The Tempering of Martensite: Part Two

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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:

1. 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.
2. By separate nucleation within the ferrite matrix, usually on dislocations inherited from the martensitic structure.
3. 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 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.