Rapid precipitation of intermetallic phases during isothermal treatment of duplex stainless steel joints produced by friction stir welding

Journal Pre-proof Rapid precipitation of intermetallic phases during isothermal treatment of duplex stainless steel joints produced by friction stir welding Igor J. Marques, Flavio J. Silva, Tiago F.A. Santos PII:

S0925-8388(19)34416-0

DOI:

https://doi.org/10.1016/j.jallcom.2019.153170

Reference:

JALCOM 153170

To appear in:

Journal of Alloys and Compounds

Received Date: 12 August 2019 Revised Date:

20 November 2019

Accepted Date: 23 November 2019

Please cite this article as: I.J. Marques, F.J. Silva, T.F.A. Santos, Rapid precipitation of intermetallic phases during isothermal treatment of duplex stainless steel joints produced by friction stir welding, Journal of Alloys and Compounds (2019), doi: https://doi.org/10.1016/j.jallcom.2019.153170. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.

Rapid precipitation of intermetallic phases during isothermal treatment of duplex stainless steel joints produced by friction stir welding

Igor J. Marques1, Flavio J. Silva1, Tiago F. A. Santos1,*

1

Department of Mechanical Engineering, Universidade Federal de Pernambuco,

Avenida da Arquitetura, s/n, Cidade Universitária, 50750-550, Recife, PE, Brazil *[email protected]

Abstract The propensity of duplex stainless steels to form secondary phases that harm their main properties is a metallurgical problem that continues to be studied. Friction stir welding has shown excellent results when applied to duplex stainless steels, enabling the formation of joints that maintain the austenite-ferrite balance, without the presence of secondary phases. Welding processes cause significant microstructural reformulation in the joint microstructure, so the properties and metallurgical behavior of welded joints differ significantly from those of the corresponding base metals. The formation of secondary phases in duplex stainless steels occurs by diffusional transformation processes that are dependent on the transformation time, with the kinetics being influenced by the previous microstructure. Although friction stir welding makes it possible to obtain joints free from secondary phases, it is important to study the propensity of welded joints to subsequently form secondary phases, in order to be able to understand the behavior of these joints when they are subjected to hot operations, such as weld repairs. In this work, friction stir welded joints were treated isothermally at 850 °C and were analyzed using scanning electron microscopy, energy-dispersive X-ray spectroscopy, Ferritoscope measurements, and X-ray diffractometry. The results revealed that the friction stir welded joints presented much faster formation of intermetallic phases, compared to the base metals, with the appearance of regions showing advanced stages of intermetallic phase precipitation after only 5 min exposure at 850 °C.

1

Keywords: Duplex stainless steels; sigma phase; chi phase; friction stir welding; thermal treatment; phase transformation.

Introduction Duplex stainless steels (DSS) are used in applications that require a combination of high corrosion resistance and excellent mechanical performance, such as in the gas and oil, chemical, paper and cellulose, food, and seawater desalination industrial sectors [1,2]. The properties of these steels are the result of a biphasic microstructure composed of austenite and ferrite in balanced proportions, with an absence of secondary phases in the microstructure [3,4]. The chemical composition of DSS typically includes the elements Fe, Cr, Ni, Mo, Mn, and N, sometimes with the addition of Cu and W [1]. DSS are high-alloyed steels and their wide range of chemical components favors the formation of secondary phases when exposed to temperatures between 300 and 1050 °C [5,6]. The important secondary phases reported in DSS include α’, σ, χ, carbide, and nitride phases. In the work of Farias et al., it was found that over 50% of the DSS failures reported in the literature occurred due to the precipitation of secondary phases, especially the σ phase [7], which is a tetragonal phase essentially composed of Fe and Cr, formed at temperatures between 600 and 950 °C. This phase has been most extensively studied in DSS, because it can develop large volumetric fractions and has strong deleterious effects on corrosion resistance [8], fracture resistance, and ductility [9,10]. The χ phase is typically associated with the σ phase, since the former is frequently a predecessor in microstructures that evolve to formation of the σ phase [11– 13]. The χ phase, which presents Fe36Cr12Mo10 stoichiometry, has been reported to occur at temperatures between 700 and 900 °C. However, this phase is metastable in commercial duplex stainless steels and is consumed by formation of the σ phase, as the system approaches thermodynamic equilibrium [11,14]. The chemical compositions of the σ and χ phases present higher levels of Cr and Mo, compared to the α and γ phases [13]. During the formation of these phases, the α phase is consumed and there is concentration of gammagenic elements, resulting in the emergence of a secondary austenite depleted in Cr (γ’), For this reason, formation of the sigma phase is usually described as a eutectoid reaction: α → σ + γ’ [11]. 2

