materials

Article Relationship between Microstructure and Corrosion Behavior of Martensitic High Nitrogen Stainless Steel 30Cr15Mo1N at Different Austenitizing Temperatures

Zhouhua Jiang 1, Hao Feng 1, Huabing Li 1,* ID , Hongchun Zhu 1, Shucai Zhang 1, Binbin Zhang 1, Yu Han 1, Tao Zhang 2 and Dake Xu 2

1 School of Metallurgy, Northeastern University, Shenyang 110819, China; [email protected] (Z.J.); [email protected] (H.F.); [email protected] (H.Z.); [email protected] (S.Z.); [email protected] (B.Z.); [email protected] (Y.H.) 2 School of Materials Science and Engineering, Northeastern University, Shenyang 110819, China; [email protected] (T.Z.); [email protected] (D.X.) * Correspondence: [email protected]; Tel.: +86-24-8368-9580; Fax: +86-24-2389-0559

Received: 29 June 2017; Accepted: 25 July 2017; Published: 27 July 2017

Abstract: The relationship between microstructure and corrosion behavior of martensitic high nitrogen stainless steel 30Cr15Mo1N at different austenitizing temperatures was investigated by microscopy observation, electrochemical measurement, X-ray photoelectron spectroscopy analysis and immersion testing. The results indicated that finer Cr-rich M2N dispersed more homogeneously than coarse M23C6, and the fractions of M23C6 and M2N both decreased with increasing austenitizing temperature. The Cr-depleted zone around M23C6 was wider and its minimum Cr concentration was lower than M2N. The metastable pits initiated preferentially around coarse M23C6 which induced severer Cr-depletion, and the pit growth followed the power law. The increasing of austenitizing temperature induced fewer metastable pit initiation sites, more uniform element distribution and higher contents of Cr, Mo and N in the matrix. In addition, the passive film thickened and Cr2O3, Cr3+ and CrN enriched with increasing austenitizing temperature, which enhanced the stability of the passive film and repassivation ability of pits. Therefore, as austenitizing temperature increased, the metastable and stable pitting potentials increased and pit growth rate decreased, revealing less susceptible metastable pit initiation, larger repassivation tendency and higher corrosion resistance. The determining factor of pitting potentials could be divided into three stages: dissolution of M23C6 ◦ ◦ (below 1000 C), dissolution of M2N (from 1000 to 1050 C) and existence of a few undissolved precipitates and non-metallic inclusions (above 1050 ◦C).

Keywords: martensitic high nitrogen stainless steel; austenitizing temperature; microstructure; pit initiation; passive film; pit growth

1. Introduction Corrosion is diagnosed as one of the main factors for premature bearing failures in many aviation applications, particularly in aircraft engines and helicopters [1,2]. The conventional high carbon martensitic stainless steels, such as 440C and BG42, possess some corrosion resistance. However, the existence of coarse Cr-rich eutectic carbides (M7C3) deteriorates their fatigue capabilities and corrosion resistance [2]. As an important alloying element, nitrogen can significantly improve the corrosion resistance and mechanical properties of stainless steels, and has been widely used in austenitic and duplex stainless steels [3–6]. Various mechanisms, including theories of ammonia production [7], surface enrichment [8], anodic segregation [9] and salt film formation [10], have been proposed to explain the effect of nitrogen on corrosion resistance of nitrogen-containing steels.

Materials 2017, 10, 861; doi:10.3390/ma10080861 www.mdpi.com/journal/materials Materials 2017, 10, 861 2 of 19

For martensitic stainless steels (MSSs), nitrogen in solid solution could also enhance their corrosion resistance [11,12]. Besides, nitrogen is beneficial to improve the hardenability and avoid the segregation of eutectic carbides in MSSs [2]. The partial substitution of nitrogen for carbon could also increase the thermodynamic stability of solid solution, toughness and plasticity of MSSs [13]. However, due to the low nitrogen solubility (normally less than 0.08 wt %) in martensitic steels at atmospheric pressure [11,14], it is difficult to obtain MSSs with high nitrogen content by traditional methods, such as electric arc furnace (EAF) melting, vacuum induction melting (VIM) and argon oxygen decarburization (AOD) refining. Therefore, few studies about martensitic high nitrogen stainless steels were reported [12,15–17]. In recent years, with the development of pressure metallurgy, series of MSSs with high nitrogen content (0.3–0.5 wt %), such as CRONIDUR steels, have been invented using pressurized electroslag remelting (PESR) [11]. For example, CRONIDUR 30 has been used successfully in demanding applications, such as bearings for aviation turbine and cryogenic rocket turbopumps [18]. It is well known that austenitizing and tempering strongly influence the microstructure and properties of MSSs [19]. In general, the austenitizing temperature determines the amount and distribution of undissolved carbides and retained austenite [20], and has an important influence on the mechanical properties and corrosion resistance [21]. Meanwhile, decarburization, grain coarsening and the formation of δ-ferrite should be avoided in selecting a suitable austenitizing temperature [19,22]. Several studies have been carried out about the effect of austenitizing treatment on the microstructure and corrosion resistance of nitrogen-free or low nitrogen martensitic stainless steels [20,21,23]. The results indicated that the reduction of undissolved Cr-rich carbides and the improvement in homogeneity of Cr distribution with increasing austenitizing temperature improved the corrosion resistance of steels. The existence of Cr-depleted zones in matrix adjacent to Cr-rich precipitates acted as preferential sites for metastable pits. In addition, Lu et al. [20,23] pointed out that the increased austenitizing temperature contributed to forming more protective and well-structural passive films on 13 wt % Cr-type MSSs. However, few studies can be found about the influence of austenitizing temperature on microstructure, composition of passive film, pit initiation and propagation of martensitic high nitrogen stainless steels. Martensitic high nitrogen stainless steel 30Cr15Mo1N with 0.44 wt % nitrogen has been developed successfully in Northeastern University using pressure metallurgy method. The microstructure, mechanical and corrosion properties of friction stir welding 30Cr15Mo1N were investigated in our previous study, and the results revealed that the weldment exhibited lower hardness and better corrosion resistance [24]. The present work aims to reveal the relationship between microstructure and corrosion behavior of 30Cr15Mo1N at different austenitizing temperatures, and promote the development and application of martensitic high nitrogen stainless steels. The investigation of characterization of passive film, pitting initiation and propagation, etc. of such steel could afford a broad overall view on pitting corrosion, and is beneficial for understanding the corrosion mechanism of martensitic high nitrogen stainless steels.

2. Materials and Methods

2.1. Material and Heat Treatment The experimental martensitic high nitrogen stainless steel 30Cr15Mo1N was melted in a 25 kg pressurized induction furnace (Jinzhou Yuanteng Electric Furnace Technology Co., Ltd., Jinzhou, China). During the smelting process, the vacuum carbon-deoxidization together with adding nickel-magnesium (Ni-Mg) master alloy and metallic cerium was used for deoxidization and desulfurization. Then, nitrogen was introduced to the steel via the reaction, N2 (g) → 2[N], at the gas-melt interface under nitrogen pressure of 0.4 MPa. Finally, the molten steel was poured into the ingot mold and solidified under nitrogen pressure of 1.5 MPa to avoid the formation of nitrogen bubbles and porosity. The chemical composition of the present material is shown in Table1. Firstly, the ingot was Materials 2017, 10, 861 3 of 19 diffusionMaterials 2017 annealed, 10, 861 at 1270 ◦C for 10 h, then hot forged into 110 mm × 30 mm plate in the temperature3 of 19 range of 1050–1150 ◦C. Afterwards, the plate was annealed at 875 ◦C for 5 h, then furnace cooled to temperature range of 1050–1150 °C. Afterwards, the plate was annealed at 875 °C for 5 h, then furnace 700 ◦C and kept for another 3 h, followed by cooling to 600 ◦C in furnace at the speed of 1 ◦C/min, cooled to 700 °C and kept for another 3 h, followed by cooling to 600 °C in furnace at the speed of finally air-cooled to room temperature (Figure1). The plate was cut into several parts, then austenitized 1 °C/min, finally air-cooled to room temperature (Figure 1). The plate was cut into several parts, then at 900 ◦C, 950 ◦C, 1000 ◦C, 1020 ◦C, 1050 ◦C, 1100 ◦C and 1150 ◦C, respectively, for 1 h followed by austenitized at 900 °C, 950 °C, 1000 °C, 1020 °C, 1050 °C, 1100 °C and 1150 °C, respectively, for 1 h oil quenching. followed by oil quenching. Table 1. Chemical composition of martensitic high nitrogen stainless steel 30Cr15Mo1N (wt %). Table 1. Chemical composition of martensitic high nitrogen stainless steel 30Cr15Mo1N (wt %).

CC CrCr MoMo NN Mn Mn Si Si Ni Ni S P P O O 0.310.31 15.1715.17 1.031.03 0.440.44 0.440.44 0.520.52 0.070.07 0.002 0.002 0.011 0.011 0.0015 0.0015

FigureFigure 1.1. Schematic ofof heatheat treatmenttreatment procedure.procedure.

2.2.2.2. ThermodynamicThermodynamic CalculationsCalculations TheThe variationvariation ofof phasephase fractionsfractions andand compositioncomposition withwith temperaturetemperature atat equilibriumequilibrium werewere generatedgenerated usingusing Thermo-CalcThermo-Calc withwith TCFE7TCFE7 databasedatabase (Thermo-Calc(Thermo-Calc SoftwareSoftware AB,AB, Solna,Solna, Sweden).Sweden).

