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Vol. 8. Issue 5.
Pages 3936-3949 (September - October 2019)
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Vol. 8. Issue 5.
Pages 3936-3949 (September - October 2019)
Original Article
DOI: 10.1016/j.jmrt.2019.07.002
Open Access
Remarks on the evolution and performance of the different austenite morphologies at the simulated HAZs of a 2205 duplex stainless steel
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E.V. Moralesa,b,
Corresponding author
valencia@uclv.edu.cu

Corresponding author.
, J.A. Pozoc, L. Olayab, E. Kassabb, J.A.C. Poncianod, K. Ghavamie, I.S. Bottb
a Physics Department, Central University of Las Villas, Santa Clara, VC, CP 54830, Cuba
b Chemical and Materials Engineering Department, Pontifical Catholic University of Rio de Janeiro/PUC-Rio, Rua Marques de S. Vicente 225, Gavea, Rio de Janeiro, RJ, CEP 22541900, Brazil
c Welding Research Center, Central University of Las Villas, Santa Clara. VC, CP 54830, Cuba
d COPPE, Department of Metallurgical and Materials Engineering, Corrosion Laboratory, Ilha do Fundão, UFRJ, IF-B, Caixa Postal 68505, CEP21941-972, Rio de Janeiro, RJ, Brazil
e Civil Engineering Department, Pontifical Catholic University of Rio de Janeiro/PUC-Rio, Rua Marques de S. Vicente 225, Gávea, Rio de Janeiro, RJ, CEP 22453-900, Brazil
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Tables (5)
Table 1. Chemical composition (wt.%) of the DSS (UNS S31803).
Table 2a. Austenite (Vγ) and ferrite (Vδ) volume fractions at different cooling times. Austenite volume fractions taking into account their morphologies. Heat input (H), the ξIGA, ξGBA, and ξWA are the ratios between the volume fraction of the different austenite morphologies and the total austenite volume fraction.
Table 2b. Microstructural parameters of the simulated HAZ at different cooling times. Average grain size of ferrite (d), average width of the GBA (AGBA), average length of the WA laths (lWA), average radius of the IGA precipitates (RIGA) and average width of the WA laths (AWA).
Table 3. Pitting corrosion behavior in the as-received DSS (UNS S31803) and simulated HAZ of the same steel at 60 °C for different Δt12/8.
Table 4. Diffusive distances of chromium and nitrogen atoms in the ferrite matrix during the continuous cooling between 1200 °C and 500 °C.
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Abstract

Some features of the austenite micro structural evolution have been reviewed at the simulated HAZs (Heat Affected Zones) in a duplex stainless steel. The evolution of the different morphological forms of austenite was quantified in the microstructures for each simulated heat input. A detailed study of the micro structural parameters in the present study, allowed justifying the performance of the volume fractions corresponding to the three austenite morphologies on the HAZ mechanical properties and pitting corrosion resistance. Transmission electron microscopy showed the predominance of specific orientation relationships between the small intragranular austenite precipitates and ferrite considering the nitrogen content in the alloy and the transformation route.

Keywords:
Duplex stainless steel
Intragranular austenite
HAZ
Optical microscopy
Transmission electron microscopy
Mechanical properties
Full Text
1Introduction

Duplex stainless steels (DSSs) are widely used as structural materials in the chemical, petrochemical, marine and paper industries because of the attractive combination of their mechanical properties, weldability and corrosion resistance in various types of environments [1–5]. These attractive qualities are due to a balance between the ferrite/austenite (δ/γ) phases in their microstructures [6,7]. During welding, this phase ratio (δ/γ) tends to deviate from 1:1 at the Heat Affected Zone (HAZ) influencing the mechanical properties and the pitting corrosion resistance. These microstructural changes that occur at the HAZ during welding are induced by the phase transformations during cooling [8]. Therefore, a high heat input (slow cooling rate) used during welding promotes a high reformed austenite fraction at the HAZ giving a more favourable phase balance [9]. However, it may also promote the precipitation of harmful intermetallic phases such as σ and χ, which will lead to a severe reduction in toughness [10,11]. On the contrary, the “quenched-in” nitrides (Cr2N and CrN) can precipitate at lower heat input depending of the cooling rate and nitrogen content in the alloy [12–14].

The effect of different heat inputs on the microstructure and on the localized corrosion behavior of the simulated HAZ in the UNS S31803 duplex stainless steel (DSS) has been widely discussed in the literature. Most articles about duplex steels explain the mechanical behavior and resistance to localized corrosion in the HAZ considering only the total austenite and ferrite volume fractions. Few articles describe the influence of the cooling rate on the change of the different austenite morphological forms in the HAZ. Thus, Yang et al. [8] have shown that by increasing the cooling time between 800º and 500 °C (Δt8/5), the volume fractions of the three morphological forms of austenite (grain boundary austenite (GBA), intragranular austenite (IGA) and Widmanstätten austenite (WA)) in the 2205 DSS HAZ increased with predominance of the GBA. Muthupandi et al. [15] have affirmed that shorter cooling times in this HAZ generated a greater IGA fraction than GBA. Liou et al. [16] showed that the amounts of WA and IGA increased, although the GBA in the HAZ was almost fully replaced by the WA at longer cooling times in a similar DSS (UNS S31803) with 0.096 wt.% of N. On the contrary, Liao [14] and Wang et al. [17] reported that at longer cooling times both the GBA and IGA were observed, being the GBA volumetric fraction the major in the HAZ. Koruda et al. [18] established that the GBA was thickened and the formation of IGA was promoted by increasing the cooling time. A recent article has analyzed the effect of input heat (or cooling time) on austenite microstructural evolution of the simulated HAZ. This paper shows that with the increase in heat input (longer cooling times), austenite gradually changes from GBA to WA and IGA paying attention to the WA decomposition [19].

