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VANADIUM APPLICATION TECHNOLOGY THE USE OF VANADIUM IN STEEL A Selection of papers Presented at The Vanitec International Symposium Held in Beijing, China 13-14 October 2001 V A N I T E C Vanadium International Technical Committee Winterton House, High Street, Westerham, Kent, TN16 1AQ England Tel: (0)1959 563400 Fax: (0)1959 562563 EmaI: [email protected] Web Site: www.vanitec.org
Transcript
Page 1: The Use of Vanadium

VANADIUM APPLICATION TECHNOLOGY

THE USE OF VANADIUM IN STEEL A Selection of papers Presented

at The Vanitec International Symposium Held in Beijing, China

13-14 October 2001

V A N I T E C Vanadium International Technical Committee Winterton House, High Street, Westerham, Kent, TN16 1AQ England Tel: (0)1959 563400 Fax: (0)1959 562563 EmaI: [email protected] Web Site: www.vanitec.org

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Contents The Use of Vanadium 1 P. S. Mitchell A New Role for Microalloyed steels Adding Economic Value 24 Michael Korchynsky Development and Use of Vanadium in Micro-Alloyed Reinforcing Bar 35 D. Russwurm Vanadium Microalloyed Forging Steels 59 G. Krauss Precipitation and Grain Refinement in Vanadium-containing Steels 74 S. Zajac The Evolution of Microstructure During Thin Slab Direct Rolling Processing in Vanadium Microalloyed Steels 94 Y. Li, D. N. Crowther, P. S. Mitchell and T. N. Baker The Effects of Microalloying Elements on Cracking During 115 Continuous Casting D. N. Crowther The Influence of Vanadium-Microalloying on the Weldability of Steels 146 P. H. M. Hart Vanadium Microalloying in Steel Sheet, Strip and Plate Products 165 Robert J. Glodowski

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THE USE OF VANADIUM

P. S. Mitchell

Chairman, VANITEC Winterton House, High Street, Westerham, Kent, TN16 1AQ, England.

1. Introduction

Vanadium is the 17th most commonly occurring element in the earths’ crust and finds wide use as an alloying element in steels, titanium based alloys for aerospace applications, catalysts, ceramics, chemicals and probably, in the future, in batteries both for the storage of electricity generated during off peak times and as power packs for electrically powered vehicles. However, as can be seen from Figure 1,(1) by far the majority of vanadium consumed, presently, is used by the steel industry and it is on giving a broad indication of this use that this paper will concentrate.

2. The Use of Vanadium in Steel

Figure 2 shows the range of steel products where vanadium is used by the steel industries of Germany, Japan and the U.S.A. It appears from this figure that the pattern of use in the three countries is different, only tool steels being consistently categorised in all three. There are, however, other less obvious similarities in the pattern of use. Structural steels recorded in Germany are almost certainly recorded within HSLA / carbon steels in the U.S.A. and in “others” in Japan. Similarly linepipe steels recorded in Japan are almost certainly recorded within special structural steels in Germany and within HSLA steels in the U.S.A. The main point to note, however, is that when vanadium is added to steel it results in a benefit to steel production or to an improvement in steel properties or to both, ultimately leading to a reduction in the cost of producing or of using the steel or of both.

Figure 3 shows vanadium consumption in the World and in the West up to the end of 2000. From this it is clear that consumption has been growing strongly since the early 1990’s. Indeed, world consumption of vanadium in 2000 was 84% higher than in 1993 and was, in fact, a record. Similarly, in the West, consumption of vanadium has increased by almost 70% since 1993. The record world consumption, to some extent, reflects the record world steel production of 826 million tonnes observed in 2000. Dividing vanadium consumption by steel production we see that the unit consumption (or the intensity of consumption) of vanadium, in 2000, was also at a record level of just over 0.047kgV / T steel in the world as a whole and 0.053kgV / T steel in the West (Figure 4). The world unit consumption has increased by approximately 60% since 1993 while Western unit consumption has increased by almost 50% in the same period. These increases represent new consumption over and above that which would be expected from the increase in steel production, and reflect the efforts of workers in research and development, worldwide, into the use of vanadium in steel.

I will now examine some of the technical attributes of the use of vanadium, which have led to this significant increase in consumption of vanadium. Many of these attributes will be covered in more detail by other speakers at this conference.

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3. Technical Attributes of the Use of Vanadium in Steel

3.1 Steel Production

In steel, vanadium forms stable compounds with both carbon and nitrogen and it is in the way in which these elements interact with vanadium which determines many of the properties of vanadium containing steels.

3.2 Solubility of Vanadium Compounds in Austenite and Ferrite

Vanadium carbides, nitrides (and carbonitrides) exhibit solubility in both austenite (2-6) and in ferrite (7-9) (Figure 5a & b). In austenite, the solubility of VC is the highest of all of those shown, while that of VN is lower than that of VC and is more equal to that of NbC, NbN and TiC. It is also interesting to note that in the group shown, TiN has the lowest solubility in austenite.

Vanadium carbide also exhibits higher solubility in ferrite than that of the other micro-alloy carbides and nitrides shown. It is important to note that the solubility of all the carbides and nitrides shown reduce as the temperature falls and that the solubility in ferrite is significantly lower than that in austenite.

3.3 Precipitation of Vanadium Compounds in Austenite and Ferrite

The reduction in solubility with decreasing temperature and on transforming from austenite to ferrite leads to the possibility of precipitation of vanadium compounds in steel. However, because of their relatively high solubilities vanadium compounds tend not to precipitate in austenite until relatively low temperatures are reached, usually in the presence of high levels of vanadium and nitrogen (or carbon) and normally in the presence of deformation. Figure 6

shows that in a 0.05%C – 1.2%Mn steel containing 0.115%V and 0.006%N, deformed continuously in compression, the “nose” of the precipitation start curve was below 900°C and the incubation period was relatively long at 30 seconds, or more. It has been suggested (11) that the start of such precipitation can be enhanced by the presence of a suitable substrate and Figure 7 (12) shows an example where VN has been precipitated as a cap on existing MnS. However, in most modern clean steels unless special rolling schedules and/or chemical compositions are used such a process is unlikely to make a significant contribution to precipitation of VN in austenite. Even when such special processes are adopted, it is likely that less than 10% of the available vanadium will precipitate, during rolling, in austenite.

This is advantageous as it means that, in the majority of steels, most, if not all, of the vanadium added to the steel is likely to remain in solution up to the start of transformation from austenite to ferrite. Thus, by far the majority of precipitation in vanadium containing steels takes place during (13, 14, 15) and after transformation,(4) giving rise to precipitation strengthening. The precipitates, which form during transformation, tend to form in rows, while those, which form afterwards, tend to be more randomly dispersed and to have both a smaller particle diameter and interparticle spacing than the row precipitates.

Examples of coarse austenitic precipitation, row precipitation and general precipitation in ferrite can be seen in Figure 8a and b. Some effects of this precipitation behaviour will now be considered.

3.4 Continuous Casting of Vanadium Containing Steels – Hot Cracking

During continuous casting, one of the most common types of defect, which can occur, is that known as transverse cracking. This defect occurs in microalloyed steels, particularly niobium

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microalloyed steels, and results from lack of ductility in the region of and just above the temperature at which austenite transforms to ferrite, during slab cooling. If a stress, such as that applied at the straightener of the continuous casting machine, is applied while the temperature of the steel is within the critical region, cracking can occur. This cracking is normally transverse to the casting direction and the cracks are frequently observed to be in association with reciprocation marks. Recent work,(16) using bend tests cooling down from the casting temperature, has examined the crack susceptibility of vanadium and niobium-containing steels. Figure 9a & b shows a comparison of the extent of cracking obtained in 0.1%V steel, with that obtained in 0.03% Nb steel. Clearly there is a greater degree of cracking in the latter case. These cracks were found to propagate mainly along the austenite grain boundaries. These boundaries were frequently, but not always, decorated with ferrite, indicating that the cracking, as noted above, was associated with the transformation. Figure 10 illustrates the effect of testing temperature on the length of the longest cracks for five steels including C-Mn, C-Mn-0.1%V, C-Mn-1%V-0.018%N, C-Mn-0.03%Nb and C-Mn-0.1%V-0.03%Nb steels. The shortest cracks were found to be in the C-Mn and C-Mn-0.1%V steels, indicating that these two steels should exhibit similar behaviour during continuous casting and that they should be able to be cast relatively crack-free. The longest cracks and the widest range of temperature over which cracking occurred were found to be in the 0.03%Nb steel, indicating that this steel would be the most difficult to cast and obtain crack-free slabs. The other two steels fell between these two extremes, indicating an intermediate degree of difficulty in obtaining crack-free slabs. It does appear, however, that vanadium containing steels are less likely to exhibit transverse cracking than those, which contain niobium alone.

3.5 Rolling of Vanadium Containing Steels

As previously noted, the solubility of vanadium carbide in austenite is significantly greater than that of vanadium nitride.(2-6) However, even in the case of the nitride the solubility in austenite is quite high. Figure 11 depicts the equilibrium solubility temperature for steels containing different vanadium and nitrogen contents. From this, for a steel containing 0.1%V and 0.02%N, a relatively high combination, the equilibrium solution temperature is only 1098°C. This relatively low solution temperature permits the use of energy efficient low soaking temperatures with little or no loss of precipitation strengthening capability in vanadium-containing steels. This has proved to be particularly important in the new process of thin slab casting, where the temperature in the equalisation furnace is typically in the range 1050-1150°C, but could also be important in the rolling of reinforcing bar where high furnace pushing rates are desirable.

One potential drawback of these relatively low solution temperatures, especially if soaking temperatures are not reduced, is that austenite grain coarsening can occur during reheating. As has been widely demonstrated, such coarsening can be controlled by an addition of approximately 0.01%Ti, although the fact that this will use nitrogen which would have been used for precipitation strengthening needs to be recognised. Furthermore, the possibility that the presence of TiN, as a substrate, will encourage precipitation of VCN, at high temperatures also needs to be taken into account.

Because vanadium compounds tend to remain in solution during rolling and vanadium only exhibits a small solute drag effect, vanadium steels recrystallise during rolling, even down to relatively low temperatures.(18) Consequently, by the process of deformation, recovery and recrystallisation on a falling temperature scale, it is possible to produce austenite grains, of

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high surface area/volume ratio, which transform to fine ferrite on subsequent cooling. In Figure 12 (19) the ferrite grain size of vanadium microalloyed steels, transformed from fully recrystallised austenite of high surface area/volume ratio, is similar to that obtained for the same steels transformed from austenite, which contained some deformation. Furthermore, at high surface area/volume ratio the grain sizes obtained for the vanadium-containing steels (4-5µm) were similar to those obtained for niobium steels transforming from deformed austenite.

A further effect of the recrystallisation behaviour exhibited by vanadium microalloyed steels is that their rolling loads are similar to those for carbon-manganese steels when measured in the same temperature range.(20,21) Examples of the effect of rolling temperature on the flow stress during rolling of <3mm thick strip of C-Mn and C-Mn-V steels are given in Figure 13. The flow stresses were calculated from rolling loads and were normalised to account for differences in true strain and strain rate. Also shown in Figure 13 is the well documented increase in flow stress, which occurs in Nb containing steels. In the figure, the flow stress at 860°C of the 0.03% Nb containing steel was some 84% higher than that of the 0.09% V containing steel. These differences were attributed to the differing recrystallisation behaviour of the steels during rolling. This difference in the recrystallisation behaviour of V and Nb steels is particularly important when rolling thin hot rolled coil.

Finally, because the recrystallised austenite grain size of vanadium-containing steels appears to exhibit little variation over a fairly wide range of temperature,(22) the properties of such steels are relatively insensitive to changes in finish rolling temperature.(23,24) Figure 14 (23)

shows that as the finish rolling temperature increased from 870°C to 1050°C, in steel with vanadium content 0.05-0.23% and nitrogen content 0.009-0.014%, there was no significant change in yield strength and a relatively small effect on impact transition temperature. In this work it was also suggested that the absence of Widmanstatten ferrite in vanadium-containing steels, finished rolled at high temperature, assisted in maintaining impact properties. Thus, in the rolling mill, vanadium microalloyed steels are relatively user friendly and their properties tend to be relatively insensitive to changes in rolling conditions.

3.6 Transformation from Austenite to Ferrite in Vanadium Containing Steels

The transformation of austenite to ferrite and the changes, which accompany this transformation, are amongst the most important factors to be considered during steelmaking. As has already been noted, in vanadium containing steels, precipitation of VCN can occur during this transformation. However, it is equally important to consider the effects of vanadium on the transformation and the transformation products, which result.

Figure 15a (25) shows the transformation start temperature for a steel containing 0.1% V with the relatively coarse austenite grain size of 150µm. A comparison with a C/Mn and a C/Mn/0.03% Nb steel is also given. Clearly the transformation start temperatures of the C/Mn and C/Mn/V steels are similar, that of the vanadium containing steel being, if anything, slightly the higher, while the transformation temperature of the C/Mn/Nb steel is 50 – 60°C below that of the other two steels. Furthermore, on differentiating the cooling curves of the three steels (Figure 15b) a clear difference between the C/Mn and C/Mn/V steels, on one hand, and the C/Mn/Nb steel, on the other hand, can be seen. Not only does the Nb steel transform at a lower temperature than the other two steels, its rate of transformation from austenite to ferrite is also faster. This has a significant effect on the microstructure which forms, particularly at the high cooling rates associated with quenching or with welding, with the C/Mn and C/Mn/V steels likely to form different, less acicular, microstructures from that of the C/Mn/Nb steel.

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Recently, there has been significant interest in the effects of vanadium on the formation of intra-granular ferrite during transformation. That second phase particles (TiO2) can act as substrates for the nucleation of such ferrite have been well known, in weld metals, for many years. The presence of intra-granular ferrite in the coarse grained heat affected zone of vanadium containing steel weldments has also been reported.(26) More recently (27,28) Japanese workers have indicated that, as has already been mentioned, by the of adoption of suitable rolling schedules which encourage VN precipitation in austenite, it is possible to promote the formation of intra-granular ferrite in vanadium-containing steels and that this, in turn, leads to grain refinement. Figure 16 (27) shows a comparison of the isothermal transformation behaviour at 700°C and 650°C, of a V-N containing and a niobium containing steel. The microstructure of the vanadium containing steel is finer than that of the niobium containing steel and, particularly at 650°C, there is a significantly greater proportion of intra-granular ferrite in the vanadium steel than in the niobium steel.

Another indication of the effect that vanadium and nitrogen levels can have on refining the ferrite grain size is shown in Figure 17 (29). Increasing the V.N product gave significant reduction in ferrite grain size. What proportion of this increase was due to normal grain refinement during rolling and what was due to intra-granular ferrite is, however, unknown. Nevertheless, that it is possible to produce fine ferrite grain size in vanadium containing steels is clearly shown.

A final effect of vanadium on transformation behaviour has also been noted in hyper eutectoid steels.(30) In the absence of vanadium in these steels, grain boundary cementite tends to form in continuous films around the prior austenite grain boundaries, with the centres of the grains being pearlite. This renders such steels as being particularly difficult to draw. The addition of vanadium appears to break up these continuous films (Figure 18), resulting in a series of cementite islands in a matrix of ferrite. This may have a beneficial effect on the drawability of rod and wire manufactured from such steel.

3.7 Strength and Toughness

In vanadium containing HSLA steels, the two main factors affecting strength and toughness are ferrite grain size and precipitation strengthening. As has already been demonstrated there is a clear relationship between vanadium (and nitrogen) level and ferrite grain size and the ferrite grain size remains reasonably constant over a fairly wide range of finish rolling temperatures. That increasing vanadium and nitrogen levels also increases both yield strength and UTS, via precipitation strengthening, can be seen in Figure 19.(29) In this figure the effects of vanadium and nitrogen have been combined into a V.N product. The yield strength increased from a level of 350MPa to a level of 600MPa as the V.N product increased from 0 to 0.002. In the same interval the level of precipitation, or dispersion, strengthening increased from 0 to 230MPa, giving an average increase of 115 MPa / 0.001% increase in V.N product. While this vector may be affected by parameters such as cooling rate and transformation temperature it is broadly correct for steels with a wide range of carbon content.

Refining the ferrite grain size tends to improve toughness in HSLA steels and a typical grain size vector is 11.5°C / d -½mm-½ .(4) On the other hand, increasing precipitation strengthening increases the impact transition temperature by about 3-4°C for every 10Mpa increase in yield strength.(31,32) The net effect of these changes on impact transition temperatures is shown in Figure 20. It must be emphasized that the Charpy vee-notch impact test pieces used to compile this diagram were 5mm x10mm in cross section i.e. sub standard, half size. Nevertheless it can be clearly seen that even with high levels of precipitation strengthening (600MPa YS, 230MPa precipitation strengthening) in a steel with a finish rolling temperature

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of 850 / 900°C and an aim coiling temperature of 600°C, the level of impact transition temperature was excellent at -50°C and that this improved down to –80 / -100°C as the level of precipitation strengthening reduced. It is also worth noting that, in vanadium containing steels, even with high nitrogen content, no strain ageing due to nitrogen is observed, providing the V:N ratio is maintained at or above the stoichiometric ratio of 4:1.(33)

3.8 Weldability

Arguably the two most important regions in a weld heat affected zone (HAZ) are the coarse grained region, close to the fusion boundary, and the intercritically reheated heat affected zone some distance from the fusion boundary. In these regions the properties which receive most attention are the hardness and toughness.

The hardness of the coarse grained HAZ, in the as welded condition, tends to increase with increasing vanadium level, Figure 21,(34) the increase being similar for weld cooling times between 800°C and 500°C of 12 seconds and 55 seconds.

Despite this increase in hardness vanadium can have a beneficial effect on toughness in the coarse grained HAZ (Figure 22).(34) Increasing the vanadium level from 0.0% to 0.16% led to a 50°C improvement in Charpy vee notch toughness and was accompanied by a modest, 10°C, increase in CTOD transition temperature.

In this work it was noted that the addition of vanadium appeared to promote intra-granular ferrite in the coarse grained HAZ microstructure. (Figure 23) The development of this microstructure is thought to have had a significant beneficial effect on toughness in the coarse grained HAZ.

In the intercritically reheated HAZ the addition of vanadium up to 0.1%V appears to have had little or no detrimental effect on toughness, although an addition of 0.05%V may be beneficial.(35) As can be seen in Figure 24 it is the presence of the intercritically reheated HAZ, which contains islands of M-A phase (Figure 25), which matters. Thus, it can be stated with reasonable confidence that in terms of HAZ hardness and toughness vanadium containing steels are readily weldable.

4. Conclusions

1. Most of the vanadium produced is consumed by the steel industry and this consumption has increased significantly in recent years.

2. The properties of vanadium containing steels are largely governed by the interactions, which take place between vanadium, nitrogen and carbon and the ways in which these interactions affect both precipitation and the transformation from austenite to ferrite.

3. Vanadium containing steels tend to be relatively easy to continuously cast, require low reheating temperatures, can be rolled with no significant increase in rolling load and their properties are relatively insensitive to finish rolling temperature.

4. Vanadium containing steels obtain their properties from a combination of fine grain size and precipitation strengthening, the presence of vanadium promoting both strong, tough, intra-granular ferrite as well as precipitation of vanadium carbides and nitride.

5. Vanadium containing steels have good weldability and the presence of vanadium can improve the heat affected zone toughness.

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REFERENCES

1. Mineral Industry Surveys, Vanadium, U.S. Geological Survey, March 2001. 2. G. Frohberg and H. Graf, Stahl und Eisen, 1960, 80, p. 539. 3. L. A. Erasmus, J. Iron & Steel Inst., 1964, 202, p. 32. 4. K. J. Irvine et al, J. Iron & Steel Inst., 1967, 205. P. 161. 5. K. Narita, Trans. ISIJ, 1975, 15, p. 145. 6. K. Bungardt et al, Archiv. f. d. Eisenhuttenwesen, 1956, 27, p. 6. 7. H. Sekine et al, Trans. ISIJ, 198, 8, p. 101. 8. K. A. Taylor, Scripta Met. & Mat., 1995, 32, p. 7. 9. H. Ohtani and M. Hillert, Calphad, 1991, 15, p. 25. 10. Akbed and J. J. Jonas, HSLA Steels Technology & Applications, 1983, p. 149. 11. H. Satoh et al, 41st MWSP Conf., 1999, p. 911. 12. S. Zajac, Swedish Institute for Metals Research, Work for Vanitec. 13. A. D. Batte and R. W K Honeycombe, JISI, 1973, 211, p. 184. 14. R. W. K. Honeycombe, Trans. AIME, 1976, 7A, p. 915. 15. F. A. Khalid and D. V. Edwards, Materials, Science & Technology, 1993, 9, p. 384. 16. D. N. Crowther et al, Microalloying in Steels, 1998, p. 469. 17. B. Mintz et al, Intl. Materials Review, 1991, 36, 5, p187. 18. L. J.Cuddy, Thermomechanical Processing of Microalloyed Austenite, 1981, p. 129. 19. T. Siwecki et al, Microalloying ’95, 1995, p. 197. 20. P. Repas, Microalloyed HSLA Steels, 1998, p. 9. 21. M. de Lisi et al, Journees Siderurgiques, ATS, Paris, Dec., 1992. 22. Yang-Zheng Zheng et al, HSLA Steels Technology & Applications, 1983, p. 85. 23. J. M. Chilton & M. J. Roberts, Met. Trans. A, 1980, 11A, p. 1717. 24. P. S. Mitchell et al, Intl. Symp., Low Carbon Steels for the 90’s, p. 339-341. 25. D. N. Crowther, Corus Group (UK), Work for Vanitec. 26. A. M. Sage and P. H. M. Hart, Metallurgy, Welding and Qualification of Microalloyed

(HSLA) Steel Weldments, AWS, Miami, 1990, p. 806 27. T. Kimura et al, Int. Sym. on Steel for Fabricated Structures, ASM, 1999, p. 165. 28. H. Satoh et al, 37th MWSP Conf. Baltimore, Oct 1999, p. 911. 29. Y. Li et al, University of Strathclyde, Work for Vanitec. 30. K. Han et al, Mat. Sci. Eng., 1995, A190, p. 207. 31. T. Gladman et al, Microalloying 75, New York, 1976, p. 72. 32. K. A. Taylor and S. S. Hansen, Met. Trans. A, 1991, 22A, p. 2359. 33. D. Russwurm & P. Wille, Microalloying’95, Pittsburgh, USA. June 1995, p. 377. 34. P.S. Mitchell et al, Microlloying’95, Pittsburgh, USA, June, 1995, p. 149. 35. Y. Li et al, ISIJ International, 2001, Vol. 41, p. 45.

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Fig. 1 The consumption of vanadium, by end use, in the USA, in 2000.

Fig. 2 The consumption of vanadium, by end use, for Germany, Japan and USA.

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Fig. 3 Vanadium consumption since 1970.

Fig. 4 Change in specific consumption of vanadium since 1970.

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Fig. 5 Solubility of microalloy carbides and nitrides in a) austenite and b) ferrite.

b)

a)

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Fig. 6 Dynamic PTT Curves for 0.115% V and 0.035% Nb Steels

Fig. 7 SEM micrograph showing VN precipitated as a cap on MnS. (Zajal)

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Fig. 8 Examples of precipitation a) coarse precipitation in austenite and row precipitation

which occurred during transformation from austenite to ferrite, b) general precipitation in ferrite.

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a)

b)

Fig. 9 Bend tests a) 0.1% vanadium steel tested at 850°C, b) 0.03% niobium steel tested at 792°C.

Fig. 10 Effect of temperature on the length of the largest crack observed during hot ductility testing on a range of steels

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Fig. 11 The equilibrium solution temperature of vanadium nitride in austenite

Fig. 12 Dependence of ferrite grain size on the austenite grain boundary area per unit volume. Data points are for Ti-V and V microalloyed steels. Curves refer to Nb microalloyed steels. (Siwecki)

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Fig. 13 Mean flow stress as a function of hot rolling temperature at the finishing stand. (De Lisi)

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Fig. 14 Effect of finish rolling temperature on mechanical properties of vanadium steels. Soak temperature 1175°C. (Chilton and Roberts)

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Fig. 15 Effect of vanadium and niobium on transformation behaviour a) transformation start

temperature, b) transformation temperature and rate of transformation.

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Fig. 16 The number density of ferrite grains during isothermal transformation at 700°C and

650°C.

Fig. 17 The influence of VxN on the ferrite grain size, (end cool temperature, 550-650°C).

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Fig. 18 TEM bright field image showing the grain boundary carbides of 1.05%C-0.60%Mn-

0.23%Si-0.15%V steel partially transformed at 650°C for 3 seconds and then quenched. (Edmonds)

Fig. 19 The effect of vanadium and nitrogen on yield strength and UTS.

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Fig. 20 Relation between impact transition temperature and VxN.

Fig. 21 Effect of vanadium on the change in maximum HAZ hardness in the as-welded

condition.

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Fig. 22 Effect of vanadium on the fracture toughness of 0.12%C-1.6%Mn steels multipass welded at 2 KJ/mm (∆8/5 = 12 secs) in the as-welded condition.

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C-Mn

C-Mn-0.1%V

C-Mn-0.03%Nb Fig. 23 Simulated HAZ microstructure, cool rate = (800-500) 7°C/s, X500.

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Fig. 24 The effect of hardness on the 50J ITT of the ICGCHAZ in V and Nb steels.

Fig. 25 Typical SEM micrograph showing islands of M-A phase.

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A New Role for Microalloyed Steels − Adding Economic Value

Michael Korchynsky

Consultant in Metallurgy. U.S. Vanadium Corporation A Subsidiary of Strategic Minerals Corporation, Pittsburgh. Pennsylvania.

Abstract

Microalloyed (MA) steels have matured during the past 40 years into an important class of high-strength structural materials. Their cost-effectiveness has been enhanced by the growth of electric-arc-furnace (EAF) steelmaking and the thin-slab-casting process. A recent project involving an ultra-light-steel auto body (ULSAB) concluded that high-strength-steels are the materials of choice for the automotive industry. This project showed that replacing cheaper carbon steels with high-strength steels allowed automakers to reduce the weight of an auto body at the same or at potentially lower costs. The same economic principles can be applied to other applications.

The strengthening effects of vanadium make microalloyed steels particularly suited for high-strength-steel applications. By effectively combining grain refinement and precipitation hardening, vanadium maximises the strengthening process and is compatible with current steel-processing technology.

To dramatise the cost effectiveness of these high-strength-steels in potentially new applications, a series of demonstration projects is needed involving the cooperation of steel producers, fabricators, and users. In many applications, the decision to replace plain-carbon steel with higher strength vanadium-bearing microalloyed steel can be shown to improve the profitability of both the steelmaker and the steel user.

1. Introduction: Use of Microalloyed Steels Reduces Costs

The development of microaIloyed steels, including their alloy design, processing, and applications, covers the last four decades.1-4) During this period, microalloyed, high-strength, low-alloy (HSLA) steels became an indispensable class of structural steels. Their ability to achieve final engineering properties in as hot-roIled conditions eliminated the need for heat treatments, such as normalising. Yield strengths ranging up to 550 to 600 MPa can be attained through small additions (less than 0.1%) of selected carbonitride formers without requiring costly alloying elements. The resulting cost-effectiveness of microalloyed steels led to the successful displacement of heat-treated steels in applications such as truck side rails and telescoping crane booms. Recent technological developments in steel melting and hot rolling further reduced the cost and enhanced the competitiveness of microalloyed steels.

Despite these improvements, the total consumption of microalloyed steel is currently estimated to be only 10 to 15% of the world's steel production (i.e., 80 to 120 million tons per year). This tonnage is about evenly distributed between flat and long products. As a result, there is plenty of room for growth. A major jump in the usage of microalloyed steels should have strong economic benefits for both steel producers and steel users.

2. Microalloying: Complimentary Strengthening Mechanisms

Hot-rolled plain-carbon steel is the most popular material used in construction. Its strength can be increased by raising the carbon content. In fact, its strength is proportional to the carbon equivalent (CE) which is essentially the combined effect of the carbon and manganese content of steel, based on the formula: [CE = %C + %Mn/6]. While raising the

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carbon equivalent increases strength, it also drastically reduces other engineering properties such as ductility, toughness, and weldability. Since welding is irreplaceable as a method of fabrication, the carbon-equivalent mechanism of steel strengthening cannot be used in many applications requiring weldability.

As shown in Figure 1, the success of microalloyed steels is due to complimentary strengthening mechanisms, specifically grain refinement and precipitation hardening.5) Precipitation hardening increases strength but may contribute to brittleness. Grain refinement increases strength but also improves toughness. As a result, grain refinement counteracts any embrittling caused by precipitation hardening.

In practice, grain refinement can be achieved during hot rolling by the interaction between microalloying elements (niobium, vanadium, or titanium) and hot deformation. During the allotropic transformation, ferrite nucleates on austenitic grain boundaries. Maximum grain refinement can be achieved by increasing the austenitic grain-boundary area. This can be accomplished by either producing fine grains of austenite through repeated recrystallization between passes6) or by flattening non-recrystallized austenite grains into "pancakes". The first process is generally used for vanadium steels and the second for columbium (niobium) steels.7)

Grain refinement may be further enhanced by accelerating cooling after the completion of hot rolling. The undercooling of austenite enhances the rate of ferrite nucleation and slows down the rate of growth. A combination of these two factors contributes to the formation of smaller grains.

Significant strengthening is obtained by the precipitation of microalloying elements appearing as carbonitrides (or carbides) in ferrite.8.9) Since their solubility in ferrite is much less than in austenite, there is strong supersaturation which provides the driving force for precipitation. The most desirable are those microalloys, which contribute to both grain refinement and precipitation hardening. The combined effect of these two strengthening mechanisms may provide as much as 70% of the yield strength, accounting for the remarkable cost-effectiveness of microalloyed steels.

Because these two dominant strengthening mechanisms operate in microalloyed steels, their carbon content (or CE) may be very low. A yield strength of 550 MPa can be obtained in a steel containing only 0.04 to 0.06% carbon.10) This low-carbon content contributes to excellent weldability.

3. A New Function for Hot Rolling: Optimising Material Properties

Traditionally, the main objective of the hot-rolling process was to change the geometry of a slab or billet to meet the dimensions of the final product. For this purpose, the temperature of rolling was not well controlled, with the unwritten rule being "the hotter, the better".

Accomplishing the "miracle" of converting ordinary carbon steel into a sophisticated HSLA steel requires an understanding of the evolution of the austenite microstructure during hot rolling. At high temperatures, the size of a recrystallized grain after each deformation pass depends on the initial grain size, temperature, and the amount of deformation.11) The tendency of grains to coarsen between passes can be prevented by precipitated particles within the grain boundaries. Finely dispersed titanium nitrides (TiN) formed by titanium additions as low as 0.005 to 0.007% effectively prevent grain coarsening. When the rolling temperature is low enough to prevent recrystallization, austenite is flattened into a "pancake". The temperature at which this occurs depends on the type of microalloy. Recrystallization is suppressed at a much higher temperature in niobium (columbium) steels than in vanadium steels.7)

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Grain refinement may be further enhanced by the intra-granular nucleation of ferrite in austenite.l2) The precipitation of vanadium nitride in austenite provides the most effective intra-granular nucleation of ferrite. The highest rate of carbonitride precipitation in ferrite occurs at 600°C, which is the customary sheet coiling temperature.

4. Making Nitrogen A Friend, Not A Foe

Two new developments in steelmaking and steel processing − the growth of the electric arc furnace (EAF) and processing by thin slab casting have contributed to further cost reduction in the production of microalloyed steels.13)

EAF steelmaking is growing rapidly worldwide because it is less capital intensive than the conventional processes used by integrated steel producers. Virtually all new steelmaking capacity added either by mini-mills or integrated producers, uses electric arc furnaces. Soon, 50% of the world's steelmaking or about 400 million tons annually will be made in these facilities.

In a scrap based EAF practice, the nitrogen content is 70-100 ppm or 2 to 3 times higher than that typical of the basic oxygen or BOF practice. The nitrogen level of steels made in an EAF can be reduced by modifying the slag practice or changing the feed stock. Both these methods can increase costs.

Free nitrogen in solution in ferrite, has serious detrimental effects such as aging and brittleness. During concasting, excessive nitrogen may increase possible transverse or longitudinal cracking. Fears are also frequently expressed about the detrimental effects of nitrogen on weldability.

However, the harmful effects of nitrogen may be neutralised by nitrogen binding elements which acting as scavengers, remove nitrogen from solid solution in ferrite. Aluminium and titanium are effective scavengers; however, niobium (columbium) is not an effective nitrogen-binding element in high strength, low alloy steels. In niobium steels, niobium carbonitrides are only present when the carbon to nitrogen ratio ranges between 1:1 and 4:I. Thus, the effect of niobium depends on the nitrogen content of the steel.

Among the various microalloying elements, vanadium has a unique dual effect on nitrogen.14) Vanadium not only neutralises nitrogen by forming VN compounds but also uses nitrogen to optimise the precipitation reaction. Enhanced nitrogen increases the supersaturation in ferrite and promotes a more active nucleation of V(C,N) particles, as shown in Figure 2. Consequently, the interparticle distance is reduced (Figure 3) and the strengthening effect of precipitation is increased. In the presence of nitrogen, less vanadium is needed to achieve the desired yield strength. As a result, vanadium effectively converts nitrogen, previously considered an impurity, into a valuable alloy that helps strengthen steel, as shown in Figure 4.

The pioneering efforts of the Nucor Steel Corporation15) in commercialising thin slab casting have dramatically changed the economics of hot band production. The revolutionary effect of this new process can be compared to two previous developments which have changed the economics of steel production: the switch of steelmaking from open hearths to a basic oxygen (BOF) converter and the replacement of ingot casting by continuous casting.

