48f1b39297

Nucleation Mechanisms in Ductile Iron
T. Skaland
Elkem ASA, Foundry Products, Kristiansand, Norway
Copyright © 2005 American Foundry Society
ABSTRACT
The present paper reviews different mechanisms for graphite nucleation in ductile iron (DI) and how these are affected by the
inoculation process. Theories describing the fundamentals of graphite formation are given and the strengths and weaknesses
of each theory discussed. Effects of key elements in the nucleation process, such as silicon (Si), calcium (Ca), strontium (Sr),
barium (Ba), aluminum (Al), magnesium (Mg), cerium (Ce), sulfur (S), oxygen (O) and nitrogen (N) are described and
discussed, and the importance of non-metallic heterogeneous compounds in the iron such as sulphides, oxides, nitrides and
silicates are considered. Studies of nucleation and growth of graphite are shown, and the complex interaction between the
Mg treatment and the inoculation process is described. The importance of the crystal structure and the stability of the nuclei
to become a potent site for graphite formation are reviewed, and examples of potent and non-potent nucleation sites are
shown. The paper arrives with a more comprehensive understanding of graphite nucleation in DI, and explains the key
difference between Mg and the other three elements Ca, Sr and Ba in this respect.
INTRODUCTION
Ductile irons (DIs) are iron-carbon-silicon alloys where the chemical composition is adjusted to ensure that carbon (C) will
precipitate as graphite spheroids during solidification. The C content is typically between 3-4% and the silicon (Si) content
between 2-3%, which gives a eutectic solidification temperature of about 1165ºC (2129ºF).
It is evident that one of the most important stages of the iron founding process is the economic production of liquid iron and
its metallurgical treatments in preparation for pouring into the mold. This involves maintaining compositional and
temperature control over the liquid during melting and holding in order to achieve the correct condition of the iron, the
correct graphitizing potential and the correct state of the nodularizing and inoculation processes in order to ensure a sound
casting of the desired structure and the required properties.
Magnesium (Mg) is the most common spheroidizing element used in the DI production, and it is usually added in
multicomponent alloy form with Si, calcium (Ca), rare earths, etc. Such alloys are balanced to reduce the reaction violence,
to promote graphite spheroidizing, to neutralize the effect of impurities on graphite morphology and to control the matrix
structure. The most common materials for nodularizing DI are ferrosilicon alloys containing about 45% Si, from 3-12% Mg
and various levels of Ca and rare earths (cerium [Ce], lanthanum [La], etc.).
Inoculation is a means of controlling the structure and properties of cast iron by minimizing undercooling and increasing the
number of graphite nucleation events during solidification. An inoculant is a material added to the liquid iron just prior to
casting that will provide a suitable phase for nucleation of graphite nodules during the subsequent cooling (Patterson, 1978).
Traditionally, inoculants have been based on graphite, ferrosilicon or calcium silicide. The most common inoculants today
are ferrosilicon based alloys containing small and controlled quantities of elements such as Ca, aluminum (Al), barium (Ba),
strontium (Sr) , zirconium (Zr), Ce, titanium (Ti), bismuth (Bi), etc. (Elliott, 1988).
HETEROGENEOUS NUCLEATION THEORY
Heterogeneous nucleation of graphite is an important aspect of cast iron metallurgy (Minkoff, 1983). The classic model for
heterogeneous nucleation is shown schematically in Fig. 1. Here the graphite phase (G) grows from the nucleant (N), and the
geometry of the graphite phase is a segment of a sphere of radius (r) and an angle of contact (θ). The interfacial energies
between the three phases graphite (G), nucleant (N), and liquid (L) are γGN, γGL, and γNL, respectively. The following
relationship exists between the interfacial energies:
NLGNGL γγθγ =+cos Equation 1
Proceedings of the AFS Cast Iron Inoculation Conference, September 29-30, 2005, Schaumburg, Illinois
13

Fig.1. This is a schematic representation of heterogeneous nucleation.
The change in free energy, ∆G, accompanying the formation of a graphite nucleus with this configuration is given by:
NLGNGNGNGLGLVG AAAGVG γγγ −++∆−=∆ ( ) ⎟
⎠
⎞
⎜
⎝
⎛
+∆−= GLV rGrf γππθ 23
4
3
4
Equation 2
where VG is the volume of solid graphite, ∆GV is the free energy of graphite formation, AGL and AGN are the area of the
graphite-liquid and graphite-nucleant interfaces, respectively, and f (θ) is the so-called shape-factor, defined as:
( ) ( )( )
4
cos1cos2 2
θθ
θ
−+
=f Equation 3
The critical radius of the stable nucleus, r*, is found by differentiating equation 2 with respect to r and equating to zero:
V
GL
G
r
∆
−=
θγ sin2
* Equation 4
The corresponding value of the critical free energy barrier, ∆G*, is then given by:
( )
( )
( )θθ
πγ
f
T
C
f
G
G
V
GL
2
1
2
3
3
16
*
∆
=
∆
=∆ Equation 5
where ∆T is the undercooling, and C1 is a kinetic constant which is characteristic of the system under consideration.
When θ = 0 the graphite nucleus will completely wet the substrate, which implies that there is no energy barrier to
nucleation. The nucleation rate
•
N (the number of graphite nuclei formed per unit time and volume) is, in turn, interrelated to
∆G* through the following equation (Elliott, 1988):
( )
⎥
⎦
⎤
⎢
⎣
⎡ ∆+∆
−=
•
kT
GG
NN D
V
*
expν Equation 6
where v is a frequency factor, NV is the total number of heterogeneous nucleation sites per unit volume, and ∆GD is the
activation energy for diffusion of atoms across the interface of the nucleus. Since ∆GD is negligible compared with ∆G* in
liquids, the nucleation rate of graphite is determined by ∆G*.
The value of ∆G* (or ∆T) depends, in turn, on the crystallographic disregistry between the substrate and the nucleated solid.
The disregistry can be defined as δ = (∆a0/a0) where ∆a0 is the difference between the lattice parameter of the substrate and
the nucleated solid for a low-index plane, and a0 is the lattice parameter for the nucleated phase. A mean factor representing
planar lattice disregistry can be calculated as follows (Bramfitt, 1970):
( ) 100
3
% 321
×
++
=
δδδ
δ Equation 7
where δ1, δ2, and δ3 are the disregistries calculated along the three lowest-index directions within a 90º quadrant of the planes
of the nucleated solid and the substrate.
In practice, the undercooling, ∆T, increases in a parabolic manner with increasing values of the planar lattice disregistry (δ)
(as shown in Fig. 2) (Turnbull, 1952). Since the undercooling during solidification of DI varies typically from 2-10°C ([36-
50ºF] depending on the section size) the results in Fig. 2 suggest that planar lattice disregistry between the inoculant and the
graphite is in the order of 3-10%. (Minkoff, 1983). Such low values are characteristic of coherent/semi-coherent interfaces.
14