Although the above papers describe qualitatively as the different cooling rates affect the volume fractions of the three austenite morphological forms in the HAZ for the DSS (UNS S31803), the relationship between the cooling time and the volume fractions of the different austenite morphologies is still not clear.

There are also no reports in the literature of whether any of these austenite morphologies in the HAZ microstructure contributes with greater or lesser intensity to the mechanical properties and localized corrosion resistance.

Recently, much attention has been devoted to the character of the interfaces between ferrite and austenite in these HAZ microstructures which extensively influences in their properties and resistance to the localized corrosion. In this sense, Karlsson and Börjesson [20] has reported that random orientation relationships (ORs) between the IGA precipitates and ferrite is more likely for shorter cooling times and compositions promoting nucleation at low temperatures in duplex stainless steels. On the contrary, other authors [21] found that there was a higher tendency towards interfaces with specific ORs when the cooling time was decreased. Then, the study of the specific ORs between the small IGA precipitates and ferrite for current nitrogen content in the alloy and different transformation routes in the HAZ calls attention.

The purpose of this work is to delve in the characterization of the austenite fractions with different morphologies in the HAZ microstructure at different cooling rates. Then, to analyze the influence of the austenite microstructural evolution in simulated HAZs on some mechanical properties and resistance to localized corrosion. Also, in the present work the predominance of interfaces with specific ORs between the small IGA precipitates and ferrite was assessed according to the composition and phase transformation route in these simulated HAZs.

2Experimental procedure

A commercial DSS (UNS S31803) available in the form of an extruded tube of 200 mm outer diameter and 8.15 mm wall thickness was used. The chemical composition of this alloy is listed in Table 1. Samples, longitudinal and transversal to the rolling direction of the as-received steel, prior to weld simulation, were mechanically polished and electrolytically etched in a 30% KOH solution at 6 V for 20 s. The microstructures of these samples were examined by optical microscopy (OM) and scanning electron microscopy (SEM).

Table 1.

Chemical composition (wt.%) of the DSS (UNS S31803).

Elements  Cr  Ni  Mo  Mn  Si  Cu 
UNS S31803  0.026  22.6  5.23  3.2  0.85  0.025  0.004  0.49  0.13  0.12–0.13 

A Gleeble 3800 thermo-mechanical simulator performed the HAZ simulations using cylindrical bars of 101.5 mm in length and 6 mm of diameter. These bars had a reduction at the central part corresponding to the region where the HAZ was simulated as it is shown in Fig. 1. This reduction is made to guarantee that samples break in the region corresponding to the simulated HAZ during the tensile test at room temperature. The applied thermal cycles are schematically shown in Fig. 2. The programmed thermal cycles in the Gleeble simulator are in correspondence with the Rykalin-2D model.

Fig. 1.

Sub-dimensioned sample for the tensile test where a reduction was performed in the region corresponding to the HAZ.

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Fig. 2.

The Gleeble simulated thermal cycles corresponding to different heat inputs (H). The heating rate was 350 °C/s and the corresponding cooling rates between 1350 °C and 500 °C were 13.8 °C/s(A), 7.8 °C/s(B), 5 °C/s(C) and 3.5 °C/s(D), respectively.

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A peak temperature of 1350 °C was reached at a rate of 350 °C/s. A holding time of 2 s at this temperature followed by cooling at different rates up to 500 °C to simulate different heat inputs and then water quenched. The typical temperature range accepted for the austenite reformation and precipitation of intermetallic compounds is between 1200–800 °C [8] in these DSSs. However, the cooling time from 800 to 500 °C (Δt8/5) is relatively easier to be measured more accurately than that from 1200 to 800 °C (Δt12/8). Thus, in the following we will only refer to Δt12/8 that can be calculated by measuring the Δt8/5 according to Eq. 1 in Ref. [8].

The microstructures of the simulated HAZs were examined using OM (ZEISS Axioplan2 Imaging), SEM (JSM-6510LV JEOL microscope operated at 20 kV) and transmission electron microscopy (TEM-Jeol 2010 operated at 200 kV). The (δ/γ) phase ratios and the austenite morphological forms in the HAZ microstructures were obtained processing 30 optical images for each heat input using the AxioVision 4.8.2 and Fiji softwares.

Thin-foil samples for TEM were prepared by cutting a 200-μm-thick slice from the sectioned HAZ samples, punching 3-mm diameter disks from the slice, and electropolishing the disks in a Tenupol-5 apparatus, using a solution of 10% perchloric acid and 90% ethylic alcohol at 0 °C and 20 V.

