F03-9 RECENT DEVELOPMENT OF DUPLEX STAINLESS STEELS J.-O. Nilsson, G. Chai, U. Kivisäkk R&D Centre, Sandvik Materials Technology, Sweden Abstract Recent development of duplex stainless steels is described. The advent of SAF 2707 HD, a 27Cr-7Ni-5Mo-0.4N duplex steel, shows that it is possible to reach a PRE-value close to 50 without sacrificing the fabricability. Modern methods for simulating the interaction between ferrite and austenite intimates that the steels of tomorrow may be optimized with respect to mechanical as well as corrosion properties. Methods under development presented here are multi-scale modelling of plastic deformation and high resolution electrochemical techniques. Introduction Duplex stainless steels (DSS) were first described by Bain and Griffiths in 1927 but it was not until the 1930’s that duplex stainless steels (DSS) became commercially available. About 80 years have passed since the first discovery but DSS are still under development. The interest in DSS in recent years derives from the high resistance of newly developed high alloy DSS to chloride induced corrosion. As a matter of fact it is the combination of several properties such as corrosion resistance, mechanical properties, weldability and price that makes the DSS unrivalled in many applications [1, 2]. Despite the attractiveness of DSS they have limitations. 475°C-embrittlement sets an upper limit to the temperature range recommended during service. Improper welding or production may cause precipitation of σ-phase or chromium nitrides resulting in deteriorated mechanical properties and/or corrosion properties. The endeavour to design gradually more corrosion resistant DSS provides a driving force to add more chromium, molybdenum and nitrogen, all of which destabilize the microstructure and promote formation of precipitates. The conflict between microstructural stability on one hand and the incentive to add more alloying elements on the other is a challenge to the designer of the alloys of tomorrow. The characteristic features of DSS, whether it is plastic deformation or corrosion, derive from the interplay between the two constituents ferrite and austenite. With the aid of modern computational tools it has become possible to predict microstructures with great precision and also simulate plastic deformation in a two-phase material such as a DSS. We therefore have powerful tools for simulating this interaction as a means of optimising corrosion and mechanical properties. Trends in the development of DSS Two trends in the development of DSS may be identified; one towards lean nickel-poor DSS and one towards highly alloyed so called super duplex stainless steels (SDSS). One advantage of lean DSS is the paucity of nickel, the price of which is high and also shows enormous fluctuations. 585
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F03-9
RECENT DEVELOPMENT OF DUPLEX STAINLESS STEELS
J.-O. Nilsson, G. Chai, U. Kivisäkk
R&D Centre, Sandvik Materials Technology, Sweden
Abstract
Recent development of duplex stainless steels is described. The advent of SAF 2707 HD, a 27Cr-7Ni-5Mo-0.4N duplex steel, shows that it is possible to reach a PRE-value close to 50 without sacrificing the fabricability. Modern methods for simulating the interaction between ferrite and austenite intimates that the steels of tomorrow may be optimized with respect to mechanical as well as corrosion properties. Methods under development presented here are multi-scale modelling of plastic deformation and high resolution electrochemical techniques.
Introduction
Duplex stainless steels (DSS) were first described by Bain and Griffiths in 1927 but it was not until the 1930’s that duplex stainless steels (DSS) became commercially available. About 80 years have passed since the first discovery but DSS are still under development. The interest in DSS in recent years derives from the high resistance of newly developed high alloy DSS to chloride induced corrosion. As a matter of fact it is the combination of several properties such as corrosion resistance, mechanical properties, weldability and price that makes the DSS unrivalled in many applications [1, 2]. Despite the attractiveness of DSS they have limitations. 475°C-embrittlement sets an upper limit to the temperature range recommended during service. Improper welding or production may cause precipitation of σ-phase or chromium nitrides resulting in deteriorated mechanical properties and/or corrosion properties. The endeavour to design gradually more corrosion resistant DSS provides a driving force to add more chromium, molybdenum and nitrogen, all of which destabilize the microstructure and promote formation of precipitates. The conflict between microstructural stability on one hand and the incentive to add more alloying elements on the other is a challenge to the designer of the alloys of tomorrow. The characteristic features of DSS, whether it is plastic deformation or corrosion, derive from the interplay between the two constituents ferrite and austenite. With the aid of modern computational tools it has become possible to predict microstructures with great precision and also simulate plastic deformation in a two-phase material such as a DSS. We therefore have powerful tools for simulating this interaction as a means of optimising corrosion and mechanical properties.
Trends in the development of DSS
Two trends in the development of DSS may be identified; one towards lean nickel-poor DSS and one towards highly alloyed so called super duplex stainless steels (SDSS). One advantage of lean DSS is the paucity of nickel, the price of which is high and also shows enormous fluctuations.
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However, if a high resistance to pitting corrosion is required significant amounts of the ferrite stabilizers chromium and molybdenum have to be used which sets a lower limit as regards the nickel concentration. The SDSS UNS S32750, S32760 and S32520 with a PRE-value of about 42 were introduced more than 15 years ago. It was thought that SDSS with a PRE-value well above 42 were remote but very recently a SDSS having a PRE-umber close to 50 has been launched. Existing SDSS have shown some limitations in high temperature sea water. Therefore, there has been a need on the market of a DSS with improved resistance to pitting corrosion. SAF 2707 was developed to meet this need and provided a leap in performance. The nominal composition of this alloy together with SAF 2507 is shown in Table 1 below. Table 1. Nominal chemical composition of two super duplex stainless steels
Grade UNS Cmax Cr Ni Mo N PRE SAF 2507 S32750 0.03 25 7 4 0.3 42 SAF 2707HD S32707 0.03 27 7 5 0.4 49 The comparison in Figure 1 shows that the pitting resistance is significantly improved in SAF 2707 compared to SAF 2507. The tests used in this comparative study was a modified version of the ASTM G48 test and a crevice corrosion test in 6% FeCl3 according to the MTI-2 procedure [3]. Critical pitting temperatures (CPT) can also be measured using potentiostatic tests at +600mV. The CPT as a function of the concentration of sodium chloride in the range 3-25% is shown in Figure 2. It is quite apparent that SAF 2707 HD is superior to SAF 2507 in the entire concentration range.
