established
(i.e., imposed), then, as a consequence, the
chemical potentials of carbon in the gas phase,
according to any of the three equilibria (Eq
8a–c) are equal (Ref 29, 30). Thus, any further
considerations concerning which of the three
equilibria is first established or not and the consequences of changes in the gas composition by
reactions (Eq 13 and 14) are unnecessary.
Thus, it is possible to nitrocarburize with chosen chemical potentials, separately for nitrogen
and carbon, in the gas atmosphere. This has been
demonstrated in Ref 29, which can be considered
as the outcome of a development departing from
Ref 30 to 32. In this case, the nitriding potential
can still be defined as given by Eq 6 and given a
value accordingly, as equilibrium (Eq 2) is the
only nitriding pathway in the gas atmosphere
considered. To use the notion carburizing potential is not meaningful. Only if one of the possible
carburizing reactions is the operating pathway
for carburizing, one could define the carburizing
potential by the corresponding formulation (one
of the three Eq 12a, b, or c). That such a case
occurs in reality is uncertain (see discussion following Eq 14).
The logical consequence of this discussion
would be to simply characterize a nitrocarburizing atmosphere, as discussed here and having
a constant composition establishing the equilibria (Eq 13 and 14), as a nitrocarburizing atmosphere of specific chemical potentials of
nitrogen and carbon. At present, this approach
appears to be the only feasible way toward controlled nitrocarburizing and, until today (2013),
has been established only in the laboratory
(Ref 29).
It should be recognized that the chemical
potentials of nitrogen and carbon, as set for the
gas atmosphere, are only equal to the chemical
potentials of nitrogen and carbon in the specimen
(at its surface) if local equilibrium at the surface
of the specimen has been established. This
may take a relatively long time, especially in
nitrocarburizing treatments, and leads to complicated microstructural developments of the
626 / Nitriding and Nitrocarburizing of Steels
compound layer composed (largely) of iron carbonitrides (see Section 9.2).*
Of course, more side reactions than considered previously may be considered as possible
in a nitrocarburizing atmosphere. Thus, for possibly produced gas components, in particular
O2, C2H2 (acetylene), C2H4 (ethylene), C2H6
(ethane), and HCN, it can be shown that their
partial pressures are so low that their formation
and thus their effect on the gas composition and
the chosen values of the chemical potentials of
carbon and nitrogen are negligible (Ref 29).
NH3 $ Nads þ3=2 H2
8. Local Equilibria and Stationary
States
Local equilibrium at the gas-solid interface
(Fig. 5) can be established under the following
conditions:
8.1 Gas-Solid Interface. The imposition of
a fixed chemical potential of nitrogen, and, in
the case of nitrocarburizing, also a fixed chemical potential of carbon, at the surface of the
specimen/component implies that the gas composition in the furnace should not change from
the desired inlet composition. Thus, in the case
of nitriding, any thermal dissociation of ammonia in the furnace must be avoided (or the
degree of this dissociation should be known
precisely, and be controllable, at the surface
of the specimen/component). If the gasatmosphere equilibrium is at 1 atm and at a
temperature above approximately 350 C
(623 K), ammonia would be practically fully
dissociated. Hence, a stationary gas atmosphere
is inappropriate for controlled nitriding.
Because this dissociation is a relatively slow
process but catalytically activated by the presence of iron (the specimen and possibly the furnace walls; see Fig. 4 in Ref 24), this thermal
dissociation in the case of nitriding iron-base
specimens/components can be made negligible
by application of a sufficiently large (linear) gas
flow rate in the furnace. However, if the nitriding
potential must be very large (i.e., for an NH3-H2
gas mixture implying a fraction of NH3
approaching 100%), even a tiny amount of thermal dissociation of ammonia causes the real
nitriding potential to deviate distinctly from the
one calculated from the gas composition at the
furnace inlet. Such a situation is not expected
for the nitriding of iron-base materials but happens, for example, upon nitriding nickel to
produce a Ni3N compound layer requiring an
extremely large nitriding potential (e.g.,
performing nitriding with a gas composed of pure
ammonia at the gas inlet, which corresponds with
rN = 1 at the gas inlet, Eq 6) (Ref 33).
Now, take a closer look at the establishment
of equilibrium (Eq 2). At the surface of the
specimen, ammonia molecules are adsorbed
and dissociate by stepwise removal of hydrogen
atoms (Ref 34). This leads to adsorbed individual nitrogen atoms at the surface, Nads:
Dissolution of nitrogen only occurs accord-
* The carbon solubility of a ferritic pure-iron substrate is so
low (<0.1 at.% C; Table 1) that the carbon uptake by
nitrocarburizing is effectively confined to the compound layer.
(Eq 15)
Next, two routes can be followed by the
adsorbed nitrogen atoms. They can dissolve
into the solid substrate, which is the desired
effect:
Nads $ ½N
(Eq 16)
or they recombine at the surface and desorb,
which counteracts nitriding (Fig. 5):
Nads þ Nads $ N2 "
(Eq 17)
ing to Eq (16).
The diffusion of dissolved nitrogen from the
surface to larger depths is negligible.
The equilibrium according to Eq 16 is estab-
lished, which can be taken as certain because
the rate of nitrogen uptake according to Eq
16 is very fast compared to the rates of the
reactions corresponding with Eq 15 and 17.
The equilibrium according to Eq 15 is established (i.e., the forward and backward reactions according to Eq 15 are equal).
Under these conditions, the chemical potential
of nitrogen in the gas atmosphere (related oneto-one to the nitriding potential; Eq 4, 5b,
and 6), as determined from the NH3 and H2
contents in the gas atmosphere, can be equal
to the chemical potential of nitrogen in the solid
substrate at its surface.
It should be recognized that, in principle,
as long as transport of nitrogen from the surface
into the bulk of the specimen occurs, a genuine
equilibrium at the surface cannot occur, albeit it
can be approached very closely; in this sense,
the term local equilibrium is somewhat misleading. In the beginning stage of nitriding a nitrogen-less specimen, the diffusion of dissolved
nitrogen from the surface to larger depths can
be significant, compared to the rate of nitrogen
uptake according to Eq 15 and 16, due to the
occurrence of a pronounced gradient for the dissolved nitrogen content in the solid at the surface
(cf. Fick’s first law). Then, upon gradual saturation of the substrate, a gradual increase of the dissolved nitrogen content at the surface occurs until
the dissolved nitrogen content at the surface
equals the equilibrium value (Fig. 6).
However, if the recombination and desorption of the adsorbed nitrogen according to
Eq 17 occurs with a nonnegligible rate** compared to the rate of nitrogen uptake according
to Eq 15 (assuming the equilibrium, Eq 16, is
always established), then the chemical potential
** The nitriding reaction according to Eq 17 can be
neglected due to its comparatively low rate (Ref 35) and
the usually low partial pressure of nitrogen, pN2 (usually <1
atm), in the gas atmosphere.
of the dissolved nitrogen at the surface of the
specimen is smaller than that determined from
the NH3 and H2 contents in the gas atmosphere
(cf. Eq 2), which leads to the consideration of
two different cases:
If no significant diffusion of dissolved nitro-
gen into the bulk of the specimen occurs, a
time-independent situation (at constant temperature and constant pressure) at the surface
is established as a consequence of equal
rates of nitrogen uptake (Eq 15; equilibrium,
Eq 16, is always established) and nitrogen
desorption (Eq 17), and thus, although no
equilibrium of the surface of the solid with
the gas atmosphere according to only Eq
2 is established, a so-called stationary state
at the gas-solid interface has been realized,
characterized by a smaller content of dissolved nitrogen in the solid substrate at the
surface than would occur if equilibrium at
the gas-solid interface would prevail according to only Eq 2. This state of affairs is illustrated in Fig. 5 (Ref 24, 34).
If significant diffusion of dissolved nitrogen
into the bulk of the specimen does occur,
as at the start of nitriding a nitrogen-less
specimen, a similar situation as described
in the preceding paragraph for the absence
of significant recombination and desorption
of adsorbed nitrogen emerges; a gradual
increase of the concentration of dissolved
nitrogen occurs. However, in contrast with
the case described in the preceding paragraph, the eventual concentration of dissolved nitrogen in the solid at the surface
does not equal the equilibrium value but is
smaller than that. Using kinetic data from
Ref 34 to 36 for the reactions considered
here, such an evaluation of the dissolved
nitrogen (surface) concentration for an a-Fe
Fig. 5
Concentration of dissolved nitrogen in the solid
(at the surface), cN, vs. the nitrogen-uptake rate,
dcN /dt. If the nitrogen-uptake rate according to Eq 16 is
relatively very fast and the desorption according to Eq
17 can be neglected, then equilibrium at the surface of
the solid is established, if the forward and backward
reactions according to Eq 15 are of equal magnitude;
that is, the dynamic equilibrium as given by Eq 15 is
then established at the surface of the solid. However, if
the rate of nitrogen desorption according to Eq 17 cannot
be neglected, such a local equilibrium is not realized.
Instead, a stationary state occurs that is given by equality
of the (net) nitriding rate due to (net) decomposition of
NH3 and the denitriding rate due to nitrogen desorption
from the surface. Source: Ref 24 as adapted from Ref 34
Fundamentals of Nitriding and Nitrocarburizing / 627
foil of 500 mm thickness, nitrided at 580 C
(853 K), and with rN = 0.104 atm1/2 is
shown in Fig. 6 (the first such calculation
was presented in Ref 37, neglecting the role
of recombination and desorption of the
adsorbed nitrogen atoms, Eq 17; this contribution was included in a later work, Ref 38,
leading to, for the conditions considered, only
marginal differences, in accordance with the
results shown in Fig. 7). A detailed description
of the calculation procedure and a summary of
the available kinetic data is provided in the
Appendix of Ref 23 (which is the first presentation in the English literature of these rateconstant data as provided in the original
Ref 28, 34, and 35, written in German; also a
drastic correction of a rate constant as
provided in Ref 28 is presented there).
The tendency to recombination and desorption of Nads becomes significant above approximately 580 C (853 K) for pure iron (ferrite) as
substrate. The difference between the equilibrium concentration, cN,eq, and the lower stationary value of the concentration, cN,st, at the
surface of the nitrided component increases
with increasing temperature and nitriding
potential (Fig. 7). Then, the result obtained for
the content of dissolved nitrogen in the solid
at the surface, that is, cN,st, does not correspond
with the chemical potential of nitrogen in the
gas atmosphere, as determined from the NH3
and H2 contents in the gas atmosphere and as
expressed by the nitriding potential (Eq 4–6).
Clearly, in view of the results shown in
Fig. 7, at the usual nitriding temperatures
(500 to 580 C, or 773 to 853 K), local equilibrium at the surface of at least pure iron can
be supposed, provided the diffusion rate of the
absorbed nitrogen to larger depths is negligible
as compared to the nitrogen-uptake rate, that is,
only at later stages of nitriding (Fig. 6). At
higher temperatures and/or at relatively high
nitriding potentials, the contribution of the
recombination and desorption according to Eq
17 can no longer be neglected. For example,
at a nitriding temperature of 580 C (853 K, a
nitriding temperature applied in practice), the
nitriding potential should not be larger than
6.1 atm1/2 (pertaining to a nonexceptional
75%NH3-25%H2 gas mixture), in order that
the difference of cN,eq and cN,st is smaller than
1%. Under the latter conditions, nitrogen ferrite
is likely not the equilibrium phase at the surface
of the specimen/component, but assuming that
the rate constants for Eq 15 to17 for iron nitride
at the surface of the specimen/component do
not differ much from those for pure iron (no
such data for the iron nitrides are available),
the preceding sentence suggests that nitriding,
under conditions such that, in particular, e-iron
nitride occurs at the surface of the specimen/
component, can be associated with the occurrence of a stationary state at the surface already
at temperatures distinctly lower than 580 C
(853 K; see the experimental results presented
in Fig. 3 of Ref 24).
The development of the nitrogen concentration-depth profile in an a-Fe foil of 500 mm thickness upon
nitriding in an NH3-H2 gas mixture of rN = 0.104 atm1/2 at 580 C (853 K). Three competing reactions
are considered here: (1) nitriding according to Eq 15 (equilibrium, Eq 16, is assumed to be realized instantaneously;
see the text and the caption for Fig. 5); (2) diffusion of absorbed nitrogen into the initially unsaturated solid;
and (3) recombination of adsorbed nitrogen and its subsequent desorption according to Eq 17. The rate equations
and the rate constants (as given by and as derived from data in Ref 35 and 36) for these calculations have
been described in detail in the Appendix of Ref 23, departing from original treatments in Ref 28 and 34 to 36. In
view of the temperature and nitriding potential pertaining to the case considered here, the difference between the
value of the dissolved nitrogen content at the surface according to local equilibrium, cN,eq, and the lower stationary
value of this concentration, cN,st, as observed after homogenization of the solid has been realized is marginal
(cf. Fig. 7).
