The simplest version of analyzes the effects of alloying elements on iron-carbon alloys
would require analysis of a large number of ternary alloy diagrams over a wide temperature
range. However, Wever pointed out that iron binary equilibrium systems fall into four main
categories (Fig. 1): open and closed γ-field systems, and expanded and contracted
γ-field systems. This approach indicates that alloying elements can influence the
equilibrium diagram in two ways:
- by expanding the γ-field, and encouraging the formation of austenite over wider
compositional limits. These elements are called γ-stabilizers.
- by contracting the γ-field, and encouraging the formation of ferrite over wider
compositional limits. These elements are called α-stabilizers.
The form of the diagram depends to some degree on the electronic structure of the alloying
elements which is reflected in their relative positions in the periodic classification.
Figure 1. Classification of iron alloy phase diagrams: a. open γ-field;
b. expanded γ-field; c. closed γ-field
Eisenhüttenwesen, 1928-9, 2, 193)
Class 1: open γ-field. To this group belong the important steel
alloying elements nickel and manganese, as well as cobalt and the inert metals ruthenium,
rhodium, palladium, osmium, iridium and platinum. Both nickel and manganese, if added in
sufficiently high concentration, completely eliminate the bcc α-iron phase and
replace it, down to room temperature, with the γ-phase. So nickel and manganese
depress the phase transformation from γ to α to lower temperatures (Fig. 1a),
i.e. both Ac1 and Ac3 are lowered. It is also easier to obtain
metastable austenite by quenching from the γ-region to room temperature, consequently
nickel and manganese are useful elements in the formulation of austenitic steels.
Class 2: expanded γ-field. Carbon and nitrogen are the most
important elements in this group. The γ-phase field is expanded, but its range of
existence is cut short by compound formation (Fig.1b). Copper, zinc and gold have a similar
influence. The expansion of the γ-field by carbon, and nitrogen, underlies the whole
of the heat treatment of steels, by allowing formation of a homogeneous solid solution
(austenite) containing up to 2.0 wt % of carbon or 2.8 wt % of nitrogen.
Class 3: closed γ-field. Many elements restrict the formation of
γ-iron, causing the γ-area of the diagram to contract to a small area referred
to as the gamma loop (Fig. 1c). This means that the relevant elements are encouraging the
formation of bcc iron (ferrite), and one result is that the δ- and γ-phase
fields become continuous. Alloys in which this has taken place are, therefore, not amenable
to the normal heat treatments involving cooling through the γ/α-phase
transformation. Silicon, aluminium, beryllium and phosphorus fall into this category,
together with the strong carbide forming elements, titanium, vanadium, molybdenum and
Class 4: contracted y-field. Boron is the most significant element of
this group, together with the carbide forming elements tantalum, niobium and zirconium.
The γ-loop is strongly contracted, but is accompanied by compound formation
The distribution of alloying elements in steels. Although only binary
systems have been considered so far, when carbon is included to make ternary systems the
same general principles usually apply. For a fixed carbon content, as the alloying clement
is added the y-field is either expanded or contracted depending on the particular
With an element such as silicon the γ-field is restricted and there is a
corresponding enlargement of the α-field. If vanadium is added, the γ-field is
contracted and there will be vanadium carbide in equilibrium with ferrite over much of the
ferrite field. Nickel does not form a carbide and expands the γ-field. Normally
elements with opposing tendencies will cancel each other out at the appropriate
combinations, but in some cases anomalies occur. For example, chromium added to nickel in
a steel in concentrations around 18% helps to stabilize the γ-phase, as shown by
18Cr8Ni austenitic steels.
One convenient way of illustrating quantitatively the effect of an alloying element on
the γ-phase field of the Fe-C system is to project on to the Fe-C plane of the
ternary system the γ-phase field boundaries for increasing concentration of a
particular alloying element. For more precise and extensive information, it is necessary
to consider series of isothermal sections in true ternary systems Fe-C-X, but even in
some of the more familiar systems the full information is not available, partly because
the acquisition of accurate data can be a difficult and very time-consuming process.
Recently the introduction of computer-based methods has permitted the synthesis of
extensive thermochemical and phase equilibria data, and its presentation in the form,
for example, of isothermal sections over a wide range of temperatures.
If only steels in which the austenite transforms to ferrite and carbide on slow cooling
are considered, the alloying elements can be divided into three categories:
- elements which enter only the ferrite phase
- elements which form stable carbides and also enter the ferrite phase
- elements which enter only the carbide phase.
In the first category there are elements such as nickel, copper, phosphorus and silicon
which, in transformable steels, are normally found in solid solution in the ferrite phase,
their solubility in cementite or in alloy carbides being quite low.
The majority of alloying elements used in steels fall into the second category, in so
far as they are carbide formers and as such, at low concentrations, go into solid solution
in cementite, but will also form solid solutions in ferrite. At higher concentrations most
will form alloy carbides, which are thermodynamically more stable than cementite.
Typical examples are manganese, chromium, molybdenum, vanadium, titanium, tungsten and
niobium. Manganese carbide is not found in steels, but instead manganese enters readily
into solid solution in Fe3C. The carbide-forming elements are usually present
greatly in excess of the amounts needed in the carbide phase, which are determined
primarily by the carbon content of the steel. The remainder enters into solid solution in
the ferrite with the non-carbide forming elements nickel and silicon. Some of these
elements, notably titanium, tungsten, and molybdenum, produce substantial solid solution
hardening of ferrite.