2.3.2.3. MicrostructureMicrostructure CharacterizationsCharacterizations ToTo revealreveal thethe phasephase constitution,constitution, thethe specimensspecimens forfor X-rayX-ray diffractiondiffraction (XRD)(XRD) analysesanalyses withwith thethe dimensionsdimensions ofof 1212 mmmm ×× 10 mm × 66 mm mm were ground with SiC paper to 2000-grit2000-grit andand mechanicallymechanically polishedpolished with 2.5 μµm diamond paste, paste, then then electropolished electropolished at at 20 20 V Vin in the the electrolyte electrolyte composed composed of of25 25vol vol % %perchloric perchloric acid acid and and 75 75 vol vol % % ethanol. ethanol. The The XRD XRD tests tests were were carried carried out out on an X-rayX-ray diffractometerdiffractometer (D/Max-2400,(D/Max-2400, Rigaku,Rigaku, Tokyo,Tokyo, Japan) Japan) with with Cu Cu K αKαradiation radiation under under 50 50 kV, kV, 180 180 mA mA and and 2θ ranging2θ ranging from from 40◦ to40° 101 to ◦101°at a at scanning a scanning speed speed of 2◦ of/min. 2°/min. The The volume volume fraction fraction of retained of retained austenite austenite was determinedwas determined by Equation by Equation (1) [25 (1)]: [25]: 1.4Iγ VA = 1.4Iγ (1) VA =Iα + 1.4Iγ (1) Iα + 1.4Iγ where Iγ and Iα are the integrated intensities of (111)γ and (110)α, respectively. whereTo I observeγ and Iα are the the influence integrated of austenitizing intensities of temperature (111)γ and (110) on precipitationα, respectively. distribution, specimens with dimensionsTo observe ofthe 12 influence mm × 10 of mm austenitizing× 6 mm were temper groundature with on SiC precipitation paper to 2000-grit distribution, and mechanically specimens polishedwith dimensions with 2.5 µ mof diamond12 mm paste,× 10 thenmm etched× 6 mm by Vilella’swere ground reagent with (consisting SiC pa ofper 1 g picricto 2000-grit acid, 10 and mL hydrochloricmechanically acid polished and 100 with mL 2.5 ethanol). μm diamond Microstructure paste, then and etched element by distributionVilella’s reagent were (consisting performed usingof 1 g fieldpicric emission acid, 10 scanningmL hydrochloric electron microscopeacid and 100 (FE-SEM, mL ethanol). Ultra Microstructure Plus, Carl Zeiss, and Oberkochen, element distribution Germany) andwere electron performed probe using microanalyzer field emission (EPMA, scanning JXA-8530F, electron JEOL, microscope Tokyo, Japan) (FE-SEM, equipped Ultra withPlus, wavelength Carl Zeiss, dispersiveOberkochen, spectrometer Germany) (WDS).and electron The FE-SEM probe microanalyzer and EPMA were (EPMA, operated JXA-8530F, with an JEOL, accelerating Tokyo, voltage Japan) ofequipped 15.0 kV, with and wavelength the probe current dispersi ofve EPMA spectrometer was 5.0 (WDS).× 10−8 TheA. The FE-SEM area percentageand EPMA ofwere precipitates operated −8 waswith quantified an accelerating using voltage Image-Pro of 15.0 Plus kV, software and the Version probe 6.0current (Media of EPMA Cybernetics was 5.0 Inc., × 10 Rockville, A. The MD,area USA).percentage The microstructure of precipitates and was lattice quantified information using of the Image-Pro precipitates Plus were software investigated Version by transmission 6.0 (Media electronCybernetics microscopy Inc., Rockville, (TEM, MD, JEM-2100F, USA). The JEOL, microstructure Tokyo, Japan) and at lattice 200 kV, information and the Cr of concentration the precipitates at were investigated by transmission electron microscopy (TEM, JEM-2100F, JEOL, Tokyo, Japan) at 200 kV, and the Cr concentration at the precipitate–matrix interface was analyzed using energy dispersive spectrometer (EDS) attached to TEM. For TEM foil preparation, thin disks with diameter

Materials 2017, 10, 861 4 of 19 the precipitate–matrix interface was analyzed using energy dispersive spectrometer (EDS) attached to TEM. For TEM foil preparation, thin disks with diameter of 3 mm were cut from the specimen austenitized at 900 ◦C and twin-jet polished using the electrolyte composed of 8 vol % perchloric acid and 92 vol % ethanol at −20 ◦C.

2.4. Electrochemical Measurement The specimens for electrochemical measurements were mounted in epoxy resin with an exposed area of 10 mm × 10 mm and abraded with SiC paper to 2000-grit. The test medium was 3.5 wt % NaCl solution at 30 ± 0.5 ◦C prepared with analytically pure NaCl and deionized water. The potentiodynamic polarization measurement was performed using a Gamry Reference 600 potentiostat (Gamry, Warminster, PA, USA) with a three-electrode system, consisting of a platinum counter electrode, a saturated calomel electrode (SCE) reference electrode and the specimen as the working electrode. Prior to the measurement, the working electrode was polarized at −1.0 VSCE for 5 min to eliminate the surface oxide layer formed in the air, then stabilized at open circuit potential (OCP) for 30 min. After that, potentiodynamic polarization was performed from −0.3 V below OCP toward the positive direction at a scan rate of 0.333 mV/s, and test was terminated when current density reached 1 mA/cm2.

2.5. Passive Film Analysis The specimens for passive film growth were mechanically abraded and polished. After removing the surface oxide layer by cathodic polarization, the passivation potential, which was chosen based on the potentiodynamic polarization curves, was applied for 1 h to promote the growth of passive film. The passive film analysis was then performed using X-ray photoelectron spectrometer (XPS, ESCALAB 250, Thermo Scientific, Waltham, MA, USA) with Al Kα (1486.6 eV) X-ray source. The binding energies were calibrated relative to C 1s peak at 284.6 eV. The argon ions sputtering (base pressure: 1.33 × 10−5 Pa, energy: 2 kV, current: 2.0 µA/cm2) over an area of 2 mm × 2 mm were performed for 10 s to clear the contamination layer. The curve fitting was performed via XPSPEAK 4.1 software by referencing to a database [26].

2.6. Immersion Tests The pit growth kinetics was evaluated using immersion testing and metallographic examination. After being abraded with 2000-grit SiC paper, the specimens austenitized at 900 ◦C, 1000 ◦C and 1100 ◦C ◦ were immersed in 6 wt % FeCl3 solution at 50 ± 0.5 C according to ASTM G48-11 (Method A) [27]. After being immersed for 3 h, 6 h, 12 h and 24 h, the specimens were taken out of the solution, cleaned using deionized water and dried using cold air. Then the specimens were cleaned in accordance with ISO 8407-2009 [28] to remove the corrosion products. Finally, the variation of pit depth was estimated using laser scanning confocal microscope (LSCM, LEXT OLS4100, Olympus, Tokyo, Japan).

3. Results

3.1. Thermodynamic Calculations The variation of phase fractions with temperature in martensitic high nitrogen stainless steel 30Cr15Mo1N using Thermo-Calc software is presented in Figure2. There existed two kinds of precipitates, i.e., M23C6 and hexagonal close-packed (hcp) phase. According to previous work by Mehtedi et al. [13] and Kaluba et al. [15], the hcp phase was preliminarily inferred to be Cr-rich M2N. The contents of M23C6 and M2N both decreased with the increasing of temperature, whereas the dissolution of M23C6 occurred ahead of M2N due to the higher diffusivity of carbon respect to nitrogen [13,15]. In contrast with traditional carbon alloyed MSSs, the existence of M2N with higher thermal stability in 30Cr15Mo1N enhanced the required austenitizing temperature in order to obtain uniform distribution of elements. In addition, Figure3 shows the variation of phase composition with Materials 2017, 10, 861 5 of 19

temperature. The Cr and Mo contents in M23C6 and M2N decreased, and the Cr, Mo, C and N contents inMaterials matrix 2017 increased, 10, 861 with increasing austenitizing temperature. 5 of 19 Materials 2017, 10, 861 5 of 19

Figure 2. Variation ofof phasephase fractionsfractions withwith temperaturetemperature inin 30Cr15Mo1N.30Cr15Mo1N. Figure 2. Variation of phase fractions with temperature in 30Cr15Mo1N.

(a) (b) (a) (b)

(c) (c)

Figure 3. Variation in phase composition of: (a) M23C6; (b) M2N; and (c) matrix with temperature Figure 3. Variation in phase composition of: (a) M23C6; (b) M2N; and (c) matrix with temperature Figurecalculated 3. Variation by Thermo-Calc. in phase composition of: (a)M23C6;(b)M2N; and (c) matrix with temperature calculated by Thermo-Calc. 3.2. Microstructure Characterizations 3.2. Microstructure Characterizations 3.2.1. X-ray Diffraction and Hardness Tests 3.2.1. X-ray Diffraction and Hardness Tests Figure 4 plots the XRD patterns of 30Cr15Mo1N in annealing and austenitizing conditions. The Figurediffraction 44 plotsplots peaks thethe of XRDXRD body-centered patternspatterns ofof cubic 30Cr15Mo1N30Cr15Mo1N (bcc) phase inin were annealingannealing observed andand in austenitizingaustenitizing all specimens. conditions.conditions. However, Theit is diffractionimpossible peakspeaks to separate ofof body-centered body-centered the ferrite cubic (cubicα) and (bcc) (bcc) martensite phase phase were were (α observed’) observedphases in owing allin specimens.all tospecimens. the identical However, However, lattice it is itconstants is impossible and structureto separate [20,22]. the ferrite Thus, ( αRockwell) and martensite hardness (α was’) phases measured owing using to the HRS-150D identical latticetester constants(Jvjing Precision and structure Instrument [20,22]. Manufacturing Thus, Rockwell Co., hardnessLtd., Shanghai, was measured China) with using Rockwell HRS-150D C scale tester to (Jvjingidentify Precision these two Instrument phases, as Manufacturing shown in Figure Co., 5. Ltd., It is Shanghai,noted that China) the hardness with Rockwell value of C annealed scale to identifyspecimen these was two significantly phases, as lower show thann in Figurethose of 5. theIt is austenitized noted that theones. hardness Therefore, value the of bcc annealed phases specimendetected bywas XRD significantly spectra were lower identified than those to beof theferrite austenitized in annealing ones. condition Therefore, and the martensite bcc phases in detected by XRD spectra were identified to be ferrite in annealing condition and martensite in

Materials 2017, 10, 861 6 of 19

impossible to separate the ferrite (α) and martensite (α’) phases owing to the identical lattice constants and structure [20,22]. Thus, Rockwell hardness was measured using HRS-150D tester (Jvjing Precision Instrument Manufacturing Co., Ltd., Shanghai, China) with Rockwell C scale to identify these two phases, as shown in Figure5. It is noted that the hardness value of annealed specimen was significantly lower Materials 2017, 10, 861 6 of 19 Materialsthan those 2017 of, 10 the, 861 austenitized ones. Therefore, the bcc phases detected by XRD spectra were identified6 of 19 to beaustenitizing ferrite in annealing condition condition [20]. With and the martensite increasing in of austenitizing austenitizing condition temperature, [20]. Withthe peaks the increasing for γ phase of austenitizing temperature, the peaks for γ phase were enhanced, indicating increased content of retained austenitizingwere enhanced, condition indicating [20]. increased With the contentincreasing of ofretained austenitizing austenite. temperature, Due to the the dissolution peaks for ofγ phaseM23C6 austenite. Due to the dissolution of M C and M N, the contents of carbon and nitrogen in matrix were wereand M enhanced,2N, the contents indicating of carbon increased and23 content nitrogen6 of 2 inre tainedmatrix austenite.were raised Due with to theincreasing dissolution austenitizing of M23C6 andraisedtemperature, M with2N, the increasing thereby contents depressing austenitizing of carbon theand temperature, martensitic nitrogen in therebytransformation. matrix depressing were raised the with martensitic increasing transformation. austenitizing temperature, thereby depressing the martensitic transformation.