Fig. 6
Fig. 7
The difference between the value of the dissolved nitrogen content at the surface according to local
equilibrium, cN,eq, and the lower stationary value of this concentration, cN,st (as illustrated in Fig. 5), as a
function of temperature for various nitriding potentials. Obviously, these results hold for a homogeneous solid (here:
nitrogen ferrite) having a nitrogen concentration at all depths equal to that in the solid at the surface. These
calculated results thus pertain to the competition of (cf. the caption of Fig. 6) nitriding according to Eq 15
(Equilibrium, Eq 16, is assumed to be realized instantaneously; see text and the caption of Fig. 5) and recombination
of adsorbed nitrogen and its subsequent desorption according to Eq 17. As shown in the Appendix of Ref 23, an
analytical equation relating cN,eq and cN,st can be given.
628 / Nitriding and Nitrocarburizing of Steels
Fig. 8
Schematic illustration of the progressive microstructural stages of compound-layer formation
and evolution upon nitriding a-Fe. (a) Nucleation of g0
nitride at the surface followed by its growth by nitrogen
supply via the g0 nuclei formed, but in particular also by
nitrogen supply via the ferrite that surrounds them, as an
efficient bypass possibility using the much faster
diffusion of nitrogen through ferrite than through the
nitride. (b) At the same time, e phase may develop on
top of the g0 particles before these have produced, by
lateral growth, a closed layer. (c) As a result, a double
e/g0 layer has formed, and further growth can only be
established by nitrogen transport through both sublayers.
(d) Decomposition of (especially e) iron nitride
underneath the surface of the compound layer occurs for
extended stages of nitriding, leading to development of
pores filled with N2 gas at grain boundaries and even
within the grains. (e) Coalescence of pores at grain
boundaries (usually oriented more or less perpendicular
to the surface) leads to channels at grain boundaries in
contact with the outer nitriding atmosphere. Source: Ref
19, 39, and 41
The previously mentioned consideration has
as a consequence that the phase boundaries presented in the potential Lehrer diagram (Fig. 4b)
above approximately 580 C (853 K) (and
above lower temperatures at high nitriding
potentials as correspond to the development of
iron nitrides) are incorrect, because the experimentally determined nitriding potential for the
phase boundary (at a certain temperature) corresponds to a stationary state rather than a
(local) equilibrium. The true nitriding potential,
characterizing the chemical potential of the
hypothetical gas atmosphere in equilibrium
with the solid (at the phase boundary) and given
by the NH3 and H2 contents of that hypothetical
gas atmosphere, will be smaller than the experimentally determined one (see the results shown
for the a/g phase boundary in Fig. 4c and the
discussion at the end of the Section 4). The
effect may be considerably larger for the g0 /e
phase boundary, as indicated in the previous
paragraph.
8.2 Solid-Solid Interface. Upon nitriding of
an iron substrate, various iron-nitrogen phases
may develop in the nitrided surface-adjacent
zone of the substrate. In particular in the case
of pure iron, these phases can occur as a layered
structure, with the order of iron-nitrogen phases
corresponding with a decrease in the local
chemical potential of nitrogen in the solid with
increasing depth beneath the surface, in accordance with the iron-nitrogen phase diagram
(Fig. 2) and the activity and potential (Lehrer)
diagrams (Fig. 4) (see Section 9).
At the interfaces between two such iron-nitrogen phases (layers), a situation described as
local (metastable; see the following) equilibrium (cf. the previous discussion) can occur if
the compositions of both phases at the interface
comply with those pertaining to the metastable
equilibria presented by the phase diagram (note
that realizing true thermodynamic equilibrium
in the solid, underneath the surface, requires
decomposition of these phases in Fe and N2, as
discussed in Section 3 in this article). The accumulation or depletion of nitrogen at either side
of the solid-solid interface implies that local
equilibrium can only be maintained if the mobility of the interface is infinitely large, so that the
redistribution of nitrogen, to restore the local
equilibrium situation at the interface, is instantaneous. In this case and if the accumulation/
depletion of nitrogen at the interface is due to
diffusional transport of nitrogen in the (layer)
system, the phase transformation at the interface
is a diffusion-controlled one (Ref 14). Evidently, during growth of the nitrided zone,
implying a net inward transport of nitrogen, also
across the interface considered, and recognizing
the real, finite rate of interface migration, genuine equilibrium cannot be established at the
interface, but it can be closely approached.
In the case of nitrocarburizing, the occurrence of local equilibrium at a solid-solid interface implies that the nitrogen concentrations
at both sides of the interface and the carbon
concentrations at both sides of the interface
must correspond with the metastable (ternary)
phase diagram. To indicate one of the possible
complications with two diffusing components,
it is conceivable that local equilibrium at a
solid-solid interface could be realized for one
of the components, whereas this may not be
possible for the other component.
The evidence until now suggests that at the
solid-solid interfaces in the nitrided/nitrocarburized zone of iron-base specimens/components, local equilibrium is closely approached.*
9. Microstructural Development of
the Compound Layer
9.1 Microstructural Development of the
Compound Layer during Nitriding. A schematic presentation of the development of the
compound layer upon nitriding a-Fe is provided
in Fig. 8 (Ref 39–41); micrographs of the twophase compound layer are shown in Fig. 9 for
cases of less (Fig. 9a, lower temperature,
shorter time) and more (Fig. 9b, higher temperature, longer time) porosity. At the start of the
nitriding process, nitrogen is dissolved in the
ferrite matrix. At the moment that the solubility
of nitrogen in ferrite (at the surface) equals and
becomes larger than that compatible with the
a/g0 equilibrium, g0 iron nitride must form thermodynamically. This need not occur at the start
of nitriding; an apparent incubation time for
nitride formation occurs (Ref 37, 38). After a
certain time, the nitrogen concentration in the
solid at the surface has risen to a level allowing
the development of g0 iron nitride as a result of
the competition of:
Nitrogen supply to the surface from the gas
via dissociation
Removal of dissolved nitrogen from the
surface by diffusion into the interior of the
specimen/component
Removal of nitrogen adsorbed at the surface
by recombination and desorption from the
surface (see discussion in Section 8.1)
The g0 nitride particles nucleated at the surface
initially do not form a closed layer (Fig. 8a). They
grow largely via supply of nitrogen from the ferrite that surrounds them, rather than by nitrogen
diffusion through the particles, because the diffusion of nitrogen through ferrite is much faster
than that through the nitride. Yet, after some
time, lateral growth of the nitride particles at the
surface leads to a closed layer. Depending on
the nitriding potential applied, before a closed
layer has formed, e iron nitride may have
* This assumption of the occurrence of local equilibria at
solid-solid interfaces in diffusion zones is often taken for
granted, and the compositions in both phases at the interface
have then been interpreted as defining phase boundaries in
the corresponding phase diagram. However, this cannot be
generally true, and therefore, phase boundaries determined
in this way and taken up in published phase diagrams may
be flawed.
Fundamentals of Nitriding and Nitrocarburizing / 629
Cross sections of nitrided pure a-Fe (light optical micrographs, after etching in 1.0 vol% Nital containing
0.1 vol% HCl). (a) Nitrided at 550 C (823 K) for 5 h at rN = 2.37 atm1/2. (b) Nitrided at 560 C (833 K)
for 20 h at rN = 2.37 atm1/2. The dark spots and stripes observed in the e sublayer for the specimen corresponding
to (b) are due to porosity developing during nitriding in the iron nitride underneath the surface, due to its
metastability. Inter- and intragranular pores can form. Source: Ref 41; cf. Fig. 3 and Section 3
Fig. 9
nucleated on top of the g0 particles. After closure
of the nitride surface layer, further growth of the
e/g0 compound layer requires diffusion of nitrogen through the layer from top to bottom.
As discussed in Section 3 of this article,
the iron nitrides beneath the very surface are
not in thermodynamic equilibrium; they tend
to decompose in Fe and N2. This leads to porosity upon prolonged nitriding, especially in the
oldest part of the compound layer, which is also
the most nitrogen-rich part of the compound
layer and the part adjacent to the surface
(Fig. 8d and e; see also Fig. 3 and 9b). Only
in the surface-adjacent region (where the e
phase occurs) is porosity visible with a light
microscope. The g0 iron nitride, underneath
the e iron nitride and not in contact with the
nitriding atmosphere, tends to decompose as
well and develop porosity, which usually cannot be observed with a light microscope but at
higher magnification with an electron microscope. This is also the case for the nitrogen ferrite, constituting the diffusion zone underneath
the g0 iron nitride.
9.2 Microstructural Development of the
Compound Layer during Nitrocarburizing.
The microstructural evolution of the compound
layer produced upon nitrocarburizing ferritic
iron is much more complicated than in the case
of nitriding. Of course, this relates to the occurrence of the simultaneous inward diffusion of
two components, nitrogen and carbon, and the
more complex metastable (see Section 3) phase
equilibria of the ternary Fe-C-N system. Until
now, only a few papers have appeared in the
literature presenting a more or less systematic
study of the microstructural changes accompanying compound-layer growth during nitrocarburizing (Ref 31, 42–45). The following
summarizes the common, but still limited, basis
of understanding, based largely on results presented in Ref 31 and 45.
Local equilibria appear to be established at
the interphase boundaries in the compound
layer, whereas, at the same time, a local equilibrium (or a stationary state) does not prevail
at the gas-solid interface, at least not in the
beginning of the nitrocarburizing. As a matter
of fact, such a conclusion also holds for the
case of nitriding (see Section 9.1). The occurrence of local equilibria within the nitrocarburized zone implies that the chemical potentials
of nitrogen and carbon (and iron) vary in a continuous manner through the nitrocarburized
zone, from the surface to the nonaffected core.
This can be associated with the formation of
phases in a sequence and of compositions as
prescribed by a so-called diffusion path at constant temperature in the (metastable) phase diagram (Ref 46, 47).
Thus, for the case of nitriding, pertaining to
the binary metastable iron-nitrogen phase diagram, the order of phases (from top to bottom
of the nitrided zone) would then be e/g0 /a, as
shown in Fig. 8 and 9 and discussed in Section
9.1 (provided the nitriding potential of the
nitriding atmosphere is high enough to allow
formation of e nitride). Draw a horizontal line
in Fig. 2 at the nitriding temperature and proceed from the right (corresponding to relatively
high chemical potential of nitrogen) to the left
(corresponding to relatively low chemical
potential of nitrogen); also see the discussion
in Sections 3 and 4.
To determine on this basis the sequence and
compositions of the carbonitride phases in the compound layer developing upon nitrocarburizing,
a diffusion path should be drawn in an isothermal
section of the Fe-C-N phase diagram for the
temperature at which the nitrocarburizing is
carried out. The diffusion path is defined by
the course of the (laterally, i.e., at constant
depth) averaged composition in the direction
perpendicular to the original interface of the
diffusion couple concerned—here, perpendicular to the surface of the specimen/component.
Not only do thermodynamics govern the
course of the diffusion path, but kinetics can
be decisive as well.
As follows by application of Gibbs’ phase rule,
at constant temperature and constant pressure, in
the case of the binary system iron-nitrogen, only
one phase can be present over a certain depth range
covering a certain variation in (the laterally averaged) composition (if local equilibrium prevails).
The presence of two phases in contact represents
a nonvariant situation; the two phases are in contact only at a specific depth, that is, at a specific
set of values for the chemical potentials of the
components iron and nitrogen, and at a specific
set of values of the compositions of both phases
in contact. This is compatible with Fig. 8 and 9 discussed in Section 9.1. Note that as a consequence,
the phases e and g0 occur as sublayers in that order,
from top to bottom, in the compound layer developing upon nitriding iron.
Similarly, it follows that in the compound
layer at constant temperature and constant pressure, in the case of the ternary system Fe-C-N,
one or, at most, two phases can be present over
a certain depth range covering a certain variation in the laterally averaged composition (if
local equilibrium prevails). The presence of
three phases in contact represents a nonvariant
situation; the three phases are in contact only
at a specific depth, that is, at a specific set of
values for the chemical potentials of the components iron, carbon, and nitrogen, and at a specific set of values of the compositions of the
three phases in contact. Hence, only the thermodynamics of the system can already induce
a complicated microstructure of the compound
layer in the case of nitrocarburizing. Note that
a sequence of sublayers (with each phase represented by a sublayer, as is the case for nitriding
of iron; see above discussion) is not expected in
the case of nitrocarburizing of iron, already
because two phases can be in contact over a
certain depth range, which does represent a
(local) equilibrium situation.