The thin slab casting process converts in-line liquid steel into a marketable product.l6) The process incorporates a series of steps that contribute to either cost reductions or to property improvements. The rapid solidification in the mould accounts for the small size of globular inclusions, which do not elongate during hot rolling. This promotes isotropic properties, such as bendability, in longitudinal or transverse directions. Near net-shape dimensions of the slab

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(50 - 70 mm) facilitate rolling to an aim thickness of 1 mm (or less), allowing hot rolled steel to economically replace cold rolled sheet. In-line processing permits the slab to be directly charged into the rolling mill, contributing to energy savings. The amount of deformation per pass is 2 to 3 times higher than that on a hot strip mill rolling thick slabs. Excellent microstructure and properties are obtained in a 15-mm thick strip for a total deformation of less than 4:1.

Because of lower hot rolling costs, the market share for hot bands produced by thin slab casting is being increased at the expense of high cost integrated producers. In developing the concept of replacing carbon steels with microalloyed steels, we will limit our choice initially to strip made by thin slab casting technology.

5. Ultra-Light Steel Auto Body: Quantifying the Economic Benefits of Microalloyed Steel

The pressure to reduce the weight of automobiles and the ever-present threat from light-weight materials such as aluminium or magnesium led to the creation of an international consortium of steel producers whose goal was to produce a lighter-weight auto body. Over 30 steel producers jointly sponsored an ambitious project: Ultra Light Steel Auto Body (ULSAB).17) The objectives of the project were three-fold: (1) design a stronger and safer auto body, compared to best models available, (2) lower the weight, and (3) keep costs the same or less than auto bodies being built today. All three goals have been successfully attained. Three factors contributed to the success of the project: novel design concepts, material selection, and new fabrication methods.

In the area of materials, the most important change was the replacement of inexpensive carbon steel with higher value HSLA steels. More than 90% of the ULSAB structure used high strength steel ranging in yield strength from 210 to 420 MPa. One half of the steel used had 350 MPa yield strength. Both cold and hot rolled sheet, 0.65 to 2.0 mm in thickness, have been used. The use of steel stronger than 420 MPa was minimal.

HSLA steel emerged as the material of choice for modem automobile design. For a cost conscious automobile industry, the use of more expensive HSLA steel was found to be economically attractive as a replacement for cheaper carbon steel.

Years ago, a GM executive made a controversial statement: “What is good for General Motors is good for America." Today, we may paraphrase this slogan: "What is good for ULSAB may be good for many steel processing industries." The ULSAB project demonstrated that the competitiveness of steel hinges on the following parameters: engineering properties and fabricability, weight reducing potential, and cost.

Thanks to technology advances and the successful adaptation of a series of cost reduction steps, microalloyed steels have all the necessary attributes to successfu11y replace inefficient and often higher cost carbon steels in such areas as construction, transportation, and machine building.

6. Weight Reduction: The Key to Adding Economic Value

Microalloyed, high strength, low alloy steels may have yield strengths that are 2 to 3 times higher than hot rolled, weldable carbon steels. The weight reduction achievable through substitution depends not only on the difference in strength but also on the mode of loading. For straight loading in tension, the weight reduction is proportional to the difference in strength. An increase in yield strength by a factor of two may reduce the weight of steel by two - a situation found in concrete reinforcing bars. The range of weight savings is shown in Figure 5.

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For other types of loading (e.g., bending), a two fold increase in strength may contribute to a weight reduction of 34% or more. Considering a safety factor, one may as same as the HSLA steel, being twice as strong as carbon steel, may reduce the weight by at least 25%. The following simplified calculation illustrates the economic value obtained through substitution:

It is evident that the substitution is economically attractive for both the producer and user. The producer may enjoy twice as much profit for the value added microalloyed steel. The user pays $39 less for the material. His additional benefits include ease of fabrication, improved overall properties (e.g., toughness, ductility), and lower transportation cost. The range of cost savings is shown in Figure 6.

7. Vanadium: Offering the Lowest Cost Per Unit of Strength

In selecting the most economical microalloyed steel, the following factors should be considered: (1) alloy design that maximises the strengthening effect of the two cost effective mechanisms grain refinement and precipitation hardening), (2) low processing cost affected by reheat and finishing temperatures, and (3) compatibility with the typical nitrogen content of electric furnace steels.

Of the three microalloying elements - niobium, vanadium, and titanium - only vanadium contributes to cost reductions in all three areas listed above. By incorporating vanadium in an alloy design, both grain refinement and precipitation hardening will be fully utilised, providing up to 70% of the total yield strength. In addition, vanadium permits the use of the lowest cost hot rolling practice. Because of the high solubility of vanadium nitride (VN) in austenite, a low reheating temperature (1150°C) may be used. This contributes to energy savings. At the same time, the finishing temperature may be high (900 to 1000°C), since

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grain refinement is obtained by repeated recrystallization. Finally, vanadium not only neutralises the ill effects of "free" nitrogen; it uses this "impurity" as a valuable alloying element - a unique characteristic among microalloying elements.

There are virtually unlimited sources of vanadium on earth.18) Vanadium is extracted from vanadium bearing ores (Australia), as a by-product of steel production from iron ore containing both vanadium and titanium (South Africa, Russia, and China), from spent catalysts (U.S.A.), and from oil products (U.S.A.). At present, the existing industrial capacity for vanadium production exceeds the demand. With new sources of supply on the horizon, shortages are not likely to occur in the foreseeable future.

8. Unlocking the Commercial Benefits

Based on the example in the ULSAB project, it can be seen that substituting a higher value steel in an existing application brings economic and performance benefits. One way of achieving this objective is through the use of demonstration projects. These demonstrations must be conducted on carefully selected products with a full understanding of the potential benefits for the steel producer, fabricator, and the user.

As an example, microalloyed steel might be substituted for commonly used carbon steel in the production of spirally welded water pipe. For steel producers, this substitution would increase their share of value-added products, which command a higher profit margin. For the fabricator, the ease of welding a low carbon (0.04-0.06% C) microalloyed steel provides a strong incentive in the form of reduced labour costs. In addition, the user will benefit not only from lower material cost but improved strength and fracture resistance for safe operation.

To deliver these benefits, each demonstration project must fully document and evaluate the service performance and provide a detailed cost analysis. The project should also assure that mandatory standards and specifications are met. Finally, the results of the demonstration project should be incorporated into pertinent engineering specifications and brought to the attention of designers and engineers through educational and promotional efforts.

9. New Opportunities for Producers and Users

The selection of high strength steels as the material of choice by the automotive industry suggests new opportunities for microalloyed steels in a variety of applications. Their high strength compared to hot rolled carbon steels offers an opportunity for significant weight reduction. This lower weight more than offsets the slightly higher unit cost of microalloyed steels, adding economic value to both steel producers and steel users. For steelmakers, selling more value-added high strength steels improves their competitive position against light-weight materials and allows increased profit margins. For users, value-added high strength steels offer lower material costs, better product performance, and reduced fabrication and transportation costs.

The feasibility of satisfying all engineering needs with less steel is also advantageous for the national economy. It reduces the pressure of capital intensive demands for replacing obsolete traditional steelmaking facilities. In addition, any shortfall in the capacity of integrated producers will be balanced by cheaper and more flexible electrical furnaces.

The path from a theoretically attractive concept of substitution to commercial reality requires a dedicated effort, mainly by the steel industry. Nevertheless advances achieved during the past 40 years in the science, technology, and applications of microalloyed, high strength steels make them ready for this economic challenge.

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In fact, the importance of weight reduction and greater emphasis on value-added products are squarely in line with the trends of the steel industry as reflected in the following quotations:

"The steel industry's weight paring campaign is an investment in the future." (Al Wrigley, American Metal Market, June 26, 2000)

"Innovative technology will help us to produce the high-value-added steels, critical to our success and the success of our customers." (Paul J. Wilhelm, President, U .S. Steel, New Steel, June 2000, page 56)

10 Conclusion

The new role of microalloyed steels in adding economic value through replacement of hot rolled carbon steels is feasible today. Potential economic benefits to the steel industry and its customers could reach billions of dollars annually.

The competitive advantages of microalloyed steel compared to hot rolled carbon steel include superior fabricability, weight reduction by at least 25%, and lower overall cost. To fully exploit these benefits requires vision and dedication on the part of the steel industry and its customers. The economic advantages to both producers and users - typical for a "win- win" situation, provide a strong incentive for initiating a drive in this direction.

REFERENCES

1. Symposium: Low-Alloy, High-Strength Steels; Nuremberg; May 21-23, 1970; The Metallurgy Companies.

2. International Conference: Proceedings; Micro- alloying '75; Washington, D.C.; Union Carbide Corporation.

3. International Conference: Technology and Applications of High-Strength Steels; Philadelphia, Pa.; October 1983; American Society for Metals.

4. International Conference: Microalloying '95; Pittsburgh, PA; June 11-14, 1995; Iron and Steel Society.

5. F. B. Pickering: ""Physical Metallurgy and the Design of Steel"; Applied Science Publishers, London (1978).

6. T. Siwecki et al.: "Recrystallization Controlled Rolling of HSLA Steels"; Microalloying '95; June 11-14, 1995; Iron and Steel Society.

7. L. J. Cuddy: "The Effect Of Microalloy Concentration of Austenite during Hot Deformation"; Conference Proceedings: Thermo-mechanical Processing of Microalloyed Austenite; The Metallurgical Society, Pittsburgh, PA, (1981).

8. S. Zajac et al.: "The Role of Nitrogen in Microalloyed Steels"; Microalloying '95; June 11- 14, 1995; Iron and Steel Society.

9. R. Lagneborg et al.: "Influence of Processing Route and Nitrogen Content on Microstructure Development and Precipitation Hardening of V - Microalloyed HSLA Steels"; Conference Proceedings: Thermomechanical Processing of Microalloyed Austenite; The Metallurgical Society, Pittsburgh, PA, (1981).

10. M. Korchynsky: "New Steels for New Mills"; Scandinavian Journal of Metallurgy; 28, (1999), pp. 40- 45.

11. T. Siwecki and G. Engberg: "Recrystallization Controlled Rolling of Steels;" Proceedings: Thermomechanical Processing in Theory, Modelling, and Practice; ASM International, Stockholm, (1996).

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12. T. Kimura et al.: "Heavy Gauge H-Shapes with Excellent Seismic-Resistance for Building Structures Produced by Third Generation IMCP;" International Symposium Proceedings: Steel for Fabricated Structures; Cincinnati, Ohio; ASM, (1999).

13. M. Korchynsky: "Cost Effectiveness of Micro- alloyed Steels;" International Symposium Proceedings: Steel for Fabricated Structures; Cincinnati, Ohio; ASM, (1999).

14. R. Lagneborg at al.: "The Role of Vanadium in Microalloyed Steels;" Scandanavian Journal of Metallurgy, 28, (October 1999).

15. P. J. Lubensky et al.: "High Strength Steel Processing via Direct Charging Using Thin Slab Technology;" Microalloying '95; June 11-14, 1995; Iron and Steel Society.

16. M. Korchynsky and S. Zajae: "Flat Rolled Products from thin-Slab Technology; Technological and Economical Potential;" Proceedings: Thermomechanical Processing in Theory, Modelling, and Practice; ASM International, Stockholm, (1996).

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Fig. 1 The strength of hot-rolled steel depends on the five mechanisms shown above. In high

strength low alloy steels, microalloying elements such as vanadium and nitrogen provide up to 70% of the strength by contributing to grain refinement and precipitation hardening.5)

Fig. 2 Increasing the nitrogen content promotes nucleation, forming smaller vanadium-nitride particles.8)

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Fig. 3 Reducing the particle diameter of precipitates from 4 to 2 nm gives eight times the number of precipitates in a given volume of steel. The larger number of small precipitates gives more efficient strengthening by reducing interparticle spacing,

Fig. 4 In vanadium steels, nitrogen acts as a valuable alloy that helps increase yield strength.14)

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Fig. 5 The weight of components can be reduced substantially by substituting high-strength steel for low-strength carbon-manganese steel (SSAB Sheet Steel Handbook).

Fig. 6 Because of the cost effectiveness of microalloyed steels, the weight reduction more than offsets the difference in prices between microalloyed and carbon-manganese steels (SSAB Sheet Steel Handbook).

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Development and Use of Vanadium in Micro-Alloyed Reinforcing Bar

D. Russwurm Dr.-Ing. habil.

Managing Director of German Institute for Reinforcing Steel

Introduction: Importance of Reinforcing Steel

Reinforcing steel is one of the main steel products worldwide produced and used. In

1998 the world consumption of reinforcing steel was about 95 Mio tons. Compared

with the total steel production at that time about 13 % of steel was dedicated for the

reinforcement of concrete. A comparison between cement and reinforcing steel for

the globe and for Europe is shown in Fig. 1.

0

200

400

600

800

1000

1200

1400

1600

1989

1990

1991

1992

1993

1994

1995

1996

1997

Prod

uctio

n, M

io to

ns

World Cement World Reinforcing Steel (Bars)

020406080

100120140160180200

1989

1990

1991

1992

1993

1994

1995

1996

1997

Pro

duct

ion,

Mio

tons

Europe EU CementEurope EU Reinforcing Steel (Bars)

Fig.1. World and European Production of Cement and Reinforcing Steel.

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By the way the consumption of cement – the second of the construction materials in

the compound system “steel – concrete” is round about 1,6 billion tons. This results

in a portion of ca. 20 kg of reinforcing steel per m3 of concrete.

This makes it clear that reinforcing steel is strongly depending on the amount of

reinforced and prestressed concrete which is without doubt the leading construction

method in all continents of the world.

But the consumption differs from country to country. Fig. 2 shows the consumption

per capita and year of reinforcing steel of some European countries.

0

20

40

60

80

100

120

S F GB NL N B DK P E D Irl I GR SF A L CH

Con

sum

ptio

n of

Rei

nfor

cing

Ste

el P

er C

apita

, Kg kg Meanvalue

The differences are caused by the design-codes – some are reinforcement friendly,

some not – the topographical situation – expensive constructions e.g. in traffic

buildings – and the preference for a construction method.

Generally can be said that for different reasons as reinforced concrete will remain for

this and the next generation will remain the leading construction method. But will be

steel also remain the dominating reinforcement. I am sure you have heard a lot of

non-ferrous reinforcements as carbon and glass or a Aramid.

The advantages of steel are evident. Some of them are listed up in the following

figures.

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“Reasons for Steel as Optimal Reinforcement for Concrete”

• Load transfer has to be performed from concrete to reinforcement and vice versa over a limited development length

• The reinforcement has to have a high modulus of elasticity in order to gain a high degree of stiffness for the construction as a whole.

• The interaction between concrete and reinforcement must be in terms of chemical and physical phenomena free of disadvantages.

• The reinforcement must have the appropriate delivery forms, shape and length which fit with the various constructions.

• Spacious interfaces, e.g. sheets are less suitable in comparison with rod-formed elements which are offered in a great variety and a large range to match the calculated values of reinforcement section.

• The reinforcement has also to match with the shape of the construction; it must be flexible and easily bendable.

• The reinforcement must be capable to be joint, either by overlap or particular joining techniques as welding, forming mechanical connections.

• The reinforcement has to resist without significant deterioration of damage the rough conditions during transport, storage, bundling and placing on job site. Minor damages should not reduce significantly the performance characteristics.

• The reinforcement of prestressed constructions (prestressing steel) has to assure that no sudden and brittle collapse due to corrosion attack takes place.

• The reinforcement has to provide to the construction a sufficient fatigue resistance because concrete is not resistant to dynamic loading.

• The reinforcement has to attribute to the construction a sufficiently high ductility behaviour.

• The reinforcement has to resist shear forces as well as tension and compression forces.

• The relaxation of prestressing elements has to be suitable low. • The reinforcement must offer its performance characteristics in a sufficient

range of temperatures (- 60 °C to + 80 °C). In case of extreme low or high temperatures (fire) the behaviour of the reinforcement must be predictable.

• The quality level of the reinforcement has to be usually such to be able to compensate minor imperfections in the execution of the construction.

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Some disadvantages of the non-ferrous products are also obvious (Fig. 5):

Fig. 3. Comparison Steel – Non Ferrous Products

Reinforcing Steel

Non Ferrous Reinforcement

- absolute brittle material

- low-E-modulus

- not bendable

Nevertheless what we need is

- a high quality steel

- a steel fit for usual handling at the manufacturers and on job site

- and a steel which is perfectly adjusted to all design methods of the civil

engineers.

About this steel I want to speak to you today. The structure of my presentation is

the following:

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Development and Use of Vanadium in Micro-Alloyed Reinforcing Bar Advantages of High-Strength-Steel Introduction: Importance of Reinforcing Steel Mainly used types of reinforcing steel Important performance characteristics Production methods

Cold working Heat treatment Micro-alloying

Economical Aspect Metallurgical Aspects Economic advantages of high strength steel Savings due to high yield reinforcement Ecological Aspects Summary

Mainly used types of reinforcing steel

To describe very briefly the mainly used typed of reinforcing steel I refer to the

leading standards on this field. These are:

USA: ASTM 615 / ASTM 706

UK: BS 4449

Germany: DIN 488

Europe: Draft of prEN 10080 / prEN 10081

USA:

Concerning reinforcing steel the US are not at the frontier of the development. The

standard includes two grades with the minimum yield levels of

300 N/mm2 and 400 N/mm2.

Generally the bars are deformed that means with high bond properties. This grades

are not weldable and have no specific ductility requirements apart from bendability

and no specifications concerning fatigue.

The second standard which covers weldable grades is up to date also with the

grades (yield strength) 550 / 540 and 420 N/mm2. The consumption of these grades

is far below ASTM 615 grades.

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United Kingdom:

The BS 4449 gives the specifications for the British reinforcing steel. The main grade

is a steel with minimum yield of 460 N/mm2. This steel has to be weldable, is

provided with high bond and can be used under cyclic loading.

Germany:

The German standard DIN 488 knows two grades with yield strength 420 and 500

N/mm2. Since round about 16 years exclusively the minimum yield 500 N/mm2 is

used with great success in all applications as

- bars

- coils

- wire fabric

- lattice girders.

Other European Countries:

The range of yield strength is spread between 400 to 700 N/mm2. nearly all grades

are weldable.

European Standardisation:

The work on European Standard for reinforcing steel is still going on. Due to an

intervention of the European Commission we will probably get a harmonised

standard without defining steel grades.

As chairman of the committee ECISS TC 19 I did not succeed to implement only 3

grades with yield strength 500, respectively 450, which would have been a great

success in comparison with the existing 47 grades in the different member states.

A summary shows what the up-to-date steel grades have:

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Standard Yield Strength

N/mm2

ASTM 615

ASTM 706

BS 4449

DIN 488

Europe (general)

400

550 / 420

460

500

400 to 700

Weldable steel (generally in Europe)

Fatigue resistant (generally in Europe)

I mention this important features because they play an important role for the

production routes later have to be discussed.

Important performance characteristics

The reinforcing steels have no use for themselves they depend from their application

in concrete.

The usual design codes differ under the aspect of using the reinforcement according

to their systems of structural analysis. There are different methods of structural

analysis which require different properties of the reinforcement (Fig. 8):

linear elastic analysis linear elastic analysis with limited moment redistribution plastic methods of analysis non-linear analysis earth-quake design with energy dissipation fatigue resistance

Corresponding to the applied method the requirements of the reinforcing steel differ.

Additionally the serviceability limit state of a building has to be taken into account as

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- crack formation, limitation of crack width

- limitation of deformations

Further on durability of a construction, in modern words the sustainability touches the

required properties of steel.

In order to obtain a good rationalisation in producing solid cages of reinforcement

and to perform splices weldability is of greatest interest.

How the requirements interact with the desires of the designers is shown in Table I:

Basis of Demands Resulting Requirements of Reinforcing Steel

Competitiveness of reinforced concrete

Yield as high as possible

Structural analysis: linear elastic and moment redistribution plastic methods non linear earthquake design fatigue

high yield high yield and ductility high yield and high ductility high yield and high ductility sufficiently high yield in combination with highest mobility requirements high yield and fatigue resistance

Serviceability limit state Yield strength adjusted to crack control and acceptable deformations

Sustainable constructions Corrosions resistance

Handling of reinforcement Bendability, Rebending, Weldability

Some requirements have to be explained and defined.

Yield strength is obviously clear. Ductility needs to be clarified.

In Fig. 4 a stress-strain-curve gives the answer:

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Fig. 4. Stress-strain Curve showing Ductility

Designers use two parameters to define ductility Rm/Re and Agt.

Now the different reinforcing steels provide various values.

Among them 3 groups have been formed (Table II):

Table II

Rm/Re [-]

Agt [%]

Type of steel

Normal ductile steel 1,05 2,5 cold worked

High ductile steel 1,08 5,0 hot rolled (heat treated, micro-alloyed)

Earth-quake steel 1,15 (Re,act/Re??? < 1,3)

8,0 hot rolled (heat treated, micro-alloyed)

Fatigue resistance is defined by an SN-line (Fig. 5):

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Fig. 5. Fatigue Resistance Defined by SN Line

Nx log N

m=k1

m=k2

log ∆σ

∆σRsk constNm =⋅∆σ

Normally only one point (point of intersection) is given in the standards, expressed by

the stress range and the number of load cycles. Sufficiently high values for

reinforcing steel are round about 200 N/mm2 at 2.106 load cycles (see Table III):

Product Form Bars and de-coiled rods

Wire Fabrics Requirement of Quantile Value

Fatigue stress range (N.106) *) (N/mm2) with an upper limit of not more than 0,6 fy

150

100

10

*) This fatigue requirement is not required for predominantly static loading. If higher values are shown by testing and approved by an appropriate authority, the design values (table 6.3) may be modified. Such testing should be in accordance with EN 10080.

Weldability for reinforcing steel is a complex problem: the number of splices are

limited: but the types are very sensible for load transfer. The production routes for

reinforcing steel have to be also respected. (See Table IV)

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Table IV. Permitted Welding Methods and Examples of Application

(see EN ISO 17760)

The definition of weldability for reinforcing steel is worldwide meanwhile uniform, it

consists of

- a carbon equivalent value

- limitations of pass percentage of some elements

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Table V shows the most common values

Table V. Chemical composition (% by mass)

Carbon b

(max) Sulphur (max)

Phosphorus (max)

Nitrogen c

(max) Carbon

Equivalent Value b (max)

Cast analysis

0.22 0.05 0.05 0.012 0.50

Product analysis

0.24 0.055 0.055 0.014 0.52

a Max 0.80% by mass Cu permitted b It is permitted to exceed the maximum values for carbon by 0.03% by mass, provided that the

carbon equivalent value is decreased by 0.02% by mass. c Higher nitrogen contents are permissible if sufficient quantities of nitrogen binding elements are

present.

Please, note that concerning micro alloying elements there is some freedom in the

application in combination with nitrogen.

Finally I mention bendability which is a basic demand of all manufacturers. It has to

be seen in dependency of the notch effects of the ribs which are responsible for the

bond behaviour.

Everybody knows that all constructions have to be safe. The safety of a bridge and a

building has in any case to be given even if overloading and exceptional cases like

explosions etc. take place.

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The safety philosophy currently mainly used in design are semi probabilistic methods

(Fig. 6)

Fig. 6. Design Point and Reliability Index B According to the First Order Reliability Method (FORM)

From the part of the building materials one expects significant contributions to the

safety of a construction. The mathematical interpretation of this contribution is the so-

called

partial-safety-factor γs.

Based on a log-normal distribution of the property e.g. the yield strength the formula

reads as follows:

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48

partial-safety-factor

( )[ ]

value average

deviation standard = v

tcoefficien quantile = k3,8 = 0,8 = exp

βα

βαγ kvs ⋅−⋅=

where v is the coefficient of variation, the quotient between standard deviation and

average value.

xsv =

It is clear that the larger v is, the higher the γs-value becomes and the risk increases.

Low values of v provide constructions with an additional “hidden” safety which more

than one believes will be used.

Another property of minor importance with reinforcing steel is corrosion resistance.

This problem has to be subdivided into two aspects

- corrosion before placing and pouring concrete

- corrosions of steel embedded in concrete.

In the first case slight corrosion is not dangerous, it improves bond; pitting corrosion

has to be avoided.

Much more problematic is corrosion in concrete. One tries to avoid that by

requirements for concrete density and concrete cover. Under exceptional condition

as marine surroundings stainless steel, galvanised or Epoxy-coated steel is

demanded.

In the next chapter I will show you in which advantageous way these basic and

important performance characteristics depend of micro-alloying because of the fact

that there are clear interactions with the metallurgy of the steel.

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Production methods

The very first steels used as reinforcement were profiles or rounds which are hot

rolled ferritic, perlitic steels on the basis of carbon and manganese.

Rather soon engineers detected that cold working improves the resistance, but

decreases dramatically the ductility; for simple design no problem.

Some twenty years ago microalloying in grade 400 and 500 N/mm2 appeared on the

market. Parallel to that heat treated bars were developed.

The current situation concerning product and delivery form is shown in Table VI

Table VI

Product Delivery Form

Sizes mm

Yield strengthN/mm2

Ductility Weldable Production Route

Bars 12 ÷ 40

12 ÷ 40

400 ÷ 450

400 ÷ 600

high

high

-

+

hot rolling

micro alloyed, heat treated

Wires < 14

< 14

≈ 500

≈ 500

normal

high

+

+

cold worked

hot rolled, stretched

Wire fabric < 14

< 14

≈ 500

≈ 500

normal

high

+

+

cold worked

hot rolled, stretched

Bars > 40

12 ÷ 40

≈ 500

400 ÷ 500

high

extreme

+

+

combination of micro alloying and heat treatment

mainly micro alloyed eventually heat treated

Cold working of large sizes has been important in the past but disappeared totally

due to economical reasons.

For bars currently two routes are in competition: micro alloying and heat treatment.

There are three fields of production where heat treated material cannot be used:

- small sizes produced in fast-running rolling mills for wires, speed > 60 m/sec

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- large sizes (> 40) in grades > 500

- extreme ductility requirements.

A survey how the micro alloying production route influences the performance

characteristics of reinforcing steel is given in Table VII.

Table VII

Performance Characteristics

Favourable Reasons to Use Micro-Alloyed Reinforcing Steel

Yield strength Perfect way of increasing yield by precipitations and grain size

Ductility Rm/Re

Agt

Rm/Re ≥ 1,15 independent form size

Elongation at maximum load ≥ 10%

Fatigue Reduction of notch-effect at tib-basis due to fine grain

improvement due to homogenous structure – no embedded martensite / ferrit

Weldability Improved particularly at butt-welds due to homogenous resistance across the section

Safety Coefficient of variation small due to smaller standard deviation leads to higher safety

Bendability Lower notch effects at ribs results in smaller mandrills for bending

Corrosion resistance Amount of corrosion before concreting and later is small due to rather dense layer of mill scale on the surface

Behaviour in case of fire Smaller decrease of yield depending of temperature

(significant ≥ 700 °C; cold worked 400 °C; heat treated: 550 / 600 °C

Machining and mechanical splices

Machining (lather operating) possible; splices with threads do not reduce resistance

Mechanical defects on surface

No influence on resistance, notches are softened by fine grain size

If one regards all performance characteristic and all handling practice micro alloyed

reinforcing steel has superior behaviour in comparison with reinforcing steel

produced according other production routes.

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Economical Aspect

Reinforcing steel as a mass product has to be cheap. Further on the production

should not be linked with large numbers of defective heats: a zero-defect production

is the precondition for economical and efficient production.

Rebars after leaving the rolling mill have to be fit for use without finishing process.

This means that for bars a discontinuous cold working process after hot rolling as

Rippentorstahl (twisted bars) did it – is completely out.

A usual production of reinforcing steel is done in so-called mini-mills with a

production of 0,5 to 1,5 million tons annually. These mills use electric arc furnace

process with ladle metallurgy continuous casting systems and usual mill trains. High

investment for additional equipment is not desirable. Further an immediate start of a

new product should be possible without long transition and practising time.

A complete installation of a heat treating equipment for all sizes of bars and round 1

million tons per year has costs of 3,5 million $.

One needs a sufficient large space between the last stand and the cooling bed.

Additional scissors for bars of high shear resistance are necessary and the abrasion

of conveyers etc is rather high. One needs also the installation of a temperature

control system, that means additional quality control.

Using the micro alloying process one is able to switch over from traditional hot rolled

low yield bars to high yield weldable and ductile steel.

Metallurgical Aspects

The production of EAF-steels – particularly for reinforcing steel – is expected to rise

significantly. These steels are inherently high in nitrogen for which in most standards

upper limits (120 ppm) are given. The standard permit on the other hand be exceed

the above mentioned value if sufficiently high degrees of nitrogen binding elements.

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Practise has shown that among the N-binding elements Vanadium is the most

effective to use nitrogen as a strategic alloying element to obtain best mechanical

properties. Own experiences e.g. with Titanium give an extremely high standard

deviation due to the rolling conditions. The optimised properties result from fine ferrite

grains if multiple deformation steps in the austenite zone take place. Further

precipitation strengthening can also be achieved in the steels after ferrite formation.

Thus of the different nitrogen binding elements in consideration, Vanadium offers the

best and most efficient possibility for eliminating harmful effects from nitrogen and

substantially helps to use nitrogen as a strategic alloying.

The effect of vanadium and nitrogen can be enhanced by the use of a nitrided

vanadium alloy.

If one compares a steel with the same level of vanadium obtained by FeV or the use

of nitrided vanadium alloy the yield increases by more than 120 N/mm2 in using VN.

This results logically also in a decrease of the cost of production. The advantage of

adding N-enriched nitrided vanadium alloy is caused by the fact that it promotes the

precipitation of vanadium which has a significant influence on the decrease of wasted

vanadium dissolved in the matrix.

As already explained in Chapter “Production Methods” the optimised demanded

performance characteristics for reinforcing steel have to be seen in close connection

with small grain size and precipitation of fine V (C,N). By that particularly the

properties like ductility and fatigue resistance as well as weldability are positive

influenced.

The common production process is not negatively influenced that means no specific

changes the technology are necessary. In comparison with heat treated steel no

expensive installations have to be financed.

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Economic advantages of high strength steel Savings due to high yield reinforcement

A very particular and most important aspect in this context is the improvement of the

economy of reinforced concrete without reducing the safety requirements.

The use of micro alloying is logically linked with high yield strength. This tendency

joins the trend towards higher resistance in concrete.

It is very difficult to give a simple figure for the increase in economy caused by a

certain increase in yield strength.

The reason is that all constructions are different.

Some years ago in Germany we switched over from grade 420 to grade 500. There

for I am able to give you some realistic and well experiences date of this transition. It

was in 1985 when we started a new product standard. If exists till now and contains –

as already mentioned – grade 420 and 500. Economically thinking people checked

the advantages and disadvantages and I was charged to perform an investigation.

The result from those days – which are still valuable – I will show you.

Theoretically the increase from 420 to 500 means an improvement of round about

16%. Due to the fact that in the different constructions different design rules are

dominating. For that reasons I subdivided the constructions into six groups

- office buildings / residential buildings

- industrial constructions (buildings)

- tunnels, subways

- bridges of all sizes

- high-rise buildings

- ware houses

It was possible to compare the amount of reinforcement for this six groups.

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Average values for the savings were as shown in Fig. 7.

Average round about 10 %

office buildings/residential buildings10,6

industrial constructions(buildings)10,0

tunnels, subway5,5

bridge

2,2

bridge

6,4

high-rise building

11,7

ware-house

13,8

Theoretical 16 %

It means that approximately two third of the theoretical expected saving can be

realized. The basis for this considerations are the design codes which are in use till

now in Germany.

Additional savings are linked with this reduction of the masses of reinforcement:

- better constructions: more space for concrete

- lower costs for manufacturing reinforcement: less bending, cutting, transport

- faster execution of works

The upper limit for the yield in reinforced concrete depends on different parameters.

Taking into account European design codes a yield strength between 480 and 550

fits best.

How the change from a grade 420 to a grade 520 affects the design of a column is

shown in Fig. 8.

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Fig. 8. Load-Moment Graph Showing Savings in Reinforcement

This figure shows indirectly the savings in reinforcement by means of a load-moment

interaction graph.

Another very important feature of increasing yield is given by environmental

considerations.

The consumption of energy, waste of water amount of alloys and air pollution and

output of CO2 is almost independent whether a grade 400 or 500 is produced!

This aspect gains more and more significance.

Ecological Aspects

The advantage of using high yield strength reinforcement has been pointed out.

Increasing the yield and consequently the design values a lower energy consumption

and pollution output is evident per used ton of steel.

If a production route on the basis of EAF is chosen quite a series of ecological

advantages can be obtained:

- the use of scrap saves iron or

- the melting process demands less energy in comparison with ore-reduction

- some residual elements in steel scrap can be used, as Mn.

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How progressive the EAF route concerning inputs, outputs and energy is can be

documented my an investigation executed of 7 European steel mills (Fig. 9).

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Summary

Reinforcing steel is a quantitative outstanding steel product.

Due to its use as reinforcement in the compound system reinforced concrete it has to

provide a large number of performance characteristics. Among them yield strength,

ductility, weldability, bendability, workability, fatigue and other are of main importance

and depend directly upon the metallurgy of the steel.

Therefore the production route plays an important role. Moreover the characteristics

are mutually depending and it is wise to elaborate an optimisation.

This optimisation can successfully be reached in applying the micro alloying process

on the basis of nitrided vanadium alloy.

This statement is correct for both steel making processes electric furnace and

converter.

The technology is well proven and simple in its application. No main changes in the

equipment are necessary as well as additional cost effective investment.

Due to the perfect homogeneity of precipitations and grain size a high quality level

can be reached which is of highest importance for the construction industry.

For economic reasons higher yield strength should be intended: This brings benefits

to the building industry and is technically without any risk because even with high

yield strength no decrease in the other important properties is given.

There are some fields where micro-alloyed reinforcing steel is unbeatable. On the

other hand I recommend the researchers to extend their activities to some new

sectors as

- corrosion protection in concrete

- high ductile straightened wires.

