JOURNAL OF MATERIALS SCIENCE LETTERS 12 (1993) 1144-1147

Sigma-phase formation and its prevention in duplex stainless steels

S. ATAMERT, J.E. KING Department of Materials Science and Metallurgy, University of Cambridge, Pembroke Street, Cambridge CB2 3QZ, UK

Duplex stainless steels (DSS) are gaining popularity for use in corrosive environments because of their excellent stress corrosion cracking resistance in CO2/C1- environments and their high yield strength. The high strength of these steels relies on the presence of a two-phase mixture of austenite and ferrite. Resistance to corrosion is mainly derived from high levels of Cr, Mo and N, but in many cases galvanic coupling is also believed to be beneficial.

There is a growing concern in welding of DSS that slow cooling conditions can result in the formation of various solid-state transformation products in addi- tion to austenite and ferrite. The presence of transformation products such as sigma, nitrides, laves, etc., is detrimental to both the corrosion performance and the mechanical properties. Among these precipitates, sigma-phase is of particular in- terest. It is a thermodynamically stable phase below 900 C and due to its high hardness it reduces toughness significantly, resulting in brittle weld- ments. This problem is generally overcome in the production of plates by quenching from high solu- tion heat-treatment temperatures (e.g. 1150C). However, in welding, continuous cooling is often unavoidable and hence the cooling rate is of primary importance in determining whether sigma-phase forms. Fortunately, the formation of sigma-phase is relatively sluggish because of the large tetragonal unit cell of 32 atoms.

Despite the sluggishness of the transformation, the formation of sigma-phase can be encouraged if cooling conditions and microstructural variations favour nucleation. For example, alloying additions, such as Cr, Mo and W, accelerate the formation of sigma-phase by enlarging the sigma-phase field on the time-temperature-transformation (TTT) dia- gram (Fig. 1). As recently developed super duplex stainless steels (SDSS) are particularly enriched with these elements, sigma-phase formation can be prom- oted. In addition, the presence of the ferrite in austenitic stainless steels has been shown to speed sigma-phase formation, indicating that DSS are, in general, more susceptible to sigma-phase formation than austenitic stainless steels.

This work is a part of a systematic investigation which aims to establish the relationship between the weldment microstructures and properties of DSS [1, 2]. In this study we aimed to establish optimum welding conditions to prevent the formation of sigma-phase in SDSS weldments.

SDSS (Zeron 100) with a composition of (wt %) Fe-24.8Cr-7.3Ni-3.6Mo-0.65W-0.66Cu-0.19Si-

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Effec! of increasing S i~~a ICr, Mo, W "

Time =

Figure 1 Schematic diagram showing the effect of alloying addi- tions on the position of the C-curve for the sigma-phase transformation.

0.7Mn-0.21N was supplied by Weir Material Ser- vices, Manchester. The 12.5 ram-thick plate had been rolled and solution-treated at 1150 C followed by water quenching. Some small plates were fur- nace-cooled to stimulate the formation of the sigma- phase.

Hounsfield-type tensile test specimens were pre- pared from both quenched and furnace-cooled plates and pulled using a 50 kN capacity screw- driven Schenk testing machine at a strain rate of 1 mm min-1. Fracture surfaces were examined using a Camsean $4 operated at 30 kV. The other experi- mental techniques used in this work were described elsewhere [2, 3].

The microstructure of the as-received quenched plate was found to be a mixture of banded austenite particles, aligned parallel to the rolling direction, in a ferrite matrix (Fig. 2a). It is apparent that the quenching rate is sufficient to suppress the formation of other phases such as carbides, nitrides or interme- tallics, as confirmed by transmission electron micro- scopy in [3]. On the other hand, a plate that was cooled in a furnace showed extensive precipitation (etched black) in the ferrite and at ferrite-austenite interfaces (Fig. 2b). Transmission electron micro- scopy (TEM) and diffraction analysis of these particles confirmed that they were sigma-phase which nucleates preferentially at austenite-ferrite and ferrite-ferrite interfaces (Fig. 2c). It has been proposed that the nucleation of sigma-phase at interfaces is favoured because of a crystallographic orientation relationship between the austenite and sigma phases [4]. However, Fig. 2c illustrates that sigma-phase behaves like a grain boundary allotrio- morph covering both 7/7 and 7/6 grain boundaries.