Tensile tests were performed at a crosshead speed 3 mm/min at room temperature. The above cylindrical bars (Fig. 1) were sampled at the longitudinal section of the tube. Vickers microhardness measurements were carried out on the different morphological forms of austenite and ferrite grains randomly in the cross-sections of specimens. The specimen surfaces were carefully polished before indenting. A load of 0.025 kg (HV0.025) was applied during 10 s.

The localized corrosion resistance evaluation was carried out through of anodic potentiodynamic polarization measurements in a glass three-electrode cell containing 3.5 wt.% NaCl solution at temperatures of 25 and 60 °C. A platinum wire and a saturated calomel electrode (SCE) were used as the counter and reference electrodes, respectively. The specimens of DSS (UNS S31803) acting as working electrode were embedded in epoxy resin with an exposed working area of 0.25 cm2. Prior to each electrochemical measurement was conducted, the working electrode was ground and polished, degreased with ethanol, rinsed with distilled water and dried. Also, before each corrosion test, the open circuit potential (OCP) was recorded for 30 min to stabilize the corrosion potential. The potentiodynamic polarization measurements were carried out at a scan rate of 1 mV/s from the OCP to 1000 mV SCE above the OCP. To ensure the reproducibility of the results, experiments were repeated at least three times under the same experimental condition.

The critical pitting temperature (CPT) was obtained according to the ASTM G 150-99 standard [22]. The 3.5 wt.% NaCl solution was heated from 0 °C to higher temperatures at a rate of 1 ± 0.3 °C/min. A constant potential of +750 mV SCE was applied while the current was monitored during the temperature scan. CPT measurements of the same specimen were carried out at least five times, and the average data were chosen as CPT in this paper.

3Results3.1Microstructural characterization of the HAZ by optical microscopy

In this study, a statistical analysis of the processed digital images is applied to quantify the different austenite morphologies in 30 optical micrographs for each programmed heat input (or Δt12/8). The evaluation of the reformed austenite fractions with different morphologies is carried out for each ferrite grain in the optical micrographs. Each micrograph collected the information of approximately 5–10 ferritic grains depending on the cooling rate. This provided information of about 200 ferritic grains oriented randomly. Thus, the uncertainty by considering as IGA precipitates the WA laths (or needles) that have been cut normal to its growth direction on the micrograph plane is reduced. Also, a determined aspect ratio among the austenite precipitates within the ferrite grains was taken into account. Besides, the fragments from the WA laths formed at lower temperatures were also considered as IGA precipitates. We emphasized that the OM and SEM analysis did not detect intermetallic phases in the simulated HAZ microstructures at the tested cooling times.

The WA selection in the micrographs was done following the criterion that the major axis or growth direction of the WA laths is not random in each ferrite grain. As it can be seen in Fig. 3, all the WA laths are almost identically oriented, growing from a particular austenite grain boundary with approximately the same shape, or within the grain itself.

Fig. 3.

Original and digital processed optical images of the simulated HAZs in the DSS (UNS S31803) with different cooling times, Δt12/8, and the corresponding heat inputs H. In these digital processed images, the IGA precipitates are in red, the WA laths are in blue and the GBA is in white-grey color. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

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Fig. 3 shows the processed images of the different morphological forms of austenite at the HAZ microstructures of the DSS (UNS S31803) at different cooling times or heat inputs. Tables 2a and 2b show the austenite and ferrite volume fractions as well as the austenite volume fractions taking into account their morphologies (a), and other microstructural parameters of the simulated HAZ at different cooling times (b).

Table 2a.

Austenite (Vγ) and ferrite (Vδ) volume fractions at different cooling times. Austenite volume fractions taking into account their morphologies. Heat input (H), the ξIGA, ξGBA, and ξWA are the ratios between the volume fraction of the different austenite morphologies and the total austenite volume fraction.

H(KJ/mm)  Δt8/5(s)  Δt12/8(s)  Vγ(%)  Vδ(%)  VIGA(%)  VGBA(%)  VWA(%)  ξIGA  ξGBA  ξWA 
1.5  43  15  28.2  71.8  8.9  14.0  5.3  31.4  49.7  19.0 
76  25  32.4  67.6  9.0  15.8  7.6  27.7  48.8  23.5 
2.5  118  40  34.2  65.8  10.0  16.0  8.2  29.3  46.7  23.9 
170  60  34.6  65.4  11.9  13.6  9.0  34.7  39.1  26.2 
Table 2b.

Microstructural parameters of the simulated HAZ at different cooling times. Average grain size of ferrite (d), average width of the GBA (AGBA), average length of the WA laths (lWA), average radius of the IGA precipitates (RIGA) and average width of the WA laths (AWA).

H (kJ/mm)  Δt8/5 (s)  Δt12/8 (s)  d (μm)  AGBA(μm)  lWA(μm)  RIGA(μm)  AWA(μm) 
1.5  43  15  104.6  8.9  28.5  2.1  3.1 
76  25  136.3  13.2  32.2  2.6  4.2 
2.5  118  40  160.8  14.4  33.3  2.8  4.2 
170  60  184.5  14.1  36.4  3.1  4.3 

The values shown in above tables have been depicted in Fig. 4(a) and (b) for a better appreciation about the evolution of the reformed austenite and ferrite volume fractions with the cooling time at the simulated HAZ. It can be seen from Fig. 3 that the amount of intragranular austenite (IGA) is concentrated and abounds in grains where the WA has not a strong presence.