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Vanadium Microalloyed Forging Steels

George Krauss

University Emeritus Professor Colorado School of Mines Metallurgical Consultant

Evergreen, Colorado 80439, USA

1. Introduction

Many applications for structures, machine and vehicle parts require high strength and good fatigue resistance. Traditionally, steels for these applications have been produced by forging and then heat treating to produce martensitic microstructures. The hardened microstructures are then tempered, at low temperatures, if ultrahigh strengths and moderate toughness are required, or are tempered at high temperatures, if moderate high strength and high toughness are required.1,2) For moderate high strength products, vanadium microalloyed steels are now widely used.3-5) For example, Table I, from papers given in an international symposium, lists applications for which microalloyed forging steels were used in 1987.4) The moderate high strengths of vanadium microalloyed steels are produced during cooling after forging and as described below, no further heat treatment is required. The microstructures, in contrast to the martensite of quenched and tempered steels, consist of ferrite and pearlite precipitation strengthened by vanadium carbonitride precipitate dispersions.3-5)

The Use of microalloying in steels is based on the addition of small amounts of vanadium, niobium and/or titanium,5) typically on the order of 0.1 to 0.2 mass pct or less. Other elements, such as aluminium for grain size refinement and boron for hardenability, and elements residual from steelmaking, may of course also present in steels in small amounts, but such elements and their effects are generally considered to be outside of microalloying technology.

Microalloying was first applied, in the 1960s and 1970s, to flat-rolled, low-carbon steel for higher strengths, in which low-temperature controlled rolling and niobium additions combined to prevent austenite recrystallisation, and thereby promoted very fine ferrite grain sizes with excellent combinations of strength, between 300 and 500 MPa, and toughness.6) The need for higher strengths, combined with forging and higher carbon contents, however, produced a much different set of conditions for microalloying in forging steels. As a result microalloying in forging steels developed widely only in the 1980s. Vanadium became the primary microalloying element, and microstructures consisting largely of pearlite strengthened by vanadium precipitates replaced the largely ferritic structures of low-carbon sheet and plate products in which niobium was the primary microalloying element.

This paper describes in detail the processing, microstructure and properties of medium-carbon, vanadium-containing forging steels.

2. Processing Considerations

Forging steel technology is based on thermomechanical forming of bar steels. Today the bar steels are almost universally produced by electric arc furnace melting of scrap steel, ladle

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refining, continuous casting of blooms or billets, and hot rolling to final bar diameters. The shaping of the bars to complex shapes, as required by some of the applications illustrated in Table 1, is accomplished by forging at necessarily high temperatures. In view of the high temperatures required for forming of complex shapes, controlled low-temperature hot deformation as applied to flat rolled products is not feasible. Fortunately, the high temperatures of forging make possible the solution of vanadium carbonitrides present as a result of bar steel processing, and as a result, the vanadium is available for the formation of fine strengthening precipitates on cooling from the forging temperature.

The strengthening produced by the vanadium enhances the strength of medium carbon steels to where it competes with highly tempered heat treated steels of the same hardness. A major benefit, therefore, of microalloyed forging steels is the fact that hardening heat treatment handling and equipment costs can be eliminated. Fig. 1 shows the processing steps associated with quench and tempered forging steels, and Fig. 2 shows the processing steps for microalloyed forging steels and for cold finished bar steels, another technique of increasing the strength of bar steels. The direct cooled microalloyed steels have significantly fewer processing steps than do the heat treated forgings.

3. Alloying and Microstructural Considerations

The slow air cooling of forgings produces base microstructures that consist of ferrite and pearlite. The higher the carbon content, the greater the volume fraction of pearlite that forms and the higher the strength of the as-cooled microstructure. Gladman et al.7) have developed the following equation for the yield strengths of ferrite-pearlite microstructures in carbon steels:

σ ys(MPa) = 15.4 ( f a3

1

[2.3 + 3.8 (%Mn) + 1.13d 21

− ] + (1- f a31

)[11.6 +

0.25SP21

− ]) + 4.1(%Si) + 27.6(%N) (1)

Where f a is the volume fraction of ferrite, Sais the pearlite interlamellar spacing in mm,

and d is the linear intercept grain diameter in mm. The first term of the equation relates to the factors that determine the strength of the ferrite component of the microstructure and the second term relates to the strengthening due to the pearlite component of the microstructure. As the ferrite content of the microstructure decreases, the yield strength increases with increasing pearlite content, and as shown, the interlamellar spacing of the ferrite and cementite in pearlite determine the strength of the pearlite.

Fig. 3 shows a plot of the measured strengths versus the calculated strengths according to equation [1] of a number of steels with ferrite-pearlite microstructures and various microalloying additions. All of the steels, except those with vanadium additions fit the expected yield strengths for various combinations of ferrite and pearlite and nominal concentrations of Mn and Si. The vanadium-containing steels showed consistently higher strengths than estimated only from their ferrite and pearlite contents. This enhanced strengthening of ferrite and pearlite microstructures provides the basis of the enhanced strength of vanadium Microalloyed forging steels.

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The increased strength of the vanadium-containing steel is related to the formation of fine vanadium-carbonitride precipitates by a mechanism that is termed Interphase precipitation.8) The carbonitride particles form during cooling on the interfaces between austenite and ferrite, and as the ferrite grows, rows of precipitate particles are imbedded in the ferrite. Fig. 4 shows an example of such precipitation in 0.2% C-0.14% V steel. The precipitate particles are quite fine, on the order of 10 nm in size, and are not resolvable in the light microscope. The vanadium carbonitride particles form in both equiaxed ferrite grains and in the pearlitic ferrite.5)

The effectiveness of a microalloyed precipitation strengthening mechanism depends on the complete solution of any microalloy precipitates at a forging temperature. The formation of a microalloy precipitate is assumed to take place by a reaction of the form:

M + N = MN (2)

Where M represents a transition element such as V, Nb or Ti in substitutional solid solution in austenite, N represents a non-metal such as C or N in interstitial solid solution in austenite, and MN represents the carbide or nitride precipitate that forms in the austenite when the two types of elements combine. The temperature dependence of such a reaction is given by an equation of the form:

ln [M][N] = A − TB (3)

Where [M][N] is the solubility product, A and B are constants for a given reaction, and T is the temperature in K. Equation [3] shows that the solubility product increases with increasing temperature and the various concentrations of M and N that can make up the solubility product at a given temperature. For example, plots of solubility for VN, NbN, and TiN as a function of microalloying element and nitrogen content for selected temperatures are shown in Figures 5, 6, and 7, respectively.5) For a given curve, concentrations to the left and down represent complete solubility in the austenite and concentrations to the right and up represent concentrations where solubility is exceeded and a microalloy nitride precipitate forms in the austenite.

Fig. 5 shows that VN dissolves more rapidly with temperature than does NbN, as shown in Fig. 6. Fig. 7 shows that the solubility of TiN in austenite is very low, even at 1300ºC, a temperature, which should not be exceeded during forging. These solubility relationships have a major impact on the way microalloying elements are used in forging steels. The vanadium nitrides and carbonitrides dissolve readily at forging temperatures and therefore vanadium is available for precipitation on cooling. The niobium nitrides and carbonitrides have lower solubility and require higher temperatures for complete solution. This if forging temperatures are difficult to control or vary, some of the niobium carbonitrides may remain undissolved, and precipitation strengthening response may vary, a characteristic which tends to limit the use of Nb in forging steels.

The very low solubility of Ti in austenite, even at high temperatures, prevents its use as a precipitation strengthening addition to forging steels. However, because TiN particles are stable at very high temperatures, they are effective in pinning grain boundaries. Therefore Ti additions, on the order of 0.01%, which cause the precipitation of fine TiN particles in austenite, are very effective in maintaining fine austenitic grain sizes during forging. It is important not to add too much Ti in order to prevent the formation of coarse TiN particles in liquid steel. Coarse particles are ineffective in restraining austenite grain growth and are

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detrimental to toughness and fracture. Fig. 8 ranks austenite grain coarsening behaviour of steels containing V, Al, Nb and Ti and shows that grain coarsening in Ti-containing steels does not occur at temperatures around 1200ºC or higher.9)

4. Mechanical Properties of Microalloyed Forging Steels

In order to achieve moderately high strengths in microalloyed steels with ferrite-pearlite microstructures, medium-carbon steels that form large volume fractions of pearlite are typically used for forgings. Thus higher carbon steels provide higher strengths, but because of the sensitivity of pearlite to cleavage fracture, the higher carbon steels have lower toughness, as discussed below. Figures 9 and 10 show yield and ultimate tensile strengths as a function of carbon content for plain carbon steels and steels either microalloyed with vanadium or with vanadium plus niobium.10) Vanadium levels were 0.15 mass pct and niobium, when present, was at 0.04 mass pct. The strengths of all steels increase with increasing carbon content, and the vanadium steels have significantly higher strengths that do the plain carbon steels. Additions of Nb increased strengths over those achieved in steels with only vanadium additions, but as noted above, the reduced solubility of Nb compared to V makes the solution and effectiveness of Nb difficult to control. Yield and ultimate tensile strengths correlate well with hardness, as shown in Fig. 11

Compared to highly tempered quench and tempered steels, microalloyed steels at the same hardness have much lower toughness. Fig. 12 compares CVN energy absorbed as a function of temperature for a quench and tempered 4140 steel and two vanadium microalloyed steels.11) The microalloyed steels with ferrite-pearlite microstructures have lower transition temperatures and absorb lower energies at low temperatures because of sensitivity to cleavage fracture. However, the two types of steel have comparable fatigue resistance, as shown in Fig. 13. Thus, if impact conditions are not present, as for example in the operation of crankshafts, microalloyed forging steels are effective, economical replacements for heat treated quench and tempered steels with much higher alloy contents.

Fig. 14 shows that the 27 Joule impact transition temperature of plain carbon and microalloyed forging steels increases with increasing carbon content, and that correlating with the higher strength of the microalloyed steels, the impact transition temperature increases in the microalloyed steels.10) The vanadium-containing steels have lower transition temperatures than do the steels containing niobium, but nevertheless, impact transition temperatures, which mark the transition between ductile and brittle fracture, still tend to be above room temperature. The low toughness of pearlitic microalloyed steels is consistent with the susceptibility of pearlitic microstructures to cleavage fracture1) and the fact that any strengthening mechanism, other than that based on fine ferritic grain sizes, raises impact transition temperatures.5)

Another approach to increase toughness is to add titanium as a microalloying element. As discussed earlier in this paper, Ti is a very strong nitride forming element, and titanium nitride particles are stable at the high temperatures of forging. The titanium nitride particle dispersions prevent austenite grain growth, and as a result more grain boundaries are available for pearlite nucleation and pearlite colony size is reduced. The finer pearlite colony size increases impact toughness.5)

Yet another approach to increasing toughness of microalloyed forging steels is to increase sulfur content.12) Ochi et al.13) have shown that in vanadium-microalloyed steel ferrite nucleates and grows on manganese sulfide particles. Vanadium nitride and carbides precipitate on the sulfides and stimulate the nucleation of ferrite. As a result, ferrite forms not

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only on widely spaced austenite grain boundaries in steels forged at high temperatures, but also within austenite grains. The intragranularly nucleated ferrite breaks up the structure of coarse colonies of pearlite, effectively decreasing pearlite colony size and improving toughness. Fig. 15 shows examples of intragranular ferrite formation on MnS particles in a steel containing 0.39 pct C, 0.059 pct V, and 0.094 pct S.12)

5. Summary

Microalloying additions of vanadium to medium-carbon steels effectively increase the strength of forged steels without subsequent heat treatment after cooling of the forgings. Enhanced strengthening is accomplished by the precipitation of fine vanadium carbonitride particles in the ferrite of direct cooled ferrite-pearlite microstructures. The highest strengths, approaching ultimate tensile strengths of 1000 MPa, are produced in base microstructures with large volume fractions of pearlite. As a result, the resistance to cleavage fracture during impact loading is low. Toughness is increased by lowering carbon content and reducing the amount of pearlite in the microstructure, titanium additions to refine austenite grain size and pearlite colony size, or by increasing sulfur content, which in vanadium-containing steels stimulates the formation on intragranular ferrite and effectively reduces pearlite colony size.

REFERENCES

1. G. Krauss: Steels Heat Treatment and Processing Principles, ASM International, Materials Park, Ohio, 1990

2. G. Krauss: “Deformation and Fracture in Martensitic Carbon Steels Tempered at Low Temperatures”, Metallurgical and Materials transactions A, 32A, 2001, pp. 861-877.

3. C. J. Van Tyne, G. Krauss and D. K. Matlock: Fundamentals and Applications of Microalloying Forging Steels, Editors, The Minerals, Metals & Materials Society, Warrendale, Pennsylvania, 1996.

4. G. Krauss and S. K. Banerji: Fundamentals of Microalloying Forging Steels, The Metallurgical Society Inc., Warrendale, Pennsylvania, 1987.

5. T. Gladman: The Physical Metallurgy of Microalloyed Steels, The Institute of Materials, London, England, 1997.

6. I. Tamura, H. Sekine, T. Tanaka and C. Ouchi: Thermomechanical Processing of High Strength Low-alloy Steels, Butterworths, London, England, 1988.

7. T. Gladman, I. D. McIvor and F. B. Pickering: “Some Aspects of the Structure Property Relationships in High-carbon Ferrite-Pearlite Steels”, Journal of the Iron & Steel Institute, 210, 1972, pp. 916-930.

8. R. W. K. Honeycomber: “Fundamental Aspects of Precipitation in Microalloyed Steels”, in HSLA Steels: Metallurgy and Applications, J. M. Gray et al., Editors, ASM International, 1986, pp. 243-250.

9. G. R. Speich, J. Cuddy, C. R. Gordon and A. J. DeArdo: “Formation of Ferrite from Controlled-rolled Austenite”, in Phase Transformations in Ferrous Alloys, A. R. Marder and J. I. Goldstein, Editors, The Metallurgical Society, Warrendale, Pennsylvania, 1984, pp. 341-389.

10. Y. Sawada, R. P. Foley, S. W. Thompson and G. Krauss: “Microstructure-Property Relationships in Plain-carbon, and V and V + Nb Microalloyed Medium-carbon Steels”, in 35th Mechanical Working and Steel Processing Conference Proceedings, Vol. XXXI, ISS-AIME, Warrendale, Pennsylvania, 1994, pp. 263-286.

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11. P. B. Babu, D. R. Gromer, D. J. Lingenfelser and G. P. Shandley: “Design for Fracture Resistance in Microalloyed Steel Components”, in Reference 2, pp. 389-424.

12. B. G. Kirkby, P. LaGreca, C. J. Van Tyne, D. K. Matlock and G. Krauss: “Effect of Sulfur on Microstructure and Properties of Medium-carbon Microalloyed Bar Steels”, SAE Technical Paper 920532, 1992.

13. T. Ochi, T. Takahashi and H. Takada: “Improvement of the Toughness of Hot Forged Products Through Intra-granular Ferrite Formation”, in 30th Mechanical Working and Steel Processing Conference Proceedings, Vol. XXVI, ISS-AIME, Warrendale, Pennsylvania, 1988, pp. 65-72.

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Table 1 Application for microalloyed forging steels

Crankshafts

Connecting Rods

U-Bolts for Leaf Springs

Steering Knuckle Supports

Antisway Bars

Induction Hardened Gears

Drive Couplings

Fasteners

Pistonshafts

Axle Shafts

Suspension Arms

Transmission Shafts

Wheel Hubs

Steering Arms

Axle Beams

Pipe Fittings

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Fig. 1 Temperature-time processing schedules for producing quench and tempered forgings.

Fig. 2 Temperature-time schedules for producing direct cooled microalloyed forgings and cold finished bars.

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Fig. 3 Observed and calculated yield strengths for steels with ferrite-pearlite microstructures and various microalloyed element additions.7)

Fig. 4 Fine vanadium carbonitride precipitates in ferrite of an 0.2 wt.% C- 0.14 wt.% V steel air cooled from 1200°C. Dark field transmission electron micrograph, courtesy of S. W. Thompson, Colorado School of Mines.

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Fig. 5 Solubility limits of vanadium and nitrogen at various austenitizing temperatures. From Galdman.5)

Fig. 6 Solubility limits of niobium and nitrogen at various austenitizing temperatures. From Galdman.5)

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Fig. 7 Solubility limits for titanium nitrogen in liquid steel and in austenite. From Galdman.5)

Fig. 8 The effect of various microalloying elements on austenite coarsening.9)

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Fig. 9 Yield strength as a function of carbon content and microalloying. From Sawada et al.10)

Fig. 10 Tensile Strength as a function of carbon content and microalloying. From Sawada et al.10)

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Fig. 11 Yield and tensile strengths of steels with ferrite-pearlite microstructures as a function of hardness. From Sawada et al.10)

Fig. 12 CVN energy absorbed as a function of temperature for quench and tempered 4140 steel and two vanadium microalloyed steels. From Babu et al.11)

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Fig. 13 Comparison of fatigue behaviour of quench and tempered and microalloyed steels. From Babu et al.11)

Fig. 14 27 J transition temperature as a function of carbon content and microalloying. From Sawada et al.10)

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Fig. 15 Intra-granular ferrite formed in MnS particles. The ferrite is surrounded by pearlite, which appears black in the light micrograph. From Kirby et al.12)

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PRECIPITATION AND GRAIN REFINEMENT IN VANADIUM – CONTAINING STEELS

Stanislaw Zajac Swedish Institute for Metals Research

Drottning Kristinas väg 48, S-114 28 Stockholm, Sweden

Tel: +46 8 440 4884 E-mail: stanislaw.zajac @simr.se

Abstract

The present work has concentrated on the roles of vanadium, nitrogen and carbon in controlling the precipitation of V(C,N) in austenite and ferrite and their effects on; (i) grain refinement by promoting the formation of intragranular ferrite, and (ii) precipitation strengthening by interphase and random precipitation.

The degree of precipitation strengthening of ferrite at a given vanadium content depends on the available quantities of carbon and nitrogen. It is concluded that nitrogen is a very reliable alloying element, increasing the yield strength of V-microalloyed steels by some 6 MPa for every 0.001% N, essentially independent of processing conditions. The role of carbon in precipitation strengthening is complicated. The present results have shown that the precipitation strengthening of V-microalloyed steels increases significantly with the total C-content, ~5.5MPa/0.01% C. The explanation is that the metastable equilibrium between ferrite and undercooled austenite greatly increases the solubility of carbon in ferrite, in the times available during transformation, thereby contributing to profuse nucleation of V(C,N) particles. This effect of carbon is particularly significant for medium carbon steels typically used for hot rolled bars and sections.

The experimental results strongly indicate that vanadium can by effectively used not only for precipitation strengthening but also for ferrite grain refinement. It was shown that vanadium contributes to the formation of two types of intragranulary nucleated ferrite; polygonal (idiomorph) ferrite and acicular (sideplate) ferrite. Intragranular polygonal ferrite nucleates on VN particles that grow in austenite during isothermal holding or slow cooling throughout the austenite range. Acicular ferrite microstructure forms in V-microalloyed steels during isothermal transformation at lower temperatures (~450°C). The acicular ferrite microstructure was obtained in V-microalloyed steels containing high, medium or very low nitrogen levels. This suggests that vanadium on its own can promote the formation of the acicular ferrite microstructure.

1. INTRODUCTION

Vanadium is best known as an eminent element for strong and easy controllable precipitation strengthening. The principal reason for this is the relatively large solubility product of its carbo-nitrides resulting in a lower solution temperature and a larger capacity to dissolve them at elevated temperatures. A special feature of V as compared to Nb is that its nitride is much less soluble than its carbide and this gives N a very important role in V-steels, especially in their precipitation

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strengthening. In order to maximise the precipitation strengthening it is necessary to understand the roles of nitrogen and carbon in formation of high volume fractions of finely dispersed carbonitrides. Previous data strongly indicate that the role of nitrogen is clear1. It has been demonstrated that the yield strength increases almost linearly with increasing nitrogen content for given vanadium and carbon levels making N an eminent choice for strong and easily controllable precipitation strengthening2. Carbon content has usually been considered not relevant to precipitation strengthening when the precipitation occurs in ferrite. This was deduced from the fact of very restricted solubility of carbon in ferrite (which is normally supposed to be independent of the total carbon content in the steel). Most published literature does not suggest that differences in carbon contents in the range for structural steels (0.04-0.3%) should affect significantly the response of vanadium in these steels3. However, this viewpoint must be revised in the light of the later work2 which have shown that the precipitation strengthening increases significantly with total C-content of the steel. Increasing C-content delays the pearlite formation and thereby maintains for a longer time the higher content of solute C in ferrite corresponding to the austenite/ferrite equilibrium as compared to that of ferrite/cementite, allowing more nucleation of V(C,N) particles and accordingly a more dense precipitation.

Recent studies at SIMR4 and literature data5-7 strongly suggest that vanadium can also by effectively used for ferrite grain refinement. There is also evidence that vanadium promotes the formation of acicular ferrite microstructure steels8. It was suggested that the VN particles which precipitated inside austenite grains during/or after hot rolling show strong potential for nucleation of intragranular ferrite. Although, vanadium does not readily precipitate in austenite, the precipitation process can be enhanced with increasing nitrogen in the steel or by plastic deformation (strain-induced precipitation). In the case of precipitation in ferrite, the Baker-Nutting (B-N) orientation relationship is observed between the V(C,N) particles and matrix9. This fact may be very important for the nucleation of ferrite on cooling as the interfacial energy between ferrite and vanadium nitride is very low for the B-N relationship. In fact, intragranular ferrite idiomorphs were observed to nucleate at vanadium nitrides with the B-N orientation7. These observations suggest that intragranular ferrite can nucleate on VN and maintain coherent, low energy interfaces with respect to vanadium nitride. Thus, it was concluded that the main factor governing the formation of intragranular ferrite is the presence of VN precipitates in austenite which can develop coherent, low energy, interphase boundaries with ferrite10. The precise mechanism by which vanadium additions may enhance the nucleation rate of ferrite is not known7,10,11, but this new role of vanadium can be extremely important for thermo-mechanical processing as grain refinement leads to a significant increase in strength which is accompanied by a marked improvement in toughness.

The aim of the present paper is to review the present knowledge on the roles of V, N and C in precipitation strengthening as well as on grain refinement by promoting the formation of intragranular ferrite.

2. THERMODYNAMIC CONSIDERATIONS

2.1 Solubility and composition of V(C,N) in microalloyed steels

To achieve the desired metallurgical states, a detailed knowledge of the solubilities of the microalloy carbides and nitrides is required, together with a knowledge of their precipitation behaviour. An understanding of the role of V as well as C and N can be gained from the solubility data summarised in Figs 1-3 from a recent evaluation of the thermochemical parameters for multi-component systems of HSLA steels12.

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Fig. 1 shows Thermo-Calc calculations of the equilibrium precipitation of V(C,N) in steels microalloyed with vanadium. Also shown in this figure is the mole fraction of nitrogen in carbo-nitrides at various temperatures and for various nitrogen contents from hyperstoichiometric levels to zero. These results show that vanadium starts to precipitate in austenite as almost pure nitride and the precipitation start temperature depends strongly on the nitrogen level. For 0.10%V and 0.03%N the precipitation starts at ~1110°C and the precipitation start temperature decreases to below 950°C at 0.003%N. When the nitrogen is about to be exhausted there is a gradual transition to form mixed carbo-nitrides, Fig. 1(b). Enhanced precipitation during the transformation is a result of the solubility drop of the vanadium carbo-nitrides associated with the transformation from austenite to ferrite at a given temperature. A significant feature of the solubility of VC in austenite is that this is considerably higher than the solubility of VN, suggesting that vanadium carbide will not precipitate in austenite.

Fig. 1 Precipitation of nitrides, nitrogen rich carbonitrides and carbides in 0.10% V steels at various nitrogen contents.

The precipitation of carbo-nitrides in multiple microalloyed steels with V, Nb, Ti and Al is shown in Fig. 2(a), where the soluble components are expressed as a function of temperature. The mole fraction of Ti, Nb and V in M(C,N) is also included in this figure. As expected, the major part of the precipitate at the highest temperatures is TiN whilst further precipitation at lower temperatures is predominantly Nb(C,N). Thus, the primary precipitates to be formed in austenite are composite (Ti,Nb)-nitrides. Thermodynamic calculations imply that at the high temperature of 1200°C, primary (Ti,Nb,V)N particles contain ~20% of Nb and ~5% of V for multiple microalloyed steels. It is clear from this figure that the volume fraction of microalloy nitrides at high temperatures will be considerably larger in the Ti-Nb-V steel than in the single microalloyed steels. It may also be seen that AlN (with close packed hexagonal structure) has little or no solubility for other microalloying elements and starts to precipitate at approximately 1200°C in the presence of 0.035% Al, Fig. 2(a).

Another important aspect of multiple microalloying is that the niobium or vanadium tied up as (Ti,Nb,V)N particles is not available for subsequent thermo-mechanical treatment. Fig. 2(b) shows that a significant fraction of added niobium may remain undissolved even at high reheating temperatures and cannot contribute to retardation of recrystallisation and/or precipitation strengthening. Vanadium, on the other hand, exhibits higher solubility and is therefore normally fully available for precipitation strengthening in ferrite.

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Fig. 2 Precipitation of nitrides and nitrogen-rich carbonitrides in multiple microalloyed steel (a) and the mole fraction of Ti, Nb and V in M(C,N) (b).

2.2 Thermodynamic Driving Force for Precipitation

The precipitation process proceeds at a perceptible rate only if there is a driving force, that is, a free energy difference between the product and parent phases. This driving force enters the steady state nucleation rate in a central way and must be known with some accuracy if nucleation rates are to be calculated, or even estimated. The chemical driving force for nucleation of VN and VC in V-microalloyed steels, determined from the HSLA database are illustrated in Fig. 3. It can be seen that, as the temperature is decreased the driving force increases monotonically and changes slope after the austenite to ferrite transformation. The dominating effect of N on the driving force is clearly seen in Fig. 3(a).

(a) (b)

Fig. 3 Chemical driving force, ∆Gm/RT, for precipitation of VC and VN in 0.12% V steel.

For hard precipitates such as V(C,N) bypassing of dislocations is expected to occur by bowing between the particles (Orowan-mechanism) under all practical conditions. This means that the precipitation strengthening effect is determined by the interparticle spacing. In precipitation reactions, the decisive factor which minimises the interparticle spacing is the rate of nucleation since

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it determines the particle density. Hence, the strong effect of N in increasing this driving force, Fig. 3(a), is the explanation of the well-known increase in strength by N in V-microalloyed steels. It should be noted, however, that the driving force can also be raised by V-addition but the effect per concentration unit is much less, Fig. 3(b). Moreover, V-additions increase alloying costs significantly whereas the usage of N at the present levels is free.

3. PRECIPITATION OF V(C,N) IN AUSTENITE AND FERRITE

3.1 Kinetics of VN precipitation in austenite

The precipitation-temperature-time diagram for VN in undeformed (recrystallised) austenite is shown in Fig. 4 for the 0.12%V steel. The experimental results yield the well-known “C-curve” associated with the kinetics of VN precipitation. The curves indicate precipitated vanadium at the level of 0.001%, 0.005%, 0.01% and 0.02%. The location of the C-curve on time-temperature diagram is very important in regards to processing since precipitates will only form when the processing time-temperature profile passes through the C-curve and the effectiveness of VN in controlling transformation reactions is directly related to the amount and distribution of the precipitate which is subsequently related to the location of the C-curve.

775

800

825

850

875

900

925

950

975

1000

1025

1.E+01 1.E+02 1.E+03 1.E+04 1.E+05 1.E+06

Time, s

Tem

per

atu

re,

°C

0.010.001

0.010.001

0.001

0.005

0.005

0.1%C-0.12%V-0.0082%N

Precipitated V (wt%)

0.02

Fig. 4 Precipiation-time-temperature diagram for VN in undeformed austenite.

It is clearly seen from Fig. 4 that the precipitation process of VN in undeformed austenite is very sluggish. After holding for 1 hour at 850°C less than 10% of the equilibrium amount of V, which can precipitate at this temperature, was precipitated in VN. The maximum rate of precipitation of VN in austenite was evaluated to be in the range of 850-875°C.

3.2 Precipitation of V(C,N) in Ferrite

As regards strengthening, the effective V-carbonitrides are those formed in ferrite during the latest passage through the austenite-ferrite transformation. At equilibrium the circumstances are such that a certain, small portion of the vanadium in microalloyed steel should precipitate in austenite, especially if the contents of vanadium and nitrogen are high. However, the kinetics of V(C,N) formation in austenite are sluggish and for processing at finishing temperatures higher than 1000°C and for normal steel compositions virtually all vanadium will be available for precipitation in ferrite.

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Precipitation of V-carbonitrides can occur randomly in ferrite in the wake of the migrating austenite-ferrite (γ-α) boundary – general precipitation – or by interphase precipitation characterised by the development of sheets of particles parallel to the γ/α-interface formed repeatedly with regular spacing. Many investigations have shown that for compositions typical of structural steels the general precipitation takes place at lower temperatures, typically below 700oC, and the interphase precipitation at higher temperatures.

It is now a well established fact, demonstrated in several electron microscopy studies1,2, that V(C,N)-particles are also formed in the pearlitic ferrite. Because of the lower transformation temperatures of pearlite this type of precipitation is usually finer. Both interphase and general precipitation have been observed.

3.2.1 Interphase precipitation Fig. 5 shows the typical morphology of interphase precipitation of V(C,N) in 0.04-0.10%C, 0.13%V steels. Already from its appearance one can conclude that such a microstructure is formed in sheets parallel to the γ/α-interface by repeated nucleation of particles as the transformation front moves through the austenite.

Fig. 5 Electron micrographs of the selected steels isothermaly transformed at 750°C for 500 s, (a) 0,0051%N (b) 0,0082%N (c) 0,0257%N (d) 0,0095%N, 0,04%C. The effect of transformation temperature (e) and N-content (f) on intersheet spacing of V(C,N) interphase precipitation in 0.04-0.10%C, 0.12%V, 0.005-0.025%N steels.

At high transformation temperatures, ~800°C, the interphase precipitation consists of irregularly spaced, and often curved sheets of V(C,N)-particles1. With decreasing temperatures the incidence of curved rows of precipitates diminishes and the dominant mode is regularly spaced, planar sheets of particles. From about 700oC the interphase precipitation is commonly found to be incomplete, and random precipitation from supersaturated ferrite after the γ/α-transformation takes over progressively with decreasing temperature. A characteristic feature of interphase precipitation is that it becomes more refined at lower temperatures, Fig. 5(e). Fig. 5(f) shows that the intersheet spacing is affected considerably by the nitrogen content of the steel. As shown there, it is diminished to almost one third at 750oC on increasing the nitrogen content from 0.005 to 0.026%.

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Raising the carbon content above 0.10% C, will similarly decrease the A3 temperature of the steel and suppress interphase precipitation. According to the present investigation, interphase precipitation was visible within proeutectoid ferrite or pearlitic ferrite only at transformation temperatures very close (within 40-50°C) to the eutectoid point. At lower transformation temperatures, formation of solute supersaturated proeutectoid ferrite and pearlitic ferrite occurs instead with random distributions of V(C,N).

3.2.2 The mechanism of interphase precipitation

The mechanism of interphase precipitation has been the subject of considerable discussions. The models that have been proposed to explain the phenomenon fall broadly into two categories; ledge mechanisms and models based on solute diffusion control. Honeycombe and coworkers were among the first to study interphase precipitation more profoundly13,14. They suggested that interphase particles form heterogeneously on γ/α-boundaries thereby pinning their migration normal to the boundary. Local breakaway leads to formation of mobile ledges. The ledges move sideways while the remaining part of the released boundary is stationary and enables repeated particle nucleation to occur, forming a new sheet. Hence, in this mechanism the intersheet spacing will be determined by the ledge height. From this follows one of the main drawbacks with the ledge mechanism, viz. to produce a credible explanation of the observed variation of the intersheet spacing with temperature, and steel composition, especially N, V and C. It is hard to see how these parameters should generate a corresponding variation of the ledge height.

VV

V

VV

V

A

B

C

N

ABC

γγ

Vc0

Vc

αVc

γα /

Vc

/VCN

Vcα

λλ

xo

Fig. 6 Left figure shows schematically how the γ/α-interface bows out, expands sideways and reaches eventually material with sufficient V for renewed precipitate nucleation to occur. The transfer of V boundary diffusion to the lower precipitate row is indicated. The right figure shows the V-content profile in a cross section, see eq(1).

Among the models based on diffusion control the solute-depletion model due to Roberts15 is the most prominent and promising one and has been recently further developed by Lagneborg and Zajac16. The quantitative description of the solute-depletion model considers a ferrite grain growing into austenite where the growth is controlled by carbon diffusion in austenite while maintaining local equilibrium at the interface, Fig. 6. At a given point of this growth the interplay is analysed between the nucleation of V(C,N)-particles in the γ/α-interface, the accompanying growth of V-depleted zones around the precipitates, and the continued migration of the γ/α-boundary away from the precipitate sheet. The following relationship for the intersheet spacing, λ, has been suggested;

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*/0

/*/

1

2γα

αγα

λVV

VCN

VV

b

V

cc

cc

K

SD

−−

= (1)

where S is the distance from the point of α-nucleation to the location of the γ/α-interface at time t, b

VD the boundary diffusion of V, VCNVc /α the V-content of α in the α/VCN interface, */ γα

Vc the

critical V-content in the γ/α interface for nucleation of interphase precipitation, K1 is the proportionality factor for the ferrite growth.

It was shown that volume diffusion of V cannot explain the observed intersheet spacings and that a faster diffusion process is required. Hence, a new mechanism was put forward where the transport of V occurs by boundary diffusion in the moving α/γ-interfaces. The model exhibits good agreement with observed values of intersheet spacing and their dependences on C-, V-, N-contents and temperature, as illustrated in Fig. 7. The computations shown in Fig. 7 also predict that the intersheet spacing around 700oC will fall below the approximate size of observed precipitates. Physically, this implies that the growth rate of the γ/α-interphase exceeds that of the V-depleted zone. Hence, the γ/α-boundary escapes, leaving the ferrite in its wake supersaturated with respect to V(C,N).

0

50

100

150

200

250

300

680 700 720 740 760 780 800

Temperature, °C

Inte

rsh

eet

spac

ing

, nm

0,10C-0,06V-0,010N, calculated

0,10C-0,06V-0,010N, experimental0,10C-0,12V-0,010N, calculated

0,10C-0,12V-0,010N, experimental0,04C-0,12V-0,010N, calculated0,04C-0,12V-0,010N, experimental

Fig. 7 Experimentally measured and computed intersheet spacings in interphase precipitation of a 0.06-0,13%V steel as a function of the transformation temperature and carbon and nitrogen contents.

