Solid State Phenomena Vols. 172-174 (2011) pp 350-355 Online available since 2011/Jun/30 at www.scientific.net © (2011) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/SSP.172-174.350
Duplex Stainless Steel Microstructural Developments as Model Microstructures for Hot Ductility Investigations Guilhem Martin1, a, Muriel Véron1, b, B. Chéhab2, c, R. Fourmentin2, d, J.D. Mithieux2, e, S.K. Yerra3, f, L. Delannay3, g, T. Pardoen3, h and Y. Bréchet1,i 1
SIMAP (Laboratoire de Science et Ingéniérie des MAtériaux et Procédés), Groupe Physique du Métal, Domaine Universitaire, 1130 rue de la piscine, 38402 Saint Martin D’Hères, France 2
ARCELORMITTAL STAINLESS Research Center BP 15, 62230 Isbergues, France
3
Institute of Mechanics, Materials and Civil Engineering, Université catholique de Louvain, Place Sainte Barbe 2, B-1348 Louvain-la-Neuve, Belgium a
[email protected],
[email protected], c
[email protected],
[email protected], e
[email protected],
[email protected], g
[email protected],
[email protected], i
[email protected]
Keywords: Duplex Stainless Steel, Microstructural Developments, Austenite Morphologies, Hot Ductility
Abstract. Duplex stainless steels (DSS) are alloys made of ferrite and austenite, with a proportion of each phase around 50%. Their main advantage in comparison with other austenitic and ferritic stainless steels is the attractive combination of high strength and corrosion resistance together with good formability and weldability. Unfortunately, DSS often present a poor hot workability. This phenomenon can stem from different factors associated to the balance of the phases, the nature of the interface, the distribution, size and shape of the second phase, and possibly also from difference in rheology between ferrite and austenite. In order to determine the specific influence of phase morphology on the hot-workability of DSS, two austenite morphologies (E: Equiaxed and W: Widmanstätten) with very similar phase ratio have been generated using appropriate heat treatments. It was checked that the latter treatments generate stable microstructures so that subsequent hot mechanical tests are performed on the microstructures of interest. One microstructure consists of a ferritic matrix with austenitic equiaxed islands while the other microstructure is composed of a ferritic matrix with Widmanstätten austenite. The latter morphology corresponds to the morphology observed in as-cast slabs. 1. Introduction The Duplex Stainless Steels (DSS) are defined as a family of stainless steels consisting of a twophase microstructure involving δ-ferrite and γ-austenite. Exceptional combinations of strength and toughness together with good corrosion resistance under critical working conditions designate DSS a suitable alternative to conventional austenitic stainless steels [1]. Such characteristics make these steels widely used in a variety of applications, particularly in petroleum and gas industries as well as in chemical industries. Unfortunately, the relatively poor hot workability of these alloys makes the industrial processing of flat products particularly critical. Cracking of the coils during hot rolling along the edges is frequently reported [2]. As a consequence, additional operations like grinding, discontinuous processing or scraping are often required, leading to increased manufacturing costs. The high temperature deformation behaviour of the DSS depends on several phenomena. Many of these phenomena have already been investigated: phase proportion [3], All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of TTP, www.ttp.net. (Universite de Grenoble UJF INP, Saint Martin d'Heres Cedex, IP 193.48.255.141-25/11/11,13:09:45)