0
20
40
60
80
100
120
CPT G48C CCT MTI-2
Te
mp
era
ture
°C
SAF 2507
SAF 2707
Figure 1. Critical temperature assessed using modified G-48A and MTI-2 testing.
Figure 2. CPT obtained in the concentration range 3-25% NaCl.
Material modelling of micro mechanical behaviour in DSS
Since the austenitic phase and the ferritic phase have different chemical, physical and mechanical properties, these phases behave differently at the microstructural level. Each phase respond differently to the environments such as corrosion, thermal cycle and loading. For example, the load sharing between the individual phases during loading is different due to the 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 micro stresses that maintain their equilibrium among subsets of grains of different orientations [4]. These residual micro stresses can have great effects on SCC, yielding and damage of the material, and consequently affect their strength, deformation and fracture behaviour [5, 6]. Understanding the micromechanical
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reactions is therefore important for the application of the duplex stainless steels and for alloy design and material development. It is difficult to measure the stress-strain behaviour of the individual phases in DSS by the conventional mechanical testing methods due to its fine and heterogeneous microstructure. In-situ diffraction methods using X-ray, synchrotron and neutron are now used to analyze the load sharing, stress interaction between phases and grains and consequently the micro stress-strain behaviour of DSS [4, 7]. Hardness is a measure of the material resistance to plastic deformation. This indicates that the hardness (micro or nano) method can also be used to estimate plastic deformation hardening rate if the size of the austenitic and ferritic phase is sufficient [6]. In recent years, multi-scale material modelling has gained much interest from the researchers in the field of material mechanics. This type of modelling offers the possibility to study the behaviour of single phases, single grains and load sharing between the phases in a multi-phase material. The basic idea in multi-scale material modelling is that the a priori homogenized macro-scale material model is replaced by the homogenized response of a representative volume element (RVE) as shown in Figure 3. Multi-scale material modelling uses micro-scale crystal plasticity and continuum models [6]. Figure 4 shows the results of the multi-scale material modelling for SAF 2507 bar material in as delivered condition; 2507AD during static tensile testing. The ferritic phase is a stronger phase at a total strain less than about 3% and then becomes a softer phase with increasing strain. These observations are similar to the results from the experimental observations as shown in [6].
Figure 3. Multi-scale material modelling of duplex stainless steels
Fatigue is a progressive process. The early stage of fatigue damage is the permanent substructural and microstructural changes (strain localization) and creation of microscopic cracks. Fatigue damage is indicated by the formation of persistent slip bands (PSB) on the meso-micro scales and the subsequent crack initiation. Although much work has been done concerning two-phase or multi-phase metals, it is not clear in which phase or how fatigue damage occurs in these metals. Multi-scale material modelling provides the possibility to study the behaviour of single phases, single grains and load sharing between the phases in a multi-phase material as shown [6]. During cyclic strain loading, DSS materials usually have hardening and then softening processes. The simulations using the multi-scale material modelling show that hardening and softening processes also occur in the austenitic and ferritic phases [6], but behave differently. The ferritic phase has a shorter cyclic hardening period and lower hardening rate compared with the austenitic phase. In this paper, micro material damage is defined as the formation of slip bands in
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the individual phases. Figure 5 shows the micro damage behaviour in 2507AD due to cyclic loading. The accumulated effective plastic slips are mainly in the ferritic phase, but can also be observed in the austenitic phase. This can be explained by the fact that the damage in 2507AD may start in the austenitic phase but finally dominates in the ferritic phase since the weaker phase is the first to become damaged. This indicates that damage and crack initiation in a two-phase alloy depend not only on the initial strength of the individual phases, but also their deformation hardening behaviour. The final damage and crack initiation may occur in the weakest phase.
True strain
Truestress
2507AD
γ
α
DSS
Figure 4. Stress versus strain curves for 2507AD by multi-scale material modelling.
Figure 5. Representative volume element shows the simulated effective accumulated plastic slip in the 28th cycle for 2507AD. Lighter areas correspond to areas with a high degree of plastic slip.
Corrosion properties
The corrosion properties of duplex stainless steel depend upon the chemical composition as well as the degree of homogeneity of alloying elements. In an entirely austenitic stainless steel the distribution of elements is very homogeneous. However, a complication arises in DSS, in which chromium and molybdenum are partitioned to the ferrite and nitrogen is partitioned to the austenite. As a consequence, the PREN value [8] and the associated resistance to pitting may become notably different in the two phases. This problem in DSS can be circumvented by choosing an annealing temperature at which PREN are equal in austenite and ferrite, whereby equal pitting resistance in ferrite and austenite ensues [9]. As mentioned before -phase can be formed in DSS leading to a decrease in corrosion resistance. However, a finite amount of -phase is required to reduce the pitting corrosion significantly. In a previous investigation [10] the influence of various amounts of -phase on the pitting corrosion behaviour of Sandvik SAF 2507 and Sandvik SAF 2906 was made. It is shown that about 1% of -phase is necessary to significantly deteriorate the pitting corrosion resistance of Sandvik SAF 2507 and Sandvik SAF 2906. Qualitatively similar results have been reported by others for super duplex stainless steels [11]. In seawater Sandvik SAF 2507 with 6% -phase has passed a test at 35ûC [12]. Similar results have been found for UNS S32760 for which pitting was observed in chlorinated seawater at 35ûC with 6% -phase while no pitting was observed at 1.5% [13]. Modern electrochemical techniques offer a means of investigating the corrosion properties of DSS. The advantage of these techniques is that the potential (Scanning probe force microscopy
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(SKPFM)), or current distribution (electrochemical scanning tunnelling microscopy technique (EC-STM)), can be mapped with µm resolution. Since an AFM is used for the SKPFM this enables magnetic force microscopy (MFM) to be used to identify in which phase the measurements are performed. Hence, the corrosion properties of each phase in the DSS as well as galvanic interactions can be studied. The general corrosion properties of DSS have been studied in 1 M and 4 M H2SO4 with 1 M NaCl with both EC-STM and SKPFM. The austenite was found to be more noble than the ferrite and consequently more pronounced dissolution of ferrite was observed [14, 15].