The model predicts also that the spacing is directly proportional to the ferrite growth, or to the degree of transformation. In fact, the model predicts that in the early stages of transformation the ferrite growth will be too rapid for V(C,N)-nucleation to take place and it is only later, when the growth rate has declined sufficiently, that the condition of interphase nucleation will be fulfilled. In the first part of the γ/α-transformation the ferrite will therefore be supersaturated with respect to V(C,N) directly behind the moving γ/α-interface and hence general precipitation will occur. Smith and Dunne17 made exactly this statement based on their microscopical observations and without the assistance of a model. Beside this, however, no attention seems to have been paid in the literature to this essential feature of interphase precipitation. The general statement often made, that a large variation of precipitation modes are observed, is probably largely a reflection of this specific characteristic of the phenomenon.

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3.2.3 General precipitation

For compositions typical of V-microalloyed steels, 0.10%C-0.10%V, general precipitation of V(C,N) occurs in a temperature range from about 700oC and below. As already discussed in the previous section this transition from interphase to general precipitation can be predicted by the solute-depletion model. Randomly distributed V(C,N) particles are most common in V-steels transformed at 550-650°C at all carbon contents. The fact that these particles had formed in the ferrite after transformation and not on the inter-phase boundary during transformation is evident both from their random dispersion and from the existence locally of several different variants of the B-N orientation relationship9.

Experimentally it is well established that VN has a considerably lower solubility and much larger chemical driving force for formation than VC, both in ferrite and austenite. This larger chemical driving force makes N-rich V(C,N) the preferred precipitation as long as there is sufficient nitrogen in the matrix, as has been demonstrated both by thermodynamic analysis, Figs 1 and 3 and by experiment18. It is only when the nitrogen content falls below about 0.005% that the initial V(C,N) starts to increase its C-content, Fig. 8.

Fig. 8 Thermodynamic calculations of the variation in the composition of VCxNy with content of N remaining in solution during precipitation in ferrite18.

Fig. 9 Growth of V(C,N)-precipitates after transformation at 650°C (a) as a function of the N-content of the steel, and (b) as a function of holding time.

A technically very important finding is that the precipitate size diminishes considerably with increasing nitrogen content in the steel, as is shown in Fig. 9. This is accompanied by a concurrent increase of precipitate density, although that is more difficult to measure quantitatively. These effects must be accounted for by an increased nucleation frequency as a result of the larger chemical driving force for precipitation of nitrogen-rich V(C,N). For the temperature and holding times in the experiments of Fig. 9 supersaturation with respect to V(C,N) still remains. Therefore the observed difference in precipitate growth between high and low-N steels cannot be explained by differences in coalescence. Instead, it must be interpreted as resulting from the denser particle nucleation in the high-N material, thereby producing an earlier soft impingement of V-denuded zones and so slowing down the precipitate growth. In numerous investigations of V-microalloyed steels it has been shown

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that the observed precipitation strengthening originates seemingly only from N-rich V(C,N), despite the fact that there is abundant V to combine with the carbon dissolved in ferrite. The normal explanation of this behaviour has been that when the first formed N-rich precipitates have consumed practically all nitrogen the chemical driving force for formation of C-rich V(C,N) is too small for profuse precipitation to occur and hence no added strength is observed. However, this issue is complex, and recently it has been shown that under certain conditions the C-content in the steel can add significantly to precipitation strengthening, as shown below.

The recent results of detailed measurements of particle sizes on steels having different carbon and nitrogen contents are summarised in Fig. 10, after isothermal transformation at 650°C. As indicated in this figure both nitrogen and carbon cause a decrease in the particle size and at the same time an increase in volume fraction, so giving the lowest planar spacing at a given transformation temperature. For steel having 0.005% nitrogen and 0.1% carbon the particles were relatively sparse with an average diameter of ~ 11 nm. They became smaller and more dense with increasing nitrogen content, reaching ~ 6 nm at 0.022% nitrogen. An increase in carbon content from 0.10% to 0.22% causes significant refinement (about a factor of two) of the vanadium-rich particles at all nitrogen levels. A higher density of smaller V(C,N) suggests that carbon play an important role in precipitation of V(C,N). This agrees with the suggestion that metastable carbon content increases the supersaturation levels achieved in ferrite and the intensity of V(C,N) precipitation is higher.

0

2

4

6

8

10

12

0 0.005 0.01 0.015 0.02 0.025 0.03

wt% N

PA

RT

ICL

E D

IAM

ET

ER

, n

m

0.22C-0.12V

Isoth. trans. 650°C/500s

(lowdensity)

(highdensity)

0.10C-0.12V

Fig. 10 Electron micrographs showing V(C,N) precipitation after isothermal transformation at 650°C in 0,12%V-0,013%N steels with different C-contents; (a) 0.04% C, (b) 0.10% C and (c) 0.22% C , and the effect of N and C on the size and density of V(C,N)-precipitates during isothermal transformation at 650°C for 500s (d).

4. PRECIPITATION STRENGRHENING

As was stated above vanadium is one of the most widely used precipitation strengthening additives in microalloyed steels. A modest addition of 0.10%V can bring about a strength increase beyond 200 MPa, and in special cases even up to 300 MPa, Figs. 11 and 1219. These figures show another essential feature of V-steels, viz. that nitrogen in small contents adds significantly to precipitation strengthening. Similarly, enhanced cooling through the γ/α-transformation and afterwards augments the particle hardening, Fig. 12. The results of Fig. 11 show also that normalising at 950oC does not allow all V to go into solution completely and hence produces only modest strengthening.

(d)

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Fig. 11 Effect of processing method on the precipitation strengthening derived from V(C,N) for 0,12%C-0,35%Si-1,35%Mn-0,09%V-0,02%Al steel19.

Fig. 12 Influence of cooling rate on the strength contribution from V(C,N) precipitates. Base composition as for Fig. 1118.

The technically very important effect of N on the strengthening of V-steels is illustrated in Fig. 13 for different ageing temperatures. This strong effect of N has sometimes been interpreted such that only VN forms as precipitates, and that the remaining V does not combine with C, possibly due to the larger solubility and lower chemical driving force for VC. A physically more correct account would be as follows. In the present material with hard non-penetrable V(C,N)-precipitates the particle strengthening occurs by the Orowan mechanism – bowing of dislocations between particles – and in that case the decisive parameter is the interparticle spacing in the slip plane. In turn, that is determined by the density of the precipitates. This density is controlled by the nucleation frequency and the number of nuclei it creates until the supersaturation has diminished so that nucleation dies. The essential parameter governing the variation in nucleation is the chemical driving force for V(C,N)-precipitation which, as was shown above, depends on the available quantities of both C and

Fig. 13 Effect of N, V and transformation temperature on the precipitation strengthening in 0,1%C-V-N steels after isothermal ageing at different temperatures for 500s.

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N. It is true, however, that it varies relatively strongly with the N-content, but carbon may also significantly influence the chemical driving force.

Fig. 14 Deduced values of precipitation strengthening for isothermally transformed V-steels (650°C/500s). (a) as a function of N-content (b) as a function of C-content.

Recent investigations2,20 have demonstrated very clearly that the precipitation strengthening of V-steels increases significantly with the C-content of the steels. Plots of the observed strengthening ∆Rp against the C-content, and for comparison also N-content, are given in Fig. 14. These results show that carbon raises ∆Rp by ~ 5.5MPa for every 0.01%C in the steel, whereas nitrogen raises it by ~ 6 MPa for every 0.001%N. This important effect of carbon has not previously been noticed. The reason is presumably that there has always been an uncertainty in the procedure of evaluating the precipitation strengthening as to how the contribution of pearlite should be deduced and deducted from the measured yield strength. This procedure was improved in the present case, and furthermore nano-hardness measurements were performed in order to unequivocally demonstrate the precipitation hardening in the ferrite, Fig. 15. These results demonstrate very clearly that the precipitation strengthening of V-steels increases significantly with the C-content of the steels. The role of carbon in precipitation strengthening of V-steels is discussed below.

2

2.5

3

3.5

4

4.5

5

0 0.05 0.1 0.15 0.2 0.25 0.3

wt% C

NA

NO

-HA

RD

NE

SS

OF

F

ER

RIT

E,

GP

a

0.005%N

0.025%N

0.022%N

0.005%N

0.12% V

Isothermal transformation650°C/500s

Fig. 15 The effect of carbon and nitrogen contents on the nano-hardness of ferrite in 0.12%V.

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4.1 The Role of Carbon in Precipitation Strengthening

Carbon content has usually been considered not relevant to precipitation strengthening when the precipitation occurs in ferrite. This was deduced from the fact of very restricted solubility of carbon in ferrite (which is normally supposed to be independent of the total carbon content in the steel). Most published literature does not suggest that differences in carbon contents in the range for structural steels (0.04-0.3%) should affect significantly the response of vanadium in these steels3. However, it was suggested recently that the effective carbon for precipitation in ferrite may be much greater in the times available during phase transformation20. Fig 16(a) shows the considerable difference in solubility of C in ferrite when governed by the γ/α and cementite/α equilibria below A1. The solubility at 600oC is 5 times larger in γ/α than cementite/α. This metastable equilibrium between ferrite and undercooled austenite can greatly increase the solubility of carbon in ferrite thereby contributing to profuse nucleation of V(C,N) particles. The driving forces for nucleation of V(C,N) calculated for two different levels of dissolved carbon content equivalent to the equilibration ferrite+pearlite or metastable ferrite+austenite are shown in Fig. 16(b). In the extreme case of 250 ppm carbon in ferrite the driving force for nucleation increases by about 25-30% in comparison with the equilibrium content at 650°C. This demonstrates that the nucleation of V(C,N) is indeed a function of the carbon content dissolved in ferrite. In reality the carbon content in ferrite should lie somewhere between these values depending on transformation kinetics, mainly phase boundary mobility.

The total C-content affects the kinetics of the transformation of γ to α and an increase in total carbon content displaces the pearlite formation to longer times as was shown in Fig. 16(c). The implication of this is that a larger chemical driving force for V(C,N)-precipitation will remain longer and hence promote profuse nucleation. Consequently, denser V(C,N) precipitation nucleates in ferrite, as was shown in Fig. 10 for 0.22%C steels. This role of carbon in precipitation strengthening is particularly significant in steels with higher carbon contents where the metastable condition is more extreme and prolonged.

3.5

4

4.5

5

5.5

6

0 50 100 150 200 250 300

Nitrogen content in ferrite, ppm

DR

IVIN

G_F

OR

CE

fo

r V

(C,N

)

equilibrum carbon content in ferrite

250 ppm carbonin ferrite

(a) (b) (c)

Fig. 16 (a) Solvus lines for C in ferrite in equilibrium with cementite and austenite. (b) Chemical driving force, ∆Gm/RT, for precipitation of VC and VN in 0.12% V steel. (c) Effect of equilibrium and metastable carbon content in ferrite on driving force for precipitation of V(C,N) at 650°C.

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In practice the metastable α-γ equilibrium solubilities may not be so large as shown in Fig. 16(a) since diffusion of carbon in austenite occurs away from the interface tending to lower the concentration there. This will be accentuated if the interface migration is hindered by some effect such as solute drag of substitional elements. In the extreme case where boundary mobility is very low and full equilibration of carbon can take place both in ferrite and austenite the new metastable carbon levels can be calculated with the aid of Thermo-Calc (activities of carbon in ferrite and austenite are equal). Fig. 17 shows such calculation for two different carbon contents and demonstrates that the level of carbon dissolved in ferrite which governs the nucleation of V(C,N) is indeed a function of the total carbon content. In reality the carbon content in ferrite should lie somewhere between these values and the limit given in Fig. 16(a) (i.e. ~ 0.024 %C at 650°C) depending on phase boundary mobility.

The above situation continues until pearlite is nucleated at which point a new equilibrium with a reduced carbon activity is established. So fast is the diffusion of carbon in ferrite that the transition can take place in less than one second for typical microstructures. The great reduction in carbon activity at this stage is likely to virtually eliminate further nucleation of V(C,N) although some continued growth of existing particles may be expected.

The results in Fig. 17(a) also predict that the metastable carbon content in ferrite depends on the degree of transformation. In the early stage of transformation the carbon content is only slightly higher than the equilibrium level. As degree of transformation increases, ferrite is progressively enriched with carbon, from the shrinking austenite, so producing more nucleation of V(C,N) particles and accordingly a more dense precipitation. This was confirmed from micro-hardness measurements of partially transformed samples of low N steels, as shown in Fig. 17(b). In these ways the positive influence of total carbon content on the precipitation strengthening contribution can be understood.

0

0.005

0.01

0.015

0.02

0.025

0 10 20 30 40 50 60 70 80 90 100

% Transformed ferrite

wt%

CA

RB

ON

0.22%C

0.10%C

ferrite

(a)

150

160

170

180

190

200

210

220

230

240

250

0 50 100 150 200 250 300 350 400 450 500 550

TIME at 650°C, sec

MIC

RO

-HA

RD

NE

SS

OF

F

ER

RIT

E,

HV

15g

0.22C-0.12V-0.022N

0.22C-0.12V-0.0015N

IT at 650°C

(b)

Fig. 17 Values of the metastable carbon content in ferrite as a function of volume fraction of transformed ferrite at 650°C (a) and micro-hardness of ferrite as a function of transformation time (transformed ferrite) in high N and low N steels (b).

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5. GRAIN REFINEMENT

5.1 Formation of Intragranular Ferrite in V-Steels

The effect of V on the intragranular ferrite formation was examined on specimens isothermally treated at 850°C for 1 hour and partially transformed at the temperatures 700-450°C followed by gas quenching to room temperature. The resulting microstructures of the 0.22%C-0.12%V-0.013%N steel are shown in Fig. 18. High transformation temperatures resulted in largely reconstructive transformation at the prior austenite grain boundaries. The microstructure produced on a partial transformation at 700°C and gas quenching is shown in Fig. 18(a). It consisted of coarse grain boundary ferrite. Within the prior austenite grains a mixture of acicular ferrite with a second phase was formed as a result of the secondary transformation at a lower temperature during gas quenching. The microstructure obtained in his steel after a partial transformation at 650°C followed by gas quenching is shown in Fig. 18(b). Comparing to the specimen transformed at 700°C there was less grain boundary ferrite at the prior austenite grain boundaries, but a high density of intragranulary nucleated polygonal ferrite inside the prior austenite grains. Some of the intergranulary nucleated ferrite crystals were elongated with sharp edges which is probably characteristic of the faster cooling after the interrupted transformation. Quenching from 600°C produced a thinner band of grain boundary ferrite and higher number of intragranulary nucleated ferrite crystals, Fig. 18(c). There is also a change in the morphology of the intragranular ferrite which become less polygonal and more sideplate. However, it is not clear whether this plate morphology of ferrite was associated with reconstructive growth mechanism during isothermal transformation at this temperature or whether it formed at lower temperatures after the gas quench. There were also some ragged ferrite plates growing from the grain boundary ferrite allotriomorphs.

(a) 700°C (b) 650°C (c) 600°C (d) 550°C (e) 450°C

Fig. 18 Microstructures of 0.22%C-0.12%V-0.013%N steel austenitised at 1150°C, cooled at a rate of 2°C/s to 850°C, held at this temperature for 1 hour and partially isothermally transformed between 700-450°C prior to gas quenching to room temperature.

a

e

c b

d

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Fig. 18(d) shows the microstructure resulting from a partial isothermal transformation at 550°C followed by gas quenching. There was very little grain boundary nucleated allotriomorphs which a very thin skeleton of ferrite at the prior austenite boundaries. Within the prior austenite grains a high density of small ferrite crystals was nucleated. A large proportion of this intragranular ferrite was in the form of thin elongated plates, some of which were observed to have grown from particles. The microstructure produced on isothermal transformation at 450°C is shown in Fig. 18(e). Almost no grain boundary nucleated allotriomorphic ferrite is seen. The general matrix microstructure formed at this temperature was very fine acicular ferrite.

5.2 Number Density of Grain Boundary and Intragranular Ferrite

Metallographic examination of these specimens was carried out to determine the amounts of grain boundary ferrite, intragranular polygonal ferrite and intragranular acicular ferrite as a function of the transformation temperature. Measurements were made of the number density of ferrite grains to quantitatively describe the refining effect of intragranular ferrite formation. Fig. 19 summarises the effect of isothermal transformation temperature on the number density of grain boundary ferrite and intragranular ferrite. It can be seen that the grain boundary ferrite started to form directly below the transformation start temperature and the number density of grain boundary ferrite grains increased slightly with decreasing the transformation temperature. More complex behaviour was observed for intragranular ferrite. There was no intragranulary nucleated ferrite at higher transformation temperatures, above about 650°C. The temperature range where the large density of intragranular polygonal ferrite was formed was narrow, between 650-600°C. On further lowering the isothermal transformation temperature, below 600°C, there was a tendency for the intragranular ferrite to change morphology and grow as ferrite sideplates. The number density of intragranular ferrite grains at 550°C was higher than that for 600°C.

The number density and morphology of the ferrite grains changed considerably after isothermal transformation at 450°C. Cooling to 450°C led to the formation of exclusively acicular ferrite structure with a more than 4 times higher density of acicular grains. The tendency for grain boundary ferrite formation was reduced and it was difficult to estimate the number density of grain boundary ferrite grains as they have a similar acicular morphology.

0

20

40

60

80

100

120

400 500 600 700 800

Temperature, °C

Nu

mb

er o

f g

rain

bo

un

dar

y fe

rrit

e/10

0um

or

intr

agra

nu

lar

ferr

ite/

100u

m2 Intragranular ferrite

Grain boundary ferrite

acicular ferrite

Fig. 19 Number density of grain boundary ferrite and intragranular polygonal and acicular ferrite grains as a function of transformation temperature.

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5.3 Effect of Vanadium on the Formation of Acicular Ferrite.

The experimental results described above indicate that intragranular acicular ferrite was the dominant microstructure after isothermal transformation at 450°C. This treatment was repeated for steels with low and high nitrogen contents in order to verify the role of N as well as VN in the formation of acicular ferrite in V-steels. The resulting microstructures are shown in Fig. 20. It is clearly seen that a similar acicular ferrite structure was obtained in the V-microalloyed steel with a very low N content of 0.0015%N and high N content of 0.025%N. In the low N steel no VN precipitates were expected to grow in austenite. This result confirms the suggestions of He and Edmonds8 that there exists an effect of vanadium on the formation of acicular ferrite microstructure even in low nitrogen steels.

Fig. 20 Microstructure of low N-steel, and high N-steel isothermally transformed at 450°C.

5.4 Analysis of Nucleation Sites

Nucletion of intragranular ferrite and the morphology of nucleating precipitates/inclusions was examined in several of V-containing steels using both scanning and transmission electron microscopy. Fig. 21 shows a micrograph of the 0.10%C-0.12%V-0.025%N steel, which was water quenched from 630°C after ~5% transformation. In the intragranular regions, a number of inclusions exhibited ferrite growing with different morphologies. There were loops, blocks and caps of ferrite coating part of the inclusion surfaces as well as wedges and ferrite growing with more of a plate type of morphology. The formation of very long plates from individual inclusions was also observed.

A number of spherical non-metallic inclusions were present in the investigated steels. The majority of these inclusions were identified by EDS as manganese sulphides, Fig. 21. Other inclusions which were identified were oxides, mainly silicon-manganese oxides. Vanadium-rich particles, presumably nitrides were always visible in connection with inclusions. The fact that VN particles had formed in the austenite prior to transformation and nucleated ferrite crystals is clearly evident from SEM micrographs.

0.1C-0.12V-0.025N 0.22C-0.12V-0.0015N

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Fig. 21 Intragranular ferrite nucleus in 0.10%C-0.12%V-0.025%N steel, water quenched from 630°C after ~5% transformation.

6. SUMMARY

The present work has concentrated on two strengthening mechanisms in V-microalloyed steels: (i) grain refinement by promoting the formation of intragranular ferrite, and (ii) the role of nitrogen and carbon in precipitation strengthening by interphase and random precipitation of V(C,N) in ferrite.

The experimental results strongly indicate that vanadium can by effectively used not only for precipitation strengthening but also for ferrite grain refinement. It was shown that vanadium contributes to the formation of two types of intragranulary nucleated ferrite; polygonal (idiomorph) ferrite and acicular (sideplate) ferrite. Intragranular polygonal ferrite nucleates on VN particles which grow in austenite during isothermal holding or slow cooling throughout the austenite range. The intragranular polygonal ferrite forms in the narrow temperature range, between 650-600°C for the investigated compositions. Acicular ferrite microstructure forms in V-microalloyed steels during isothermal transformation at lower temperatures (~450°). The acicular ferrite microstructure was obtained in V-microalloyed steels containing high, medium or very low nitrogen levels. This suggests that vanadium on its own can promote the formation of the acicular ferrite microstructure.

Vanadium is an effective and easy controllable precipitation strengthening element. The degree of precipitation strengthening of ferrite for a given vanadium content depends on the available quantities of carbon and nitrogen. It was confirmed that nitrogen is a very reliable alloying element, increasing the yield strength of V-microalloyed steels by some 6 MPa for every 0.001% N, essentially independent of processing conditions. Experimentally it is demonstrated that the V(C,N)-precipitation becomes denser and finer with increasing N-content. Carbon content, on the other hand, has usually been considered not relevant to precipitation strengthening when the precipitation occurs in ferrite because of the very small carbon content in solution in ferrite at equilibrium. The present results have shown that the precipitation strengthening of V-microalloyed steels increases significantly with the total C-content, ~5.5MPa/0.01% C. The explanation is that the C-content of the steel delays the pearlite (ferrite+cementite) formation and thereby maintains the higher, metastable C-content in ferrite given by the austenite/ferrite equilibrium for a longer time. This effect of carbon is particularly significant for medium carbon steels typically used for hot rolled bars and sections.

MnS VN

α

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7. ACKNOWLEDGEMENTS

The present endeavour has been supported financially by the U.S. Vanadium Corporation /Stratcor/ and the Vanadium International Technical Comittee /VANITEC/. It is a pleasure to acknowledge the stimulating and challenging discussions with Michael Korchynsky of U.S. Vanadium and Peter Mitchell of VANITEC. Thanks are also due to Bevis Hutchinson and Rune Lagneborg for valuable discussions.

REFERENCES

1. S. Zajac, T. Siwecki and M. Korchynsky: “Importance of Nitrogen for Precipitation Phenomena in V-Microalloyed Steels”, in Proc. Conf. Low Carbon Steels for the 90`s, (eds. R Asfahani and G. Tither), ASM/TSM, Pittsburgh, USA, 1993, pp.139-150.

2. R. Lagneborg, T.Siwecki, S. Zajac and B. Hutchinson, “The Role of Vanadium in Microalloyed Steels”, Scandinavian Journal of Metallurgy, Vol 28, No. 5, October 1999, pp.1-241.

3. W. Roberts, “Recent Innovations in Alloy Design and Processing of Microalloyed Steels”, 1983 Int. Conf. on Technology and Applications of High Strength Low Alloy (HSLA) Steels, Philadelphia, PA, Oct. 3-6, 1983, pp. 67-84.

4. S. Zajac, T. Siwecki, B. Hutchinson, L-E. Svensson and M. Attlegård, ”The Influence of Plate Production Processing Route, Heat Input and Nitrogen on the HAZ Toughness in Ti-V Microalloyed Steel”, Swedish Institute for Metals Research, Report IM-2764, (1991).

5. T. Kimura, A. Ohmori, F. Kawabata, and K. Amano ”Ferrite Grain Refinement through Intra-granular Ferrite Transformation VN Precipitates in TMCP of HSLA Steel”. Thermec 97, (1997) pp.645-651.

6. T. Kimura, F. Kawabata, K. Amano, A. Ohmori, M. Okatsu, K. Uchida and T. Ishii, “Heavy Gauge H-Shapes with Excellent Seismic-Resistance for Building Structures Produced by the Third Generation TMCP”, CAMP-ISIJ, (1999) pp.165-171.

7. F. Ishikawa and T. Takahashi, ”The Formation of Intragranular Ferrite Plates in Medium-Carbon Steels for Hot-Forging and its Effect on the Toughness”, ISIJ International, (1995) pp.1128-1133.

8. K. He and D.V. Edmonds, “The Formation of Acicular Ferrite in C-Mn-V Steels”, 19th ASM Heat Treating Society Conf., Cincinnati, OH, (1999) pp.519-525.

9. T. N. Baker, "Precipitation and Transformation in Steels", Acta Metall., 23, (1973, p. 261.

10. M. Enomoto, “The Mechanisms of Ferrite Nucleation at Intragranular Inclusion in Steels”, International Conference on Thermomechanical Processing of Steel & Other Material, Edited by T. Chandra and T. Sakai, The Minerals, Metals & Materials Society, (1997) pp.427-433.

11. F. Ishikawa, T. Takahashi and T. Ochi, “Mechanism of Intragranular Ferrite Nucleation”, Solid-Solid Phase Transformations, Eds. W.C. Johnson, J.M. Howe, D.E. Laughlin and W.A. Soffa, The Minerals, Metals & Materials Society, (1994) pp.171-176.

12. S. Zajac, “Thermodynamic Model for the Precipitation of Carbonitrides in Microalloyed Steels”, Swedish Institute for Metals Research, Report IM-3566, Jan. 1998.

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13. A. T. Davenport and R. W. K. Honeycombe, “Precipitation of Carbides at Austenite/Ferrite Boundaries in Alloy Steels”, Proc. Roy, Soc. London 322 (1971) pp. 191-205.

14. R. W. K. Honeycombe, “Transformation from Austenite in Alloy Steels”, Metall.Trans.A, 7A (1976) pp. 915-936.

15. W. Roberts, Hot Deformation Studies on a V-Microalloyed Steel, Swedish Institute for Metals Research, Report IM-1333, 1978.

16. R. Lagneborg and S. Zajac, “A Model for Interphase Precipitation in V-Microalloyed Structural Steels”, Metall. and Mat. Trans. Vol. 32A, January 2001, pp.39-50.

17. R. M. Smith and D. P. Dunne, Structural Aspects of Alloy Carbonitride Precipitation in Microalloyed Steels, Materials Forum 11 (1988) pp.166-181.

18. W. Roberts, A. Sandberg, The Composition of V(C, N) as Precipitated in HSLA Steels Microalloyed with Vanadium, Swedish Institute for Metals Research, Report IM-1489 Oct. 1980.

19. W. Roberts, A. Sandberg, and T. Siwecki, Precipitation of V(C,N) in HSLA Steels Microalloyed with V, Proc. Conf. Vanadium Steels, Krakow, Oct. 1980, Vanitec, pp.D1-D12.

20. S. Zajac, T.Siwecki, W. B. Hutchinson, R. Lagneborg., “Strengthening Mechanisms in Vanadium Microalloyed Steels Intended for Long Products”, ISIJ International, 38 (1998) pp.1130-1139.

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The Evolution of Microstructure During Thin Slab Direct Rolling Processing in Vanadium Microalloyed Steels

Y. Li1), D. N. Crowther2), P. S. Mitchell3) and T. N. Baker1)

1) Metallurgy and Engineering Materials Group, Department of Mechanical Engineering,

University of Strathclyde, Glasgow, G1 1XJ, UK 2) Corus Group, Swinden Technology Centre, Moorgate, Rotherham, S60 3AQ, UK 3) Vanitec, Winterton House, High Street, Westerham, Kent, TN16 1AJ, UK

Abstract

The evolution of microstructure during a simulation of the thin slab direct rolling process has been studied on two low carbon steels, microalloyed with V-N and V-Ti-N. The steels were examined using optical microscopy, analytical transmission electron microscopy (TEM) and energy dispersive x-ray (EDAX).

After the 4th rolling pass, in a five pass schedule, the initial coarse austenite grain size (≈ 1 mm) was reduced to about 50 µm in Steel V-N and 22 µm in Steel V-Ti-N. The average ferrite grain size in the final strip was slightly smaller in Steel V-Ti-N (4.8-6.6 µm) than in Steel V-N (5.3-7.2 µm). For Steel V-N, VN was only observed after 1050°C equalization, but it was not found after 1200°C and 1100°C equalisation. For Steel V-Ti-N, V-Ti(N) particles formed during casting and during equalization for all the equalization temperatures (1200°C, 1100°C and 1050°C). AlN particles precipitated in Steel V-N only during 1050°C equalization and were often associated with MnS or MnS and VN. No AlN was detected in Steel V-Ti-N. Fine V containing precipitates (<10 nm) were observed in the final strip for both of the steels, but the frequency of the fine particles was lower in Steel V-Ti-N than in Steel V-N. The fine precipitates in the final strip make a major contribution to dispersion strengthening. High strength (LYS ≈ 460-560 MPa) with good toughness and good ductility were achieved in the steels, which are competitive to similar products made by conventional controlled rolling. However, the addition of Ti to the V-N steel decreased the yield strength due to formation of V-Ti(N) particles in austenite, which reduced the amounts of V and N available for subsequent V rich fine particle precipitation in ferrite.

KEY WORDS: Vanadium and vanadium-titanium microalloyed steel, thin slab direct rolling, equalization temperature, microstructure, properties.

1. Introduction

Thin slab direct rolling (TSDR) technology has changed the economics of steel production, because of low capital investment and higher productivity. The new technology has gained rapid worldwide acceptance and will continue to expand. Thin slab direct rolling process is especially attractive for production of microalloyed steels. High strength (>450 MPa) with good toughness and good ductility has been achieved in the vanadium microalloyed steels produced by thin slab direct rolling.1-2)

The thin slab direct rolling process route when compared with the conventional, thick-slab cold-charge rolling (CCR) implies some significant differences in the microstructural development. With TSDR, the slabs (30 to 80 mm) are thinner compared with 200 to 250 mm in CCR. Casting speeds are higher, 3.5-6.0 m/min in TSDR compared with 0.75-1.25 m/min in CCR. The thin slabs in TSDR solidify much more quickly, resulting in less segregation and a more homogenous microstructure.3) The rapid solidification of thin slabs also affects the morphology of inclusions and

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precipitates.4-6) In steels utilizing titanium technology for restriction of grain coarsening of austenite, the faster post solidification cooling rate is beneficial to ensure a critical size range dispersion of TiN particles.

CCR processing commences with cold charging of slabs, which must be reheated after casting and cooling to room temperature, before hot rolling. With TSDR processing, the slabs, still hot from the caster, are direct hot charged to an equalization furnace without cooling to ambient temperatures as in CCR, leading to a saving in energy. However, the relatively coarse, as-cast microstructure is retained during the TSDR process, whilst for CCR, the as-cast microstructure is significantly refined as a result of the slab cooling through the transformation temperature and then being reheated prior to rolling. A coarser austenite grain has been reported after casting and equalization in TSDR processing. Also a smaller total strain (deformation) is available in TSDR, therefore, a complete rescheduling of the rolling pass and their reductions at pre-set temperatures is necessary compared with CCR process. 7-8)

A further difference concerns the microalloying, which can create important changes in steel properties. In the TSDR process, the as-cast austenite microstructure prior to rolling is generated at a temperature in excess of 1450°C, while the equilibrium solubility of microalloy carbonitrides is very much greater than that at the soaking temperature used in the CCR process. Most TSDR processing chooses steels with carbon content less than 0.065 wt.% to avoid the peritectic reaction and subsequent segregation. Also in the TSDR process, the as-cast austenite prior to rolling may be more highly supersaturated with respect to microalloying elements than the reheated austenite in the CCR process. This can affect subsequent microstructural development during processing.

The addition of Nb to HSLA steel can give considerable strengthening, but when Nb is present in continuously cast HSLA steels, slab surface cracking, especially in the transverse direction, is a well documented observation.9) Attempts to produce acceptable surface finishes in Nb microalloyed steels have not been completely successful to date.7,10) This is associated with the precipitation of Nb compounds in a manner similar to that responsible for the ductility trough found during hot tensile testing of CCR processed steels in the temperature range from 750°C to 925°C.1,11) For this reason, the use of V additions has been explored.10) Ti additions have been used widely in HSLA steels to control austenite grain size. However, it has been reported that when Ti is present in a steel which contains other microalloying elements such as V and Nb, the Ti addition changes the precipitation of V and Nb, and results in a lower yield strength.12-18)

Most of the work done so far in the development of the thin slab direct rolling technique has been concerned with the process parameters and there are no reports in the literature on a detailed study of the evolution of microstructure over the entire thin slab direct rolling process. The aim of this project was to investigate microstructural development in V-N and V-Ti-N microalloyed steels during the thin slab direct rolling process.

2. Experimental Methods

2.1 Materials

The study was carried out on two low carbon steels with 0.1%V or 0.1%V and 0.008%Ti. Both steels contained about 0.02% N. The chemical compositions of the steels are given in Table 1. The thin slab direct rolling processing was simulated at Corus Group, Swinden Technology Centre and a schematic diagram representing the process adopted in the present work is shown in Fig. 1. The steels were melted in air, and cast into moulds to produce 50 mm thick ingots. The typical cooling rate at the mid thickness position of the ingots was 3.5°C/s. The ingots were equalized in a furnace set at temperatures of 1050°C, 1100°C or 1200°C for 30-60 minutes prior to rolling. After equalization, the ingots were rolled on a laboratory reversing mill into 7 mm strips by 5 passes,

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which gave a total reduction of 86%. A typical interpass time was 6s. After the 4th pass, the strips were held until a temperature of approximately 870°C was reached, and the holding times were approximately 25-40s. Finish rolling temperatures varied from 850°C to 880°C and the total rolling times were in the range of 75-90 s. After rolling, the strip was cooled under water sprays to simulate run-out table cooling. The end cool temperature of the strip ranged from 540 to 720°C. Following cooling, the strips were immediately put into a furnace set at 600°C and slow cooled (average cooling rate between 600-400°C was 35°C/h) to simulate coiling. Samples were quenched after casting, after equalization, after 4th rolling pass and after coiling, to follow the evolution of microstructure. The temperatures for the steels at different stages in the processing are given in Table 2.