Welding is essential for the large-scale use of steels. However, welding operations cause severe microstructural reformulations and the possible occurrence of metallurgical problems. Conventional welding processes involve melting of the material, which can lead to problems such as heat cracking, hydrogen-induced cracking, and segregation of alloying elements. An additional problem concerns microstructural changes during welding of DSS, since the exposure of these steels to high temperatures may not only favor the formation of secondary phases, but also lead to excessive formation of ferrite, with these changes greatly altering the properties of the welded joint in the stir zone (SZ), compared to the base metal (BM) [15]. For satisfactory formation of DSS welded joints, it is recommended to avoid very high or very low thermal inputs, because the thermal input determines the rate of cooling of the molten zone [15,16]. High cooling rates tend to lead to excessive formation of ferrite and precipitation of nitrides, while low cooling rates favor the formation of intermetallic phases [15]. Friction stir welding (FSW) is a solid state process that, in addition to avoiding the typical problems associated with fusion welding [16], produces narrower heat affected zones (HAZ) during welding. Applied to DSS, FSW favors maintenance of the α/γ balance in the weld metal [17], besides enabling microstructures to be obtained without the presence of intermetallic phases and HAZ, resulting in improved mechanical properties [18–20] and corrosion resistance [21–23]. Despite these positive aspects, there have been no reports concerning the way that the microstructural reformulation induced by FSW could affect the precipitation kinetics of secondary phases in operations at temperatures exceeding 280 °C. The σ and χ phases are formed in reactions involving their nucleation and growth, so their formation depends on the duration of exposure to the temperature range that promotes stability of these phases [24]. Several factors influence the kinetics of formation of intermetallic phases, including the microstructure [25,26], crystallographic characteristics [27,28], and chemical composition [29]. The kinetics of solid state isothermal transformations is typically analyzed using the Avrami mathematical model (Equation 1) [30]. The Avrami model was established in a series of papers published by Melvin Avrami [30–32], in addition to works by Kolmogorov [33], Johnson and Mehl [34], Cahn [35], Wert and Zener [36], and others [37,38]. The Avrami model predicts exponential progression of the transformation, where the constants K and n describe the reaction kinetics. The constant K is related to the temperature of the transformation 3

process and to microstructural aspects that influence the rates of nucleation and growth. The constant n, commonly called the Avrami constant, is related to the dimensionality of the nucleation and growth sites [39].

Vv (t ) = 1 − e− ( Kt )

n

(1)

In DSS, the main sites for intermetallic phase nucleation are the α/γ interfaces and triple points [26], so even brief microstructural changes, such as microstructure refinement, can lead to significant changes in the kinetics of σ phase formation [25]. Some studies have proposed that the kinetics of formation of the σ phase can be described by the Avrami equation in two stages [40–42]. The kinetic behavior in two stages is due to alteration between the mechanisms controlling phase formation [41,42], or to the influence of the χ phase in the transformation process [40]. Besides the problem of microstructure development during welding, the weld metal should be evaluated considering transformations that could occur during exposure to high temperatures, such as during welding repairs. Since microstructural changes affect the formation of intermetallic phases, severe microstructural reformulation, as occurs in welding operations, may cause significant alteration of the kinetics of formation of intermetallic phases. Sato and Kokawa showed that welding using autogenous GTAW could retard the formation kinetics of intermetallic phases, which was attributed to the development of coherent α/γ interfaces, according to the Kurdjumov-Sachs orientation relation (K-S OR) [27]. It is known that friction stir welding favors the formation of intermetallic phases in austenitic stainless steels [43]. However, previous studies by Santos et al. [17,18], Sarlak et al. [22], and Saeid et al. [20] found no evidence of intermetallic phase formation during welding. Other studies have found that the FSW process can lead to the formation of intermetallic phases in the microstructure in cases of high thermal input, as reported by Sugimoto et al. [44], or in materials that are more liable to formation of such deleterious phases [17,21]. The present work investigates the kinetics of formation of intermetallic phases in the SZ of DSS welded by FSW, submitted to isothermal treatment, and evaluates the effect of microstructural reformulation induced by the FSW technique on the kinetics of intermetallic phase formation.

4

Materials and Methods The materials used were 6 mm plates of UNS S32205 and superduplex UNS S32750 and UNS S32760. The chemical compositions of the as-received materials, provided by the manufacturers, are shown in Table 1 [45,46].

UNS

Table 1. Chemical compositions (wt.%) of the steels used in this work. C Si Mn Cr Ni Mo W Cu N P

