The physical metallurgy of duplex stainless steels J.-O. Nilsson and G. Chai Sandvik Materials Technology, R&D Centre, S-81181 Sandviken, Sweden

Abstract. Duplex stainless steels are unrivalled in a large number of applications in the temperature range -50°C to 250°C

where a combination of corrosion resistance and mechanical strength is required. The key to the unique properties is the duplex

structure and the synergistic interaction between the two phases. However, the duplex structure also renders them inherently

sensitive to phase transformations that may lead to reduced toughness and/or corrosion problems. Correct production, heat

treatment and welding require a thorough knowledge of the relationship between micro structural phenomena and properties.

Consequently, it is essential for steel producers to convey this knowledge to manufacturers and end-users. Much attention in his

paper is paid to pathological phenomena in duplex steels such as the precipitation of undesirable phases of various types. In

conjunction with this, the conditions for avoiding these phase transformations are identified. It is therefore hoped that this review

paper provides a platform for the appropriate use of this family of steels.

1 Introduction The characteristic features of duplex stainless steels (DSS) result from an interaction between ferrite and austenite. While ferrite is rather brittle it merits from being resistant to stress corrosion cracking. As opposed to ferrite, austenite is inherently tough but suffers from being very prone to cracking under stress corrosion conditions. Furthermore, austenite work hardens much more than ferrite implying that the load sharing between austenite and ferrite varies with the degree of plastic deformation. The properties of early DSS were interpreted as a result of the relative contributions of ferrite and austenite in terms of a linear law of mixture. It is now apparent that several features of DSS result from a rather complex interaction between the phases1,2,3,4. In fact, there are synergetic effects that lead to corrosion resistance and mechanical behaviour that cannot be predicted from the properties of the constituents alone. A significant part of this paper will be devoted to this synergism. The paper also addresses the precautions that have to be taken in producing, fabricating, and manufacturing various types of components of DSS. Therefore, the pathological aspects of DSS will be analyzed in some detail. Since the advent of the first duplex alloys 453E (26Cr-4Ni) and 453S (26Cr-5Ni-1.5Mo) launched by Avesta Jernverk in 1930 considerable efforts have been made to improve and optimize DSS. The first alloys were far from optimized in terms of corrosion resistance as the corrosion resistance in the two phases could be very different, therefore leading to selective corrosion of one phase. Furthermore, the fraction of ferrite was quite high leading to moderate toughness. DSS of today contain ferrite and austenite in almost equal fractions. Moreover, the resistance to pitting corrosion is close to equal in ferrite and austenite. Two innovations have contributed to the optimization of DSS. Firstly, process metallurgical techniques have been developed allowing the precise control of nitrogen up to the solubility limit. Secondly, a deeper understanding of the thermodymic interplay between alloying elements has led to the development of computer programs that are invaluable implements in alloy development.

The first DSS to be optimized using Thermocalc was SAF 2507 in which equal pitting resistance in the two phases was obtained by calculating the solution heat treatment temperature at which equal PRE-values was achieved. Such methods have revolutionized alloy development and are now used routinely among steel producers. Rather recently multi scale modelling has been employed to model the plastic deformation of two-phase materials. DSS provide an interesting challenge as ferrite and austenite show very different plastic behaviour. Ferrite is often the strongest component during the beginning stage of deformation but, due to work hardening in austenite, this phase shares gradually more load as deformation proceeds. There is, therefore, a complex interplay between the two phases during plastic deformation. Multi scale modelling enables an investigation of this interaction in detail, thereby offering a tool for optimizing DSS with respect to plastic deformation. The role of techniques based on multi scale modelling and thermodynamics in obtaining a deeper understanding of DSS will be discussed in the paper. Moreover, results from modern micro analytical techniques in characterizing the microstructure down to the atomic level will be discussed with reference to alloy development.

2 Microstructure-Characterization 2.1 Phase diagrams There are many possible representations of the Fe-Cr-Ni system. Phase diagrams may today be computed on a routine basis using computer based thermodynamics. It is instructive to take an example from a specific alloy for a more general discussion about phase transformations of DSS. The details may vary from one alloy to another but the principles are the same. The usefulness of a phase diagram is illustrated by an isopleths diagram computed using Thermocalc5 for Cr-7Ni-4Mo-0.28N using chromium as a variable (Fig 1). The composition of SAF 2507 is indicated by a broken vertical line. The steel solidifies ferritically after which austenite is formed until, below about 1300°C, an entirely duplex structure is formed. As the phase fractions of ferrite and austenite may be calculated the relative fractions can be controlled by selecting the appropriate heat treatment temperature (see Fig 2). Moreover, the laws of thermodynamics can be used to control the heat treatment so as to give equal pitting resistance in ferrite and austenite. In fact, SAF 2507 was the first DSS to be designed and optimized using thermodynamics. Computer-based thermodynamics is now a standard tool employed by steel-makers in designing new alloys.

Fig 1. Phase diagram of SAF 2507 computed using the thermodynamic computer program Thermocalc. The composition of SAF 2507 is indicated by a dashed line.

Fig 2. Phase fractions of ferrite, austenite, σ-phase, χ-phase and Cr2N as a function of temperature in SAF 2507. Computed using Thermocalc.

The ideal structure of a DSS is of course a purely ferritic-austenitic structure. However, in practice, this is quite difficult to achieve. In particular, this is so when the dimensions are large or during welding when the cooling rate cannot be controlled fully. If the DSS is a highly alloyed one the situation is aggravated. The cooling path down to room temperature through phase fields in the equilibrium diagram is perilous. Various obstacles in terms of phase transformations have to be by-passed. As shown in Fig 1 phases such as σ-phase, Cr2N and χ-phase are thermodynamically stable at temperatures below about 1000°C. The list of secondary phases observed in DSS is much longer but Thermocalc and other programs of the same kind are unable to predict all of them as they are not sufficiently well described thermodynamically. It is even quite likely that some of the listed phases are non-equilibrium phases. For instance, R-phase is most frequently found in the early stages of precipitation at around 700°C and is rarely found after longer ageing times. It is therefore suggested that precipitation of R-phase is energetically favourable because of the well-defined orientation relation with the ferritic matrix ⟨0001⟩R//⟨111⟩α after which σ-phase gradually becomes the dominating phase due to a lower free energy. The list below (Table 1) is a summary of phases observed experimentally in DSS. Of these, M23C6 is of limited interest owing to the paucity of carbon in modern DSS and, hence, absence of carbides. However, copper is used as an alloying element in certain DSS to improve the mechanical strength and machinability. Also improved corrosion resistance in sulphuric acid solutions have been observed6. As the solubility of copper decreases rapidly below about 700°C copper precipitates are often formed in the supersaturated ferrite. This has been shown to refine the structure by promoting the nucleation of austenite7.