As an example, a schematic presentation of
the microstructural development, in particular
the phase constitution, of the compound layer
obtained upon nitrocarburizing a-Fe at 580 C
(853 K) is provided in Fig. 10 and described
subsequently (Ref 45). {The nitrocarburizing
conditions for these experiments involved a
630 / Nitriding and Nitrocarburizing of Steels
gas mixture composed of 15.44 vol% NH3,
57.59 vol% H2, 20 vol% CO, and 6.61 vol%
N2; the gas flow rate was 13.5 mm s1, as calculated for room temperature (Ref 45). If no
side reactions occur in the gas atmosphere
(but see next sentence), a nitriding potential
for this gas mixture can be given as rN =
0.35 atm1/2 (Eq 6). The carburizing potential
for this gas mixture would be infinite; however,
because of side reactions in the gas atmosphere,
an effective carburizing potential/chemical
potential of carbon operates (Ref 48).}
In the following discussion of the microstructural development, the phase constitution is presented as a sequence of sublayers containing
one or two phases (see previous discussion)
indicated by their symbols and separated by
the symbol “/”; thus, the notation e/e + g0
θ
indicates that at the surface, an e sublayer occurs
which is followed at some depth by a dual-phase
e + g0 sublayer (see stage 6 in Fig. 10), and that
underneath the last sublayer, the substrate occurs
but is not separately indicated.
At the start of nitrocarburizing, a single-phase
layer of cementite (y) forms at the surface
(stage 1). The e carbonitride phase develops subsequently at the layer/substrate interface, and a
ε/ε+γ ⬘
(1)
θ
ε
γ⬘
α-Fe + [N] + [C]
θ/ε
α-Fe + [N] + [C]
(6a)
ε
(2a)
θ
α-Fe + [N] + [C]
ε/ε+γ ⬘
ε
γ⬘
θ/θ+ε
ε
(2b)
θ
α-Fe + [N] + [C]
α-Fe + [N] + [C]
(6b)
θ+ε/θ/θ+ε
(3)
ε/γ ⬘
ε
θ
ε
α-Fe + [N] + [C]
ε/θ+ε/ε
γ⬘
α-Fe + [N] + [C]
θ
(4)
ε
(7)
α-Fe + [N] + [C]
γ⬘
ε
γ⬘
ε
(5)
α-Fe + [N] + [C]
(8)
α-Fe + [N] + [C]
Schematic illustration of the progressive microstructural stages of compound-layer formation and evolution upon nitrocarburizing a-Fe. The results pertain specifically to a
treatment temperature of approximately 580 C (853 K). The process starts with the formation of a carbon-rich phase (cementite) and proceeds in the direction of nitrogenrich phases along diffusion paths, as sketched for successive nitrocarburizing times in Fig. 11. The following successive stages have been indicated: (1) single-phase cementite (y)
layer, (2a) y/e double layer, (2b) y/y + e layer, (3) y + e/y/y + e layer, (4) e/y + e/e layer, (5) single-phase e layer, (6a) e with some g0 developing in regions close to the interface
with the substrate, (6b) e/e + g0 layer, (7) e/g0 double layer, and (8) single-phase g0 layer. Source: Ref 45
Fig. 10
Fundamentals of Nitriding and Nitrocarburizing / 631
y/e double layer (stage 2a) or a y/y + e double
layer (stage 2b) results. The amount of e phase
in the compound layer then increases strongly
by growth of e into the substrate and also in the
opposite direction, by conversion of y into e
(stages 3 and 4; that such y ! e conversion can
occur was shown for the first time in Ref 49). This
growth phase ends with a single-phase e (compound) layer (stage 5). Continued nitrocarburizing induces emergence of (the carbon-poor) g0
phase close to the interface with the substrate,
and thus, an e/e + g0 double layer is formed
(stages 6a and 6b), which is succeeded by an e/
g0 double layer (stage 7). Thereafter, the amount
of g0 phase in the compound layer increases until
a single-phase compound layer results (stage 8).
These results, as obtained in Ref 45, have
been presented here with the recognition that
580 C (853 K) is a nitrocarburizing temperature often applied in practice, and indeed,
microstructures as presented as stages 5 and 6
in Fig. 10 have frequently been observed in
practice after treatment times of, say, 2 to 4 h
at the temperature concerned (580 C, or
853 K). Stage 8 has been observed after a treatment time of 24 h, which is much longer than
treatment times applied in practice.
For all types of nitrocarburizing atmospheres
(including salt baths and plasmas), it cannot be
claimed that precisely the same sequence of
microstructures as presented in Fig. 10 occurs
in the compound layer (Ref 39). However, the
general conclusion from observations as well
as the previous discussion is that the microstructural evolution of the compound layer
developing upon nitrocarburizing proceeds
from a carbon-rich phase (cementite, y) into
the direction of carbon-poorer and nitrogenricher phases (e and g0 ).
The change (with treatment time) of the composition and the phase constitution at the surface
of the compound layer immediately makes clear
that neither a local equilibrium nor a stationary
state (see Section 8) occurs at the gas-solid interface, at least not for the largest part of the period
of time in the nitrocarburizing experiments pertaining to Fig. 10, implying that this certainly
holds for the nitrocarburizing treatments in (commercial) practice. The initial development of
cementite (not generally recognized, because this
cementite at the interface with the substrate
disappears in the subsequent stage of the
process) was first reported in Ref 43 for salt bath
nitrocarburizing and in Ref 39 for gaseous nitrocarburizing. This occurrence of cementite can
be discussed as follows.
The rate of carbon transfer from the nitrocarburizing medium, that is here from CO, is
much higher than that of nitrogen transfer, that
is, from NH3 (Ref 28). Further, the solubility of
carbon in the substrate (ferrite) is very small
and, in any case, much smaller than that of
nitrogen (Table 1). Then, recognizing that the
diffusion coefficients of carbon and nitrogen
in ferrite do not differ strikingly (Ref 50), it follows that the ferrite substrate at the surface is
saturated with carbon much faster than with
nitrogen, leading to the initial formation of
cementite (which has a negligible solubility for
nitrogen) at the surface of the substrate and not
of a nitrogen-rich(er) carbonitride that appears to
comply with local equilibrium at the surface or
with a stationary state at the surface (see the results
for long(er) treatment times as indicated in
Fig. 10). Upon continued treatment, the ferrite substrate at the substrate/compound-layer interface
becomes gradually enriched in nitrogen (by diffusion of nitrogen through the grain boundaries of
the cementite, Ref 51), and then the e phase can
nucleate there (Fig. 10, stage 2).
Whereas the composition and phase constitution at the surface are controlled by kinetics, it
can be shown that, given this constraint, the microstructure within the compound layer, dependent
on the composition and phase constitution at the
surface, is controlled by thermodynamics: local
equilibria prevail at the solid-solid phase interfaces within the compound layer. The way to make
this likely is by drawing diffusion paths (see the
beginning of this section) in the metastable ternary
Fe-C-N phase diagram (Ref 52) at the treatment
temperature (Fig. 11). For the various stages
shown in Fig. 10, suggestions for diffusion paths,
representing the change of the lateral average
composition and phase constitution as a function
of depth, have been made (schematically) in
Fig. 11 (Ref 45). That this is possible indicates that
the depth dependence of the lateral average composition and the thermodynamics of the Fe-C-N
system (fully) govern the microstructure within
the compound layer.
This section is concluded with a paragraph
about a special microstructural feature that can
be observed in particular if larger fractions of
ammonia occur in the gas atmosphere. In that
case, distinct porosity can develop in the compound layer (see Section 3), in particular in
the e (carbo)nitride phase/(sub)layer. Such a situation has not been considered in the previous
discussion and in Fig. 10. After these pores,
preferentially nucleated at grain boundaries,
have coalesced and induced channel formation
along these grain boundaries, they can have
direct contact with the outer nitrocarburizing
atmosphere. Then, a preferential uptake of carbon has been observed to take place via these
channel walls at some depth beneath the surface
(Fig. 12a), leading eventually to cementite
formation at the channel walls (Fig. 12b); even
a cementite sublayer has been observed to
develop subsequently (Ref 42). One can
Diffusion paths in the isothermal section of the Fe-C-N phase diagram at 580 C (853 K) for various stages of
compound-layer development upon nitrocarburizing a-Fe at 580 C (853 K) and as shown in Fig. 10. On this
basis, the time-dependent microstructural evolution of the compound layer can be illustrated for the case that local
equilibrium at solid/solid interfaces occurs, which appears to be the case for nitrocarburizing (and nitriding; see text).
A diffusion path represents, at a given time, the course of the lateral gross composition and the phase constitution,
going from the top to the bottom of the compound layer; see the solid lines indicated with arrows for those parts of
the diffusion paths where a continuous change of the gross composition as a function of depth occurs, and the
dashed lines for depths where a jump in the gross composition takes place. The time-dependent change of the gross
composition at the surface of the compound layer is represented by the dotted line indicated with an arrow. It should
be noted that the isothermal section of the Fe-C-N phase diagram shown here is that given in Ref 52, on the basis of
experimental data, which is incompatible with the calculated isothermal section according to the CALPHAD database
(2008). Source: Ref 45
Fig. 11
632 / Nitriding and Nitrocarburizing of Steels
Fig. 12
Carbon uptake through “open” grain boundaries/channels in contact with the outer nitrocarburizing
atmosphere. (a) Preferential carbon uptake is observed to take place via the channel walls at some
distance from the outer surface (in this case, nitrocarburizing took place in a CO-containing gas atmosphere: 3 vol%
CO; 53 vol% NH3, and 44 vol% H2 at 570 C, or 843 K). (b) The carbon enrichment in the e phase at these depths
can eventually lead to cementite formation. The associated transformation of the original e phase to cementite (y)
adjacent to the channel walls leads to fine pores (filled with N2 gas), because the cementite has a very low solubility
for nitrogen. Source: Ref 39
speculate about the origin of this phenomenon.
It can be suggested that the outer gas atmosphere can penetrate the channels, after these
have established an opening at the outer surface, and that at the fresh channel walls the
kinetics of carbon uptake may be much faster
than that of nitrogen uptake, as discussed earlier
for the formation of cementite at the surface of
the substrate at the start of nitrocarburizing.
The driving force for the decomposition of the
e phase, giving rise to pore/channel formation,
is largest close to the outer surface, because
the nitrogen content of the e phase is largest
there. Thus, the initial presence of a large
amount of Nads at the fresh channel walls may
be most pronounced at the channel walls close
to the outer surface, which thus could effectively block the occurrence of a gas-solid reaction (Eq 8a) at these depths at the channel walls
close to the outer surface.
9.3 Microstructural Development of the
Compound Layer in the Presence of Alloying
Elements. Steels subjected to a nitriding treatment often possess alloying elements having
distinct chemical affinity for nitrogen (see Section 10). The question emerges how the presence of these alloying elements influences the
formation of the compound layer. To provide
an answer, one must distinguish between a
group of alloying elements that has an apparently strong-to-intermediate interaction with
nitrogen (alloying elements belonging to this
group are titanium, vanadium, and chromium)
and a group of alloying elements that has an
apparently weak interaction with nitrogen
(alloying elements belonging to this group are
aluminum, molybdenum, and silicon). (For the
classification of strong, intermediate, and weak
Me-N interaction, which expresses a net outcome of the precipitation promoting change of
chemical Gibbs energy and the precipitation
obstructing misfit energy, see Section 5.4 in
Ref 53; see also Section 11.2 in this article).
Strong interaction implies that upon the start
of nitriding, an immediate precipitation of MeNn
nitride particles in the ferritic matrix is invoked.
Then, as soon as the amount of nitrogen dissolved in the ferrite matrix at the surface of the
specimen exceeds the solubility limit of nitrogen, a compact (largely) iron-nitride compound
layer develops at the surface of the specimen.
During growth of this iron-nitride compound
layer, the MeNn nitride particles, which already
have precipitated in the matrix (the diffusion
zone), are incorporated in the compound layer
by overrunning them (Ref 41, 54–56).
Underneath the iron-nitride compound layer,
iron-nitride developments along grain boundaries
of the matrix can often be observed (Fig. 13).