Fig. 2. The characteristic undercooling vs. planar lattice disregistryi is graphed. (From D. Turnbull and R.
Vonnegut, Industrial Engineering Chemistry, vol 44, 1952)
Some scientists have proposed that graphite formation in cast iron may be resulting from homogeneous nucleation. Turnbull,
however, investigated the magnitude of undercooling, ∆T, required for homogeneous nucleation by means of small droplet
experiments (Turnball, 1952). The magnitude of required undercooling is found to be about 20% of the melting temperature
before homogeneous nucleation occurs. This means that undercooling in excess of 250°C (482ºF) would be required in DI
for homogeneous nucleation of graphite to be initiated. Consequently, homogeneous nucleation is rarely encountered in the
solidification processing of such liquid irons. If homogeneous nucleation occurs in cast iron, this would happen anyway at
undercoolings well below the metastable iron-carbide equilibrium temperature, thus resulting in fully carbidic
microstructures.
In conventional DI production, there will always be a number of non-metallic inclusions present in the treated (deoxidized
and desulphurized) liquid iron as dispersed heterogenities throughout the metal volume. As described above, heterogeneous
nucleation sites having the best planar lattice fit to graphite nucleate at only a very few degrees undercooling. Even
heterogenities having a very poor crystallographic fit to graphite as well as non-crystalline (amorphous) heterogenities,
eventually act as nucleation sites according to the Bramfitt model (Bramfitt, 1970). This occurs anyway at some 30 to 50
degrees undercooling at the maximum for the worst possible mismatch between the graphite and the heterogeneity.
THEORIES FOR GRAPHITE NUCLEATION MECHANISMS
Traditionally, cast iron inoculants are based on ferrosilicon, graphite or calcium silicide, the former being the most common
(Patterson, 1978; Hughes, 1980). Since pure Si and ferrosilicon are found to be ineffective as inoculants, their nucleation
potency depends on the presence of minor elements such as Ca, Al, Zr, Ba, Sr, Ti, etc. in the alloys (Dawson, 1961; Dawson,
1966; Kanetkar, 1984; Lownie, 1963; McClure, 1957; Mickelson, 1967). At present, the role of these minor elements are
partly understood, but still complex matters related to formation of different types of nucleation sites in DI remains to be
understood completely.
Several theories that exist in the literature explain the phenomena of heterogeneous nucleation of graphite in solidifying cast
iron. In the following, some of the most established theories are described and discussed.
THE GAS BUBBLE THEORY
According to Karsay, graphite tends to crystallize onto any given surface or imperfections such as cracks, pinholes,
inclusions, etc. (Karsay, 1976). The gas bubble theory states that graphite can form only if its crystallization is protected by
the presence of some sort of phase boundary. The needed phase boundaries are provided by the presence of carbon
monoxide bubbles in the melt. The carbon monoxide bubbles are very finely dispersed in the melt and their size is less than
10 µm. Karsay presented the gas bubble theory as illustrated in Fig. 3.
Karsay’s gas bubble theory is in principle based on the presence of carbon monoxide bubbles (Karsay,1976). However, in
industrial DI heats, strong deoxidizers, such as Mg, rare earths (RE), Ca, etc., are added that will effectively tie up and
neutralize any oxygen (O) in the form of dissolved O or as CO gas. Various gases such as hydrogen (H), nitrogen (N), and
CO are however found in DI castings as internal defects and voids. These are often covered on the inside by graphite linings.
However it is highly unlikely that a complete graphite nodule will extend into the entire volume of a gas bubble, since this
eventually would have to involve diffusion of C through the graphite shell. Under normal conditions, there should be no
driving force for C diffusion through solid graphite. Partly solidified and quenched irons should then also reveal partly filled
gas bubbles, which normally would never be observed in DI under any circumstances.
15

Fig. 3. Karsay’s gas bubble theory is illustrated—(A) gas bubble,(B) graphite, (C) melt and (D)austenite (From S. I.
Karsay, Ductile Iron I: Production, Quebec Iron & Titanium Corp.,1976).
THE GRAPHITE THEORY
The early theories for heterogeneous graphite nucleation are based on the assumption that the graphite nucleation occurred
epitaxially from other graphite particles contained in the iron melt (Boyles, 1947). Eash extended these ideas to Si-based
inoculants by proposing that their effectiveness is due to the formation of Si-rich regions around the dissolving particles
within which the solubility of C is sufficiently reduced to promote graphite precipitation (Eash, 1941). Later, Feest showed
that this assumption is not correct, since the dissolution time of ferrosilicon in liquid iron is just a matter of seconds, and that
graphite tends to form at the interface between the dissolving particle and the liquid (Feest, 1983). They therefore modified
Eash’s model by proposing that these seed crystals will be preserved in the melt down to the eutectic temperature, provided
that Sr or Ba is present in sufficient amounts to prevent redissolution of the graphite ((Eash, 1941; Kayama, 1979).
One weakness of the graphite theory and the assumption of small crystalline graphite particles, being preserved in the liquid
iron for extended times, is the conflict with the well established fact that graphite in the form of crystalline recarburizers
readily dissolves in liquid iron. Graphite recarburizers are typically added in sizes of millimeters, and will dissolve within
seconds or a few minutes. Graphite based nucleation sites in a solidifying iron would be in the sizes of microns, and their
dissolution time would consequently be very short. There is no question that graphite would be the ideal nucleation site for
graphite itself. However, it can be argued whether the thermodynamic stability of micron sized graphite particles above the
liquidus temperature would withstand its own dissolution characteristics for the entire fading time of inoculation.
THE SILICON CARBIDE THEORY
Following the dissolution of ferrosilicon in liquid iron, Wang and Fredriksson observed that silicon carbide crystals and
graphite particles are formed in the melt close to the dissolving ferrosilicon particles (Wang, 1981; Fredriksson, 1984). They
also observed that these transient particles redissolve readily after the inoculation treatment. No oxide or sulphide particles
are detected. Based on their experimental observations, a theory developed and calculations were performed in order to
explain the nucleation of graphite and the fading mechanism. A salient assumption in Wang and Fredriksson’s model is the
existence of an inhomogeneous distribution (local supersaturation) of C and Si in the melt subsequent to the SiC dissolution
which provides the necessary driving force for homogeneous nucleation of graphite (Wang, 1981; Fredriksson, 1984). The
fading effect is thus explained by a homogenization of the melt with respect to Si and C through convection and diffusion.
One weakness of the SiC theory for graphite nucleation is that the recognized critical role of elements like Ca, Sr and Ba in
the FeSi inoculant cannot be explained by this theory. Another weakness of the SiC theory is the assumption of local
supersaturation of C and Si due to restricted convection and diffusion. Both C and Si are recognized for having very high
diffusivity in liquid iron, and heat convection in hot metal is also recognized for being quite significant. It is therefore
unlikely that dissolving SiC particles in liquid iron would be capable of maintaining a supersaturation of C and Si throughout
liquid metal processing and into the solidification. Furthermore, the observation of SiC and graphite, surrounding a partly
dissolved FeSi inoculant particle, is most likely resulting from the experimental quenching technique itself, forcing transient
SiC and C out of solution in the Si-rich metal during quenching.
16

