Effect of phosphorus content on the mechanical, microstructure and corrosion properties of supermartensitic stainless steel

Effect of phosphorus content on the mechanical, microstructure and corrosion properties of supermartensitic stainless steel

Author’s Accepted Manuscript Effect of phosphorus content on the mechanical, microstructure and corrosion properties of supermartensitic stainless ste...

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Author’s Accepted Manuscript Effect of phosphorus content on the mechanical, microstructure and corrosion properties of supermartensitic stainless steel C.A.D. Rodrigues, R.M. Bandeira, B.B. Duarte, G. Tremiliosi-Filho, A.M. Jorge www.elsevier.com/locate/msea

PII: DOI: Reference:

S0921-5093(15)30476-7 http://dx.doi.org/10.1016/j.msea.2015.10.013 MSA32864

To appear in: Materials Science & Engineering A Received date: 4 August 2015 Revised date: 3 October 2015 Accepted date: 5 October 2015 Cite this article as: C.A.D. Rodrigues, R.M. Bandeira, B.B. Duarte, G. Tremiliosi-Filho and A.M. Jorge, Effect of phosphorus content on the mechanical, microstructure and corrosion properties of supermartensitic stainless s t e e l , Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2015.10.013 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. 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.

Effect of Phosphorus Content on the Mechanical, Microstructure and Corrosion Properties of Supermartensitic Stainless Steel

C. A. D. Rodriguesa,*, R. M. Bandeiraa, B. B. Duartea, G. Tremiliosi-Filhoa, A. M. Jorge Jr.b

a

Instituto de Química de São Carlos - Universidade de São Paulo

Av. Trabalhador São-carlense, 400, CP: 780, 13560-970, São Carlos, SP, Brasil b

Universidade Federal de São Carlos, Departamento de Engenharia de Materiais, Rodovia

Washington Luiz, Km 235, CP: 676, 13560-270 São Carlos, SP, Brasil

* Corresponding author at Grupo de Eletroquímica, Instituto de Química de São Carlos Universidade de São Paulo, Av. Trabalhador São-carlense, 400, CP: 780, 13560-970, São Carlos, SP, Brasil. E-mail address: [email protected] (C. A. D. Rodrigues).

Abstract In this study, the effect of high and low phosphorus contents on the mechanical and corrosion properties of two supermartensitic stainless was evaluated as a function of tempering temperature (570, 620, 670 and 690 °C). The best heat treatment condition (tempered at 620 °C) was determined by conducting potentiodynamic polarization tests in natural seawater. At all temperatures, the two tempered steels presented nearly the same values of hardness, yield strength (0.2%), ultimate tensile strength and elongation. A high phosphorus content promoted the formation of nanoprecipitates of CrP4 between the retained austenite (3%) and martensitic grain boundaries. This structure significantly improved the impact toughness. The steels also presented similar pit potentials (0.290 V). Nevertheless, chromium phosphide (CrP4) shifted the anodic corrosion curve toward a nobler region and produced a passive region that was more stable against corrosion.

Keywords: Supermartensitic stainless steels, Phosphorus, Heat treatment, Mechanical properties, Microstructural characterization, Pitting corrosion.

1.

Introduction

Supermartensitic stainless steels (SMSSs) have been increasingly used in several industries, such as the chemical, pulp and paper, and oil and gas industries, for various purposes, such as in chemical tanks, tankers, pollution-control equipment, control equipment, and onshore and offshore tubing, due to lower production costs compared to those of duplex grades of stainless steel [1-3]. SMSSs combine high strength, elevated toughness, good weldability and corrosion resistance for applications in marine environments [4, 5]. The chemical composition of these steels is based on chromium (11 - 13 wt %), nickel (4 - 7 wt %) and molybdenum, (1 - 3 wt %) with trace amounts of nitrogen ( 0.002 wt %), carbon ( 0.02 wt %), phosphorus and sulfur (C, P and S  0.003 wt %) [6]. The mechanical property values of SMSSs typically fall within the following ranges: 25 – 32 HRC for hardness, 650 - 750 MPa for yield strength (0.2 %), 880 - 950 MPa for tensile strength, elongations at rupture of up to 20 % and impact energies of up to 150 J [3, 7]. The pitting corrosion potential of SMSSs ranges from 0.18 to 0.300 V, and when microalloyed with Al, Nb, Ti, V, and W, the microstructure of the steels will be improved in terms of grain refinement and precipitation; hence, the mechanical and corrosion properties will be improved [8, 9]. SMSSs always present a small amount of retained austenite in a martensitic matrix structure, which is generally obtained by applying an appropriate thermal treatment and achieving a suitable chemical composition balance. The amount of retained austenite can