Table 1. List of phases observed experimentally in DSS

Type of

precipitate

Nominal

chemical

formula

Temperature

range of

stability, °C

Space group Lattice parameter,

nm

Reference

number

Austenite (γ) - - mFm3 0.358-0.362

Ferrite (α) - - m3Im 0.286-0.288

σ-phase Fe-Cr-Mo 600-1000 P42/mnm a=0.879, c=0.454 8 Chromium

nitride Cr2N 700-900 mP 13 a=0.480, c=0.447 9

χ-phase Fe36Cr12Mo10 700-900 mI 34 a=0.892 10 R-phase Fe-Cr-Mo 550-700 3R a=1.090, c=1.934 11 π-phase Fe7Mo13N4 550-600 P4132 12

τ-phase Not determined 550-650 Fmmm a=0.405, b=0.484, c=0.286

13

ε−Cu Cu 300-550 mFm3 a=0.361 M23C6 - 950-1050 mFm3 a=1.056-1.065 14

However, σ-phase and Cr2N play an increasingly important role as the risk of formation becomes higher in modern highly alloyed DSS. Adding molybdenum, chromium and nitrogen is a mixed blessing as these elements increase the resistance to pitting corrosion but, at the same time, create a material that is increasingly difficult to produce. There are indeed fundamental limits as to how much molybdenum, chromium and nitrogen that can be added to DSS without forming Cr2N and σ-phase as the corresponding phase fields will expand. However, it is quite likely that the production facilities and the billet dimensions will be limiting in practice as the production window in a space defined by heat treatment temperature and cooling rate shrinks as more molybdenum, chromium and nitrogen is added (See Fig 3). This is a Scylla-and-Charybdis type of situation. While σ-phase (and other intermetallic phases) tends to be a problem during insufficient cooling chromium nitrides behave in an opposite way. In particular, non-equilibrium intragranular nitrides (different from the ones indicated in the TTT-diagram in Fig 6) are likely to form when the cooling is so rapid that there is insufficient time for nitrogen to escape from ferrite. Such a situation may occur during welding followed by cooling when the ferrite eventually becomes supersaturated with nitrogen and the diffusion distance shrinks and becomes much smaller than the dimension of the ferrite grains. The example given in Fig 4 shows intragranular non-equilibrium nitrides formed in SAF 2507 heat treated at 1300°C for 20 min and subsequently quenched to ambient temperature in

water. It is quite evident in this micrograph that the number density of nitrides diminishes close to the phase boundary. This is expected as the nitrogen in a narrow zone close to the phase boundary has had sufficient time to escape to the adjacent austenite whereas the nitrogen in the grain interior is trapped. For comparison chromium nitrides formed during isothermal ageing at 850°C are shown in Fig 5. As shown here, isothermally formed nitrides are often associated with a front of secondary austenite advancing from the primary austenite between the pinning nitrides.

Fig 3. Schematic diagram showing the conditions under which

σ-phase and chromium nitrides are formed in a space defined by temperature and cooling rate. While σ-phase tend to form when

the cooling is slow Cr2N are likely when the cooling is rapid.

Fig 4. Ferrite grain in which chromium nitrides (arrowed) have formed during water-quenching from 1300°C. SAF 2507 in

SEM. Note that there is a zone depleted of chromium nitrides at the phase boundary (Sandvik archive).

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700

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900

1000

0,001 0,01 0,1 1 10 100

Ageing time, h

Temperature, °C

R-phase

σ-phase

χ-phase

Cr2N

475°C-

embrittlement

Fig 5. Secondary austenite formed at an austenite/ferrite phase

boundary in SAF 2507 after ageing for 3 min at 900°C. The advancing phase boundary is pinned by chromium nitrides, which have formed simultaneously with secondary austenite

(Sandvik archive).

Fig 6. Experimentally determined TTT-diagram of SAF 2507 showing the C-curves of Cr2N, χ-phase, σ-phase, R-phase and

spinodal decomposition. The C-curve of R-phase is dashed as it is supposed to be a precursor of the more stable σ-phase3.

2.2 Intermetallic phase and chromium nitrides The following is an example of the sometimes rather complicated interplay between secondary phases in DSS. The early formation of χ-phase in the temperature range 800-900°C turns out to be a precursor of σ-phase. This is indicated in the TTT-diagram shown in Fig 6 representing SAF 2507. Ferrite and χ-phase are both cubic with lattice parameters differing by almost exactly a factor of 3. As a consequence, χ shows a cube-cube orientation relationship with the ferritic matrix in combination with small coherency strains and, therefore, nucleates very easily. This is reflected in the diffraction pattern shown in Fig 7. As opposed to χ−phase σ-phase, with a tetragonal crystal structure, nucleates with

some difficulty. Based on observations in SAF 2507 it seems very probable that σ-phase is often the most stable phase thermodynamically as χ-phase is sometimes absent after long-term ageing. However, χ apparently serves as a nucleation site for σ-phase (Fig 8), thereby enhancing the kinetics of the precipitation process. It so happens that tungsten favours the formation of χ-phase15. The faster kinetics of intermetallic phase (χ and σ) formation observed in tungsten-alloyed DSS is therefore in part explicable in terms of the catalytic effect of χ-phase. Three factors contribute to this; the thermodynamic effects on the free energy of formation16, the effect of χ-phase on nucleation, and the fact that the increased stability range allow intermetallic phases to form at higher temperatures where the diffusion is faster. This has practical implications, in particular during welding. Experience has shown that the use of tungsten-alloyed filler metals requires lower heat inputs to avoid intermetallic phase formation leading to more passes during multi-pass welding and, consequently, lower productivity17.