0261-8028 1993 Chapman & Hall

Figure 2 Optical microstructures of (a) as-quenched base plate and (b) furnace-cooled base plate; (c) TEM micrograph showing the precipitation of sigma-phase at grain boundaries (arrows). The diffraction pattern (inset) shows that the grain boundary phase is the sigma-phase; the zone axis of the pattern is [21 1]sigma.

This in turn suggests that nucleation is heterogen- eous in nature and does not depend strongly on the crystallographic orientation relationships between the phases. In addition, the preferential growth of sigma-phase into the ferrite cannot be explained on the basis of crystallographic orientation since, as sigma-phase adopts a particular orientation relation- ship with respect to the austenite, it should also be crystallographically related to the ferrite, within the same grain, because of the Kurdjumov-Sachs orien- tation relationship between the austenite and ferrite phases. It has also been claimed that preferential growth of the sigma-phase into the ferrite is mainly a result of the higher Cr and Mo concentrations in the ferrite matrix [4]. A-fundamental reason why the sigma-phase preferentially grows into the ferrite is that the ferrite is thermodynamically metastable at temperatures where the sigma phase precipitates so that the ferrite phase should decompose into an equilibrium state. The evidence for this comes from recent investigations which have shown that an equilibrium structure is a mixture of sigma and austenite phases in the temperature range 650-900 C [5]. It should be noted that the composi- tions of the ferrite and sigma phases are very similar, indicating that growth of sigma-phase does not require long-range diffusion of substitutional alloy- ing elements, such as Cr and Mo. The growth of sigma-phase is likely to be controlled by small-scale atomic rearrangement of b c c atoms into a tetra- gonal lattice in a manner similar to the in situ transformation of one carbide phase into another [6]. The absence of any intragranular precipitation of sigma-phase confirms that nucleation is hetero- geneous and a rate-controlling step is the nucleation stage as suggested previously [7].

Although long-range diffusion of Cr and Mo may not be essential for the nucleation and growth of sigma-phase, small compositional variations can influence the formation of sigma. For example, sigma-phase grows preferentially near grain bound- ary austenite allotriomorphs rather than intragranu- lar austenite in the weld metal, as shown in Fig. 3 [5]. This is probably a result of an extensive partitioning of Cr and Mo into the ferrite phase near

Figure 3 An optical micrograph of a weld metal aged 8 days at 800 C. TOe dark etched phase is ferrite, lightly etched phase is austenite and sigma-phase appears light. (Original source of the micrograph [7]).

the grain boundary austenite, consistent with recent observations [3].

The effect of sigma-phase is of considerable interest since its presence has been associated with the degradation of mechanical properties. Charpy room-temperature impact energy for the microstruc- ture in Fig. 2b is found to be as low as 7 J. Tensile test data agree with impact tests results, showing that the presence of sigma-phase reduces the total elongation to failure from 40% in the base plate to 7%. Both yield strength and ultimate tensile strength showed a small increase in the presence of sigma-phase. The fracture path in samples contain- ing sigma-phase appears to involve both fracture of brittle sigma-phase and interfacial failure as opposed to a typical ductile failure in an as-quenched sample (Fig. 4).

