CommunicationsDiscussion of “Precipitation Behaviorand Its Effect on Strengthening of anHSLA-Nb/Ti Steel”*

H.-J. KESTENBACH, S.S. CAMPOS, J. GALLEGO, andE.V. MORALES

In a recent article, Charleux et al.[1] attempted to identifythe type of carbonitride particles responsible for precipita-tion strengthening of commercial hot strip steels. In partic-ular, the authors presented microstructural evidence point-ing to the posttransformation precipitation of carbonitrideparticles on ferrite dislocations as the principal strengthen-ing mode. This mode, however, has never been encounteredin a series of detailed transmission electron microscopy(TEM) studies on commercial hot strip products conductedby the present authors during recent years.[2,3,4] It is there-fore the aim of the present discussion to show that car-bonitride precipitation strengthening in commercial hot stripsteels is not associated with ferrite dislocations, and to lookfor an alternative interpretation of the experimental evidencepresented by Charleux et al.

Several examples considered typical for the TEM mi-crostructure of commercial hot strip steels are presented inFigure 1. All of them reveal the presence of at least someferrite dislocations. However, no interaction between thesedislocations and the carbonitride particles can be recog-nized. The examples were selected from previous studieson five different microalloyed steels whose chemical com-positions and yield strength values are shown in Table I. Itshould be noted that at least two of the steels (NbTi-2 andNbTi-3) are rather similar to the one investigated byCharleux et al.,[1] employing the same microalloying ele-ments (Nb and Ti) and presenting comparable levels of yieldstrength.

Among the five micrographs included in Figure 1, in-terphase precipitation can be observed in four of them, asshown in Figures 1(a), (c), (d), and (e), while carbonitrideparticles that have nucleated in austenite appear in two ofthem, as shown in Figure 1(b) and (c). In each case, theorigin of the particles was identified by determining theirorientation relationship with respect to the surrounding fer-rite matrix.[5] It should be noted that many of the particlesin Figure 1 are fine enough to give precipitation strength-ening,[6] considering that mean particle sizes ranged fromabout 6 nm in Figure 1(a) to 2 nm for the interphase pre-

cipitation in Figure 1(c). Interphase precipitation is gener-ally acknowledged to be an important source of precipita-tion strengthening in commercial microalloyed steels,[7,8]

while the strengthening contribution of carbonitride pre-cipitation in austenite has been defended by the presentauthors[2–4,9] but is not generally recognized in the litera-ture.[10,11] No example of posttransformation precipitationof carbonitride particles in supersaturated ferrite has beenincluded in Figure 1, because, in commercial microalloyedhot strip steels during coiling, microalloy elements thathave remained in solution precipitate preferentially on veryclose “pre-existing” carbonitride particles rather than nu-cleate additional particles on farther-away ferrite disloca-tions.[2,4] On the other hand, the precipitation of new par-ticles in supersaturated ferrite can be of technologicalimportance in special circumstances such as after welding.In this case, rapid cooling from the austenitizing temper-ature as well as the absence of plastic deformation in austen-ite prevent the formation of pre-existing particles, leavingthe nucleation of carbonitride particles on ferrite disloca-tions as the only precipitation mode during a postweldingheat treatment.[12,13]

In their article, Charleux et al.[1] present their experimentalevidence for carbonitride precipitation on ferrite dislocationsin Figure 6. Very similar particle distributions have beenobserved by the present authors for carbonitride precipita-tion in austenite, where particles are known to nucleate onthe deformation substructure generated during hot rolling.An example is shown in Figure 2, where the same carboni-tride reflection has been used to generate particle contrastboth under bright-field conditions (Figure 2(a)) and underdark-field conditions (Figure 2(b)). As can be deduced fromthe position of the objective aperture (double exposure in-dicated by the arrow) in Figure 2(c), this reflection does notobey the Baker–Nutting orientation relationship and musttherefore have come from carbonitride particles that havenucleated in austenite.