Modelling of microstructures
The advent of thermodynamically based computer programs such as Thermo-Calc provided powerful tools for developing new alloys during the 1980’s. In fact, SAF 2507 was the first alloy ever to be developed and optimised using computerized techniques that later became known as Thermo-Calc [16]. The main achievement was to define a combination of temperature and composition that led to equal PRE and consequently equal pitting resistance in both phases. The development of SAF 2507 therefore provides a milestone in alloy development, not only within Sandvik but in the steel industry in general. The techniques have since been refined and developed further to include also DICTRA [17], a computer based tool by which diffusion controlled phase transformations can be modelled. Both programs are dependent upon experimental data such as activities, equilibrium tie lines, solubilities, diffusivities and surface energies. It is very often the case that the experimental data are uncertain and therefore limit the accuracy of the calculations. Surface energy is a parameter that is known for being difficult to measure experimentally with accuracy. As a consequence coarsening processes are difficult to model with accuracy. Fortunately, new tools are available for calculating surface energies. Using ab initio calculations based on density functional theory surface and interfacial energies can now be calculated with a precision that is far better than experimental methods can offer. Results from such calculations will provide new and more reliable input data to programs like Thermo-Calc and DICTRA and will therefore contribute to more accurate modelling of materials behaviour in the future.
Future prospects
Although DSS have been produced since the beginning of the 1930’s new DSS emerge continually. The trend in alloy development has been to increase the concentrations of chromium, molybdenum and nitrogen so as to improve the resistance to pitting corrosion. Also copper has been added to some DSS to enhance the resistance to general 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. There is also evidence that copper promotes spinodal decomposition of ferrite [18]. It is, therefore, 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 with a PRE-number close to 50 have been produced, thus confirming that alloys considered visionary a decade ago have now become a reality.
Acknowledgements
This paper is published with permission from Sandvik Materials Technology. The support from Prof. Olle Wijk is gratefully acknowledged.
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References
[1] H.D. SOLOMON and T.M. DEVINE, Proc. Conf. DSS (ed. R.A. Lula), Materials Park, OH, ASM, (1984) , pp. 693-756
[2] J. CHARLES, Proc. Conf DSS ’91, Les Ulis, France, Les Editions de Physique, (1991), pp. 3-48.
[3] K. GÖRANSSON, M.-L. NYMAN, M. HOLMQUIST AND E. GOMES, Sandvik SAF 2707 HD, Internal lecture no. S-51-63, 2006.
[4] N. JIA, R. Lin PENG, Y.D. WANG, G.C. CHAI, S. JOHANSSON, G. WANG, and P.K. LIAW, Acta Mater., 54, (2006), pp. 3907-3916.
[5] G. CHAI and R. LILLBACKA, (2006), Key Engineering Materials 324-325, (2006), p. 1117.
[6] R. LILLBACKA, G. CHAI, M. EKH, P. LIU, E. JOHNSON and K. RUNESSON, Acta Mater., 55, 2007, pp. 5359-5368.
[7] P.R LIN, J. GIBMEIR, S. EULERT, S. JOHANSSON and G.C. CHAI, Materials Science Forum 524-525, (2006), p. 847.
[8] G. HERBSLEB, Werkst. Korros., 33, (1982), pp. 33-34. [9] H. VANNEVIK, J.-O. NILSSON, J. FRODIGH and P. KANGAS, Trans. ISIJ 36, (1996),
pp. 807-812. [10] P. KANGAS and J.-O. NILSSON, Proc. Stainless Steel World 05 Conf., Maastricht (2005),
KCI Publishing BV, Zutphen, The Netherlands (2005). [11] A. TURNBULL, P.E. FRANCIS, M.P. RYAN, L.P. ORKNEY, A.J. GRIFFITHS AND B.
HAWKINS, Corrosion 58 , (2002), p. 12. [12] S. SOLTANIEH, M. KLOCKARS, U. KIVISÄKK and P. EKLUND, Stainless Steel World
15, (2002), p. 28. [13] R. FRANCIS and G.R. WARBURTON, Proc. Stainless Steel World 99 Conf., Hague
(1999), KCI Publishing BV, Zutphen, The Netherlands (1999), p.711 [14] M. FERMENIA, J. PAN, and C. LEYGRAF, Journal of Electrochemistry Society 149,
(2002) p. B187 [15] M. FERMENIA, C. CANALISA, J. PAN, and C. LEYGRAF, Journal of Electrochemistry
Society 150, (2003) p. B274. [16] B SUNDMAN, B. JANSSON and J-O ANDERSSON, Calphad, 9, (1985), pp 153-190 [17] J-O. ANDERSSON and J. ÅGREN, J. Appl. Phys., 72, (1992), p. 153. [18] H.D. SOLOMON and L.M. LEVINSON, Acta Metal., 26, (1978), pp. 429-442.