Fig. 4.

(a) Volume fractions of the reformed austenite and ferrite respectively at different cooling times. Also, the corresponding austenite volume fractions according to their morphologies for different cooling times, Δt12/8. (b) Volume fractions of the austenite morphological forms with respect to the total austenite volume fraction for different cooling times, Δt12/8.

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Table 2a and Fig. 4(a) show the increase of the austenite fraction with increasing the cooling time; however, the rate of the increase of this austenite fraction is smaller for larger cooling times. It is observed in this Fig. 4(a), that the GBA volume fraction was predominant for all studied cooling rates, Δt12/8 between 15 and 60 s, and its predominance was greater for the shorter cooling times. Also, this GBA volume fraction reached a maximum for cooling times between 30 and 40 s according to Δt12/8. In Table 2a and Fig. 4(a) is appreciated that the intragranular austenite volume fraction is greater than that of the Widmanstätten austenite (WA) with increasing the cooling time. Then, as shown in Fig. 4(b) the ξGBA is reduced and the ξIGA is increased for higher austenite volume fractions.

Table 2b shows the increase of the average ferrite grain size. This ferritic grain growth was accompanied of the corresponding increase of the average length of the WA laths. Also (Table 2b), the average width of the GBA increases significantly at the shorter cooling times (between Δt12/8 = 15 s and Δt12/8 = 25 s), while for longer cooling times the average width of the GBA do not show substantial changes. The average width of the WA laths presents a similar behavior to the GBA width, since the average width of the WA laths did not change for longer cooling times. The average size of the IGA precipitates had a monotonous growth with the increase of the cooling time.

As the heating rate and holding time at the peak temperature were constant for all simulations, the increase of the average ferrite grain size (Table 2b) was due to the corresponding cooling times between 1350 °C and 1200 °C.

3.2Microstructural characterization of the HAZ by transmission electron microscopy

Thin foils corresponding to the HAZ samples with the shorter and longer cooling times (Δt12/8 = 15 and 60 s) were analyzed by TEM. Many small IGA precipitates (or fragmented WA laths), corresponding to approximately ten random ferrite grains for both cooling times, were carefully observed at different crystallographic orientations resulting in the predominance of boundaries with specific orientation relationships (ORs). Fig. 5 shows one of these intragranular austenite precipitates corresponding to the simulated HAZ of the DSS (UNS S31803) where the Δt12/8 was 15 s (the shorter cooling time). In this figure, the same precipitate was observed with two different tilts. The bright field image (BF) in Fig. 5(a), was observed with the beam near to the [01¯1]δ whereas in Fig. 5(b), this precipitate was oriented close to the edge-on view. This lath-shaped intragranular austenite exhibited a major planar facet which is denoted as the habit plane (HP), and several side facets where a curved interface between the HP and the major side facet is observed. This small precipitate (˜2.5 μm in length and 0.9 µm in wide) seems to have nucleated from the dislocations forest in the matrix. An array of structural dislocations spaced between 9 and 12 nm located in the broad faces of this precipitate can be observed.

Fig. 5.

(a) Bright field image (BF) of an IGA precipitate. The misfit dislocations can be observed at the broad faces (tilt 1.1°), (b) the same precipitate of IGA showing the major planar facet and the interwoven dislocations in the ferrite matrix (tilt 6.3°). The diffraction patterns of austenite and ferrite obtained with the same tilt conditions.

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The diffraction patterns of both phases obtained with the same observation conditions showed that this IGA precipitate has an specific orientation relationship close to the Nishiyama–Wasserman (N–W) variant, [21 1¯ ]γ//[0 1¯ 1]δ, with ferrite. Also all precipitates carefully analyzed had an aspect ratio very different from the unit (laths-like morphology).

TEM careful analysis of twenty (20) different grains observed in thin foils of samples from the simulated HAZ where the Δt12/8  = 15 s, showed precipitation of the Cr2N only in four grains indicating that the Cr2N precipitation was very scarce. It can be observed (Fig. 6) that all the Cr2N needles are oriented approximately in the same direction. These needles have 200–400 nm in size and the interparticle spacings are smaller than 0.2 μm.

Fig. 6.

TEM bright field (BF), dark field (DF) and diffraction patterns (DP) of Cr2N in two random ferrite grains corresponding to Δt12/8 = 15 s. The dark fields were obtained with the diffracted beams g = (0-11-1) and (01-11), respectively.

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At longer cooling times, Δt12/8 = 60 s, another striking feature in the TEM images can be observed at the simulated HAZ of this DSS. This feature is shown in Fig. 7, where the dislocation structure in some small ferrite grains has the cross-stitch form.

Fig. 7.

Dislocation structure in a small ferrite grain corresponding to a simulated HAZ of the DSS(UNS S31803) at the longest cooling time (Δt12/8 = 60 s). Diffraction pattern (δ DP) of this small ferrite grain.