2.2 Microstructural Examination

The microstructures in the steels were examined using optical microscopy and analytical transmission electron microscopy (TEM). Samples were taken from stages A, B, C and D of the process shown in Fig. 1. Samples were cut from ¼ thickness of the ingots or ¼ width of the strips. For optical microscopy, samples were etched in a 2% nital solution after polishing. For transmission electron microscopy, carbon extraction replicas were prepared in a four-step procedure by etching the polished specimen surface in 2% nital, coating the surface with a thin film of carbon, stripping the thin film in 5% nital and then cleaning the film in both alcohol and distilled water. The austenite grain size in the as-cast ingots and the specimens after the 4th pass, and the ferrite grain size in the final strip, were measured using a linear intercept method. Precipitates in the steels at different stages during the process were studied on carbon extraction replicas using an analytical Philips EM-400 TEM with an Energy Dispersive X-ray (EDAX) attachment. The composition of the precipitates was analysed using EDAX. The size of the precipitates was measured on TEM micrographs using the image analyser with software of Image-Pro Plus.

2.3 Tensile and Charpy Tests

To assess mechanical properties in the final strip, duplicate, transverse, full thickness tensile test pieces with width of 12.5 mm and a gauge length of 50 mm were tested. Longitudinal 10×5 mm Charpy test pieces (2mm V notch) were tested in accordance with BS EN 10045 to produce a complete transition curve.

3. Results

3.1 Microstructure

3.1.1 After Casting (Stage A)

The microstructure in the quenched as-cast ingot consisted of equiaxed grains near the slab surface of approximately 300 µm in diameter, columnar grains with lengths in the range from 5 to 10 mm, and a central zone of coarse, equiaxed grains19). Fig. 2 shows an example of equiaxed grains in the central zone of the as-cast ingot. For both Steel V-N and Steel V-Ti-N, the average prior austenite grain size in the ¼ thickness position of the ingots was about 1 mm. For Steel V-N, only alumina particles were found. The alumina particles formed during solidification and the size of the particles was 0.5 to 1.5 µm. For Steel V-Ti-N, the main precipitation was in the form of irregularly shaped

particles (Fig. 3), which contained both vanadium and titanium, and the TiVV+

ratio was about 0.5.

The size of the irregularly shaped particles ranged from 40-200 nm, and the particles usually were

present in rows. A small number of Ti rich (TiVV+

≈ 0.07) cuboid particles (Fig. 4) were also

observed. Most of the Ti rich particles had a size ranged from 50 to 350 nm, but a few of the

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particles were about 1 µm. The third type of particle in Steel V-Ti-N at this stage was large

dendritic particles (Fig. 5). The core and arm of the dendrites had same TiVV+

ratio, which was 0.3-

0.4.

3.1.2 After Equalization (Stage B)

After equalization, the microstructure and prior austenite grain size were similar to that in the as-cast ingots. For Steel V-N, no carbide or nitride precipitates were identified in the specimens after 1100°C or 1200°C equalization.2) VN was only observed in the specimen after 1050°C equalization. VN particles precipitated mainly along prior austenite grain boundaries (Fig. 6). The size of VN particles was from 10 to 80 nm and the average size was 24 nm. A few of VN particles were associated with MnS. AlN associated with MnS (Fig. 7), or MnS and VN, were also found in the specimen after 1050°C equalization.2)

For Steel V-Ti-N, spherical particles precipitated during equalization at 1200°C and were randomly

distributed in the matrix. The size of the spherical particles was from 20 to 50 nm and the TiVV+

ratio was about 0.6. Cruciform particles (Fig. 8a) were the main precipitate type observed during 1050°C and 1100°C equalizations. The EDAX spectrum from the cruciform particle is shown in Fig. 8b. For 1050°C equalization, the length of the cruciform arm was from 20 to 50 nm and the

TiVV+

ratio was 0.7. For 1100°C equalization, the length of the cruciform arm was from 30 to 150

nm and the TiVV+

ratio was 0.6. Complex particles in the matrix (Fig. 9) were observed after

1050°C equalization. Most of the complex particles contained V, Ti and S and a few also contained Mn. Dendritic particles in the specimen after 1050°C equalization had different compositions of

core and arm. The core was enriched with Ti and the TiVV+

ratio was about 0.45, while the arm had

a higher TiVV+

ratio, which was about 0.7-0.9.

3.1.3 After 4th Pass Rolling and Holding (Stage C)

A fine and uniform prior austenite grain structure was observed in the specimens after the 4th pass for both of the steels (Fig. 10) and this implied that recrystallisation occurred, and repeated recrystallisation eliminated the initial coarse, as cast structure. The average prior austenite grain size for the steels is given in Table 3. Steel V-Ti-N had a smaller prior austenite size than Steel V-N. No major precipitation occurred between the start of rolling and the end of the 4th pass in both of the steels. For Steel V-N, no strain or interphase precipitation was observed. The VN particles precipitated in the specimen after 1050°C equalization had grown during rolling. The particle size ranged from 30 to 80 nm and had an average size of 55 nm. For Steel V-Ti-N, strain induced fine cuboid particles were observed in the samples for 1100°C and 1200°C equalizations. The average

size of the strain induced particles was 10 to 20 nm and the TiVV+

ratio was 0.6-0.7. Cuboid

particles precipitated in rows or individually in Steel V-Ti-N for 1050°C equalization. The size of

the particles was 20-80 nm and the TiVV+

ratio was about 0.7. The particles observed in the

previous stages have grown and agglomerated during and/or after rolling.

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3.1.4 After Coiling (Stage D)

The microstructure in the final strip consisted of ferrite and pearlite, Fig. 11. The average ferrite grain size for the steels is given in Table 3. Steel V-Ti-N had a finer ferrite grain size than Steel V-N. Fine precipitates in the matrix, with a size range 2-20 nm, were observed in both of the steels (Fig. 12a). The fine particles were mainly VN20) in Steel V-N and V rich V-Ti nitrides in Steel V-Ti-N. An EDAX spectrum from the fine particles in Steel V-Ti-N is shown in Fig. 12b and the

weight ratio ofTiVV+

for the fine particles was 0.85-0.95. The frequency of the fine particles was

lower in Steel V-Ti-N than in Steel V-N, especially for Steel V-Ti-N equalized at 1050°C. It is considered that these fine particles make a major contribution to dispersion strengthening.

A summary of the precipitates, which occurred in the steels at different stages during the processing, is given in Table 4.

3.2 Tensile Properties and Charpy Toughness

Tensile and Charpy tests were carried out on the final strip and the results are given in Table 5. Lower yield strength in the range from 461 to 557 MPa, tensile strengths of 571 to 664 and elongations of 18 to 27% were achieved in the final strip, together with a good Charpy toughness. The 13J impact transition temperature was from –40 to −120°C. The mechanical properties of the steels are competitive to the similar products made by conventional controlled rolling. For example, steel strips (6.3 mm) with 0.08-0.17% C, 0.03-0.14%V and 0.006-0.022%N produced by controlled rolling with similar rolling conditions had yield strength of 345 to 550 MPa and elongation of 26 to 29%.21)

The equalization temperature had little or no significant influence on the yield strength and Charpy toughness of the steels. Figs. 13-14 show the effects of end cool temperature on yield strength and Charpy toughness. There was an increase in yield strength as the end cool temperature was reduced. However, Charpy 13J impact transition temperature was not strongly influenced by end cool temperature. At a similar end cool temperature, Steel V-Ti-N had lower yield strength than Steel V-N, but improved Charpy toughness.

4. Discussion

4.1 Microstructure

The initial coarse austenite (≈ 1 mm) was refined by the 4 rolling passes. The fine and uniform austenite grains in the strip after 4th pass rolling implied that recrystallization occurred and the repeated recrystallization eliminated the initial coarse, as cast structure. The average austenite grain size (20-22 µm) in Steel V-Ti-N was finer than that (40-50 µm) in Steel V-N. This is due to precipitation of V-Ti nitride particles in austenite at higher temperatures. These V-Ti nitrides had no effect on recrystallization, but could have prevented austenite grain growth. The final strip had a fine ferrite grain size, which was relatively insensitive to the equalization temperature. The average ferrite grain size (4.8-6.6 µm) in Steel V-Ti-N was also finer than that (5.3-7.2 µm) in Steel V-N due to a smaller austenite grain size before γ−α transformation.

Solution temperatures of the carbides and nitrides, and weight percent of precipitates at various temperatures in austenite for the steels were calculated using the thermodynamic based software (ChemSage). The solubility data recommended by Turkdogan22) have been used in the present study. The calculated results are shown in Figs. 15 and 16 for the steels. It can be seen that the precipitation of V was different in the two steels as a result of Ti addition. The thermodynamic calculation also indicted that no carbon was out of solution at a temperature above 900°C.

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The calculations indicated that the precipitates present in austenite above 900°C should be nitrides for both of the steels. According to the calculations, if AlN is present, then the solution temperatures of VN and AlN in austenite should be 1101°C and 1343°C, respectively, for Steel V-N. As expected, VN precipitation was only observed in the specimen after 1050°C equalization in Steel V-N.20) However, the solution temperature of AlN was higher than the equalization temperatures, but AlN particles only precipitated on substrates in the specimens after 1050°C equalization in Steel V-N. The absence of AlN precipitation in the specimens after 1100°C and 1200°C equalization may be due to the fact that AlN has a close packed hexagonal structure and nucleates with some difficulty in austenite. The nucleation of AlN is controlled more by kinetics than thermodynamics. AlN particles observed in the steel were associated with MnS particles. According to the solubility data of MnS,23) MnS particles should form before AlN, and therefore MnS particles may act as nucleation sites for AlN particles. No major precipitation occurred in Steel V-N between the start of rolling and the end of the 4thPass rolling, despite of the rolling temperature being below the calculated solution temperature of VN in austenite.

Steel V-N and Steel V-Ti-N had similar amounts of V and N, but the ChemSage calculation showed that the temperature for V starting to precipitate in austenite (1444°C) in Steel V-Ti-N was much higher than that (1101°C) in Steel V-N due to the Ti present in Steel V-Ti-N. Also when V and Ti are both present in the steel, the stable microalloyed precipitate in austenite would be the mixed V-Ti nitride. The calculation was in agreement with the experimental results. V-Ti nitrides were found in the as-cast specimen and in the specimens after all the equalization temperatures in Steel V-Ti-N. Also strain induced fine V-Ti nitrides formed during rolling.

No AlN was observed in Steel V-Ti-N. The absence of AlN in Steel V-Ti-N may be due to the high misfit of AlN in austenite compared with V-Ti compounds, and a consequent increase in surface energy. Also, the solution temperature of V-Ti nitrides in austenite (1444°C) was higher than the solution temperature of AlN in austenite (1238°C). There was less N left for formation of AlN in Steel V-Ti-N due to the prior formation of V-Ti nitrides in austenite at higher temperatures, so that the driving force for nucleation of AlN in austenite would be lower in Steel V-Ti-N than that in Steel V-N. This would be another reason for the absence of AlN in Steel V-Ti-N.

According to the ChemSage calculations, the equilibrium precipitate in austenite for Steel V-Ti-N should be richer in Ti at higher temperatures and richer in V at low temperatures (Fig. 17). Therefore, the Ti rich V-Ti(N) particles, would be expected to commence nucleation at high temperatures in the austenite phase, and grow on cooling, as processing proceeded. In fact, the thermodynamic calculations were in agreement with the EDAX results (Fig. 17). The V-Ti(N)

particles in the as-cast specimens had a TiVV+

ratio of 0.1-0.5, the V-Ti(N) particles forming during

equalization had a TiVV+

ratio of 0.45-0.8, and the fine (2-20 nm) particles in the final strip had a

TiVV+

ratio of 0.85-0.98. The limits of the dashed lines for dendrites and irregularly shaped

particles represent the solution temperatures of TiN (1413°C) and cuboids (1100°C).

4.2 Tensile Properties and Charpy Toughness

The level of dispersion strengthening (σP) can be approximately estimated by subtracting the components of yield strength (σy) due to ferrite lattice fraction stress (σo), solid solution strengthening (σs) and ferrite grain size strengthening (σg) from the measured value using a modified version of the Hall-Petch equation, with the units in MPa.

σP = σy − (σo + σs + σg) (1)

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σo = 45 MPa 24) (2)

σs = 84(Si) + 32(Mn)+680(P)+38(Cu)-43(Ni) 25-26) (3)

σg = 18.1 d 21

− 27) (4)

The dispersion strengthening could also contain strengthening contributions due to other strengthening mechanisms such as dislocation strengthening.

The relationships between σy, σs, σP and equalization temperature are given in Fig. 18. It can be seen that Steel V-Ti-N showed a reduction in yield strength and dispersion strengthening compared with Steel V-N at given equalization temperatures, especially at 1050°C. However ferrite grain size strengthening was slightly higher in Steel V-Ti-N than that in Steel V-N.

Fig. 19 shows the influences of end cool temperature on dispersion strengthening. Dispersion strength decreased as the end cool temperature was increased, and at a similar end cool temperature, Steel V-Ti-N showed significantly lower dispersion strength than Steel V-N. Two major factors, which control the dispersion strengthening, are volume fraction and mean size of the fine particles in the steels. Higher volume fraction and smaller mean size of the particles could result higher dispersion strengthening. For Steel V-Ti-N, the stable V-Ti nitrides, which precipitated in austenite at higher temperature, limited austenite grain growth during or after rolling, but they were too big to have a significant contribution to the strength of the final strip. Instead those large particles removed a significant fraction of V and N from solution before the austenite to ferrite transformation and reduced the amount of V and N available to precipitate in ferrite to provide dispersion strengthening. Thermodynamic calculations using ChemSage indicted that for Steel V-N, no carbonitride precipitates were predicted to form during 1100°C and 1200°C equalizations. VN particles, which precipitated during the 1050°C equalization contained 11% V. For Steel V-Ti-N, V-Ti nitride particles started to precipitate at a very high temperature (1444°C). After equalization, the particles in Steel V-Ti-N removed a significant amount of V from solution (7% V for 1200°C equalization, 21% V for 1100°C equalization and 30% V for 1050°C equalization), which left less V to precipitate in ferrite. This resulted a lower volume fraction of fine precipitates in Steel V-Ti-N compared with Steel V-N. The mean size of the fine particles was mainly dependent on the end cool temperature. The fine particles were smaller in the samples with the lower end cool temperatures. For a similar end cool temperature, there was no obvious difference in the particle size between the two steels. The loss of dispersion strengthening resulted a lower yield strength in Steel V-Ti-N. However, the reduction in dispersion strengthening and the slight refinement of ferrite grain size have improved the Charpy properties in Steel V-Ti-N.

The results in the present study are in accordance with previous work.12-18) Wang15) showed that in the thermomechanically controlled rolled (TMCR) condition, a titanium addition of 0.015 % to steels containing 0.08%V and 0.008%N resulted in a loss of yield strength of approximately 10-40 MPa. Crowther and Morrison16) also reported that in the vanadium steels, as was previously observed by He and Baker for niobium steels,28-29) there was some grain refinement associated with a titanium addition, but the loss of dispersion strengthening out-weighed the benefits to strength expected from the reduced grain size. Despite grain refinement and a reduction in dispersion strengthening for the V-Ti steels, the impact properties of the steels were not always improved significantly and it is suggested that this may reflect the embrittling effect of large TiN particles (>1µm), which are often present in the titanium-treated steels. Swedish work, however, has shown that by using a lower reheating temperature and a modified rolling schedule, attractive properties can be developed in V-Ti steels.12, 14)

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5 Conclusions

The evolution of microstructure during the simulated thin slab direct rolling has been studied in two low carbon V-N and V-Ti-N microalloyed steels. The results obtained are summarized as follows:

1. No carbide or nitride precipitation was identified in Steel V-N after 1100°C or 1200°C equalization. However, AlN and VN were observed in Steel V-N after 1050°C equalization.

2. The precipitation of V in Steel V-Ti-N was modified significantly because of V-Ti(N) particles formed in austenite at higher temperatures as a result of Ti addition to the steel.

3. The initial coarse austenite (≈1mm) was refined by 4 rolling passes due to rapid recrystallization. The final strip had a ferrite grain size of 4.8-7.2 µm.

4. Precipitation of complex V-Ti(N) particles in austenite had no effect on recrystallization, but restricted growth of recrystallised austenite grains during and after rolling.

5. Fine V(C,N) and V-Ti(C,N) precipitates in the range 2-10 nm in the final strip, make a major contribution of >90 MPa to dispersion strengthening.

6. The addition of Ti to the V-N steel decreased the yield strength. The reduction of yield strength was due to formation of V-Ti(N) particles in austenite, which removed a significant fraction of V and N from solution, and thereby reduced the amount of V and N available to precipitate in ferrite to providing dispersion strengthening.

7. The present study showed that a combination of high strength with good toughness and good ductility can be achieved in vanadium microalloyed steels produced by a simulated thin slab direct rolling processing. The mechanical properties of the steels are competitive to the similar products made by conventional controlled rolled processing.

REFERENCES

1) D. N. Crowther, P.S. Mitchell and W. B. Morrison: in Proceedings of the Int. Conf. 39th Mechanical Working and Steel Processing, Iron & Steel Soc. Of AIME, Warrendale, PA, USA, (1998), 839.

2) Y. Li, D. N. Crowther, P. S. Mitchell and T. N. Baker: in Proceedings of the Int. Conf. HSLA Steels’ 2000, Xian, China, (2000), 326.

3) R. Kaspar and O. Pawelski: in Proceedings of the Int. Conf. METEC Congress, VDEH, Dusseldorf, Germany, (1994), 390.

4) M. Korchynsky, Scandinavian Journal of Metallurgy, 28(1998), 40. 5) A. M. Sage, R. C. Cochrane and D. S. Howse: in Proceedings of the Int. Conf. Processing

Microstructure and Properties of Microalloyed and Other Modern High Strength Low Alloy Steels, Iron & Steel Soc, Pittsburgh PA, USA, (1991), 443.

6) V. Leroy and J. C. Herman: In Proceedings of the Int. Conf. Microalloying 95, Iron & Steel Soc, Pittsburgh PA, USA, (1995), 213.

7) P. J. Lubensky, S. L. Wigman and D. J. Johnson: in Proceedings of the Int. Conf. Microalloyed ’95 Iron & Steel Soc, Pittsburgh PA, USA, (1995), 225.

8) R. Kasper, N. Zentarat and J. C. Herman: Steel Research, 65(1994), 279. 9) K. Kunishige and N. Nagao, ISIJ International, 29(1989), 940. 10) T. Nilsson: In Proceedings of the Int. Conf. Tech. & Applic., ASME, (1983), 253. 11) B Mintz and J M Arrowsmith: Met. Tech., 6(1979), 24.

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12) W. Roberts: in Proceedings of Int. Conf. Technology and Applications of HSLA Steels, Metals Park, OH, USA, (1984), 33.

13) R. Lagneborg, T. Siwecki, S. Zajac and B. Hutchinson: Scandinavian Journal of Metallurgy, 28(1999), 186.

14) T. Siwecki, A. Sanderg and W. Roberts: in Proceedings of Int. Conf. Technology and Applications of HSLA Steels, Metals Park, OH, USA, (1984), 619.

15) S. C. Wang: J. Mat. Sci. (1990), 187. 16) D. N. Crowther and W. B. Morrison: Titanium Technology in Microalloyed Steels, ed. by T. N.

Baker, IoM, London, (1997), 44. 17) R. L. Bodnar and S. S. Hansen: in Proceedings of Int. Conf. 35th Mechanical Working and Steel

Processing, Pittsburgh, Iron and Steel Society, (1993), 495. 18) S. Zajac, T. Siwecki, W. B. Hutchinson and R. Lagneborg: ISIJ International, 38(1998), 1130. 19) D. N. Crowther, Y. Li, T. N. Baker, M. J. W. Green and P. S. Mitchell, in Proceedings of Int.

Conf. Thermomechanical Processing of Steels, IOM, London, (2000), 527. 20) J. Wilson, A. J. Craven: Y. Li, T. N. Baker, in Proceedings of Int. Conf. EMAG 2001, Dundee,

UK, (2001). 21) J. D. Grozier: in Proceedings of Int. Conf. Microalloying 75, New York, (1977), 241. 22) E. T. Turkdogan: Iron & Steelmaker, 16 (1989), 61. 23) L. D. Frawley and R. Priestner: Materials Science Forum, 284-286(1998), 485. 24) A. Cracknell and N. J. Petch: Acta. Met., 3(1955), 186. 25) F. B. Pickering and T. Gladman: ISI Special Rep. 81(1963), 181. 26) W. B. Leslie: Metall. Trans., (1972), 5. 27) W. B. Morrison and J. A. Chapman: Rosenhain Centen Conf. Contrib of Phy Metall to Eng

Pract, the Royal Society, London, (1976), 289. 28) K. J. He and T. N. Baker: Mat. Sci. Eng., A169 (1993), 53. 29) K. J. He and T. N. Baker: Titanium Technology in Microalloyed Steels, ed. by T. N. Baker,

IoM, London, (1997), 115.

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Table 1 The chemical compositions of the steels. (mass%)

Steel C Si Mn P S Cr Mo Ni

V-N 0.068 0.37 1.44 0.013 0.004 0.09 0.02 0.07

V-Ti-N 0.079 0.40 1.40 0.013 0.007 0.09 0.02 0.07

Steel Al B Cu N Nb Ti V O

V-N 0.03 <0.0005 0.07 0.022 <0.005 <0.005 0.10 0.0096

V-Ti-N 0.021 <0.0005 0.07 0.020 <0.005 0.008 0.10 0.0058 Table 2 The processing conditions for the steels.

Steel Steel V-N Steel V-Ti-N

Furnace entry T, °C 899 980 932 957 1018 916

Equalization T, °C 1050 1100 1200 1050 1100 1200

Equalization time, min 60 32 30 41 47 43

Start hold T, °C 843 937 1015 963 954 1020

Finish rolling T, °C 837 854 882 886 873 840

End cool T, °C 602 720 646 609 537 643 Table 3 Prior austenite grain size after 4th pass rolling and ferrite grain size in the final strip.

Steel Steel V-N Steel V-Ti-N

Equalization T, °C 1050 1100 1200 1050 1100 1200

Prior Austenite Grain, µm 40 40 50 20 22 22

Ferrite Grain Size, µm 6.2 5.3 7.2 6.0 4.8 6.6

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Table 4 Summary of the precipitates which occurred in the steels at different stages during the processing.

Steel V-N Steel V-Ti-N (R=

TiVV+

)

As Cast

1) Al2O3, 0 5-1.5 µm 1) Al2O3

2) V-Ti(N)

(a) Irregular shaped, 40-200 nm, R ≈ 0.5

(b) Cuboid, 50-350 nm, R ≈ 0.07

(c) Dendrites, R ≈ 0.3-0.4

Equalization Temperature 1050°C 1100°C 1200°C 1050°C 1100°C 1200°C

After Equalization

1) MnS 50-150 nm

2) CuS 50-250 nm

3) AlN+MnS

4) VN 10=80 nm

1) MnS 50-150 nm

2) CuS 50-250 nm

1) MnS 50-150 nm

2) CuS 50-250 nm

1) V-Ti(N)

(a) Cruciform 20-50 nm R≈0.7

(b) Complex with V, Ti, S, Mn

1) V-Ti(N)

(a) Cruciform 30-150 nm R≈0.6

1) V-Ti(N)

(a) Spherical 20-50 nm R≈0.6

After 4th Pass

1) VN 30-80 nm

1) V-Ti(N)

Cuboids 20-80 nm R≈0.7

1) V-Ti(N)

Strain induced cuboid 10-20 nm R≈0.6-0.7

1) V-Ti(N)

Strain induced cuboid 10-20 nm R≈0.6-0.7

Final Strip 1) V(C,N) 2-20 nm

2) AlN+VN

V(C,N) 2-20 nm

V(C,N) 2-20 nm

V-Ti(C,N) 2-20 nm R≈0.95

V-Ti(C,N) 2-20 nm R≈0.85

V-Ti(C,N) 2-20 nm R≈0.9

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Table 5 Tensile and Charpy properties.

Tensile Properties

Steel Steel V-N Steel V-Ti-N

Equalization T, °C 1050 1100 1200 1050 1100 1200

Av. LYS (MPa) 557 527 489 459 522 461

Av. UYS (MPa) 570 492 482 526 462

Av. UTS (MPa) 664 644 631 578 609 571

Av. EL (%) 24 24 19 26 18 27

Charpy Toughness

Steel Steel V-N Steel V-Ti-N

Equalization T, °C 1050 1100 1200 1050 1100 1200

J@, −20°C 43 35 27 71 43 45

13J ITT, °C −85 −45 −40 −120 −90 −100

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Fig. 1 Schematic diagram showing the thin slab direct rolling processing adopted in this study.

Fig. 2 Optical micrograph showing the microstructure in the central zone of the as-cast ingot.

Fig. 3 TEM micrograph showing the irregularly shaped particles on the prior austenite grain

boundary in Steel V-Ti-N.

Melting

Casting Equalization

Rolling Cooling Coiling

A B C D

50 nm

500 µm

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Fig. 4. TEM micrograph showing the Ti rich particles in Steel V-Ti-N.

Fig. 5 TEM micrograph showing the dendritic particle in Steel V-Ti-N.

Fig. 6 TEM micrograph showing VN particles in Steel V-N after equalization at 1050°C.

200 nm

200 nm

100 nm

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Fig. 7 TEM micrograph showing an AlN particle associated with a MnS particle in Steel V-N

after equalization at 1050°C.

Fig. 8a TEM micrograph showing a cruciform particle in Steel V-Ti-N after equalization at

1100°C.

Fig. 8b EDAX spectrum from the cruciform in Fig.8a.

100 nm

50 nm

AlN

MnS

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Fig. 9 TEM micrograph showing a complex particle in Steel V-Ti-N after equalization at 1050°C.

Fig. 10 Optical micrograph showing the microstructure in the strip after 4th pass rolling for

Steel V-N.

Fig. 11 Optical micrograph showing the ferrite and pearlite microstructure in the final strip

for Steel V-N.

100 nm

10 µm

10 µm

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Fig. 12a TEM micrograph showing fine and coarse precipitates in the final strip for Steel V-N.

Fig. 12b EDAX spectrum from the fine particles in Steel V-Ti-N.

100 nm

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Fig. 13 Effect of end cool temperature on yield strength.

Fig. 14 Effect of end cool temperature on Charpy toughness.

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Fig. 15 Calculated precipitation in Steel V-N.

Fig. 16 Calculated precipitation in Steel V-Ti-N.

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Fig. 17 Weight ratio of TiVV+

for the precipitates in Steel V-Ti-N.

Fig. 18 Relation between strength and equalization temperature.

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Fig. 19 Effect of end cool temperature on dispersion strength.

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The Effects of Microalloying Elements on Cracking During Continuous Casting

D. N. Crowther

Corus Group, Swinden Technology Centre, Moorgate, Rotherham, S60 3AQ, UK

SUMMARY

To ensure the appropriate quality in finished products, it is important that defects in continuously cast products are minimised. As the use of hot charging and thin slab rolling coupled with direct rolling becomes more common, it is increasingly important to produce defect free continuously cast product, as inspection and repair in these situations becomes more difficult.

Of the many types of defect in continuously cast products, only transverse surface cracking is strongly influenced by the presence of microalloying elements. Nb has a particularly strong detrimental effect, and Nb additions of as low as 0.01% can promote cracking. For V steels with <0.005%N, transverse cracking does not appear to occur, although at high levels of V and N (0.15%V, 0.02%N), transverse cracking has been reported.

It is believed that transverse cracks form in the mould, and propagate later in the continuous casting process, particularly during the straightening process. Microalloyed steels can exhibit low ductility over certain temperature ranges, and when the straightening process is carried out in this low ductility region, cracking can occur. In this respect, Nb has a strong effect in deepening the ductility trough, and extending it to higher temperatures. This behaviour is due to the presence of Nb(CN) precipitates, which promote low ductility failures, and retard recrystallisation. The effect of V on hot ductility is much less marked, and only at high levels of V and N does their ductility approach that found in Nb steels. V additions to Nb steels appear to slightly improve hot ductility, by promoting coarser precipitates. The effects of Ti on hot ductility are complex and still not completely understood.

Transverse cracking may be minimised by appropriate selection of steel composition, such as minimising Nb, replacing Nb by V and N combinations, or by making V additions to Nb steels. Machine operating conditions such as secondary cooling strategy are also important in avoiding transverse cracking. By selecting straightening temperatures, which are outside the temperature range of low hot ductility, cracking can be reduced.

1. INTRODUCTION

During the production of continuously cast products, it is very important to avoid both surface and internal defects, as otherwise expensive and time consuming slab or bloom/billet repair operations are required, or defective final product may be produced. The production of defect free continuously cast products is becoming ever more important as the use of hot charging and direct rolling from thin slab casters increases. In these situations, inspection and repair of defects in continuously cast slab becomes more difficult, and the production of defect free continuously cast products is of vital importance.

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Some high strength, microalloyed steels are particularly prone to some types of continuous casting defect, and there have been a number of excellent reviews on the topics of hot ductility and defects in this type of continuously cast product.1-3) In such steels, it is found that the type of microalloying elements used, and the overall composition of the steel, is very important in controlling the amount of defects.

The objectives of this report are to review briefly the effects of the microalloying elements V, Nb and Ti on the formation of defects in continuously cast products, the mechanisms by which these microalloying elements influence defects, and to identify possible ways in which defect-free continuously cast products can be produced.

2. DEFECTS IN CONTINUOUSLY CAST PRODUCTS

2.1 Classification of Defects in Continuously Cast Products

In Figures 1 and 2, the defects found in continuously cast products are schematically illustrated, based on a classification system devised by the International Iron and Steel Institute.4) Figure 1 illustrates surface defects, and Figure 2 illustrates internal defects.

2.2. Influence of Composition on Defects in Continuously Cast Products

2.2.1 General

Of the many types of defect in continuously cast product shown in Figures 1 and 2, only transverse surface cracks are known to be strongly influenced by the microalloying elements V, Nb and Ti. Other elements also influence transverse cracking, and their influence is discussed further in section 2.2.2.

Some of the other types of surface defect, such as longitudinal surface cracks, are influenced by composition, particularly C (0.07-0.18% being prone to longitudinal cracking) S, P, and Mn/S ratio,5,6) increased P and S, and decreased Mn/S ratio leading to increased cracking.

Internal crack formation is also influenced by composition, and again C, S and P are particularly important.7)

2.2.2 Transverse Surface Cracks

Nb Steels

There are numerous reports in the literature stating that Nb additions promote the formation of transverse cracks in continuously cast slab.8-13) The amount of Nb required for transverse crack formation appears to be very low, and cracking has been reported to increase sharply for Nb additions of as little as 0.01%,10) Figure 3. Most authors have reported that for Nb containing steels, increasing Al contents also produced increased cracking,9-12) Figure 4. However, other authors have reported no influence of Al on transverse cracking in Nb steels.8) Figure 4 also shows that other factors as well as composition influence transverse cracking; in this case, one casting machine was performing significantly better than another for the same composition. The influence of machine variables on transverse cracking will be discussed in subsequent sections.

Increased N also promotes transverse cracking in Nb containing steels,9,14) but this can be minimised if N is kept below 0.004%.14)

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The C content has a very important influence on transverse cracking, and C contents within the range 0.10-0.17% are particularly prone to transverse cracking.14)

Hannerz9) reports increased transverse cracking in Nb steels with higher S levels. However, at very low S levels (<0.005%), transverse cracking was reported to increase in Nb containing steels.14)

For Nb steels, additions of 0.2-0.3% Cu and Ni have been reported to promote transverse cracking.14) There have been conflicting reports concerning the influence of Ca on transverse cracking in Nb steels. Ca additions have been reported to reduce transverse cracking Nb steels.14,15) However, calcium silicide has been associated with uneven oscillation marks, which in turn promotes cracking.14)

Elements reported to reduce transverse cracking in Nb containing steels include Ti,8,14) P,10) Ce and Zr.15) Ti additions of 0.02-0.04% were required to reduce transverse cracks, but 0.15% Ti was required to completely eliminate the cracks.14)

The previous results relate to conventional thick (i.e. >225mm) continuously cast slab. However, there are also reports of Nb leading to transverse cracking in thin slabs of 50mm thickness.16)

V Steels

In contrast to Nb steels, there are few references in the literature to V containing steels being associated with transverse cracking. Patrick and Ludlow14) report that at N < 0.005%, V has little effect on transverse cracking. However, at high N levels (0.02%), transverse cracking can occur in 0.15%V steels.14,15) However, when cast as thin, 50mm slab, VN steels are reported to have better surface quality than Nb steels.16)

Ti Steels

There are few reports on the effects of Ti alone on transverse cracking. Williams17) has reported that no slab scarfing was required for steels containing 0.01-0.06%Ti, suggesting that these levels of Ti did not result in any slab surface defects. However, for these steels, C contents were <0.09%, and Mn 0.3-0.8%.

Other Steels

Cryogenic steels containing 9% nickel have also been reported to be prone to transverse cracking.15)

2.3 Description of Transverse Cracks

The general features of transverse surface cracks are illustrated in Figure 1. Examples of transverse cracks in slabs are shown in Figure 5. Transverse cracks may be formed on the broad face, narrow face, or corner of continuously cast slab, but are not always apparent to visual inspection unless the slab surface is dressed. They are usually associated with the depression of oscillation marks, and are predominantly found on the top slab surface. The cracks themselves can be several 10s of mm in length in extreme cases, and generally follow austenite grain boundaries. The cracks are partially oxidised, but there is little de-carbonisation, and little oxidation of the inner end of the crack.8)

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Examination of fracture surfaces of transverse cracks has indicated intergranular fracture surfaces, with ductile dimples initiating at a variety of particle types, but predominantly MnS and AlN.15) Hater8) also reported AlN on fracture surfaces.