typically be controlled by a single or double tempering treatment [10, 11]. Additionally, the formation of -ferrite phase within the martensite structure depends mostly on the cooling rates used in the heat treatment procedure and incorporating sufficient amounts of Mo and Ni [12]. Inappropriate heat treatments may significantly change the microstructure (at the nanoscale), reducing the weldability and mechanical and corrosion properties of these steels. Retained austenite, intermetallic compounds and nanosized δ-ferrite are often present, which cannot be observed by conventional, low-resolution microscopy. The mechanical and/or corrosion properties of SMSSs are strongly dependent on achieving a suitable balance of chemical composition. It is known that elements found in low concentrations modify the microstructure of the steels and consequently damage the aforementioned properties. For example, a high carbon content (C > 0.02 wt %) leads to the preferential formation of carbide precipitates of the type M7C3 or M23C6 (M = Cr, Fe, Ni and Mo) at existing austenite grain boundaries, a process also known as sensitization [13]. Sensitization generally causes Cr depletion in the vicinity of martensite lath grain boundaries and is associated with high strength or low toughness values; in addition, the process makes steels much more susceptible to intergranular corrosion in aggressive environments [14, 15]. A high nitrogen content (N > 0.002 wt %) induces the formation of carbon nitrides (CN), and when steels are microalloyed, nitrides of the type XN (X = Ti, Nb, V) are formed, which reduces corrosion [16-18]. A high chromium content (Cr > 12.5 %wt) leads to -ferrite phase formation within the martensitic structure, which is considered the most detrimental process to the mechanical and corrosion resistance of SMSSs [19-21]. A high nickel content (Ni > 5.5 wt %) promotes a higher retention of austenite during martensite formation (austenite  martensite), but an appropriate nickel

content improves mechanical and corrosion properties [22, 23]. A high molybdenum content (Mo > 2.0 wt %) leads to the formation of fine precipitates of Fe2Mo collectively known as Laves phase. These precipitates are considered detrimental because they reduce the mechanical and corrosion properties of SMSSs and are hardly following simple and double tempering [11]. Other factors can help improve the softness of SMSSs when alloyed with low carbon and nitrogen contents to help improve their weldability, and maintaining low sulfur and phosphorous contents helps improve toughness when steels are impact tested at low temperature. Temper brittleness can be detected by impact tests and is usually attributed to traces of impurities, such as sulfur and phosphorus. This type of brittleness occurs when steels are tempered at low temperature (350 – 550 °C) or during slow cooling through this temperature range [12, 24, 25]. High sulfur contents lead to the formation of manganese sulfide in stainless steels, which is undesirable due to the formation of cracks under straining conditions, thereby reducing corrosion resistance [26, 27]. A high phosphorus content has been widely recognized to be detrimental, mainly due to the segregation of phosphorus to grain boundaries after the element has already been fully rejected from the retained austenite; this segregation may lead to embrittlement and decreases corrosion resistance [28, 29]. Therefore, maintaining low sulfur and phosphorus contents has been an effective way to avoid other problems such as intergranular cracking and brittleness [30]. Because of the limitations inherent to the studies reported in the literature to date, in the current work, we evaluated the effect of high and low phosphorus contents on the microstructure, mechanical properties (at different tempering temperatures), and corrosion properties of SMSSs to determine the best heat treatment conditions.

Experimental procedure

2.1.

Materials and methods The steels were produced by Villares Metals Research and Development Center