THE SALT-LIKE CARBIDE THEORY
In a classical paper on the nature of the graphite nuclei, Lux considers both homogeneous and heterogeneous nucleation of
graphite (Lux, 1964). He concludes that the elements Ca, Sr and Ba form salt-like carbides of the CaC2 type in liquid iron,
and that a direct epitaxial transition from the CaC2-lattice to the graphite lattice is possible without major changes in the
lattice dimensions. Under such conditions, the interfacial energy between the nucleus and the substrate is sufficiently low to
allow for extensive graphite nucleation at small undercoolings during solidification. The concept of the salt-like carbide
nucleation theory by Lux is illustrated in Fig. 4 (Lux, 1964).
Fig. 4. Epitaxial growth of graphite on a CaC2-crystal is illustrated.(From B. Lux, Modern Casting, vol 45, 1964).
However, particles of CaC2 have never been observed in the microstructure of inoculated DI. The thermodynamic stability of
CaC2 crystals as heterogeneous substrates having to survive in sulphur (S) and O containing liquid iron throughout holding
and pouring is also highly questionable. In competition with available S and O in commercial irons, it is unlikely that the
inoculating active elements, such as Ca, Sr and Ba, prefer to combine with C, forming such salt-like carbides. The sulphides
and oxides of these elements are significantly more stable and thus more favorable than forming compounds with C. The
salt-like carbide theory still offers an interesting approach from a crystallography point of view. It also attempts to give an
explanation to the important role of Ca, Sr and Ba in the inoculation process. The theory is however questionable from a
thermodynamic standpoint.
THE SULPHIDE/OXIDE THEORY
Several investigators have suggested that the graphite nucleation occur on sulphide, oxide or nitride particles, which are
formed after the addition of the inoculant (Gadd, 1984; Jacobs, 1974; Muzumdar, 1972; Muzumdar, 1973; Naro, 1970; Sun,
1983). Lalich and Hitchings confirmed this hypothesis by demonstrating the importance of non-metallic inclusions (Lalich
and Hitchings, 1976). They found that compounds of magnesium calcium sulphide act as nucleation sites for graphite
nodules in DI treated with Mg ferrosilicon alloys. They concluded that the majority of nodules in DI are associated with non-
metallic inclusions and that graphite growth in some instances is also related to the shape and distribution of these inclusions.
Inclusions in graphite nodules extracted from cast iron have been investigated by different techniques in order to determine
the identity of the catalyst particles. These techniques include both electron diffraction pattern analysis and X-ray
microanalysis (Deuchler, 1962; Rosenstiel, 1964; Zeedijk, 1965). The investigation by Jacobs is directed to determine the
nature of nuclei and detect possible changes in their chemical composition and crystal structure after treatment of iron with
Mg ferrosilicon (Jacobs, 1974). The subsequent inoculation treatment included the use of commercial Sr-FeSi alloy.
Different series are carried out in order to clarify the effects of elements such as Al and Sr on the inclusion characteristics.
These results are interesting for cast iron in general, since the examination revealed evidence of a duplex substrate structure
consisting of a sulphide core surrounded by an oxide shell. The different constituent phases are found —(Ca,Mg)- and
(Sr,Ca,Mg)-sulphides in the core, and (Mg,Al,Si,Ti)-oxides in the outer shell. Moreover, Jacobs observed that inclusions
embedded in the iron matrix contained the same constituent elements as those detected in the nodule centers, and that the
typical size of the particles is about 1 µm (Jacobs, 1974).
THE SILICATE THEORY
In an investigation of the inoculation mechanisms in DI, Skaland put particular emphasis on the aspects of heterogeneous
nucleation of graphite at inclusions (Skaland, 1993). It showed that the majority of the inclusions in ductile cast iron are
primary or secondary products of the Mg treatment (e.g. MgS, CaS, MgO⋅SiO2, and 2MgO⋅SiO2). After inoculation with
(X,Al)-containing ferrosilicon (X denotes Ca, Sr or Ba), hexagonal silicate phases of the XO⋅SiO2 or the XO⋅Al2O3⋅2SiO2
type form at the surface of the oxide inclusions, probably through an exchange reaction with MgO. The presence of these
phases enhances the nucleation potency of the inclusions with respect to graphite. In particular, the (001) basal planes of the
17

crystals are favorable sites for graphite nucleation, since these facets allow for the development of coherent/semi-coherent
low-energy interfaces between the substrate and the nucleus.
Figure 5a shows a TEM (Transmission Electron Microscope) examination of a silicate nucleus in DI. Figure 5b shows X-ray
mapping images of the Mg, Ca, Al and Si distribution in an inclusion, while Fig. 5c shows a schematic representation of a
heterogeneous nucleation site for graphite in DI.
(a) (b) (c)
Fig. 5. These depict—(a) TEM examination of silicate nucleus in DI, (b) STEM X-ray images showing the
distribution of Mg, Ca, Al and Si in inclusions after inoculation with a (Ca,Al) containing ferrosilicon and (c)
schematic illustration of heterogeneous nucleation site in D. (From T. Skaland, Metallurgical Transactions A,
vol 24A, 1993).
Skaland also gives a theory for the fading mechanisms of inoculation (Skaland, 1993). This is explained by a general
coarsening of the inclusion population with time, which reduces the total number of catalyst particles for graphite in the melt.
A theoretical analysis of the reaction kinetics gives results which are in close agreement with experimental observations.
NATURE OF NON-METALLIC INCLUSIONS
Non-metallic inclusions of varying composition have been observed in the iron matrix and at the centers of graphite nodules
by a number of investigators. Table 1 gives a summary of different element combinations and phases detected in DI
inclusions.
In the periodic table of elements, the group IIA-elements Mg, Ca, Sr and Ba are of specific interest in DI production, since
they are all strong sulphide and oxide formers and are typically added deliberately through ferroalloys. In the following,
possible reactions between these elements and C, Si, S, O and N are discussed.
SULPHIDES
The pure sulphides of the group IIA-elements are all of the face center cubic NaCl-structure type, and are characterized by
similar lattice parameters and high melting points. In cast iron melts these sulphides are among the most stable non-metallic
compounds. Hence, sulphides should form in preference to oxides. This conclusion is in close agreement with the results of
Jacobs who found that the inclusions consisted of a sulphide core surrounded by an oxide shell (Jacobs, 1974). Sulphides of
the group IIA-elements are also found to be a vital ingredient in the nucleus of graphite nodules by several researchers, as
shown in Table 1. Table 2 gives a summary of crystal structures, melting points and standard free energies of formation for
the group IIA-sulphides.
From the numerous literature sources, there should be no question that sulphides of the group IIA-elements do exist in the
core of graphite nodules in DI. Several scanning electron microscope (SEM) investigations have revealed in particular the
presence of Mg and S in the core of graphite nodules. It is therefore reasonable to expect that MgS and also CaS and other
sulphides are important ingredients in the heterogeneous nucleation sites for graphite. It is also recognized in the foundry
industry that the addition of Mg to cast iron is contributing to desulphurizing and subsequently to the growth of nodular
graphite morphologies. Table 2 shows that MgS and the other sulphides of the group IIA-elements, i.e. CaS, SrS, and BaS,
have very similar crystal structures, lattice parameters and stability. Mg additions to cast iron are however not recognized for
contributing to the nucleation of graphite nodules, while the other three elements Ca, Sr and Ba are very well recognized for
being nucleating elements when added through ferrosilicon inoculants. Because of the physical similarities between the
group IIA-sulphides, it is therefore unlikely that the existence of sulphides alone can explain the remarkable difference
between Mg and the other three elements in the graphite nucleation process.
18