S

S32205

0.02

0.30

1.80

22.5

5.40

2.80

------

------

0.16

0.030

0.001

S32750

0.02

0.25

0.78

24.9

6.88

3.79

------

0.34

0.26

0.023

0.001

S32760

0.02

0.35

0.64

25.2

7.00

3.70

0.62

0.62

0.23

0.024

0.002

The FSW joints were produced using a Transformation Technologies RM-1a system operated with a rotation speed of 200 rpm, advance speed of 100 mm/min, and uniaxial force of 37 kN. The welding tool used was 40% PCBN in a W-Re matrix, with a conical pin (6 mm length, 10 mm maximum diameter, and shoulder diameter of 25 mm). Details of the welding process can be found in previous studies concerning the process and the associated microstructural evolution [17,18], but which did not consider the kinetics of the joint under post-welded conditions. The welded joints and the corresponding base materials were sectioned to produce specimens suitable for heat treatment in a muffle furnace. The treatment was performed isothermally at 850 °C for times between 5 min and 2 h, followed by cooling in a water bath with ice. This isothermal treatment temperature was selected because it provided the fastest precipitation of intermetallic phases in these steels [42,47], as shown in preliminary tests. The atmosphere was not controlled during the isothermal treatments. The samples were prepared using water-resistant abrasive papers (80, 120, 220, 220, 320, 600, 1000, and 1500 mesh), followed by polishing with 3 µm diamond paste. The faces prepared and analyzed were the one normal to the lamination, in the case of the BM, and the top face, in the case of the SZ. Microstructural characterization of the samples was performed by scanning electron microscopy (SEM), using the backscattered electron (BSE) signal for contrast differentiation of the chemical components of the phases formed. The BM was analyzed using a Hitachi TM-3000 instrument, while the SZ was characterized using a Zeiss EVO MA 15 instrument. Chemical microanalysis by energy dispersive X-ray spectroscopy (EDS) employed a TESCAN MIRA3 SEM-FEG system coupled to an 5

Oxford Instruments X-act EDS analyzer. The SEM BSE images were submitted to digital image analysis for quantification of the volumetric fractions of the phases, with the images being processed using a Gaussian filter and classified into three classes by thresholding. The volumetric fraction of ferrite was determined using a Fischer MP30 ferritoscope. The microstructural crystallography was characterized using electron backscatter diffraction (EBSD) maps, obtained using a Philips FEG XL 30 SEM, with preprocessing using TSL OIM Analysis software, followed by analysis using the MTEX toolbox [48]. The EBSD maps were only obtained for the samples that were not submitted to the isothermal treatment. X-ray diffraction (XRD) analysis was performed using Cu Kα1 fluorescence radiation (1.5406 Å), with scanning in the 2θ range between 35° and 105°, with scanning speed 1°/min (step size of 0.02°, exposure time of 1.2 s), voltage of 40 kV, and current of 30 mA. Theoretical powder diffraction patterns were calculated using pymatgen [49].

Results Friction stir welding The FSW procedures resulted in welded joints with full penetration and no defects. After welding, the regions with significantly different microstructural characteristics were identified, considering the stir zone (SZ), the stir zone advancing side (SZ-AS), the stir zone retreating side (SZ-RS), the thermomechanically affected zone advancing side (TMAZ-AS), and the base metal (BM). It was not possible to identify the heat affected zone (HAZ) in the microstructures obtained after the FSW. Figure 1 shows EBSD maps of the BM and the corresponding SZ. Considerable differences can be seen among the SZ, SZ-AS, and SZ-RS, caused by the asymmetry of the FSW process. In all cases, there was significant grain refinement, compared to the BM, but the degree of refinement was highest in the SZ-AS, followed by the SZ and the SZ-RS. The microstructures resulting from the FSW showed no secondary phases, except in the case of the UNS S32760, for which there was evidence of the formation of carbides and intermetallic phases, as observed for the SZ in previous studies [17,18,21]. The earlier studies provide further details concerning the microstructure and mechanical properties [18], microstructural development and recovery mechanisms [17], and corrosion resistance of the joints [21,50]. 6

Figure 1. Phase maps obtained by EBSD for the BM and SZ of the UNS S32205, UNS S32750, and UNS S32760 stainless steels without thermal treatment.

Microstructural analysis of the isothermally treated samples The microstructural characterization of the isothermally treated samples enabled the identification of from two to four phases in all the situations analyzed. Figure 2(a) shows the microstructure of the UNS S32205 treated isothermally at 850 °C for 1 h, indicating the points selected for EDS analysis. In the area analyzed, it was possible to identify the presence of four phases, distinguished by the chemical contrast provided by the backscattered electron signal, resulting in intensities ranging from white to gray. For duplex stainless steels, the phases typically formed in heat treatments between 700 and 900 °C, identified by SEM-BSE with colors ranging from the lightest to the darkest, are the χ, σ, γ, and α phases, as reported elsewhere [12,13]. Figure 2(b) shows the EDS spectra acquired at the points indicated in Figure 2(a). Table 2 shows the quantification results for the EDS spectra shown in Figure 2(b). The EDS analyses confirmed that the microstructure was composed of α, γ, χ, and σ phases. The γ phase presented a higher Ni content, compared to the other phases, while the χ phase presented a peak for Mo 7

that was significantly more intense, compared to the other phases. The σ phase showed the highest Cr content, together with a significant Mo content (although lower than that of the χ phase). The α phase was identified by the presence of alphagenic material and presented a higher Fe concentration, compared to the intermetallic phases. The EDS results confirmed that the phases identified following isothermal treatment at 850 °C were χ, σ, γ, and α, with colors from lightest to darkest, as expected and in agreement with the literature [12,13].

Figure 2. (a) Microstructure of the BM of UNS S32205 treated isothermally at 850 °C for 1 h, with EDS spectra acquired at the points indicated (SEM-BSE analysis of the polished surface); (b) EDS spectra obtained at the points indicated in (a).