(Villares Metals S.A., Brazil). To facilitate the distinction between samples containing low and high amounts of phosphorus, the samples will be hereafter called SMSS - P and SMSS + P, respectively. Square cast ingots measuring 140 mm2 in surface area and 300 mm in thickness and weighing 50 kg were hot-forged to intermediate 85 mm2 square billets and hot-rolled into round bars with a diameter of 29 mm. The chemical compositions of the two steels are listed in Table 1, which were determined by atomic absorption spectroscopy. The amounts of chemical elements are in accordance with the standards of the American Society for Testing and Materials (ASTM A 751), except for the high content of phosphorus. The experiments performed to study the solid-state transformation in the steels were carried out in an Adamel-Lhomargy-DT-1000 dilatometer. The thermal cycles were executed as follows: (i) On heating: a fixed rate of 0.33 °C/s was used to reach the austenitizing temperature of 900 °C, which was maintained for 5 min. (ii) On cooling: a rate of 0.2 °C/s was used up to room temperature. As a result, both steels exhibited similar dilatation heating curves and both showed a phase transformation from martensite to austenitic between 690 and 750 °C; furthermore, upon cooling, martensite was formed between 230 and 160 °C, as shown in Fig. 1. Finally, the measured AC1 temperature as 685 °C, which was automatically determined by the software tool of the dilatometer system.

The AC1 value served as the basis for selecting the following tempering conditions. The two steels were heated to 1000 °C, held at this temperature for 45 min, water-quenched, tempered at 570, 620, 670 and 690 °C, and then air-cooled (tempered condition). All samples were conventionally polished by progressive grinding with wet abrasive SiC sandpapers with grit numbers of 200, 400, 600, 800, 1200, 1500, and 2000 and then polished using 20 m chromium oxide powder. The SMSS structure was revealed by etching in Vilella’s reagent. Microstructural characterization was carried out using an optical microscope (OM) (Zeiss – Axiotech), a scanning electron microscope (SEM-LEO440 or HRSEM-FEIMagellan 400 L) coupled to an energy dispersive X-ray spectroscopy (EDS) system (OXFORD ISIS LINK 300), a transmission electron microscope (TEM), and scanning transmission electron microscopy (STEM) (FEI-TECNAI G2 F20) operated at 200 KV and coupled to an EDS system (EDAX). Selected area electron diffraction (SAED) patterns were interpreted the JEMS software program [31]. The crystal structures of the possible phases identified by the JEMS software were verified using the Inorganic Crystal Structure Database (ICSD) [32]. The retained austenite in the tempered samples was characterized by conventional X-ray diffraction (XRD-Rigaku Rotaflex - Japan) using Cu K radiation. The volume fraction (vol. %) of the austenite phase was quantified using the Rietveld refinement procedure based on the X-ray peak intensities (Maud software) [33]. The mechanical properties of the samples were evaluated by hardness measurements and tensile and Charpy impact tests conducted at -60 °C and room

temperature in accordance with the ASTM-E18, ASTM-A 370 and ASTM 23-04 standards, respectively. The corrosion resistances of the steels were determined electrochemically in natural seawater by obtaining anodic polarization curves using an Autolab potentiostat (model VGSTAT-302). Natural seawater was collected from the coastal area of Cabo Frio, Rio de Janeiro, Brazil. The chloride concentration, conductivity and pH of the seawater were 19.20 g L-1, 55.73 mS cm-1 and 8.0, respectively. Disc-shaped electrodes of the SMSS samples measuring 15 mm in diameter and 5 mm in thickness were assembled in a PTFE holder, with 0.5 cm2 of each specimen exposed to the test solution. The samples were then ultrasonically cleaned in acetone and ultrapure water for 10 min in each solvent. The auxiliary electrode was a 2 cm2 platinum sheet, and the reference electrode was a saturated silver-silver chloride electrode (Ag/AgCl/NaCl(sat)). The open-circuit potential (OCP) was measured for 60 min in the seawater. It was observed that, for all samples, the OCP reached a nearly n almost stable steady state value between -0.15 and -0.25 V vs. Ag/AgCl/NaCl(sat) after 13 min. Thus, a 15 min interval was adopted as the stabilization time to reach the OCP before starting the potentiodynamic polarization tests. The working electrode was polarized at -0.8 V (potential below the open-circuit potential), and anodic potentiodynamic polarization tests were begun at a scan rate of 1 mV s-1. All experiments were carried out in a climatized room at 26±1 °C.

3.

Results

3.1.

