In Steel Heat Treating in the New Millenium: An Internationals Symposium in Honor of Professor George Krauss. ASM International, 2000, pp. 515-519

Lath (Packet) Martensite Revisited

Gareth Thomas Department of Materials Science ans Mineral Engineering,

University of California, Berkeley, CA 94720-1760, USA, and Materials Science Division, Lawrence Berkeley National Laboratory,

Berkeley, CA 94720, USA

Greg Kusinski Department of Materials Science ans Mineral Engineering, University of California, Berkeley, CA 94720-1760, USA

Abstract

The structure of packet lath martensite is reviewed to reemphasize that its structure is two-phase with interlath untransformed austenite forming a multilayer with the martensitic (α′ ) laths. Some implications of this structure for structural steel applications, especially dual-phase steels, are reiterated.

Introduction

It is a great pleasure to participate in this Symposium honoring George Krauss, who has made so many contributions to our understanding of the physical metallurgy of steels. In this short paper the objective is to review the detailed substructure of lath or packet martensite, which is formed at relatively high Ms temperature (≥ 350°C). This terminology was originally defined in the classification of Krauss and Marder [1] in 1971 and has been frequently described in detail in several publications (e.g., refs. 2, 3). The excellent properties attainable in steels with packet/lath martensitic structures (carbon ≤ 0.35wt%) are summarized in refs. 4-11.

Characterization of Packet/Lath Martensite

Mader and Krauss [1] showed that packet martensites consist of dislocated laths (α′ ) and form in steels when the Ms transition temperature is above about 350°C (see Fig. 1). This temperature is of course very strongly dependent on composition--especially carbon. From the known values of the Ms temperature-composition relationships (e.g., Fig. 2), it is possible to design a composition [4] which ensures obtaining the

packet lath microstructure rather than the brittle twinned plate martensites which form at lower Ms transformation temperatures. Since the γ→α′ transition occurs with the K-S crystallographic relationships, there is a maximum of only 4 variants of the habit planes ({111}), and thus only a maximum of 4 packets (orientations) of α′ for each γ grain. If the γ grain size is reduced by controlled rolling, then the resultant α′ packet size is also reduced. Thus the α′ microstructure is refined by refining the prior austenite, e.g., in controlled rolling. McMahon and Thomas [2] first showed that the dislocated structures at martensitic lath boundaries (α′ ) were in fact defining thin microlayers of retained austenite. The identification of this retained or untransformed martensite requires careful electron microscopy and diffraction; a "classical" analysis is shown in Fig. 1. This interlath austenite was also revealed by high resolution lattice imaging electron microscopy from which it was suggested that there was considerable carbon enrichment at the α′ /γ interfaces. Further detailed work using convergent beam electron diffraction [3] confirmed this result. However, the final confirmation of this enriched/stabilized interlath austenite was obtained by the powerful atomic resolution method of field atom probe spectroscopy described in this volume by G. Smith and reported in refs. 3 and 7. An example of such analysis is shown in Fig. 3.

Transformation Mechanism

The above analyses strongly suggest that the transformation path is as follows: At the Ms temperature austenite is converted by shear into dislocated laths of martensite (α′ ) and the dislocations

sweep along at the γ/α′ transformation interface, picking up carbon (and other solutes present in the steel), reducing the dislocation mobility until at Mf no further transformation occurs. These events are schematically shown in Fig. 4. Thus, the austenite is stabilized by this "entrapment" and interfacial stabilization by solutes, especially carbon. The atom probe analyses (e.g., Fig. 3) show that the carbon content at this α′/γ interface can be more than an order of magnitude larger than the carbon content of the alloy as a whole. The resultant microstructure is an array of laths of α′ interspersed with thin sheets of stabilized austenite forming a multilayered structure (Fig. 1). The significance of this structure in relation to temper martensite embrittlement has been described by Sarikaya et al. [8]. Heat treatments of α′/γ in the range 300°C–500°C result in the austenite decomposition to interlath carbides (now the structure is similar to that of lower bainite), causing embrittlement (transgranular with respect to prior austenite)—the so-called "temper martensite embrittlement" (TME) as shown in Fig. 5.

Fig. 2 Schematic showing the calculated effect of some alloying elements on the Ms temperature. Lath martensite only forms for Ms ≥ 350° C; below this the martensite becomes twinned, plate-like (see also ref. 1)

Fig. 3 Analysis (atomic resolution) by field atom probe spectroscopy of lath martensite in Fe/0.3C/3Cr/2Mn quenched steel. Only C analysis is shown--note the very high values at the α′ /γ interfaces which are heavily dislocated (see Fig. 4) (courtesy S. J. Barnard, G.D.W. Smith, M. Sarikaya and G. Thomas, Scripta Metall., 15, 387-393 (1981). The other alloying elements also concentrate at these interfaces.

Fig. 1 Transmission electron microscopy (TEM) analyses of packet lath martensite (α′ ). The dark field image of austenite reflections reveals interlath untransformed austenite, i.e., lath martensite is actually a two-phase microstructure (α′ /γ). The α’ laths are heavily dislocated.