Table 2. Quantification results for the EDS spectra shown in Figure 2, with the phases identified based on the values obtained. Concentration (wt.%) Phase Spectrum Fe Cr Ni Mo Mn Si identified 66.44 22.13 6.58 2.42 1.99 0.45 1 γ 65.08 26.60 3.06 3.21 1.63 0.42 2 α 57.36 29.89 2.70 7.60 1.90 0.54 3 σ 57.19 24.63 3.12 12.74 1.72 0.61 4 χ

Figure 3(a) shows micrographs of the BM of the UNS S32205 treated isothermally at 850 °C for 10, 30, and 60 min. The images showed that after 10 min of isothermal treatment, there was the formation of small precipitates of χ and σ phases dispersed 8

along the α/γ interfaces. With continuing isothermal treatment, there was growth of the intermetallic phases towards the interior of the ferritic sites and a progressive decrease of the ferrite volumetric fraction. After 1 h of isothermal treatment, a significant volumetric fraction of intermetallic phases had developed within the microstructure, occupying a substantial portion of the previous ferritic sites. Figure 3(b) shows micrographs of the SZ of the UNS S32205 treated isothermally at 850 °C for 5, 10, and 30 min. After only 5 min of treatment, the SZ microstructure had developed substantial amounts of intermetallic phases. Notably, formation of the intermetallic phases was fastest in the SZ-AS, followed by the SZ and the SZ-RS, which was most evident in comparison with the other regions, where the ferritic sites were preserved for longer during the isothermal treatment. The SZ-AS presented only slight microstructural alteration during continuing isothermal treatment, with progressively fewer ferritic sites available for phase precipitation, while the alterations in the microstructures of SZ and SZ-RS were much more evident with continuing isothermal treatment. The main intermetallic phase present in the SZ, under all the conditions analyzed, was the σ phase. The χ phase was only identified at points dispersed throughout all regions of the sample treated isothermally for 5 min.

9

Figure 3. Micrographs of the UNS S32205 treated isothermally at 850 °C: (a) BM after 10, 30, and 60 min; (b) SZ after 5, 10, and 60 min. SEM-BSE analyses of the polished surface.

Figure 4(a) shows micrographs of the BM of the UNS S32750 treated isothermally at 850 °C for 5, 10, and 30 min. After only 5 min, most of the ferrite of the microstructure had been consumed by formation of the σ phase. As the isothermal treatment time increased, there was only brief evolution of the volumetric fraction of the σ phase, which continued to consume the remainder of the ferritic matrix. For all the thermal treatment conditions, it was possible to identify small particles of χ phase dispersed in the microstructure between the σ and α sites. Figure 4(b) shows micrographs of the SZ of the UNS S32750 treated isothermally at 850 °C for 5, 10, and 30 min. As in the case of the BM, after only 5 min of isothermal treatment, the SZ of the UNS S32750 already presented widespread formation of

10

intermetallic phases. The σ phase developed in the previously ferritic flow lines of the FSW microstructure, while the χ phase was identified next to the remaining ferritic sites. It was notable that for all regions of the SZ, the fraction of ferrite remaining after 5 min of isothermal treatment was lower than for the BM of the UNS S32750. After 10 min of isothermal treatment, it was only possible to identify a brief decrease of the ferritic sites, which was more pronounced in the SZ and the SZ-RS. After 30 min of treatment, only small ferritic sites remained dispersed throughout the SZ, notably in the SZ-AS, where there was a marked scarcity of ferritic sites.

Figure 4. Micrographs of the UNS S32750 treated isothermally at 850 °C for 5, 10, and 30 min: (a) BM; (b) SZ. SEM-BSE analyses of the polished surface.

11

Figure 5(a) shows micrographs of the BM of the UNS S32760 treated isothermally at 850 °C for 5, 10, and 30 min. There was no evidence of the formation of intermetallic phases after 5 min of treatment, with the microstructure presenting the lamination features typical of DSS composed of the α and γ phases. After 10 min, there was clear evidence of formation of moderate amounts of the σ and χ intermetallic phases, which precipitated and grew around the ferritic grains and at the α/γ interfaces. After 30 min, the microstructure presented massive formation of intermetallic phases that occupied almost all the previously ferritic region, with some islands of ferrite still visible, surrounded by intermetallic phases. After 30 min of isothermal treatment, there was still a substantial fraction of χ phase, while the σ phase accounted for almost all the volumetric fraction of intermetallic phases. Figure 5(b) shows micrographs of the SZ of the UNS S32760 treated isothermally at 850 °C for 5, 20, and 30 min. Different to the BM, the SZ of UNS S32760 presented widespread formation of σ and χ phases after only 5 min. As the isothermal treatment progressed, the contrast between the σ and χ phases became more pronounced, due to the concentrations of the alloying elements characteristic of each phase. Ferritic islands well dispersed in the microstructure could only be identified in the microstructure resulting from 5 min of isothermal treatment. After 20 min, the ferritic islands became scarcer, except for small particles of ferrite dispersed in the SZ-RS. After 30 min, the particles of α phase became scarce in all regions of the SZ, while large quantities of χ phase particles could be seen dispersed in the sites of σ phase.

12

Figure 5. Micrographs of UNS S32760 treated isothermally at 850 °C: (a) BM after 5, 10, and 30 min; (b) SZ after 5, 20, and 30 min. SEM-BSE analyses of the polished surface.