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DETECTION OF THE 475ºC EMBRITTLEMENT IN A LEAN DUPLEX STAINLESS STEEL USING THE ELECTROCHEMICAL
POTENTIODYNAMIC REACTIVATION (EPR) TEST
C. Fosca, J. Sakihama
Pontificia Universidad Católica del Perú, Peru
Abstract
The duplex stainless steel UNS S32101 (LDX 2101®) is a new leaner DSS that provide a high
mechanical resistance with a corrosion behavior, in most cases, better than the traditional austenitic
stainless steels. Although DSS are very competitive alloys, they are susceptible to precipitation of
secondary phases as the spinodal decomposition of the ferrite, when they are exposed at
temperatures between 300º-600ºC. The – ´ spinodal decomposition of ferrite can increase the
hardness of the DSS but reduce strongly its toughness and corrosion resistance.
Microstructure changes for each aging condition were characterized by Electrochemical
Potentiodynamic Reactivation EPR test in order to achieve a non destructive method to detect on
service detrimental aging conditions in this alloy. The EPR test was carried out using an appropriate
electrolyte composition (H2SO4 with addition of KSCN) at 20ºC. The reactivation potential and scan
rate were selected to improve more sensitivity to the microstructural changes.
Introduction
LDX 2101® (EN 1.4162, UNS S32101) is a new low alloyed (lean) Duplex Stainless Steels with
low addition of nickel in order to reduce the cost. To assure an adequate phase balance in the
microstructure (50% austenite, 50% ferrite), the austenite stability effect of the nickel is replaced
with additions of manganese and nitrogen. The mechanical resistance of this alloy is comparable
with the DSS 22%Cr-5%Ni (EN 1.4462, UNS S32205) and the corrosion properties are in general
better than for austenitic 304 (EN 1.4301)
The unique combination of mechanical properties, corrosion resistance and low cost, make this alloy
an excellent choice for many applications for which the traditional austenitic stainless steel are
usually employed.
Its relative low alloy content in comparison with others DSS brings the additional advantage to be
less sensitive to the secondary phase precipitation, when these materials are heated in the range of
700º- 900ºC (sigma phase, Chi-, R- phases, carbides) and in the range of 300º - 600ºC ( - ´ spinodal
decomposition of ferrite). All of these phase precipitations produces a strong reduction of the
material toughness, known as thermal aging embrittlement of DSS.
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Many researches have investigated and developed different methods to detect and also to quantify
the thermal embrittlement in the DSS [1,2,3,4]. Due to the presence of precipitated phases promoted
not only the embrittlement but also the localized corrosion of the DSS, electrochemical techniques
can be used to detect microstructural changes in these alloys [5].
The Electrochemical Potentiodynamic Reactivation (EPR) test, originally developed to detect the
intergranular corrosion of austenitic stainless steel, has been used also to detect the susceptibility to
intergranular corrosion of DSS due secondary phases [6,7,8] but there are scarce published results
about the use of this electrochemical technique to detect the 475ºC embrittlement of DSS [9]. The
aim of this article is to study the use of the EPR technique on the detection of the α−α´ aging
embrittlement in a new lean duplex stainless steel.
Experimental Procedure
As mentioned before, the material studied was the LDX 2101® duplex stainless steel. The alloy was
supplied in the form of a 6 mm plate in as received condition, with a chemical composition
described in table 1. The samples were aged at 475 ± 5°C during different times (4, 8, 16, 24, 48 and
72hrs) and cooled with water at room temperature.
Table 1. Chemical composition of the LDX 2101® duplex stainless steel.
Isothermal “short” ageing treatments of specimens, were carried out at temperatures 550-850°C
for 15-90 min and “long” treatment were carried out at 670 °C for 15-200 h.
The volume fractions of ferrite and austenite in a solution treated sample have been measured on
3 longitudinal and 3 transversal sections (20 fields for each section) by image analysis on light
micrographs at 200×, after etching with the Beraha’s reagent (reaction time, 10s).
The martensite which has been detected by OM and SEM, after etching with Beraha’s reagent
and by X ray diffraction (CrK radiation).
Different phases have been observed by SEM examination of polished samples, using the
backscattered electron (BSE) signal, on the basis of atomic number contrast effect: the ferrite
appears slightly darker than austenite, while the secondary phases would appears lighter. The
SEM operated at 25 kV; the BSE detector was set to maximize the atomic number contrast,
allowing ferrite, austenite and other phases to be identified.
Instrumented Charpy–V impact specimens were prepared in the standard form of 10×10×55 mm.
Impact test was carried out at room temperature, on samples treated at 550-650-750 and 850°C
for 15-45-90 and 120 min.
Results and discussion
Solution treated material
The banded structure of elongated γ islands is observed in the longitudinal section, while the
isotropic structure of ferrite and austenite grains is displayed on the transverse section. No
secondary phases were detected. The values of volume fractions of ferrite and austenite,
measured on longitudinal and transverse sections (200×), are reported in Table 2.
Table 2. Austenite (!) and ferrite (α) % vol. in longitudinal and transverse sections.
%2101 %2304 α %2101 α %2304
Longitudinal 50 ±2 56 ±2 50 ±2 53 ±3
Transverse 46±3 44±3 54±3 47±5
Table 3 reports chemical composition of ferrite and austenite measured with EDS-analysis,
expressed as partition coefficients.
The Ni and Mn austenite enrichment and Cr ferrite enrichment are evident, the partition
coefficients are quite similar to that observed in the common Cr-Ni-Mo grades.
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Table 3. Austenite and ferrite compositions. (Wt %, EDS)
!/ 2304 !/ 2101
Cr 1.23 1,14
Mn 0.75 0,84
Ni 0.59 0,62
Heat-treated samples
Microstructure of 2101 grade DSS
The microstructure has been investigated mainly on not-etched specimens by SEM-BSE; the
ferrite is darker than austenite. At the temperature of 600°C, for treatment time < 40 min no
precipitation of secondary phases has been detected, for longer times some small dark particles
were detected at the ferrite grain boundaries. They were analyzed by SEM-EDS (close to the
resolution limit) and an enrichment of Cr was observed at the grain boundaries so the precipitates
were identified as chromium nitrides.
