Maraging steels may feature as suitable grain refiners in HSLA steels;

9.2.3 Maraging steels may feature as suitable grain refiners in HSLA steels;

examples include AlN, Nb(CN), V(CN), (NbV)CN,

A serious limitation in producing high-strength steels TiC and Ti(CN). The solubility of these particles in the is the associated reduction in fracture toughness. Car-

austenite decreases in the order VC, TiC, NbC while bon is one of the elements which mostly affects the

the nitrides, with generally lower solubility, decrease toughness and hence in alloy steels it is reduced

in solubility in the order VN, AlN, TiN and NbN. to as low a level as possible, consistent with good

Because of the low solubility of NbC, Nb is per- strength. Developments in the technology of high-alloy

haps the most effective grain size controller. However, steels have produced high strengths in steels with very

Al, V and Ti are effective in high-nitrogen steels, Al low carbon contents ⊲<0.03%⊳ by a combination of

because it forms only a nitride, V and Ti by forming martensite and age-hardening, called maraging. The

V(CN) and Ti(CN) which are less soluble in austenite maraging steels are based on an Fe– Ni containing

than either VC or TiC.

between 18% and 25% Ni to produce massive marten- The major strengthening mechanism in HSLA steels site on air cooling to room temperature. Additional

is grain refinement but the required strength level is

300 Modern Physical Metallurgy and Materials Engineering

obtained usually by additional precipitation strength- ening in the ferrite. VC, for example, is more soluble in austenite than NbC, so if V and Nb are used in combination, then on transformation of the austen- ite to ferrite, NbC provides the grain refinement and VC precipitation strengthening; Figure 9.3 shows a stress – strain curve from a typical HSLA steel.

Solid-solution strengthening of the ferrite is also possible. Phosphorus is normally regarded as deleteri- ous due to grain boundary segregation, but it is a pow- erful strengthener, second only to carbon. In car con- struction where the design pressure is for lighter bodies and energy saving, HSLA steels, rephosphorized and bake-hardened to increase the strength further, have allowed sheet gauges to be reduced by 10 – 15% while maintaining dent resistance. The bake-hardening arises

Figure 9.3 Stress–strain curves for plain carbon, HSLA and from the locking of dislocations with interstitials, as

dual-phase steels .

discussed in Chapter 7, during the time at the temper- ature of the paint-baking stage of manufacture.

volume fraction of hard phase increases with a corre- sponding decrease in ductility; about 20% volume frac-

tion of martensite produces the optimum properties. Much research into the deformation behaviour of spe-

9.2.5 Dual-phase (DP) steels

The dual phase is produced by annealing in the (˛ C ) region followed by cooling at a rate which ensures ciality steels has been aimed at producing improved

that the -phase transforms to martensite, although strength while maintaining good ductility. The con-

some retained austenite is also usually present leading ventional means of strengthening by grain refinement,

to a mixed martensite– austenite (M – A) constituent. solid-solution additions (Si, P, Mn) and precipitation-

To allow air-cooling after annealing, microalloying hardening by V, Nb or Ti carbides (or carbonitrides)

elements are added to low-carbon – manganese– silicon have been extensively explored and a conventionally

steel, particularly vanadium or molybdenum and chro- treated HSLA steel would have a lower yield stress

mium. Vanadium in solid solution in the austenite of 550 MN m , a TS of 620 MN m and a total

increases the hardenability but the enhanced harden- elongation of about 18%. In recent years an improved

ability is due mainly to the presence of fine carboni- strength – ductility relationship has been found for low-

tride precipitates which are unlikely to dissolve in carbon, low-alloy steels rapidly cooled from an anneal-

either the austenite or the ferrite at the temperatures ing temperature at which the steel consisted of a

employed and thus inhibit the movement of the austen- mixture of ferrite and austenite. Such steels have

ite/ferrite interface during the post-anneal cooling.

a microstructure containing principally low-carbon, The martensite structure found in dual-phase steels fine-grained ferrite intermixed with islands of fine

is characteristic of plate martensite having internal martensite and are known as dualphase steels. Typi-

microtwins. The retained austenite can transform to cal properties of this group of steels would be a TS

martensite during straining thereby contributing to the of 620 MN m , a 0.2% offset flow stress of 380 MN

increased strength and work-hardening. Interruption m and a 3% offset flow stress of 480 MN m with

of the cooling, following intercritical annealing, can

a total elongation ³28%. lead to stabilization of the austenite with an increased The implications of the improvement in mechan-

strength on subsequent deformation. The ferrite grains ical properties are evident from an examination of

(³5 µ m) adjacent to the martensite islands are gen- the nominal stress – strain curves, shown in Figure 9.3.

erally observed to have a high dislocation density The dual-phase steel exhibits no yield discontinuity

resulting from the volume and shape change associ- but work-hardens rapidly so as to be just as strong

ated with the austenite to martensite transformation. as the conventional HSLA steel when both have been