This phenomenon can be explained as follows:
Segregation at grain boundaries of the alloy-
ing element Me, already in the unnitrided
condition, leads to preferred precipitation
of Me-nitride at grain boundaries in the diffusion zone below the compound layer. As
a consequence, Me-depleted regions adjacent to such grain boundaries occur. The
prevailing supersaturation of nitrogen in
these regions then, in the absence of Me,
can lead to the development of iron nitride
adjacent to such grain boundaries.
The initially developed nanosized, largely
coherent MeN precipitates (see Section
11.2), in the so-called continuous precipitation (CP) region, may coarsen via a discontinuous coarsening reaction (Ref 14, 57),
which results in the development of a lamellar, discontinuously coarsened (DC) microstructure (Ref 58–60). The ferrite matrix
surrounding the continuous, nanosized,
largely coherent precipitates can contain
much more excess nitrogen (see Section 11.3)
than the ferrite matrix in the DC region.
Thus, the occurrence of DC is accompanied
by release of a huge amount of excess nitrogen that either locally enhances the nitrogen
supersaturation of ferrite, leading to the precipitation of nitride along grain boundaries
and also at the CP/DC interface, or associates at the grain boundaries under formation
of pores filled with N2 gas, coalescence of
which leads to the development of open
grain boundaries/cracks (refer to the discussion on pore formation in the iron-nitride
compound layer in Sections 3 and 9.1). The
penetration of the outer nitriding atmosphere
through the cracks opened to the surface
then leads to the development of iron nitride
along the crack faces (see discussion of
Fig. 12).
Weak interaction implies that, upon the start
of nitriding, not an immediate but a (very) slow
precipitation of MeNn nitride particles can
occur. Such a delayed precipitation reflects the
pronounced misfit strain field that is associated
with the development of MeNn precipitates. In
this case, a competition between the slow precipitation of MeNn and the development of iron
nitride can occur. Consider the case that nitriding takes place under conditions which allow
the development of g0 iron nitride. The solubility of Me in g0 iron nitride can be small, as, for
example, holds for aluminum and molybdenum.
Then, precipitation/development of g0 iron
nitride occurs either after a partitioning of Me
in the ferrite matrix, as by a preceding precipitation of Me as MeNn, has been realized, to
make possible the development of Me-free g0
iron nitride, or the g0 iron nitride is forced to
nucleate and grow with Me dissolved in it.
The difficulty in the precipitation of both MeNn
and g0 iron nitride allows the absorbed nitrogen
to diffuse deeply into the specimen, leading to a
(unusually) high nitrogen supersaturation of the
ferrite matrix. Eventually, such nitrogen supersaturation brings about the development of g0
iron nitride, with Me dissolved in it, across the
depth range of high-nitrogen supersaturation. As
a result, a peculiar, platelike morphology of g0 ,
Fundamentals of Nitriding and Nitrocarburizing / 633
iron-carbonitride layer upon nitrocarburizing is
kinetically controlled by the inward diffusion
of the interstitial components. Of course, this
does not hold for the beginning stage of layer
growth as long as no local equilibrium or stationary state has been realized at the surface
of the layer (see discussion in Section 8.1). This
has an important consequence for the nucleation of (carbo)nitrides at the surface of the
specimen/component, which is discussed first
for the case of nitriding pure iron and the development of an iron-nitride layer at the surface in
the following section.
The nitrogen-uptake rate is the outcome of
competing processes:
The dissociation process (Eq 2)
The recombination and desorption process
(Eq 17) at the gas-solid interface
The diffusion process in the solid substrate
(see Sections 8.1 and 9.1)
Fig. 13
Cross section of nitrided Fe-Me alloy specimen with a “strong” Me-N interaction (see text). Here, Me = V;
Fe-4at.%V specimen nitrided at 580 C (853 K) for 4 h with rN = 0.8 atm1/2 (light optical micrograph;
after etching in 2 vol% Nital). g0 nitride has formed not only as a layer at the surface but also along the “open” grain
boundaries (g.b.) in the matrix and along the interface between the zones of continuous precipitation (CP) (of VN)
and of discontinuous coarsening (DC). Source: Ref 56
Fig. 14
Cross sections of nitrided Fe-Me alloy specimens with a “weak” Me-N interaction (see text). (a) Fe-4.7at.%Al
specimen nitrided at 500 C (773 K) for 10 min with rN = 1.73 atm1/2. (b) Fe-1at.%Mo alloy specimen
nitrided at 480 C (753 K) for 2 h with rN = 0.7 atm1/2. Note the unusual plate-type morphology of the developed g0
iron nitride (light optical micrographs, cross sections etched in 2 vol% Nital). Source: Ref 61 and 62, respectively
deeply penetrating the specimen, occurs (for an
example, see Fig. 14) (Ref 56, 61, 62).
The results discussed in the preceding paragraph pertain to iron-base alloys, with Me as an
element of weak interaction with nitrogen, in
recrystallized condition. If these materials have
been deformed (e.g., by cold rolling), at the start
of nitriding immediate precipitation of MeNn can
occur (defect/dislocation-facilitated nucleation),
and a compact compound layer of (largely) iron
nitride develops at the surface that incorporates
the already precipitated MeNn particles (by overrunning them) (Ref 62). This is just as in the case
for Me as an alloying element experiencing a
strong interaction with nitrogen.
It is concluded that compound layers not
only can be avoided (by proper selection of
the nitriding potential, which should not exceed
a critical value), but that they can be microstructurally and morphologically modified by
incorporating in the ferritic matrix, in dissolved
state, alloying elements Me having a weak
Me-N interaction (Me = aluminum, molybdenum, and silicon) next to elements Me having
a strong Me-N interaction (Me = titanium,
vanadium, and chromium).
10. Kinetics of Compound-Layer
Growth
It is often taken for granted that growth of
an iron-nitride layer upon nitriding or an
If rate-constant and diffusion-coefficient
data are available, the joint result of these
processes can be calculated numerically (see
the Appendix of Ref 23). Hence, recognizing
that local equilibrium, or a stationary state, is
not established instantaneously at the surface
in the case of nitriding, the nitrogen concentration-depth profiles cannot be calculated
straightforwardly on the basis of (analytical
or numerical) solutions of only Fick’s second
law.
An example of the evolution of the nitrogen
concentration-depth profile with time at constant temperature for the case of nitriding
pure iron is shown in Fig. 6. Indeed, the surface
concentration of nitrogen gradually increases
to attain a constant value only after pronounced
nitriding. This behavior, for the normal nitriding conditions considered here, is dictated
by the competition between the dissociation
of NH3 at the surface and the inward
diffusion of nitrogen; the effect of the recombination of adsorbed nitrogen at the surface
and its subsequent desorption as nitrogen gas
is negligible at temperatures below 580 C
(853 K) and not very high nitriding potentials
(Fig. 7).
The effect shown in Fig. 6 has as consequence that, even if the nitriding potential and
temperature predict the development of, for
example, g0 iron nitride, then nucleation of this
nitride at the surface can only occur after distinct nitriding time has passed; g0 nucleation
can occur (at the earliest) at the moment
that the surface concentration of nitrogen
surpasses the value compatible with the a/g0
equilibrium. Hence, an incubation time for
iron-nitride formation occurs. This effect was
first recognized and quantitatively predicted
and experimentally verified in Ref 37. In the
calculations of Ref 37, the contribution of
the recombination and desorption process at
the surface was neglected. In a later work
(Ref 38), this effect was taken into account;
the difference in resulting incubation time for
634 / Nitriding and Nitrocarburizing of Steels
iron-nitride nucleation at the surface is only
a few percent at most, in accordance with the
previous discussion.
Only after a closed (carbo)nitride layer has
formed at the surface and a stationary state or
local equilibrium situation occurs at the surface
of the layer may interstitial diffusion processes
in the (carbo)nitride layer control the rate of
layer growth.
Surprisingly little information on interstitial
diffusion in (growing) iron-nitride layers on ferrite is available in the literature (Ref 63–66).
This holds even more so for (growing) ironcarbonitride layers on ferrite (Ref 67). Especially in the case of iron-carbonitride compound layers, the microstructure can be so
complicated, as exemplified by the presence of
(if local equilibrium prevails, at most) two
phases over a certain depth range in the layer
(see Section 9.2), that a more or less straightforward analysis of diffusion processes is
obstructed. Hence, researchers aiming at characterizing the kinetics of layer growth, on the
basis of parameters as diffusion coefficients,
look for geometries that are laterally invariable.
This is guaranteed, in principle, for the binary
iron-nitrogen system; if the two nitrides e and
g0 occur in the compound layer upon nitriding
an a-Fe substrate, then the e phase is present
as a sublayer on top of a g0 sublayer, on top
of the a-Fe substrate (see Fig. 8 and 9 and Section 9.1). This microstructure of planar, parallel
e and g0 sublayers is the one most frequently
investigated (apart from the case of a single g0
layer on top of an a-Fe substrate). For the ternary Fe-N-C system, such a dual-sublayer
microstructure is only possible for restricted
ranges of chemical potentials of nitrogen and
carbon in the nitrocarburizing atmosphere
and limited ranges of temperature and time
(and pressure); for example, see stage 7 in
Fig. 10. The only work performed until now
devoted to diffusion analysis of both nitrogen
and carbon upon nitrocarburizing a-Fe followed this strategy (Ref 67).
The diffusional flux of only nitrogen in
an iron-nitride layer growing on ferrite can
be characterized by a single diffusion coefficient: the intrinsic diffusion coefficient of nitrogen, DN. (Note that at the normal nitriding and
nitrocarburizing temperatures, the iron in the
system can be considered as immobile.) Thus,
the flux of nitrogen, JN, and the gradient of
the nitrogen concentration in the nitride layer,
dcN/dx, are related by DN according to Fick’s
first law:
JN ¼ DN dcN
dx
dcj
dck
Dkj (Eq 19)
Jk ¼ Jkk þ Jkj ¼ Dkk dx
dx
with k = N, j = C, and with k = C, j = N. So, instead
of only one intrinsic diffusion coefficient, now
four intrinsic diffusion coefficients are required
to describe the fluxes of nitrogen and carbon:
DNN, DNC, DCC, and DCN. The coefficient DNC
describes the contribution to the diffusional
transport of nitrogen due to the concentration gradient of carbon, and DCN describes the contribution to the diffusional transport of carbon due to
the concentration gradient of nitrogen. Each
intrinsic diffusion coefficient is given by the
self-diffusion coefficient of the diffusing component considered, Dk , multiplied by the so-called
thermodynamic factor, ykj (Ref 46):
Dkj ¼ Dk ykj
(Eq 20)
In principle, both the self-diffusion coefficients and the thermodynamic factors depend
on concentration. These thermodynamic factors
express the thermodynamic interaction of nitrogen and carbon, leading to the diffusional cross
effects as outlined in Eq 19. It appears that in
the case of simultaneous diffusion of nitrogen
and carbon in iron-carbonitride, these contributions, as exposed by DNC and DCN, can be very
distinct (see the following).
The growth of an e/g0 double layer on top of
a ferrite substrate during nitrocarburizing can
be described by considering the (mass) balances
that describe the shifts of the e/g0 interface and
the g0 /a interface, by infinitesimal distances dx
and dz, respectively, resulting from the diffusive fluxes arriving at and departing from the
interfaces (Fig. 15) (Ref 63, 67):
0
e=
0 =e
d þ Wk;e
Growth of e sublayer: ck ck
e=gas
0 =e
¼ Jk
Jk
dt
(Eq 21a)
0
=a
a= 0
Growth of 0 sublayer: ck ck
dð ÞþWk; 0
0
=e
a= 0
dt
¼ Jk Jk
(Eq 21b)
I=II
with ck as the concentration of component k
I=II
in phase/sublayer I at the I/II interface, Jk as
the flux of component k in phase/sublayer I at
the I/II interface, and Wk,I as the amount
of component k necessary to maintain the
concentration-depth profile of component k in
sublayer I. This set of equations (four in the
case of simultaneous diffusion of nitrogen and
carbon) reduces to one equation if g0 ironnitride layer growth is considered.
On the basis of Eq 21, the diffusion
coefficients are determined by 0making layergrowth
rate measurements (ve/g = dx/dt and
0
vg /a = d(z – x)/dt) and concentration-depth
profile measurements. Results presented in
the literature pertain to experiments where, for
the single g0 layer or for the e/g0 bilayer, and
at constant temperature, parabolic growth is
(Eq 18)
Such descriptions of diffusional mass transport become essentially more complicated if
more than one diffusing component must be
considered. Thus, for the case of simultaneous
diffusion of nitrogen and carbon in an ironcarbonitride layer, Fick’s first law must be
written as:
Fig. 15
Schematic concentration-depth profile of the interstitial component k in the compound layer of the I/II (here:
e/g0 ) double-layer morphology. The concentration profiles in both sublayers have been taken linearly. The
dark-gray area represents the amount of component k per unit area cross section (perpendicular to the surface of the
specimen) to be accumulated in sublayer II to realize a shift of the I/II interface by a distance dx into the sublayer II.