The same grain boundaries precipitation was observed after soaking times longer than 40min at
650°C, while at 750°C the first grain boundary precipitation has been detected after a 20 min
treatment (Figure1a) and can still be observed after 20 h (Figure1b). Increasing the temperature,
particles became larger and the precipitation occurs also at the α-γ boundaries (Figure 1b).
(a) (b)
Figure 1. SEM-BSE micrographs of sample treated at 750 °C for 45 min (a) and 20 h (b)
The shortest times for grain boundary carbide precipitation lies in the temperature range
650-750°C, as already observed [3].
No σ and χ phases have been detected, neither for very long thermal treatments in the 650-900
temperature range. This could be related to the low Ni and Mo contents.
In addition the Ni content may induce the instability of the austenite, as suggested in previous
researches, which report of a probable transformation to martensite during cold forming (1).
Moreover the martensite formation has been confirmed (2) in some low-Ni DSS after cold
rolling and annealing (1040°C, air quenched).
We have detected different amount of martensite laths (Figures 2 and 3) in treated and rapidly
quenched from 750-850°C samples. Different cooling rates have no significant effect on the
amount of final martensite.
The X-ray diffraction spectra evidenced that the ferrite peaks increase as the amount of
martensite increases.
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Figure 2. SEM-BSE and OM micrograph (750 °C, 25 min, WQ): martensite laths.
Microstructure of 2304 grade DSS
The microstructure of the 2304 grade DSS is not affected by the heat treatment at 600°C and no
secondary phases or alpha-ferrite spinodal decomposition have been noted. A moderate
precipitation of carbo-nitrides has been detected after long treatments (100 hours) at 670°C and
after 45 min at 750°C.
In Figure 3 the SEM-BSE micrographs of the specimen treated for 90 min and 20 hours are
reported. The carbo-nitride precipitation is evident just below the austenite grain boundary, as it
has moved towards the austenite (ferrite) giving the precipitation inside the austenite grains.
A similar grain boundary precipitation was observed after long times (10 hours) treatment at
850°C but the kinetics are slower than at 750°C. On the other hand the main effect of the heat
treatment in the range 600-850°C is the increasing of the austenite volume fraction, with values
ranging from 44±1% of the as received sample to 62±2% of the sample treated at 850°C for
15 hours, accompanied by a decreasing of the Cr content in the austenite and by its increasing in
the ferrite. This Cr enrichment and the very low amount of Mo seem to stabilize the ferrite and
the secondary phases formation is not favoured.
The austenite of this grade appears to be stabile, indeed no austenite to martensite transformation
has been detected after heat treatment. Probably the Ni (4.3%) and N (0.18%) contents are high
enough to stabilize the austenite avoiding the structural transformation evidenced in the
2101 grade DSS.
Figure 3. SEM-BSE micrographs: 750°C for 90min(left) and 20 hours(right)
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Impact toughness
The effect of heat treatments on toughness of both the grades was studied by Charpy impact tests
carried out at room temperature.
As shown in Figures 4 and 5, the steels have different impact toughness properties: the 2101
grade has the ductile and fragile behaviour, while the 2304 has no the fragile behaviour.
The 2101 is ductile until 20-40 minutes of isothermal treatment at 600-650°C, corresponding to
first stages of the carbides-nitrides precipitation, where the impact energy drops down at about
50 J. The critical times for precipitation lie around 750°C, in good agreement with (3).
However the impact energy is never lower than 30 J, also after very long soaking times, of
several hours. The presence of some laths of martensite on the impact toughness has not yet been
investigated. The 2304 grade DSS is always ductile, with the impact energy values never lower
than 200 J.
At this stage of the research we may conclude that the presence of nitrides at the austenite grain
boundaries has no remarkable effects on the toughness of the 2304 steel.
The sample was treated at 550°C (1), 650°C (2), 750°C (3), 850°C (4), for 15 min (A),
45 min (B), 90 min (C), 120 min (D).
DUPLEX 2101
0
50
100
150
200
250
300
350
1A 1B 1C 1D 2A 2B 2C 2D 3A 3B 3C 3D 4A 4B 4C 4D
Heat Treatment
Imp
act
En
erg
y [
J]
(Geo
metr
ic T
est)
Figure 4. Impact energy of 2101 versus time/temperature of treatment
DUPLEX 2304
0
50
100
150
200
250
300
350
1A 1B 1C 1D 2A 2B 2C 2D 3A 3B 3C 3D 4A 4B 4C 4D
Heat Treatment
Imp
act
En
erg
y [
J]
(Geo
metr
ic T
est)
Figure 5. Impact energy for 2304 grade DSS versus time/temperature of heat treatment
Conclusions
Some results about the study of two duplex stainless steels with different low nickel contents
were presented:
- The relatively low nickel and molybdenum contents make the precipitation of
intermetallic phases more sluggish than in conventional duplex stainless steels, and no
sigma related phases precipitation has been detected, also after long time isothermal
aging treatments, in both the grades
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- Precipitation at the grain boundaries of chromium nitrides has been observed after
isothermal treatment in the temperature range 600-750°C, with different kinetics
- The austenite of the 2101 type DSS is quite instable, and a diffuse transformation
austenite-martensite has been evidenced, while the austenite of the 2304 DSS is more
stable and no martensite has been detected.
- The impact toughness after solution treatment is very good in both grades,
- The impact energy after isothermal treatment in the 2101 grade is never lower than 30 J,
while in the 2304 is never lower than 200 J,
- General corrosion properties in chloride environments are quite similar to that of
austenitic AISI 304 grades.