Figure 4. XRD patterns of 30Cr15Mo1N in anne annealingaling and austenitizing conditions. Figure 4. XRD patterns of 30Cr15Mo1N in annealing and austenitizing conditions.

In addition, from the local amplification of XRD spectra from 47°◦ to 58°◦ in Figure 4, weak peaks InIn addition, addition, from from the local amplification amplification of XRD spectra from 47° 47 to 58° 58 inin Figure Figure 44,, weakweak peakspeaks for M23C6 and M2N were found in annealed and austenitized at 900 ◦°C specimens. With increasing for M23C6 and M2N were found in annealed and austenitized at 900 C specimens. With increasing foraustenitizing M23C6 and temperature M2N were found to 950 in °C, annealed the peaks and for austenitized M23C6 and M at2 N900 weakened, °C specimens. and then With disappeared increasing austenitizing temperature to 950 ◦C, the peaks for M C and M N weakened, and then disappeared austenitizingabove 1000 °C. temperature However, theto 950 non-appearance °C, the peaks forof peak M2323Cs 6for and austenite M22N weakened, or precip anditates then in XRDdisappeared profiles above 1000 ◦C. However, the non-appearance of peaks for austenite or precipitates in XRD profiles abovemeant 1000the absence °C. However, of these the phases non-appearance or the conten ofts peak weres lowerfor austenite than the or detection precipitates limit in [29]. XRD profiles meantmeant the absence of these phases or the conten contentsts were lower than the detection limit [[29].29].

Figure 5. Variation of retained austenite content and hardness with austenitizing temperature. Figure 5. Variation of retained austenite content and hardness with austenitizing temperature. Figure 5. Variation of retained austenite content and hardness with austenitizing temperature. 3.2.2. Microscopy Observation 3.2.2. Microscopy Observation The morphologies of 30Cr15Mo1N austenitized at 900 °C, 950 °C, 1000 °C, 1050 °C, 1100 °C and 1150 The°C are morphologies illustrated in of Figure 30Cr15Mo1N 6. The microstructure austenitized at consisted 900 °C, 950 of scattered°C, 1000 °C,precipitate 1050 °C, particles, 1100 °C and i.e., 1150Cr-rich °C Mare23 Cillustrated6 and M2N, in distributingFigure 6. The in microstructure lath martensite consisted and retained of scattered austenite. precipitate As shown particles, in Figure i.e., 6 Cr-richand Table M23 C2,6 andthe fractionM2N, distributing of precipitates in lath decreased martensite significantly and retained as austenite. the austenitizing As shown temperature in Figure 6 andincreased Table to2, 1000 the °C.fraction When of the precipitates austenitizing decreased temperature significantly was above as 1050the austenitizing °C, just a few temperatureundissolved increasedprecipitates to 1000and non-metallic°C. When the inclusionsaustenitizing existed temp eraturein the matrix. was above In addition, 1050 °C, justno aδ -ferritefew undissolved or coarse precipitateseutectic carbides and non-metallicwere observed inclusions among all existed the sp inecimens the matrix. austenitized In addition, in the norange δ-ferrite of 900–1150 or coarse °C, eutecticwhich is carbides consistent were with observed the Thermo-Calc among all result the sp (Figureecimens 2). austenitized in the range of 900–1150 °C, which is consistent with the Thermo-Calc result (Figure 2).

Materials 2017, 10, 861 7 of 19

3.2.2. Microscopy Observation The morphologies of 30Cr15Mo1N austenitized at 900 ◦C, 950 ◦C, 1000 ◦C, 1050 ◦C, 1100 ◦C and 1150 ◦C are illustrated in Figure6. The microstructure consisted of scattered precipitate particles, i.e., Cr-rich M23C6 and M2N, distributing in lath martensite and retained austenite. As shown in Figure6 and Table2, the fraction of precipitates decreased significantly as the austenitizing temperature increased to 1000 ◦C. When the austenitizing temperature was above 1050 ◦C, just a few undissolved precipitates and non-metallic inclusions existed in the matrix. In addition, no δ-ferrite or coarse eutectic carbides were observed among all the specimens austenitized in the range of 900–1150 ◦C, which is consistentMaterials 2017 with, 10, 861 the Thermo-Calc result (Figure2). 7 of 19

Figure 6. Morphologies of 30Cr15Mo1N austenitized at: (a) 900 °C; (b) 950 °C; (c) 1000 °C; (d) 1050 °C; Figure 6. Morphologies of 30Cr15Mo1N austenitized at: (a) 900 ◦C; (b) 950 ◦C; (c) 1000 ◦C; (d) 1050 ◦C; (e) 1100 °C; and (f) 1150 °C. (e) 1100 ◦C; and (f) 1150 ◦C.

Table 2. Area percentage of precipitates in 30Cr15Mo1N austenitized at different temperatures. Table 2. Area percentage of precipitates in 30Cr15Mo1N austenitized at different temperatures. Temperature (°C) 900 950 1000 1050 1100 1150 Area percentageTemperature of precipitates (◦C) (%) 7.75 900 ± 0.71 6.07 950 ± 0.96 3.94 1000 ± 0.99 1.59 1050 ± 0.63 1.23 1100± 0.48 0.63 1150± 0.40 Area percentage of precipitates (%) 7.75 ± 0.71 6.07 ± 0.96 3.94 ± 0.99 1.59 ± 0.63 1.23 ± 0.48 0.63 ± 0.40 The distribution of Cr, C and N via EPMA in Figure 7 shows that the precipitates were rich in Cr, C and N, and the quantity of N-rich particles was much lower than that of C-rich ones. The distribution of Cr, C and N via EPMA in Figure7 shows that the precipitates were rich in Cr, The N-rich particles were finer and distributed more homogeneously than the C-rich ones. With the C and N, and the quantity of N-rich particles was much lower than that of C-rich ones. The N-rich increasing of austenitizing temperature, the precipitates dissolved and the distribution of Cr, Mo, C particles were finer and distributed more homogeneously than the C-rich ones. With the increasing of and N became more homogeneous. Besides, the C content in lath martensite at 1000 °C and 1100 °C austenitizing temperature, the precipitates dissolved and the distribution of Cr, Mo, C and N became (indicated by arrows in Figure 7) was remarkably elevated. The extra dissolved C in the martensite more homogeneous. Besides, the C content in lath martensite at 1000 ◦C and 1100 ◦C (indicated by would result in the increased internal martensite lattice stress, furthermore increasing the hardness, arrows in Figure7) was remarkably elevated. The extra dissolved C in the martensite would result in strength [19] and deteriorating the corrosion resistance [30]. the increased internal martensite lattice stress, furthermore increasing the hardness, strength [19] and Figure 8 shows the TEM results of 30Cr15Mo1N austenitized at 900 °C. Spherical precipitates in deteriorating the corrosion resistance [30]. the grain interior were observed, as shown in Figure 8a,b. Based on the selected area electron Figure8 shows the TEM results of 30Cr15Mo1N austenitized at 900 ◦C. Spherical precipitates in diffraction (SAED) patterns, the precipitates were identified as Cr-rich M23C6 and M2N, respectively. the grain interior were observed, as shown in Figure8a,b. Based on the selected area electron diffraction The M2N was finer than M23C6, which is consistent with the EPMA results (Figure 7). The high- (SAED) patterns, the precipitates were identified as Cr-rich M C and M N, respectively. The M N magnification images and Cr concentration at the precipitate–matrix23 6 interface2 are illustrated in Figure2 was finer than M23C6, which is consistent with the EPMA results (Figure7). The high-magnification 8c–e. The Cr content of M2N was obviously higher than that of M23C6, and the Cr-depleted zones were images and Cr concentration at the precipitate–matrix interface are illustrated in Figure8c–e. The Cr observed in the interface between M23C6/M2N and matrix. It is noteworthy that the widths of content of M2N was obviously higher than that of M23C6, and the Cr-depleted zones were observed Cr-depleted zones around M23C6 and M2N were about 13.75 nm and 10.16 nm, respectively, and the in the interface between M C /M N and matrix. It is noteworthy that the widths of Cr-depleted minimum Cr contents were23 6about2 13.52 wt % and 15.78 wt %, respectively. Therefore, the zones around M23C6 and M2N were about 13.75 nm and 10.16 nm, respectively, and the minimum Cr Cr-depletion induced by coarse M23C6 was severer than M2N.

(a) (b) (c)

2 μm 2 μm 2 μm

Materials 2017, 10, 861 7 of 19

Figure 6. Morphologies of 30Cr15Mo1N austenitized at: (a) 900 °C; (b) 950 °C; (c) 1000 °C; (d) 1050 °C; (e) 1100 °C; and (f) 1150 °C.

Table 2. Area percentage of precipitates in 30Cr15Mo1N austenitized at different temperatures.

Temperature (°C) 900 950 1000 1050 1100 1150 Area percentage of precipitates (%) 7.75 ± 0.71 6.07 ± 0.96 3.94 ± 0.99 1.59 ± 0.63 1.23 ± 0.48 0.63 ± 0.40

The distribution of Cr, C and N via EPMA in Figure 7 shows that the precipitates were rich in Cr, C and N, and the quantity of N-rich particles was much lower than that of C-rich ones. The N-rich particles were finer and distributed more homogeneously than the C-rich ones. With the increasing of austenitizing temperature, the precipitates dissolved and the distribution of Cr, Mo, C and N became more homogeneous. Besides, the C content in lath martensite at 1000 °C and 1100 °C (indicated by arrows in Figure 7) was remarkably elevated. The extra dissolved C in the martensite would result in the increased internal martensite lattice stress, furthermore increasing the hardness, strength [19] and deteriorating the corrosion resistance [30]. Figure 8 shows the TEM results of 30Cr15Mo1N austenitized at 900 °C. Spherical precipitates in the grain interior were observed, as shown in Figure 8a,b. Based on the selected area electron diffraction (SAED) patterns, the precipitates were identified as Cr-rich M23C6 and M2N, respectively. The M2N was finer than M23C6, which is consistent with the EPMA results (Figure 7). The high- magnification images and Cr concentration at the precipitate–matrix interface are illustrated in Figure 8c–e.Materials The2017 Cr, 10 content, 861 of M2N was obviously higher than that of M23C6, and the Cr-depleted zones 8were of 19 observed in the interface between M23C6/M2N and matrix. It is noteworthy that the widths of Cr-depleted zones around M23C6 and M2N were about 13.75 nm and 10.16 nm, respectively, and the minimumcontents were Cr aboutcontents 13.52 were wt % about and 15.78 13.52 wt wt %, respectively.% and 15.78 Therefore, wt %, respectively. the Cr-depletion Therefore, induced the by coarse M C was severer than M N. Cr-depletion23 6 induced by coarse M223C6 was severer than M2N.