Characteristics of the utilized samples

Figs. 2 (a and b) show OM images of the SMSS + P and SMSS - P steels, respectively. Both steels presented very fine martensite after tempering at 620 °C. Measurements indicated grain sizes of ~13 µm for the SMSS + P steel and ~20 µm for the SMSS - P steel. In the remaining samples, after performing other tempering treatments, similar tempered martensite microstructures were observed, without significant changes in the grain size when observed by OM. Figs. 3(a and b) show the X-ray diffraction (XRD) patterns obtained for the SMSS + P and SMSS – P steels, respectively, tempered at 570, 620, 670 and 690 °C. The XRD patterns for the samples treated at 570 °C present characteristic martensite peaks. However, the retained austenite peaks appear only in the SMSS - P steel. However, both steels present retained lamellae austenite between thin laths of martensite, as shown in Figs. 4(a and b), indicating that the amount of retained austenite was less than 3 vol. % in the SMSS + P sample treated at 570 °C. In contrast, the XRD patterns of the samples treated at all other tempering temperatures (620, 670 and 690 °C) reveal that martensite and retained austenite were produced in both steels. The volume fraction of retained austenite was calculated using the Maud software program (Fig. 5(a)). As previously indicated, even though the SMSS + P steel tempered at 570 °C did not show characteristic peaks for this phase, the phase did appear in the micrographs shown in Fig. 4. Thus, we believe that, in under the tempering conditions, the amount of this phase was less than 3 vol. %, below the detection limit of the X-ray

technique [26]. Fig. 5(a) also that the SMSS - P steel possessed ~3.1 vol. % retained austenite at 570 °C, and when tempered at 620 °C, both steels showed similar volume fractions of the phase, ~3.3 and 3.1 vol. %, respectively, for the SMSS - P and SMSS + P samples. However, the amounts of retained austenite increased with temperature, and when tempered at 670 °C, the SMSS - P steel presented a volume fraction of ~ 5.5 vol. % and the SMSS + P steel presented a volume fraction 4.5 vol. %; moreover, when tempered at 690 °C, the SMSS - steel showed a volume fraction of ~12.0 vol. % and the SMSS + P steel a volume fraction of ~10.0 vol. %.

3.2.

Mechanical properties

Fig. 5(b) shows the variation in hardness as a function of the tempering temperature. The initial average hardness values (as tempered at 570 °C) were quite similar, with average values of ~27.40 and ~26.20 HRC, respectively, for the SMSS + P and SMSS – P steels. At higher tempering temperatures, the average hardness values decreased slightly, increased and decreased, for temperatures of 620, 670 and 690 °C, although the observed values were similar for both steels. Fig. 6 compares the variation and absolute values of the tensile properties and toughness for the SMSS + P and SMSS - P steels obtained at room temperature after quenching and tempering at 570, 620, 670 and 690 °C. Under these conditions, both steels presented similar variations in the 0.2 % yield strength (YS) and ultimate tensile strength (UTS) and even for the percent elongation (A-%). In terms of absolute values, the SMSS + P steels presented higher values for YS and UTS than did the SMSS - P steel and similar A-

% values when tempered at 570 and 620 °C. However, after tempering at 670 and 690 °C the situation was inverted. The values of YS and UTS became similar for both steels, and the A-% values increased slightly, with the highest value observed for the SMSS + P steel. Fig. 6(d) compared the impact energy (toughness) values obtained at room temperature and -60 °C for both steels in the as-quenched and tempered conditions (570, 620, 670 and 690 °C). As shown, when tested at room temperature or at -60 °C, both steels presented a greater increase in toughness at 620 °C, and the highest toughness values were observed for the SMSS + P steel under any testing condition at tempering temperatures below 670 °C, at which the SMSS - P steel began to acquire higher values. It is also worth noting that at -60 °C and when tempered below 620 °C, the SMSS + P steel acquired higher toughness values than did the SMSS - P steel tested at room temperature. 3.3.

Corrosion properties

The samples with more and less phosphorous were submitted to corrosion tests by potentiodynamic polarization. The samples selected for these tests were heat treated at 1000 °C/45 min/water and 620 °C/2 h/air because they presented a suitable combination of tensile properties and toughness (see Fig. 6). Fig. 7 shows polarization curves corresponding to the SMSS + P and SMSS - P steels. The polarization curves present a decrease in the corrosion current density for the SMSS + P steel relative to that of the SMSS – P steel. The corrosion current densities of the SMSS – P and SMSS + P steels were 1.32 x 10-7 and 6.82 x 10-8 A cm-2, respectively. The pitting potential of the SMSS + P steel was 0.290 V, lower than that of the SMSS – P steel, which presented pitting potential of 0.310 V. Thus, we can affirm that there was no significant difference between the pitting

potentials of the two steels. The corrosion potential was observed to be approximately 0.290 V for both steels. The passive region (-0.200 V to 0.300 V) for the SMSS + P sample (passivation current density of approximately 2 x 10-6 A cm-2) suggested higher stability than the same region for the SMSS – P sample (passivation current density of approximately 9 x 10-5 A cm-2). For both steels, which did not exhibit a repassivation process, when the potential was reversed (i.e., reverse scan), high anodic currents at potentials of up to approximately 0.08 V were observed. This significant current drop is associated with the shape of the very open curve, i.e., the curve return characteristics without passivation.