There have also been reports of examinations of what are believed to be the earliest stages of the initiation of transverse cracks. Examination of break-out shells has indicated small cracks below oscillation marks formed in the mould section.12) There have been other reports of fine subsurface cracks prior to straightening,8) and observations of internal cracks formed in segregated regions immediately below oscillation marks.13)

2.4 Summary

Of the many types of defect encountered in continuously cast products, only transverse surface cracks are strongly influence by the presence of microalloying elements. These cracks can form on the broad face, narrow face or at the corners of continuously cast slab, and can be many 10s of mm in length. The presence of Nb greatly promotes the formation of transverse cracks, but V at low N levels has no effect, although combinations of 0.15%V and 0.02%N have been reported to lead to transverse cracking. There are no reports of Ti alone leading to transverse cracks, and Ti additions to Nb steels can be beneficial in reducing transverse cracking.

3. FORMATION MECHANISMS FOR TRANSVERSE CRACKS

3.1 Introduction

For crack formation to occur there must be an applied stress combined with the inability of the material to support this stress. To understand the various crack formation processes in continuous casting therefore requires a knowledge of the sources of stress, and also high temperature properties of the material, particularly ductility. Also it should be noted that crack formation need not proceed uniformly; there may be distinct crack initiation and crack propagation phases.

3.2 Stresses and Strains during Continuous Casting

Stresses can arise from a large number of different causes during continuous casting, and the subject has been reviewed by Lankford.18) Stresses may arise due to transformation effects, thermal effects (variable heat transfer within the mould, temperature gradients within slabs, effects of cooling water sprays, contact with rollers, etc.), friction between strand and mould, bulging of the strand caused by ferrostatic pressure, mechanical effects due to misalignment of the casting machine, and straightening strains.

The observation of numerous, large transverse cracks in the final straightened slab, together with the fact that these are often most numerous on the top surface of the slab (i.e. the surface which is in tension during straightening) suggests that there is much crack propagation induced by the stresses experienced during the straightening process.

3.3 High Temperature Ductility

Cracking is much more likely to occur if regions of low ductility are present. Several techniques are available to study hot ductility under the conditions relevant to continuous casting, and these are discussed in Appendix 1. It is possible to identify 4 distinct regions of low ductility under test conditions relevant to continuous casting. These four regions are illustrated on a schematic hot ductility curve, Figure 6:

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Region I - Embrittlement by Incipient Melting

Region IIa - Embrittlement by Second Phase Particles - (Mn,Fe)S

Region IIb- Embrittlement by Second Phase Particles - Nb(CN), AlN, V(CN)

Region III - Embrittlement by Transformation

Region I occurs at high temperatures, typically 20-50°C below the mean solidus temperature. Fracture surfaces are characterised by inter-dendritic failure and the presence of particles such as MnS. This region of low ductility is associated with incipient melting at inter-dendritic and grain boundaries, and is important in the formation of many types of defect in continuously cast products, such as longitudinal surface cracking. The segregation of elements such as S to inter-dendritic regions during solidification is important to this type of failure.

This region of low ductility may be responsible for the initiation phase of transverse surface cracks, as small subsurface cracks have been observed associated with oscillation marks.12,19) The oscillation marks themselves are regions in which high degrees of segregation of elements such as S, P and Mn19) can occur. Heat transfer to the mould in the vicinity of the oscillation mark is also reduced, which will tend to keep temperatures high, and within the brittle zone.

Region II occurs over the approximate temperature range 1200-900°C, depending on composition and test conditions, and fracture surfaces are typically along austenite grain boundaries, and sometimes show the presence of second phase particles, with ductile dimples around these second phase particles. These low ductility regions are associated with precipitates - (Mn,Fe)S for Region IIa and Nb(CN), V(CN), Ti(CN) and AlN for Region IIb.

The distinction between regions IIa and IIb is determined by the stability of the different particle types. Type IIa low ductility is only apparent at quite high strain rates; at lower strain rates, or when there is an extended hold prior to testing, ductility is good.18,20,21) On the other hand, Type IIb ductility loss is worse as strain rate decreases, Figure7. Type IIa ductility loss is strongly dependent on composition, particularly Mn/S ratio Figure 8. It has also been suggested that IIa ductility loss is due to the precipitation of liquid FeS particles, and reduction of grain boundary decohesion due to S segregation.18) Transverse cracking is usually associated with high strength microalloyed steels, with high Mn contents, and therefore high Mn/S ratios. The strain rates during the processing of continuously cast slabs are also too slow for Type IIa ductility loss to occur, and this suggests that type IIa ductility loss is not responsible for transverse crack formation.

Type IIb ductility loss is initiated by austenite grain boundary sliding, which encourages crack formation at grain boundaries,22) and the presence of second phase particles such as Nb(CN), V(CN) or AlN. These particles have two major roles; they can delay the onset of recrystallisation, and they can reduce the strain required for fracture.

The high temperature end of this ductility trough is believed to be associated with the onset of recrystallisation.23) If recrystallisation can occur prior to failure, any developing grain boundary cracks become isolated, and further propagation is not possible. It is well known that the microlloying elements Nb and V can delay recrystallisation, either in solution or as precipitates, and this retardation of recrystallisation is believed to be responsible for

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extending the Type IIb ductility trough to higher temperatures. However, in this respect, V is much less effective than Nb in delaying recrystallistion.

The presence of microalloy precipitates can also reduce the strain to fracture by a number of possible mechanisms: precipitate free zones are often observed adjacent to austenite grain boundaries, and this may lead to strain concentration at the grain boundary; the particles (or groups of particles) at the grain boundaries may act as crack initiation sites; or general matrix precipitation can lead to an increase in strength, and an overall reduction in ductility.1-3) The proposed mechanism for low ductility failures in the presence of Nb and V carbonitrides is illustrated in Figure 9.

Region III occurs over the approximate temperature range 900-600°C, depending on composition, and if Type II low ductility is present, these two ductility troughs can merge together. Fracture surfaces are characterised by intergranular failures, and the facets of the individual grains are often associated with void formation around second phase particles. It is believed that this region of low ductility is associated with the austenite to ferrite transformation. On cooling below the transformation temperature, ferrite formation commences at austenite grain boundaries, leading to the formation of films of ferrite around the austenite grains. At temperatures within the transformation range, ferrite is softer than austenite (Figure10) and so when deformation commences, strain is concentrated within the ferrite at grain boundaries, and the processes of ductile failure, i.e. void nucleation at second phase particles, and the growth of these voids, continues within the ferrite film.10,22,24) Thus on a microscopic scale, fracture can be described as ductile, but overall the failure is brittle. The mechanism for this type of fracture is illustrated in Figure 9.

The high temperature end of the ductility trough is associated with the start of transformation, and is thus determined by composition and processing conditions. There appears to be a good relationship between the temperature at ductility starts to fall and the Ar3 temperature, the transformation temperature measured during cooling Figure 11.11) It has also been suggested that the temperature at which ductility starts to fall is very close to the equilibrium transformation temperature Ae3, rather than the Ar3, as the deformation process accelerates the transformation kinetics.25)

Ductility recovers at lower temperatures because the volume fraction of ferrite is higher, and the strain distribution between austenite and ferrite becomes more uniform. At lower temperatures, the strength differential between austenite and ferrite is also less, which will again contribute to a more uniform distribution of strain between austenite and ferrite. For ductility to recover completely, it appears that approximately 50% of the austenite must have transformed to ferrite, Figure 12.11)

Microalloying elements can influence the position of this type of ductility trough through their influence on transformation temperature. For example, the presence of Nb in solution prior to transformation is known to reduce transformation temperatures, and has also been shown to reduce the temperature at which the type III ductility trough occurs, Figure. 13.1) In this respect V also has a lesser effect, as quite large additions of V are required to depress the transformation temperature significantly. Microalloys can also deepen this type of ductility trough when they are present in the form of precipitates.1,26) The precipitates may act as nucleation sites for voids within the thin ferrite films, or reduce the ductility of the thin ferrite films by retarding recovery processes.

3.4 Formation Mechanism of Transverse Cracks

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It has been suggested that the earliest stages of transverse crack formation occur in the mould, and are associated with segregation in the vicinity of oscillation marks.13) The low melting point of these regions, coupled with higher temperatures due to reduced heat transfer to the mould, lead to hot tearing. Further evidence of the importance of events in the mould to transverse cracking is suggested by the strong effect of carbon on transverse cracking. When the C content is such that some peritectic solidification can occur, transverse cracking increases and it has been suggested that this is due to transformation strains during solidification. There is no evidence to show that microalloying elements influence this stage of transverse crack formation. There have been reports suggesting that Nb additions refine the as-cast grain size, which should help to reduce transverse cracking, but these beneficial effects must be overshadowed by the detrimental effect of Nb on hot ductility.27)

Although the early stages of transverse crack formation may be in the mould, there is evidence to suggest that these defects become larger and more numerous under the application of stresses from various sources below the mould, particularly those encountered during slab straightening.12) When these stresses occur in the temperature range over which ductility is poor, transverse cracking is severe, and as hot ductility is strongly influenced by microalloys, this is the proposed mechanism by which microalloying elements effect transverse cracking. As well as having a role in the nucleation of transverse cracks, oscillation marks would also tend to favour the propagation of cracks, in that grain size may often be coarse beneath the oscillation mark, and the notch like geometry will also tend to concentrate stresses.

3. 5 Summary

In summary, there is evidence to suggest that transverse cracks initiate at high temperatures in the mould, and this initiation is associated with oscillation marks. However, subsequent propagation of these cracks continues at lower temperatures is a result of the application of further strains, particularly during slab straightening. When stresses are applied in the regions of low temperature ductility which occur due to the precipitation of microalloy carbides and nitrides, and the presence of the austenite to ferrite transformation, severe transverse cracking can occur.

4. THE INFLUENCE OF MICROALLOYING ELEMENTS ON HOT DUCTILITY

4.1 Introduction

In previous sections, the role of microalloying elements, and particularly Nb, in promoting transverse cracks was described, and the importance of high temperature properties, and particularly low levels of ductility, in leading to crack formation, was highlighted. In this section, the effect of the different microalloying elements on hot ductility will be compared, and the mechanisms by which they influence hot ductility will be discussed in more detail.

4.2 Composition Effects

V Steels

There have been several studies of the influence of V on hot ductility using hot tensile tests,9,

11, 14, 15, 26, 28- 31) and typical results are illustrated in Figure14. The results from all the various studies are consistent in that V additions of up to 0.1% at low N contents (<0.005%) have only a very slight detrimental effect on hot ductility by broadening the ductility trough. At higher N levels, the effect of V additions becomes more marked, and the ductility trough

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becomes deeper and broader. In fact, a good relationship can be constructed between the product VxN, and the depth and breadth of the ductility trough, Figure15. It should also be noted that in this example it is only at the highest VxN product, 0.1%Vx0.01%N, that hot ductility approaches that of a 0.028%Nb steel.

In other reports where direct comparisons between the hot ductility of V steels and Nb steels have been made under the same test conditions,11,29,30) the ductility of the V containing steels is superior to that of the Nb steels. For example, in the work of Fu,30) a steel containing 0.16%V and 0.011%N had superior hot ductility to a 0.039%Nb steel. The differences in hot ductility behaviour were attributed to differences in precipitation, in that the V steel exhibited little VN precipitation, whilst the Nb steel showed copious precipitation of NbCN, as a result of the different solubilities of the two precipitate types. Mintz28) has also shown that VN precipitates tend to be coarser than NbCN precipitates under processing conditions similar to those found in continuously cast slab, and hence less detrimental to hot ductility. However, when comparing V-Ti steels and Nb-Ti steels, it has been found that both steel types had similar hot ductility.32)

In a study using a hot bend test to simulate thin slab casting,26) additions of 0.1%V were found to have no effect on hot ductility for an N content of 0.007%. As the N level was increased to 0.02%, ductility did decrease, but not to the extent observed in a 0.04%Nb steel, Figure 16.

Nb steels

There have been a very large number of studies of the effect of Nb on hot ductility, and the activity in this field is probably related to the perceived detrimental effect of Nb on slab surface quality.9-11, 23, 24, 29, 30, 33-37) These results can be summarised by saying that Nb additions deepen and broaden the ductility trough to extend to higher temperatures. Nb additions of as little as 0.017% had a detrimental effect, and ductility continues to deteriorate up to at least 0.074%. Typical results are shown in Figure 17.

Al additions to Nb containing steels deepened and broadened the ductility trough,10,15,36) as did increasing N contents.24,29) When expressed in terms of AlxN, the combination of increased Al and/or N was also detrimental to the hot ductility of Nb containing steels.11)

There are conflicting reports as to the influence of P, with some workers reporting a slight beneficial effect of increased P levels,10,29) whilst others report no influence.24)

When examining the influence of Ti additions on the hot ductility of Nb steels, care must be taken to ensure that the thermal cycle is appropriate to continuous casting conditions. It is common practice when performing hot ductility tests to carry out a solution treatment at a high temperature, prior to cooling down to and testing at a lower temperature (See Appendix 1). Whilst this is acceptable for many steels, for steels containing Ti, it can lead to the formation of a fine austenite grain size, due to the grain boundary pinning effects of TiN precipitates which are stable to high temperatures. Thus there are several reports of the apparent benefits of Ti additions to hot ductility, but when the finer grain sizes of these steels are taken into account, the benefits are not as apparent,29,30) Figure 18. Thus for simulating the continuous cast condition for Ti containing steels, it is particularly important that the test pieces are melted in-situ, as described in Appendix 1. When only tests from samples melted in-situ are considered, Ti additions to Nb steels have only a very small beneficial effect on hot ductility, or even a detrimental effect, and it is believed that this is due to the influence of

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Ti:N ratio in controlling precipitate size.3) If compositions are carefully chosen to ensure that precipitate size is maximised, then good ductility can be achieved in Nb-Ti steels. The picture is complicated further still be cooling rate effects; at relatively slow cooling rate of 25°C/min, Ti can give a large improvement the hot ductility of Nb steels, again due to the formation of coarse precipitates.38)

As with Ti, the effect of S on the hot ductility of Nb steels also depends upon the thermal cycle used in the test. For steels reheated to a solution temperature prior to testing, S has little effect on hot ductility.39) However, for in-situ melted test pieces increasing S levels have a detrimental effect on hot ductility, as more S is taken into solution to precipitate on grain boundaries.40)

There are several reports that V additions to Nb steels improve hot ductility, compared with a steel containing only Nb, and this may be due to the formation of coarser (V,Nb)(C,N) precipitates in this type of steels.26,41)

Ti Steels

As discussed in the previous section, when examining the influence of Ti on hot ductility, it is important to consider the austenite grain size. In some reports, the apparent benefit to hot ductility of Ti additions is due to a refinement of grain size. It is most appropriate to evaluate the influence of Ti on hot ductility using samples melted in-situ, as this technique produces approximately similar grain sizes for Ti and Ti free steels.

There are relatively few reports looking at the influence of Ti additions to C-Mn-Al steels after in-situ melting,38,42) and relationship between Ti and hot ductility appears to be complex. In situations where large TiN precipitates can form, such as at slow cooling rates or high values of TixN, hot ductility may be slightly improved by Ti additions, but for conditions which generate large volume fractions of fine TiN particles, such as a stoichiometric ratio of Ti:N in low N steels, hot ductility can deteriorate with Ti additions, Figure 19.

4.3 The Influence of Microalloy Precipitation on Hot Ductility

From the discussions in section 3.3 on the occurrence of ductility troughs, and the previous section on the influence of composition on these ductility troughs, it is clear that the precipitation of microalloy carbides and nitrides has a crucial role to play in determining the depth, position and width of the ductility trough, through their influence on dynamic recrystallisation, strain to fracture and transformation. The influence of precipitates on dynamic recrystallisation is dependent on their size and volume fraction, large volume fractions of fine particles having the greatest effect on delaying recrystallisation. Similarly, large volume fractions of fine precipitates are likely to increase strength, and hence reduce ductility.

In Figure 20, equilibrium volume fractions of carbonitride precipitates have been calculated for the steels described in references 28 and 41, and compared with reduction of area values at 850°C in a hot tensile test. Testing was carried out under the same conditions for all the steels, and hence the results are directly comparable. There is a decrease in ductility as the equilibrium volume fraction of precipitates increases, but at a given volume fraction of precipitates, ductility is much lower for Nb steels in comparison with V and V-Nb steels. However, it is difficult to measure experimentally particle volume fractions, and it is unclear whether equilibrium precipitate volume fractions are achieved in hot tensile tests, or indeed under continuous casting conditions.

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The imposition of strain, either during a tensile test or during continuous casting, rapidly increases the rate of precipitation, and this type of precipitation occurring simultaneously with deformation is known as dynamic precipitation. Studies of dynamic precipitation kinetics in Nb and V steels43) indicate typical “C” curve behaviour (Figure 21) with the time for completion of precipitation being dependent on temperature. For a steel containing 0.035%Nb, and a steel containing 0.11%V and 0.006%N, precipitation kinetics were similar at 875°C, but the time required for completion of precipitation was approximately 10 mins. The time for the tensile test under the conditions observed in references 28 and 41 was 2-4 minutes, and this suggests that precipitation was incomplete during the time taken for the test, and that the equilibrium volume fractions were not achieved. In recent work, Banks44) has attempted to relate predicted dynamic precipitation kinetics to the hot ductility curve of microalloyed steels with some success. Figure 22 shows a good relationship for Nb steels between Tn, the predicted temperature at which the rate of dynamic precipitation is a maximum, and the temperature for a reduction of area value of 50%. However, the detailed agreement between precipitation model and experimental results in this work does depend on factors such as the hold time prior to testing in the hot tensile test, and the total level of interstitial elements.

The temperature at which ductility begins to recover also appears to be influenced by the volume fraction of precipitates. Figure 23 shows how the calculated equilibrium precipitate solution temperature varies with the temperature at which reduction in area recovers to 75%, for the same data that was used to construct Figure 20. There is a general trend for steels with higher solubility temperatures to have higher ductility recovery temperatures, although the solution temperature is always higher than the ductility recovery temperature. The work of Banks44) shown in Figure 22 shows a similar trend for Nb steels. As Nb is less soluble in austenite than V, this may explain the tendency for the Nb steels to have wider ductility troughs.

As well as precipitate volume fraction, there is clear evidence that the size of Nb(CN) precipitates influence the likelihood of transverse crack formation, and it has been shown that finer precipitates produce increased levels of cracking,10) Figure 24. Similarly, there is evidence that in hot ductility tests, that finer precipitates result in lower ductility,3) Figure 25. There are several observations that Nb(CN) precipitates observed in hot tensile tests tend to be finer than V(CN) precipitates.26,41) The reason for this may lie in the greater amounts of V in solution, which will promote particle coarsening, and also the higher diffusion coefficient of V in austenite compared with Nb. Thus if it is assumed that precipitate volume fractions for Nb and V steels are similar, the results in Figure 20 may be explained by differences in precipitate sizes, with the V steels having coarser precipitates.

4.4 Summary

Hot ductility is strongly influenced by composition, and the addition of Nb is especially detrimental to hot ductility, extending the ductility trough to higher temperatures, and deepening the ductility trough. The effect of V is much less severe and at N levels of 0.005%, V additions of 0.1% can be made with very little detrimental effect on hot ductility. As N is increased in V steels ductility deteriorates, but even at 0.11%V and 0.01%N, the ductility of a 0.028%Nb steel is still worse. Also, V additions to an Nb containing steel appear to slightly improve hot ductility. The differing effects of V and Nb may be explained by the generally coarser precipitates observed in V containing steels. The effects of Ti are complex, and depend on the Ti:N ratio.

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5. STRATEGIES FOR DEALING WITH TRANSVERSE CRACKS

5.1 Introduction

From the previous sections, it is apparent that that if stresses and strains are introduced during the continuous casting process over certain critical temperature ranges for which ductility is low, transverse surface cracks can occur in continuously cast products. There are many steps, which can be taken to minimise the likelihood that these cracks will form, but it may not be possible to completely eliminate them, in which case some form of slab repair operation is required. The following sections discuss methods to minimise cracking, and the implications of slab repair prior to final rolling.

5.2 Techniques for Crack Minimisation

5.2.1 Control of Composition

It is evident from the above discussion that composition, and particularly the use of microalloying elements, can strongly influence transverse cracking by their influence on hot ductility. This suggests that to minimise cracking a steel composition should be chosen which maximises hot ductility bearing in mind the final product requirements. The following guidelines should help to maximise hot ductility and minimise transverse cracking:

Choose C and alloy additions to avoid peritectic solidification, and particularly avoid 0.1-0.13%C

Minimise Nb

Use V or V/N combinations to replace Nb

Minimise Al

Minimise N

Make V additions to Nb steels

Consider Ti additions

5.2.2 Machine Operation

Mould Heat Transfer

Thermal stresses and surface structure can be influenced by heat transfer in the mould. It is important to consider the type of mould powder used, particularly viscosity, and ensure good and uniform fluxing.

Mould Oscillation

The importance of oscillation marks to transverse crack formation has already been mentioned, and the depth of these oscillation marks can be reduced by increasing the mould oscillation frequency and/or decreasing the stroke to reduce the heal time. Increasing the oscillation frequency has been shown to reduce the incidence of transverse cracking.12) Deep, irregular oscillation marks may also be formed due to poor mould level control and other factors.

Secondary Cooling

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The secondary cooling strategy is very important to minimise transverse cracking. In the previous sections, it has been shown that there is a wide ductility trough associated with microalloyed steels, and if slab straightening is carried out within this ductility trough, transverse cracking can result. If slab straightening is carried out at temperatures either above or below this temperature range, cracking should be minimised. Both these different cooling strategies (“soft” cooling and “hard” cooling) have been used on various machines around the world, with some success in reducing cracking.11,12) When a “soft” cooling strategy is used, it is important to keep the entire cross section of the slab above the critical temperature, including the slab corners, which are typically colder than the broad face. This has encouraged the installation of devices to maintain high temperatures in the slab corner region in some plants. A steep temperature gradient through the slab thickness is also desirable using this cooling strategy, to minimise the penetration of surface cracks which may form in cold spots.

For “hard” cooling strategies, it is important to maintain all cooling nozzles; blocked cooling nozzles may lead to localised regions of the slab having temperatures within the critical range. A “hard” cooling strategy may also lead to subsurface crack formation:11) the distance between these cracks and the slab surface must be such that they are not exposed during subsequent reheating operations. However, “hard” cooling practices may increase thermal stresses.

It should be noted that the use of these cooling strategies requires a knowledge of the temperature range over which low ductility exists, and this temperature range may not necessarily correspond to those obtained in a hot tensile test, as discussed in Appendix 1.

A variant on the use of different cooling strategies which has been used to minimise crack formation during the rolling of hot charged slabs is the use of slab quenching.45) This technique rapidly chills the slab surface layers below the transformation temperature, leading to the development of a fine grain structure at the surface. This fine grain structure then restricts the formation and propagation of cracks, which may have formed otherwise during the subsequent rolling process.

Non-uniform secondary cooling can promote thermal stresses, and hence lead to cracking. This requires good nozzle design and maintenance, and preferably the use of air-mist cooling.

Mechanical Stresses

Mechanical stresses can be introduced by poor alignment of the components of the machine, and from many other sources, but of most relevance to transverse cracking is the straightening operation. Straightening may be carried out over a single point, or by multi-point straightening. The effect of these two types of straightening operation on transverse cracking are not clear, but there are reasons to suppose that multi-point bending will not improve hot ductility: strain rate will be reduced, which will reduce hot ductility; more time will be allowed for dynamic precipitation; and at least for Nb containing steels, stress relaxation between each bending point seems to be minimal.46)

5.3 Repair of Slab/Final Product

If slab cracking cannot be avoided by any of the above methods, the only options to avoid losses in the final product are to repair the continuously cast product prior to rolling. In the most extreme case, slab repair could involve machine scarfing of top and bottom broad faces, removing several mm from each face, together with scarfing of the narrow faces. After

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machine scarfing, it may be necessary to inspect the slabs, and remove any remaining defects by hand scarfing.

These slab repair operations are expensive and time consuming. A recent example from a European plant suggested that machine scarfing of top and bottom broad faces, followed by inspection and hand scarfing of any remaining defects added a cost of $600,000 to a production volume of 100,000 tonnes. This figure will of course vary from plant to plant, but is still likely to be a considerable sum. Nevertheless, this additional cost is preferable to having to reject the final product. In the above example, prior to the introduction of scarfing, costs due to a 10% rejection level in the final product were approximately $1,600,000.

6. CONCLUSIONS

1. The only type of defect in continuously cast products reported to be influenced by microalloying elements are transverse surface cracks.

2. Nb additions are reported to have a strong influence in promoting transverse cracks, but V additions at low levels of N do not lead to cracking. Ti additions in themselves do not appear to produce transverse cracks, and Ti additions to Nb steels reduce transverse cracking.

3. Transverse cracks can occur at the slab broad face, narrow face or corner, and can be several mm in depth. They follow austenite grain boundaries, and the fracture surfaces are covered with particles such as MnS and AlN.

4. Transverse cracks are believed to initiate in the mould, but propagate during the straightening process.

5. Several regions of low ductility exist when steels are tested at high temperatures under conditions simulating those experienced during continuous casting. Nb, and to a much lesser extent V, extend the depth and width of this ductility trough. The ductility trough is associated with the precipitation of microalloy carbonitrides and the transformation from austenite to ferrite.

6. To minimise transverse cracking, careful attention should be given to the choice of steel composition: Nb should be minimised, and consideration given to replacing Nb with V and N.

7. Machine operating conditions should be optimised, and in particular the straightening temperature should be chosen to be outside the region of low ductility.

REFERENCES

1. Y. Maehara et. al., “Surface Cracking Mechanism of Continuously Cast Low Carbon Low Alloy Steel Slabs” Mat. Sci. and Technol., 1990, 6, 793-806.

2. B. Mintz: “Hot Ductility of Steels and It’s Relationship to the Problem of Transverse Cracking During continuous Casting”, Int. Mat. Rev., 1991, 36, 187-217.

3. B. Mintz: “The Influence of Composition on the Hot Ductility of Steels and to the Problem of Transverse Cracking” ISIJ Int. 1999, 39, 833-855.

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4. “Continuous Casting of Steel 1985 - A Second Study” , International Iron and Steel Institute, 1986.

5. J. K. Brimacombe and K. Sorimachi: “Crack Formation in the Continuous Casting of Steel”, Metall. Trans., 1977, 8B, 489-505.

6. W. R. Irving and A. Perkins: “Basic Parameters Affecting the Quality of Continuously cast Slabs”, Int. Conf. on Continuous Casting, Biaritz, 1976, 36.

7. H. Vom Ende and G. Vogt: JISI, 1972, 210, 889-894. 8. M. Hater et. al., “Results From a Curved Mould Continuous Casting Machine Making

Pipe and Plate Steel” , “Open Hearth Proceedings”, AIME, 1973, 202-217. 9. N. E. Hannerz: “Critical Hot Plasticity and Transverse Cracking in Continuous Slab

Casting with Particular reference to Composition” Trans ISIJ, 1985, 25, 149158. 10. B. Mintz and J. M. Arrowsmith, “Hot-ductility behaviour of C-Mn-Nb-Al Steels and its

relationship to crack Propagation During the straightening of Continuously Cast Strand”, Met. Technol., 1979, 6, 24-32.

11. N. Bannenberg et. al., “Procedures for Successful Continuos Casting of Steel Microalloyed with Nb, V Ti and N”, Microalloying 95, 83-94.

12. N. A. McPherson and R.E. Mercer: “Continuous Casting of Slabs at BSC Ravenscraig Works”, Ironmaking and Steelmaking, 1980, 167-179.

13. S. Harada et. al., “A formation Mechanism of transverse Cracks on CC Slab Surface”, ISIJ Int., 1990, 30, 310-316.

14. B. Patrick and V. Ludlow: “Development of Casting Practices to Minimise Transverse Cracking in Microalloyed Steels” Rev. Metall. 1994, 91, 1081-1089.

15. T. H. Coleman and J. R. Wilcox: “Transverse Cracking in Continuously Cast HSLA Slabs-Influence of Composition”, Mat. Sci. and Technol., 1985, 1, 80-83.

16. P. J. Lubensky et. al., “High Strength Steel Processing Via Direct Charging Using Thin Slab Technology”, Microalloying 95, 225-233.

17. J. G. Williams: ”Titanium Microalloyed Hot rolled Strip Steels - Production, Properties and Applications” in “HSLA Steels-Technology and Applications” Philadelphia, 1984, ASM 261-275.

18. W. T. Lankford, “Considerations of Strength and Ductility in the Continuous-Casting Process”, Metall. Trans., 1972, 3, 1331-1357.

19. S. Tanaka et. al., “Formation Mechanism of Surface Cracks Along The Oscillation Mark”, Trans. ISIJ, 1981, B-350

20. G. A. Wilber et. al., “The effects of Thermal History and Composition on the Hot Ductility of Low Carbon Steels”, Metall., Trans., 1975, 6A, 1727-1735.

21. Y. Yasumoto et. al., “Effects of Sulphur on the Hot Ductility of Low Carbon Steel Austenite”, Mat. Sci. and Technol., 1985, 1, 111-116.

22. H. G. Suzuki et. al.: ”Embrittlement of Steels Occurring in the Temperature Range from 1000 to 600°C”, Trans ISIJ, 1984, 24, 169-177.

23. J. R. Wilcox and R. W. K. Honeycombe:” Hot Ductility of Nb and Al Microalloyed Steels Following High Temperature Solution Treatment”, Met. Technol., 1984, 11, 217-225.

24. C. Ouchi and K. Matsumoto “Hot Ductility in Nb-bearing High Strength Low-alloy Steels”, Trans. ISIJ, 1982, 22, 181-189.

25. B. Mintz, R. Abushosa and A. Cowley: “An Analysis of the Hot Ductility Curve in Simplae C-Mn Steels” to be published in Mat. Sci., and Technol.

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26. D. N. Crowther, M. J. W. Green and P. S. Mitchell “The Influence of Composition on the Hot cracking Susceptibility During Casting of microalloyed Steels processed to Simulate Thin Slab Casting Conditions”, “Microalloying in Steels”, Materials Science Forum, Vols. 284-286, .1998) 469-476.

27. H. Zhang et. al.: “Effect of Niobium on Continuous Casting Solidification Structures of HSLA Steels” in “HSLA Steels Metallurgy and Applications”, Beijing 1985, 445-453.

28. B. Mintz and R. Abushosa, “Influence of Vanadium on Hot Ductility of Steel” Ironmaking and Steelmaking, 1993, 20, 445-452.

29. B. Mintz and J. M. Arrowsmith, “Influence of Microalloying Additions on Hot Ductility of Steels”, in “Hot Working and Forming Processes”, The Metals Society, 1980, 99-103.

30. J. Y. Fu et. al., “Hot Ductility of Continuously Cast Microalloyed Steels” in “Processing , Microstructure and Properties of High Strength, Low Ally Steels”, Pittsburgh, 1987, 27-38.

31. E. Schmidtmann and M. Merz: “Effect of Cooling Conditions and Strain Rate on High Temperature Properties of Structural Steels in Continuous Casting”, Steel Research, 1987, 58 191-196.

32. L. P. Karjalainen, H. Kinnunen and D. Porter: “Hot Ductility of Certain Microalloyed Steels Under Simulated Continuous Casting Conditions” in: “Microalloying in Steels”, Materials Science Forum, Vols. 284-286, .(1998) 477-483.

33. G. Bernard: “Study of Susceptibility to cracking of Continuously cast Steels using Hot Ductility Tests”, Rev. Metall., 1978, 75, 467-480.

34. R. Abushosa, R. Vipond and B. Mintz, “Influence of Sulphur and Niobium on Hot Ductility of As Cast Steels”, Mat. Sci and Technol., 1991, 7, 1101-1107.

35. L. Zhen, Z. Hongtao, W. Baorong: “Effect of Niobium on Hot Ductility of Low C-Mn-Steel Under continuous Casting Simulation Conditions” Steel Research, 1990, 61, 620-623.

36. P. Sricharoenchai, C. Nagasaki and J. Kihara: “Hot Ductility of High Purity Steels Containing Niobium” ISIJ Int., 1992, 32, 1102-1109.

37. H. G. Suzuki, S. Nishimura and S. Yamaguchi “Characteristics of Hot Ductility in Steels Subjected to Melting and Solidification”, Trans. ISIJ, 1982, 22, 48-56.

38. R. Abushosa, R. Vipond and B. Mintz, “Influence of Titanium on the Hot Ductility of As-Cast Steels”, Mat. Sci and Technol., 1991, 7, 613-621.

39. B. Mintz, J. R. Wilcox and D. N. Crowther: “Hot Ductility of Directly cast Steels” Mat. Sci. and Technol.,1986, 2, 589-594.

40. B. Mintz and R. Abushosa: Mat. Sci. and Technol., 1992, 8, 171. 41. B. Mintz and R. Abushosa: “The Hot Ductility of V, Nb/V and Nb Containing Steels” in:

“Microalloying in Steels”, Materials Science Forum, Vols. 284-286, .(1998) 461-468. 42. O. Commineli, R. Abushosa and B. Mintz: “Influence of Titanium and Nitrogen on Hot

Ductility of C-Mn-Nb-Al Steels” Mat. Sci. and Technol., 1999, 15, 1058-1068. 43. M. G. Akben, I. Weiss and J. J. Jonas: “Dynamic Precipitation and Solute Hardening in a

V Microalloyed Steel and Two Nb Steels Containing High Levels of Mn” Acta Metall., 1981, 28, 111-121.

44. K. M. Banks, A. P. Bentley and A. Koursaris: “A Precipitation Model for Predicting Hot Ductility Behaviour in Microalloyed Steels”, 42nd Mechanical Working and Steel Processing Conference, 2000, 329-340.

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45. A. Carboni, D. W. Ruzza and S. L. Feldbauer: ”Quenching for Improved Direct Hot Charging Quality” Iron and Steelmaker, 1999, 30-42

46. 45.Y. Maehara, H. Tomono and Y. Ohmori: “Stress Relaxation During Hot Deformation of Austenite”, Trans. ISIJ, 1987, 27, 499-505.

47. D. N. Crowther and B. Mintz: “Influence of Grain Size and Precipitation on Hot Ductility of Microalloyed Steels”, Mat. Sci. and Technol., 1986, 2, 1099-1105.

48. G. X. Liu and W. Dahl: “The Influence of Temperature and Strain Rate on the Stress-Strain Behaviour of “in-situ solidified” Steel During tensile Test”, Steel Research, 1989, 60, 221-229.