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orientation relationship [4], partitioning of alloying elements, strain partitioning [5-7], static and dynamic recrystallization and recovery of the constituting phases [8, 9]. However it is very seldom that only one parameter is varied thus making the results difficult to interpret. In general, aswrought materials show a better hot ductility than as-cast materials [10]. In order to determine the influence of phase morphology on the hot workability, it is necessary to start from the same metallurgical state, and, hence to generate different austenite morphologies with a steady phase ratio. In this study, two different austenite morphologies are investigated. Particular attention is paid to microstructural evolutions leading to different austenite morphologies starting from the same metallurgical state. Good understanding of nucleation and growth during the δ→γ phase transformation allows identifying appropriate heat treatments 2. Experimental Procedure 2.1 Material A typical duplex stainless steel with the composition shown in Table 1 has been investigated in the present work. This steel was supplied by the company ArcelorMittal in hot rolled condition with an austenite volume fraction about 50 %. In the transverse direction, austenite presents an equiaxed shape whereas austenite is elongated in the rolling direction (see Fig. 1). Table 1 : Chemical composition of the investigated material, mass contents in [%]
A)
Cr
Ni
Mo
Mn
Si
Cu
P
C
N
S
22.21
3.02
1.09
2.92
0.30
0.65
0.03
0.02
0.19
0.0005
B)
Fig. 1. Microstructures in the as-received conditions. White = γ-austenite, Black=δ-ferrite. (A) Transverse direction. (B) Rolling direction. 2.2 Microstructural Characterization Techniques The specimens were ground down on abrasive paper (silicon carbide) to 1200 grit, and polished with a diamond paste down to 3 µm. The final polishing step was performed with a 1 µm alumina powder. Samples were then etched for metallographic observations and characterizations. Two reagents were employed: the Beraha chemical selective etchant (100mL H2O, 30mL HCl and 11,5g K2S2O5) or a 40% aqueous sodium hydroxide electrolytic solution under a tension of 3V. The microstructural observations were carried out using optical microscopy and scanning electron microscopy (SEM, Zeiss Ultra 55). Electron backscattered diffraction (EBSD) was used to determine crystallographic orientation relationships. Prior to the etching, the distribution of Cr, Ni and Mo in the austenite grains was analyzed by electron microprobe analysis. Volume fractions of each phase, lath and grain size distributions were determined by image analysis using Aphelion software. Specimen preparation for EBSD required further steps. After the 1 µm alumina polishing, samples were mechanically polished in colloidal silica with a 0.04 µm particle size for 2 min.
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3. Microstructural development The goal is to develop microstructures of either lath (Widmanstätten microstructure: W) or equiaxed γ-austenite (equiaxed microstructure: E) in a δ-ferrite matrix which is stable at 1050°C, temperature at which hot ductility will be studied. 3.1 Understanding of the as-cast slab microstructure The as-cast microstructure consists of a ferritic matrix with allotriomorphic austenite on the grain boundaries or Widmanstätten austenite laths inside the grains (Fig. 2). The schematic equilibrium phase diagram (Fig. 2) provides qualitative indications about the origin of the microstructure: from the liquid, alloy solidifies into δ-ferrite. During cooling, austenite precipitates by a nucleation and growth mechanism. Allotriomorphic austenite nucleates at existing δ-ferrite grain boundaries. Widmanstätten austenite laths nucleate from the allotriomorphic austenite with a Kurdjumov-Sachs orientation relationship [11]. A)
δprimary γ1-allotriomorphic
δprimary
δprimary
A) D) B)
δprimary
B) γ1-allotriomorphic
C) γ2-Widmanstätten
C)
γ2-Widmanstätten
γ1-allotriomorphic
Fig. 2. Origin of the as-cast slab microstructure. At high temperature DSS is entirely ferritic (A). During cooling, austenite precipitates at existing δ-ferrite grain boundaries: γ1-allotriomorph (B) and inside the δ-ferrite matrix with a lathy morphology: γ1-Widmanstätten (C). Thermal history explains the the as-cast slab microstructure (D).
A)
B) 25
γ
δ γ
wt % Cr
24
δ
γ
δ
γ
23 22 21 20 0
50
100
150
200
250
300
350
Distance (µm)
Fig. 3. Composition profile in wt% Cr. Composition profile across an austenite lath shows that austenite growth is diffusion-controlled with Cr partition between γ and δ (Fig. 3). Characterization of the as-cast slab reveals significant heterogeneity of the austenite content and the lath mean thickness (Fig. 4) along the thickness of the slab.