Discussion

After the tempering treatments, the microstructure was composed of very fine martensite, whose grain size remained nearly constant and ranged from ~13 µm for the SMSS + P steel to ~20 µm for the SMSS – P steel. On the other hand, the amount of retained austenite and the grain size thereof increased with the tempering temperature, even though the amounts of retained austenite were similar for both steels and slight more and larger grains were observed for the SMSS - P steel. However, a likely more important factor was the rate of transformation. In fact, although the volume fractions of austenite were similar, they increased with a rate of ~0.02 vol. %/oC until approximately 670 °C and thereafter with a rate of ~0.3 vol. %/oC, i.e., one order of magnitude. The tempering temperature of 690 °C is above the Ac1 temperature (685 °C). It is well known that at Ac1, retained austenite is formed and reaches large volume fractions [11, 19, 34]. However, the

concentrations of austenite-stabilizing elements (C, N and Ni) decrease and lead to an unstable austenite phase within a fresh martensite structure [34, 35]. The above mentioned features may explain the observed microhardness behavior, i.e., a nearly constant grain size, fewer and smaller retained austenite grains produced at nearly the same level of microhardness (~26 HRC) for the tempering temperature range of 570 to 670 °C. However, after an abrupt increase in the amount and grain size of retained austenite at 690 °C, the microhardness decreased to ~22 HRC, which was certainly due to the massive presence of retained austenite, a much softer phase than martensite. The observed variation in hardness as a function of tempering temperature is similar to (but less appreciable than) the variation in the YS and UTS values for both steels. Remarkable differences could be observed between the YS and UTS values when tempered at 570 and 620 °C, where the SMSS + P steel displayed higher YS and UTS values than the SMSS - P steel but similar A-% values. This behavior may be explained by the fact that, in both steels, the tempering temperature of 620 °C was not sufficient to eliminate the elevated deformed state, as can be readily observed in Fig. 8, which shows the presence of a high dislocation density in both steels. Moreover, as the small difference in grain size cannot explain by itself the behavior regarding YS and UTS, probably there is a synergetic contribution of grain size and substructure in the P-rich sample. In fact, the fine precipitation induced a more deformed state in the sample SMSS + P by pinning dislocations, leading to a fastest recrystallization in the SMSS+P sample and so reducing grain sizes. This last statement can be confirmed by Figure 9 where the bright-field STEM image shown in Fig. 9 (a) clearly demonstrates ultrafine martensite-lath morphologies with recrystallized grains in the interior and the presence of subgrains in the SMSS + P steel. The deformed state is confirmed by the

presence of characteristic streaks in the diffraction spots in the SAED pattern presented in Fig. 9 (b). The presence of subgrains in only the SMSS + P steel may have been due to the higher phosphorus content, which allowed producing fine precipitation and also affected the kinetics of formation of the retained austenite. During tensile tests, smaller grain sizes together with precipitation, which will pin dislocations, will increase the work-hardening and leading YS and UTS to higher levels than in the P-poor sample. The white circles in Fig. 9 (a) represent regions magnified in Figs. 9 (c and d). Furthermore, Fig. 9 (c) presents a bright-field STEM image showing phosphorus-rich nanoprecipitates (point 1), and Fig. 9 (d) presents a dark-field STEM image showing nanoprecipitates measuring approximately 8 nm, which occurred preferentially in dislocation lines. The SAED pattern in the inset of Fig. 9 (d) shows a diffraction of such precipitates together with matrix spots. Indexing of the SAED pattern revealed the presence of chromium phosphide (CrP4) along the [3, 1, 1] zone axis, which was confirmed by EDX analysis (point 1 as indicated, not shown here); this result indicates a higher content of phosphorous (0.07 wt. %) compared with that of the martensite matrix (point 2 as indicated, not shown here), where phosphorous was not detected. According to Rowcliffe and Nicholson [36], nanoprecipitates of chromium phosphide are nucleated only in grain boundaries or dislocations in austenitic stainless steels. The precipitates are present after solution treatment and clearly depend on the density and distribution of the dislocations. There are two reasons why dislocations are good sites for the nucleation of chromium phosphide precipitates. First, the dislocations are able to attract a collection of much smaller phosphorus atoms, and second, the dislocations are able to accommodate part of the misfit between the precipitates and the matrix. The foregoing discussion can also explain the behavior observed at higher