Table 1. Summary of Element Combinations and Phases Detected in Inclusions
Phase Literature Reference
MgS
CaS
SrS
CeS
LaS
Askeland (1969), Askeland (1970), Jacobs (1976), Lalich (1976),
Mercier (1969), Warrick (1966),
Jacobs (1976), Lalich (1976), Mercier (1969)
Jacobs (1976)
Warrick (1966)
Warrick (1966)
MgO
SiO2
MgO⋅SiO2
2MgO⋅SiO2
(Mg,Al)3O4
(Mg,Al)SiO3
(Mg,Ca,Al)SiO3
CaO⋅Al2O3⋅2SiO2
Fe2O3
Fe2SiO4
Mg-Al-Si-Ti-O
CeO2
Askeland (1969), Askeland (1970), Askeland (1972), Francis
(1979), Heine (1966)
Askeland (1970), Askeland (1972), Heine (1966), Heine (1966)
Askeland (1970), Askeland (1972), Skaland (1993)
Askeland (1969), Askeland (1970), Askeland (1972), Skaland
(1993)
Latona (1984)
Latona (1984)
Latona (1984)
Skaland (1993)
Askeland (1972), Francis (1979)
Trojan (1968)
Jacobs (1976)
Francis (1979)
MgSiN2
Mg3N2
Mg2.5AlSi2.5N6
Mercier (1969)
Wittmoser (1952)
Igarishi (1998), Solberg (2001)
Mg3P2 Wittmoser (1952)
Table 2. Summary of Crystal Structures, Melting Points, and Standard Free Energies of Formation for the
Group IIA-Sulphides (From I. Barin, Thermochemical Data of Pure Substances, VCH Verlagsgesellshaft, 1989)
Phase Space
group
Crystal
system
Lattice
parameter
(Å)
Melting
point, TM
(°C)
Free
energy,
∆GF
*)
MgS Fm3m cubic 5.191 2000 -232
CaS Fm3m cubic 5.696 2450 -380
SrS Fm3m cubic 6.020 2000 -370
BaS Fm3m cubic 6.386 2227 -356
*) Standard free energy of formation at 1327C (1600K).
CARBIDES
The carbides CaC2, SrC2 and BaC2 also reveal the NaCl-structure type and have similar lattice parameters. They are
supposed to be metastable in liquid iron, but it is uncertain whether these phases can actually be formed from ferrosilicon
alloy additions in liquid iron. In fact, and in opposition to sulphides and oxides, such carbides have never been detected
experimentally as nucleation sites for graphite, although they, from a theoretical standpoint, are considered to be favorable
(Lux, 1964). It is however interesting to note that Mg does not form any known compound of the MgC2-type. This might
explain why Mg is only used as a spheroidizing agent in DI and not as an inoculant. Table 3 gives a summary of crystal
structures, melting points and standard free energies of formation for the group IIA-carbides.
Group IIA-Carbides (From I. Barin, Thermochemical Data of Pure Substances, VCH Verlagsgesellshaft, 1989)
Phase Space
group
Crystal
system
Lattice
parameter
(Å)
Melting
point, TM
(°C)
Free
energy,
∆GF
*)
MgC2 --- --- --- --- ---
CaC2 Fm3m cubic 5.86 2300 -106
SrC2 Fm3m cubic 6.24 -93
BaC2 Fm3m cubic 6.56 -96
19

OXIDES
The group IIA-elements also form stable oxides in liquid iron; therefore, nodularizers and inoculants based on these elements
are known to be effective deoxidizers. Table 4 gives information on crystal structures, melting points and standard free
energies of formation for the pure oxides. As with the sulphides of group IIA, the oxides also show very little difference in
crystal structure and stability between Mg and the other elements Ca, Sr and Ba. Thus, the pure oxides alone do not offer any
good explanation as to why Mg is only active as a spheroidizing agent in DI, while the other elements are effective also as
inoculating agents.
Group IIA-Oxides (From I. Barin, Thermochemical Data of Pure Substances, VCH Verlagsgesellshaft, 1989)
Phase Space
group
Crystal
system
Lattice
parameter
(Å)
Melting
point, TM
(°C)
Free
energy,
∆GF
*)
MgO Fm3m cubic 4.215 2832 -401
CaO Fm3m cubic 4.811 2927 -466
SrO Fm3m cubic 5.140 2665 -428
BaO Fm3m cubic 5.539 2013 -395
*) Standard free energy of formation at 1327C (1600K)
In the following, different types of more complex oxide inclusions that contain Mg, Ca, Sr or Ba as constituent elements are
considered. Since these elements are typically added via ferrosilicon alloys, a convenient basis for the discussion of oxide
inclusions is the ternary systems, XO–Al2O3–SiO2 where X denotes Mg, Ca, Sr or Ba. The specific interest in this respect is
the MgO–Al2O3–SiO2 and the CaO–Al2O3–SiO2 ternary systems, since a variety of different phases (silicates and aluminates)
may form, depending on the deoxidation and inoculation practices applied.
THE MgO–AL2O3–SiO2 SYSTEM
Non-metallic inclusions containing MgO as one component may form during the Mg treatment. The pure MgO–Al2O3–SiO2
system is similar to the MnO–Al2O3–SiO2 and FeO–Al2O3–SiO2 systems. A projection of the liquidus surface of the former
system is given in Fig. 6.
Fig. 6. Phase diagram for the system MgO–Al2O3–SiO2 includes also projections of different liquidus
surfaces. (From R. Kiessling, Non-Metallic Inclusions in Steel, book no. 194, the Metals Society, 1978).
20