On the other hand, in Figure 6 of their article, Charleuxet al.[1] present two electron diffraction patterns thatidentify the relevant carbonitride reflections as having aBaker–Nutting orientation relationship with the surround-ing ferrite matrix. Such diffraction patterns must have con-vinced the authors that the micrographs of Figure 6, andin particular the micrograph in Figure 6(a), which repre-sents the as-coiled hot strip steel, should contain car-bonitride particles nucleated in ferrite. However, the de-tection of heterogeneous precipitation of small particleson dislocations may not always be an easy task becausea clear distinction between dislocation and particle con-trast is required. Due to the complex nature of dislocationcontrast, which, for example, oscillates as a function ofdepth in the sample, dislocations that are sloping throughthe foil may show dotted images by themselves, with thefrequency of dots (or of other periodic image details) beingcontrolled by the appropriate extinction distance.[14]

Examples of such periodic dislocation contrast are shownin Figure 3. Some of the dislocation images in Figures3(a) and (c), identified by arrows, could be interpreted asshowing the precipitation of small particles on the dis-location line. In such a case, however, the absence ofparticles can be verified when changing the electron

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*M. CHARLEUX, W.J. POOLE, M. MILITZER, and A. DESCHAMPS:Metall. Mater. Trans. A, 2001, vol. 32A, pp. 1635-47.

H.-J. KESTENBACH, Professor, S.S. CAMPOS, Graduate Student, arewith the Department of Materials Engineering, Federal Unversity of SãoCarlos, 13565-905 São Carlos, SP, Brazil. Contact e-mail: dhjk@power.ufscar.br. J. GALLEGO, Assistant Professor, is with the Department ofMechanical Engineering, State University of São Paulo, 15385-000 IlhaSolteira, SP, Brazil. E.V. MORALES, Associate Professor, is with theDepartment of Physics, Central University of Las Villas, Villa Clara, Cuba.

Discussion submitted May 31, 2002.

diffraction conditions by a small sample tilt, as shown inFigures 3(b) and (d).

In the presence of strongly diffracted matrix beams, dislo-cation images have a typical width of between 0.2 jg and0.4 jg, where jg is the appropriate extinction distance for theoperating reflection that produces two-beam contrast.[14] Forlow-order ferrite reflections, the smallest extinction distanceis 27 nm,[15] which means that the image width of a disloca-tion must be expected to reach 5 to 10 nm. On the other hand,

precipitation strengthening in microalloyed steel has beenestimated to come from particles of about 3 nm in diameter.[6]

Clearly, such particles would be masked by the proper dis-location image. Excluding the case of g·b 5 0 (for whichthe dislocation may be out of contrast), it is therefore neces-sary to assure the absence of any strong matrix reflection toprove the presence of small precipitates on dislocations.

Returning once more to Figure 6 of the article by Charleuxet al.,[1] a rather strong matrix reflection (light area) can be

1014—VOLUME 34A, APRIL 2003 METALLURGICAL AND MATERIALS TRANSACTIONS A

Table I. Chemical Compositions (Weight Percent) and Yield Strength Values for Five Commercial Microalloyed Hot StripSteels Presented in Figure 1

Steel C Mn Si P S A1 Nb Ti V YS (MPa)

Nb 0.07 0.68 0.01 0.012 0.009 0.04 0.04 — — 310NbTi-1 0.05 0.55 0.02 ND ND 0.02 0.02 0.06 — 332NbTi-2 0.12 1.20 0.33 0.023 0.008 0.05 0.06 0.05 — 534NbTi-3 0.11 1.54 0.28 0.026 0.007 0.01 0.04 0.11 — 638NbTiV 0.14 1.38 0.25 0.018 0.007 0.07 0.04 0.04 0.03 599

Fig. 1—TEM micrographs of carbonitride particles and dislocations in five commercial microalloyed hot-rolled steels. Steels (a) Nb, (b) NbTi-1, (c) NbTi-2,(d) NbTi-3, and (e) NbTiV. Original magnification 75.000 times.