Due to the rapid formation of intermetallic phases, the residual ferrite fraction could be used as an effective parameter to follow the progress of the α → σ + γ’ transformation, with a smaller fraction of ferrite indicating greater phase transformation. Therefore, ferritoscope measurements were performed in order to provide confirmation of the SEM results. The ferritoscopy results (Figure 6) indicated that the α phase volumetric fraction decreased fastest in the SZ, due to the formation of intermetallic phases. For the UNS S32205 and S32750, the ferrite fraction clearly decreased fastest in the SZ-AS, followed by the SZ and the SZ-RS. For the UNS S32760, the ferritoscope results indicated that decomposition of the α phase occurred fastest in the center of the SZ, followed by the SZ-AS and the SZ-RS. A further observation was that for the UNS S32760, the results for the SZ-AS and the SZ-RS were initially statistically equivalent, while continuing 13

isothermal treatment led to the rate of decomposition of the SZ-AS becoming similar to that of the central SZ.

Figure 6. Ferritoscope measurements results, showing the evolution of the ferrite volumetric fraction during isothermal treatment at 850 °C.

Figure 7 shows the XRD results for the SZ of the samples treated isothermally at 850 °C, confirming that the phases identified by SEM and EDS analyses corresponded to the α, γ, σ, and χ phases. No other phases were detected, indicating that either there were no other phases present, or that they were present at volumetric fractions below the detection limit of the technique (around 2%) [51].

14

Figure 7. XRD patterns for the SZ of samples treated isothermally at 850 °C for 10 min (UNS S32205) or 5 min (UNS S32750 and S32760) and calculated XRD powder diffraction patterns of α, γ, σ and χ phases.

Crystallographic orientation at the α/γ interfaces Sato and Kokawa investigated the formation of σ phase in duplex stainless steel joints welded by autogenous GTAW [27]. A strong association was found between the deviation of the Kurdjumov-Sachs orientation relation (∆θK-S), between neighboring grains of α and γ, and the tendency for precipitation at the interface. In the GTAW welding, neighboring α and γ grains tended to have low ∆θK-S, which was suggested to explain retardation of the kinetics of σ phase formation in the weld metal. In later work, Haghdadi et al. conducted five parameters characterization for the α/γ interfaces of DSS cooled continuously to 900 °C [28]. It was reported that evaluation of the deviation between disorientation of the grains and the K-S orientation relation, as performed by Sato and Kokawa, would result in an incomplete analysis, because planes close to the 15

orientation relation might not be orientated in the direction of the corresponding α/γ interface. Haghdadi et al. considered a population of interfaces with ∆θK-S < 10° and found that for 86% of the interfaces with low ∆θK-S that presented σ phase precipitation, the coherent plane pair {111}γ//{110}α was not parallel to the corresponding interface. In both of these previous studies, ∆θK-S was calculated by measuring the deviation between the disorientation between the α and γ grains and the disorientation of the ideal K-S OR. This approach for determining ∆θK-S only provides the deviation between the nearest {111}<110>γ//{110}<111>α pairs of planes and directions. However, it is possible that there could be pairs of planes and directions with low ∆θK-S that are closer to being parallel with the direction of the interface. Figure 8(a) shows the example of an interface with two pairs of planes close to the K-S OR, highlighting that the pair of planes with smaller ∆θK-S presents a greater slope, relative to the direction of the interface. In order to quantify the deviation of a pair of planes relative to the interface, measurement was made of an angle φ indicating the rotation required for a hypothetical average plane between the planes {111}γ and {110}α to be oriented in the direction of the α/γ interface, as shown in Figure 8(b). Figure 9 shows histograms of ∆θK-S for the pairs of planes with smaller ∆θK-S, for the BM and SZ regions studied. It can be seen that the microstructure resulting from the FSW presented the development of α/γ interfaces with more random ∆θK-S, with the FSW tending to reduce the fraction of interfaces with low ∆θK-S. In addition, the 2D histograms of ∆θK-S and φ were analyzed in order to identify possible trends in the angle of deviation between the {111}γ//{110}α planes and the interface. Figure 10 shows 2D histograms of ∆θK-S and φ for the SZ and BM of the UNS S32205. For both the BM and the SZ, φ presented a homogeneous distribution, indicating that analysis of the histograms shown in Figure 9 was sufficient to infer that after the FSW, there was a decrease of the fraction of interfaces with low ∆θK-S. For the SZ-AS, SZ-RS, and TMAZ-AS, the distributions of ∆θK-S were similar to those for the corresponding SZ.

16

Figure 8. (a) Example of an interface with two pairs of planes close to the K-S OR; (b) illustration of the measurement of φ and ∆θK-S.

Figure 9. Histograms of the distribution of the deviation from the Kurdjumov-Sachs orientation relation (∆θK-S) at the α/γ interfaces of the microstructures of the base metals and the stir zones.

Figure 10. 2D histograms of the distributions of ∆θK-S and φ for the BM and SZ of the UNS S32205.