49. P. Deprez, J. P. Bricout and J. Oudin: “A New Tensile Test on in situ Solidified Notched Specimens: Hot Ductility Analysis of Continuous Casting Steels” , J. Mat. Proc. Technol., 1992, 32, 325-334.

50. T. Revaux et. al., “In Situ Solidified Hot Tensile Test and Hot Ductility of Some Plain Cabon Steels and Microalloyed Steels”, ISIJ International, 1994, 34, 528-535.

51. B. Mintz, J. M. Stewart and D. N. Crowther: “The Influence of Cyclic Temperature Oscillations on Precipitation and Hot Ductility of a C-Mn-Nb-Al Steel”, Trans. ISIJ, 1987, 27, 959-964.

52. G.I.S.L. Cardoso, B. Mintz and S. Yue: “Hot Ductility of Aluminium and Titanium Containing Steels With and Without Cyclic Temperature Oscillations”, Ironmaking and Steelmaking, 1995, 5, 365-377.

53. M. Suzuki et. al., “Simulation of Transverse Crack Formation on Continuously Cast Peritectic Medium Carbon Steel Slabs”, Steel Research, 1999, 70, 412-419.

54. B. Mintz, R. Abushosa and D. N. Crowther, “Influence of Small Additions of Copper and Nickel on Hot Ductility of Steels”, Mat. Sci. and Technol., 1995, 11, 474-481.

55. H. G. Suzuki, S. Nishimura and S.Yamaguchi: “Physical Simulation of The Continuous Casting of Steels”, Proc. Conf. on Physical Simulation Techniques for Welding, Hot forming and Continuous Casting, 1988, Canmet, Ottowa.

56. B. Mintz and S. Yue: “The Hot Tensile test For Assessing the Likelihood of Transverse Cracking During Continuous Casting of Steel”, 34th Mechanical Working and Steel Processing Conference, 1992, 391-398.

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APPENDIX 1

TECHNIQUES FOR ASSESSING HOT DUCTILITY AND THEIR RELEVANCE TO CONTINUOUS CASTING

A1.1. Types of Hot Ductility Test

Hot ductility is most commonly assessed using an elevated temperature tensile test. There is no standardised test, and many different procedures have been adopted by different researchers. Typically the thermal cycle used will involve heating to a solution temperature in the range 1200-1350°C to produce a coarse austenite grain size and dissolve any microalloy precipitates, followed by cooling to a test temperature at a rate to simulate that experienced at the surface of a continually cast product (typically 60°C/min), followed by straining to failure at a strain rate of 10-3 to 10-4 s-1, to simulate that experienced during the straightening of continuously cast slab.

The exact parameters used for the thermal cycle can all have an influence on hot ductility. The solution temperature may effect the austenite grain size, and it is known that coarser austenite grain sizes will reduce hot ductility.47) The cooling rate to test temperature can be important, and there are results that show that increased cooling rates can reduce hot ductility. The strain rate is also important, and reducing strain rates for microalloyed steels reduce hot ductility,22) Figure 26.

To simulate more closely the continuously cast condition, some tests melt the test piece in-situ prior to testing.22, 34, 39, 48-50) It is particularly important to use this type of test when evaluating the effects of Ti and S on hot ductility as complete dissolution will not be achieved by a solid state heat treatment.

Rather than employing continuous cooling to a test temperature, some workers have employed more complex thermal cycles in an attempt to more accurately simulate the complex temperature patterns experienced at the surface of a continuously cast product.51 52) Cooling cycles employing temperature oscillations have been shown to have a marked effect on hot ductility in some situations, by promoting precipitation of AlN.

Attempts have been made to simulate oscillation marks in tensile tests, by machining a notch into the sample prior to testing.49,53) In this way the effect of notch size and geometry on hot ductility was assessed.

The test atmosphere may also be significant in some cases. The majority of tests are carried out under vacuum or in an inert atmosphere to avoid oxidation. However, some tests have been carried out in air so as to be able to reproduce hot shortness effects due to copper.54)

Less commonly, high temperature bend tests have been used to assess hot ductility.18,26) These tests typically cast a small scale laboratory ingot, remove the ingot from the mould whilst still hot, and deform it in a three point bend test. Ductility is assessed either qualitatively or quantitatively by examining any cracks on the surface. This technique has the advantage of being able to accurately reproduce continuously cast structures, including long columnar grains and microalloy precipitation, but is a slow and expensive test. Even this test does not simulate the geometry and segregation characteristics associated with oscillation marks, which are known to be crack initiation sites for transverse cracks.

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A1.2 Relevance of Hot Ductility to Transverse Cracking

In qualitative terms, hot ductility as measured in laboratory tests appears to bear some relationship to the observations of transverse cracking; the detrimental effects of Nb, N and Al on transverse cracking are all reflected in reductions in hot ductility. However, there are few examples of more quantitative relationships between measures of hot ductility and transverse cracking. Bannenebrg11) established a relationship between the number of transverse cracks per slab and the reduction of area in a hot tensile test. Above a reduction of area value of 75%, no cracked slabs were observed. Suzuki et. al.,55) have suggested a value of 60% reduction of area to avoid slab cracking, whilst Mintz and Yue56) suggest 30-40% is more realistic. It can therefore be seen that there is considerable discrepancy between the suggested values, and it is likely that such a value can only be ascribed to specific tensile test conditions and slab assessment methods.

As well as the depth of the ductility trough, the temperature at which it occurs is also significant; if slab straightening can be carried out outside the temperature range of low ductility, then transverse cracking may be avoided. Considering the upper temperature range, the simple hot tensile test is unlikely to give a reliable indication of this temperature as dynamic recrystallisation occurs during tensile testing giving a marked increase in ductility. However, the strains encountered at the surface of continuously cast slab are only ~1-2%, which is insufficient for recrystallisation to occur. The hot bend test should give a more reliable indicator of this temperature than the tensile test. Considering the lower end of the ductility trough, recovery of ductility is associated with the completion of a certain proportion of transformation (Bannenberg11) has suggested 50%), and therefore the hot tensile test should give a more reliable indication of this temperature. However, the higher strains encountered during the tensile test may again give misleading temperatures, due to the influence of strain in transformation, and the bend test is again more likely to give accurate indications.

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1. Transverse corner cracks; 2. Longitudinal corner cracks; 3. Transverse cracks; 4. Longitudinal cracks (broad face); 5. Star cracks; 6. Deep osciuation marks; 7. Pinholes; 8. Macro inclusions.

Fig. 1 Surface defects in continuously cast products.4)

1. Internal corner cracks; 2. Side haltway cracks; 3. Centreline cracks; 4. Centreline segregation; 5. Porosity; 6. Halfway cracks; 7. Non-metallic inclusions, clusters; 8. Sub-surface ghost lines; 9. Shrinkage cavity; 10. Star cracks, diagonal cracks; 11. Pinholes; 12. Semi-macrosegregation.

Fig. 2 Internal defects in continuously cast products.4)

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Fig. 3 Influence of Nb on transverse cracking.10)

Fig. 4 Influence of soluble Al and machine on plate rejections due to slab transverse cracking in C-Mn-Nb-Al steels.12)

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Fig. 5 Transverse cracking in (a) scarfed slab showing transverse cracking and (b) partially rolled slab showing severe transverse edge.12)

Fig. 6 Schematic illustration of types of ductility trough.

a) b)

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Fig. 7 Regions of low ductility due to (a) precipitation of carbides/nitrides and (b) sulphides.1)

Fig. 8 Effect of Mn/S ratio on hot ductility.20)

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Fig. 9 Mechanisms of high temperature intergranular failure in low alloy steels.1)

Fig. 10 Dependence of maximum stress on temperature.

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Fig. 11 Influence of transformation temperature Ar3 on temperature at which ductility loss begins.11)

Fig. 12 Relationship between temperature for ductility recovery and temperature for 50% transformation.11)

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Fig. 13 Influence of Nb on transformation temperature and position of ductility trough.1)

Fig. 14 Influence of V and N on hot ductility.28)

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Fig. 15 Relationship between VxN and (a) depth of ductility trough at 850°C and (b) width of ductility trough.28)

Fig. 16 Influence of V, N and Nb on length of largest crack in a hot bend test.26)

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Fig. 17 Influence of Nb on hot ductility and strength.24)

Fig. 18 Influence of a 0.04% Ti addition to the hot ductility of a Nb steel containing 0.004%N for similar austenite grain sizes.29)

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Fig. 19 Influence of Ti on the hot ductility of as cast Al steels.38)

Fig. 20 Influence of calculated equilibrium precipitate volume fraction on reduction of area at 850°C.28,41)

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Fig. 21 Dynamic precipitation kinetics for V and Nb steels.43)

Fig. 22 Relationship between temperature for maximum precipitation rate, Tn, precipitate dissolution temperature and temperature for 50% reduction in area Nb steels.44)

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Fig. 23 Influence of precipitate equilibrium solution temperature for 75% reduction of area.28,41)

Fig. 24 influence of precipitate size and spacing on hot ductility and transverse cracking in Nb steels.2)

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Fig. 25 Influence of particle size on hot ductility for Ti and Nb-Ti steels.3)

Fig. 26 Influence of strain rate on ductility in Nb steels.22)

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The Influence of Vanadium-Microalloying on The Weldability of Steels

P. H. M. Hart

Technology Group Manager

TWI Limited, Granta Park, Great Abington, Cambridge CB1 6AL

1. INTRODUCTION

The commencement of microalloyed steels was approximately half a century ago with the initial drivers of reducing material costs and increasing strength. Although improved weldability was not amongst these very first incentives for the development of microalloyed steels the opportunities that microalloying gives for improving weldability by reducing carbon content was fairly quickly recognised and became another major driving force for their development and understanding. However, it was not until possibly the first major international conference on these steels, MA75 that a substantial public hearing of the topic occurred. In this paper the aim will be to focus attention on the current information concerning weldability aspects related to vanadium-microalloyed steels for structural, pressure vessel and linepipe applications and will consider the relevant weldability aspects of solidification cracking, reheat cracking, hydrogen cracking during fabrication, HAZ hardening, and HAZ and weld metal toughness. In recent years vanadium additions have also been applied to the Cr-Mo steels for elevated temperature use, and this too will be considered. Finally some consideration will be given to the so-called 'vanadium effect' in promoting intragranular nucleation, particularly during welding, in C:Mn microalloyed steels.

2. WELDABILITY TOPICS

2.1 Solidification Cracking

Past experience has certainly shown that in relation to arc welding the benefits that microalloying additions can bring to steel composition in terms of a reduction in carbon content, have been overwhelmingly beneficial. A systematic study of the influence of compositional variables on solidification cracking was carried out some years ago by TWI1) and the results incorporated into British and now European Standards for arc welding of ferritic steels. The work was done using submerged arc welding which usually involves considerable dilution of the parent steel and therefore allowed compositional effects of parent steels to be evaluated. A formula was devised which relates crack susceptibility, in terms of units of crack susceptibility, (UCS) to the composition of the weld metal and is: UCS = 230C+190S+75P+45Nb-12.3Si-5.4Mn-1.

The equation is valid for the following range of weld metal compositions 0.08-0.23%C, 0.010-0.050%S, 0.010-0.045%P, 0.15-0.65%Si, 0.45-1.6Mn, 0-0.07%Nb. It was concluded that the presence of up to 1%Ni, 0.5%Cr, 0.4%Mo, 0.07%V, 0.3%Cu, 0.02%Ti, 0.03%Al, 0.002%B had no significant influence on the risk of cracking. One study2) which examined the effect of vanadium by employing submerged arc welding of pipe containing 0.16%C and 1.4%Mn, showed that vanadium additions apparently reduced the risk of solidification cracking.

While solidification cracking in arc welds is nowadays uncommon, principally because of the lowered C & S levels, as a weldability problem it has reappeared in high speed beam welds,

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particularly laser welds. It would be interesting to see if the reported beneficial effect of vanadium in submerged arc welding of pipe also applies to laser welding, which is now being considered for pipe welding.

2.2. Reheat Cracking

Reheat, or stress relief, cracking is generally limited to low alloys steels containing at least two of the elements Cr, Mo, V and B. However, as the problem is limited to a few steels there has been little attempt to develop universal formulae for this type of cracking. Two examples of formulae that have been used are:

i) ∆G =Cr+3.3Mo+8.1V (3)

where ∆G<0 indicates low susceptibility and

ii) PSR = Cr+Cu+2Mo+10V+7Nb+5Ti-2 (4)

when PSR<0 indicates low susceptibility

As the cracking is intergranular, other elements, which influence temper embrittlement can also affect the risk of cracking and so the problem is not solely related to the level of alloying elements present. Certainly with respect to structural steels, the tendency to avoid postweld heat treatment, means that this is a rare and decreasing problem.

2.3. HAZ Hydrogen Cracking

2.3.1 C:Mn Steels

During the early days of development of microalloyed steels, HAZ hydrogen cracking was a significant weldability problem, not least because carbon levels were still relatively high and hydrogen levels from many consumables, either as-received, or because of poor storage and handling, were also relatively high. Both of these influencing factors, particularly in the last 10-15 years have beneficially improved, as carbon levels of steels have continued to fall and hydrogen levels in welds significantly reduced as a result of improvements in consumables. The risk of the problem occurring, is still however, related to the risk of producing hardened predominantly martensitic microstructures in the heat affected zone and material sensitivity to the problem is still generally assessed by a compositional characterising parameter generally referred to as a carbon equivalent formula. These parameters are often linked into methods for devising minimum preheat to avoid the problem. The most comprehensive of these schemes today are probably those used (i) in Britain,5,6) devised by TWI, (ii) in Japan often referred to as the CEN method,7) (iii) that incorporated within AWS D1.1 and (iv) the more recent method originating within Germany.8) These use the following compositional characterising parameters.

i. CE = C+6

Mn + 5

VMoCr ++ +15

CuNi +

ii. CEN = C+ f(C) {24Si +

6Mn +

5VMoCr

20Ni

15Cu ++

++ }

where f(C) = 0.75 + 0.25tanh[20(C-0.12)]

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iii CE = C+6

SiMn + +5

VMoCr ++ +15

CuNi +

and Pcm= C+30Si +

20Mn +

60Ni +

20Cr +

15Mo +

10V +5B

iv. CET = C+10

MoMn + + 20

CuCr + + 40Ni

Most of these show vanadium to have some small adverse effect on the risk of HAZ hydrogen cracking, although it is interesting to note that the more recent German study which applies to steels up to 0.18%V did not find it necessary to include this element in the compositional characterising parameter. The study by Hart and Harrison9) in the ‘80s, on a series of experimentally produced carbon manganese nickel vanadium molybdenum steels, found that the critical cooling time for the avoidance of HAZ cracking could be related to the steel composition by the expression:

log ∆t 800-500°C (crit) = 3.7(C+10Mo

40Ni

6V

13Mn

+++ ) – 0.31

The risk of cracking characterised in this expression is a summation of two separate factors. One is the hardenability of the steel, which describes the risk of forming hardened and susceptible microstructures, the other is the inherent susceptibility of the hardened microstructures often expressed as the critical hardness for cracking. Hart and Harrison found that the influence of vanadium was of a different sign in respect of these two factors. For hardnesses around the critical hardness level for that study (~350HV) vanadium had a decreasing effect on steel hardenability (see next section), which by itself would lead to a reduction in risk of cracking.

However this was offset by the finding that the critical hardness was decreased (ie susceptibility increased) by vanadium, as indicated by the equation:

HV= 283.3+668.1 (C+42Mn -

4V +

24Mo )

In subsequent regression analysis of a much larger database of C:Mn type steels, many microalloyed with Nb, Hart et al,10) produced a similar equation for the compositional characterising parameter:

HVcrit = 207+692(C+100Mn3 +

25Si4 +

50Mo3 +

25Cu3 -

25Cr2 +

25Al17 +

3P5 -

4V -

3Nb4 )

As discussed by Hart et al since the critical microstructures for cracking tend to be martensitic, the strong dependency of HVcrit on carbon is not surprising. Moreover they noted that there were several studies which have shown that the final contraction stress, when a weld cools out, is related to the steel composition, being lower for steels with lower transformation temperatures. Given that, the trend for HVcrit to reduce with increasing C and CE levels, is also not surprising.

In considering the cracking risk of a particular element its effect should be viewed in respect of the increase of likely alloying levels, and since much steel development is against a specific strength level, comparison should be made on alternative alloying to produce the

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same strength level. Hart and Harrison9) showed that compared to manganese, albeit in normalised steel, a given strength increase could be achieved for a smaller increased risk of cracking by using V rather than Mn alloying.

2.3.2 Low Alloy Steels

Particularly in the Cr-Mo steels for elevated temperature service, vanadium additions have been quite prominent in steel development e.g. modified 9Cr1Mo steels, so called P91, the vanadium modified 2.25Cr1MoV steels and the 3CrMoVTiB steels. Being more recent developments less work has been carried out on the effect of vanadium although for P91 type compositions Panton-Kent11) has shown that the risk of weld metal hydrogen cracking is no higher than for the straight 9Cr1Mo composition. For the 2¼Cr1Mo type steels there is data12) to suggest that V-additions may have distinct benefits as far as hydrogen embrittlement in service are concerned, and this seems to be linked to greater trapping and lower diffusivities in the V-modified steels. Indeed other work13) has indicated that the problem of disbonding at cladding interfaces can be reduced in V-modified steels. The possible consequence of a reduced hydrogen diffusivity on welding procedures for fabricating v-modified Cr-Mo steels, in particular the application of dehydrogenation heat treatments is the subject of a forthcoming joint industry project at TWI.

2.4 HAZ Hardening Behaviour

Clearly the influence of microalloying in HAZ hardening behaviour is of importance because of the role of this on the risk of HAZ hydrogen cracking as described above, although, this aspect has become of less importance because of the decreased risks of HAZ hydrogen cracking in general, also referred to above. However, HAZ hardening is also important, perhaps increasingly so, from the point of view of avoidance of environmental cracking and in particular that which is HAZ hardness related and specifically those types caused by hydrogen embrittlement either from cathodic protection or sulphide stress corrosion cracking. These forms of environmental cracking are prevented by controlling HAZ hardness to specified limits which for fully sour service will typically be 250HV max.

As well as looking at the risk of HAZ hydrogen cracking, the study by Hart and Harrison9) also developed equations for prediction of HAZ hardness levels in the range 250-450HV. As can be seen from Table 1 this shows an interesting trend for the influence of vanadium. For the higher hardness levels of 325 HV and above, which will tend to be associated with the more rapid cooling conditions and therefore shorter thermal cycles, it can be seen that vanadium reduces the hardenability, while for the softer microstructures of 300 down to 250HV the influence of vanadium was of an increasingly positive effect. This change in sign of coefficient for vanadium is not apparent from other work examining HAZ hardness, for example that of Suzuki14), Yurioka15) and Duren16) but this may be because of their lower levels of vanadium studied and the fact that they did not develop expressions for each specific hardness level.

The reason for the change of sign is not known but may be due directly from an effect of vanadium on the γ→α transformation. It may also be attributed in part to the effect of vanadium as V(C/N) in providing some pinning of austenite grain boundaries and thus reducing austenite grain growth in the high temperature HAZ. Vanadium only has a weak effect in this respect and thus the effect would be more likely to be observed in fast thermal cycles when there is least time for dissolution of these particles and subsequent grain growth. At the lower hardness levels, which in general would be achieved by longer cooling times, some positive contribution to hardness through a precipitation hardening mechanism might

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be expected and therefore provide the explanation for the increasingly positive magnitude of vanadium.

2.4 HAZ Toughness

There are relatively few studies that have specifically looked at the effect of increasing vanadium content on HAZ toughness. However, the generality of these for the as-welded condition, is that increasing vanadium up to levels of about 0.15% (mostly the upper level studied) either have no effect on the as-welded toughness, or can show useful improvements. The pattern of results is not always as this, since there have sometimes been different trends between studies of simulated HAZ and welded joints and between results produced from Charpy and CTOD testing. However, this author prefers to concentrate on information from welded joints in the strong belief that this is most likely to provide the most realistic representation of what is actually occurring.

Results reported by Mitchell et al17) showed a strong improvement in Charpy toughness of multipass welds at 2kJ/mm in 0.12% C-steels, but hardly any change in CTOD transition temperature, see Fig.1. In a study of 2-pass welds in a series of lower (0.07%) carbon steels Hart and Sage18) found that at 0.05-0.10%V there was some improvement in the grain coarsened HAZ toughness of the second pass, see Fig.2. Mitchell et al17) using their own data, and data produced by Crowther18) and Wang,19) looked at the effect of heat input on as-welded toughness of vanadium steels and this information is shown in Fig.3. It indicates some clear differences between vanadium microalloyed and niobium and niobium+vanadium microalloyed steels, in that the vanadium microalloyed steels are more tolerant to increased heat input and when toughness is studied in terms of the CTOD transition temperature, interestingly an improvement in behaviour can be observed.

The above studies are essentially where the steel composition had just been differing in respect of the vanadium content and therefore provide the clearest opportunity to study the effect of vanadium on HAZ toughness. However, interesting trends can sometimes be found in steels in which other elements are not constant. For example in the large project reported by Harrison and Hart20) they showed that when studying the sub-critically reheated grain coarsened HAZ vanadium-bearing steels produced a better 0.1mm CTOD transition temperature, for a given hardness, compared to vanadium-free and niobium-bearing steels, see Fig.4. Interestingly this pattern was observed for both as-welded and postweld heat treated (2 hours at 590°C) conditions. Li et al21) studying the intercritically reheated grain coarsened (ICGC) HAZ also found that a 0.05% V steel gave better ICGC HAZ fracture toughness than either a plain C:Mn or a C:Mn:0.3Nb steel, although at 0.10% V toughness decreased associated with an increased M-A fraction.

In assessing HAZ toughness of carbon manganese and microalloyed steels consideration must be given to the influence of nitrogen. To the author’s knowledge probably all the studies that have investigated the effect of nitrogen on HAZ toughness in C:Mn or C:Mn:Nb steels have tended to show that increasing nitrogen decreases the HAZ toughness and this has been generally ascribed to the presence of free nitrogen in the HAZ. In contrast to this, the position in C:Mn:V steels is different in that there have been fewer studies but that the majority of these have tended to show tolerance of increased nitrogen rather than an adverse effect on HAZ toughness. Hannerz22) found little or no effect on HAZ toughness of increasing nitrogen from 60 to 130ppm. Mitchell et al17) found increasing nitrogen from 80 to 170 ppm in multipass welds to also have little or no effect and this also mirrored the experience described by Lagneborg et al19) for multipass welds. However, Lagnaborg also reported on single pass welds made at the high heat input of 8.0kJ/mm when an adverse effect was observed.

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As part of the effort to reduce fabrication costs there is an increasing trend to reduce the number of occasions when it is necessary to carry out postweld heat treatment by demonstrating that adequate as-welded joint toughness can be achieved. In this respect the beneficial effects that alloying with vanadium can bring to improving as-welded toughness are significant. Nevertheless, there will be some circumstances where postweld heat treatment is necessary and the behaviour of vanadium-containing steels needs to be considered. As can be see from Fig.1 (a) the effect on Charpy toughness shows that for vanadium levels of up to about 0.06% toughness is about equal to or better than that in the as-welded condition. For higher levels of vanadium some deterioration in as-welded toughness is noted. In contrast to the Charpy behaviour, the CTOD transition temperature data for the multipass welds in Fig.1 (b) show substantial improvement as a result of postweld heat treatment, the CTOD transition temperature improving by 30-50°C. For thicker section (50mm) higher heat input (5kJ/mm) welds, Hart and Mitchell24) found for a 0.12C, 0.10V steel that the 40J Charpy transition temperature improved from –24 to –39°C, while the CTOD at –10°C showed only a small fall from 0.46 to 0.30mm. This mixed pattern of behaviour would be a summation of effects of tempering processes and precipitation hardening during postweld heat treatment. The extent of these would also depend on the condition of the prior microstructure, itself a reflection of whether the weld was made at medium or high heat input.

2.5 Weld Metal Toughness

The influence of parent steel microalloying on weld metal toughness, for arc welding, is really only of concern for those processes that dilute large amounts of the parent material into the weld metal and the principal arc welding process capable of doing this is submerged arc welding. The position for this process in respect of vanadium, and niobium, was reviewed by Dolby25) some years ago and although the situation showed the usual degree of complexity, it was apparent that vanadium additions were slightly better tolerated than niobium additions. Moreover he noted that vanadium additions were able to provide some degree of promotion of acicular ferrite type microstructures which in general can improve the toughness of weld metals.

Since that review interest has grown in the use and application of power beam processes, which in their preferred mode of use i.e. welding autogenously without filler, dilution is obviously very high. Microstructural development in power beam weld metals is different to that in most of the arc welding processes since in the absence of fluxes and active shielding gases the weld metal oxygen content is very much lower, and similar to that of the parent steel. In this respect some work by Abson and Francis-Scrutton26) is of interest since it showed in the low oxygen content prevailing in TIG welds, analogous to that in power beam welds, vanadium additions of 0.04% to the weld (by dilution from the parent steel) provided useful increases in the acicular ferrite fraction of 20-30%.

3. DISCUSSION

3.1 Vanadium and Weld Zone Microstructures

Particularly in relation to the toughness of welded joints it would be apparent from the above that there are several recorded examples of where vanadium additions to C:Mn steels can either produce some useful improvement or the hardening tendency can be tolerated. Probably the most common theme which runs through the various explanations that have been put forward for these observations relate to changes in the type of transformed microstructure that are produced by vanadium additions. These are changes from a

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microstructure dominated by grain boundary nucleated products as in C:Mn and particularly C:Mn:Nb steels, to microstructures with a significant degree of intragranular nucleated products, often referred to as intragranular ferrite typical in V-microalloyed steels. The advantage of this change for toughness is that it tends to reduce the effective colony size of bainite or ferrite with aligned second phase, thereby reducing the effective cleavage crack size and improving the ductile to brittle transition temperature. This “vanadium-effect” on transformed microstructure appears to be able to be active in both HAZ and weld metal microstructure development. Indeed the author has seen it in several instances where it can be contrasted with a coarse type of microstructure that typically develops in its absence. Figure 5 shows the effect in a high heat input (5kJ/mm) weld HAZ from the work of Harrison and Hart20) and is believed to have contributed to the better performance of the vanadium steels in Fig.4. The effect has also been observed in both laser and electron beam welds in vanadium-containing steels as evidenced by the microstructures in Fig.6 & 7. Finally, the greatest contrast, epitomising the vanadium-effect, is perhaps in Fig.8 which shows the influence of a high level of vanadium on weld metal produced from cellulosic coated electrodes, as part of a development programme in the late 70’s and early 80’s.27) From a weldability point of view, the vanadium-effect is a great potential advantage but it would appear that more work would need to be done to give a greater understanding of how it occurs, so that the effect can be maximised.

3.2 Intragranular Microstructural Development

The development of intragranular acicular ferrite in weld metals has now long been recognised to be substantially associated with the presence of non-metallic inclusions, produced in substantial amounts in most arc weld metals. Their precise composition however has often been considered an important aspect to allow them to be effective particle microstructural controlling agents. It’s worth asking whether the effect of these particles in weld metals, works solely through a nucleation role for the gamma to alpha transformation, or whether there is another, additional or alternative explanation/mechanism that can explain the vanadium effect, which at least in parent steels, is unlikely to be through an influence on non-metallic inclusions. The current information in the literature that might aid the understanding of the “vanadium-effect” is sparse but an early brief paper by Uchino et al28) and Hasegawa et al29) looking at mechanisms of ferrite grain refinement in normalised high niobium vanadium steels, suggests that it might arise from the effect of interphase precipitation of VN at the gamma alpha interface, preventing ferrite grain growth as well as from the precipitation of VN within austenite grains. Ochi et al30) reported they had found that VC, precipitated on VN, itself precipitated on MnS particles provided effective sites for intragranular nucleation of ferrite. More recently still, Kimura et al31) have described how they believe that VN particles are able to produce a finer ferrite grain size in heavy gauge H-shapes. They believe VN precipitated at both gamma grain boundaries and grain interiors provide effective nucleation sites for ferrite. The nucleation of ferrite in an allotriomorphic form on VN precipitates at prior gamma grain boundaries is certainly consistent with the blocky grain boundary ferrite which in the author’s experience is characteristic of a vanadium containing weld metal and very well emphasised in the microstructure in Fig.8b.

In part the development of acicular ferrite should be seen as a competition between transformation events at and near the austenite grain boundary and intragranular nucleation events. More specifically, a competition between two alternative sequences of events. The first being stable growth of the allotriomorphic ferrite as the temperature decreases, allowing the remaining austenite to transform by intragranular nucleation and growth. The second sequence is that of the α/γ boundary of the allotriomorphic ferrite becoming unstable and

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changing to the more rapid Widmanstätten ferrite growth mechanism which sweeps over the austenite grain before any intragranular nucleation and growth can occur. The work of Cuddy and Lalby32) on HAZ microstructure development in TiN is relevant to this. They noted that in a TiN steel Widmanstätten ferrite growth from the grain boundary allotriomorphic ferrite was markedly restricted, compared with that in a Ti free reference steel, and that blocky, intragranular ferrite was promoted instead. They believe that fine TiN particles retarded the rapid edgewise growth of the Widmanstätten plates and that this allowed time for ferrite to nucleate intragranularly, aided by larger TiN particles. So its worth considering whether the vanadium effect occurs more because of what vanadium does either directly or through VN interface precipitation, on the stability of the α/γ boundary associated with the allotriomorphic ferrite, thus allowing intragranular nucleation to occur, or whether it has a more direct effect on enhancing intragranular nucleation sites. He and Edmonds33) have recently reported strongly producing the vanadium acicular ferrite effect in some high purity experimental C:Mn steels. They observed segregation of vanadium both at prior austenite grain boundaries and within austenite grain interiors, and have suggested34) that Fe-V clusters might provide suitable sites for embryonic acicular ferrite nucleates.

Clearly a greater understanding of the mechanism of the V-producing acicular ferrite effect is needed so that it can be actively applied to allow the consistent production of acicular ferrite in both the HAZ and weld metal and hence give improvement in joint toughness.

4. CONCLUDING REMARKS

Just as with other microalloying routes vanadium-microalloying, by facilitating reductions in steel carbon content, can provide improvements in steel weldability. Moreover, there is now clear evidence that vanadium-microalloying can bring additional benefits, especially to weld joint toughness, by virtue of its influence on microstructural development. Specifically the promotion of intragranular decomposition of austenite allows the development of finer effective grain sizes with the concomitant improvement in resistance to cleavage fracture.

In contrast to niobium, vanadium-microalloying is most effective with enhanced nitrogen levels, yet there is understandable concern with this, based on the established adverse effect of nitrogen on weld joint toughness in C:Mn and C:Mn:Nb steels. More data are needed on the effect of nitrogen in vanadium microalloyed steels, but at least for lower heat input welding available data shows little or no adverse effect, possibly because of the beneficial effects of vanadium on HAZ and weld metal microstructures.

Power beam welding, especially laser welding, is gaining in areas of application and provides significant opportunities for applying and maximising the ‘vanadium effect’ to both HAZ and fused zone microstructural development.

REFERENCES

1. Bailey N and Jones S B: “The solidification cracking of ferritic steel during submerged arc welding”, Weld. J, 1978, 57, 217s-213s.

2. Mandel’berg S L, Rybacov A A and Sidorenko B G: “Resistance of weld joints in steel tubes to hot cracking”, Automatic Welding, 1972, 25, 3, p1-5.

3. Nakamura H et al: Proceedings of the 1st International Conference on Fracture 2, Sendai, Japan, 1965, p863.

4. Ito Y and Nakanishi M: IIW document X-668-72.

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5. Bailey N, Coe F R, Gooch T G, Hart P H M, Jenkins N and Pargeter R J: “Welding steels without hydrogen cracking”, Abington Publishing, 1973.

6. British Standard specification for arc welding of carbon and carbon-manganese steels, BS 5315: 1984 (see also BS EN 1011-2:2001).

7. Yurioka N and Kasuya T: “A chart method to determine necessary preheat in steel welding”, Welding in the World, 33, 5, p327-334.

8. Stahl-Eisen – Werkstoffblatten 088 1993, Weldable fine grained structural steels, Guidelines for processing, particularly for fusion welding (see also BS EN 1011-2:2001).

9. Hart P H M and Harrison P L: “Compositional parameters for HAZ cracking and hardening in C-Mn steels”, Weld. J, Oct 1987, 66, 10, p310s-322s.

10. Hart P H M, Matharu I S and Jones A R: “The influence of reduced carbon equivalent on HAZ cracking in structural steels”, 7th Int. Conf on Offshore Mechanics and Arctic Engineering, Houston, February 1988.

11. Panton-Kent R: TWI Journal 1, 3, 1992, pp334-364. 12. Courdreuse L et al: Proc International Conference on Interaction of steels with Hydrogen

in Petroelum Industry Pressure Vessel Service. Paris, France, March 1989, MPC, pp51-59.

13. Shimomura J et al: ISIJ International, 31 1991, pp379-386. 14. Suzuki H: 1983 IIW Doc. IX-1279-83. 15. Yurioka N, Ohshita S and Tamerhiro H: March 1981 AWRA Symposium, Pipeline

Welding in the ‘80’s, p1-15. 16. Duren C and Lorenz K: 1982 IIW Doc. IX-B-11-82. 17. Mitchell P S, Hart P H M and Morrison W B: “The effect of microalloying on HAZ

toughness”, Microalloying ’95, Pittsburgh June 1995. 18. Crowther D N: British Steel Contract Report for Vanitec, March 1994. 19. Wang G R, North H T and Lewis G K: Weld. J., January 1969, 1, p14s-22s. 20. Harrison P L and Hart P H M: “Relationships between HAZ microstructure and CTOD

transition behaviour in multipass C:Mn steel welds”, 2nd Int Conf on Trends in Welding Research, ASM International, Gatlinburg, Tennesee, May 1989.