(a) (b) (c)

Materials 2017, 10, 861 2 μm 2 μm 2 8μ ofm 19

Cr Cr Cr Materials 2017, 10, 861 8 of 19

Cr Cr Cr

C C C

C C C

N N N N N N

FigureFigure 7. Element 7. Element mapping mapping of of30Cr15Mo1N 30Cr15Mo1N austenitized austenitized at:at: ( a(a) )900 900 °C; °C; (b () b1000) 1000 °C; °C; and and (c) 1100 (c) 1100 °C °C Figure 7. Element mapping of 30Cr15Mo1N austenitized at: (a) 900 ◦C; (b) 1000 ◦C; and (c) 1100 ◦C by EPMA. by EPMA.by EPMA.

(a) (a) ((bb)

Figure 8. Cont.

Materials 2017,, 10,, 861861 9 of 19 Materials 2017, 10, 861 9 of 19 (c) (c) (e) (e)

(d) (d)

Figure 8. TEM results of 30Cr15Mo1N austenitized at 900 °C: (a,b) morphologies and SAED patterns; Figure 8. TEM results of 30Cr15Mo1N austenitized at 900 ◦C: (a,b) morphologies and SAED patterns; Figure(c ,8.d) TEMhigh-magnification results of 30Cr15Mo1N of precipitates austenitized in (a,b); and at ( e900) Cr °C: content (a,b) profilesmorphologies at the precipitate–matrix and SAED patterns; (c,d) high-magnification of precipitates in (a,b); and (e) Cr content profiles at the precipitate–matrix (c,d) high-magnificationinterface (The inset figure of precipitates is the high magnif in (a,bication); and of ( eCr) Cr contents content in Cr-depletedprofiles at thezones). precipitate–matrix interface (The inset figure is the high magnification of Cr contents in Cr-depleted zones). interface (The inset figure is the high magnification of Cr contents in Cr-depleted zones). 3.3. Electrochemical Measurement 3.3. Electrochemical The potentiodynamic Measurement polarization curves of 30Cr15Mo1N in annealing and austenitizing conditions are presented in Figure 9a. There were obvious passive ranges where the current density The potentiodynamicpotentiodynamic polarization curves of 30Cr15Mo1N30Cr15Mo1N in annealingannealing andand austenitizingaustenitizing conditionsremained are almost presented stable. in However, Figure9 a.several There current were transients obvious passiverelating to ranges the initiation where and the repassivation current density conditionsof metastable are presented pits [20,31] in can Figure be observed 9a. There in passive were regions.obvious The passive current ranges density where increased the dramatically current density remained almost stable. However, several current transients relating to the initiation and repassivation remainedabove almost the pitting stable. potential, However, indicating several the currentstable pitt transientsing corrosion relating occurred. to the As initiationshown in Figure and repassivation 9b, with of metastablethe increasing pits [[20,31]of20 austenitizing,31] can be observed temperature, in passive both the regions. metastable The and current stable dens densitypittingity potentials increased increased, dramatically aboveindicating the the pitting pitting the potential, potential, pit initiation indicating indicating became the themore stable stable diffic pitt pittingingult corrosionand corrosion the corrosion occurred. occurred. resistance As shown As shownwas in Figureenhanced. in Figure 9b, with 9b, withthe increasingIn the addition, increasing of the austenitizing difference of austenitizing between temperature, me temperature,tastable both and the stable both metastable pitting the metastable potent and stialsable also and pitting increased, stable potentials pitting which meant potentialsincreased, increased,indicatingthe repassivation indicatingthe pit initiation thetendency pit initiation ofbecame metastable became more pitting morediffic becameult difficult and larger andthe [32]. thecorrosion corrosion resistance resistance was was enhanced. enhanced. In addition, the difference between metastablemetastable and stable pitting potentialspotentials also increased, which meant the repassivation(a) tendency tendency of of metastable metastable pitting pitting became became (b)larger larger [32]. [32].

(a) (b)

Figure 9. (a) Potentiodynamic polarization curves of 30Cr15Mo1N; and (b) change in metastable and stable pitting potentials with austenitizing temperature.

The morphologies of the largest corrosion pits on 30Cr15Mo1N austenitized at 900 °C, 1000 °C Figureand 1100 9. (°Ca) Potentiodynamicafter potentiodynamic polarization polarization curves are of 30Cr15Mo1N;shown in Figure and 10. (b )The change size inand metastable depth of andpits Figure 9. (a) Potentiodynamic polarization curves of 30Cr15Mo1N; and (b) change in metastable and decreased with the increasing of austenitizing temperature. As shown in Figure 10a,c, the pits in stable pitting potentials withwith austenitizingaustenitizing temperature.temperature. specimens austenitized at 900 °C and 1000 °C were presented as holes with collapsed lacy cover. The Theexistence morphologies of these pits of thewas largest related corrosion to passivatio pitsn onand 30Cr15Mo1N undercutting austenitizednear the pit mouthat 900 [31].°C, 1000In °C Theaddition, morphologies several micro-pits of the largest(indicated corrosion by arrows pits in onFigure 30Cr15Mo1N 10a) nucleated austenitized and grew at at the 900 bottom◦C, 1000 of ◦C and 1100 °C after potentiodynamic polarization are shown in Figure 10. The size and depth of pits and 1100 ◦C after potentiodynamic polarization are shown in Figure 10. The size and depth of pits decreased with the increasing of austenitizing temperature. As shown in Figure 10a,c, the pits in decreased with the increasing of austenitizing temperature. As shown in Figure 10a,c, the pits in specimens austenitized at 900 °C and 1000 °C were presented as holes with collapsed lacy cover. The specimens austenitized at 900 ◦C and 1000 ◦C were presented as holes with collapsed lacy cover. The existence of these pits was related to passivation and undercutting near the pit mouth [31]. In addition, several micro-pits (indicated by arrows in Figure 10a) nucleated and grew at the bottom of

Materials 2017, 10, 861 10 of 19 existence of these pits was related to passivation and undercutting near the pit mouth [31]. In addition, Materials 2017, 10, 861 10 of 19 several micro-pits (indicated by arrows in Figure 10a) nucleated and grew at the bottom of primary pit in specimen austenitized at 900 ◦C. From the high magnification micrograph (image inserted primary pit in specimen austenitized at 900 °C. From the high magnification micrograph (image into Figure 10a), there existed a mass of particles at the bottom of pit, which were confirmed to be inserted into Figure 10a), there existed a mass of particles at the bottom of pit, which were confirmed rich in Cr according to the EDS spectrum (Figure 10b). The combination of chemical composition to be rich in Cr according to the EDS spectrum (Figure 10b). The combination of chemical composition and high-magnification morphology indicated that these particles originated from the uncorroded and high-magnification morphology indicated that these particles originated from the uncorroded precipitates in the matrix. When the austenitizing temperature increased to 1000 ◦C, just several precipitates in the matrix. When the austenitizing temperature increased to 1000 °C, just several particles existed at the pit bottom because of the dissolution of precipitates (Figure 10c). As to the particles existed at the pit bottom because of the dissolution of precipitates (Figure 10c). As to the specimen austenitized at 1100 ◦C, almost no particles existed in shallow dish-typed pit (Figure 10d), specimen austenitized at 1100 °C, almost no particles existed in shallow dish-typed pit (Figure 10d), and the pit was significantly shallower than those austenitized at lower temperatures. and the pit was significantly shallower than those austenitized at lower temperatures.

(a) particles (b)

4 μm EDS 2 μm

micro-pits lacy cover

40 μm (c) (d)

4 μm 4 μm dish-typed pit

lacy cover

20 μm 20 μm Figure 10. Morphologies of the largest corrosion pits in 30Cr15Mo1N austenitized at different Figure 10. Morphologies of the largest corrosion pits in 30Cr15Mo1N austenitized at different temperatures after potentiodynamic polarization: (a) 900 °C; (b) EDS spectrum and high-magnification temperatures after potentiodynamic polarization: (a) 900 ◦C; (b) EDS spectrum and high-magnification of particles in (a); (c) 1000 °C; and (d) 1100 °C. of particles in (a); (c) 1000 ◦C; and (d) 1100 ◦C. To clearly present the metastable pit initiation sites, a specimen austenitized at 900 °C was ◦ mechanicallyTo clearly abraded present and the polished. metastable Then, pit it initiation was observed sites, using a specimen FE-SEM austenitized after potentiodynamically at 900 C was mechanicallypolarized to the abraded metastable and polished. pitting potential. Then, it wasAs shown observed in Figure using FE-SEM11a, there after existed potentiodynamically obvious ditches polarizedaround large-sized to the metastable precipitates, pitting whereas potential. no obviou As showns ditches in Figure around 11a, small-sized there existed precipitates obvious ditches were aroundfound. EDS large-sized spectra precipitates,in Figure 11b whereas indicated no the obvious large and ditches small around precipitates small-sized were Cr-rich precipitates M23C6 were and found.M2N, respectively. EDS spectra Therefore, in Figure 11theb indicatedditches around the large M23 andC6 were small confirmed precipitates to werebe the Cr-rich preferential M23C6 andmetastable M2N, respectively.pit initiation Therefore,sites, which the might ditches propagate around to M form23C6 stablewere confirmedpits. to be the preferential metastable pit initiation sites, which might propagate to form stable pits.