It is evident that a number of silicate phases may form as a result of reactions between MgO and SiO2, including:
)(enstatite22 SiOMgOSiOMg ⋅→+
e)(forsterit22 22 SiOMgOSiOMgO ⋅→+
Enstatite can exist in three different modifications, and is a common reaction product in Mg treated DI slags (Askeland,
1972; Kiessling, 1978). MgO in MgO⋅SiO2 can be completely substituted by FeO, but pure FeO⋅SiO2 is not stable at normal
pressures. MgO can also be replaced by CaO up to about 50 wt%. In addition, enstatite may dissolve as much as 10 wt%
Al2O3. Because of a faceted growth morphology, enstatite tend to form characteristic angular shaped inclusions in DI, which
means that they can easily be identified by means of optical microscopy.
The other Mg silicate type, forsterite, may also exist in liquid iron. This phase may contain varying amounts of other oxides
in solid solution. For example, CaO is widely soluble in forsterite, since the crystal structure of 2MgO⋅SiO2 is similar to that
of γ-2CaO⋅SiO2. In addition to enstatite and forsterite, a variety of other phases have been detected in DI, including
(Mg,Al)3O4, (Mg,Al)SiO2 and (Mg,Al,Ca)SiO3 together with complex Ca and aluminum silicates and pure silica (Latona,
1984).
THE CaO-Al2O3-SiO2 SYSTEM
Ca is, from a technical standpoint, insoluble in liquid iron. Nevertheless, a small solubility has been reported by Sponseller
and Flinn, who found that pure iron could dissolve up to 0.032 wt% Ca at 1600C (2912F) (Sponseller and Flinn, 1964).
Slightly higher values were observed in the presence of Al, C, nickel (Ni) and Si. Ca is the most common trace element in
ferrosilicon inoculants. Consequently, due to the low solubility and the high affinity to O, Ca could play a role in the
graphite nucleation process by entering the deoxidation products at some later stage of the process. However, in steel
inclusions, CaO is not present as a separate phase, since it reacts readily with other oxides to form complex calcium silicates
and aluminates. This may also be the case in DI. Figure 7 shows the ternary CaO–Al2O3–SiO2 phase diagram and
projections of the different liquidus surfaces.
Fig. 7. Phase diagram for the system CaO–Al2O3–SiO2 includes also projections of different liquidus
surfaces. (From R. Kiessling, Non-Metallic Inclusions in Steel, book no. 194, The Metals Society, 1978).
Referring to the phase diagram, CaO can combine with silica according to the following reactions:
ite)(wollaston22
SiOCaOSiOCaO ⋅→+
)(rankinite2323 22 SiOCaOSiOCaO ⋅→+
)(bredigite22 22 SiOCaOSiOCaO ⋅→+
(alite)33 22 SiOCaOSiOCaO ⋅→+
Several modifications of the calcium silicates are also known, but the transformations between the different polymorphic
forms are complex and not fully understood.
21

CaO⋅SiO2 may dissolve varying amounts of MnO, FeO and Al2O3, but not MgO. Inclusions with a composition
corresponding to CaO⋅SiO2 are commonly observed in steel deoxidized with CaSi. These inclusions may contain up to about
10 wt% Al2O3 in solid solution (Kiessling, 1978).
CaO can also combine with alumina to form a number of different phases, including:
3232 OAlCaOOAlCaO ⋅→⋅
3232 22 OAlCaOOAlCaO ⋅→+
3232 66 OAlCaOOAlCaO ⋅→+
3232 33 OAlCaOOAlCaO ⋅→⋅
Oxide inclusions containing CaO⋅Al2O3 and CaO⋅2Al2O3 are all common deoxidation products. Some of the group IIA
silicates and aluminates are summarized in Table 5. This table also includes data for the inclusion crystal structures, lattice
parameters, and melting points (if known).
Table 5. Selected Crystallographic and Thermodynamic Data for Some Possible Silicates and Aluminates in
Liquid Iron Containing Mg, Ca, Sr and Ba (From I. Barin, Thermochemical Data of the Pure Substances, VCH
Verlagsgesellshaft, 1989; R. Kiessling, Non-Metallic Inclusions in Steel, book no. 194, The Metals Society,
1978; R. H. Rein and J. Chipman, Transactions of the Metals Society, AIME, vol 233, 1965)
Phase Space Group Crystal System Lattice
Parameters [Å]
TM
[°C]
∆GF*)
[kJ/mol]
MgO⋅SiO2 Pbca orthorhombic 18.2/8.86/5.204 1577 -1060
2MgO⋅SiO2 Pmnb orthorhombic 4.76/10.20/5.99 1898 -1491
CaO⋅SiO2 P1 hexagonal 6.82/19.65 1125-1544 -1184
SrO⋅SiO2 hexagonal 7.127/10.115 1580 -1186
BaO⋅SiO2 hexagonal 7.500/10.467 1605 -1180
MgO⋅Al2O3⋅2SiO2 - - - - -
CaO⋅Al2O3⋅2SiO2 P63/mcm hexagonal 5.113/14.743 1550 -3022
SrO⋅Al2O3⋅2SiO2 hexagonal 5.25/7.56
BaO⋅Al2O3⋅2SiO2 hexagonal 5.304/7.789 (1380)
MgO⋅6Al2O3 - - - - -
CaO⋅6Al2O3 hexagonal 5.54/21.82 1850
SrO⋅6Al2O3 P63/mmc hexagonal 5.589/22.07 1500
BaO⋅6Al2O3 P63/mmc hexagonal 5.607/22.90 (1400)
Four intermediate ternary phases exist in the ternary CaO–Al2O3–SiO2 system. Their stoiciometric compositions are as
follows:
CaO⋅Al2O3⋅2SiO2 (anorthite)
2CaO⋅Al2O3⋅SiO2 (gehlenite)
2CaO⋅2Al2O3⋅5SiO2 (corderite)
3CaO⋅Al2O3⋅3SiO2 (grossularite)
Bruch has studied the composition of Ca inclusions in steel after CaSi-deoxidation (Bruch, 1965). The mean composition of
these CaO-containing inclusions is situated within the dotted area of Fig. 8. Crystalline ternary phases in the CaO–Al2O3–
SiO2 system are rarer than glassy phases. The most common of the crystalline phases is anorthite (CaO⋅Al2O3⋅2SiO2), which
is the only ternary phase within the dotted area of Fig. 8. Anorthite undergoes four different transformations down to room
temperature and the stable high temperature modification is the hexagonal α-anorthite. Crystallographic data and melting
point for the α-anorthite phase are given in Table 5. Note that crystalline ternary phases of the gehlenite, corderite and
grossularite type are not common deoxidation products in liquid steel or cast iron.
22

Fig. 8. A summary of the mean composition (in wt%) of different calcia containing inclusions found in steel
after CaSi-deoxidation is illustrated. (From R. Kiessling, Non-Metallic Inclusions in Steel, book no. 194,The
Metals Society, 1978).
NITRIDES
In a study by Igarishi and Okada, the existence of a nitride phase, containing Mg, Al and Si taking part in the graphite
nucleating process, is demonstrated (Igarishi and Okada,1998). In a recent study by Solberg , the crystal structure and
composition of nuclei for graphite spheroids in DI containing small amounts of Mg and traces of Al are studied (Solberg,
2001). The particles are identified as Al–Mg–Si nitrides, having a trigonal super-lattice crystal structure derived from a
hexagonal Bravais lattice, with parameters a = 0.544 nm and c = 0.482 nm. The parameters of the fundamental cell are af =
0.314 nm and cf = 0.482 nm, deviating only 1-3% from the parameters of hexagonal AlN. Based on the compositional
analysis, the chemical formula of the nitride is suggested to be Mg2.5AlSi2.5N6. No such nitride phase containing all the
elements Mg, Al and Si are known from earlier literature.
Figure 9 shows a bright field image from TEM, a typical EDX (Energy Dispersive X-ray)-spectrum from the particle and the
crystallographic unit cell for this nitride phase. Solberg found, however, that apart from the fact that graphite also has a
hexagonal lattice, there is no obvious crystallographic similarity between the nitride particle and graphite (Solberg, 2001).
Thus, the nucleating power of the nitride does not seem to be associated with its crystal structure, but with the fact that it is a
heterogeneity in the melt. The simple existence of this Mg containing nitride phase may also explain why N gas defects are
rarely observed in treated DI. The Mg addition assists in effectively neutralizing N.
(a) (b) (c)
Fig. 9. These illustrate (a) bright field image of nitride particle from TEM, (b) EDX-spectrum and (c) four
super-lattice unit cells viewed along the (001)]-axis. (From J. K. Solberg and M. I. Onsoien, Materials Science
and Technology, vol 17, 2001).
23