(a) (b) (c)

(e)(d)

noted in the upper left-hand corner of Figure 6(a). This ma-trix reflection cannot be explained by the diffraction patternshown in the lower right-hand corner of that figure, becausethe objective aperture position that has selected a 200 car-bonitride reflection is too far away from the 200 matrix re-flection (after discarding oxide reflections, the diffraction pat-tern can easily be indexed as representing a ferrite 012reciprocal lattice plane). It must therefore be questionedwhether the diffraction pattern shown in the lower right-handcorner of Figure 6(a) could have been generated from someadjacent area where another ferrite grain may have existed.It should be noted that, in the case of microalloyed steels,which are known for their small ferrite grain size, crossinga grain boundary during dark-field observation may happenrather easily, especially when using a carbonitride reflection,which does not illuminate the ferrite matrix. If this had hap-pened, the particles shown in Figure 6(a) could have beenformed in austenite. The reflection used for the particle con-trast in Figure 6(a) would continue to be a 200 carbonitridereflection, and the strong matrix contrast present in the upperleft-hand corner of Figure 6(a) could have come from a 110

ferrite reflection that may have passed through the same ob-jective aperture. It should be realized that a 200 carboni-tride reflection may come close to a 110 ferrite reflection onthe diffraction pattern because of similar lattice spacings(0.223 and 0.218 nm for the NbC and TiC planes, respec-tively, against 0.203 nm for the ferrite planes).

It is therefore suggested that the dotted contrast shownin Figure 6(a) of the article by Charleux et al.[1] may ei-ther represent carbonitride particles that were nucleated inaustenite or the weak-beam contrast of particle-free ferritedislocations.

This work received financial support from the São PauloState Agency FAPESP through Grant No. 99/11101-3. Re-search scholarships and assistantships were provided bythe Brazilian National Research Council (CNPq) to H-JK,by FAPESP to SSC and by the Brazilian Federal AgencyCAPES to JG and EVM. The authors are grateful to COSIPAfor supplying the steel samples.

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Fig. 2—Carbonitride particles in steel NbTi-3, nucleated on the austenite deformation substructure during hot rolling. (a) TEM bright-field micrograph,(b) dark-field micrograph of same area, and (c) diffraction pattern indicating carbonitride orientation. Magnification 150,000 times in (a) and (b).

(a)

(b) (c)

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Fig. 3—Dotted dislocation images in steel NbTiV. (a) and (b) Same area under bright-field conditions, and (c) and (d) other area under dark-field condi-tions. Orientation changes from (a) to (b) and from (c) to (d) as indicated by the diffraction patterns. Magnification 170.000 times.

(a) (b)

(d)(c)

REFERENCES

1. M. Charleux, W.J. Poole, M. Militzer, and A. Deschamps: Metall.Mater. Trans. A, 2001, vol. 32A, pp. 1635-47.

2. A. Itman, K.R. Cardoso, and H.-J. Kestenbach: Mater. Sci. Technol.,1997, vol. 13, pp. 49-55.

3. S.S. Campos, J. Gallego, E.V. Morales, and H.-J. Kestenbach: in HSLASteels ‘2000, L. Guoquan, W. Fuming, W. Zubin, and Z. Hongtao,eds., Metallurgical Industry Press, Beijing, 2000, pp. 629-34.

4. S.S. Campos, E.V. Morales, and H.-J. Kestenbach: Metall. Mater.Trans. A, 2001, vol. 32A, pp. 1245-48.

5. A.T. Davenport, L.C. Brossard, and R.E. Miner: J. Met., 1975, vol. 27,pp. 21-27.

6. T. Gladman, D. Dulieu, and I.D. McIvor: Microalloying 75, UnionCarbide Corporation, New York, NY, 1977, pp. 32-55.

7. W.B. Morrison: J. Iron Steel Inst., 1963, vol. 201, pp. 317-25.8. T. Gladman: The Physical Metallurgy of Microalloyed Steels, The

Institute of Materials, London, 1997, pp. 298-303.9. H.-J. Kestenbach and J. Gallego: Scripta Mater., 2001, vol. 44, pp. 791-

96.10. L. Meyer, C. Strassburger, and C. Schneider: in HSLA Steels: Metallurgy

and Applications, J.M. Gray, T. Ko, Z. Shouhua, W. Barong, and X. Xis-han, eds., ASM INTERNATIONAL, Metals Park, OH, 1986, pp. 29-44.

11. A.J. DeArdo: 8th Process Technology Conf., ISS, Warrendale, PA,1988, vol. 8, pp. 67-78.

12. H.-J. Kestenbach: Mater. Sci. Technol., 1997, vol. 13, pp. 731-39.13. J. Bosansky, D.A. Porter, H. Aström, and K.E. Easterling: Scand. J.