Fig 7. Electron diffraction pattern obtained from χ-phase showing the cube-cube ⟨001⟩χ//⟨001⟩α orientation relation with ferrite and

the lattice parameter mismatch of about 33%. TEM (Sandvik archive)

Fig 8. Nucleation of σ-phase is facilitated in the presence of χ-phase. The example shows σ-phase in a phase boundary between ferrite (right), austenite (left) and a χ-phase precipitate (Sandvik

archive) 2.3 Secondary austenite The austenitic component in DSS may be called primary austenite as it is formed immediately after the ferrite has solidified. However, austenite may also be formed inadvertently at relatively low temperature, i. e. after the duplex structure has been established. Typical conditions are those prevailing during multipass welding when reheating of underlying weld beads may lead to formation of so called secondary austenite. Isothermal ageing at a temperature below the duplex phase field in the equilibrium diagram (if such a situation ever occurs in practice) may also lead to secondary austenite formation. As shown in Fig 2 the equilibrium phase fraction of austenite in SAF 2507 increases from the onset of austenite formation at about 1300°C to about 1000°C. The principles are similar also for other DSS but are shown here with a specific example for simplicity. Consider a DSS with an already established duplex structure with given volume fractions. If subjected to a temperature where the equilibrium volume fraction of austenite is higher the DSS will experience a driving force for forming additional austenite, i. e. secondary austenite. At least two morphologies of secondary austenite may be identified; Intergranular as in Fig 9 and Intragranular as in Fig 5. The latter is the result of isothermal ageing while the former may occur during multipass welding as a result of repeated heating. 2.4 Spinodal decomposition

Below about 500°C there is a miscibility gap in the iron-chromium phase diagram (Fig 10). As a consequence ferrite decomposes into two ferritic phases of slightly different lattice parameters. The phenomenon is termed spinodal

decomposition in academic research18. It is also called 475°C-embrittlement as embrittlement ensues, a circumstance well known to steel-makers. Similar observations in other systems had puzzled scientists for a long time but it was not until 1961 when Hillert19 published his seminal work on spinodal decomposition that it was explained theoretically. The free energy as a function of composition exhibits two minima, implying that the system can gain energy by decomposing into iron-rich and chromium-rich ferrite. This requires so called uphill diffusion, i. e. diffusion against the concentration gradient. As opposed to classical precipitation spinodal decomposition involves no nucleation stage. In practice, the phenomenon limits the long-term use of DSS to temperatures below about 250°C. Above this temperature embrittlement takes place, the higher the temperature the higher the rate of embrittlement. The subject has been reviewed in depth by Guttmann20. An example representing experimental data obtained from investigations of SAF 2507 base material is shown in Fig 11. Each dot represents an individual test. Red dots (blue dots) represent specimens with a toughness value below (above) 27J, which is regarded as a critical value. As an example, SAF 2507 shows severe embrittlement at 350°C after about 1 year of ageing as seen in Fig 11. In this context it should be pointed out that, although spinodal decomposition in ferrite leads to brittleness as measured in an impact test, the ductility in a tensile test performed at a much lower strain rate is often satisfactory. This implies that the material is fully applicable provided shock loads can be eliminated. The brittleness is related to the increase in hardness of the ferrite, which, in turn, is associated with the strain fields generated by the phase separation. Spinodal decomposition may be influenced by composition. It has been shown by Mössbauer spectroscopy that copper and nickel enhance the formation of chromium-rich ferrite21, a fact taken as evidence that decomposition is enhanced. However the effect seems to be rather weak. Also molybdenum seems to have an accelerating effect on the decomposition as shown by field ion microscopy investigations22. It would be highly desirable to find an alloying element capable of suppressing the decomposition significantly. This would raise the allowable service temperature of DSS and open up new applications. However, up to date there seems to be no such element, at least not among the most common alloying elements.

Fig 9. Structure of as-welded 25Cr-10Ni-4Mo showing

Widmannstätten type of secondary austenite in a reheated weld bead. Multipass welding using GTAW. SEM23.

Fig 10. Phase diagram of the Fe-Cr system showing a miscibility gap below about 500°C. Computed using Thermocalc.

It has been shown quite recently that spinodal decomposition may be influenced also by plastic deformation24. Immediately outside the spinodal line phase separation in ferrite may be changed by cold work from nucleation and growth to one that is truly spinodal in nature. It has been suggested that this is due to dislocation stress fields interacting

with the strain fields generated by decomposition. The micrographs shown in Figs 12 a and b are examples of decomposition in SAF 2507. Undeformed material shows classical nucleation and growth at 500°C while spinodal decomposition takes place in material plastically deformed to 15% strain and subsequently heat treated at 500°C. The mottled contrast arising from phase separation becomes visible when and only when the ferrite crystal is oriented in the ⟨001⟩-direction.

200

300

400

500

600

700

0,1 1 10 100 1000 10000 100000

Ageing time, h

Ageing temperature, °C

< < < < 27J

> 27J

Fig 11. Impact toughness of SAF 2507 base material as a function of temperature and ageing

time. Red dots represent unacceptably low toughness values below 27J. The nose of the C-curve appears at about 450°C25.

Fig 12a. SAF 2507 in undeformed condition aged for 3h at 500°C. Chromium rich particles have formed by classical

nucleation and growth24

Fig 12b. SAF 2507 plastically deformed to 15% strain aged for 48h at 500°C. Mottled contrast visible when viewed along (001) is indicative of spinodal decomposition24.The deformation has

caused a transition from classical precipitation to spinodal decomposition24

The phase separation may also be imaged directly using atom probe field ion microscopy in which three-dimensional images may be produced (3-DAP). In this instrument individual atoms on the specimen surface are imaged and analyzed in a time-of-flight spectrometer. Analyzing adjacent layers successively enables the production of three-dimensional images, an example of which is shown in Fig 13. In this image chromium is shown in blue and iron in green contrast.