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The microstructural heterogeneity through the slab thickness can be explained by the process conditions and more specifically by the difference of cooling rate between the skin and the interior of the slab: the faster the cooling rate, the thinner the austenite laths. This microstructure turns out to be too heterogeneous for performing reproducible hot ductility investigations, hence motivating to generate a more homogeneous microstructure. A) D)
C)
B)
50 45
B)
65
60
40
C)
Columnar zone
Transition zone
A)
%γ
Slab skin
D)
35
Heart slab
30
55
25 20
50 15 10
%γ eγ
45
5
Mean austenite lath thickness eγ (µm)
70
0
40 0
10
20
30
40
50
60
70
80
90
100
110
Position in the slab thickness (mm)
Fig. 4. Illustration of the as-cast slab heterogeneity through thickness. Blue = gradient of austenite content and Green = gradient of lath thickness. (A) Slab skin. (B) Columnar zone. (C) Transition zone. (D) Centre of the slab. 3.2 Widmanstätten Microstructure Generation In order to obtain a homogeneous Widmanstätten microstructure at 1050°C, the following heat treatment was used (HTW.1): solution-treatment at 1380°C for 15 minutes in δ single-phase region, cooling to 1050°C in the δ+γ phase region (temperature of the future hot mechanical test), holding for 8 hours at 1050°C, and finally water quenching (WQ) to room temperature (Fig. 5.A). The thermal stability of this lath microstructure was assessed by annealing at 1050°C for 20 minutes. Neither the volume fraction of austenite, nor the lath size evolved significantly (Fig. 5.B). If a specimen follows a similar thermal history with a shorter annealing at 1050°C (HTW.2, see Fig. 5.C), a homogeneous microstructure is obtained but the stability test reveals that the microstructure is not stable. If the heating rate is not sufficiently quick, the driving force for the nucleation of austenite will be huge and consequently a fine secondary austenite nucleates in the δ-ferrite matrix (see encircled regions in Fig. 5.D).
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Microstructure generated
Stability test
1380°C / 15 min
A)
C)
-15°C/min 1050°C / 20 min
1050°C / 20 min Water quenching
As-wrought material
% γ = 49 ± 3 %
% γ = 52 ± 3 %
1380°C / 15 min
B)
D)
-15°C/min 1050°C / 20 min
1050°C / 8h
Water quenching
As-wrought material
eγ = 29 µmv
eγ = 29 µm
% γ = 50 ± 3 %
% γ = 51 ± 3 %
Fig. 5. Schematics of the HTW and resulting microstructures. White = γ, Black = δ. (A) HTW.1. (B) Stability test of HTW.1. (C) HTW.2. (D) Stability test of HTW.2. 3.3 Equiaxed Microstructure Generation To generate an equiaxed microstructure (E) that is thermally stable at 1050°C, from the as-received material, annealing at 1250°C (temperature required to dissolve quickly the fine secondary austenite, chromium nitrides, sigma or chi phase which could be present initially in the as-received condition) was followed by water quenching (WQ). In the transverse direction, an isotropic microstructure with a volume fraction of austenite close to the equilibrium fraction (% γeq= 35%) was obtained (Fig. 6.A). 1250°C / 4h
Water quenching (WQ)
1050°C
1h
%γ = 36 ± 3%
42h
24h
WQ
WQ
WQ
%γ = 51 ± 3%
%γ = 49 ± 3%
%γ = 51 ± 3%
As-received material
A)
B)
C)
D)
Fig. 6. Schematics of the HTE.1 and resulting microstructures. White = γ, Black = δ. (A) 1250°C/4h + WQ. (B) 1250°C/4h + WQ + 1050°C/1h + WQ. (C) 1250°C/4h + WQ + 1050°C/24h +WQ. (D) 1250°C/4h + WQ + 1050°C/42h + WQ. As hot ductility is characterized at 1050 °C, the aim is to generate such equiaxed austenite morphology at 1050°C. First, the HTE.1 (Heat Treatment for generating an Equiaxed austenite) was performed (Fig. 6). The resulting microstructure was satisfactory only after a very long annealing at 1050°C (Fig. 6.D). When annealing at 1050°C was shorter (Fig. 6.B & .C), fine secondary austenite appears inside the δ-ferrite matrix. Increasing the duration of the annealing at 1050°C promotes coarsening which suppresses fine secondary austenite leading to the desired microstructure. However such a treatment is very long: 50h (Fig. 7.A). In order to generate homogeneous equiaxed austenite morphology after only 5h of heat treatment, a slow cooling down from 1250°C up to 1050°C was performed. This promoted growth stages at the expense of nucleation (Fig. 7.A).