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3.3Mechanical properties

Microhardness measurements of the different austenite morphologies, and ferrite matrix corresponding to the simulated HAZ of the DSS (UNS S31803) were performed at the different cooling times using a load of 0.025 kg (HV0.025), Fig. 8, which indentations had the ''pile-up'' morphology. The hardness of the WA laths was higher for the shorter cooling times (Fig. 8), while the hardness of the GBA showed a slight increase at longer cooling times. The ferrite hardness did not experiment significant changes and only a small increase of hardening was recorded at longer cooling times. In all microhardness measurements and independently of the austenite morphologies, the austenite hardness values were always higher than the hardness values measured in ferrite at the different cooling times (Fig. 8).

Fig. 8.

Vickers microhardness (HV0.025) of the different austenite morphologies and ferrite for the simulated HAZ microstructures in the DSS (UNS S31803).

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Fig. 9(a) shows the tensile tests realized at room temperature on samples corresponding to the HAZs with different cooling times. In each test, three HAZ samples with the same cooling time were broken to obtain the average yield and tensile strengths (YS and TS). Fig. 9(b) shows the average yield and tensile strengths of the simulated HAZ at different cooling times. We emphasized that the area below the curves in the tensile tests presented no appreciable changes (Fig. 9(a)) and the yield and tensile strength values did not suffer significant changes despite the growth of the ferrite grain size with increasing the cooling time (Fig. 9(b)).

Fig. 9.

(a): Tensile tests realized in samples corresponding to the HAZs with different cooling times. (a): 1.5 kJ/mm (or Δt12/8 = 15 s); (b): 2 kJ/mm (or Δt12/8 = 25 s); (c): 2.5 kJ/mm (or Δt12/8 = 40 s) and (d): 3 kJ/mm (or Δt12/8 = 60 s). In this Figure, σ is the stress (MPa) and ε is the tensile strain (%). (b): The yield and tensile strengths (YS-TS) in MPa, corresponding to the simulated HAZs at different cooling times (Δt12/8) for the DSS (UNS S31803).

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3.4The pitting corrosion test

The pitting corrosion resistance in the as-received DSS (UNS S31803) and simulated HAZ with Δt12/8 = 15 s and Δt12/8 = 60 s for the same steel was measured. Fig. 10 shows the potentiodynamic polarization curves in 3.5 wt.% NaCl solution at 25 °C (a) and 60 °C (b). Temperature of 60 °C is above the CPT of all the tested HAZs. Fig. 10(c) shows the average critical pitting temperatures for the as-received DSS (UNS S31803) and the simulated HAZs of this steel with Δt12/8 = 15 s and Δt12/8 = 60 s.

Fig. 10.

Potentiodynamic polarization curves in 3.5 wt.% NaCl solution at 25 °C (a) and 60 °C (b) in the as-received DSS(UNS S31803) and simulated HAZ of the same steel with Δt12/8 = 15 s and Δt12/8 = 60 s. (c) CPT for the corresponding as-received DSS(UNS S31803) and simulated HAZ samples (Δt12/8 = 15 s and Δt12/8 = 60 s).

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The room temperature tests (25 °C), Fig. 10(a), showed a wide passive region from −200 mV to 800 mV with no increase of the current density for the DSS (UNS S31803) in as-received condition and in the simulated HAZ with different heat inputs (or Δt12/8 as cooling time). This behavior indicates an excellent localized corrosion resistance in the NaCl solution at 25 °C.

Different pitting corrosion resistance is observed when temperature is increased to 60 °C (above the CPT). An abrupt increase of the current density occurs in all three conditions, indicating loss of the integrity of the passive film and the presence of localized corrosion (Fig. 10(b)). Lower pitting corrosion resistance occurs in specimens with simulated HAZ, since they present a lower pitting potential, (Epit), and a lower pit nucleation resistance, (ΔE), in the polarization tests (Table 3). The ΔE is the difference between Epit and Ecorr (ΔE = Epit − Ecorr) and represents the amplitude of the passive domain. So that, a higher ΔE indicates higher resistance to localized corrosion. According to Table 3, the simulated HAZ with shorter cooling time, (Δt12/8 = 15 s), presents a lower pitting corrosion resistance than the simulated HAZ with longer cooling time, (Δt12/8 = 60 s). Fig. 10(c) shows the average CPT values measured in the as-received and simulated HAZs with shorter and longer cooling times in the DSS (UNS S31803). It is appreciated that the average CPT values in the simulated HAZs decreased in more than 10 °C respect to the as-received duplex steel CPT value. Average CPT values are: 56.8 ± 0.1 °C for the as-received duplex steel, 47.6 ± 0.2 °C for the simulated HAZ with Δt12/8 = 60 s and 44.2 ± 0.1 °C for the simulated HAZ with Δt12/8 = 15 s. The CPT results showed a decrease in the localized corrosion resistance in samples where the HAZs were simulated.

Table 3.

Pitting corrosion behavior in the as-received DSS (UNS S31803) and simulated HAZ of the same steel at 60 °C for different Δt12/8.