Fig 13. Three-dimensional image obtained in the atom probe field ion microscope showing phase separation in the ferrite into

chromium-rich (blue) and iron-rich (green) regions. The material is a 28Cr-7Ni-2Mo-0.37N DSS weld metal aged for 243h at 450°C (Courtesy M. Hättestrand, Sandvik)

In summary it should be emphasized that there seems to be no efficient means of suppressing spinodal decomposition in DSS. However, the process can be enhanced by adding copper, nickel and molybdenum. There is also evidence that it is favoured by plastic deformation. From this we may conclude that the concentration of copper, nickel and molybdenum should be kept low and cold work should be avoided.

3 Micro mechanical behaviour in duplex stainless steels DSS is usually described as a mixture of a hard ferrite and a soft austenite. However, its strength is even higher than that of the corresponding ferritic material as shown in Fig 14. High strength of DSS is generally attributed to their fine grain structure. The synergetic effect is also an important factor, but has been less discussed in the literature.

0

200

400

600

800

1000

0 10 20 30 40 50 60

Strain, %

Stress, MPa

Austenite, 316L

Ferrite, 4C54

Duplex, SAF 2205

Fig 14. Comparison of the tensile behaviour of DSS, ferritic and austentic

stainless steels.

3.1 Micro yielding behaviour In DSS containing almost equal amounts of austenite and ferrite inhomogeneities on a micro-scale play an important role. This may lead to micro selective behaviour under service26. The different chemical and physical properties of austenite and ferrite have also led to complexities in materials design and development. Using modern computer technology, the pitting corrosion resistance in duplex stainless steels can be optimised by calculating the conditions under which equal PRE values are obtained27. Compared with this successful achievement, the work on the optimization of SCC resistance and strength of duplex stainless steels is quite limited, but some exploratory studies have been conducted recently using advanced in-situ diffraction measurements and multiscale material modelling28,29,30,31. Since the load sharing between the individual phases during loading is different owing to differences in the modulus of elasticity and deformation hardening rate of the individual phases, strong inter-phase reactions will also result in the formation of microstresses that maintain their equilibrium among subsets of grains of different orientations29. These residual micro stresses can have great effects on SCC, yielding and damage of the material, and, consequently, affect their strength, deformation and fracture behaviour26,28. Understanding the micromechanical reactions is therefore important for successful application of DSS, alloy design and materials development. Recently, some in-situ diffraction methods using X-ray, synchrotron and neutron radiations have been developed to analyze the load sharing, stress interaction between phases and consequently the micro stress-strain curves of the austenite and ferrite in DSS29,30,31. Fig 15 shows the setup of an in-situ X-ray diffractometer used to measure the stress versus strain behaviour of the individual phases. The in-situ X-ray diffraction experiments were carried out on a Seifert X-ray diffractometer. A compact test rig was used to provide uniaxial loading to the specimen. Two loading-unloading cycles were applied. During each cycle, the tensile stress was increased stepwise from about 0 MPa to a given strain and then decreased in steps to 0 MPa as shown in Fig 16.. At each stress level, while recording the applied load via a built-in load cell on the test rig, the phase specific stresses were measured by X-ray diffraction using the sin2ψ-method30 on the α-211 and γ-220 reflections, respectively.

Fig 15. X-ray equipment for performing the in-situ tests, setup and specimen.

The microyielding behaviour of the individual phases in two phase alloys can be described by the phase-specific stress versus the macroscopic strain curves. Fig 16 shows two examples. Fig 16a shows the macro deformation behaviour of the bulk material and micro deformation behaviour of the individual phases determined by in-situ X-ray diffraction method in Sandvik SAF 2507 in as-delivered condition pre-strained to 0.1% (2507AD)28. Without loading,

the ferritic phase has a small tensile residual stress and the austenitic phase has a compressive stress of 50 MPa. Both phases start with elastic deformation as expected. However, yielding commences in the austenitic phase at an early stage as shown by a small deviation of the phase-specific stress (about 230 MPa), and then a large deviation of the phase-specific stress (about 425 MPa) from a linear distribution. The former phenomenon can be attributed to the first yielding of the crystal plane of the austenitic phase29. Yielding in the ferritic phase starts to occur at a higher load (about 565 MPa). This reflects the fact that the austenitic phase has a higher deformation hardening rate. Yielding in the bulk material appears at a macro-stress of about 545 MPa. The above-mentioned phenomenon was also confirmed by an in-

situ neutron diffraction method under a compression test29. The results of the in-situ neutron diffraction measurements also show that the yielding at the crystal planes of the individual phases is selective. Fig 16b shows the macro deformation behaviour of the bulk material and micro deformation behaviour of the individual phases determined by in-situ X-ray diffraction in one SAF 2507 duplex stainless steel in as-delivered condition (2507AD) prestrained 8%. Owing to deformation hardening during pre-straining the austenitic phase now starts to yield at a stress of 805 MPa. However, the ferritic phase starts to yield at a stress of 555 MPa. This means that there is no deformation hardening taking place in the ferritic phase during pre-straining. Moreover, this shows the complexity of the deformation behaviour of a duplex stainless steel due to the stress interaction between two phases. In Fig. 16, we can also note the residual stresses remaining after the deformation cycle. For the material in the as-received condition, the residual stresses in the two phases after the deformation cycle can be totally different from the original ones, and can change from compressive stresses to tensile stresses, and vice versa (Fig 16a). For the heavily strained material the residual stresses in the two phases before and after the loading show only marginal changes (Fig. 16b).

-100

0

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700

0,00 0,20 0,40 0,60 0,80 1,00

Strain (%)

Specific phase stress (MPa)

austenite loading

ferrite loading

DSS loading

unloading

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0

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0,0 0,1 0,2 0,3 0,4 0,5

Strain (%)

Specific phase stress (MPa)

austenite loading

austenite unloading

ferrite loading

ferrite unloading

DSS loading

DSS unloading

Fig 16. Influence of strain on the phase-specific stresses in the austenitic and ferritic phases in Sandvik SAF 2507; (a). Material in as

received condition, (b). Material pre-strained to 8%.

