Tempering of alloy steels
The 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 increased.
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.
Alloying elements 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.
Alloying elements 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 grains grow.
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.
The metallic 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.
This secondary 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.
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 the grains.
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 relatively simple.
Tempering 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 tempering.
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.
This secondary 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.
List of Articles - Knowledge Base