Materials 2017, 10, 861 11 of 19

Materials 2017, 10, 861 11 of 19

(a) (b) Materials 2017, 10, 861 11 of 19 Spectrum for M2N

(a) ditches (b) Spectrum for M2N M2N Spectrum for M C ditches 23 6

M2N

M23C6 Spectrum for M23C6

0.5 μm

M23C6 Figure 11. (a) Morphology of pit initiation sites; and (b) EDS spectra of precipitates in 30Cr15Mo1N Figure 11. a b austenitized( ) Morphology at 900 °C. of pit initiation sites;0.5 μ andm ( ) EDS spectra of precipitates in 30Cr15Mo1N austenitized at 900 ◦C. 3.4. XPSFigure Results 11. (a ) Morphology of pit initiation sites; and (b) EDS spectra of precipitates in 30Cr15Mo1N austenitized at 900 °C. 3.4. XPS ResultsXPS analysis was performed to determine the effect of austenitizing temperature on the XPS3.4.composition XPS analysis Results of passive was performed films. Prior to to the determine tests, the specimens the effect were of polarized austenitizing at 0 mV temperatureSCE (about 200 mV on the higher than OCP) for 1 h to promote the growth of passive films. Figure 12 shows Cr 2p3/2 and Fe 2p3/2 compositionXPS ofanalysis passive was films. performed Prior toto thedetermine tests, the the specimens effect of austenitizing were polarized temperature at 0 mV onSCE the(about spectra of sputtering surface of passive film for 10 s. The Cr 2p3/2 spectra were split into Cr(OH)3 200 mVcomposition higher than of passive OCP) films. for 1 Prior h to to promote the tests, thethe growthspecimens of were passive polarized films. at Figure0 mVSCE 12 (about shows 200 CrmV 2p 3/2 (576.9 eV), Cr2O3 (575.6 eV) and Cr metal (573.7 eV), and the Fe 2p3/2 spectra were divided into FeOOH and Fehigher 2p3/2 thanspectra OCP) offor sputtering 1 h to promote surface the grow of passiveth of passive film for films. 10 s.Figure The 12 Cr shows 2p3/2 Crspectra 2p3/2 and were Fe split2p3/2 into (711.2 eV), Fe3O4 (708.6 eV) and Fe metal (706.5 eV). The binding energies of Cr metal slightly Cr(OH)spectradecreased3 (576.9 of sputtering eV),with Crthe2O surfaceincreasing3 (575.6 of eV) passiveof andaustenitizing Crfilm metal for 10 (573.7 temperature,s. The eV), Cr 2p and 3/2which spectra the Fe was 2pwere 3/2consistent splitspectra into were withCr(OH) dividedthe3 into FeOOH(576.9dissolution eV), (711.2 Cr of2 OCr-rich3 eV), (575.6 Fe precipitates eV)3O 4and(708.6 Cr metalwhich eV) (573.7 andhad Fehigher eV), metal and binding the (706.5 Fe energies. 2p eV).3/2 spectra The With binding were increasing divided energies austenitizing into FeOOH of Cr metal slightly(711.2temperature, decreased eV), Fe the3O with4 peaks (708.6 the for eV) increasing Cr andmetal Fe weakened, ofmetal austenitizing (706.5 demonst eV). temperature,ratingThe binding the thickening energies which of was ofthe Cr passive consistent metal film slightly [33]. with the dissolutiondecreasedThe components of Cr-richwith theof precipitates passive increasing films which ofon 30Cr15Mo1Naustenitizing had higher batemperature, bindingsed on the energies. fittingwhich data Withwas (Figure increasingconsistent 12) are with austenitizing listed the in dissolution of Cr-rich precipitates which had higher binding energies. With increasing austenitizing temperature,Table 3. The the peaksCr2O3 content for Cr metaland ratio weakened, of Cr3+ to demonstrating the sum of Fe2+ the and thickening Fe3+ increased of the with passive increasing film [33]. temperature, the peaks for Cr metal weakened, demonstrating the thickening3+ of the passive film [33]. The componentsaustenitizing oftemperature. passive films The on variation 30Cr15Mo1N in content based of on Cr the2O3 fittingand Cr data in (Figure the passive 12) are films listed in The components of passive films3+ on 30Cr15Mo1N based on the fitting data (Figure 12) are listed in Tabledemonstrates3. The Cr O thatcontent Cr2O3 and and Cr ratio were of Crenriched3+ tothe by enhancing sum of Fe the2+ andaustenitizing Fe3+ increased temperature. with increasing Table 3. The2 Cr3 2O3 content and ratio of Cr3+ to the sum of Fe2+ and Fe3+ increased with increasing 3+ austenitizingaustenitizing temperature. temperature. The variationThe variation in content in content of Cr2 Oof3 andCr2O Cr3 andin theCr3+ passive in the filmspassive demonstrates films 3+(a) that Crdemonstrates2O3 and Cr thatwere Cr2O enriched3 and Cr3+ were by enhancing enriched by the enhancing austenitizing the austenitizing temperature. temperature.

(a)

(b)

(b)

(c)

(c)

Figure 12. X-ray photoelectron spectra of Cr 2p3/2 and Fe 2p3/2 recorded from the passive films after sputtering for 10 s on 30Cr15Mo1N austenitized at (a) 900 °C; (b) 1000 °C; and (c) 1100 °C.

Figure 12. X-ray photoelectron spectra of Cr 2p3/2 and Fe 2p3/2 recorded from the passive films after Figure 12. X-ray photoelectron spectra of Cr 2p3/2 and Fe 2p3/2 recorded from the passive films after sputtering for 10 s on 30Cr15Mo1N austenitized at (a) 900 °C; (b) 1000 °C; and (c) 1100 °C. sputtering for 10 s on 30Cr15Mo1N austenitized at (a) 900 ◦C; (b) 1000 ◦C; and (c) 1100 ◦C.

Materials 2017, 10, 861 12 of 19

Materials 2017, 10, 861 12 of 19 Table 3. Component of passive films on 30Cr15Mo1N austenitized at 900 ◦C, 1000 ◦C and 1100 ◦C. Table 3. Component of passive films on 30Cr15Mo1N austenitized at 900 °C, 1000 °C and 1100 °C. Component of Passive Films (at %) Specimens Component of Passive Films (at %) Cr3+/(Fe2+ + Fe3+) 3+ 2+ 3+ SpecimensFe 3O4 FeOOH Cr2O3 Cr(OH)3Cr /(Fe + Fe ) Fe3O4 FeOOH Cr2O3 Cr(OH)3 ◦ 900 C900 °C 12.51 12.51 4.68 4.68 6.97 6.97 4.52 4.52 0.24 0.24 1000 ◦C 11.39 2.49 8.47 6.34 0.29 1000 °C 11.39 2.49 8.47 6.34 0.29 1100 ◦C 11.92 2.56 9.04 5.33 0.31 1100 °C 11.92 2.56 9.04 5.33 0.31

FigureFigure 1313 showsshows thethe X-rayX-ray photoelectronphotoelectron spectra spectra of of N N 1s 1s and and Mo Mo 3p 3p3/23/2 recorded from the passive filmsfilms afteraftersputtering sputtering for for 0 0 s ands and 10 10 s. Ons. On the the outmost outmost surface surface of passive of passive films, films, peaks peaks representing representing NH3 (399.8NH3 (399.8 eV), Cr eV),2N (397.2Cr2N (397.2 eV), CrN eV), (396.8 CrN eV)(396.8 for NeV) 1s for and N MoO 1s and2 (396.1 MoO eV)2 (396.1 for Mo eV) 3p for3/2 wereMo 3p detected.3/2 were Afterdetected. sputtering After sputtering for 10 s, the for spectra 10 s, the exhibited spectra peaksexhibited for Crpeaks2N, CrNfor Cr and2N, MoOCrN2 and, whereas MoO2, peaks whereas for NHpeaks3 disappeared. for NH3 disappeared. As the austenitizing As the austenitizing temperature temp increased,erature increased, the intensity the intensity of Cr2N inof theCr2N passive in the filmspassive decreased films decreased while the while intensity the intensity of CrN increased, of CrN increased, which is which associated is associated with the with dissolution the dissolution of M2N precipitatesof M2N precipitates at higher at austenitizing higher austenitizing temperature. temperature.

(a)

(b)

(c)

FigureFigure 13.13. X-ray photoelectron spectraspectra ofof NN 1s1s andand MoMo3p 3p3/23/2 recordedrecorded from from the ou outmosttmost and sputtering surfacessurfaces of the the passive passive films films for for 10 10 s son on 30Cr15Mo1N 30Cr15Mo1N austenitized austenitized at: at:(a) (900a) 900°C; ◦(bC;) 1000 (b) 1000 °C; and◦C; ◦ and(c) 1100 (c) 1100 °C. C.

3.5.3.5. Immersion Tests

TheThe pittingpitting corrosioncorrosion andand generalgeneral corrosioncorrosion occurredoccurred duringduring immersionimmersion teststests inin 66 wtwt %% FeClFeCl33 atat 5050 ±± 0.50.5 °C,◦C, thus both of them should be considered.considered. The general corrosioncorrosion ratesrates forfor 30Cr15Mo1N30Cr15Mo1N austenitizedaustenitized atat 900900 ◦°C,C, 10001000 ◦°CC andand 11001100◦ °CC areare 23.53 23.53 mm/year, mm/year, 4.104.10 mm/yearmm/year andand 1.331.33 mm/year,mm/year, respectively.respectively. The The specimen specimen austenitized austenitized at 900at 900◦C contained°C contained massive massive Cr-rich Cr-rich precipitates precipitates and the and lowest the chromiumlowest chromium content incontent the matrix, in the thus matrix, having thus the havi highestng the general highest corrosion general rate. corrosion With the rate. increasing With the of austenitizingincreasing of temperature,austenitizing the temperature, chromium contentthe chromium in matrix content increased, in matrix inducing increased, the enhanced inducing general the corrosionenhanced resistance.general corrosion The evolution resistance. of maximum The evolut pit depthion of considering maximum generalpit depth corrosion considering on specimens general corrosion on specimens austenitized at 900 °C, 1000 °C and 1100 °C after being immersed for 3 h, 6 h, 12 h and 24 h is shown in Figure 14. It has been reported that the pit growth follows Equation (2) [34,35]:

Materials 2017, 10, 861 13 of 19 austenitized at 900 ◦C, 1000 ◦C and 1100 ◦C after being immersed for 3 h, 6 h, 12 h and 24 h is shown inMaterials Figure 2017 14, .10 It, 861 has been reported that the pit growth follows Equation (2) [34,35]: 13 of 19

n dmax = kt (2) dmax = ktn (2) µ µ −n where ddmaxmax (μ(m)m) is isthe the maximum maximum pit depth, pit depth, k (μm/hk ( −nm/h) and n) are and constants,n are constants, and t (h) is and thet immersion(h) is the immersiontime. It can time. be seen It can that be the seen maximum that the maximum pit depth pitincreased depth increased with increasing with increasing immersion immersion time, which time, whichfollowed followed the power the powerlaw. With law. the With increasing the increasing of austenitizing of austenitizing temperature, temperature, the maximum the maximum pit depth pit depthdecreased decreased significantly. significantly. The fitted The parameters fitted parameters in Figure in Figure14 showed 14 showed that the that values the valuesof n merely of n merelyvaried varied(approximately (approximately 0.5) and 0.5) the and values the values of k of decreasedk decreased significantly significantly with with increasing increasing austenitizing temperature, revealing aa lowerlower pitpit growthgrowth raterate forfor specimenspecimen austenitizedaustenitized atat higherhigher temperature.temperature.