A

B

0.5 0.5 µµµµµµµµmm

0.5 0.5 µµµµµµµµmm

Fig. 4 Schematic to illustrate the transformation of γ to α′ and the stabilization of untransformed austenite between α′ laths. That the failure in the TME condition is along the axes of the laths (i.e., transgranular with respect to prior austenite grains) is verified by an in-situ straining experiment in a 1 meV electron microscope shown in Fig. 6, where the fracture path is along the α′ laths with little or no plastic zone developed at the crack tip.

Relation to Properties

The effect of the γ→α′ transformation substructures on strength and toughness received considerable attention in our early studies, and is reviewed in Ref. 4. More recently, emphasis has been placed on dual-phase steels containing ferrite and the α′

martensite described above, and obtained by rapid quenching from the intercritical α+γ phase field.

Fig. 5 Plot showing temper martensite embrittlement in experimental Fe/3Cr/1Mn/0.3C steel. Insets show DF images of transition from retained austenite films to embrittling M3C films at martensite lath boundaries. (Courtesy M. Sarikaya et al., Met. Trans. 14A, 1121 [1983]).

The utilization of on-line (Fig. 7) dual-phase steel processing allows the design of composite ferrite/lath martensite microstructures whereby the advantages of the martensite phase are optimized, while the less desirable features of this phase are simultaneously mitigated by the presence of the other constituent phase. The size, morphology, distribution, shape and volume fraction of the martensite phase critically control the mechanical properties, especially fracture and fatigue, of such steels. As a result, these structures offer a degree of metallurgical flexibility that is absent in conventionally processed alloys. Low carbon martensite/ferrite structures give high strength, cold formability, improved low temperature ductility, and improved corrosion resistance in concrete [6]. Their greatly improved corrosion resistance may permit the elimination of coatings, with considerable potential for cost savings. These steels can be

(a) Scheme to show “entrapment” of stabilizeduntransformed austenite between laths in the martensitepackets

γγγγ γγγγ LathMs

Solute enrichedaustenite stabilizedby solutes/strain

Martensitedislocations sweep

up solutes

(b) “Final State”

Dislocated γ/α Interface with high % solutes

LathMs

γγγγ

produced on line in a hot mill by controlled rolling and quenching, as shown schematically in Fig. 7. One of the practical problems to be solved in a hot mill using controlled rolling is that rapid quenching is needed [9] so as to convert the austenite phase into the desired lath martensite. Knowledge of the transformation kinetics is thus needed.

At a recent ASM Conference on Heat Treatment we reported [10] for the first time the evaluation of the TTT and CCT kinetics for dual-phase steels (2%Si-0.08%C) and these results demonstrated that cold water quenching from the finish rolling temperatures (in the α+γ range) is sufficient to transform γ into α′ for material of thickness 12 mm or larger, as has been shown commercially [11]. Such practice is clearly economical since no further heat treatment is needed. Thus, steels containing lath martensite have attractive possibilities for future development.

Fig. 6 BF/DF TEM at 800 kV showing in situ deformation to fracture of lath martensitic steel after tempering at 380° C; cracks follow laths.

Fig. 7 Schematic showing on-line rolling (finish in the α+γ field) and quenching to obtain “dual-phase” steels. Actually the structure is 3-phase: α, α’ and γ (see Fig. 1).

Acknowledgments

This work was supported by the Director, Office of Energy Research, Office of Basic Energy Sciences, Materials Sciences Division of the U.S. Department of Energy under Contract No. DE-AC03-76SF00098. I also acknowledge the enormous contribution made by my graduate students, post-docs and research colleagues towards the understanding and utilization of packet, lath, martensites.

References

1. G. Krauss and A. R. Marder, Met. Trans., 2, 2243-2257 (1971).

2. J. McMahon and G. Thomas, in Proc. Third Int. Conf. on Strength of Metals and Alloys, The University of Cambridge (U.K.), 1973, pp. 180-184.

3. E.g., see the papers by G. Thomas et al., in Proc. Int. Conf. on Solid-Solid Phase Transformations, H. J. Aronson, ed., ASM, 1981, pp. 999, 1421.

4. G. Thomas, Iron & Steel Int., 46, 451-461 (1973). 5. G. Thomas, in Frontiers of Materials Technology, M.

A. Meyers and O. T. Inal, eds., Elsevier, 1985, pp. 89-100.

6. G. Thomas, Trans. Indian Inst. Metals, 49, 127-132 (1996).

7. S. J. Barnard, G.D.W. Smith, M. Sarikaya and G. Thomas, Scripta Metall., 15, 387-393 (1981).

8. M. Sarikaya, A. K. Jhingan and G. Thomas, Met. Trans., 14A, 1121-1132 (1983).

9. G. Thomas in Physical Metallurgy of Direct Quenched Steels, K. A. Taylor, S. W. Thompson and F. B. Fletcher, eds., TMS, 1993, pp. 265-278.

10. G. Kusinski and G. Thomas, in Proc. 1st Auto-motive Heat Treating Conf., Puerta Vallarta, Mexico, R. Colas, K. Funatani and C. A. Stickles, eds., ASM International, 1999, pp. 17-25.

11. Private communications: UCIN, Spain and Mukhand Corporation, India.

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