References
[1] C.D. Van Lelyveld, A. Van Benekom, Mater. Sci. Eng., A 204 (1996) p. 229.
[2] J.M. Hauser et al., Stainless Steel ’99, Science and Market, Sardinia, Italy, (1999), p. 85.
[3] P. Johansson, M. Liljas, AvestaPolarit Corrosion Management and Application
Engineering, 24 (2002) p. 17.
[4] H. Sieurin, R. Sandstrom, E.L. Westin, Met. Mat. Trans. A, 37A (2006) p. 2975
[5] I. Mészáros, Physica B, 372 (2006) p. 181
[6] I. Mészáros, P.J. Szabó:, Journal of NDT&E International, 38 (2005) 7, p. 517
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DUPLEX STAINLESS STEEL WELDS: RESIDUAL STRESS DETERMINATION, MAGNETIC FORCE MICROSCOPY AND
SUSCEPTIBILITY TO INTERGRANULAR CORROSION
B. Gideon1, L. Ward2, D.G. Carr3, O. Muransky3
1ARV Offshore, Thailand, 2RMIT University, Australia, 3Australian Nuclear Science and
Technology Organization (ANSTO), Australia
Abstract
A section of a Duplex Stainless Steel (DSS) pipeline girth weld (single vee joint configuration) was systematically analyzed to determine the stress / strain levels and correlation with susceptibility to IGC within austenite and ferrite phases in the weld cap, fill and root region. Stress / strain levels were determined by means of Neutron Diffraction techniques. Magnetic Force Microscopy (MFM) were used to determine the size, shape and distribution of the austenite and ferrite within the various regions. ASTM A262 and a modified Double Loop Electrochemical Potentiokinetic Reactivation (DL-EPR) Test methods were used to assess the susceptibility to IGC. A clear variation of stress/strain was evident between the austenite and ferrite in the base material, HAZ and weld from the neutron diffraction results obtained. The results of the weld metal from MFM shows the formation of both a finer and coarse structure within the weld metal, which is dependent on the level of undercooling. The values for Ir/Ia and Qr/Qa in the DL-ERP test results revealed that the fill area had the highest level of susceptibility to IGC.
Keywords: Duplex Stainless Steel, Intergranular Corrosion, Double Loop Electrochemical Potentiokinetic Reactivation, Magnetic Force Microscopy, Neutron Diffraction
Introduction
Welding of Duplex Stainless Steel (DSS) is particularly difficult with respect to maintaining a ferrite–austenite ratio close to 50:50. Rapid cooling effects associated with weld thermal cycles, often results in ferrite contents in the weld metal in excess of 50% may result in the loss of strength and increased susceptibility to IGC. The weld structure and the austenite / ferrite phase ratio are largely influenced by weld heat inputs and the cooling rates. The aim of this study is to conduct a detailed analysis of a girth welded sections of a DSS pipeline, as a function of heat input and type of weld, in terms of the residual stress by neutron diffraction, metallurgical analysis by means of magnetic force microscopy, and to assess the susceptibility to IGC by means of ASTM A262 and a modified Double Loop Electrochemical Potentiokinetic Reactivation (DL-EPR) test.
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Experimental Procedures
Welding Conditions
The DSS linepipe wall thickness was 10mm with a 200mm diameter of UNS 31803. Full details are listed in Table 1. Table 1. Chemical composition of pipe and filler material.
Note; Creq = Cr+1.37Mo+1.5Si+2Nb+3Ti and Nieg = Ni+22C+0.31Mn+14.2N+Cu
Manual Gas Tungsten Arc Welding (GTAW) technique was performed as detailed in Table 2. Upon completion of welding, the weld was subjected to Non Destructive Testing. Table 2. Weld Condition – Single Vee joint configuration
Weld
Pass Travel Speed Heat Input
mm/min J/min
1 (weld root) 51.00 1474.71
2 (weld fill) 123.00 883.12
3 (weld fill) 66.00 1745.45
4 (weld fill) 64.00 1788.00
5 (weld cap) 64.00 1685.63 Average 1515.38
Residual Stress Measurements
Residual strain measurements were made using neutron diffraction with a wavelength of 1.40Å on TASS (The Australian Strain Scanner) at the Australian Nuclear Science Technology Organization (ANSTO), Strains were measured in the three directions - longitudinal, transverse and normal (L,T and N) to the welding direction. These were the assumed principal stress directions. The measurement of residual elastic strain by monochromatic neutron diffraction relies on the use of Bragg’s law to relate the lattice spacings, dhkl, to the angle of diffraction 2 hkl associated with the diffraction peak labeled by Miller indices hkl at a fixed wavelength. Strain was calculated from the selected planar atomic spacing for ferrite and austenite at discrete locations in the weldment using Eq. 1.
00 /)( hklhklhklhkl ddd −=ε (1)
The calculation of the residual strains requires the knowledge of an appropriate reference lattice spacing 0
hkld . This is problematic in welds where there is a possibility of redistribution of alloying
elements, and secondly, inhomogeneous plastic deformation across the weld will generate relatively strong intergranular stresses in DSS. This problem was addressed by cutting a companion slice 2mm thick from the weld and cutting slits every 2mm across it‘s length in order to relieve the macroscopic residual stress field. Thus, the reference measurements 0
hkld represented the lattice
spacing as a function of position relative to the weld centre and included any effects of alloy diffusion and intergranular stresses.