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Finally, the stability of the microstructure generated by HTE.2 was assessed by annealing at 1050°C for 20 minutes (Fig. 7.B). Microstructure generated A)
1250°C / 4h -5°C/min 1050°C / 1h Water quenching (WQ)
Stability test B) 1050°C / 20min WQ
As-received material
%γ = 52 ± 3%
%γ = 50 ± 3%
Fig. 7. Schematics of the HTE.2 and resulting microstructures. White = γ, Black = δ. (A) HTE.2. (B) Stability test of HTE.2. 4. Conclusion Good knowledge of phase transformation mechanisms guides the optimization of microstructures. Starting from a given metallurgical state, it is possible to generate a variety of microstructures with different phase morphologies. By controlling temperature, duration of annealing sequences and cooling rate, alloys with a desired phase proportion, and desired size and shape of the microstructural constituents can be produced. The high temperature tearing resistances of the W and E microstructures were determined using the Essential Work of Fracture concept [12]. The E microstructure turns out to be almost twice more resistant to hot ductile tearing than the W microstructure. The origin of this difference could be accounted for several parameters, especially the difference of solute elements partitioning between the two microstructures and the nature of the interfaces (Work in progress). References [1] J. Charles, Steel Research International, 79 (2008) 455. [2] H.Y. Lieu, Y.T. Pan, R.L. Hsieh and W.T. Tsai, Journal of Materials Engineering and Performance, 10 (2001) 231. [3] A. Iza-Mendia and I. Gutierrez, Factors affecting the hot workability of duplex stainless steels, Duplex Stainless Steels 2007, Grado Italy, 2007. [4] A. IzaMendia, A. PinolJuez, I. Gutierrez and J.J. Urcola, 1st International Conference on Ceramic and Metal Matrix Composites (CMMC 96), San Sebastian, Spain, Sep 09-12 1996. [5] C. Pinna, J.H. Beynon, C.M. Sellars and M. Bornert, Conference on mathematical modelling in metal processing and manufacturing, Ottawa, Ontario, Canadian Institute of Mining, 2000. [6] L.E. Hernandez-Castillo, J.H. Beynon, C. Pinna and S. van der Zwaag, Steel Research International, 76 (2005) 137. [7] L. Duprez, B.C. De Cooman and N. Akdut, Zeitschrift Fur Metallkunde, 93 (2002) 236. [8] E. Evangelista, H.J. McQueen, M. Niewczas and M. Cabibbo, Canadian Metallurgical Quarterly, 43 (2004) 339. [9] A. Dehghan-Manshadi and P.D. Hodgson, Journal of Materials Science, 43 (2008) 6272. [10] A. Iza-Mendia, A. Pinol-Juez, J.J. Urcola and I. Gutierrez, Metallurgical and Materials Transactions a-Physical Metallurgy and Materials Science, 29 (1998) 2975. [11] E.F. Monlevade and I.G.S. Falleiros, Hume-Rothery Symposium on Structure and Diffusional Growth Mechanisms of Irrational Interphase Boundaries, Charlotte, NC, Mar 15-17 2004. [12] T. Pardoen, F. Hachez, B. Marchioni, P.H. Blyth and A.G. Atkins, Journal of the Mechanics and Physics of Solids, 52 (2004) 423.
Solid-Solid Phase Transformations in Inorganic Materials 10.4028/www.scientific.net/SSP.172-174
Duplex Stainless Steel Microstructural Developments as Model Microstructures for Hot Ductility Investigations 10.4028/www.scientific.net/SSP.172-174.350 DOI References [2] H.Y. Lieu, Y.T. Pan, R.L. Hsieh and W.T. Tsai, Journal of Materials Engineering and Performance, 10 (2001) 231. doi:10.1361/105994901770344665 [7] L. Duprez, B.C. De Cooman and N. Akdut, Zeitschrift Fur Metallkunde, 93 (2002) 236. doi:10.1007/s11661-002-0026-4 [8] E. Evangelista, H.J. McQueen, M. Niewczas and M. Cabibbo, Canadian Metallurgical Quarterly, 43 (2004) 339. doi:10.1007/s11661-004-0130-8 [9] A. Dehghan-Manshadi and P.D. Hodgson, Journal of Materials Science, 43 (2008) 6272. doi:10.2355/isijinternational.48.208 [12] T. Pardoen, F. Hachez, B. Marchioni, P.H. Blyth and A.G. Atkins, Journal of the Mechanics and Physics of Solids, 2 (2004) 423. doi:10.1016/S0022-5096(03)00087-5
