tempering of alloys steel
addition of alloying elements to steel has a substantial effect
on the kinetics of the transformation, and also of the pearlite
reaction. Most common alloying elements move the TTT curves to longer
times, with the result that it is much easier to "miss"
the nose of the curve during quenching. This essentially gives higher
hardenability, since martensite structures can be achieved at slower
cooling rates and, in practical terms, thicker specimens can be
made fully martensitic.
Alloying elements have also been shown to have a substantial effect
in depressing the Ms temperature. In this section, we will examine
the further important effects of alloying elements during the tempering
of martensite, where not only the kinetics of the basic reactions
are influenced but also the products of these reactions can be substantially
changed, e.g. cementite can be replaced by other carbide phases.
Several of the simpler groups of alloy steels will be used to provide
examples of the general behavior.
The effect of alloying elements on the formation of iron carbides
The structural changes during the early stage of tempering are difficult
lo follow. However, it is clear that certain elements, notably silicon,
can stabilize the carbide to such an extent that it is still present
in the microstructure after tempering at 400°C in steels with
1-2% Si, and at even higher temperatures if the silicon is further
While the tetragonality of martensite disappears by 300°C in
plain carbon steels, in steels containing some alloying elements,
e.g. Cr, Mo, W, V, Ti, Si, the tetragonal lattice is still observed
after tempering at 450°C and even as high as 500°C. It is
clear that these alloying elements increase the stability of the
supersaturated iron-carbon solid solution. In contrast manganese
and nickel decrease the stability.
also greatly influence the proportion of austenite retained on quenching.
Typically, a steel with 4% molybdenum, 0.2% C, in the martensitic
state contains less than 2% austenite, and about 5% is detected
in a steel with 1% vanadium and 0.2% C. On tempering each of the
above steels at 300°C, the austenite decomposes to give thin
grain boundary films of cementite which, in the case of the higher
concentrations of retained austenite, can be fairly continuous along
the lath boundaries. It is likely that this interlace cementite
is responsible for tempered martensite embrittlement, frequently
encountered as a toughness minimum in the range 300-350°C, by
leading to easy nucleation of cracks, which then propagate across
the tempered martensite laths.
can also restrain the coarsening of cementite in the range 400-700°C,
a basic process during the fourth stage of tempering. Several alloying
elements, notably silicon, chromium, molybdenum and tungsten, cause
the cementite to retain its fine Widmanstatten structure to higher
temperatures, either by entering into the cementite structure or
by segregating at the carbide-ferrite interfaces. Whatever the basic
cause may be, the effect is to delay significantly the softening
process during tempering. This influence on the cementite dispersion
has other effects, in so far as the carbide particles, by remaining
finer, slow down the reorganization of the dislocations inherited
from the martensite, with the result that the dislocation substructures
refine more slowly. The cementite particles are also found on ferrite
grain boundaries, where they control the rate at which the ferrite
In plain carbon
steels cementite particles begin to coarsen in the temperature range
350-400°C, and addition of chromium, silicon, and molybdenum
or tungsten delays the coarsening to the range 500-550°C. It
should be emphasized that up to 500°C, the only carbides to
form are those of iron. However, they will take varying amounts
of alloying elements into solid solution and may reject other alloying
elements as they grow.
The formation of alloy carbides: secondary hardening
A number of the familiar alloying elements in steels form carbides,
which are thermodynamically more stable than cementite. It is interesting
to note that this is also true of a number of nitrides and borides.
Nitrogen and boron are increasingly used in steels in small but
significant concentrations. The alloying elements Cr, Mo, V, W and
Ti all form carbides with substantially higher enthalpies of formation,
while the elements nickel, cobalt and copper do not form carbide
phases. Manganese is weak carbide former, found in solid solution
in cementite and not in a separate carbide phase.