DDS(UNS S31803)  Ecorr [mV]  Epit [mV]  ΔE [mV] 
As-received  −217 ± 5  440 ± 9  657 
HAZ (Δt12/8 = 15 s)  −138 ± 6  261 ± 3  399 
HAZ (Δt12/8 = 60 s)  −159 ± 5  272 ± 6  431 
4Discussion4.1Evolution of the different austenite morphologies at the HAZ

Energetic criteria for heterogeneous nucleation in the solid state, has established that during continuous cooling of the DSS (UNS S31803) HAZ, the first reformed austenite is the allotriomorphic austenite (GBA) [8,14–18,23]. Depending on the cooling rate, this allotriomorphic austenite will redistribute the N and /or the substitutional alloying elements with ferrite. This first GBA changes the chemical composition of the matrix reducing the austenite-forming elements and consequently the sites for the austenite nucleation. Therefore, a higher driving force or undercooling is necessary for the subsequent nucleation of the reformed austenite. At lower temperatures (1000–800 °C) the Widmanstätten austenite is formed from the allotriomorphic GBA basically. In sequence the WA grows along the invariant line of transformation for each randomly oriented ferrite grain (considered as single crystal [24]) having orientation relationships close to the Kurdjumov–Sachs (K–S) or Nishiyama–Wasserman (N–W) with the ferrite [24–27]. Nucleation of this WA can take place in a wide interval of temperatures according to its chemical composition [28].

The WA nucleated at higher temperatures (or at longer cooling times) can better redistribute the chemical elements than that formed at lower temperatures. This WA nucleated at low temperature remains saturated in Cr and Mo and becomes unstable. This unstable WA at low temperatures, breaks down into small pieces of diamond-shape austenite [19,28], and could be identified as IGA in the micrographs. The fragmentation of these WA laths formed at lower temperatures [28] could contribute to the increase of the intragranular austenite volume fraction respect to the WA fraction at the tested cooling times. Fig. 11 shows a WA side-plate during the fragmentation process at Δt12/8 = 15 s.

Fig. 11.

TEM bright field image. Fragmentation in progress corresponding to a WA side-plate at Δt12/8 = 15 s.

(0.38MB).

At even lower temperatures (900–750 °C) and after the WA nucleation (as can be seen in Fig. 3), the IGA nucleates within the ferrite grain. These small IGA precipitates required a greater undercooling to nucleate as a consequence of a lower nitrogen content in ferrite and little availability of heterogeneous nucleation sites in the matrix. Then, the IGA precipitates growth is dictated by the phase transformation mechanism. Greater undercooling promotes a shear transformation, which ultimately leads to the directional growth of austenite nuclei along the (K–S)/N–W interfaces. Consequently, the overall fraction of specific ORs between the small IGA precipitates and ferrite becomes higher at lower transformation temperatures than in the HAZ where the phase transformation occurs at a higher temperature, either due to a lower cooling rate and /or higher saturation in nitrogen content. Therefore, interfaces with random ORs, between the IGA precipitates and ferrite will be more likely at higher transformation temperatures. The above can explain the predominance of interfaces with specific ORs especially in the simulated HAZ with the shorter cooling time in the present DSS. Fig. 5(a) shows the TEM bright field image of one IGA precipitate where the structural dislocations accommodate the misfit between the δ/γ lattices. According to the Frank and van der Merwe theory [29,30], these misfit dislocations arranged periodically reveal certain specific ORs between the austenite and ferrite lattices.

The increase of the austenite fraction with increasing the cooling time (Table 2a and Fig. 4(a)) in this work is in agreement with the literature [8,9,15–17,20,24]. However, the decrease of the increasing rate of the austenite fraction for longer cooling times can be due to the decrease of the nitrogen supersaturation in the present alloy (0.12–0.13 wt.%) and by reduction of the favorable nucleation sites. It should be observed that by considering the austenite morphologies, the volume fractions of different morphologies have shown significant deviations with respect to the fractions previously considered [8,15,16,19]. As the GBA is the first reformed austenite, the behavior showed in Fig. 4(a and b) seems to be logical because at shorter cooling times (programmed in this study), the ferrite grain size is smaller and a greater δ/δ boundary area and nitrogen content will be available to nucleate the GBA; reason for which the growth of the GBA predominates. Increasing the cooling time, increased the ferrite grain size and therefore the GBA nucleation sites diminish. Also, the supersaturation of nitrogen decreases given by the formation and growth of the other austenite morphological forms. All this contributes to a smaller GBA fraction and a stopping of the GBA average width at these longer cooling times (see Table 2b). Then, the ξGBA must diminish for higher austenite fractions as can be seen in Fig. 4(b).

As the WA needles or laths do not appear commonly in isolated form but in groups from the allotriomorphs GBA or inside the ferrite grain, a small ferrite grain size will contribute with a greater amount of favorable sites for the WA nucleation [9,24,31]. But, as a consequence of these smaller ferrite grain sizes; both, the large number of the WA nuclei and the predominance of GBA formed at higher temperatures, will consume the available nitrogen making difficult the growth of these small WA needles. Indeed, as can be appreciated in Fig. 3 and Table 2b, that greater ferritic grain sizes favor the growth of the WA in agreement with [32]. This is because at longer cooling times (lower undercooling), the driving force for nucleation of the IGA precipitates decreases so that the growth of the WA will be less interrupted. The relative growth of the IGA fraction for the longer cooling times is therefore associated with the decomposition of a certain WA fraction before formed. In addition, the monotonous increase in size of the IGA precipitates is a direct consequence of a better diffusion of the austenite-forming elements at longer cooling times.