Dislocations are also usually evident around retained deformed by about 5%. In contrast to ferrite– pearlite

austenitic islands due to differential contraction of the steels, the work-hardening rate of dual-phase steel

ferrite and austenite during cooling. increases as the strength increases. The absence of

Some deformation models of DP steels assume both discontinuous yielding in dual-phase steels is an advan-

phases are ductile and obey the Ludwig relationship, tage during cold-pressing operations and this feature

with equal strain in both phases. Measurements by sev- combined with the way in which they sustain work-

eral workers have, however, clearly shown a partition- hardening to high strains makes them attractive mate-

ing of strain between the martensite and ferrite, with rials for sheet-forming operations. The flow stress and

the mixed (M – A) constituent exhibiting no strain until tensile strength of dual-phase steels increase as the

deformations well in excess of the maximum uniform

Modern alloy developments 301 strain. Models based on the partitioning of strain pre-

dict a linear relationship between yield stress, TS and volume fraction of martensite but a linear relationship is not sensitive to the model. An alternative approach is to consider the microstructure as approximating to that of a dispersion-strengthened alloy. This would be appropriate when the martensite does not deform and still be a good approximation when the strain differ- ence between the two phases is large. Such a model affords an explanation of the high work-hardening rate, as outlined in Chapter 7, arising from the interaction of the primary dislocations with the dense ‘tangle’ of dislocations generated in the matrix around the hard phase islands.

Several workers have examined DP steels to deter- mine the effect of size and volume fraction of the hard phase. Figure 9.4 shows the results at two dif- ferent strain values and confirms the linear rela-

⊲f/d⊳ 1/2 predicted by the dispersion-hardening the- ory (see Chapter 7). Increasing the hard phase volume fraction while keeping the island diameter constant increases the work-hardening rate, increases the TS but decreases the elongation. At constant volume frac- tion of hard phase, decreasing the mean island diameter produces no effect on the tensile strength but increases the work-hardening rate and the maximum uniform elongation (Figure 9.5). Thus the strength is improved by increasing the volume fraction of hard phase while the work-hardening and ductility are improved by reducing the hard phase island size. Although dual-phase steels contain a complex microstructure it appears from their mechanical behaviour that they can be considered as agglomerates of non-deformable hard particles, made up of martensite and/or bainite and/or retained austenite, in a ductile matrix of ferrite. Consistent with the dispersion-strengthened model, the

Figure 9.4 Dependence of work-hardening rate on (volume fraction f/particle size) 1 /2 for a dual-phase steel at strain values of 0.2 and 0.25 (after Balliger and Gladman, 1981) .

Bauschinger effect, where the flow stress in compres- sion is less than that in tension, is rather large in dual- phase steels, as shown in Figure 9.6 and increases with increase in martensite content up to about 25%. The Bauschinger effect arises from the long-range back- stress exerted by the martensite islands, which add to the applied stress in reversed straining.

The ferrite grain size can give significant strengthen- ing at small strains, but an increasing proportion of the strength arises from work-hardening and this is inde- pendent of grain-size changes from about 3 to 30 µ m. Solid solution strengthening of the ferrite (e.g. by sil- icon) enhances the work-hardening rate; P, Mn and

V are also beneficial. The absence of a sharp yield point must imply that the dual-phase steel contains a high density of mobile dislocations. The microstruc- ture exhibits such a dislocation density around the martensite islands but why these remain unpinned at ambient temperature is still in doubt, particularly as strain-ageing is significant on ageing between 423 and 573 K. Intercritical annealing allows a partitioning of the carbon to produce very low carbon ferrite, while aluminium- or silicon- killed steels have limited nitro- gen remaining in solution. However, it is doubtful whether the concentration of interstitials is sufficiently low to prevent strain-ageing at low temperature; hence it is considered more likely that continuous yielding is due to the residual stress fields surrounding second- phase islands. Two possibilities then arise: (1) yielding can start in several regions at the same time rather than in one local region which initiates a general yield process catastrophically, and (2) any local region is prevented from yielding catastrophically because the glide band has to overcome a high back stress from the M – A islands. Discontinuous yielding on ageing at higher temperatures is then interpreted in terms of the relaxation of these residual stresses, followed by classical strain-ageing.

In dual-phase steels the n value ³0.2 gives the high and sustained work-hardening rate required when

stretch formability is the limiting factor in fabrica- tion. However, when fracture per se is limiting, dual- phase steels probably perform no better than other steels with controlled inclusion content. Tensile failure of dual-phase steels is initiated either by decohesion of the martensite– ferrite interface or by cracking of the martensite islands. Improved fracture behaviour is obtained when the martensite islands are unconnected, when the martensite– ferrite interface is free from pre- cipitates to act as stress raisers, and when the hard phase is relatively tough. The optimum martensite con- tent is considered to be 20%, because above this level void formation at hard islands increases markedly.