NUCLEATION OF GRAPHITE AT INCLUSIONS
There seems to be general agreement in the literature that nucleation of graphite in DI occurs heterogeneously from particles
contained in the melt. The problem has mainly been to determine the nature of these heterogenities, i.e. their origin,
composition, surface characteristics, stability, etc. The SEM micrograph in Fig. 10 shows a graphite nodule associated with a
non-metallic inclusion that contains a magnesium sulphide core and an outer shell of complex magnesium silicates. It should
be noted that the presence of these phases is not a sufficient criterion for graphite formation, since Mg treatment of DI
generally is not regarded as an efficient means of nucleating high numbers of graphite nodules. Modification of the inclusion
surface chemistry by additions of minor elements through the inoculants is always required to achieve a high nodule density.
Consequently, the key to a better understanding of the microstructure evolution in DI lies primarily in the recognition of the
important difference between nodularizing (Mg treatment) and inoculation when it comes to graphite nucleation.
(a) (b)
Fig. 10. SEM micrographs show evidence of graphite nucleation at a complex duplex magnesium sulphide
and silicate inclusion: (a) graphite nodule with nucleus in the core and (b) larger magnification of core.
(From T. Skaland, Metallurgical Transactions A, vol 24A, 1993).
BRAMFITT’S PLANAR LATTICE DISREGISTRY MODEL
As stated earlier in this paper, the interfacial energy at the nucleating interface (γGN) is the controlling factor in heterogeneous
nucleation. For fully in-coherent interfaces, γGN would be expected to be of the order of 0.5-1 J/m2
. However, this value will
be greatly reduced if there is epitaxy between the inclusions and the graphite nucleus, which results in a low lattice
disregistry between the two phases. In general, assessment of the degree of atomic misfit between the graphite (G) and the
nucleant (N) can be done on the basis of Bramfitt’s planar lattice disregistry model (Bramfitt, 1970).
[ ]
( ) [ ]
[ ]
%100
cos
3
13
1 ⎟⎟
⎟
⎠
⎞
⎜⎜
⎜
⎝
⎛ −
= ∑
= i
N
i
G
i
N
uvw
uvwuvw
i d
dd α
δ Equation 8
where [uvw]N = a low-index direction in (hkl)N;
[uvw]G = a low-index direction in (hkl)G;
d[uvw]N = the inter-atomic spacing along [uvw]N;
d[uvw]G = the inter-atomic spacing along [uvw]G;
α = the angle between the [uvw]N and the [uvw]G.
In practice, the undercooling ∆T (which is a measure of the energy barrier against heterogeneous nucleation) increases
monotonically with increasing values of the planar lattice disregistry δ, as shown earlier in Fig. 2. This means that the most
potent catalyst particles are those that provide a good epitaxial fit between the nucleant and the graphite embryo.
The characteristic irregular shape of the inclusion in Fig. 5a indicates faceted growth morphology. Faceted growth occurs as
a result of anisotropy in the growth rates between high-index and low-index crystallographic planes (Kurz and Fisher, 1989).
If the former type of planes are the fastest growing ones, these planes grow out soon, leaving a faceted crystal delimited
solely by low-index planes.
GRAPHITE
SULPHIDE
SILICATE
NUCLEUS
24

NUCLEATION OF GRAPHITE AFTER MAGNESIUM TREATMENT
The probable crystal growth morphologies of enstatite (MgO⋅SiO2) and forsterite (2MgO⋅SiO2), as the main deoxidizing
substrates formed after Mg treatment, are shown in Fig. 11. Included in Fig. 11 is also a sketch of the lattice arrangement at
the interface between the (100)-plane of MgO⋅SiO2 and the (001)-plane of graphite. This orientation relationship conforms to
growth of graphite along the pole of the basal plane perpendicular to the inclusion surface, which is the normal growth mode
of graphite in DI. By only considering the position of the corner atoms in the orthorhombic unit cell (see sketch in Fig. 11c,
the planar lattice disregistry between graphite, enstatite and forsterite is calculated from Equation 8 for a wide spectrum of
orientation relationships. The results from these computations are summarized in Table 6.
(a) (b) (c)
Fig. 11. Diagrams illustrate nucleation of graphite at enstatite (MgO⋅SiO2) and forsterite (2MgO⋅SiO2)—(a)
crystal form of MgO⋅SiO2, (b) crystal form of 2MgO⋅SiO2 (From P. Ramdohr and H. Strunz, Lehrbuch der
Mineralogie, Ferdinand Enke Verlag, 1978), (c) details of lattice arrangement at graphite/ MgO⋅SiO2 interface
(From T. Skaland, Metallurgical Transactions A, vol 24A, 1993).
Table 6. Calculated Planar Lattice Disregistry between Enstatite, Forsterite and Graphite for Different
Orientation Relationships (From T. Skaland, Metallurgical Transactions A, vol 24A, 1993).
Inclusion Phase Orientation Relationship*)
Lattice Disregistry
(100)I ⏐⏐ (001)G 10.2 %
Enstatite (010)I ⏐⏐ (001)G 8.7 %
MgO⋅SiO2 (001)I ⏐⏐ (001)G 5.9 %
(110)I ⏐⏐ (001)G 12.3 %
(111)I ⏐⏐ (001)G 10.1 %
(100)I ⏐⏐ (001)G 9.9 %
Forsterite (010)I ⏐⏐ (001)G 24.3 %
2MgO⋅SiO2 (001)I ⏐⏐ (001)G 15.5 %
(101)I ⏐⏐ (001)G 25.5 %
(111)I ⏐⏐ (001)G 29.7 %
*) I = inclusion, G = graphite
It is evident from the data in Table 6 that chances of obtaining a small planar lattice disregistry between graphite, MgO⋅SiO2
or 2MgO⋅SiO2 are rather poor, which means that the energy barrier against heterogeneous nucleation is correspondingly high.
Hence, these phases, which are primary reaction products of the Mg treatment, would not be expected to act as favorable
nucleation sites for graphite during solidification. This is also in agreement with general experience of Mg treatment not
providing efficient nucleation of graphite.
The Mg treatment however, provides an important basis for the subsequent inoculation. Formation of a high number of small
magnesium sulphides, oxides, silicates and nitrides dispersed throughout the iron is expected to facilitate the formation of
heterogeneous and potent nucleation sites that settle onto the surface of the Mg treatment reaction products. A calm and
gentle Mg treatment is recognized to provide better conditions for the subsequent inoculation than a more violent and reactive
treatment process. The treatment process reactivity decides whether numerous small and dispersed nucleation sites or, on the
contrary, larger agglomerates of slag clusters are to be formed after the treatment. Thus, even though the Mg treatment does
not provide potent nucleation sites on its own, it forms a very important basis for the subsequent inoculant to settle potent
phases onto a highest possible number of sites initially produced during the Mg treatment.
25