17

Quantitative analysis of the transformation kinetics Digital image analysis of the micrographs of the BM and the SZ of the UNS S32205 was performed in order to quantify the volumetric fractions of intermetallic phases. The data obtained were fitted using the Avrami model (Equation 1). The data for the UNS S32750 and the UNS S32760 were not treated, because the values presented statistical equivalence and did not allow convergence of the Avrami model. Figure 11 shows the evolution of the volumetric fraction of the intermetallic phases during the isothermal treatment, together with the curves (obtained using the Avrami model) describing the kinetics for each of the regions analyzed. Table 3 presents the values of K, n, and R2 obtained following fitting of the data for each microstructural region. The results obtained for the quantitative analysis and the fitting of the experimental data were coherent with faster kinetics of formation of intermetallic phases in the SZ, compared to the corresponding BM. It was also observed that the regions with greater grain refinement presented faster kinetics of formation of intermetallic phases, compared to regions with coarser grains. In the case of the SZ-AS, fitting of the experimental data was not performed, due to the statistical equivalence of the data and the decreasing trend observed for longer isothermal treatment times, leading to non-convergence of the Avrami model. Nevertheless, the results discussed above suggested that the kinetics was even faster in the SZ-AS, which was corroborated by the rapid consumption of ferritic sites (Figure 3) and the results of the ferritoscope measurements (Figure 6), as well as by the fact that the microstructure of this zone presented a larger volumetric fraction of intermetallic phases after only 5 min of isothermal treatment.

18

Figure 11. Evolution of the volumetric fraction of intermetallic phases in the BM and SZ of the UNS S32205 during isothermal treatment, and Avrami curves obtained by fitting of the experimental data.

Table 3. Quantification of the volumetric fraction of intermetallic phases for the UNS S32205 treated isothermally at 850 °C, and fitting of the experimental data using the Avrami model.

Volumetric fraction of intermetallic phases (%)

Time (min)

BM

SZ-RS

SZ

SZ-AS

5

--

19 ± 2

19 ± 1

25 ± 2

10

0.9 ± 0.2

20 ± 1

22 ± 2

26 ± 2

30

6.6 ± 0.5

--

--

--

60

12 ± 2

25 ± 2

24 ± 1

25 ± 1

Results of fitting using the Avrami model -1

(4 ± 1)×10-3

(7.59 ± 0)×10-3

--

1.3 ± 0.2

0.23 ± 0.03

0.30 ± 0.06

--

0.97

0.91

0.94

--

K (s )

(1.9 ± 0.2)×10

N 2

R

-4

In its usual form, as used here, the Avrami model is a generalization of several mathematical models for modeling the kinetics of transformation of solid state phases. The models that employ the Avami equation [30] include those of Johnson-Mehl [34], Cahn [35], Wert and Zener [36], and others [37,38], which use different boundary conditions for modeling the transformation kinetics. The mathematical models for describing solid state phase transition kinetics generally include a set of pre-exponential variables, represented by the variable K in the generalized form of the model. The variables that result in K include the rates of nucleation and growth of the new phase particle [30,34,35]. It should also be noted that with increase of the specific surface of 19

α/γ interfaces, there is an increase in the number of intermetallic phase nucleation sites, hence increasing the nucleation rate.

Discussion The microstructural analyses indicated that the microstructures produced by FSW had greater propensity to form intermetallic phases, compared to the microstructures of the corresponding BM. The results showed that even for the steels that developed little (UNS S32205) or no (UNS S32760) intermetallic phase fraction after 5 min exposure at 850 °C, the FSW SZ microstructures of all the materials analyzed presented advanced stages of precipitation of intermetallics after 5 min at 850 °C. The UNS S32205 steel generally presented slower kinetics of intermetallic phase formation, while the UNS S32750 and S32760 superduplex steels presented widespread intermetallic phase formation in the BM after 30 min and in the SZ after only 5 min, which could be attributed to the substantial grain refinement induced by the FSW process. The BM of the UNS S32760 presented significantly slower kinetics of intermetallic phase formation, compared to the BM of the UNS S32750, despite the fact that these materials have very similar chemical compositions. The significant difference between the BM behaviors for UNS S32760 and UNS S32750 could be explained by the presence of W and Cu in UNS S32760, with the lower diffusibility of W retarding formation of the χ phase and consequently also retarding formation of the σ phase [29]. Despite the slower transformation kinetics of the BM of UNS S32760, the SZ of this steel showed rapid formation of intermetallic phases, with almost complete consumption of the α phase in all the microstructural regions after 5 min exposure at 850 °C. The probable reason for the difference between the results for the BM and SZ of UNS S32760, after treatment for 5 min, was the pre-existence of some particles or nuclei of intermetallic phases and carbides [17,21]. Previous studies performed with this material indicated that secondary phase formation occurred during the microstructural development of the SZ and the SZRS, with the pre-existing χ phase nuclei continuing to grow and the carbide particles acting as sites for heterogeneous nucleation, in addition to the interfaces present. The previous presence of secondary phases could also explain the slower formation of intermetallic phases in the SZ-AS of the UNS S32760, compared to the central SZ, which was very different to the observed behaviors of the other steels studied. In a 20