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The average phase stress was calculated in the L,T and N directions for ferrite and austenite using the generalized Hooke’s law:
( )( )( ) ( )[ ]
( )( )( ) ( )[ ]
( )( )( ) ( )[ ]phase
Lphase
ThklphaseNhkl
hklhkl
hkl
phase
N
phaseN
phaseLhkl
phaseThkl
hklhkl
hkl
phase
T
phaseN
phaseThkl
phaseLhkl
hklhkl
hkl
phase
L
E
E
E
εενεννν
σ
εενεννν
σ
εενεννν
σ
++−−+
=
++−−+
=
++−−+
=
1211
1211
1211
_
_
_
(2)
Where !hkl and vhkl and the diffraction elastic constants for each phase. The macroscopic residual stress field was then calculated by weighing the contribution of the respective phase stresses according to Eq. 3, αγ σσσ NTLfNTLf
MacroNTL VV ,,,,,, )1( +−= (3)
The volume fraction Vf of ferrite was determined from the ASTM E562 point count method. Intergranular Corrosion Tests (IGC)
A modified ASTM A262 [1] was adopted in conjunction with a modified quantitative test method namely The Double Loop Electrochemical Potentiokinetic Reactivation (DL-ERP) test. Modified ASTM A262 Standard Practices E—copper–copper sulfate sulfuric acid test for detecting susceptibility to intergranular attack was used. The specimen was covered with copper shot and grindings and immersed in a solution of 16 wt% sulfuric acid with 6 wt% copper sulfate. The solution was then heated to its boiling point and maintained at this temperature for 48 hours. On removal from solution, the specimen was bent through 180° over a rod with a diameter equivalent to twice the thickness of the specimen instead of four times the thickness to ensure, if cracks appeared, they would be apparent by the more restrictive bending radius. The bent surface of the specimen was then examined for cracks at low magnifications in the range X5 to X20.
A modified Double Loop Electrochemical Potentiokinetic Reactivation (DL-ERP) Tests, was used as conducted by Schultz et al. [2,3,4]. The solution used was 0.5M H2SO4 + 0.001M TA (thioacetamide). TA is added to reduce the extent of ferrite dissolution. The test was conducted at 60 °C. The polarization scan was started 5 minutes after immersion of the specimen. The potential was scanned from -500 mV (SCE) to +200 mV (SCE) and back to -500 mV (SCE) at a rate of 1.67 mV/s. The ratio of the reactivation charge to the passivation charge was calculated and is shown in the results section.
Magnetic Force Microscopy Analysis
Magnetic force microscopy studies were conducted on metallographically prepared cross-sections of the welds, after grinding and polishing using 3 µm diamond paste. The Scanning Probe Microscopy from Digital Instruments at ANSTO, operating in tapping and lift modes was employed to study the topographic and magnetic features of the DSS samples. Topographic and magnetic force data were taken in the same scan. In order to produce reliable images, repeated scans in different directions were done to ensure reproducibility of the features. Various scan sizes and speeds were tested to enhance height and magnetic induced signals, thus minimizing tip hysteresis and the delay between line scans.
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Results and Discussion
The residual stress, microstructure, resulting phase transformation, mechanical properties and degree of susceptibility to IGC are discussed in detail in this section. Table 3 summarises the results of the IGC tests carried out on the welded duplex stainless steels in this study.
Table 3. Summary of the DLERP results of the of the welded duplex stainless steels
Weld
Pass DL-ERP Test
Qr/Qa Ir/Ia
1 (weld root) 0.04 0.06
2 (weld fill) - -
3 (weld fill) 0.09 0.12
4 (weld fill) - -
5 (weld cap) 0.04 0.07
Residual Stress Measurements
In the transverse direction, austenite exhibits tensile strains in the weld while the ferrite has contracted lattice spacing. As the distance from the weld centerline increases out to the HAZ (~5mm), an inversion occurs where upon ferrite strains are tensile and austenite is compressive. In the longitudinal direction, the strains for both phases are initially tensile in the weld, although the magnitudes are inverted for each phase in comparison to the transverse direction. In this direction the macroscopic residual stress field is at a maximum, due to constraint impeding contraction of the weld bead during cooling, and it is likely that this is the dominating effect. Moving out from the weld, the HAZ can be clearly distinguished from the weld as both average phase strains become uniformly tensile. This is interesting, in that this area of the weld undergoes transformation back to a completely austenitic structure before transforming partially back to ferrite [5]. In order to convert phase strain to stress (Eq.2) the diffraction elastic constants !hkl and vhkl for each phase must be known, and these in turn depend on the crystallographic texture of the weldment which varies with position from parent to weld. Given the demanding experimental requirement for the texture at each location in the weld, a best approximation of the diffraction elastic constants was chosen using the self-consistent scheme proposed by Kröner [6] for random texture. Such that,
α211E = 225.5, γ
311E = 183.5 GPa, and αν 211 = 0.28 , γν 311 = 0.31 for the ferrite and austenite phases. The
calculated phase and macro stresses (Eq.3). In the normal and transverse direction, the HAZ is strongly tensile for ferrite and compressive for austenite. These results suggest tensile ferrite regions could be susceptible to cracking in the HAZ. It is interesting to note that each phase is under very different stress states throughout the weldment. Considering the samples studied do not have the additive operational stresses normally superimposed on the residual stress field, it is quite likely that ferrite could be subjected to large tensile stresses in practice. Very high compressive stresses were estimated in the austenite phase for the transverse and normal directions, however, these stresses appear to balance by observation of the macroscopic stress field. It is a requirement for stress balance that the macroscopic stress in the normal direction tend towards zero at the surface and this generally true, however, the magnitude of the compressive stress in the austenite 2mm form the ID surface is questionable. A systematic error in the stress free reference may be a likely source of error. In the weld, the results show both austenite and ferrite to be under tensile stress in the transverse and longitudinal directions. Observation of the macroscopic stress field shows the highest stress to
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occur in the longitudinal (welding) direction as expected. In the transverse direction, where cracking is most problematic in welds, the highest ferrite phase stresses occur in the mid-thickness of the plate. Microstructural Evaluation and Magnetic Force Microscopy
Microstructural analysis for both GTAW weld conditions as shown by the magnetic force microscopy (MFM) images in Figure 1 reveals the presence of a two-phase banded structure, typical of such materials. In general, the austenite regions observed in the DSS weld metal is formed from ferrite in three modes, viz., as allotriomorphs at the prior-ferrite grain boundaries, as Widmanstätten side-plates growing into the grains from these allotriomorphs and as intragranular precipitates [7]. In the micrographs, the grain boundary allotriomorphs and Widmanstätten austenite are clearly seen. However, the austenite seen within the grain could be either intragranular precipitates or Widmanstätten austenite intercepted transverse to the long axis. These microstructures, in addition to the presence of discontinuous grain boundary austenite layers (Figure 1a) and intragranular acicular ferrite are thought to be associated with variations in transformation rates and the degree of undercooling [8]. In summary, these observed microstructures are typical of those formed under such welding conditions. The topographic image of (Figure 1d) showed a very flat surface where the only distinguishable features were some contamination particles and a few grinding scratches. From this image, it was not possible to distinguish the distribution of the ferritic and austenitic phases over the surface. On the other hand, the magnetic domain distribution presented in Figures 1a, 1b and 1c are thought to be associated with the microstructures, typical of the various DSS weld regions. The MFM technique was capable of clearly imaging the magnetic domain structure of the ferrite phase that surrounds the “islands” of austenite, appearing flat and uniform due to their paramagnetic properties. Clear bands of ferrite could be easily distinguished, but a closer look revealed other regions of ferrite that did not exhibit the more typical striped magnetic domain configuration, similar to the ferrite regions.