The light-gray area represents the amount of component k per unit area cross section (perpendicular to the surface of
the specimen) to be accumulated in sublayer II to realize a shift of the II/a interface by a distance dz. Source: Ref 63, 67
Fundamentals of Nitriding and Nitrocarburizing / 635
assumed for the (sub)layer thickness. Such parabolic growth is expected for constant surface
concentrations and constant interface concentrations and an initially fully saturated substrate
or an initially unsaturated substrate of infinite
thickness. Additional assumptions that have
been made in such analyses are, for example,
the supposition of linear concentration-depth
profiles in the sublayers and constant (i.e., independent of concentration) intrinsic or self-diffusion coefficients (Ref 63, 67).
In the case of growth of only a g0 iron-nitride
layer upon nitriding ferrite, adopting the
self-diffusion coefficient of nitrogen as independent of concentration, the experimentally
observed concave curvature of the nitrogen
concentration-depth profile in the massive g0
layer near the surface could be ascribed to the
concentration dependence of the thermodynamic factor (Eq 20) (Ref 68). If porosity
occurs near the surface of the g0 iron-nitride
layer, the penetration of the outward gas atmosphere along open grain boundaries (owing to
coalesced pores developing at the grain boundaries; see Sections 3 and 9) can contribute distinctly to such concave curvature of the
nitrogen concentration-depth profile (Ref 69).
A summary of available data on the diffusion
coefficients of nitrogen in the g0 and e ironnitride phases, at the typical nitriding temperatures, is provided in Ref 24. Data at considerably lower temperatures (360 to 400 C, 633
to 673 K) are presented in Ref 66.
The only work until now devoted to the
kinetic analysis of simultaneous diffusion of
nitrogen and carbon in e iron-carbonitride upon
nitrocarburizing ferrite is presented in Ref 67.
The window of experimental parameters (see
the introductory paragraphs of this section)
was such that an e/g0 bilayer occurred at the
surface of the nitrocarburized ferrite. The g0
phase can incorporate only little carbon
(Fig. 11); in the analysis, the g0 sublayer was
taken as stoichiometric Fe4N, the nitrogen and
carbon concentration profiles in the e sublayer
were taken linear (validated experimentally),
and the four intrinsic diffusion coefficients in
the e phase (see text following Eq 19) were
taken as independent of concentration. The
results obtained for DNN, DNC, DCC, and DCN
upon nitrocarburizing at 550 C (823 K) demonstrate that DNC is approximately as large as
DNN and that DCN equals approximately
¼DCC. Hence the “off-diagonal” diffusion coefficients, DNC and DCN, are as significant as the
“diagonal” ones, DNN and DCC. This implies
the occurrence of strong thermodynamic interaction of nitrogen and carbon, which are dissolved on the same sublattice of octahedral
interstitial sites as provided by the hexagonal
close-packed parent lattice of iron atoms
(Table 1). As an illustration, the nitrogen flux
contributions JNN and JNC and the carbon flux
contributions JCC and JCN (Eq 19) at the gas/e
interface are shown in Fig. 16; JNC can be as
large as JNN, and JCN can be as large as JCC.
11. Microstructural Development
of the Diffusion Zone
With respect to the development of the
diffusion zone, a distinction between nitriding
and nitrocarburizing need not be made; carbon,
as compared to nitrogen, does not dissolve
to a significant extent in the ferritic matrix
(Table 1). Carbon, as a species offered by a thermochemical process such as nitrocarburizing,
plays a pronounced role only in the development
of the compound layer (see Section 9.2). In other
words, the virtue of nitrocarburizing is restricted
to its effect on the development of the compound
layer of (largely) carbonitrides at the surface of
the specimen/component.
11.1 Iron Nitrides in Pure Iron and Carbon Steels. In the absence of alloying elements
with affinity for nitrogen, at the nitriding temperature no precipitation of nitrides can occur
in the diffusion zone, and hence, the absorbed
nitrogen stays in solid solution; the nitrogen
atoms reside in random distribution at octahedral interstices of the body-centered cubic
a-Fe parent lattice. Upon slow cooling after
nitriding at relatively high temperature, the g0
iron nitride, Fe4N1x, based on a face-centered
cubic iron sublattice with an ordered distribution of nitrogen on octahedral interstices
leading to a primitive cubic translation lattice,
can precipitate. Continued cooling, if the
supersaturation allows, can then lead to the precipitation of an intermediate nitride: a00 -Fe16N2,
which has a body-centered tetragonal iron
sublattice with an ordered distribution of nitrogen on octahedral interstices leading to a
body-centered tetragonal translation lattice
(Table 1; see Fig. 17). Alternatively, if the
specimen has been quenched, so that all nitrogen is still in solid solution (actually, effectively
only possible with specimens that are relatively
thin, such as foils), aging at room temperature
and at temperatures until approximately 150 to
160 C (423 to 433 K) leads to the development
of regions with a a00 -Fe16N2 crystal structure, followed at higher temperature by the emergence of
g0 -Fe4N1x. (Note that local enrichment plus
ordering of nitrogen atoms suffices for a00 no rearrangement of iron atoms is required, Ref 70; for a
Diffusive fluxes of nitrogen and carbon at the surface of the e/g0 compound layer for a nitrocarburizing time of 4 h at 550 C (823 K) as a function of the chemical potential
of carbon in the gaseous nitrocarburizing atmosphere (here represented as the carbon activity, see Sections 6 and 7) for a fixed chemical potential of nitrogen in the
gaseous nitrocarburizing atmosphere (here represented as the nitrogen activity; see Sections 4 and 7). The chemical potentials of nitrogen and carbon were controlled by the
method described in Section 7. The fluxes JNN and JNC represent the nitrogen fluxes carried by the nitrogen concentration gradient and the carbon concentration gradient,
respectively; similarly, the fluxes JCC and JCN represent the carbon fluxes carried by the carbon concentration gradient and the nitrogen concentration gradient, respectively (see
Eq 19). Source: Ref 67
Fig. 16
636 / Nitriding and Nitrocarburizing of Steels
Fig. 17
Cross section of the diffusion zone of a nitrided
a-Fe specimen. The specimen was powder
nitrided (see the “Introduction” of this article) in a box
(for 8 h at 545 C, or 818 K; Ref 3) and thereafter very
slowly cooled, leading to relatively coarse g0 and a00 iron
nitrides (the small and large precipitates visible in the
micrograph). (Normally, the a00 iron-nitride precipitates
are only visible by applying higher magnifications, as
provided by a transmission electron microscope.) Light
optical micrograph; oblique illumination, oil immersion,
after etching in 0.5 vol% Nital. Source: Ref 3
discussion of this process, see Ref 14, pages 421
to 422.)
11.2 Crystalline and Amorphous Alloying
Element Nitrides in Iron-Base Alloys. The
introduction into the matrix of alloying elements, Me, with (chemical) affinity for nitrogen
is intended to induce the precipitation of tiny,
possibly (semi)coherent, Me-nitride precipitates
in the diffusion zone, which pronouncedly
enhances the mechanical properties of the diffusion zone (as, for example, reflected by a
strongly improved fatigue resistance, Ref 71;
see also Ref 14, pages 576 to 580). The following
binary Fe-Me systems, with respect to Me-nitride
development upon nitriding, have been documented in the literature: iron-chromium (Ref 3,
58, 59, 72), iron-aluminum (Ref 73, 74), ironvanadium (Ref 60, 75), iron-titanium (Ref 76,
77), iron-molybdenum (Ref 78, 79), and iron-silicon (Ref 80–82). This list of systems is not
exhaustive, and the references given have been
restricted to those reporting largely on the microstructure developing upon Me-nitride precipitation. More references can be found in the
publications listed here, and specific ones are also
given later in this section.
The notions “strong” and “weak” interaction
of Me and N have already been touched upon
in Section 9.3. They can be discussed as follows.
The Me-N interaction in ferrite, with Me
and N as dissolved solutes, can be defined as
the ratio of energy gained (chemical Gibbs
energy) and the energy required (misfit-strain,
and interfacial, Gibbs energy) upon precipitation of MeNn nitride particles from the supersaturated ferritic Fe-Me-N matrix. For values
of such an interaction parameter for a number
of alloying elements that leads to a ranking of
these alloying elements regarding their strength
of interaction, see the examples in Table 6 in
Ref 53. This interaction parameter facilitates
the characterization of two extreme behaviors
of Me-nitride precipitation (Fig. 18) (Ref 83):
Fig. 18
Types of Me-N interaction as revealed by the emerging nitrogen concentration-depth profiles. The symbols
C, t, and z denote nitrogen concentration, nitriding time, and depth below the surface, respectively. Source:
Ref 83
Weak interaction: The nitride-precipitation
process progresses with the same rate at
every depth below the surface (for a foil of
finite thickness, see the following). A nitrogen gradient is virtually absent. This can be
formulated as follows: nitriding a foil of
finite thickness would first lead to nitrogen
saturation of the ferrite matrix throughout
the foil (i.e., homogeneous nitriding) before
the nitride precipitation occurs (with a rate
of nitrogen consumption distinctly slower
than the rate of nitrogen uptake by the specimen, such that the state of homogeneous
nitriding is effectively maintained).
Strong interaction: A surface-adjacent
region (case) develops, where all Me atoms
have precipitated as nitride. A sharp case/
core boundary occurs, and nitrogen in the
core is virtually absent.
The type of crystal structure of Me nitride
most frequently encountered is the NaCl-type
crystal structure, based on a face-centered cubic
translation lattice; this holds for TiN, VN, CrN
(see also the discussion in Appendix A of Ref
53), and also for the cubic (rock salt) AlN
(which can be favored over the equilibrium
hexagonal modification, of AlN, wurtzite;
Ref 84). The cubic Mo2N crystal structure can
be conceived as a NaCl-type crystal structure
with 50% vacancies on the nitrogen sublattice
(Ref 79). The lattice parameter of the unit cell
of these NaCl-type crystal structures for MeN
is closely equal to:
pffiffiffi
aaFe 2
where aa-Fe indicates the lattice parameter of
body-centered- cubic a-Fe. Thus a {100} habit
plane and an orientation relationship of the following type can be expected:
ð001ÞaFe ==ð001ÞMN ; ½100aFe ==½110MN
Obviously, three variants (one for each cube
plane of the matrix) of this (so-called Bain, also
called Baker-Nutting) orientation relationship
can occur. These predictions are in agreement
with the experimental observations. Then, a
coherent interface along the {100}a-Fe plane
can be expected, a linear misfit on the order of
a few percent. The linear misfit perpendicular
to the habit plane is very much larger, of the
order 40% and more. Hence, the nitrides
develop as tiny platelets, say, 10 nm long and
1 nm thick, depending on the precise nitriding
conditions (Ref 85). In agreement with this discussion, a high-resolution transmission electron
microscopy image shows the platelet faces to
be coherent, whereas (misfit) dislocations can
be detected at the platelet edges (Fig. 19).
If more than one alloying element is present,
such as Me1 = Cr and Me2 = Al (a combination
of alloying elements met in well-known nitriding steels), one may wonder whether
separate precipitation of Me1 nitride and Me2
nitride will occur or that a mixed nitride
(Me1)x(Me2)1xN will precipitate. It has been
recently shown that for Me1 = Cr and Me2 =
Al and for Me1 = Cr and Me2 = Ti, the mixed
nitride, with NaCl-type crystal structure, is preferred to precipitate (Ref 86–88). This can be
understood as follows (Ref 89):
For CrxAl1xN: The formation of the equi-
librium NaCl-type CrN precipitate is relatively fast, whereas development of the
equilibrium hexagonal AlN is relatively very
slow, owing to its large volume misfit with
the ferritic matrix. The misfit-strain energy
of the CrN precipitates can be reduced by
the uptake of aluminum. Also, because the
diffusion of chromium and aluminum in the
ferrite matrix is very slow compared to
the diffusion of nitrogen, the aluminum
atoms are “dragged” into the developing
cubic NaCl-type CrN precipitates. (Note that
the NaCl-type crystal structure is a possible
crystal structure for AlN; see previous discussion and Ref 84.) The system thus
accepts the gain of a smaller-than-maximum
Fundamentals of Nitriding and Nitrocarburizing / 637
amount of Gibbs energy, released by nitride
precipitation, as an intermediate solution;
CrxAl1xN precipitates develop.