microscopy study of the formation of intermetallic phases, Santa Cruz et al. found that the SZ-AS was the only weld region of UNS S32760 where no intermetallic phase formation was detected [21]. The slower kinetics of intermetallic phase formation for UNS S32205, compared to the other steels, was due to the lower content of alloying elements, which reduced the stability of the σ phase, relative to the α phase. The microstructural refinement induced by FSW accelerates formation of the σ and χ phases, because these phases present preferential heterogeneous nucleation at the α/γ interfaces [26]], while the specific surface of the α/γ interfaces increases with the degree of refinement, resulting in a greater quantity of nucleation sites [25]. The microstructures resulting from the FSW were considerably more refined than those of the BM, with this being the main factor that led to faster kinetics of intermetallic phase formation. The significantly faster kinetics of intermetallic phase formation in the DSS SZ showed that short exposure to high temperature may severely compromise FSW SZ performance. In addition to refinement of the microstructure, the changes of the crystallographic ratios of the α/γ interfaces after the FSW process also favored faster formation of intermetallic phases. The microstructures resulting from the FSW showed increase of the specific surface of α/γ interfaces, as well as decrease of the density of α/γ interfaces with low ∆θK-S. Hence, there was not only an increase of the quantity of nucleation sites, but also an increase of the quantity of nucleation sites with high tendency for rapid precipitation of intermetallic phases. Interfaces that follow the K-S OR show coherence between the α and γ grains [52]. The coherence at the interface reduces the energy derived from the crystallographic mismatch, only presenting energy associated with local deformation of the crystalline structures, caused by the difference between the interatomic distances in planes {111} of the austenite and {110} of the ferrite. Despite the energy associated with deformation of the crystalline lattice, the energies associated with coherent interfaces are lower than for non-coherent interfaces, so they have less propensity to act as heterogeneous nucleation sites [27]. Hence, it is unlikely that there would be a substantial fraction of interfaces that perfectly followed an orientation relation (in the case of the K-S OR, with ∆θK-S = 0°). Interfaces with small ∆θK-S values tend to show the presence of crystalline defects such as mismatches and stacking faults, with the greatest part of the interface presenting coherence [53]. As ∆θK-S increases, so does the quantity of defects, 21

so interfaces with high ∆θK-S present large numbers of defects, associated with higher energies. Therefore, after a certain value of ∆θK-S, increase of the deviation does not have a significant effect on the tendency for precipitation at the interface, considering deviations of around 15º [27], which explains why some previous works have specifically studied low deviations [28]. The results for the distributions of ∆θK-S and φ of the α/γ interfaces showed that for the conditions analyzed, φ presented a homogeneous distribution, probably because the development of the microstructures was mainly due to recrystallization processes. In the BM, there was partial recrystallization during solubilization after the cold rolling step of the manufacturing process, while the SZ underwent complete recrystallization during the FSW [17]. In processes that involve rapid cooling or solidification, there is expected to be a higher density of interfaces with low φ values, since transformations that are more distant from thermodynamic equilibrium tend to show a higher density of coherent interfaces [54]. Hence, although Haghdadi et al. provided good arguments concerning the analysis performed by Sato and Kokawa, it is expected that material welded by GTAW should present a high density of interfaces with low ∆θK-S. Therefore, it would be expected that a substantial proportion of the 80% fraction of interfaces with low ∆θKS,

analyzed by Sato and Kokawa, should present low φ. It is also possible that interfaces

with low ∆θK-S and φ determined by 2D EBSD measurements may be non-coherent interfaces, since only 3D EBSD patterns can enable three-dimensional characterization of φ, in order to confirm coherence of the interface. The 2D EBSD measurement of φ, together with ∆θK-S calculated from misorientation among the grains, can be used as an estimate of the density change of interfaces close to coherence, since the φ value and high OR deviations can indicate when the interface is not coherent. The results obtained for quantification of the phase volumetric fractions confirmed the rapid kinetics of intermetallic phase formation in the SZ, as well as the differences in the kinetics for the SZ-AS, SZ, and SZ-RS regions. The fastest kinetics was in the SZAS, followed by the SZ and the SZ-RS, which was due to more intense refinement of the microstructure in the SZ-AS [17]. The Avrami model showed good correlation with the experimental data and could therefore be considered suitable for modeling the kinetics of formation of intermetallic phases in DSS, as proposed elsewhere [14,41,42], as well as for analysis of a biphasic microstructure subjected to FSW, as presented in 22

the present work. It should also be noted that the formation of intermetallic phases in DSS involves double kinetic behavior, with one kinetic behavior at the beginning of intermetallic phase formation and another distinct kinetic behavior during the more advanced stages of transformation, due to the correlated formation of the σ and χ phases. During the initial stage of the transformation, there is a rapid increase of the volumetric fraction of intermetallic phases, due to rapid heterogeneous nucleation of the σ phase next to the χ phase. During the final stage, there is progressive slow growth of the volumetric fraction of intermetallic phases, with the early growth of nucleated σ particles being the main feature of this stage [40–42]. Fitting of the experimental data using the Avrami equation enabled description of the initial stage of the transformation, with fitting of the BM data, as well as the second stage of the transformation, with fitting of the data for the SZ microstructures.