4.2Mechanical properties analysis

The relative greater hardness recorded in the WA and the IGA precipitates at shorter cooling times (Fig. 8) cannot be explained simply by the hardening due to the substitutional elements in these precipitates. It is widely recognized that at normal welding cooling conditions (H˜0.5–2.5 kJ/mm) [33], the austenite formation is controlled by a paraequilibrium transformation, in which the diffusion of nitrogen is the controlling process rather than diffusion of the slower moving metallic elements [31,33–35]. The relative increases of hardness in these austenite precipitates at shorter cooling times must be, among another causes, due to the own transformation mechanism of the WA and IGA, which is different to the GBA formation mechanism [36]. For this reason, the hardness of the GBA showed a relative uniformity with a small increase at longer cooling times given by the partition of the alloying elements and enrichment in nitrogen. Also, it could be thought that such an increase of the hardness in the IGA and WA for the shorter cooling times was due to the interaction between the indenter and the boundaries of the small austenite precipitates. This is a typical indentation size effect that occurs whenever the indentation dimensions are comparable to the size of precipitates. In order to reduce as much as possible, the grain boundary effect and concentrate on the intragranular hardness which should have been more sensitive to the transformation mechanisms, it was therefore decided to use only larger precipitates for the microhardness measurements.

It is also noted that almost all nitrogen in the alloy partitions to austenite presented by the low density of Cr2N found in four large ferrite grains (greater than the average ferrite grain size) of the twenty carefully observed at TEM (Fig. 6). Furthermore, the ferrite hardness profile did not show an increase at the shorter cooling times, Fig. 8, as consequence of the small precipitated nitride fraction (corroborated by the TEM analysis). But, contrary, a small increment of the hardness was noted at the longer cooling times as consequence of the partition of the δ-forming elements as Cr and Mo to ferrite.

It is well accepted that when the ferrite grain size in the HAZ microstructure is increased, the yield strength of this thermal affected zone is deteriorated in a similar proportion [8,37–40]. However, taking into account the programmed cooling rates and the nitrogen content in the alloy, the strength did not undergo deterioration in correspondence with the ferrite grain size increase, as can be noted in Fig. 9(b). Also, as can be seen in Fig. 9(a), the areas below the strain-stress curves in tensile tests were practically the same. The area below the stress-strain curve is a qualitative indicator of the fracture toughness in the HAZ. So, small areas are related to low fracture toughness and large areas with more ductile fractures [41]. Therefore, it is also expected that no significant changes occur in the HAZ toughness, even when the austenite fraction increase for the tested longer cooling times. It is generally recognized that the WA affects the toughness in these HAZs [15,19]. However, the calculated WA fractions with practically stationary thicknesses at these cooling times (or heat inputs) (Table 2b) do not seem to have significantly influence on the HAZ toughness and strength.

The behavior of the yield strength (YS) and toughness in the HAZ of the DSS at these programmed (normal) cooling times is characterized by the interaction of several effects. At shorter cooling times, the relative increase in strength (in relation to the other cooling times and to the as-received material having 837 MPa as YS) must be associated with the different austenite volume fractions considering their morphological forms and with a greater ferrite fraction that has a smaller grain size. Increasing the cooling time, the ferrite grain size grew and its fraction decreased, but the yield strength and toughness did not decrease in the same proportion due to a higher fragmentation of the WA and growth of the IGA fraction. At longer cooling times, the influence of the relative hardening of ferrite, the increasing of the IGA fraction together with the increase of the entire austenite fraction and the contribution of several small ferritic grains with a dislocation structure in form of the cross-stitch [42] made possible the non-appreciable deterioration of both properties at the HAZs. It is reported without detailed calculations [19] that by increasing the heat input to above 2.5 kJ/mm, the toughness of the HAZ deteriorates as a result of an increase of the WA fraction (length and width).

4.3The pitting corrosion analysis

Nitride precipitation is considered to be responsible for preferential localized corrosion attack because of the chromium depletion adjacent to the nitrides in the ferrite matrix [43–47]. Many authors [43–47] have recommended high heat inputs to eliminate the chromium depletion by providing sufficient time for the chromium distribution to homogenize during continuous cooling. The survival or elimination of the chromium depleted zones around nitrides will depend on the interparticle spacing as well as on the diffusivity of chromium and the cooling time. To understand the influence of the recorded nitride densities on the pitting corrosion resistance at the studied HAZ microstructures, the diffusion distances (χ) of chromium and nitrogen atoms in the ferrite matrix are calculated during the continuous cooling between 1200 °C up to 500 °C for the programmed cooling rates. The used equations for these calculations, according to [14], are:

where tf − ti is the cooling time between 1200 °C and 500 °C. The D(T) is the volume diffusion coefficient and can be expressed by:

with D0 = 2.3 × 10−4 m2/s and Q = 239 kJ/mol for chromium, and D0 = 1.13 × 10-6 m2/s and Q = 83(1 −  14.03T ) kJ/mol for nitrogen, respectively [14,48]. The T temperature in Eq. 2 can be represented by: T = T0 − vt, where T0 (K) is the selected initial temperature during cooling, and v(K/s) is the cooling rate between tf − ti. Table 4 shows the diffusive distances of chromium and nitrogen atoms in the ferrite matrix during the continuous cooling between 1200 °C and 500 °C.