The results in Fig 16 show that modern duplex stainless steels can not be fully optimized from the mechanical behaviour point of review, and also provides an explanation why selective stress corrosion cracking and fatigue damage can occur in duplex stainless steels. These results provide some interesting information regarding the improvement of material quality and alloy design. Table 2 shows a summary of the results of the macro and micro yield strength of Sandvik SAF 2507 duplex stainless steels in three conditions: as received, with 8% pre-straining, and aged at 475°C for 3 hours. It seems that the yield strengths of austenite or ferrite depend not only on chemical composition, but also material state. Ageing at 475°C that has led to the occurrence of spinodal decomposion in the ferritic phase has increased the strength of the ferritic phase, but not the austenitic phase. If equation 1 is used to evaluate the macro yield strength due to the micro yielding, it can be found that the calculated macro yield strengths are lower than the real measured yielding strengths, σbulk, of the bulk materials. The factors are shown in Table 2. In fact, this is a perfect example of synergy, here explained by the contribution of grain boundaries to the yield strength of the bulk materials.

σcal = fασα+fγσγ (1)

Table 2. Results of macro and micro yielding in SAF 2507 duplex stainless steel

Fraction AR 8% CW Aged

(%) (MPa) (MPa) (MPa)

Austenite (σγ) 57.1 425 805 420

Ferrite (σα) 42.9 565 555 675

DSS (σbulk) 545 705 610

(a) (b)

σcal = fασα+fγσγ 485 682 529

factor = σbulk/σcal 1.12 1.03 1.15

3.2 Multiscale material modelling for micro yielding and deformation behaviour In recent years, multi-scale material modelling has attracted much interest from the researchers in the field of material mechanics. This type of modelling offers a possibility to study the behaviour of single phases, single grains and load sharing between the phases in multi-phase materials. The basic idea in multi-scale material modelling is that the a priori homogenized macro- induced scale material model is replaced by the homogenized response of a representative volume element (RVE), a generation of a numerical grain structure from a physical grain structure using RVE (Fig 17)26,28.

Fig 17. Generation of a numerical grain structure (b) from a physical grain structure (a) using Representative Volume Element (RVE) by Voronoi polygonization algorithms.

Multi-scale material modelling uses micro-scale crystal plasticity and continuum models26. The crystallographic directions, unit cells and slip systems are identified. The yield function and shear (Schmid) stress, viscoplasticity and slip-damage, hardening stress, and plastic and damage velocity gradients in the austenitic and ferritic phases are then evaluated. Fig 18 shows the results of multi-scale material modelling for the as received SAF 2507 (2507AD) and the aged SAF 2507 (2507HT). The ferritic phase in the as received SAF 2507 is the stronger phase at a total strain less than about 3% and then becomes the softer phase with increasing strain (Fig. 18a). The ferritic phase in aged SAF 2507 is the stronger phase in the whole strain range used (Fig 18b). These observations are consistent with the experimental observations29,30,31.

0

200

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800

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1200

0,00 0,05 0,10 0,15

True strain

True stress (MPa)

DSS

αααα

γγγγ

0

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1000

1200

0,00 0,05 0,10 0,15

True strain

True stress (MPa) γγγγ

αααα

DSS

(a) (b)

(a) (b)

Fig 18. Stress versus strain curves of super duplex stainless steel by multi-scale material modelling; (a). As received, (b). Aged at

475°C for 3 hours.

4 Some microstructure-property relations 4.1 Stress corrosion cracking Standard austenitic steels are susceptible to stress corrosion cracking (SCC). The problem can be mitigated by additions of nickel (30% or higher) but nickel is expensive and is not employed unless required. DSS are less sensitive to SCC despite the fact that the nickel content even in super DSS is moderate. The time to fracture of two austenitic stainless steels (type 304 and 316) and the lean DSS SAF 2304 in a solution of 45% MgCl2 at 150°C is plotted in Fig 19 as a function of normalised applied stress (applied stress divided by the tensile strength). Not only is the critical stress intensity (based on absolute value of stress) for stress corrosion cracking higher in DSS. The same ranking is also valid if the stress normalized with respect to the tensile strength is used an example of which is given in the redrawn diagram shown in Fig 1932. From this it can be inferred that the resistance to SCC in the presence of chlorides is not only a function of the higher strength of DSS. It is evident that the duplex structure as such and the synergy between the two phases is advantageous in providing resistance to stress corrosion cracking. This is, at least in part, explicable by the small effective grain size and crack branching taking place at phase boundaries (see Fig 20). It has been pointed out that the ranking of DSS with respect to SCC may be somewhat different in other chloride solutions33. These authors conclude that testing in NaCl is a more suitable solution for ranking various DSS because of the good correlation between alloying content and resistance to SCC

Fig 19. Time to failure versus normalized stress under stress

corrosion conditions for SAF 2304 compared with 304 and 316 austenitic steels. Results from constant load tests in aerated

45% MgCl2 at 150°C32.

Fig 20. A crack induced by stress corrosion in SAF 2507 (Sandvik archive)

4.2 Influence of σσσσ-phase on pitting corrosion

Secondary phases may influence the corrosion properties adversely. Although σ-phase is highly undesirable it is worth pointing out that DSS such as SAF 2507 and SAF 2906 can tolerate 0.1-0.5% σ-phase without noticeable effects on the pitting corrosion resistance (Fig 21). However, when the volume percentage reaches a level of about 1% a detrimental influence is observed for both alloys. A percentage of σ-phase below 0.5% is perfectly acceptable provided pitting corrosion is the main concern. It should be pointed out, however, that the impact toughness is much more vulnerable to σ-phase as shown in Fig 29. An important conclusion may be drawn from this; as low volume fractions of σ-phase (secondary phases in general) are difficult to quantify metallographically it is preferred to focus a discussion about

applicability on desired properties rather than absolute volume fractions of intermetallic phase.