NUCLEATION OF GRAPHITE AFTER INOCULATION
During inoculation with Ca, Sr or Ba containing ferrosilicon, faceted hexagonal silicate phases of the XO⋅SiO2 or the
XO⋅Al2O3⋅2SiO2 type (X denotes Ca, Sr or Ba) may form at the surface of the Mg treatment inclusion products. In particular,
the (001) basal planes of the crystals will be favorable sites for graphite nucleation, since these facets allow for formation of
coherent/semi-coherent low energy interfaces between the nucleant and the graphite, as illustrated by the examples in Fig. 12.
In fact, nearly all of the hexagonal silicate phases of the XO⋅SiO2 and the XO⋅Al2O3⋅2SiO2 type, which may form at the
surface of inclusions after the inoculation treatment, are effective catalyst substrates for graphite. Table 7 gives examples of
calculated planar lattice disregistry between graphite and the different silicates that may form during inoculation. The low
lattice disregistry explains why commercial inoculants for cast iron typically are based on either calcium, strontium or barium
as the critical reactive ingredients of the ferrosilicon inoculant.
(a) (b)
Fig.12. These illustrate the details of lattice arrangement at a nucleus/substrate interface—(a) coherent
graphite/BaO⋅SiO2 interface and (b) coherent graphite/CaO⋅Al2O3⋅2SiO2 interface. (From T. Skaland,
Metallurgical Transactions A, vol 24A, 1993).
Table 7. Calculated Planar Lattice Disregistry between Graphite and Different Inclusion Constituent Phases
That May Form During Inoculation (From T. Skaland, Metallurgical Transactions A, vol 24A, 1993).
Inclusion Phase Orientation Relationship*)
Lattice Disregistry
CaO⋅SiO2 (001)I ⏐⏐ (001)G 7.5 %
SrO⋅SiO2 (001)I ⏐⏐ (001)G 3.5 %
BaO⋅SiO2 (001)I ⏐⏐ (001)G 1.5 %
CaO⋅Al2O3⋅2SiO2 (001)I ⏐⏐ (001)G 3.7 %
SrO⋅Al2O3⋅2SiO2 (001)I ⏐⏐ (001)G 6.2 %
BaO⋅Al2O3⋅2SiO2 (001)I ⏐⏐ (001)G 7.1 %
*) I = inclusion, G = graphite.
EFFECTS OF Al, Ti, and Zr IN GRAPHITE NUCLEATION
Smickley reported that Al additions had a mild effect as an inoculant and a stronger effect as a chill reducer (Smickley,
1981). McClure conducted experiments with gray irons (GIs) exploring the inoculating effect of two ferrosilicon qualities,
one containing low levels of Al and Ca and the other one high levels of both elements (McClure, 1957). From these
experiments it is found that the alloy with low levels of Al and Ca had, for all practical purposes, no inoculation effect, while
the alloy with high levels of Al and Ca produced an iron with improved mechanical properties and reduced chill forming
propensity. A variety of proprietary inoculants are developed in order to produce more effective inoculants and more
consistent inoculation. These alloys normally, in addition to Al and Ca, contain controlled amounts of other elements.
Higher levels of Al, up to 4% in ferrosilicon based inoculants, is also found to improve inoculation effectiveness in DIs and
increase the ferrite content.
Titanium is a commonly known additive to GIs for N scavenging. Titanium is also observed to cause a change in the
hardness of iron castings (Merchant, 1972). Titanium has a very high affinity to O and S. Therefore, when adding Ti to cast
iron for N control, an element that has a higher affinity for S and O than Ti does, should also be added. Narasimhan found
that inoculation of GI with Ti effectively lowered the chill depth in low C equivalent irons (Narasimhan, 1969). Titanium
bearing inoculants are generally not recommended for DI due to the risk of interfering with the formation of graphite nodules.
The detrimental effect of Ti on the graphite morphology in DI is more pronounced as section size increases (Lownie, 1956;
Watmough, 1971). Additions of Ce or mish metal to the iron may reduce or eliminate the detrimental effects of Ti on the
26

graphite structure in DI (Pearce, 1962). Both the graphite morphology and the mechanical properties of the DI may be
restored and even improved when compared to irons without either Ti or rare earths (Naro, 1969; Sawyer, 1968).
Several investigations have reported that N in high carbon flake graphite irons may make the graphite shorter, round the end
of graphite flakes, hinder the growth of eutectic cells and stabilize the graphite, which are of benefit to some of the
mechanical properties (Hu, 1994, Koshel, 1971, Mountfort, 1966, Naro, 1970, Ruff, 1976, Wallace, 1965, Wallace, 1975).
Since the affinity between Zr and N is quite strong, very fine and dispersed nitrides are formed by adding Zr-bearing
inoculants (Quian, 1985). These Zr-nitrides may act as nuclei for graphite during the solidification and thus prevent chill
formation. The matrix of the iron will be strengthened by the dispersion of the hard ZrN particles as well. Dissolved N in
GIs is controllable by the deliberate introduction of Zr, thus reducing the risk for N fissure defects in the castings. The
possible beneficial effects of Zr in the DI nucleation process is however not well understood, since it is expected that the Mg
treatment itself will tie up and neutralize free N in the base iron forming complex Mg-Al-Si nitrides as described by Solberg
(Solberg, 2001). The formation of stable Zr-oxides as potent and additional nucleation sites for graphite, cannot however be
ruled out.
EFFECTS OF RARE EARTH METALS (REM) IN GRAPHITE NUCLEATION
The effects of Ce and other rare earth metals (REM) on microstructure and mechanical properties of DI are widely studied.
These studies generally involve the incorporation of REM in the MgFeSi nodularizer alloy. However, Amin conducted
experiments with Ce added either as a constituent in pre-alloyed MgFeSi, added separately with the MgFeSi, added with the
inoculant or added with both the MgFeSi and the inoculant (Amin, 1978). It was concluded from this extensive study of the
behavior of REM in DI that the observed effects are independent of the method of addition.
Previous investigations have shown that REM such as Ce and La can either have a beneficial or a detrimental effect on the
microstructure and properties of DI, depending on experimental conditions and additions. For example, small additions of
REM are frequently used to restore the graphite nodule count and nodularity in DIs containing subversive elements, such as
antimony (Sb) , lead (Pb), Ti, etc. (Bofan, 1984, Stefanescu, 1986, Udomom, 1985). On the other hand, REM in excessive
concentrations may lead to problems with chill formation in thin cast products and chunky graphite in larger section irons,
with subsequent degradation in the mechanical properties (Itofuji, 1990; Liu, 1989; Pan, 1994).
Several investigators reported an optimum level of REM with respect to a high nodule count and reduced carbide forming
propensity. However, the optimum rare earth content varies significantly according to different investigators. For example,
Lalich concluded that the optimum Ce level is about 0.006-0.010% for low Ce rare earths, and about 0.015-0.020% for high
Ce rare earths (Lalich, 1974). Kanetkar found a maximum nodule count at the following residual contents of REM: 0.007-
0.010% praseodymium (Pr), 0.017% neodymium (Nd), 0.018% La, 0.02% yttrium (Y) and 0.032% Ce (Kanetkar, 1984).
Onsøien reported an optimum Ce level of 0.035% and an optimum La level of 0.017% with respect to an optimum nodule
count in low sulphur DI (Onsøien, 1997).
It is expected that the formation of stable rare earth sulphides and oxides, and also oxy-sulphides play a very important role in
the heterogeneous nucleation of graphite in DI.
HIGH PURITY IRON CONDITIONS
GRAPHITE NUCLEATION IN HIGH PURITY MELTS
Heterogeneous nucleation of graphite at foreign particles is also reported in melts containing low levels of S and O (less than
0.2 ppm and 7 ppm, respectively). Dhindaw and Verhoeven studied vacuum melted high purity Fe–C–Si alloys produced
from ultra-pure zone-refined iron (Dhindaw and Verhoeven, 1980). They found from extensive SEM examinations that
impurity atoms are never detected in the nodule centers, which suggests that the graphite nucleation is not associated with
sulphide inclusions. However, commercial Si/Ca/Al-inoculants effectively increased the nodule count in the ultra-pure zone-
refined iron. It should however be noted that the maximum nodule count of about 26 per mm2
, as reported by Dhindaw and
Verhoeven is rather low compared with that normally observed in commercial DI where the nodule number density often
exceeds 300 to 400 per mm2
(Dhindaw and Verhoeven, 1980). Consequently, their investigation is not conclusive in that it
excludes the possibility of graphite nucleation on non-metallic inclusions.
GRAPHITE GROWTH IN HIGH PURITY MELTS
In a paper by Sadocha and Gruzleski, Fe–C–Si alloys, having various purities higher than those found for commercial irons,
are investigated (Sadocha and Gruzleski, 1974). It is shown that with increasing purity there is a transition from a plate-like
to a nodular morphology. Impurities present in commercial irons are thought to suppress nodular growth because they do not
allow curved crystal growth of graphite to occur, and they change the nature of the austenite-liquid mushy zone so that
27