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Figure 1. MFM Image of DSS; a) Root region, b) Fill region, c) Cap region, d) Topographical image of weld.
Intergranular Corrosion Tests
Modified ASTM A262 Standard Practices E Test
The absence of cracks on the surface of the bent specimens, even under restricted and reduced bending radius, in accordance with ASTM A262 Standard Practices E, indicates no evidence of sensitization in all weld conditions.
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Modified DL-ERP test
The test efficiency was measured by means of a response test, which was characterized by weak values of the current density ratio (Ir/Ia <1%) and the charge ratio (Qr/Qa < 1%) for nonsensitized materials, and relatively high ratio values (Ir/Ia "1%) and (Qr/Qa " 1%) for high-sensitized materials. The reverse polarization from the passive to the active region gave rise to a reactivation peak, the magnitude of which is sensitive to the degree of alloy element depletion. The susceptibility to corrosion was characterized in terms of both the ratio of the reactivation-current peak to the activation current peak as well as the ratio of the reactivation charge to the activation charge [9]. Analysis of the results shows that the fill region had a higher degree of sensitization (DOS) compared to the root and cap region of the weld. This correlates with the neutron diffraction measurements of the average phase stress, where tensile ferrite and austenite stresses were observed to be at a maximum in the fill region of the weld in the transverse direction.
Conclusions
- Microstructure of the weld metal as detailed in Figure 1 shows a typical “as weld” structure, resulting in the formation of both fine and coarse structure within the weld metal. This was thought to be associated with variations in transformation rates and the degree of undercooling.
- It was shown the MFM is a powerful tool to use for differentiating the austenite and ferrite phase in duplex stainless steel.
- The DL-ERP test results revealed that the fill area had the highest values for Ir/Ia and Qr/Qa.
- Residual stress measurements by neutron diffraction revealed that the ferrite phase stress was tensile in the HAZ and weld and appeared to be balanced by a local compressive austenite phase stresses in the normal and transverse directions. The results showed that for ferrite and austenite, a maximum tensile stress is formed in the fill section of the weld and decreases in the root and cap regions for the transverse direction.
- A correlation was observed between the stress / strain distribution in the DSS weld regions and the degree of susceptibility to IGC.
References [1] ASTM A262, Standard Practices for Detecting Susceptibility to Intergranular Attack in
Austenitic Stainless Steels. [2] S. Schultze, J. Gollner, K. Eick, P. Veit, I. Garz, “ The modified EPR test: A new tool for
examination of corrosion susceptibility of duplex stainless steel”, Duplex ’97, Paper D97-067, KCI Publishing, Zutphen, The Netherlands (1997), p. 639.
[3] A. Turnbull, P.E. Francis, A.J. Griffiths, E. Bennett, W. Nimmo, “Measurement of Corrosion Resistance of Super-Duplex Stainless Steel Welds by Electrochemical Techniques,” Eurocorr” 2000 (London, U.K.: Institute of Materials, 2000).
[4] E. Otero, C. Merino, C. Fosca, P. Fernandez, “Electrochemical Characterization of Secondary Phases in a Duplex Stainless Steel by EPR Test,” Duplex ’94, paper no. 56 (Cambridge, U.K.: TWI, 1994).
[5] B. Gideon, L. Ward, D. G. Carr “Strain Measurements by Neutron Diffraction and Characterization of Duplex Stainless Steel Welds. Duplex 2007 Conference”, paper 49 (Aquileia and Grado Italy)
[6] E. Kröner, Berechnung der elastischen Konstanten des Vielkristalls aus den Konstanten des Einkristalls, Z. Physik. 151, 1958, p. 404-418.
[7] B. Gideon, L. Ward, G. Biddle, “Testing and Characterization of Duplex Stainless Steel Welds and their Susceptibility to Intergranular Corrosion”, Eurocorr”06, (Maastricht,
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Netherlands, September 2006). [8] L. Karlsson, Welding in the World, Vol. 43, no. 5, 1999. [9] T. Amadou, C. Braham, and H. Sidhom, Metallurgical and Materials Transactions A, Vol.