For CrxTi1xN: The equilibrium precipitates
CrN and TiN have the same NaCl-type
crystal structure. The interaction parameter
Fig. 19
Vanadium-nitride precipitate (rocksalt-type crystal structure) in an a-Fe (body-centered cubic,
or bcc) matrix (high-resolution transmission electron
microscopy). At the top right corner, crystallographic
directions referring to the bcc lattice of the a-Fe matrix
are shown. The set of (110) lattice planes in the a-Fe matrix
continues as a set of (111) planes in the VN platelet,
as indicated by the black-white line contrast in the
micrograph, which traverses the matrix and the particle
(a thin platelet) in a continuous way; the interface between
the matrix and the faces of the nitride platelet is coherent.
The curvature of the lattice fringes is due to elastic
accommodation of the misfit between matrix and platelet.
Misfit dislocations occurring at the platelet extremities have
been indicated by arrows; these can be conceived as a
consequence of the misfit in directions perpendicular to the
platelet faces, as experienced at the platelet circumference,
being (very) much larger than parallel to the platelet faces
(Fe-2.2at.%V alloy nitrided for 25 h at 640 C, or 913 K).
Source: Ref 85
Fig. 20
(see previous discussion) for titanium-nitrogen is appreciably larger than that for
chromium-nitrogen. So, the driving force
for TiN to precipitate is much larger than
that for CrN to precipitate. The misfit-strain
energy of the TiN precipitates can be
reduced by the uptake of chromium. Also,
because diffusion of chromium and titanium
in the ferrite matrix is very slow compared
to the diffusion of nitrogen, the chromium
atoms are “dragged” into the developing
cubic NaCl-type TiN precipitates. The system thus accepts the gain of a smaller-thanmaximum amount of Gibbs energy, released
by nitride precipitation, as an intermediate
solution; CrxTi1xN precipitates develop.
Note the subtle distinction of contributing factors promoting the precipitation of a mixed
nitride, (Me1)x(Me2)1xN, for both systems in
the previous consideration.
The formation of the mixed nitride releases
a considerable amount of Gibbs energy. However, thermodynamically the precipitation of
the separate equilibrium nitrides is favored.
Indeed, it was shown that, after nitriding at
580 C (853 K) and by annealing at 700 C
(973 K), the metastable CrxAl1xN precipitates
in the diffusion zone become depleted of aluminum, followed by subsequent precipitation of
the released aluminum as hexagonal AlN in
the interior and at grain boundaries of the ferritic matrix in the diffusion zone (Ref 87).
A peculiar, interesting observation has been
made upon nitriding iron-base iron-silicon solid
solution. A distinct chemical driving force
exists for the formation of Si3N4 from a supersaturated Fe-Si-N solid solution. Yet, this
precipitation process is very slow, leading to
practically ideally weak nitriding kinetics
(Fig. 18, Ref 83). The very slow rate of nitride
precipitation is undoubtedly due to the very
large volume misfit of over 100% between
nitride precipitate and ferrite matrix. It was a
great surprise to observe that the nitride precipitates that eventually develop are not of crystalline nature but are amorphous (albeit of the
composition Si3N4) (Ref 80); it is extremely
rare for nature to favor the amorphous modification over the crystalline one for the precipitate in a solid-state precipitation process. For
small-sized precipitates associated with a relatively large interface/volume ratio, a relatively
low value of the energy of the interface
between the amorphous precipitate and the
crystalline ferrite matrix, compared to the interfacial energy in the case of the crystalline modification of the precipitate, can stabilize the
amorphous modification such that it is preferred
over the crystalline modification (Ref 80). The
amorphous precipitates developing at 580 C
(853 K) initially occur as bands along the ferrite
grain boundaries; at later stages, cuboidal amorphous nitride particles develop within the ferrite grains (Ref 81). The faces of the cuboidal
amorphous precipitates (Fig. 20) are parallel
to {100}a-Fe, suggesting that the interface
between amorphous Si3N4 and a-Fe preferably
forms along {100}a-Fe. At higher temperatures
(650 C, or 923 K), amorphous Si3N4 precipitates with a strangely eight-legged (octopodshaped) morphology occur (Fig. 20a, b) (Ref
82); the initially cubically shaped amorphous
particles, with a shape morphology dictated
by favorable interface energy (see previous
discussion), experience growth especially
along the <111> directions of the crystalline
ferrite matrix as a consequence of the very
large volume misfit (see earlier discussion)
and the elastically strongly anisotropic nature,
in particular at higher temperatures, of the
ferrite matrix.
Amorphous Si3N4 precipitates developing upon nitriding iron-silicon alloys. (a) At relatively low nitriding temperature, cubically shaped amorphous silicon nitride
precipitates develop; the cubical shape, with faces parallel to {100} planes of the a-Fe matrix, is governed by favorable interface energy (Fe-4.5at.%Si alloy specimen
nitrided at 600 C, or 873 K, for 40 h with rN = 0.02 atm1/2; scanning electron micrograph from cross section polished with colloidal silica solution OPS). (b) At higher
temperature, eight-legged (octapod-shaped) amorphous silicon nitride particles occur as a consequence of preferred growth along the <111> directions of the a-Fe matrix due to
the very large volume misfit and the anisotropic elasticity of the a-Fe matrix (Fe-4.5at.%Si alloy specimen nitrided at 650 C, or 923K, for 48 h with rN = 0.02 atm1/2, scanning
electron micrograph from jet-electropolished section). Source: Ref 81 and 82, respectively
638 / Nitriding and Nitrocarburizing of Steels
Another peculiar microstructural consequence of weak Me-N interaction is observed
upon nitriding iron-aluminum alloy under conditions where no compound (iron-nitride) layer
can develop. (For the developing microstructure
in case iron nitride can develop, see the description of weak interaction in Section 9.3.) Upon
nitriding of a specimen of Fe-4.65at.%Al, a
high density of microcracks appears along the
original grain boundaries in the ferrite matrix.
These microcracks arise by recombination of
nitrogen originally dissolved in the matrix. This
process happens because of the slow precipitation of hexagonal wurtzite-type AlN; the diffusion to the grain boundaries of the dissolved
nitrogen (followed by its recombination and
pore formation, which, by pore coalescence,
causes open, cracked grain boundaries) competes with the slow AlN precipitation. As a
consequence, AlN precipitates occur in the center of the grains, and precipitate-free zones are
present along the open, cracked grain boundaries. Upon continued nitriding, inward diffusion of nitrogen brings about a full nitriding
of the initially only partially nitrided grains.
Then, the microcracks become closed due to
the compressive stress that develops in the
nitrided zone—a remarkable process of selfhealing (Ref 90).
11.3 Types of Absorbed Nitrogen; Excess
Nitrogen. The amount of nitrogen taken up in
the nitrided zone, where (largely coherent)
alloying element nitride precipitates have
developed, can significantly exceed the amount
of nitrogen predicted, assuming that all Me has
precipitated as the expected MeNn nitride and
that the remaining ferrite matrix contains the
equilibrium amount of dissolved nitrogen. The
difference between the actual observed nitrogen
content and this expected value is called
“excess nitrogen.”
Detailed research has revealed that (at least)
three kinds of absorbed nitrogen can be distinguished (Fig. 21) (Ref 53, 91–96):
Type III: Nitrogen dissolved in octahedral
interstitial sites of (and throughout) the ferrite matrix (Fig. 21b). The misfit strain field
surrounding the nitride platelet of NaCl-type
crystal structure, in accordance with the
discussion in the previous section, is of
tetragonal nature, implying that the ferrite
surrounding the nitride platelet is severely
tetragonally distorted (Fig. 22). The elastic
straining of the surrounding ferrite matrix
is associated with a hydrostatic stress component of tensile nature that leads to a thermodynamically induced enhanced solubility
of nitrogen (Ref 53).
Thus, excess nitrogen is the sum of the adsorbed
nitrogen (type II nitrogen) and the surplus dissolved nitrogen (i.e., the actual amount of dissolved nitrogen minus the amount of dissolved
nitrogen in the absence of misfit stress). The
total amount of excess nitrogen taken up by the
specimen/component is by no means marginal;
type II excess nitrogen, adsorbed at the nitride/
matrix interfaces, can in practical cases be
50% of the amount of nitrogen necessary to precipitate all Me as MeN, of NaCl-type crystal
structure, and the amount of excess nitrogen
dissolved in the ferrite matrix can be of the
order of the equilibrium amount of dissolved
nitrogen. The occurrence of excess nitrogen
(thus) has a great impact on the nitriding kinetics
(see Section 12).
11.4 Nitriding Carbide-Containing Steels.
In physical metallurgy, the following rule of
thumb holds: oxides are more stable than
nitrides, which, in turn, are more stable than
carbides. Therefore, upon nitriding a quenched
and tempered steel containing alloying elements, with affinity for both carbon and nitrogen, the alloying element carbide particles,
which result from the quenching and tempering
treatment performed before nitriding, can be
Fig. 21
Schematic presentation of the three types of absorbed nitrogen. (a) Type I nitrogen is bonded to Me in the
MeN platelet (of NaCl-type crystal structure; a monolayer is shown for the case of a Bain, or BakerNutting, orientation relationship with the ferrite matrix; see text). Type II nitrogen is adsorbed at the interface between
the ferrite matrix and the MeN platelet in the octahedral interstices in the ferrite matrix in direct contact with the Me
atoms in the MeN platelet. (b) Type III nitrogen is the nitrogen dissolved in the ferrite matrix at octahedral interstitial
sites. Source: Ref 75
Type I: Nitrogen strongly bonded to the
nitride precipitates (Fig. 21). This nitrogen
cannot generally (easily) be removed by
denitriding in a reducing atmosphere (such
as pure H2 gas).
Type II: Nitrogen adsorbed at the (coherent)
interface of the nitride plate with the matrix.
For MeN of NaCl-type crystal structure, and
in view of the orientation relationship and
platelet morphology mentioned in Section
11.2, it can be anticipated that these
adsorbed nitrogen atoms reside in octahedral
interstitial sites of the surrounding ferrite
matrix opposite to the Me atoms in the
MeN platelet at the platelet/matrix interface
(Fig. 21a). For a monolayer MeN, this would
mean that the actual composition at the location of the nitride platelet can be indicated as
MeN3. The adsorbed nitrogen atoms are less
strongly bonded than the type I nitrogen
atoms and thus can generally be removed
by denitriding.
Fig. 22
Schematic view of a MeN platelet (of NaCl-type crystal structure for the case of a Bain, or Baker-Nutting,
orientation relationship with the ferrite matrix; see text) with its surrounding misfit-stress field in the ferrite
matrix. An expansion parallel to the platelet/matrix interface and a compression perpendicular to this interface are
induced by elastic accommodation of the precipitate/matrix misfit. Thus, the ferrite matrix surrounding the platelet is
tetragonally distorted. Source: Ref 75
Fundamentals of Nitriding and Nitrocarburizing / 639
replaced by nitride particles. This can occur by
transformation of the existing carbide particles
upon their reaction with nitrogen in the diffusion zone (Ref 97). This process takes place
relatively slowly and therefore occurs not only
at but also in the wake of the nitriding front
moving into the specimen/component. The
released carbon atoms can diffuse outward in
the direction of the compound layer, precipitate
as carbide (cementite) along grain boundaries
in the diffusion zone, and diffuse inward to the
unnitrided core, where a pronounced carbide
development can occur. All three effects have
been observed (Ref 97, 98). The micrograph in
Fig. 23(a) exhibits such precipitated carbides
along grain boundaries that run more or less parallel to the surface. This preferred orientation of
these carbides is caused by the presence of a
compressive residual stress parallel to the surface acting in the diffusion zone (see Section
11.18 in Ref 14). The nitrogen and carbon concentration-depth profiles shown in Fig. 23(b)
highlight the presence of carbides (at grain
boundaries) in the diffusion zone (indicated by
arrows and dashed lines in the figure) and the
development of a zone of carbon enrichment
underneath the (nitrogen) diffusion zone.
12. Kinetics of Diffusion-Zone
Growth
If no compound layer forms at the surface of the
specimen/component, the nitrogen concentrationdepth profile as it develops in pure ferrite (a-Fe)
Fig. 23
and in carbon steels not containing alloying elements with affinity for nitrogen is the outcome
of the following competing processes:
The dissociation process (Eq 2)
The recombination and desorption process at
the gas-solid interface (Eq 17)
The diffusion process in the solid substrate
(see Section 8.1 and the first paragraphs of
Section 10).