graphite cannot grow physically unhindered by austenite. Their conclusion is that nodules appear to be the basic form of
graphite growth in pure Fe–C–Si alloys, and that the nodularizing elements act solely as scavengers to remove deleterious
impurities such as S and O from the melt.
In a recent paper by Nakae, the findings of Sadocha and Gruzleski were confirmed by applying special inert atmosphere and
crucibles for melting cast iron (Nakae, 2004; Sadocha and Gruzleski, 1974)). The effect of cooling rate and S level on the
graphite morphology is investigated. It is confirmed that the formation of spheroidal graphite needs a critical cooling rate
and S level in Mg free base iron conditions. A specimen with 86 mass ppm S has no spheroidal graphite even at the cooling
rate of 1000 K/min. Spheroidal graphite appeared in a specimen with 11 mass ppm S at the cooling rate of 100 K/min, while
for specimens with 1.5 mass ppm S spheroidal graphite easily formed at the cooling rate of 40 K/min. This confirms that
nodularizing additions are not required to form nodular graphite as long as the cooling rate is high or the S content is kept
ultra low, below 11 mass ppm. However, Nakae also found that cementite structure prevailed for the ultra low S containing
irons at higher cooling rates due to the difficulty of forming graphite nucleus (Nakae, 2004).
This confirms the theory that removing impurities such as S and O are decisive in controlling nodular graphite growth, and
that Mg additions as such are not directly required for this purpose. The key role of the nodularizing addition is to neutralize
S and O and thus facilitate nodular graphite growth in an essentially S and O free environment. The findings by Nakae and
Sadocha also confirm that S and O have a role in improving the nucleation of graphite by the formation of heterogeneous
sulphide and oxide nucleation sites (Nakae, 2004; Sadocha, 1974).
Figure 13 shows the 3-dimensional growth morphology of graphite from SEM investigations (Nakae, 2004). With reducing
concentrations of S from 98 to 1.5 mass ppm, the graphite morphology changes from flake to nodular growth.
Fig. 13. Growth morphology of graphite with reducing sulphur concentrations shows the transition to
nodular growth without introduction of any nodularizer to the iron. (From H. Nakae, S. Jung, H. Inoue, H.
Shin, Proceedings of the 66
th
World Foundry Congress, 2004).
CONCLUSIONS
The following summary is given from the present paper:
• Although the subject of nucleation and growth of graphite nodules in DI has long been a topic of considerable
discussion, conflicting views are still held about the major controlling mechanisms. This situation calls for a closer
review and examination of the existing theories.
• Several theories have been developed in the past to explain the nucleation of graphite nodules during solidification of DI,
including the gas bubble theory, the graphite theory, the silicon carbide theory, the salt-like carbide theory, the
sulphide/oxide theory, the nitride theory and the silicate theory. These theories are mostly based on the assumption that
the graphite is formed as a result of heterogeneous nucleation events occurring during solidification and that minor
elements such as Ca, Ba, Sr, Al, Zr, Ti, and Ce play an important role in this nucleation process.
• The effectiveness of a substrate in promoting heterogeneous nucleation depends on the crystallographic disregistry
between the substrate and the graphite to be nucleated. In practice, the undercooling, ∆T, increases with increasing
values of the planar lattice disregistry. Since the undercooling during solidification of DI is very small, the lattice
disregistry between the nucleus and the graphite phase must also be small and comparable with that of coherent/semi-
coherent interfaces.
• A variety of different inclusions (sulphides, oxides, nitrides, and silicates) can form in the liquid state. The sulphides and
oxides of the group IIA-elements (Mg, Ca, Sr and Ba) are all of the face centered cubic NaCl-structure type, and are
28

characterized by very similar lattice parameters and high melting points. The similarity between these phases does not
explain why Mg is not acting as a potent nucleating element for graphite in DI, while the other three elements Ca, Sr and
Ba will provide potent nucleation effects.
• Additions of Mg, Ca, Sr, or Ba to Si-rich iron melts may also result in the formation of complex silicates with different
stoiciometric compositions. After Mg treatment, a wide spectrum of inclusions is present in the melt, including Mg
silicates of the enstatite and forsterite type. These phases, however, will not act as potent nucleation sites for graphite
during solidification because of their non-hexagonal (orthorhombic) crystal structures and large planar lattice
disregistries with the hexagonal graphite.
• After inoculation with Ca, Sr or Ba (and Al), containing ferrosilicon, faceted hexagonal silicate phases of the XO⋅SiO2 or
the XO⋅Al2O3⋅2SiO2 type may form at the surface of existing inclusions from the Mg treatment. Such phases will be
favorable sites for graphite nucleation, since the hexagonal facets allow for the formation of coherent/semi-coherent low
energy interfaces between the nucleant and the graphite.
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