temperatures. After being tempered at 670 and 690 °C, both steels presented similar YS and UTS values, with a significant increase when tempered at 670 °C and a large decrease when tempered at 690 °C (see Fig. 6 (a and b)). Thus, the increase in the YS and UTS values at 670 °C may be associated with a lower dislocation density and an adequate amount (approximately 4.0 %) of smaller retained austenite grains within the tempered martensite matrix. The significantly sharp decrease in the YS and UTS values at 690 °C may be associated with dislocation elimination and inadequate amount (approximately 10.0 %) of larger, unstable retained austenite grains within the martensite matrix. Similar arguments have been made for SMSSs of different chemical compositions [37, 38], but in these cases, the structure was composed of greater amounts of retained austenite (10 %) with carbide nanoprecipitates within the martensite matrix. The elongation and impact toughness values showed similar variations as functions of tempering temperature for both steels. As previously indicated, this variation is also associated with different amounts of dislocations and retained austenite. The best impact toughness of the SMSM+P steel was obtained at -60 °C and room temperature (see Fig. 6(d)) when tempered at 620 °C. In this case, the presence of CrP4 nanoprecipitates with an optimum dislocation density and finely retained austenite (approximately 4.0 %) distributed along the martensite interlath boundaries justifies the best impact toughness obtained under these conditions. Similarly, some authors have [39-42] reported that the effect of phosphorus at adequate concentrations (0.03 wt %) is beneficial to increasing the impact toughness. This effect mainly occurs when phosphorus forms intermetallic nanoprecipitates of the type XP (X = Fe, Mo, Cr) in grain boundaries, which depends on the phosphorus

content of stainless steels. This behavior may be treated as a kind type of martensitic precipitation hardening process. In addition, Song et al. [43] demonstrated that retained austenite is enriched by Cr atoms whereas P atoms are fully rejected from the retained austenite in martensitic stainless steel (low carbon–13 % Cr – 4 % Ni – 1 % Mo – 0.017 % P) using three-dimensional atom probe (3DAP) tomography. Moreover, the authors also verified that the amount of retained austenite (10 %) obtained by two-stage tempering leads to a strong enhancement in the impact toughness of their steel compared with the amounts obtained by a single-stage treatment (no retained austenite detected by XRD). Ma et al. [44, 45] also reported on a SMSS (13 % Cr – 5 % Ni – 1 % Mo – 0.022 % P) whose impact toughness was effectively improved by increasing the retained austenite content when tempered at temperatures ranging from 550 °C to 625 °C. In this case, when tempered at 625 °C, the structure was composed of 9.3 % retained austenite and Cr-rich globular nanoprecipitates measuring 25 nm (likely of the type Cr23C6) within a martensite matrix. Under these condition, the impact energy was 230 J, very similar to the value obtained for the SMSS + P steel tempered at 620 °C, which was 246 J, but with a lower amount of retained austenite (3.4 %). This result indicates that precipitate strengthening can be effective even with smaller amounts of nanosized precipitates. Thus, the condition for best balancing the strength and toughness (-60 °C and room temperature) for both steels was tempering at 620 °C, which was selected for the corrosion tests, as shown in Fig. 7. In this case, again the effect of phosphorous was remarkable. Because both steels presented the same grain size and the same amount of retained austenite (~3 %), the only difference was the precipitation of CrP4 in the SMSS + P steel,