Table 4.

Diffusive distances of chromium and nitrogen atoms in the ferrite matrix during the continuous cooling between 1200 °C and 500 °C.

H (kJ/mm)  1.5  2.0  2.5  3.0 
Δt8/5 (s)  42.5  75.5  117.9  169.8 
Δt12/5 (s)  56.9  101.1  158  227.6 
v (K/s)  12.3  6.9  4.4  3.1 
χ (μm) for Cr  2.1  2.8  3.4  4.1 
χ (μm) for N  139  186  232  278 

As can be observed in Table 4, the diffusive distances for chromium in ferrite are approximately one order higher than the average nitride interparticle spacing (0.2 μm for Δt12/8 = 15 s) measured for the shorter cooling time. Thus, the survival of the chromium depleted zones around the nitrides is little probable at the programmed cooling times. Furthermore, the diffusive distances of nitrogen between 1200 °C and 500 °C are greater than the average ferrite grain sizes, Tables 2b and 4, for all cooling times. From this, an extensive nitride precipitation at the center of the ferrite grains is very difficult. This is in agreement with the TEM analysis where few coalesced nitrides were observed in isolated larger ferrite grains. The ferrite microhardness profiles also agree with the above analysis as it has been explained already.

The lower pitting corrosion resistance at the HAZ with the shorter cooling time, Δt12/8 = 15 s, is directly related to the ratio between the volume fractions of both phases and partition of the alloying elements. This difference in the localized corrosion resistance is caused mainly by contribution of the nitrogen to the pitting resistance equivalent number (PREN) of the austenitic phase. As the largest fraction of nitrogen partitioned to austenite (little precipitation of nitrides in ferrite) without significant changes in the partition of the substitutional elements, the PREN of this austenitic phase is raised at the lower cooling time. As the ferrite volume fraction at the HAZ, with smaller PREN, predominates at this shorter cooling times, its localized corrosion resistance decrease.

5Conclusions

Based on the detailed calculation of the austenite volume fractions taking into account their morphologies and the main microstructural parameters at the simulated heat-affected zone microstructures, the following conclusions are reached:

  • -

    The three austenite morphological forms were determined quantitatively in the heat-affected zone microstructures of a duplex stainless steel (DSS) UNS S31803 for different welding cooling times. The predominance of the grain boundary austenite (GBA) fraction was at the shorter cooling times, while the intragranular austenite (IGA) and the Widmanstätten austenite (WA) fractions increased moderately at the longer cooling times, decreasing the GBA fraction.

  • -

    The average length of the WA increased in correspondence with the ferrite grain growth, while its width did not suffer appreciable changes for longer cooling times. Also, the average width of the GBA stopped its growth at these longer cooling times.

  • -

    The ferrite hardness profile did not show an increase of the hardness values at the shorter cooling times which is basically given by the not significant fractions of the coalesced nitrides into the ferrite grains. Thus, the fractions of the different morphological forms of austenite and the corresponding microstructural parameters in the heat-affected zone (HAZ) are governed by the availability of nitrogen during the δ/γ transformation at the programmed cooling rates basically.

  • -

    The WA volume fractions were between 5% and 9% at the tested cooling times in the simulated heat-affected zones (HAZs). It was found that such WA fractions did not affect the HAZ mechanical properties as the strength and toughness. The mechanical behavior of these HAZs at the programmed cooling times is characterized by the interaction of multiple effects as: ferrite grain size, volumetric fractions of ferrite and austenite with different morphologies, fragmentation of the WA, partition of both ferrite and austenite forming-elements, dislocation structure in ferrite grains, and precipitation of nitrides.

  • -

    The small IGA precipitates in these HAZs showed predominance of the specific orientation relationships with ferrite for the programmed cooling times. The largest population of these specific interfaces was at the shorter cooling times as a consequence of a greater undercooling.

  • -

    Increasing electrolyte temperature from 25 °C to 60 °C reduced the pitting corrosion resistance of the DSS UNS S31803. Anodic polarization and critical pitting temperature (CPT) measurements showed that there is a decrease in the corrosion resistance of the DSS UNS S31803 for the simulated HAZs. Lower localized corrosion resistance at the HAZ was registered for shorter cooling times due to the nitrogen enrichment of the austenitic phase and a greater amount of ferrite with very little nitrides fraction at these cooling times.

Conflicts of interest

The authors declare no conflicts of interest.

Acknowledgments

The authors wish to thank Fundação de Amparo à Pesquisa do Estado do Rio de Janeiro (FAPERJ) and Coordenação de Aperfeiçoamento de Pessoal de Nivel Superior (CAPES) in Brazil for the financial support offered by the project 09/2014 — PVE — CAPES. One of the authors (EVM) also thanks FAPERJ for financial support (Processo No E-26/201-535/218) and CAPES for the project PVE 88881.064968/2014-01). The authors would also like to thank the anonymous reviewer for useful suggestions that improved the paper.

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