50

70

90

110

0,01 0,1 1 10

Volume percentage of sigma phase

CPT

SAF 2507

SAF 2906

Fig 21. Influence of the amount of σ-phase on the critical pitting temperature in SAF 2507 and SAF 2906. The effect of σ-phase

does not become significant until a volume percentage of about 0.5% is exceeded34 4.3 Hydrogen induced SCC Duplex stainless steels (DSS) have been found to suffer from hydrogen stress induced cracking, HISC, in subsea components with a cathodic protection during the last 10 years. Hydrogen induced stress corrosion cracking is a process involving crack initiation and propagation. Crack initiation starts with the formation of hydrogen pores or voids either at the ferritic grain boundaries or the phase boundaries depending on the alloys (Fig 22). This is a process governed by hydrogen diffusion. Due to the crystal disorder at phase boundaries, diffusing hydrogen atoms can easily accumulate there leading to pore and void formation in combination with stress concentrations that further enhance the hydrogen accumulation35. When the hydrostatic pressure or stress concentration in the void reaches a critical value, the void can develop into small cracks, i. e. crack initiation (Fig 22b).

Fig 22. Crack initiation in SAF 2507-L (a) and SAF 2906 (b) during the HISC testing; (a). Pores at ferrite grain boundary; (b). Pores

and small cracks at phase boundaries36 Different crack initiation behaviour in these two duplex alloys, (SAF 2507, 25Cr7Ni4Mo0.3N and SAF 2906, 29Cr6Ni2Mo0.4N), depends on the microstructure and stress conditions. Fig 23 shows EBSD maps of the SAF 2507 (Fig 23a) and SAF 2906 (Fig 23b) bar materials after the HISC testing. SAF 2507 has a high dislocation density as evidenced by the high density of low angle grain boundaries in the ferritic phase and at the phase boundaries, while in SAF 2906 dislocations are more concentrated at the phase boundaries. This indirectly verifies the interactions between the hydrogen atoms and dislocations in these two materials. Since the applied stress in SAF 2906 was higher than the yield strength of the material, the formation of dislocation pile-ups can cause stress concentration at the phase boundaries eventually inducing HISC. This means that HISC can be the result of an interplay between hydrogen, microstructure and stress/strain in the samples.

Fig 23. Phase structure with austenite (red) and ferrite (blue) after the HISC testing. The green lines are low angle grain boundaries (2-10 degrees); (a). SAF 2507, (b) SAF 290636.

4.4 Influence of chemical inhomogenities Inhomogeneous distribution of alloying elements may lead to selective corrosion in the weakest phase. This is the case if the PRE-values differ markedly between ferrite and austenite37. Another, but rather similar situation, may occur if secondary austenite is present, examples of which are shown in Figs 24 and 25. Microanalysis of secondary austenite has shown that the concentration of chromium, molybdenum and nitrogen is low compared with primary austenite leading to a low PRE-value. This observation is supported by thermodynamic calculations38. The use of Thermocalc in modelling the formation of secondary austenite has led to a good physical understanding of this phenomenon. The presence of secondary austenite is undesirable as such regions in the microstructure are known to be sensitive to pitting corrosion attack. This susceptibility is explicable in terms of low concentrations of chromium, molybdenum and

(b) (a)

αααα

αααα αααα

αααα

αααα γγγγ

(a) (b)

αααα

αααα

γγγγ γγγγ

nitrogen in the secondary austenite. The concentration of chromium and molybdenum may be determined by EDX while nitrogen requires either EPMA (electron probe micro analysis) or SIMS (secondary ion mass spectrometry). Fig 24 shows the how secondary austenite has formed in the reheated weld bead visible to the left in the micrograph. The morphology of this austenite is of the Widmannstätten type as shown in Fig 25.

Fig 24. Cluster-like aggregates (arrowed) of secondary austenite formed in the reheated zone of a duplex weld of type 25Cr-10Ni-

4Mo. GTAW, heat input 1.0 kJ/mm and interpass temperature <150°C. Etched in Beraha solution. LOM39

Fig 25. Aggregate of secondary austenite (cf. Fig 24) where the individual secondary austenite grains have been resolved. The

Widmannstätten morphology is clearly visible. SEM39

4.5 Influence of solidification mode The preferred solidification is fully ferritic, followed by austenite formation in the ferritic matrix on further cooling. During welding, a subject treated in some detail by Karlsson40, this leads to a Widmannstätten type of morphology of the primary austenite (Fig 26). However, under certain circumstances the solidification mode is duplex. The value of the parameter Creq/Nieq, where Creq and Nieq are the chromium and nickel equivalents, respectively, seems to be critical in determining the mode of solidification during welding. Above a critical value of this parameter fully ferritic solidification occurs whereas below, a combination of duplex and ferritic solidification is observed. The critical value of this parameter has been found to be 1.841. However this value is only approximate and perhaps also a function of the overall chemical composition as different values have been reported. For instance, the value seems to be higher in super DSS. If the value is close to critical both types of solidification may be observed in the same weld. It has been shown that fully ferritic solidification is to be preferred since the resulting structure then achieves a higher toughness. The adverse effects of mixed mode solidification derive from the enrichment of chromium and molybdenum in narrow ferrite arms and high energy interfaces deviating from the usual Kurdjumov-Sachs orientation both of which enhance σ-phase formation42. This σ-phase precipitates in regions of vermicular ferrite similar in morphology to the ferrite observed in type 304 welds (Fig 27).