It would, therefore, be expected that when strong carbide forming
elements are present in steel in sufficient concentration, their
carbides would be formed in preference to cementite. Nevertheless,
during the tempering of all alloy steels, alloy carbides do not
form until the temperature range 500-600°C, because below this
the metallic alloying elements cannot diffuse sufficiently rapidly
to allow alloy carbides to nucleate.
elements diffuse substitutionally, in contrast to carbon and nitrogen
which move through the iron lattice interstitially, with the result
that the diffusivities of carbon and nitrogen are several orders
of magnitude greater in iron, than those of the metallic alloying
elements. Consequently, higher temperatures are needed for the necessary
diffusion of the alloying elements prior to the nucleation and growth
of the alloy carbides and, in practice, for most of the carbide
forming elements this is in the range 500-600°C.
hardening process is a type of age-hardening reaction, in which
relatively coarse cementite dispersion is replaced by new and much
finer alloy carbide dispersion. On attaining a critical dispersion
parameter, the strength of the steel reaches a maximum, and as the
carbide dispersion slowly coarsens, the strength drops.
Nucleation and growth of alloy carbides
The dispersions of alloy carbides which occur during tempering can
be very complex, but some general principles can be discerned which
apply to a wide variety of steels. The alloy carbides can form in
at least three ways:
In-site nucleation at pre-existing cementite particles. It has been
shown that the nuclei form on the interfaces between cementite particles
and the ferrite. As they grow, carbon is provided by the adjacent
cementite, which gradually disappears.
By separate nucleation within the ferrite matrix, usually on dislocations
inherited from the martensitic structure.
At grain boundaries and sub boundaries-these include the former
austenite boundaries, the original martensitic lath boundaries (now
ferrite), and the new ferrite boundaries formed by coalescence of
sub boundaries, or by recrystallization.
In-site nucleation at pre-existing cementite particles are a common
occurrence but because these particles are fairly widely spaced
at temperatures above 500°C, the contribution of this type of
alloy carbide nucleation to strength is very limited.
The nucleation of carbides at the various types of boundary is to
be expected because these are energetically favourable sites, which
provide paths for relatively rapid diffusion of solute. Consequently
the ageing process is usually more advanced in these regions and
the precipitate is more massive. In many alloy steels, the first
alloy carbide to form is not the final equilibrium carbide and,
in some steels, as many as three alloy carbides can form successively.
In these circumstances, the equilibrium alloy carbide frequently
nucleates first in the grain boundaries, grows rapidly and eventually
completely replaces the Widmanstatten non-equilibrium carbide within
Tempering of steels containing vanadium
Vanadium is a strong carbide former and, in steel with as little
as 0.1% V, the face-centered cubic vanadium carbide VC is formed.
It is often not of stoichiometric composition, being frequently
nearer V4C3, but with other elements in solid solution within the
carbide. Normally, this is the only vanadium carbide formed in steels,
so the structural changes during tempering of vanadium steels are
of steels containing molybdenum and tungsten
When molybdenum or tungsten is the predominant alloying element
in a steel, a number of different carbide phases are possible, but
for composition between 4 and 6 wt% of the element the carbide sequence
is likely to be: Fe3C- Mo3C- W2C.
The carbides responsible for the secondary hardening in both the
case of tungsten and molybdenum are the isomorphous hexagonal carbides
Mo3C and W2C, both of which, in contrast to vanadium carbide, have
a well-defined rod let morphology.
Complex alloy steels
The presence of more than one carbide-forming element can complicate
the precipitation processes during tempering. In general terms,
the carbide phase which is the most stable thermodynamically will
predominate, but this assumes that equilibrium is reached during
tempering. This is clearly not so at temperatures below 500-600°C.
The use of pseudo-binary diagrams for groups of steels, e.g. Cr-V,
Cr-Mo, can be a useful guide to carbide phases likely to form during
Certain strong carbide formers, notably niobium, titanium and vanadium,
have effects on tempering out of proportion to their concentration.
In concentrations of 0.1 wt % or less, provided the tempering temperature
is high enough, i.e. 550-650°C, they combine preferentially
with part of the carbon and, in addition to the major carbide phase,
e.g. Cr7C3, Mo2C, they form a separate, very much finer dispersion,
more resistant to over-ageing.
dispersion can greatly augment the secondary hardening reaction,
illustrating the importance of these strong carbide forming elements
in achieving high strength levels, not only at room temperature
but also at elevated temperatures, where creep resistance is often
an essential requirement.