The nitrogen concentration-depth profiles
generally cannot be calculated straightforwardly on the basis of (analytical or numerical)
solutions of only Fick’s second law. Numerical
calculations incorporating the processes mentioned are required (Fig. 6).
In the following, it is assumed that a stationary state or local equilibrium has been closely
attained at the surface (if nitriding/nitrocarburizing occurs under conditions that do not allow
the formation of a compound layer) or that
local equilibrium has been realized at the interface of the compound layer and the diffusion
zone in the substrate. Further, it is taken for
granted that any diffusion and uptake of carbon
in the diffusion zone can be neglected, recognizing the very small solubility of carbon in ferrite (Table 1).
If the conditions indicated in the preceding
paragraph hold, the development of the diffusion zone in pure iron and carbon steels then
is rate controlled by the inward diffusion of
nitrogen in ferrite alone. This becomes different
for the diffusion-zone development in iron-base
alloys containing alloying elements with affinity for nitrogen. In that case, the kinetics of
the precipitation process of the MeNn nitrides
can have a strong impact on the nitriding kinetics and thus on the development of the nitrogen
concentration-depth profiles. These precipitation kinetics can be controlled by the nucleation
(activation energy of nucleation), the growth
(activation energy of growth; the growth may
be interface or diffusion controlled or of mixed
mode), and the impingement mechanisms (see
Chapter 9 in Ref 14). Thus, the inward diffusion of nitrogen into the ferrite matrix is only
one of a number of processes that together control the kinetics of nitriding. Current knowledge
does not provide an encompassing model taking
into account the full complexity of the nitriding
process in alloyed ferritic matrices.
Only for the case of strong Me-N interaction
(for the definition of strong and weak Me-N
interaction, see Section 11.2) can a simple relationship between the thickness of the nitrided
region and the nitriding time at constant temperature be given. To this end, and in the
absence of a developing compound layer, the
following assumptions are made:
The nitrogen dissolved in the ferrite matrix
exhibits Henrian behavior (Ref 20). This
implies that the diffusion coefficient of
nitrogen in the ferrite matrix, DN, is independent of the dissolved nitrogen content.
The reaction of dissolved nitrogen with dissolved Me, leading to the nitride MeNn
(or the mixed nitride (Me1)x(Me2)1xNn),
takes place only and completely at a sharp
interface between the nitrided zone and the
nonnitrided core.
(a) Cementite precipitates that developed along former austenite-grain boundaries more or less parallel to the surface upon nitriding a quenched and tempered steel (see
arrows in the micrograph; light optical micrograph of cross section of quenched and tempered 24CrMo13, or En14B, steel, salt bath nitrided at 580 C, or 853 K, for 2 h;
cross section after Murakami etching that stains carbides black). (b) Nitrogen and carbon concentration-depth profiles in the diffusion zone of a quenched and tempered 24CrMo 13
(En14B) steel (salt bath nitrided at 580 C, or 853 K, for 4 h) as determined by electron probe microanalysis. The presence of carbide (cementite) at grain boundaries is revealed by the
abrupt rise of the carbon content (see the arrows and dashed lines in the figure); also, the occurrence of a carbon-rich zone underneath the nitrogen diffusion zone is exposed. Source:
Ref 98; see also Fig. 16 in Ref 98
640 / Nitriding and Nitrocarburizing of Steels
for building up the concentration profile in
the ferrite matrix of the nitrided zone is negligible in comparison to the amount of
nitrogen that is consumed at the reaction
interface.
Diffusion of Me can be neglected and is not
nitriding-rate determining.
Local equilibrium prevails at the nitriding
medium/specimen interface, so that the surface concentration of dissolved nitrogen is
equal to the lattice solubility of nitrogen,
csN , as given by the chemical potential of
nitrogen in the nitriding atmosphere.
With these assumptions and approximating the
concentration
gradient of dissolved nitrogen with
csN z, where z is the depth coordinate of the
reaction front, the amount of nitrogen (per unit
area cross section perpendicular to the diffusion
direction/specimen-surface normal) that reaches
thereaction front
in the time period dt is equal
to csN DN =z dt. This nitrogen amount must
equal the nitrogen amount required to move the
reaction front a distance dz, that is, n cMe dz,
where cMe is the Me concentration. Upon integration of the resulting differential equation for constant temperature, the following parabolic
relationship for z and t is obtained:
z2 ¼ t s
2cN DN
n cMe
(Eq 22)
An equation of this type is well known and has
been applied before in the case of internal oxidation (Ref 99).
If a compound layer develops simultaneously, the treatment dealt with here may be
applicable as well on the basis of the following
consideration. The compound layer is thin
compared to the substrate, and the compound
layer grows much slower than the depth
range covered by the nitrogen concentrationdepth profile in the substrate. Then, the consumption of part of the substrate by the growth
of the compound layer may be neglected.
Then, also, in the presence of a growing compound layer, Eq 22 can be applied if strong
Me-N interaction occurs, with z = 0 as the position of the interface of compound layer and
diffusion zone.
The validity of Eq 22 can be verified with
two examples. According to Eq 5(b), the activity of nitrogen in the solid at the surface is proportional with the nitriding potential, assuming
that local equilibrium at the surface prevails.
In the ferrite substrate, Henry’s law holds for
the dissolved nitrogen (Ref 20). Hence, csN is
proportional with rN. Then, according to the
crude model represented by Eq 22, the depth
of the nitrided front, z, must be approximately
proportional with (rN)1/2, as observed experimentally (Fig. 24) (Ref 100). Also according
to Eq 22, the squared depth of the nitrided front
is proportional with t/cMe. The concentrationdepth profiles shown in Fig. 25 pertain to an
Fe-7wt%Cr alloy specimen and an Fe-20wt%
160
Experimental nitriding depth
Linear fit
120
Nitriding depth, μm
The amount of nitrogen that is required
80
40
0
0.0
0.1
0.2
0.3
0.4
0.5
(rN)½ (atm)-¼
Fig. 24
Nitriding depth (extent of the diffusion zone),
z, as a function of the square root of the
nitriding potential, rN1/2, for Fe-7wt%Cr alloy specimens
nitrided at 580 C (853 K) for 4 h. Source: Ref 100
Cr alloy specimen nitrided under the same conditions for 7 and 15 h, respectively. According
to the proportionality t/cMe, the squared nitriding depths for these specimens should have
the ratio 4/3, which not very well agrees with
the experimental result. The discrepancy is
ascribed to the difference in the concentration
of dissolved nitrogen at the surface, exhibiting
the effect of excess dissolved nitrogen depending on the amount of alloying element nitride
precipitate (see the caption of Fig. 25).
Still confining ourselves to the idealized case
of strong Me-N interaction, two distinct modifications with respect to the aforementioned
highly simplified model are necessary. First, it
appears unrealistic to assume that all Me is precipitated instantaneously upon the arrival of
nitrogen; the dissolved nitrogen concentration
does not drop from its saturation level at one
specific depth abruptly to zero. Instead, a depth
range at the nitriding front can be discerned
over which a more or less gradual change of
dissolved nitrogen to practically nil occurs.
Then, recognizing that a certain solubility
product for Me and N in equilibrium with
MeNn holds (Ref 102), it becomes clear that
not all Me is precipitated at once upon arrival
of the nitriding front. Only when, after some
time, the dissolved nitrogen has reached its saturation level, the maximum amount of Me has
precipitated; in the meantime, the nitriding
front has progressed further into the specimen/
component. Second, the various kinds of nitrogen in the specimen have different effects on
the nitriding kinetics (Ref 101). One must discern types I, II, and III of nitrogen as discussed
in Section 11.3 (see also Fig. 21). The nitrogen
taken up in the nitride platelets (obviously) and
also the excess nitrogen adsorbed at the faces of
the nitride platelets do not contribute to the
nitrogen diffusion process; nitrogen of types I
and II is immobile nitrogen. The nitrogen dissolved in the ferritic matrix can diffuse; nitrogen of type III is mobile nitrogen. (Note that
the amount of dissolved nitrogen can be a multiple of the equilibrium amount of dissolved
Fig. 25
Nitrogen concentration-depth profiles of
nitrided Fe-7wt%Cr alloy and Fe-20wt%Cr
alloy specimens nitrided for 7 and 15 h, respectively, at
580 C (853 K) with rN = 0.1 atm1/2. The experimental
data (points in the figure) were obtained by electron
probe microanalysis. The full lines through the data are
the results of fits of the model described in Section 12 to
the experimental data, with the following results for the
fitting parameters: csN = 0.35 at.%N, b = 1.18, and KCrN
= 0.02 nm6 for the Fe-7wt%Cr specimen; and csN =
0.26 at.%N, b = 1.176, and KCrN = 0.02 nm6 for the
Fe-20wt%Cr specimen. Source: Ref 101
nitrogen in pure ferrite; dissolved excess nitrogen occurs due to the misfit strain fields surrounding the nitride platelets.) These
considerations led to the following nitriding
model, which can only be evaluated numerically (Ref 101).
The inward diffusion of nitrogen in the
ferritic matrix can be described with Fick’s second law:
dcN ðz; tÞ
d 2 cN ðz; tÞ
¼ DN dt
dz
(Eq 23)
where cN(z,t) is the nitrogen dissolved in the
ferritic matrix at depth z, at time t, and at temperature T. The formation of nitrides MeNn
removes dissolved mobile nitrogen from the
ferritic matrix. This nitrogen then becomes
trapped as immobile nitrogen. The formation
of MeNn can be described as:
Me þ nN $ MeNn
(Eq 24)
where Me and N denote alloying element and
nitrogen dissolved in the a-Fe matrix. The equilibrium constant of this reaction, K, equals
1=KMeNn , with the solubility product KMeNn
given by:
KMeNn ¼ ½Me ½Nn
(Eq 25)
where [Me] and [N] denote the concentrations
of dissolved Me and dissolved N in the a-Fe
matrix. The precipitation of MeNn will take
place at a certain location if there it holds:
½Me ½Nn > KMeNn
(Eq 26)
In solving Fick’s second law (Eq 23), it must
be tested at every location (depth z) for every
Fundamentals of Nitriding and Nitrocarburizing / 641
time (step) if the solubility product, KMeNn , is
surpassed. If this is the case, (instantaneous
strong interaction) precipitation of MeNn, at
the location considered, should be allowed for
until [Me] [N]n = KMeNn . On this basis, a
numerical finite-difference (explicit method)
solution method can be developed to solve
Fick’s second law, subject to the prevailing
boundary conditions (Ref 101). The presence
of immobile excess nitrogen (nitrogen of type
II) can be accounted for by changing the stoichiometry of the nitride particles: MeNn
becomes MeNb, with b = n + x, where x denotes
the contribution of the immobile excess nitrogen. Note that x depends on the platelet thickness (for a monolayer of MeN (n = 1), x = 2;
see Section 11.3). The presence of mobile
excess nitrogen, that is, the amount of dissolved
nitrogen in excess of the equilibrium amount
for pure a-Fe, is accounted for by adopting a
model presented in Ref 53 (Ref 103).
The effects of mobile and immobile nitrogen
can be assessed considering the results of simulations shown in Fig. 26(a). If only the existence of mobile excess nitrogen is assumed, as
expressed by higher values for csN , a significantly larger extent (depth) of the nitrided zone
occurs as compared to the absence of mobile
excess nitrogen (see dashed line versus full line
in Fig. 26a). If only the existence of immobile
excess nitrogen is assumed, as expressed by a
value of b larger than n (= 1 for the present case
of CrN precipitation), a smaller penetration
depth of nitrogen occurs (see dotted line versus
full line in Fig. 26a). The used values for csN
and b are realistic values, as follows from
the experimental results presented in Ref 58,
92–96, 101, and 103. In view of the pronounced
influences of the immobile and mobile excess
nitrogen on the nitriding kinetics, it is imperative
Fig. 26
to incorporate the presence of excess nitrogen in
any model for the nitriding kinetics.
The role of the solubility product KMeNn is
illustrated in Fig. 26(b). The transition of the
nitrided zone to the unnitrided zone (i.e., the
reaction front) becomes less sharp as the solubility product increases. Relatively large
KMeNn values imply that (at the nitriding front)
not all dissolved N reacts instantaneously with
Me to MeNn, and thus, the extent of the nitrided
zone is larger for larger KMeNn , although in
association with a more gradual transition from
the nitrided zone to the nonnitrided core of the
specimen/component. Note that the amounts
of immobile and mobile excess nitrogen can
depend on nitriding time at constant temperature, because they depend on the extent of
MeNn precipitation and on the stage of aging
(and size) of the MeNn precipitate particles;
the solubility product, KMeNn , should not depend
on nitriding time at constant temperature.