In the third category there are a few elements which enter predominantly the carbide
phase. Nitrogen is the most important element and it forms carbo-nitrides with iron
and many alloying elements. However, in the presence of certain very strong nitride
forming elements, e.g. titanium and aluminum, separate alloy nitride phases can occur.
While ternary phase diagrams, Fe-C-X, can be particularly helpful in understanding the
phases which can exist in simple steels, isothermal sections for a number of temperatures
are needed before an adequate picture of the equilibrium phases can be built up. For more
complex steels the task is formidable and equilibrium diagrams can only give a rough guide
to the structures likely to be encountered. It is, however, possible to construct
pseudobinary diagrams for groups of steels, which give an overall view of the equilibrium
phases likely to be encountered at a particular temperature.
Structural changes resulting from alloying additions. The addition to
iron-carbon alloys of elements such as nickel, silicon, manganese, which do not form
carbides in competition with cementite, does not basically alter the microstructures
formed after transformation. However, in the case of strong carbide-forming elements
such as molybdenum, chromium and tungsten, cementite will be replaced by the appropriate
alloy carbides, often at relatively low alloying element concentrations. Still stronger
carbide forming elements such as niobium, titanium and vanadium are capable of forming
alloy carbides, preferentially at alloying concentrations less than 0.1 wt%.
It would, therefore, be expected that the microstructures of steels containing these
elements would be radically altered. It has been shown how the difference in solubility
of carbon in austenite and ferrite leads to the familiar ferrite/cementite aggregates in
plain carbon steels. This means that, because the solubility of cementite in austenite
is much greater than in ferrite, it is possible to redistribute the cementite by holding
the steel in the austenite region to take it into solution, and then allowing
transformation to take place to ferrite and cementite. Examining the possible alloy
carbides, and nitrides, in the same way, shows that all the familiar ones are much less
soluble in austenite than is cementite.
Chromium and molybdenum carbides are not included, but they are substantially more soluble
in austenite than the other carbides. Detailed consideration of such data, together with
practical knowledge of alloy steel behavior, indicates that, for niobium and titanium,
concentrations of greater than about 0.25 wt % will form excess alloy carbides which
cannot be dissolved in austenite at the highest solution temperatures. With vanadium the
limit is higher at 1-2%, and with molybdenum up to about 5%. Chromium has a much higher
limit before complete solution of chromium carbide in austenite becomes difficult. This
argument assumes that sufficient carbon is present in the steel to combine with the
alloying element. If not, the excess metallic element will go into solid solution both
in the austenite and the ferrite.
In general, the fibrous morphology represents a closer approach to an equilibrium structure
so it is more predominant in steels which have transformed slowly. In contrast, the
interphase precipitation and dislocation nucleated structures occur more readily in rapidly
transforming steels, where there is a high driving force, for example, in microalloyed
The clearest analogy with pearlite is found when the alloy carbide in lath morphology
forms nodules in association with ferrite. These pearlitic nodules are often encountered
at temperatures just below Ac1, in steels which transform relatively slowly.
For example, these structures are obtained in chromium steels with between 4% and 12%
chromium and the crystallography is analogous to that of cementitic pearlite. It is,
however, different in detail because of the different crystal structures of the possible
carbides. The structures observed are relatively coarse, but finer than pearlite formed
under equivalent conditions, because of the need for the partition of the alloying element,
e.g. chromium between the carbide and the ferrite. To achieve this, the interlamellar
spacing must be substantially finer than in the equivalent iron-carbon case.
Interphase precipitation. Interphase precipitation has been shown to
nucleate periodically at the γ/α interface during the transformation. The
precipitate particles form in bands which are closely parallel to the interface, and which
follow the general direction of the interface even when it changes direction sharply. A
further characteristic is the frequent development of only one of the possible
Widmanstätten variants, for example VC plates in a particular region are all only
of one variant of the habit, i.e. that in which the plates are most nearly parallel to
The extremely fine scale of this phenomenon in vanadium steels, which also occurs in
Ti and Nb steels, is due to the rapid rate at which
the γ/α transformation takes place. At the higher transformation temperatures,
the slower rate of reaction leads to coarser structures. Similarly, if the reaction is
slowed down by addition of further alloying elements, e.g. Ni and
Mn, the precipitate dispersion coarsens.
The scale of the dispersion also varies from steel to steel, being coarsest in chromium,
tungsten and molybdenum steels where the reaction is relatively slow, and much finer in
steels in which vanadium, niobium and titanium are the dominant alloying elements and the
transformation is rapid.
Transformation diagrams for alloy steels. The transformation of austenite
below the eutectoid temperature can best be presented in an isothermal transformation
diagram, in which the beginning and end of transformation is plotted as a function of
temperature and time. Such curves are known as time-temperature-transformation, or TTT
curves, and form one of the important sources of quantitative information for the heat
treatment of steels.
In the simple case of a eutectoid plain carbon steel, the curve is roughly C-shaped with
the pearlite reaction occurring down to the nose of the curve and a little beyond. At
lower temperatures bainite and martensite are formed. The diagrams become more complex
for hypo- and hyper-eutectoid alloys as the ferrite or cementite reactions have also to
be represented by additional lines.