Figure 14. Evolution of maximum pit depth with immersion time for 30Cr15Mo1N austenitized at Figure 14. Evolution of maximum pit depth with immersion time for 30Cr15Mo1N austenitized at 900 °C, 1000 °C and 1100 °C. 900 ◦C, 1000 ◦C and 1100 ◦C. 4. Discussion 4. Discussion Based on the above results, the reason for preferential metastable pit initiation around coarse Based on the above results, the reason for preferential metastable pit initiation around coarse M23C6, and the relationship between microstructure and corrosion resistance of martensitic high M C , and the relationship between microstructure and corrosion resistance of martensitic high nitrogen23 6 stainless steel 30Cr15Mo1N at different austenitizing temperatures by analyzing pit nitrogen stainless steel 30Cr15Mo1N at different austenitizing temperatures by analyzing pit initiation, initiation, passive film and pit growth kinetics will be discussed in detail. passive film and pit growth kinetics will be discussed in detail.

4.1. Explanation of Preferential Metastable Pit Initiation Sites

The martensitic high nitrogen stainless steel 30Cr15Mo1N contained Cr-rich M23C6 and M2N, The martensitic high nitrogen stainless steel 30Cr15Mo1N contained Cr-rich M23C6 and M2N, which were different inin distribution,distribution, size,size, volumevolume fraction,fraction, chemicalchemical composition and Cr-depleted zone (Figures 2, 7 and 8). Initially, the M2N precipitated during annealing process was more coherent zone (Figures2,7 and8). Initially, the M 2N precipitated during annealing process was more coherent than M23C6 in nitrogen alloyed MSSs [18,36], resulting in the smaller size and more homogeneous than M23C6 in nitrogen alloyed MSSs [18,36], resulting in the smaller size and more homogeneous distribution of M2N than M23C6. During the austenitization process, M2N and M23C6 partially distribution of M2N than M23C6. During the austenitization process, M2N and M23C6 partially dissolved, andand the the nitrogen-induced nitrogen-induced short short range range atomic atomic ordering ordering and the and strong the Cr–N strong bond Cr–N prevented bond theprevented chromium the clustering chromium in 30Cr15Mo1Nclustering in [1630Cr15Mo,36]. Then1N in quenching[16,36]. Then process, in quenching the martensite process, inherited the martensite inherited from the austenite possessed a homogeneous element distribution. The content from the austenite possessed a homogeneous element distribution. The content of M2N was lower of M2N was lower than M23C6 because the atomic fraction of N in M2N was higher than C in M23C6. than M23C6 because the atomic fraction of N in M2N was higher than C in M23C6. Besides, the addition ofBesides, nitrogen the suppressed addition of the nitrogen precipitation suppressed of coarse the pr eutecticecipitation carbides, of coarse which eutectic was due carbides, to higher which austenite was stabilitydue to higher upon quenchingaustenite stability [36] and upon higher quenching binging energy [36] and of Cr–N higher than binging that of energy Cr–C [ 16of]. Cr–N Similar than results that haveof Cr–C been [16]. obtained Similar in results Cr15Mo1 have [16 been] and obtained SUS440A in [17 Cr15Mo1] with different [16] and nitrogen SUS440A contents. [17] with different nitrogenIt is contents. well accepted that pitting corrosion initiates from breakdown of passive film or chemical/physicalIt is well accepted heterogeneity, that pitting such as corrosion inclusions initiates and precipitates from breakdown [37,38]. The of detrimental passive effectfilm or of chemical/physical heterogeneity, such as inclusions and precipitates [37,38]. The detrimental effect of Cr-rich precipitates [39–42], such as Cr23C6, Cr2N and σ, on corrosion resistance due to Cr-depletion Cr-rich precipitates [39–42], such as Cr23C6, Cr2N and σ, on corrosion resistance due to Cr-depletion has been widely reported. In the present study, Cr-rich M23C6 and M2N were generated in the annealing process, resulting in the emerging of Cr-depleted zones in the vicinity of precipitate/matrix interface simultaneously. During the austenitization process, M23C6 and M2N partially dissolved into the matrix. As shown in Figure 8, coarse M23C6 induced wider and severer Cr-depleted zone than

Materials 2017, 10, 861 14 of 19

has been widely reported. In the present study, Cr-rich M23C6 and M2N were generated in the annealing process, resulting in the emerging of Cr-depleted zones in the vicinity of precipitate/matrix interface simultaneously. During the austenitization process, M23C6 and M2N partially dissolved into the matrix. As shown in Figure8, coarse M 23C6 induced wider and severer Cr-depleted zone than M2N, which resulted in the preferential initiation of metastable pits around M23C6 (Figure 11). On the one hand, the lower Cr content in Cr-depleted zone induced poorer stability of passive film, which was vulnerable to become initiation sites for pitting corrosion. On the other hand, the chemical composition diversity between matrix and Cr-depleted zone led to the formation of micro-cell with narrow anodic zone [36], thus increasing the pitting corrosion sensitivity.

4.2. Influence of Austenitizing Temperature on Pit Initiation As discussed in Section 4.1, metastable pits initiated around Cr-rich precipitates, preferentially the coarse M23C6 with severer Cr-depletion. The content of precipitates decreased with increasing austenitizing temperature (Figure6), which meant the number of pit initiation sites was reduced. Choi et al. [21] and Park et al. [22] reported that the pitting potential of MSSs was enhanced as the austenitizing temperature increased, which was attributed to the decrease of carbides. In the present research, the metastable and stable pitting potentials increased with increasing austenitizing temperature, revealing less susceptible metastable pit initiation and higher corrosion resistance. Moreover, the difference between metastable and stable pitting potentials at low austenitizing temperature was small, which indicates that the transition of metastable pits to stable pits occurred soon after the initiation of metastable pits and the repassivation ability was weak. With the increasing of austenitizing temperature, the difference between metastable and stable pitting potentials increased simultaneously, demonstrating that although the initiation of metastable pits occurred at relatively low potential, but the metastable pits could be transformed into stable pits only at much higher potential [32]. Additionally, it is generally believed that high temperature promotes the atomic diffusion. As shown in Figure3, with increasing austenitizing temperature, the decreased Cr and Mo contents in M23C6 and M2N and the increased Cr, Mo, C and N contents in matrix resulted in higher degree of composition homogeneity in 30Cr15Mo1N. The Cr atoms diffused from both precipitates and matrix to the Cr-depleted zones where the Cr content was the lowest. The Cr content of Cr-depleted zones was compensated to a higher extent when austenitized at higher temperature due to more rapid diffusion of Cr atoms, thus also delaying the initiation of pitting corrosion. The correlation between content of precipitates and pitting potentials of 30Cr15Mo1N austenitized at different temperatures is shown in Figure 15. The variation of pitting potentials could be divided ◦ into three stages. At the temperature below 1000 C (Stage I), although M23C6 dissolved gradually, the existence of massive M23C6 induced the relatively low and slow increase of the metastable and stable pitting potentials with the increasing of temperature. The M23C6 dissolved completely ◦ around 975 C, and the content of M23C6 decreased obviously when the temperature reached 1000 ◦C (Figure7). Accordingly, a sharp rise of metastable pitting potential (from 950 to 1000 ◦C) was observed in Figure 15. Since then, the metastable pitting potential increased stably with the increment of austenitizing temperature. When the temperature increased to 1000 ◦C, the content of precipitates decreased obviously (Figure6), and the precipitates almost disappeared at temperatures ◦ higher than 1050 C. The dissolution of M2N could be obtained from the Thermo-Calc results from 1000 to 1075 ◦C. In accordance with this, the increasing rate of stable pitting potential accelerated, particularly in the range of 1000 ◦C to 1050 ◦C (Stage II). After that, only a few undissolved precipitates existed, and the decrease rate of precipitate content reduced when the temperature was higher than 1050 ◦C (Stage III), which was consistent with the slow increasing rate of stable pitting potential. Therefore, the determining factor of the three stages for pitting potential could be dissolution of M23C6 ◦ ◦ (below 1000 C), dissolution of M2N (from 1000 to 1050 C) and existence of a few undissolved precipitates and non-metallic inclusions (above 1050 ◦C), respectively. Materials 2017, 10, 861 15 of 19 Materials 2017, 10, 861 15 of 19

Figure 15. Correlation between content of precipitates and pitting potentials of 30Cr15Mo1N Figure 15. Correlation between content of precipitates and pitting potentials of 30Cr15Mo1N austenitized at different temperatures. austenitized at different temperatures.