which may explain the position of the anodic polarization curves (see Fig. 7). In other words, although similar, the corrosion and passivation current densities were smaller and the corrosion potential was greater in the SMSS + P steel (~6.8 x 10-8 A cm-2, ~2.0 10-6 A cm-2, ~0.290 V) than in the SMSS - P steel (~1.3 x 10-7 A cm-2, ~9 x 10-5 A cm-2, ~0.310 V). According to Kimura et al. [46], the variation in the amount of retained austenite has no effect on the pitting potential. Nevertheless, retained austenite promotes the dissolution of carbon and nitrogen, preventing the precipitation of carbides or nitrides or even carbonitrides of Cr and/or Mo. This effect limits the number of sites for the nucleation of pitting corrosion. On the other hand, phosphorus is rejected by the retained austenite, hence facilitating the formation of CrP4 nanoprecipitates. However, although still contradictory in the literature, some authors [36, 47-50] have found that CrP4 is highly corrosion resistant compared with the intermetallic precipitates of Fe and Mo, thus preventing pitting initiation. However, other authors [44, 45] have reported a low pitting potential attributed to the presence of carbide nanoprecipitates in SMSS steels (13 % Cr – 5 % Ni – 1 % Mo – 0.022 % P).

4.

Conclusions

The following conclusions and comparisons were drawn in the present work: 1)

A high phosphorous content (0.017 % wt) proved effective in producing

nanoprecipitates of CrP4 between retained austenite and a finer martensitic structure. This

structure is beneficial to improving performance in terms of mechanical properties and pitting corrosion resistance after tempering at 620 °C. 2)

In particular, the phosphorous content of SMSSs can be increased by the synergism

between grain refinement, a lower amount retained austenite and the formation of CrP4 precipitates located in grain boundaries of the martensitic matrix. 3)

A low phosphorous content (0.005 % wt) produces a homogeneous structure with

slightly coarser grains and a greater amount of retained austenite with an increase in tempering temperature compared with those observed at a high phosphorus content. However, the measured values of mechanical properties fall within the expected range, and the final material exhibits good pitting corrosion resistance. 4)

Approximately 3 vol. % retained austenite did not affect the pitting potential;

however, it did produce higher corrosion resistance, which led to lower corrosion resistance between the retained austenite and lath martensite.

Acknowledgments The authors are grateful to CAPES, CNPq, and FAPESP (Brazil) for their financial support.

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Figure and Table captions:

Table 1: Chemical compositions of the SMSS samples in weight percent. Fig. 1. Dilation curves versus temperature for SMSS + P steel at heating-cooling rates of 0.33 and 0.2 °C/s, respectively. Fig. 2. OM images of (a) the SMSS + P steel tempered at 620 °C showing very fine martensite microstructure with grain sizes between 10 to 15 µm and (b) the SMSS - P steel tempered at 620 °C showing a martensite microstructure with grain sizes between 15 to 25 µm. Fig. 3. X-ray diffraction patterns of the (a) SMSS + P and (b) SMSS - P steels tempered at 570, 620, 670 and 690 °C. Fig. 4. SEM images showing the presence of retained austenite lamellae in grain boundaries (lighter-colored phase) for the (a) SMSS + P and (b) SMSS - P steels tempered at 620 °C. Fig. 5. (a) Comparison of vol. % of retained austenite in SMSS + P and SMSS - P steels tempered at 570, 620, 670 and 690 °C. (b) Variation in hardness as a function of tempering temperature. Fig. 6. Mechanical properties as a function of tempering temperature of the two steels studied: (a) yield strength (YS-0.2%), (b) ultimate tensile strength (UTS), (c) elongation (A%) and (d) Charpy toughness (J).

Fig. 7. Polarization curves for the SMSS + P and SMSS - P tempered at 620 °C in natural seawater at pH 8 and a scan rate of 1 mV s-1. Fig. 8. Bright-field STEM images showing the dislocations between the thin martensite laths of the (a) SMSS + P and (b) SMSS - P steels. Fig. 9. (a) Bright-field STEM image showing the ultrafine martensitic-lath morphologies with recrystallized grains inside and the presence of subgrains. The white circles represent regions magnified in Figs. 9(c and d). (b) SAED pattern of the martensitic matrix. (c) Bright-field STEM image showing phosphorus-rich nanoprecipitates (point 1) located in subgrains of the SMSS + P steel. (d) Dark-field STEM image showing nanoprecipitates measuring approximately 8 nm; the corresponding SAED pattern of the nanoprecipitates is shown in the upper right.

Table 1: Chemical compositions of the SMSS samples in weight percent. Steel

C

Cr

Ni

Mo

Mn

Si

S

P

N

SMSS - P

0.013

12.50

5.05

2.12

0.30

0.18

0.0014

0.005

0.001

SMSS + P

0.012

12.50

5.03

2.11

0.29

0.19

0.0013

0.017

0.001