Fig 26. Weld metal of type 29Cr-8Ni-2Mo-0.39N in the as-

received condition showing Widmanstätten type of structure. Austenite appears in yellow contrast. LOM42

Fig 27. Weld metal of type 29Cr-8Ni-2Mo-0.39N in the as-received condition showing vermicular type of structure. Austenite appears

in yellow contrast. LOM42

It would be assumed that ferrite in narrow ferrite arms is particularly sensitive to corrosion attack because of the micro structural instability. This effect may be envisaged in Fig 28 in which intermetallic phase has formed preferentially in such regions already after 90s of ageing at 850°C. Note that there is no visible precipitation in the wide ferrite regions. Note also that there are two kinds of intermetallic precipitates; χ-phase and σ-phase. The χ-phase appears in brighter contrast due to the higher concentration of molybdenum. However, the influence of the early precipitation of intermetallic phase on pitting corrosion seems to be only marginal as pointed out by Karlsson et al43. This was explicable in terms of the absence of depleted zones in the ferrite in combination with the narrowness of the depleted zone at the σ-phase/austenite interface. Because of the high concentration of chromium and molybdenum in σ-phase it is highly likely that the depletion zone in the immediate vicinity of precipitates is more vulnerable to corrosion than the precipitates themselves. In other words, the adverse effects of σ-phase precipitation seem to be of an indirect nature.

0

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0 5 10 15 20 25 30 35 40 45

Sigma phase, vol%

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Fig 28. Weld metal of type 25Cr-10Ni-4Mo aged for 90s at 850°C showing preferential precipitation of intermetallic phase in narrow ferrite arms. White arrows show σ-phase and black show χ-phase.

Backscattered electrons in SEM (Sandvik archive)

Fig 29. Evolution of hardness and impact toughness with the amount of σ-phase. Impact toughness is much more sensitive to the

presence of σ-phase than hardness (Sandvik archive)

4.6 Influence of σσσσ-phase on hardness and toughness

Hardness has been used as an indirect measure of σ-phase44. The hardness seems to obey a linear law of mixture, implying that the parameter hardness is relatively insensitive to the presence of σ-phase. Using hardness as an

indirect measure of σ-phase may give misleading results regarding its effect on properties. For instance, 4% σ-phase leads to an impact toughness below 27J but is not likely to increase the hardness significantly (Fig 28). It is therefore recommended that absolute values of the amount of σ-phase determined by image analysis be used to correlate its influence on properties.

5 Future development Although DSS have been produced since the beginning of the 1930’s new DSS emerge continually. The dominating trend in alloy development has been to increase the concentrations of chromium, molybdenum and nitrogen so as to improve the resistance to pitting corrosion. As with all remedies there are side-effects; Chromium and molybdenum both promote the formation of intermetallic phases while nitrogen is an ingredient in nitrides of the type Cr2N. As a consequence, production is becoming increasingly difficult leading to intermetallic phase formation if the cooling rate is too slow and Cr2N in the ferrite if it is too rapid. This is a Scylla-and-Charybdis type of problem; steering clear of the rock of σ-phase means being engulfed by the whirlpool of nitrides. More explicitly, it is quite obvious that the laws of nature impose fundamental limits in alloy development. However, with more sophisticated production equipment the practical limits are continually pushed forward. As an example, recently developed DSS ( e. g. SAF 2707HD and SAF 3207) with a PRE-number close to 50 have been launched, thus confirming that alloys considered visionary during the Beaune conference in 1991 are now real. It is also envisaged that further refinement of DSS will be made by modelling plastic deformation as shown in this paper. Furthermore, high resolution electrochemistry, such as Scanning Kelvin Probe Force Microscopy (SKPFM), recently presented by researchers at the Royal Institute of Technology45, has been shown to be a powerful tool for investigating corrosion on a scale down to 0.1 µm. This has enabled corrosion studies in ferrite, austenite and secondary phases leading to a deeper understanding of the relative importance of these phases during the corrosion process. Owing to the dwindling production window in highly alloyed DSS conventional production is becoming increasingly difficult. One example is when intermetallic phases form in the central part of ingots owing to too slow cooling rate. This particular problem can be circumvented, at least in part, by employing powder metallurgy. However the main problem remains; highly alloyed DSS are sensitive to intermetallic phase formation. Similarly, high concentrations of nitrogen may lead to nitride formation or the formation of pores owing to supersaturation. It is believed that the DSS of tomorrow will be optimized and tailor-made for very specific applications. Multiscale modelling of plastic deformation offers a means of modelling the load-sharing between ferrite and austenite and, consequently, optimizing the mechanical properties. Analogously, localised corrosion investigations may become an implement for optimizing the corrosion properties. Moreover, since σ-phase and chromium nitrides in highly alloyed DSS cannot be avoided entirely, we may be faced with a situation when we have to accept a certain volume fraction of secondary phases. In a discussion between supplier and user about applicability localized corrosion such as SKPFM is a potential technique for quantifying the influence of secondary phases in specific solutions. In fact, it is in the interest of the supplier to discuss the suitability of a certain DSS on the basis of the desired performance in a specific application rather than, which is very common today, being expected to assess the volume fraction of secondary phases. This is supported by the above-mentioned results showing that a small quantity of σ-phase may have little or no influence on pitting corrosion. However, such a discussion requires a dialogue for each specific application which leads to a tendency towards tailor-making of steels for well-defined niches. This paper has been devoted to the relation between micro structural features and macroscopic properties in DSS. The art of steel-making requires a deep physical understanding of phenomena occurring in steels. However, the customer, who is only interested in a material that fulfils his needs and performs satisfactorily in a given application, cares very little about the micro structural constituents. All commercial materials contain flaws to some extent and this is a fact we have to accept, at least if we require a reasonable price. A one-sided discussion about the existence (or non-existence) of a certain microscopical feature may lead to difficulties, in particular, as non-existence is virtually impossible to prove. Establishing a dialogue between supplier and user focusing on macroscopic properties is therefore desirable and appears to be much more constructive than discussions about micro structural details.

Provided they are appropriately treated, DSS are unrivalled in a large number of applications in the temperature range -50°C to 250°C where a combination of corrosion resistance and mechanical strength is required. The key to the unique properties is the duplex structure and the synergistic interaction between the two phases. However, the duplex structure also renders them inherently sensitive to phase transformations that may lead to reduced toughness and/or corrosion problems. Correct production, heat treatment and welding require a thorough knowledge of the relationship between micro structural phenomena and properties. Consequently, it is essential for steel producers to convey this knowledge to manufacturers and end-users. Much attention in his paper has been paid to pathological phenomena in DSS such as the precipitation of undesirable phases of various types. In conjunction with this, the conditions for avoiding these phase transformations have been identified. It is therefore hoped that this review paper provides a platform for the successful use of DSS.