4.3. Influence of Austenitizing Temperature on Passive Film 4.3. Influence of Austenitizing Temperature on Passive Film With increasing austenitizing temperature, the Cr-rich precipitates dissolved and the Cr content With increasing austenitizing temperature, the Cr-rich precipitates dissolved and the Cr in matrix was enhanced, promoting the thickening of the passive films and enrichment of Cr2O3 and content3+ in matrix was enhanced, promoting the thickening of the passive films and enrichment Cr (Figure 12 and3+ Table 3). Besides, the nitrogen content in matrix increased as austenitizing of Cr2O3 and Cr (Figure 12 and Table3). Besides, the nitrogen content in matrix increased temperature rose, which is beneficial to the enrichment of Cr2O3 in the passive films [10,33]. as austenitizing temperature rose, which is beneficial to the enrichment of Cr O in the passive Additionally, the interaction of nitrogen and molybdenum promoted the deprotonation2 3 of Cr-hydroxide films [10,33]. Additionally, the interaction of nitrogen and molybdenum promoted the deprotonation by inducing O-H bond stretching, which also promoted the enrichment of Cr2O3 [33,43]. On the other of Cr-hydroxide by inducing O–H bond stretching, which also promoted the enrichment of hand, compared with Cr(OH)3, Cr2O3 had higher thermodynamic stability [44] and lower point defect Cr O [33,43]. On the other hand, compared with Cr(OH) , Cr O had higher thermodynamic concentration2 3 [45], thus enhancing the stability of passive film3 and 2hindering3 pit initiation [37]. stability [44] and lower point defect concentration [45], thus enhancing the stability of passive film and Based on the deconvolution of N spectra (Figure 13), two kinds of nitrides, i.e., CrN and Cr2N, hindering pit initiation [37]. were detected. The Cr2N corresponded to the undissolved Cr-rich M2N, which had high Volta Based on the deconvolution of N spectra (Figure 13), two kinds of nitrides, i.e., CrN and Cr N, were potential [46] and hardly corroded in the corrosion process, as shown in Figure 10. Therefore,2 it could detected. The Cr N corresponded to the undissolved Cr-rich M N, which had high Volta potential [46] not consume protons2 in pits and had no positive effect on 2corrosion resistance of 30Cr15Mo1N. and hardly corroded in the corrosion process, as shown in Figure 10. Therefore, it could not consume However, nitrogen in solid solution was anodically segregated to be CrN in the passive film during protons in pits and had no positive effect on corrosion resistance of 30Cr15Mo1N. However, nitrogen pitting corrosion process [7]. The CrN could consume protons to form ammonium ions by in solid solution was anodically segregated to be CrN in the passive film during pitting corrosion Equation (3) [47]: process [7]. The CrN could consume protons to form ammonium ions by Equation (3) [47]: 2CrN + 3H2O → Cr2O3 + 2NH3 (3) 2CrN + 3H O → Cr O + 2NH (3) consequently hindering the self-catalytic process2 and2 promoting3 3 the repassivation of pits [33,48]. consequentlyWith increasing hindering austenitizing the self-catalytic temperature, process the dissolution and promoting of M2N the induced repassivation the decrease of pits of [33 Cr,482N]. intensity in passive film. Thereafter, the enhanced content of nitrogen in solid solution promoted the With increasing austenitizing temperature, the dissolution of M2N induced the decrease of Cr2N intensityincreasing in of passive CrN content film. Thereafter, in the passive the enhanced film, which content could of consume nitrogen more in solid protons solution and promoted promote the increasingrepassivation of CrNof incipient content pits. in the passive film, which could consume more protons and promote the repassivation of incipient pits. 4.4. Effect of Austenitizing Temperature on Pit Growth Kinetics 4.4. EffectAfter of the Austenitizing initiation of Temperature metastable on pit, Pit Growthit might Kinetics propagate to form a stable pit and grow self- catalytically.After the The initiation pit growth of followed metastable the pit, power it mightlaw (Figure propagate 14), and to its form rate awas stable determined pit and by grow the self-catalytically.values of n and k The. In pitthe growthpresent followed work, the the values power of law n were (Figure approximately 14), and its rate0.5 independent was determined of the by theaustenitizing values of ntemperature.and k. In the Cavanaugh present work, et al. the [48] values reported of n were that the approximately value of n equaled 0.5 independent to 0.5 when of the austenitizingpit growth was temperature. under ohmic Cavanaugh or diffusion et al. control. [48] reported The values that theof k value considerably of n equaled decreased to 0.5when with the pitincreasing growth of was austenitizing under ohmic temperature, or diffusion indicating control. The lower values pit growth of k considerably rate, which decreased was in agreement with the increasingwith the pit of morphologies austenitizing after temperature, potentiodynamic indicating polarization lower pit growth(Figure rate,10). Nitrogen which was in solid in agreement solution withcould the be pitreleased morphologies to the pits after in the potentiodynamic form of NH3 and/or polarization NH4+ by (Figure Equations 10). (4) Nitrogen and (5), in respectively, solid solution in the pitting corrosion process [33,43]:

[N] + 3H+ + 3e− → NH3 (4)

Materials 2017, 10, 861 16 of 19

+ could be released to the pits in the form of NH3 and/or NH4 by Equations (4) and (5), respectively, in the pitting corrosion process [33,43]:

Materials 2017, 10, 861 + − 16 of 19 [N] + 3H + 3e → NH3 (4)

+ − → + [N][N] + 4H + 4H++ 3e+ 3e− →NHNH44+ (5)(5) andand protons protons in in pits were partially consumed. Th Thee higher higher austenitizing austenitizing temperature temperature enhanced enhanced the the contentcontent ofof nitrogennitrogen in in solid solid solution, solution, promoting promoting the repassivation the repassivation ability ofability pits [33 of], whichpits [33], contributed which contributedto lower pit to growth lower ratepit growth and shallower rate and pits. shallower pits. BasedBased on on the the analysis analysis of of microstructure, microstructure, electroc electrochemicalhemical behavior, behavior, passive passive film, film, pit pit initiation initiation and and propagation,propagation, the the relationship relationship between between microstruc microstructureture and and corrosion corrosion resistance resistance of of 30Cr15Mo1N at at differentdifferent austenitizing austenitizing temperatures temperatures is schematically is schematically presented presented in Figure in Figure 16. When 16. When the specimens the specimens were polarizedwere polarized to high to potential high potential in chlo inride chloride solution, solution, metastable metastable pits init pitsiated initiated preferentially preferentially around aroundcoarse Mcoarse23C6 with M23 severerC6 with Cr-depletion. severer Cr-depletion. With fewer With pit initiati feweron pit sites initiation and more sites protective and more passive protective film (thicker passive 3+ andfilm enriched (thicker andin Cr enriched2O3, Cr3+ inand Cr 2CrN),O3, Cr the andspecimen CrN), austenitized the specimen at austenitized higher temperature at higher corroded temperature less severelycorroded than less that severely austenitized than that at austenitizedlower temperatur at lowere. With temperature. the proceeding With theof pitting proceeding corrosion, of pitting pits grewcorrosion, by self-catalytic pits grew by mechan self-catalyticism. The mechanism. uncorroded The Cr-rich uncorroded M23C6 and Cr-rich M2N M 23accumulatedC6 and M2N at accumulated the bottom ofat pits the in bottom specimen of pitsaustenitized in specimen at lower austenitized temperature. at lower The higher temperature. austenitizing The temperature higher austenitizing induced enhancedtemperature content induced of nitrogen enhanced in solid content solution of nitrogen in the inmatrix, solid solutionwhich promoted in the matrix, the repassivation which promoted process the andrepassivation retarded the process growth and of retardedpits, resulting the growth in shallower of pits, pits. resulting in shallower pits.

(a) (b)

FigureFigure 16. 16. SchematicSchematic models models for for corrosion corrosion mechan mechanismism of 30Cr15Mo1N of 30Cr15Mo1N austenitized austenitized at: (a at:) lower; (a) lower; and (andb) higher (b) higher temperatures. temperatures.

5.5. Conclusions Conclusions In the present work, the relationship between microstructure and corrosion behavior of In the present work, the relationship between microstructure and corrosion behavior of martensitic martensitic high nitrogen stainless steel 30Cr15Mo1N at different austenitizing temperatures was high nitrogen stainless steel 30Cr15Mo1N at different austenitizing temperatures was investigated investigated using microscopy observation, electrochemical measurement, passive film analysis and using microscopy observation, electrochemical measurement, passive film analysis and immersion immersion testing. The main conclusions could be obtained as follows: testing. The main conclusions could be obtained as follows: (1) With increasing austenitizing temperature, the fraction of precipitates decreased and retained (1) With increasing austenitizing temperature, the fraction of precipitates decreased and retained austenite increased, resulting in more homogeneous distribution and higher contents of Cr, Mo, austenite increased, resulting in more homogeneous distribution and higher contents of Cr, Mo, C and N in the matrix. The precipitates were identified as Cr-rich M23C6 and M2N, and M2N was C and N in the matrix. The precipitates were identified as Cr-rich M23C6 and M2N, and M2N was finer and distributed more homogeneously than M23C6. The Cr-depleted zone around M23C6 was finer and distributed more homogeneously than M23C6. The Cr-depleted zone around M23C6 wider and its minimum Cr concentration was lower than M2N. was wider and its minimum Cr concentration was lower than M2N. (2) The metastable pits initiated preferentially around coarse M23C6 which induced severer Cr- depletion, and the pit growth followed the power law. The dissolution of M23C6 and M2N at higher austenitizing temperature reduced the pit initiation sites. As austenitizing temperature increased, the metastable and stable pitting potentials increased and the pit growth rate decreased, revealing less susceptible metastable pit initiation, larger repassivation tendency and higher corrosion resistance. The determining factor of pitting potential could be divided into

Materials 2017, 10, 861 17 of 19

(2) The metastable pits initiated preferentially around coarse M23C6 which induced severer Cr-depletion, and the pit growth followed the power law. The dissolution of M23C6 and M2N at higher austenitizing temperature reduced the pit initiation sites. As austenitizing temperature increased, the metastable and stable pitting potentials increased and the pit growth rate decreased, revealing less susceptible metastable pit initiation, larger repassivation tendency and higher corrosion resistance. The determining factor of pitting potential could be divided into three ◦ ◦ stages: dissolution of M23C6 (below 1000 C), dissolution of M2N (from 1000 to 1050 C) and existence of a few undissolved precipitates and non-metallic inclusions (above 1050 ◦C). (3) The increasing of austenitizing temperature promoted the thickening of passive film; enrichment 3+ of Cr2O3, Cr and CrN; and higher nitrogen content in solid solution, thereby enhancing the stability of passive film and repassivation ability of pits.

Acknowledgments: This work was supported by the National Natural Science Foundation of China (grant numbers 51434004, U1435205, and 51304041) and the Fundamental Research Funds for the Central Universities (grant number N150204007). Author Contributions: Zhouhua Jiang, Hao Feng, and Huabing Li conceived and designed the experiments; Hao Feng, Hongchun Zhu and Shucai Zhang performed the experiments; Hao Feng, Binbin Zhang and Yu Han analyzed the data; and Hao Feng, Huabing Li, Tao Zhang and Dake Xu contributed to writing and editing of the manuscript. Conflicts of Interest: The authors declare no conflict of interest.