Conclusions The findings of this work showed that the welded joints of UNS S32205, S32750, and S32760 steels presented much faster intermetallic phase formation, compared to the corresponding base metals, when subjected to isothermal treatment at 850 °C. For all the materials studied, after only 5 min at 850 °C, the SZ presented advanced stages of precipitation of intermetallic phases. Therefore, for DSS, the SZ obtained by FSW is highly liable to deterioration of its properties if exposed to high temperatures. The accelerated formation of intermetallic phases after friction stir welding is associated with severe grain refinement and a decreased density of α/γ interfaces with low ∆θK-S. Crystallographic characterization of the α/γ interfaces indicated that FSW reduced the density of α/γ interfaces with low ∆θK-S. Fitting of the experimental data using the Avrami equation resulted in fits with high correlation, evidencing the significantly accelerated kinetics of formation of intermetallic phases in joints welded by FSW.

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Acknowledgments Financial support for this work was provided by FACEPE (APQ-0715-3.05/14), CNPq (181170/2016-7, 103395/2017-2, 155369/2017-2), ANP/FINEP (Ref. 2659/09). The authors are indebted to Aperam South America, Outokumpu, and Weir Materials for the donation of materials, and to IPT for the provision of SEM infrastructure. Data Availability The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study. References [1] J. Charles, Duplex Stainless Steels - a Review after DSS ‘07 held in Grado, Steel Research International. 79 (2008) 455–465. https://doi.org/10.1002/srin.200806153. [2] P. Boillot, J. Peultier, Use of Stainless Steels in the Industry: Recent and Future Developments, Procedia Engineering. 83 (2014) 309–321. https://doi.org/10.1016/j.proeng.2014.09.015. [3] Y. Guo, J. Hu, J. Li, L. Jiang, T. Liu, Y. Wu, Effect of Annealing Temperature on the Mechanical and Corrosion Behavior of a Newly Developed Novel Lean Duplex Stainless Steel, Materials (Basel). 7 (2014) 6604–6619. https://doi.org/10.3390/ma7096604. [4] L. Zhang, W. Zhang, Y. Jiang, B. Deng, D. Sun, J. Li, Influence of annealing treatment on the corrosion resistance of lean duplex stainless steel 2101, Electrochimica Acta. 54 (2009) 5387–5392. https://doi.org/10.1016/j.electacta.2009.04.023. [5] K.H. Lo, C.H. Shek, J.K.L. Lai, Recent developments in stainless steels, Materials Science and Engineering: R: Reports. 65 (2009) 39–104. https://doi.org/10.1016/j.mser.2009.03.001. [6] F. Marques, W.M. da Silva, J.M. Pardal, S.S.M. Tavares, C. Scandian, Influence of heat treatments on the micro-abrasion wear resistance of a superduplex stainless steel, Wear. 271 (2011) 1288–1294. https://doi.org/10.1016/j.wear.2010.12.087. [7] C.R. de Farias Azevedo, H. Boschetti Pereira, S. Wolynec, A.F. Padilha, An overview of the recurrent failures of duplex stainless steels, Engineering Failure Analysis. 97 (2019) 161–188. https://doi.org/10.1016/j.engfailanal.2018.12.009. [8] Q. Sun, J. Wang, H. Li, Y. Li, Y. Hu, J. Bai, P. Han, Chi Phase after Short-term Aging and Corrosion Behavior in 2205 Duplex Stainless Steel, Journal of Iron and Steel Research, International. 23 (2016) 1071–1079. https://doi.org/10.1016/S1006-706X(16)30159-5. [9] I. Calliari, G. Straffelini, E. Ramous, Investigation of secondary phase effect on 2205 DSS fracture toughness, Materials Science and Technology. 26 (2010) 81– 86. https://doi.org/10.1179/174328408X388103. [10] M. Pohl, O. Storz, T. Glogowski, σ-phase morphologies and their effect on mechanical properties of duplex stainless steels, IJMR. 99 (2008) 1163–1170. https://doi.org/10.3139/146.101738. 24

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Highlights

•

Stirred zones presented much faster intermetallic phase formation

•

Intermetallic phase formation kinetics varied along stirred zone sides

•

α/γ interfaces crystallography of weld joints favored intermetallic nucleation

Author contributions section

Igor J. Marques Conducted the heat treatments and metallographic preparing of samples; performed SEM characterization of heat treated friction stir welds and some base metal samples; performed Avrami equation fitting of data obtained by digital image analysis; postprocessed EBSD data and developed algorithms to perform the interface crystallography evaluation.

Flávio J. Silva Helped with heat treatments execution; contributed Performed XRD and Ferritoscope measurements; contributed with microstructural analysis and digital image analysis.

Tiago F. A. Santos Conceived the idea of evaluating intermetallic phase precipitation kinetics on friction stir welds; performed the friction stir welding operations; performed the EBSD characterization of the friction stir welds and base metals; conducted some SEM characterization of heat-treated base metals; supervised the findings of this work.

All the authors contributed to manuscript writing.

Declaration of interests

The authors confirm that there are no actual or potential conflict of interests that could inappropriately influence the work entitled “Rapid precipitation of intermetallic phases during isothermal treatment of duplex stainless steel joints produced by friction stir welding”, submitted to Journal of Alloys and Compounds.

MSc. candidate: Igor Jordão Marques

Prof. Dr. Flávio José da Silva

Prof. Dr. Tiago Felipe de Abreu Santos