21. Li N et al: ISIJ International, 41, 2001, pp46-55. 22. Hanners N E and Johnson Holmquist B: “Influence of vanadium on the heat affected zone

properties of mild steel”, Metal Science 1974, 8, p228-234. 23. Lagneborg R, Siwecki T, Zajac S and Hutchinson B: “The role of vanadium in

microalloyed steels”, Scandinavian Journal of Metallurgy, 1999, 28, 5. 24. Hart P H M and Mitchell P S: “The effect of vanadium on the toughness of welds in

structural and pipeline steels”, Weld. J. July 1995, p240s-248s. 25. Dolby R E: “The effect of vanadium and niobium on weld metal properties”, Proc of

Steel for Linepipe and Pipeline Fittings, The Metals Society, London, October 1981. 26. Abson D J and Francis-Scrutton N: “Influence of Al,Ti,V and O on acicular ferrite

formation in TIG weld metal”, 15th Int Offshore Mechanics and Arctic Engineering Conference, June 1996.

27. Boothby P J and Hart P H M: “Welding 0.45% vanadium linepipe steels”, Metal Construction, September 1981, 13, 9, p560-569.

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28. Uchino K, Ohno Y, Yano S, Hasegawa T and Morikawa H: “Development of high N-V type low Ceq HT50 (development of low carbon equivalent and high N-V HY50 steel plates - II)”, Transactions ISIJ, 1982, Vol 22, p405.

29. Hasegawa T, Funaki S, Morikawa H, Ohno Y and Uchino K: “Mechanism of ferrite grain refinement in normalised high N-V steels, Transactions ISIJ, 1982, Vol 22, p404.

30. Ochi T, Takahashi T and Takada H: “Improvement of the toughness of hot forged products through intragranular ferrite formation”, 30th Mechanical Working Steel Processing Conf, Iron & Steel Soc., Warrendale, USA, 1988.

31. Kimura T, Kawabata F , Amano K, Ohmori A, Okatsu M, Uchida K and Ishii T: “Heavy gauge H-shapes with excellent seismic-resistance for building structures produced by the third generation TCMP”, ASM Materials Solutions, September 1999.

32. Cuddy L J and Lalby J S: International Conf on ‘Technology & applications of high strength low alloy (HSLA) steels’, ASM Metals Congress, Philadelphia, October 1983.

33. Edmonds D V: "The promption of acicular ferrite microstrucutre in C:Mn steel by vanadium", The 4th International Conference on HSLA Steels, Nov 2000, Xi'an China, pp260-265.

34. Zhang M, He and Edmonds D V: Solid-state Phase Transformations '99 (JIMIC-3), p1477.

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Table 1 Linear regression equations for predicting cooling times (∆t800°-500°C/s) Equation of Mean Line Compositional Characterising Parameter (CCP) 1 Log ∆t (250HV) =

M (CCP1) + K C + 0.1191Mn+0.3010V + 0.0828Ni+0.1777Mo

(C + 61238MoNiVMn

+++ )

2 Log ∆t (275HV) = M (CCP2) + K

C + 0.1032Mn + 0.2480V + 0.0700Ni + 0.1508Mo

(C + 714410MoNiVMn

+++ )

3 Log ∆t (300HV) = M (CCP3) + K

C + 0.0828Mn + 0.0213V + 0.0362Ni + 0.1037Mo

(C + 10284712MoNiVMn

+++ )

4 Log ∆t (325HV) = M (CCP4) + K

C + 0.0641Mn – 0.0293V + 0.0136Ni + 0.0969Mo

(C + 10743416MoNiVMn

++− )

5 Log ∆t (350HV) = M (CCP5) + K

C + 0.0485Mn – 0.0565V + 0.0047Ni + 0.0758Mo

(C + 132131821MoNiVMn

++− )

6 Log ∆t (375HV) = M (CCP6) + K

C + 0.0508Mn – 0.0390V + 0.0020Ni +0.0836Mo

(C + 125002630MoNiVMn

++− )

7 Log ∆t (400HV) = M (CCP7) + K

C + 0.1972Mn – 0.0996V + 0.0917Ni + 0.1792Mo

(C + 611105MoNiVMn

++− )

8 Log ∆t (450HV) = M (CCP8) + K

C + 0.3743Mn – 0.3930V + 0.0240Ni + 0.3854Mo

(C + 34233MoNiVMn

++− )

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Fig. 1 The effect of vanadium on the HAZ toughness of 0.12%C-1.6%Mn steels multipass welded at 2 KJ/mm. (∆t8/5 = 12 secs)

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Fig. 2 The effect of vanadium on the temperature for 0.25 mm CTOD in the HAZ of low-carbon (<0.07%) line pipe steels welded at 5 KJ/mm, using two passes in the as-welded condition.

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Fig. 3 The effect of heat input on the as-welded HAZ toughness of V, Nb and V+Nb steels.

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Fig. 4 Influence of SCGHAZ hardness on 0.1 mm CTOD transition temperature (a) as welded condition.

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a) C:Mn:Nb steel

b) C:Mn:0.1V steel

Fig. 5 HAZ microstructures at 5 KJ/mm.

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a) 0.07C, 1.7Mn, 0.43Si, 0.026Al steel

b) 0.07C, 1.7Mn, 0.43Si, 0.16V, 0.029Al steel

Fig. 6 Fused zone microstructure of autogenous laser welds.

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a) 0.07C, 1.7Mn, 0.43Si, 0.26Al steel

b) 0.07C, 1.7Mn, 0.43Si, 0.16V, 0.029Al steel

Fig. 7 Fused zone microstructures of sutogenous welds in electron beam welds.

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a) 0.07C, 1.7Mn, 0.3Mo, 0.03Nb steel

b) 0.03C, 1.7Mn, 0.45V steel

Fig. 8 Root run weld metal microstructures of pipe welds made with E6010 type electrodes.

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Vanadium Microalloying in Steel Sheet, Strip and Plate Products

Robert J. Glodowski

STRATCOR

Process Developments in the Production of Flat Rolled Products

Flat rolled steel products, including sheet, strip and plate, have seen nearly a century of development since the first continuous rolling processes were installed. Through these years of development, the product has been continually refined, initially with emphasis on improved properties through chemistry control and subsequent heat treatments. Later, with continued pressures to improve the cost effectiveness of steelmaking, more efficient process methods have evolved to produce the desired properties at a lower cost. One of the most effective ways to reduce the cost of production of steels is to shorten the process routing, eliminating operations that add unnecessary cost. As a result, a major objective in the industry today is to produce steel as directly as possible from liquid steel to final product. Medium and thin slab casting processes along with direct charging into reheat or equalizing furnaces characterize this quickly developing technology in the production of steel sheet and plate.

Alloy design has also evolved in concert with the process technology changes to achieve the desired objectives of improved properties with minimum processing. Metallurgical response to the hot rolling process has replaced the metallurgical response to heat treatment as the principal alloy design criteria. One of the more important innovations in steel technology in recent years is the combination of microalloying along with controlled hot rolling practices. This processing technology became known as controlled rolling (CR), or thermomechanical controlled processing (TMCP). Various versions of TMCP, sometimes identified with conflicting and confusing terminology, have developed over the years.

As the term thermomechanical controlled processing implies, controlling the rolling process for optimum metallurgical response involves control of both the mechanical and thermal aspects of the rolling. Mechanical control involves the reduction pass schedule, including total reduction and the number of passes and reduction per pass. The thermal control involves the deformation temperature throughout the rolling schedule from the reheat furnace through each rolling pass. In addition, the technology exists today for close control of the cooling cycle after the final rolling pass. This technology includes laminar water cooling after rolling, first adapted to hot strip mills, and later added to plate rolling mills as well.

In plate mills, the controlled water cooling after rolling has become known as accelerated cooling (AC), or accelerated controlled cooling (ACC). In this process, the cooling rate can be controlled to optimize the austenite to ferrite/pearlite transformation, or even to quench the steel to alternate microstructures. In all cases, the technology involves controlling the rate and the extent of cooling to achieve the desired microstructure.

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Metallurgical aspects of HSLA Steels produced by Thermomechanical Controlled Processing (TMPC)

With the development of process technologies including thermomechanical controlled processing and accelerated cooling, the emphasis has been and continues to be on optimizing the alloy system to meet property requirements at low cost.1) High Strength Low Alloy (HSLA) steels have been developed using limited additions of microalloys to achieve higher strength levels in the as-rolled condition. Microalloying has generally replaced increased carbon levels as the primary strengthening process for structural steels in sheet and plate. The low toughness and poor weldability of the higher carbon grades restricted their acceptability, and therefore created the demand for alternatives. Microalloyed HSLA steels rely primarily on two strengthening mechanisms – ferrite grain refinement and precipitation strengthening – to increase strength levels beyond the base strength achieved from carbon additions for pearlite strengthening and from manganese and silicon additions for solid solution strengthening. Integrating the new processing conditions with an appropriate alloy system to optimize the grain refinement and precipitation strengthening mechanisms became the primary metallurgical objective.

• Grain Refinement

To achieve ferrite grain refinement in the final as-rolled product, several approaches are used separately or in combination.2,3) First, the austenite should be “conditioned” during the rolling process to provide the maximum amount of grain boundary surface area. One way this conditioning can be accomplished is by rolling at temperatures below the recrystallization stop temperature, thereby producing flattened or “pancaked” austenite grains. These flattened grains then promote the formation of small ferrite grains in the final product. Niobium alloyed steels using these “controlled rolling” (CR) processes were developed over 30 years ago. In the CR process, the temperature for the final rolling passes must be held below the temperature at which recrystallization will occur, the “recrystallization stop temperature”.

Recrystallization controlled rolling (RCR) was developed later. The RCR process relies on repeated interpass recrystallization to achieve a fine austenite grain size, with controlled temperatures and reduction schedules to achieve the desired results. The RCR rolling process involves finish-rolling temperatures more typical for plain carbon steels, while the CR process requires lower finish rolling temperatures. These higher rolling temperatures promote repeated recrystallization. Since niobium has been shown to raise the recrystallization temperature, vanadium microalloying was naturally more compatible with the RCR process. Vanadium does not inhibit the austenite recrystallization necessary for the RCR process.

Ferrite nucleation occurs primarily on the prior austenite grain boundaries. Both CR and RCR rolling generate an austenite with a large amount of grain boundary surface area (austenite interfacial area), promoting transformation to a small ferrite grain size. These small ferrite grains contribute to both strength (via the Hall-Petch relationship shown in Figure 1.) and to increased toughness. The effect of the ferrite grain size on the final properties of the steel is independent of the austenite conditioning process used to achieve that ferrite grain size. As long as small ferrite grains are produced, it does not matter which rolling process is used to achieve those results. At an equivalent ferrite grain size, the contribution of grain size to the final properties (strength and toughness) will be comparable. Therefore, the test of the process is the final ferrite grain size, and not the condition of the austenite that produced the ferrite.

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Accelerated cooling can be applied after finish rolling, using either the CR or RCR rolling process, to lower the austenite-ferrite transformation temperature. Lowering the transformation temperature generates smaller ferrite grains by two mechanisms. First, an increased ferrite nucleation driving force is generated because of the undercooling effect. This undercooling is the difference between the equilibrium transformation temperature and the actual transformation temperature. The higher ferrite nucleation driving force will increase the number of ferrite grains initiating at the austenite grain boundaries and other favorable sites.

After ferrite nucleation, the lower transformation temperatures will inhibit the ferrite growth rate. Interphase precipitation of grain boundary pinning microalloy constituents may also suppress the ferrite growth rates. Lowering the transformation temperature, and thereby reducing the ferrite growth rate, may also be accomplished by alloying to increase the hardenability of the steel. Typically, manganese additions are used to provide increased hardenability. However, increasing the ferrite nucleation driving force through the undercooling effect can only be achieved by using accelerated cooling.

In plain carbon steels, ferrite nucleation is normally expected to occur at the austenite grain boundaries. Ferrite nucleation may also be initiated at precipitates within the austenite grains. Work has been published by Kimura et.al.4) suggesting the precipitation of vanadium nitrides can be effective for promoting fine ferrite grains, even in larger sections where the cooling rates cannot be accelerated greatly. In this model, the fine microalloy precipitates not only reduce the recrystallized austenite grain growth rate, but also promote ferrite nucleation. The relative effectiveness of VN to enhance ferrite nucleation compared to other precipitates was explained in terms of the change in interfacial energy as shown in Figure 2. The end result is a finer ferrite grain size as a direct result of microalloy precipitation in the austenite phase prior to transformation as well as interphase precipitation during transformation from austenite to ferrite. Figure 3 is the schematic illustration used by Kimura et.al. to show a first stage of rolling to promote repeated recrystallization of the austenite. A second stage of rolling at a somewhat lower temperature is shown to promote deformation induced VN precipitation. The VN will precipitate at both the austenite grain boundaries reducing austenite grain growth, and within the interior of the grain boundaries providing ferrite nucleation sites.

• Precipitation Strengthening

To provide effective precipitation strengthening, the microalloy used should have the ability to stay in solution during the heating and rolling stages, then precipitate in fine, well-dispersed particles during or after the austenite to ferrite transformation. Vanadium has the demonstrated properties to be well suited for this purpose. The higher solubility of vanadium carbonitrides compared to the other microalloying choices available makes vanadium a strong and predictable precipitation strengthener.

Because the solubility product of VN is significantly lower than VC, nitrogen plays a significant roll in the strengthening process with vanadium.5) Also, because nitrogen has a much higher solubility in ferrite than carbon, the availability of nitrogen for VN precipitation in ferrite is quite high. For those microalloys that require carbon for effective precipitation, the amount of carbon available in ferrite can be much less than the available nitrogen. As a result, vanadium precipitation strengthening can be predictable over a relatively large range of vanadium addition levels. Most chemistry-strength models show a linear effect of vanadium additions as long as an

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appropriate amount of nitrogen is available in the steel. This allows vanadium to be used as a “trim” element that can be adjusted to meet property requirements. As grade or thickness changes are made, the desired properties can be predictably achieved by adjusting the vanadium levels. The vanadium adjustments can be made on a grade to grade basis, or even adjusting vanadium levels from heat to heat within a grade to compensate for thickness changes or variations in other alloys, including residual elements.

Figure 4 is a schematic showing the relative contribution of each of the strengthening mechanisms to HSLA steels. The relative effects of each strengthening mechanism can be visualized in this figure, illustrating the importance of both ferrite grain size control and precipitation strengthening to achieve the 550 MPa yield strength level. This example was derived from analysis of a production grade of HSLA strip, direct rolled from a 50mm cast slab.

Advantages of Vanadium for Flat Rolled HSLA Steels

Traditional microalloying systems are being challenged by the new production developments that have been commercialized in flat rolled products. These new production processes include thin slab cast and direct rolled strip steel production, and the plate mills with medium thickness slabs directly charged into reheat furnaces for final rolling. While the use of vanadium has always been a viable option, these new processes have added new reasons to consider vanadium as the alloy of choice.

Desirable characteristics of an alloy system for HSLA steels using the newer low cost steelmaking processes include the following:6,7,8)

• Compatibility with the electric furnace melting process widely used in many new facilities.

• Minimal precipitation during solidification to reduce cracking problems during casting.

• Low solution temperature (high solubility) to insure the microalloy is in solution prior to rolling.

• Precipitation strengthening should occur after finish rolling to minimize roll force requirements.

• Compatibility with high finish rolling temperatures to minimize roll wear and facilitate gauge control.

• Compatibility with RCR rolling for maximum ferrite grain refinement.

• Predictable precipitation strengthening, providing yield strengths from 350 to 550 MPa.

Vanadium microalloying clearly demonstrates the capability of meeting all of these characteristics. Many of the new mills are predominately electric furnace melting processes. The raw material is predominantly scrap steel, causing the metallic residuals to be higher than normally seen in steels produced from blast furnace iron in basic oxygen furnaces (BOF). Even more significant, the nitrogen levels are substantially higher than in the BOF steels. Because vanadium utilizes nitrogen as an integral part of the alloy system, the normally undesirable presence of the higher nitrogen can be turned into an advantage. Compatibility with higher

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nitrogen levels is a desirable characteristic of vanadium, but vanadium strengthening can be used effectively in BOF steels with managed nitrogen levels.

The castability of vanadium steels has been demonstrated in the literature,8,9) and the wide use of vanadium in use in structural grades of plate and sheet provides verification of minimal cracking problems. The low solution temperatures of the vanadium precipitates insure that the alloy is in solution even if energy saving practices dictate lower reheat temperatures and shorter soaking cycles. Additionally, the low solution temperatures allow for excellent gauge control and minimal roll wear in strip steels because the roll forces are minimized at finish rolling temperatures, typically from 900 °C to 950 °C.

With accelerated cooling, much of the vanadium precipitation will occur randomly in the ferrite after transformation. This preferred precipitation will occur at temperatures that are typical for coiling strip and light plate, near 600°C. For heavier plates where accelerated cooling is less effective, the precipitation will still occur predominately during or after transformation. Using the rolling practices previously described, final ferrite grain sizes below 5µm are readily produced in hot rolled strip. In plate steels, final grain sizes of 10 µm or less have been demonstrated.10)

Hot Rolled Sheet Applications

• Strengthening

Recent publications6,11) have reported on an evaluation of production lots of low carbon sheet steels containing various levels of manganese, vanadium and nitrogen. The steel samples were from three different steel mills, each producing steel using electric furnace melting processes, casting thin (50 mm) slabs and direct charging into equalization furnaces prior to rolling. The sheet samples were evaluated for mechanical properties, microstructure, and strain aging characteristics. Table 1 is a summary of the chemistry, mechanical properties, grain size and strain aging index for each of these steels.

Regression analysis of the mechanical property data provided a reasonable estimate of the strengthening effects of vanadium and nitrogen. Within the range of vanadium additions of this data (0.04% to 0.13%), the strengthening contribution of vanadium was 15 MPa for each 0.01% V added. For nitrogen, the estimated strengthening was 7.5 MPa for each 0.001% N added. These strengthening coefficients are valid for the data range from this data set, and can be expected to be reliable as long as the nitrogen level is sub-stoichiometric with respect to the vanadium addition. They are also consistent with the nitrogen strengthening rates reported in previous studies.5,7)

These results support the well established observations that vanadium provides more strengthening when nitrogen is available up to the stoichiometric levels of V:N. The implication for thin slab casting operations is that the nitrogen levels of electric furnace steels can be used to an advantage to optimize the cost effectiveness of the vanadium alloy addition. With appropriate management of nitrogen levels, vanadium can be effectively used in steels melted in basic oxygen processes as well.

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Ferrite grain size measurements were determined by microstructural analysis of samples using the intercept method. Figure 5 shows the ferrite grains size as a function of the sheet thickness. The dependence of grain size on sheet thickness reveals the effects of processing conditions on grain size, since both total reduction and cooling rate are directly dependent on the sheet thickness. In particular, the faster cooling rate of the thinner sheets would be expected to promote a finer ferrite grain size. The higher alloy levels, and possibly optimized processing conditions, of the higher strength sheets greater than 500 MPa resulted in finer grain sizes compared to the lower strength grades. This again demonstrates the additive effects of grain size and precipitation strengthening that allows the production of these higher strength grades. The 3 to 7 micron grain size for the vanadium microalloyed high strength grades is comparable to 3 to 8 micron results reported for ultra-high strength (750 MPa) grades using Nb and Ti along with special processing for maximum effect.12)

• Strain aging

Strain-aging test data were also reported, demonstrating the ability of vanadium to eliminate the aging effects of “free” nitrogen in these steels, even when the total nitrogen levels exceeded 200 ppm. In this study, the strain-aging index is defined as the amount of increase in flow stress after aging a prestrained sample at 100 C for 1 hour.13) The amount of prestrain used for these tests was 7.5%. The results of the strain aging index study are summarized in Figure 6. The average strain aging index (A.I.) from the results shown in Table 1 (average of 1 MPa for 16 different steel samples) are compared to average results from 5 different samples of C-Mn steels (average 39 MPa). The nitrogen level for the vanadium steels ranged from 80 to 200 ppm, while the carbon steels samples had nitrogen levels from 70 to 110 ppm.

Another recent publication14) described the development and production of 550 MPa HSLA V-N-Mo microalloyed sheet steels, rolled from 55 mm and 65 mm thin slabs above the recrystallization stop temperature (RCR rolling) and followed by accelerated controlled cooling (ACC). The typical chemistry used for these grades is shown in Table 2.

In this publication, the author describes the process development work needed to optimize the rolling practice to obtain consistently high properties of this grade. The key TMCP parameters that were monitored and controlled were described as follows:

1. Tunnel furnace discharging temperature

2. F1-F6 reduction ratios and deformation strains

3. Roll forces and flow stresses

4. Total processing time

5. Mill speeds (F1 entry and F6 exit speeds)

6. Finishing temperature

7. Coiling temperature

8. Cooling rate

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A recrystallization control rolling - accelerated controlled cooling (RCR-ACC) process technique was adopted as best suited for both the mill capabilities and the chemistry used. The author concluded that the RCR-ACC technique was an effective in producing a fine, uniform, transformed ferrite structure with a fine grain size in the ASTM range of 10.5 too 11.5. Heavy reduction passes at a higher temperature regime were determined to be critical at the start of hot rolling, providing complete recrystallization of the austenite. Short processing times (i.e. fast rolling speeds) and ACC with intensive cooling rates were also stated to be important to promote a rapid precipitation of V(C,N) and Mo2C. The performance of the steels from this process has been exceptional, safely exceeding the minimum 550 MPa yield strengths required for the grade.

Plate Applications

Vanadium has been widely used in plate steels for many years in North America, particularly for 350 MPa grade steels like ASTM A572, Gr. 50. The benefits of vanadium addition in plates are similar to those identified in sheet steels, principally the enhancement of properties, through grain refinement and precipitation strengthening in the as-rolled condition. Mitchell, et.al.10) published the results of an extensive laboratory study documenting the potential properties of 20-25 mm plates using vanadium microalloying and various thermomechanical treatments, including hot rolling, recrystallization controlled rolling (RCR) and accelerated cooling (ACC) after rolling.

A summary of the chemistry and mechanical properties of a series of RCR plates is shown in Table 3. In this case, the nitrogen levels were not enhanced beyond a normal residual level. Two different levels of Ni and Cu were investigated, representing what one might expect for BOF vs. EAF melting processes.

Because of the differences in carbon, nickel and copper levels, it was necessary to normalize the strength data to compare the strengthening effects of vanadium. This was accomplished by utilizing available equations to predict the strengthening effects of the elements included in a carbon equivalent model, specifically C, Mn, Cr, Mo, Ni and Cu. The model used was as follows:

YS = 41.4 + 575.2×CEV + (27401×Neff − 2) ×V1/2 + 419.5×t-1/2 (Eq. 1)

UTS = 74.1 + 985.1×CEV + (31125×Neff – 39) ×V1/2+ 181.5× t-1/2 (Eq. 2)

Where CEV = C + Mn/6 + (Cr + Mo)/5 + (Ni + Cu)/15

Neff = %N – Ti%/3.42

t = thickness

The effect of vanadium on the tensile properties was determined by correcting for differences in CEV. The results are shown on the Figure 7, plotted for a constant CEV value of 0.38. The strengthening rate of the vanadium additions appears to be substantially less than that determined for sheet steels and reported previously. A linear regression analysis of the yield strength data in Table 3 produced a coefficient for vanadium of about 6.5 MPa for each 0.01% V added. This value approximates the slope of the data in Figure 7 quite well. It is likely that the strengthening effects of vanadium were reduced in part because the nitrogen levels were not enhanced as they were in the sheet steel data. Also, the author determined the contribution to yield strength from

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precipitation strengthening accounted for over 10 MPa for each 0.01%V added when adjusted for the effects of carbon, solid solution strengthening and grain size.

The impact properties of these steels were all quite good, as shown in Table 3. 53J impact transition temperatures (ITT) were in the range of –70 to –95 °C. There appears to be some effect of the carbon levels, with a tendency for the 0.12 – 0.14% carbon to exhibit a marginally higher level of ITT than observed in the 0.01% carbon steels. The microstructures of these steels were all equiaxed ferrite/pearlite with a relatively uniform grain size of the order of 8 – 11 µm. Weldability tests were also performed on these steels, with excellent levels of HAZ toughness being observed at heat inputs up to 5 kJ/mm. Because of the low levels of CEV normally used in these steels, there should be no problems with cold cracking.

SUMMARY

Vanadium has proven to be a popular choice as a microalloy for flat rolled sheet and plate steels. Some of the observed advantages of vanadium as an alloy choice are as follows:

• Vanadium utilizes nitrogen as part of the alloy system. As such, vanadium can easily adapt to the increasing percentage of steels produced by Electric Arc Furnaces.

• Castability problems are minimized with vanadium compared to some other microalloy approaches.

• Higher solubility of the V(C,N) precipitates permits higher alloy levels in the steel even when restricted to the lower reheat temperatures common in modern steelmaking facilities.

• Recrystallization control rolling techniques produce refined austenitic grains at normal carbon steel rolling temperatures.

• Precipitation occurs primarily during or after transformation to ferrite, thereby not contributing to roll forces during finish rolling.

• Reasonable coiling temperatures for sheet products (580-620 °C) are compatible with normal mill equipment.

• The V and N strengthening effect in sheet and plate steels can be quantified with confidence, and the strengthening is predictable and reasonably linear over a large vanadium addition range.

• Nitrogen enhances the strengthening effect of vanadium, improving the cost effectiveness of the vanadium alloy system even when used in steel producing mills that may have inherently low residual nitrogen levels.

• Hot rolled ferrite grain size can be refined using the V-N alloy system, achieving levels competitive with other alloy systems.

• Impact toughness and weldability of vanadium bearing HSLA steels are competitive with other alloy systems when processing conditions are controlled to produce the optimum ferrite grain refinement.

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REFERENCES 1. M. Korchynsky, “Cost effectiveness of Microalloyed HSLA Steels,” Conf. Proceedings,

International Symposium on Steel for Fabricated Structures , ASM International, Cincinnati OH, Nov. 1999, pp. 139-145.

2. R. Lagneborg, “The Manufacture and Properties of As-Rolled V-containing Structural Steels”, International Symposium 2000 on Vanadium Application Technology, Guilin, China, November 2000, pp. 47-58.

3. R. Lagneborg, T. Siwecki, S. Zajac, and B. Hutchinson, “The Role of Vanadium in Microalloyed Steels,” Scandinavian Journal of Metallurgy, Vol. 28, issue 5, October 1999.

4. T. Kimura, F. Kawabata, K. Amano, A. Ohmori, M. Okatsu, K. Uchida, “Heavy Gauge H-Shapes with Excellent Seismic-Resistance for Building Structures Produced by the Third Generation TMCP,” Proc. of International Symposium on Steel for Fabricated Structures, Cincinnati, USA, 1999, pp. 165-171.

5. T. Siwecki, A. Sandberg, W. Roberts, R. Lagneborg, “The Influence of Processing Route and Nitrogen Content on Microstructure Development and Precipitation Hardening in V-microalloyed HSLA Steels,” Thermomechanical Processing of Microalloyed Austenite, (Ed. By DeArdo, Ratz, Wray) TMS-AIME, Warrendale, USA, 1982, pp. 163-192.

6. R.J. Glodowski, “Effect of V and N on Processing and Properties of HSLA Strip Steels Produced by Thin Slab Casting”, 42nd MWSP Conference Proceedings, ISS, Vol. XXXVIII, 2000, pp. 441-454.

7. P.S. Mitchell, D.N. Crowther and M.J.W. Green, “The Manufacture of High Strength, Vanadium-Containing Steels by Thin Slab Casting,” Proc. of the 41st Mechanical Working and Steel Processing Conference, Baltimore, USA, 1999, pp. 459-470.

8. P.L. Lubensky, S.L. Wigman, D.J. Johnson, “High Strength Steel Processing via Direct Charging using Thin Slab Technology,” Microalloying ’95, Iron and Steel Society Inc., Pittsburgh, PA, 1995, pp. 225-233.

9. D.N. Crowther, M.J.W. Green, P.S. Mitchell, “The Influence of Composition on the Hot Cracking Susceptibility During Casting of Microalloyed Steels Processed to Simulate Thin Slab Casting Conditions”, “Microalloying in Steels”, Materials Science Forum, Vols. 284-286, (1998) pp. 469-476.

10. P.S. Mitchell, W.B. Morrison, D.N.Crowther, “The effect of Vanadium on the Mechanical Properties and Weldability of High Strength Structural Steels”, International Symposium Low-Carbon Steels for the 90’s, ASM - TMS, Pittsburgh, PA, October 1993, pp. 337-344.

11. R.J. Glodowski, “Vanadium-Nitrogen Microalloyed HSLA Strip Steels”, HSLA Steels ‘2000, The 4th International Conference on High Strength Low Alloy Steels, Xi’an, China, Oct. 2000, pp. 313-318.

12. R.D.K. Misra, J.E.Hartmann, G.C. Weatherly, A.J. Boucek, “Role of Structure and Microstructure in the Enhancement of Strength and Fracture Resistance of Ultra-High Strength Hot Rolled Steels”, Iron and Steelmaker, June 2001, pp. 29-38.

13. J.D. Baird, “Strain Aging of Steel – a critical review,” Iron & Steel, May 1963, pp. 186-192.

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14. L.K. Chiang, “Development and Production of HSLA 80 ksi (550 MPa) Steels at Gallatin Steel”, HSLA Steels ‘2000, The 4th International Conference on High Strength Low Alloy Steels, Xi’an, China, Oct. 2000, pp. 319-325.

15. M. Korchynsky, “The Growing Role for Vanadium in HSLA Steels for Thin-Slab Casting,” 33 METALPRODUCING, Penton Publications, Nov. 1998

16. W.B. Hutchinson, “Microstructure development during cooling of hot rolled steels” Proc. of Thermomechanical Processing of Steels, Vol. 1, May 2000, pp. 233-244.

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Table 1 Chemistry and mechanical properties of production grades of vanadium microalloyed hot rolled sheet steel.6)

Product Chemistry Mechanical Properties

Steel Type

Gauge mm

C wt.%

Mn wt.%

Si wt.%

A lwt%

V wt%

N wt.%

G.S. µm

YS MPa

UTS MPa

El. %

A.I. MPa

V 15.3 0.050 0.83 0.01 0.046 0.085 0.0090 7.91 431 470 38 2

V 3.0 0.059 0.76 0.02 0.036 0.054 0.0098 5.72 404 474 29 2

V 6.3 0.056 0.79 0.04 0.039 0.057 0.0086 6.90 423 468 33 -6

V 4.7 0.054 1.10 0.02 0.043 0.066 0.0105 5.44 383 518 31 -6

V 2.1 0.051 0.83 0.03 0.034 0.057 0.0084 5.00 391 453 29 2

V-N 12.6 0.064 1.01 0.01 0.039 0.065 0.0140 9.19 467 502 36 2

V-N 7.8 0.053 0.99 0.01 0.036 0.067 0.0156 5.29 525 550 31 12

V-N 2.1 0.056 0.78 0.03 0.037 0.058 0.0136 4.30 430 512 34 3

V-N 3.8 0.058 0.86 0.02 0.037 0.086 0.0163 5.43 473 575 25 -2

V-N 6.3 0.050 0.80 0.01 0.050 0.057 0.0142 6.39 364 502 32 -3

V-N 2.0 0.054 0.99 0.19 0.024 0.050 0.0150 5.68 463 na na 1

V-N 7.6 0.051 0.97 0.20 0.022 0.046 0.0140 6.38 430 na na 3

V-N 5.7 0.053 1.62 0.35 0.026 0.130 0.0200 3.35 567 na na 8

V-N 5.8 0.054 1.57 0.32 0.022 0.130 0.0200 4.35 572 na na 5

V-N 5.7 0.052 1.58 0.36 0.026 0.120 0.0200 3.78 548 na na -1

V-N 5.7 0.057 1.62 0.35 0.025 0.130 0.0200 3.73 568 na na 1

Table 2 Typical chemistry of 550 MPa HSLA sheet steels (zz) wt.%

C Mn Si Mo V N Al 0.056-0.075 1.25-1.45 0.02-0.15 0.010-0.055 0.120-0.130 0.0150-0.0205 0.013-0.035

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Table 3 Product analysis and mechanical properties of RCR rolled 25 mm plate steel.10)

Steel C wt.%

Mn wt.%

Si wt.%

Ni wt%

Cu wt%

Al wt.%

V wt.%

N wt.%

YS MPa

UTS MPa

El. %

J 40°C

54J ITT°C

1 0.12 1.39 0.26 0.18 0.18 0.049 0.000 0.0070 364 497 42 186 -75

2 0.14 1.43 0.26 0.02 0.02 0.029 0.050 0.0050 407 519 34 215 -75

3 0.13 1.38 0.24 0.16 0.18 0.042 0.050 0.0070 421 545 39 147 -75

4 0.13 1.39 0.25 0.17 0.17 0.049 0.100 0.0070 430 545 37 140 -70

5 0.09 1.46 0.25 0.15 0.18 0.046 0.000 0.0080 359 482 47 184 -70

6 0.10 1.41 0.25 0.02 0.02 0.032 0.029 0.0055 395 464 34 267 -70

7 0.11 1.40 0.26 0.02 0.02 0.030 0.028 0.0090 402 486 37 235 -95

8 0.10 1.45 0.26 0.02 0.02 0.025 0.054 0.0052 390 480 34 283 -90

9 0.10 1.47 0.27 0.18 0.18 0.047 0.060 0.0070 375 499 44 221 -90

10 0.09 1.46 0.25 0.18 0.18 0.040 0.100 0.0070 411 515 43 231 -90

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Fig. 1 Calculated effect of ferrite grain size on Yield Strength.

Fig. 2 Change in interfacial energy and driving force for ferrite nucleation form various precipitates.4)

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Fig. 3 Thermomechanical behavior during rolling of heavy section vanadium bearing steels. 4)

Fig. 4 Contributions of each strengthening mechanism in a 550 MPa hot rolled strip.1)

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Fig. 5 Relationship of thickness and Yield Strength to ferrite grain size of hot rolled sheet steel.6)

Fig. 6 Strain aging index results of vanadium steels compared to carbon steels.

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Fig. 7 Effect of vanadium on the tensile properties of 25 mm thick RCR plates.10)


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