Results of fitting the previously described
model to experimentally determined nitrogen
concentration-depth profiles are shown in
Fig. 25 for two iron-chromium alloys of different chromium content, nitrided under the same
conditions for different times, and in Fig. 27
for an iron-vanadium alloy nitrided at different
temperatures. The model, in a correspondingly
modified form, can also be applied to the nitriding of ternary Fe-Me1-Me2 alloy specimens
(Ref 104). In all these cases, the model provides
a satisfactory fit to the experimental data. For
interpretation of the values obtained for the fit
parameters, as, for example, the value of b
and, for example, as a function of the temperature, see Ref 53, 101, and 103.
Considering the case of weak Me-N interaction, considerable complication is added to
a possible model description of the nitriding
kinetics compared to the models presented
above for the case of strong Me-N interaction.
In this case, precipitation of MeNn does not
occur instantaneously once the solubility product is surpassed locally. Nucleation and growth
as thermally activated, time-dependent processes in a state of changing (with time and
place) supersaturation and (soft) impingement
as mechanisms controlling the MeNn precipitation process must be considered (Ref 14, 81,
83, 84, 105). This is already a difficult materials
science problem for supersaturated homogeneous
specimens. The nitriding process mandates that a
comprehensive model for the nitriding kinetics
must account for the simultaneously inward diffusion of nitrogen, leading, at constant temperature,
to time and location dependencies of the supersaturation, and thus, strongly locally different timedependent kinetics of the precipitation process
occur. A feasible approach may be to adopt a relatively simple model for the precipitation kinetics
(Johnson-Mehl-Avrami-Kolmogorov models are
popular but of limited ability to describe reality,
Ref 14) and to combine such a model with a
numerical solution for the inward diffusion process of nitrogen. This is an area of research where
no results of practical importance have already
been obtained.
Epilogue
The nitriding process and its variants offer
great challenges to the materials scientist and
the materials engineer. As an illustration, and
instead of offering a list of conclusions, the
author here refers to a few major themes of
fundamental importance for scientific understanding of nature that have been highlighted
in this review:
(a) Effect of mobile and immobile excess nitrogen on the development of the nitrogen concentration-depth profile. (b) Effect of the solubility product, KCrN, on
the development of the nitrogen concentration-depth profile. The examples shown in (a) and (b) pertain to a Fe-7wt%Cr alloy sheet nitrided at 580 C (853 K)
for 7 h with rN = 0.1 atm1/2. Source: Ref 101
642 / Nitriding and Nitrocarburizing of Steels
Fig. 27
Nitrogen concentration-depth profiles of nitrided Fe-2wt%V alloy specimens nitrided at rN = 0.103 atm1/2. The experimental data (points in the figure) were obtained by
electron probe microanalysis. The full lines through the data are the results of fits of the model described in Section 12 to the experimental data. (a) Nitrided at 520 C
(793 K) for 10 h. (b) Nitrided at 550 C (823 K) for 10 h. (c) Nitrided at 580 C (853 K) for 10 h. (d) Nitrided at 600 C (873 K) for 7 h. The horizontal lines indicated with “normal
nitrogen” in the figures indicate the nitrogen content taken up if only nitrogen incorporated in the MeN (here VN) precipitates (with all vanadium precipitated) and nitrogen dissolved
in the ferrite matrix according to the equilibrium state (unstressed) would occur; the real amounts of nitrogen taken up are distinctly larger, indicating the presence of considerable
amounts of excess nitrogen. Source: Ref 103
The interpretation of equilibrium has been
put into perspective in the sense that for
the iron-nitrogen system the definition of
equilibrium has been based on the NH3 and
H2 contents of the surrounding gas atmosphere and the sole occurrence of explicitly
defined, balanced forward and backward
reactions. More than 80 years after Lehrer,
we have come to recognize the effect of
the occurrence of so-called stationary states
as a consequence of the participation of
further reactions in the exchange of nitriding
medium and object. These stationary states
could be considered as (dynamic) quilibria
as well. However, the choice of a “reference” equilibrium (here, establishment of
equilibrium, Eq 2; see Section 4), which
can alternatively be described by establishment of equilibrium (Eq 15) together with
(infinitely fast) establishment of Eq 16 (Section 8.1), in fact necessitates the introduction
of the notion of stationary state, and thereby,
the experimentally determined Lehrer diagram does not partially represent genuine
“reference” equilibria.
Multicomponent diffusion in solids is a relatively rarely investigated phenomenon. It
plays a great role in the nitrocarburizing of
iron-base specimens/components. It strikes
to find out that the first, rigorous study
devoted to the simultaneous inward diffusion of nitrogen and carbon in a Fe-C-N
phase was published only in 2013 (see
Fundamentals of Nitriding and Nitrocarburizing / 643
Section 10). The very pronounced strength
of the nitrogen-carbon interaction in an FeN-C phase is expressed by the so-called
thermodynamic factors. Knowing these will
ultimately lead to an experiment-based
description of the thermodynamics of ternary Fe-C-N phases, and thereby, much controversy in the literature about the ternary
Fe-C-N phase diagram will eventually be
removed.
The role of the alloying element nitride/
matrix misfit stress in seriously influencing
the capacity for nitrogen uptake, that is, the
emergence of excess nitrogen, has in recent
years been shown to be a general phenomenon (Section 11.3). It has not been widely
recognized that both the immobile and the
mobile excess nitrogen greatly influence the
nitriding kinetics (Section 12). Any future
successful modeling of nitriding kinetics
for practical applications will have to incorporate the occurrence of immobile and
mobile excess nitrogen.
The chemistry of steels developed for nitriding purposes has been largely determined in
an empirical way. Research of fundamental
nature is beginning to unravel the effect and
interplay of (substitutionally dissolved) alloying elements on especially the kinetics of
nitriding. As an example of such a result of
potentially pronounced practical consequence, it has been suggested recently that
combining an alloying element of so-called
strong Me-N interaction (in particular, these
alloying elements have been considered until
now as useful for nitriding purposes) with an
alloying element of weak Me-N interaction
can lead to controllable, large microstructural
and morphological modification of the compound layer (Section 9.3).
It is true, but only seemingly so, that some of
the results and insights discussed in this article
may be of no practical consequence. This
would be a misleading conclusion. The materials engineer needs fundamental knowledge that
allows the development of models and equipment to control and tune the nitriding/nitrocarburizing process. Reconsidering the state-ofthe-art as described in a review paper published
in 1997 (Ref 24), it must be concluded that our
scientific understanding has made distinct progress during the ensuing fifteen or more years
and as reviewed in this article. However, we
are still far from tuned application of the nitriding/nitrocarburizing process, notwithstanding
the many developments in practice of the last
years, including new process variants that have
improved controllability/reproducibility. However, to predict (i.e., to understand) the outcome
of the nitriding/nitrocarburizing process in
practice is not possible, other than on the
basis of empiricism/experience. Unfortunately,
this situation is not much different from that
in 1997. As a salient example, it is mentioned
that the development of a nitriding sensor, as
compared to the description in Ref 24, has not
made much progress; control and knowledge
of the chemical potential of nitrogen at the surface of the component during nitriding in
technological practice is still no easy task.
Moreover, for the case of nitrocarburizing, a
method allowing simultaneous control and
knowledge of both the chemical potential of
nitrogen and the chemical potential of carbon,
an absolute requirement for controlled and
tuned nitrocarburizing, has been published only
very recently and, for the time being, has been
applied only in the laboratory (Section 7).
There is no doubt that, even after more
than 100 years of nitriding/nitrocarburizing,
future developments in materials science and
engineering will be instrumental in bringing
about distinctly deeper understanding of
the thermodynamics and kinetics of nitriding
and nitrocarburizing, in association with
the emergence of accompanying engineering
concepts.
ACKNOWLEDGMENTS
This article could not have been written were
it not for the long-term cooperation with Prof.
Marcel Somers (Technical University of Denmark, Lyngby) and Dr. Jan Slycke (formerly
SKF Engineering and Research Centre, Nieuwegein, The Netherlands; now retired) and
the intensive cooperation with the author’s coworkers, Dr. Andreas Leineweber, Dr. Sai
Ramudu (Sairam) Meka, and Dr. Ralf Schacherl.
Dr. Sairam Meka and Dipl.-Ing Bastian Rheingans read and discussed with the author a (part
of the) first draft of this article and suggested
valuable modifications. Dr. Ralf Schacherl
assisted in the preparation of all figures; Fig. 4
(c), 6, and 7 are essentially based on calculations
performed by Dr. Minsu Jung. The author’s former Ph.D. students contributed essentially to the
understanding attained that led to this article,
including Herman Rozendaal, Marcel Somers,
and Mohammad Biglari, who were involved in
the nitriding project in Delft; and Ralf Schacherl,
Tatiana Liapina, Santosh Hosmani, Nicolas
Vives Diaz, Thomas Gressmann, Marc Nikolussi, Arno Clauss, Sai Ramudu Meka, Kyung
Sub Jung, Thomas Woehrle, Holger Selg, Benjamin Schwarz, Matei Fonovic, Holger Goering,
Maryam Akhlaghi, and Tobias Steiner, who are
former and current Ph.D. students involved in
the nitriding project in Stuttgart.
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66. T. Liapina, A. Leineweber, and E.J. Mittemeijer, Phase Transformations in IronNitride Compound Layers upon LowTemperature Annealing: Diffusion Kinetics of Nitrogen in e- and g0 -Iron Nitrides,
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67. T. Woerhle, A. Leineweber, and E.J. Mittemeijer, Multicomponent Interstitial Diffusion in and Thermodynamic Characteristics
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68. T. Woehrle, A. Leineweber, and E.J. Mittemeijer, The Shape of Nitrogen
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F.D. Tichelaar, Amorphous Precipitates
in a Crystalline Matrix; Precipitation of
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Vol 41, 1999, p 625–630
S.R. Meka, K.S. Jung, E. Bischoff, and
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83. M.H. Biglari, C.M. Brakman, M.A.J.
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On the Internal Nitriding of Deformed
and Recrystallized Foils of Fe-2 at.% Al,
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84. M.H. Biglari, C.M. Brakman, E.J. Mittemeijer, and S. van der Zwaag, The Kinetics of the Internal Nitriding of Fe-2 at.%
Al Alloy, Metall. Mater. Trans. A, Vol
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85. T.C. Bor, A.T.W. Kempen, F.D. Tichelaar, E.J. Mittemeijer, and E. van der
Giessen, Diffraction-Contrast Analysis of
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Mag. A, Vol 82, 2002, p 971–1001
86. A.R. Clauss, E. Bischoff, S.S. Hosmani,
R.E. Schacherl, and E.J. Mittemeijer,
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88. K.S. Jung, R.E. Schacherl, E. Bischoff,
and E.J. Mittemeijer, Nitriding of Ferritic
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Vol 204, 2010, p 1942–1946
89. K.S. Jung, R.E. Schacherl, E. Bischoff,
and E.J. Mittemeijer, Normal and Excess
Nitrogen Uptake by Iron-Based Fe-Cr-Al
Alloys: The Role of the Cr/Al Atomic
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90. S. Meka, S.S. Hosmani, A.R. Clauss, and
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91. H.H. Podgurski and F.N. Davis, Thermochemistry and Nature of Nitrogen Absorption in Nitrogenated Fe-Ti Alloys, Acta
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92. M.H. Biglari, C.M. Brakman, E.J. Mittemeijer, and S. van der Zwaag, Analysis
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Different AlN-Precipitate Dimensions,
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93. S.S. Hosmani, R.E. Schacherl, and E.J. Mittemeijer, Nitrogen Uptake by an Fe-V Alloy:
Quantitative Analysis of Excess Nitrogen,
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94. S.S. Hosmani, R.E. Schacherl, L.
Litynska-Dobrzynska, and E.J. Mittemeijer, The Nitrogen-Absorption Isotherm
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646 / Nitriding and Nitrocarburizing of Steels
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97. P.C. van Wiggen, H.C.F. Rozendaal, and
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98. P.F. Colijn, E.J. Mittemeijer, and H.C.F.
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100. S.S. Hosmani, R.E. Schacherl, and E.J.
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102. Y. Sun and T. Bell, A Numerical Model
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104. K.S.
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105. J.S. Steenaert, M.H. Biglari, C.M.
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