The physical metallurgy of duplex stainless steels

1 J. Charles, Duplex stainless steels ’91, 1, 3-48 (Les Ulis, France, Les Editions de Physique (1991) 2 H.D. Solomon and T.M. Devine, Duplex stainless steels, 693, Materials Park, OH, ASM (1982) 3 J.-O. Nilsson, Mater. Sci. Technol., 8, 685 (1992) 4 R. Gunn, Duplex stainless steels-Microstructure, properties and applications, ISBN 1 85573 318 8, Abington Publishing, Cambridge, UK (1997) 5 B. Sundman, B. Jansson, and J.-O. Andersson, Calphad, 9, 153 (1985) 6 J. Charles and M. Verneau, Duplex Stainless Steel 2000, 63-69, Associazione Italiana di Metallurgia (2000) 7 B. Soylu and R. W. K. Honeycombe, Mater. Sci. Techn., 7, 2, 137-145 (1991) 8 E. O. Hall and S. H. Algie, Metall. Rev., 11, 61 (1966) 9 S. Karlsson, Jernkontorets Ann., 118, 530 (1934) 10 J. Kasper, Acta Metall., 2, 456 (1954) 11 S. Rideout, W. D. Manley, E. L. Kamen, B. S. Lement, and P. A. Beck, Trans. AIME, 191, 872 (1951) 12 D. A. Evans and K. H. Jack, Acta Crystall., 10, 769 (1957) 13 A. Redjaïmia, G. Metauer, and M. Gantois, Duplex stainless steels ’91, 1, 119-126, Les Ulis, France, Les Editions de Physique (1991) 14 A. L. Bowman, G. P. Arnold, E. K. Storms, and N. G. Nereson, Acta Crystall. B, 28, 3102 (1972) 15 S. Hertzman, T. Huhtala, L. Karlsson, J.-O. Nilsson, M. Nilsson, R. Jargelius-Pettersson, and A Wilson, Mater. Sci. Technol., 13, 604-613 (1997) 16 S. Hertzman, R. Pettersson, K. Frisk, and T. Jerwin, Duplex Stainless Steel 2000, 347-354, Associazione Italiana di Metallurgia (2000) 17 C. A. Frye and D. S. Hoyt, Deep Offshore Technology Conference, Rio de Janeiro, Brazil, October 17-19, (2001) 18 J. W. Cahn, Trans AIME, 242, 166 (1968) 19 M. Hillert, Acta Met., 9, 525 (1961) 20 M. Guttmann, Duplex stainless steels ’91, 1, 79-92 (Les Ulis, France, Les Editions de Physique, 1991) 21 H. D. Solomon and L. M. Levinson, Acta Metall., 26, 429 (1978) 22 F. Danoix, S. Chambreland, J. P. Massoud and P. Auger, Duplex stainless steels ’91, 1, 111-118, Les Ulis, France, Les Editions de Physique (1991) 23 T. Huhtala, Diploma work performed at Sandvik (1993) 24 M. Hättestrand, P. Larsson, G. Chai, J.-O. Nilsson and J. Odqvist, Mater. Sci. & Eng. A, 499, 489-492 (2009) 25 A. Wilson, Internal report, Sandvik Materials Technology 26 Lillbacka R, Chai G, Eka M, Liu P, Johnson E, and Runesson S, Acta Mat, 55, p 5359 (2007) 27 Andersson J-O and Ågren J, Appl Phys, 72, p153 (1992) 28 Chai G and Lillbacka R, Key Engineering Materials, 324-325, p 1117 (2006) 29 Jia N, Lin Peng R, Wang Y D, Chai G, Johansson S, Wang G and Liaw P K, Acta Mat, 54, p 3907 (2006) 30 Lin P R, Gibmeir J, Eulert S, Johansson S and Chai G C, Materials Science Forum, 524-525, p 847 (2006) 31 Jia N, Peng R Lin, Chai G C, Johansson S and Wang Y D, J. Material Science and Engineering, A490, p 425 (2008) 32 S. Bernhardsson, Duplex stainless steels ’91, 1, 137-150, Les Ulis, France, Les Editions de Physique (1991) 33 R. F. A. Jargelius-Pettersson, J. Linder, and S. Hertzman, Duplex Stainless Steels 1997, KCI Publishing, 585-596 (1997) 34 M. Haake, Diploma work performed at Sandvik (2007) 35 M. Yoshoka, A. Ueno and H. Kishimoto, J. Soc. Mat. Sci, Japan, V, 52, 660 (2003) 36 G. Chai, S. Ronneteg, U. Kivisäkk, R. Peng Lin, and S. Johansson, Int. Steel Research, 7, 80 (2009) 37 H. Vannevik, J.-O. Nilsson, J. Frodigh and P. Kangas, Trans. ISIJ, 36, no.7, 807-812 (1996) 38 J.-O. Nilsson, L. Karlsson, and J.-O. Andersson, Mater. Sci. Technol., 11, 276 (1995) 39 A. Gregori, Diploma work performed at Sandvik (1997) 40 L Karlsson, Duplex Stainless Steels 1997, KCI Publishing, 43-58 (1997) 41 D. J. Kotecki and T. A. Siewert, Welding J., 71, 171-178 (1992) 42 A. Gregori and J.-O. Nilsson, Metall. Mater. Trans.A, 33A, 1009-1018 (2002) 43 L. Karlsson, S. Rigdahl and S. Pak, Duplex Stainless Steel 2000, 693-702, Associazione Italiana di Metallurgia (2000) 44 H. Okamoto, Application of Stainless Steels ’92, Stockholm, Jernkontoret, 360 (1992) 45 M. Fermenia, J. Pan, and C. Leygraf, Journal of Electrochemistry Society, 149, B187 (2002)

The physical metallurgy of duplex stainless steels