Three-dimensional Rosenthal heat-flow equations were used in this study to calculate thermal cycles as a function of the welding conditions. The original equations were given elsewhere [2] and are not presented here. In order to calculate a critical cooling rate to prevent sigma-phase formation, knowledge of the TTT diagrams for sigma-phase formation is essential. This information is already available in the literature for low-grade, super duplex and super ferritic stainless steels [4, 8, 9]. The results are overall in agreement that the kinetics of

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sigma-phase formation from ferrite follows classic C-curve behaviour. The position of the nose of the C-curve varies with the alloy composition as shown schematically in Fig. 1. As emphasized above, higher Cr and Mo concentrations push the C-curve towards the left-hand side, resulting in the relatively fast formation of sigma-phase. Therefore, calcula- tions were carried out for both super ferritic stainless steel and SDSS. The composition of the super ferritic stainless steel represents the ferrite phase which is enriched with Cr and Mo (29Cr and 4Mo) as a result of an excessive partitioning of alloying elements during solid-state transformation in high- heat-input welding. The positions of the nose of the C-curves of the super ferritic and duplex stainless steels are 890 C and 2.5 rain, and 900 C and 70 s respectively [4, 8, 9]. The results of the calculations are shown in Fig. 5 for three different preheat temperatures. As expected, the sigma-phase-free region depends strongly on the heat input and the preheat temperature. An approximately 0.3 kJ mm -1 lower heat input is necessary to prevent sigma-phase formation in ferrite which is enriched with respect to Cr and Mo additions.

t E 4 E

,~3 "-I

C

4 .1

~2 "r

' i i i ' i

Sigma-phase / possible / / _ =~

(a) / / Sigma-fre l ==> e region

I I , I , , , , I , , ,

0.2 0.4 0.6 0.8 Plate thickness (cm)

4.0 7 E E 3.5

"3

~_. 3.o +a

~2.5

.t.a

r~ 2.0 7"

1.5

1.0

1 " ' ' I I I ' " '

Sigma-phase poss, ble / / .

/ / Sigma'free / / region

(b) ~ ~ , I , , I , , , I , , , I . . . . 0.2 0.4 0.6 0.8

Plate thickness (cm)

3.5 ~- , ' I ' ' I ' ' ~ ' ' '

/ / _~ 2.5 S igma-phase

possible e_ 2.0 t - . _ '- /~- Sigma-free

I

1.0 (c) 0.2 0.4 0.6 0.8

Plate thickness (cm) Figure 5 Plate thickness versus heat input for three different interpass temperatures: (a) 25 C, (b) 100C and (c) 200C, showing sigma-free and sensitized regions for both (II~) SDSS (25Cr, 3.7Mo) and (A) SFSS (29Cr, 4Mo).

Figure 4 SEM micrographs of tensile fracture surfaces of parent metal: (a) furnace-cooled sample showing both fracture of the sigma-phase (arrows) and intergranular failure and (b) ductile failure in as-quenched sample.

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Acknowledgements The authors thank British Gas plc, SERC and the Fellowship of Engineering for funding and the staff at ERS (British Gas) for helpful discussions. Special thanks to K. Prosser (ERS) who initiated this particular study. Thanks are due to Professor C. J. Humphreys for the provision of laboratory facilities at the University of Cambridge. The authors are grateful to Dr N. I. A. Haddad and Dr H. K.~D. H. Bhadeshia for allowing the publication of Fig. 3 .

References 1. s. ATAMERT and J. E. KING, Z. Metallkde 82 (1991) 230. 2. S. ATAMERT, R. C. REED and J. E. KING, Mater. Sci.

Technol. in press. 3. S. ATAMERT and J. E. KING, Acta Metall. Mater. 39

(1991) 273. 4. A. J . STRUTT, G. W. LORIMER, C. V. ROSCOE and K.

J. GRADWELL, in Proceedings of Duplex Stainless Steels 86, October 1986, p. 310.

5. N . I .A . HADDAD, PhD thesis, University of Cambridge (1989).

6. R .W.K . HONEYCOMBE and A. K. SEAL, JISI 188 (1958) 9.

7. J. M. VITEK and S. A. DAVID, Weld. Res. Suppl. 65 (1986) 106-s.

8. J. CHARLES, in Proceedings of the Duplex Stainless Steels 91, October 1991, edited by J. Charles and S. Bernhardsson, p. 3.

9. G. HERBSLEB and P. SCHWABB, in Proceedings of the Duplex Stainless Steels 82, October 1982, edited by R. A, Lula, p. 15.

Received 24 August and accepted 23 November 1992

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