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However, in the base steel and the vanadium containing steels, ferrite did recrystallize during holding, and therefore, structu­ res associated with recovered deformed ferrite were only observed in these steels.

It is generally believed that partially recrystallized austenite trans­ forms to a mixed ferrite grain size. To overcome such a mixed ferrite grain size, it has often been suggested that a rolling practice is required which can give ~ uniform recrystallized austenite prior to transformation. It should be noted that this would be difficult to achieve considering the requirement for strict control over the variables which leads to the formation of partially recrystallized austenite.

Therefore, an alternative should consider a method for transforming the partially recrystallized austenite to a uniform ferrite grain size.

Studies on the transformation of austenite after various thermo-mechanical treatments suggest that, to a certain extent, the formation of mixed

ferrite grain sizes can be retarded or inhibited by controlling the transformation conditions, after finish rolling. The formation of coarse ferrite grains, — "seems to be the main cause for the transforma­ tion of partially recrystallized austenite to mixed ferrite grain sizes.

If by some mechanism an increased nucleation of ferrite can be achieved within such coarse unrecrystallized austenite, the occurrence of mixed

ferrite grain size would be minimised.

It has been shown that intragranular nucleation of ferrite occurs on the deformation bands and on sub-structural boundaries. It should be noted that these nucleation sites are not very effective due to their low’energy. Hence it is necessary to activate more of these nucleation sites by increasing deformation in the finishing passes, to produce a more defective structure within the austenite grains, and by fast cooling immediately after finish rolling. Fast cooling minimises the growth of ferrite and also increases the nucleation rate within the austenite grains. This has been clearly demonstrated in the present investigation, see Figs. 75» 76, 81 and 83, which showed the transfor­ mation of austenite to mixed ferrite grain size when isothermal trans­ formation occurs at 750°C, but when transformation occurs during conti­ nuous cooling uniform ferrite grains are formed. Isothermal transformation also represents the slow transformation which can occur during slow

cooling. In conclusion, it is suggested that if sufficient deformation is given in the finishing passes, followed by fast cooling (but not so fast that transformation to acicular products occurs), the occurrence of mixed ferrite grain sizes can be minimised or prevented.

An another mechanism which leads to the formation of a mixed ferrite grain size is through non-uniform strain introduced in the through

thickness direction during rolling. It is known that maximum deformation occurs just below the surface and decreases towards the mid-thickness of e. steel plate. Such strain variations are likely to be more prominent in thick sections rather than in thin flat-rolled products. In the present investigation such a variation of strain was observed, as shown in Fig. 141» and resulted in a fine ferrite grain size just below the surface whilst some what coarser ferrite grains occurred at the mid­ thickness.

13.3.6 A comparative effect of niobium and vanadium on the transfor­ mation of austenite. and the ferrite grain size.

austenite, and this was particularly marked for niobium than for vanadium. However, both elements forms carbides and nitrides, thus making their effect on hardenability complex, as discussed earlier. Niobium is a stronger carbide and nitride former than vanadium and therefore it us difficult to utilise to the full the strong effect of niobium on hardena­ bility.

The reason why niobium is more effective than vanadium may be associated with its larger atomic diameter which will cause it to segregate more positively to the austenite grain boundaries, and also to.its greater affinity for carbon which will cause it to decrease the diffusion rate of carbon.

After various thermo-mechanical treatments it has been shown that

additions of either niobium or vanadium refined the ferrite grain size. Vanadium refined the ferrite grain size particularly when the nitrogen content was high i.e. rJ 0.0296. A general comparison of the ferrite grain size obtained suggests that niobium refined the ferrite grain over a much more wider range of thermo-mechanical treatments than did vanadium. On the other hand, vanadium refined the ferrite grain size at low rolling temperature, due to its effect on recrystallization. Therefore, a

c.omparison of the effectiveness of niobium and vanadium on the ferrite grain refinement should be carried out at conditions giving optimum refinement for each element, rather than at the same rolling conditions. It should be noted that only a small amount of niobium or vanadium is needed to obtain the optimum ferrite grain refinement. The present results suggest that both niobium and vanadium are comparable in terms of their effect on ferrite grain refinement but the rolling practice to achieve optimum ferrite grain refinement is different for niobium and vanadium micro-alloyed steels, A possible rolling practice for niobium and vanadium steels to obtain optimum ferrite refinement, will be discu­ ssed later.

13.4 E ffe c t o f Thermo-mechanical Treatm ents and Composition V a ria b le s on the Precipitation Strengthening.

13.4*1 Niobium steels.

Niobium additions to plain carbon- steels are used to refine the ferrite (185)

grain size and also to impart precipitation strengthening' . Precipi­ tation strengthening is due to the precipitation of Nb(C,N) in ferrite formed from the transformation of austenite present at the finish rolling temperature. However, precipitation of Nb(C,N) in the austenite, which can occur due to strain inducement during thermo-mechanical treatment, does not produce strengthening. Instead, by decreasing the amount of Nb(C,N) available for subsequent precipitation in the ferrite, precipi­ tation in austenite actually decreases the precipitation strengthening

(q x 1 7 5

)

which is'observed' * . Strain induced precipitation during thermo- mechanical treatment depends on the strain, strain rate, time and. temperature parameters involved, and any treatment which leads to such precipitation will reduce the general strength level observed, but there has been little systematic work to investigate this. It is also well known that the maximum precipitation strengthening occurs at the stoichiometric Nb:(C+N) ratio^^*^^, because at stoichiometry the temperature dependence of the solubility of Nb(C,N) is a maximum. Again, however there is little systematic work to illustrate this effect,

. The precipitation strengthening, is used as a measure of preci­ pitation occurring in the austenite, the smaller the more precipi­ tation having occurred in austenite.

13.4,1.1 The effect of thermo-mechanical treatments.

Precipitation strengthening increased with increasing reheating temper­ ature, Pig. 89, due to the increased solubility of Nb(C,N). Because the lower niobium steels have less than the stoichiometric Nb:(C+N) ratio, see Fig. 113* the amount of Nb(C,N) precipitated will be controlled by the niobium dissolved in the austenite. It can be seen from Pig. 142 that for a given rolling condition, (AH) was a _P function of the piount

of niobium dissolved in the austenite at the reheating temperature* There was also a tendency for the increase in (AH) with increasing reheating temperature to be greater as the Nb: (C+N^ ratio increased up to the stoichiometric ratio#

It was apparent from Pigs. 89 and 142 that the rolling temperature influenced C^Op* Rolling at 1250°C and at 750°C produced a much larger (AH)^ value than rolling at 950°C, see Pig. 93* This was due to there being a maximum rate of strain induced precipitation of Nb(C,N) at 950°C, which therefore, precipitated more Nb(C,N) in the austenite and consequently gave less precipitation in the ferrite and a lower value of (AR)p* At rolling temperatures higher than 950°C, although the controlling diffusion rates were faster, the degree of supersaturation was lower, and it was the latter which resulted in the smaller strain- induced precipitation. At lower rolling temperatures than 950°C, the slow diffusion rates were responsible, for the slower strain-induced precipitation. Consequently, the precipitation kinetics for strain- induced Nb(C,N) showed a fC' curve with the maximum rate of precipita­ tion at about 950°C, A similar effect has been observed by other workers f n c *|7Q^‘JQ^ }

v 9 . Such * C * curve kinetics for strain induced precipitation