Metall., 1977, vol. 6, pp. 125-31.14. A. Howie and M.J. Whelan: Proc. R. Soc., 1962, vol. A267, pp. 206-30.15. P.B. Hirsch, A. Howie, R.B. Nicholson, D.W. Pashley, and M.J. Whe-

lan: Electron Microscopy of Thin Crystals, Plenum Press, New York,NY, 1965, p. 497.

Authors’ Reply

M. CHARLEUX, W.J. POOLE, M. MILITZER and A. DESCHAMPS

Over the past 2 decades, there has been considerable de-bate over the conditions under which carbonitride precipi-tation takes place in austenite.[1– 4] Analysis of this problemis not only significant to clarify the details of precipitationstrengthening mechanisms but also to establish the mecha-nisms by which Nb microalloying delays austenite recrys-tallization during hot deformation. The latter can be attrib-uted to either solute drag or precipitation.[5] The seminalpaper of Dutta and Sellars[6] outlined that Nb carbonitrideprecipitation kinetics depends critically on steel chemistryand processing conditions. It is well established that signif-icant precipitation takes place during plate rolling, which hasrelatively long processing times.[7] However, in strip rolling,the finish mill processing times are relatively short, i.e. onthe order of 10 seconds, such that there may not be suffi-cient time for precipitation to occur. Because the precipita-tion behavior is sensitive to chemistry and the details ofthe processing conditions, clarification of the precipitationsequence requires a case by case analysis.

In the analysis under discussion, we employed variousexperimental approaches to clarify the details of the precipi-tation process. Industrial hot strip material was compared tomaterial for which the industrial process was simulated inthe laboratory using hot torsion tests. Both materials werecompared with respect to yield strength and microstructureevolution during aging treatments at the coiling temperature.The results from the hot torsion simulations are an impor-tant starting point for discussion, because, in this case, the

samples could be directly cooled to room temperature, avoid-ing the coiling stage. A point that was not made sufficientlyclear in our original article[8] is that TEM investigations onthe as-quenched torsion samples showed no evidenceof anyfine precipitates. However, when these samples were reheatedto 650 °C, there was a significant increase in the yield stress(Figure 1 in Reference 8) and furthermore subsequent TEMinvestigations confirmed the presence of fine-scale precipi-tation (Figure 6(b) in Reference 8). We interpreted this asclear evidence that precipitation in ferrite plays a significantrole in the hardening of this alloy. Further, our TEM obser-vations are similar for the peak-aged condition of the simu-lated material and the as-coiled industrial material, as shownin Figure 6 of our original article.[8]

Kestenbach et al.’s discussion essentially concentrates onthe interpretation of Figure 6 in our original article. Notably,they give suggestions that what we call ferrite-nucleated par-ticles may be either austenite-nucleated particles, interphaseprecipitates, or simply weak-beam dislocation contrast. Wewill try to answer to each of these possibilities.

Austenite nucleated precipitates: First, for each dark-fieldmicrograph shown and for many others not shown in the ar-ticle, it has been checked that the distribution of precipitatesshown did correspond to actual ferrite dislocations, both byusing a dark-field weak beam contrast and by taking bright-field micrographs. For the torsion material aged for 80 hoursat 650 °C, Figure 1 shows the actual correspondence betweenthe particles and the ferrite dislocations. This figure is thegeneral case and not an exception. Second, the precipitatesshown in our study are not only extremely small but alsohighly elongated. In the case of the austenite nucleated par-ticles shown in Figure 2 of Kestenbach et al.’s discussion,the precipitates are aligned in the direction of the formeraustenite dislocations, but have a different morphology; i.e.,the precipitates have a roughly spherical shape.

Further, Kestenbach et al.suggest an alternative inter-pretation to the diffraction pattern in Figure 6 of our origi-nal article.[8] According to Kestenbach et al., the diffrac-tion spot from which the image was made could have come

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Fig. 1—Dark-field image of the torsion material aged for 80 h at 650 °Cshowing simultaneously the carbide precipitates and the ferrite dislocationson which they have nucleated.