References

1. Franz-Josef, E. An overview of performance characteristics, experiences and trends of aerospace engine bearings technologies. Chin. J. Aeronaut. 2007, 20, 378–384. 2. Berns, H.; Ebert, F.-J.; Zoch, H.-W. The new low nitrogen steel LNS—A material for advanced aircraft engine and aerospace bearing applications. In Bearing Steels: Into the 21st Century; Hoo, J.J.C., Green, W.B., Eds.; ASTM International: West Conshohocken, PA, USA, 1998; pp. 354–373. 3. Li, H.B.; Zhou, E.Z.; Ren, Y.B.; Zhang, D.W.; Xu, D.K.; Yang, C.G.; Feng, H.; Jiang, Z.H.; Li, X.G.; Gu, T.Y.; et al. Investigation of microbiologically influenced corrosion of high nitrogen nickel-free stainless steel by Pseudomonas aeruginosa. Corros. Sci. 2016, 111, 811–821. [CrossRef] 4. Li, H.B.; Jiang, Z.H.; Yang, Y.; Cao, Y.; Zhang, Z.R. Pitting corrosion and crevice corrosion behaviors of high nitrogen austenitic stainless steels. Int. J. Miner. Met. Mater. 2009, 16, 517–524. [CrossRef] 5. Ha, H.-Y.; Lee, C.-H.; Lee, T.-H.; Kim, S. Effects of nitrogen and tensile direction on stress corrosion cracking susceptibility of Ni-free FeCrMnC-based duplex stainless steels. Materials 2017, 10, 294. [CrossRef] 6. Li, H.-B.; Jiang, Z.-H.; Shen, M.-H.; You, X.-M. High nitrogen austenitic stainless steels manufactured by nitrogen gas alloying and adding nitrided ferroalloys. J. Iron Steel Res. Int. 2007, 14, 64–69. [CrossRef] 7. Levey, P.R.; Vanbennekom, A. A mechanistic study of the effects of nitrogen on the corrosion properties of stainless steels. Corrosion 1995, 51, 911–921. [CrossRef] 8. Lu, Y.; Bandy, R.; Clayton, C.; Newman, R. Surface enrichment of nitrogen during passivation of a highly resistant stainless steel. J. Electrochem. Soc. 1983, 130, 1774–1776. [CrossRef] 9. Clayton, C.R.; Martin, K.G. Evidence of anodic segregation of nitrogen in high nitrogen stainless steels and its influence on passivity. In Proceedings of the High Nitrogen Steels—HNS 88, Lille, France, 18–20 May 1988. 10. Clayton, C.R.; Halada, G.P.; Kearns, J.R. Passivity of high nitrogen stainless alloys—The role of metal oxyanions and salt films. Mater. Sci. Eng. A 1995, 198, 135–144. [CrossRef] 11. Valentin, G.; Gavriljuk, H.B. High Nitrogen Steels: Structure, Properties, Manufacture, Applications; Springer: Berlin/Heidelberg, Germany, 1999. 12. Li, H.-B.; Jiao, W.-C.; Feng, H.; Jiang, Z.-H.; Ren, C.-D. Influence of austenitizing temperature on the microstructure and corrosion resistance of 55Cr18Mo1VN high-nitrogen plastic mould steel. Acta Metall. Sin. 2016, 29, 1148–1160. [CrossRef] 13. El Mehtedi, M.; Ricci, P.; Drudi, L.; El Mohtadi, S.; Cabibbo, M.; Spigarelli, S. Analysis of the effect of deep cryogenic treatment on the hardness and microstructure of X30 CrMoN 15 1 steel. Mater. Des. 2012, 33, 136–144. [CrossRef] Materials 2017, 10, 861 18 of 19

14. Speidel, M.O. Properties and applications of high nitrogen steels. In Proceedings of the High Nitrogen Steels—HNS 88, Lille, France, 18–20 May 1988. 15. Kaluba, W.J.; Kaluba, T.; Taillard, R. The austenitizing behaviour of high-nitrogen martensitic stainless steels. Scr. Mater. 1999, 41, 1289–1293. [CrossRef] 16. Gavriljuk, V.G.; Berns, H. Precipitates in tempered stainless martensitic steels alloyed with nitrogen, carbon or both. Mater. Sci. Forum 1999, 318, 71–80. [CrossRef] 17. Shimada, T.; Yamamoto, A.; Abe, S. The effect of nitrogen content on material properties of high carbon stainless steel SUS440A. Tetsue-to-Hagane 1996, 82, 309–314. [CrossRef] 18. Trojahn, W.; Streit, E.; Chin, H.; Ehlert, D. Progress in bearing performance of advanced nitrogen alloyed stainless steel, Cronidur 30. Mater. Sci. Eng. Technol. 1999, 30, 605–611. [CrossRef] 19. Isfahany, A.N.; Saghafian, H.; Borhani, G. The effect of heat treatment on mechanical properties and corrosion behavior of AISI420 martensitic stainless steel. J. Alloys Compd. 2011, 509, 3931–3936. [CrossRef] 20. Lu, S.-Y.; Yao, K.-F.; Chen, Y.-B.; Wang, M.-H.; Shao, Y.; Ge, X.-Y. Effects of austenitizing temperature on the microstructure and electrochemical behavior of a martensitic stainless steel. J. Appl. Electrochem. 2015, 45, 375–383. [CrossRef] 21. Choi, Y.-S.; Kim, J.-G.; Park, Y.-S.; Park, J.-Y. Austenitizing treatment influence on the electrochemical corrosion behavior of 0.3C-14Cr-3Mo martensitic stainless steel. Mater. Lett. 2007, 61, 244–247. [CrossRef] 22. Park, J.; Park, Y. Effects of austenitizing treatment on the corrosion resistance of 14Cr-3Mo martensitic stainless steel. Corrosion 2006, 62, 541–547. [CrossRef] 23. Lu, S.-Y.; Yao, K.-F.; Chen, Y.-B.; Wang, M.-H.; Ge, X.-Y. Influence of heat treatment on the microstructure and corrosion resistance of 13 wt pct Cr-type martensitic stainless steel. Metall. Mater. Trans. A 2015, 46, 6090–6102. [CrossRef] 24. Geng, X.; Feng, H.; Jiang, Z.; Li, H.; Zhang, B.; Zhang, S.; Wang, Q.; Li, J. Microstructure, mechanical and corrosion properties of friction stir welding high nitrogen martensitic stainless steel 30Cr15Mo1N. Metals 2016, 6, 301. [CrossRef] 25. Leem, D.-S.; Lee, Y.-D.; Jun, J.-H.; Choi, C.-S. Amount of retained austenite at room temperature after reverse transformation of martensite to austenite in a Fe-13%Cr-7%Ni-3%Si martensitic stainless steel. Scr. Mater. 2001, 45, 767–772. [CrossRef] 26. Database for Surface Spectroscopies as XPS, AES and UPS. Available online: http://www.lasurface.com (accessed on 11 November 2016). 27. ASTM. ASTM G48–11. Standard Test Methods for Pitting and Crevice Corrosion Resistance of Stainless Steels and Related Alloys by Use of Ferric Chloride Solution; ASTM International: West Conshohocken, PA, USA, 2015. 28. ISO. ISO 8407: 2009(E). Corrosion of Metals and Alloys—Removal of Corrosion Products from Corrosion Test Specimens; International Standards Organization: Geneva, Switzerland, 2009. 29. Schlüter, K.; Shi, Z.; Zamponi, C.; Cao, F.; Quandt, E.; Atrens, A. Corrosion performance and mechanical properties of sputter-deposited MgY and MgGd alloys. Corros. Sci. 2014, 78, 43–54. [CrossRef] 30. Candelaria, A.; Pinedo, C. Influence of the heat treatment on the corrosion resistance of the martensitic stainless steel type AISI 420. J. Mater. Sci. Lett. 2003, 22, 1151–1153. [CrossRef] 31. Bai, G.; Lu, S.; Li, D.; Li, Y. Influences of niobium and solution treatment temperature on pitting corrosion behaviour of stabilised austenitic stainless steels. Corros. Sci. 2016, 108, 111–124. [CrossRef] 32. Tang, Y.; Zuo, Y.; Wang, J.; Zhao, X.; Niu, B.; Lin, B. The metastable pitting potential and its relation to the pitting potential for four materials in chloride solutions. Corros. Sci. 2014, 80, 111–119. [CrossRef] 33. Fu, Y.; Wu, X.; Han, E.-H.; Ke, W.; Yang, K.; Jiang, Z. Effects of nitrogen on the passivation of nickel-free high nitrogen and manganese stainless steels in acidic chloride solutions. Electrochim. Acta 2009, 54, 4005–4014. [CrossRef] 34. Muto, I.; Sato, E.; Ito, S. Pitting corrosion behavior of stainless steels in a marine environment and its estimation method. Zairyo-to-Kankyo 1993, 42, 714–720. [CrossRef] 35. Woldemedhin, M.; Shedd, M.; Kelly, R. Evaluation of the maximum pit size model on stainless steels under thin film electrolyte conditions. J. Electrochem. Soc. 2014, 161, E3216–E3224. [CrossRef] 36. Berns, H.; Ehrhardt, R. Carbon or nitrogen alloyed quenched and tempered stainless steels—A comparative study. Steel Res. Int. 1996, 67, 343–349. [CrossRef] 37. Liu, J.; Zhang, T.; Meng, G.; Shao, Y.; Wang, F. Effect of pitting nucleation on critical pitting temperature of 316L stainless steel by nitric acid passivation. Corros. Sci. 2015, 91, 232–244. [CrossRef] Materials 2017, 10, 861 19 of 19

38. Frankel, G. Pitting corrosion of metals—A review of the critical factors. J. Electrochem. Soc. 1998, 145, 2186–2198. [CrossRef] 39. Della Rovere, C.; Aquino, J.; Ribeiro, C.; Silva, R.; Alcântara, N.; Kuri, S. Corrosion behavior of radial friction welded supermartensitic stainless steel pipes. Mater. Des. 2015, 65, 318–327. [CrossRef] 40. Zhang, S.; Jiang, Z.; Li, H.; Feng, H.; Zhang, B. Detection of susceptibility to intergranular corrosion of aged super austenitic stainless steel S32654 by a modified electrochemical potentiokinetic reactivation method. J. Alloys Compd. 2017, 695, 3083–3093. [CrossRef] 41. Chan, K.W.; Tjong, S.C. Effect of secondary phase precipitation on the corrosion behavior of duplex stainless steels. Materials 2014, 7, 5268–5304. [CrossRef] 42. Moteshakker, A.; Danaee, I. Microstructure and corrosion resistance of dissimilar weld-joints between duplex stainless steel 2205 and austenitic stainless steel 316L. J. Mater. Sci. Technol. 2016, 32, 282–290. [CrossRef] 43. Habibzadeh, S.; Li, L.; Shum-Tim, D.; Davis, E.C.; Omanovic, S. Electrochemical polishing as a 316L stainless steel surface treatment method: Towards the improvement of biocompatibility. Corros. Sci. 2014, 87, 89–100. [CrossRef] 44. Lv, J.; Liang, T.; Wang, C. Surface enriched molybdenum enhancing the corrosion resistance of 316L stainless steel. Mater. Lett. 2016, 171, 38–41. 45. Fuertes, N.; Pettersson, R. Passive film properties and electrochemical response of different phases in a Cu-alloyed stainless steel after long term heat treatment. J. Electrochem. Soc. 2016, 163, C377–C385. [CrossRef] 46. Willenbruch, R.; Clayton, C.; Oversluizen, M.; Kim, D.; Lu, Y. An XPS and electrochemical study of the influence of molybdenum and nitrogen on the passivity of austenitic stainless steel. Corros. Sci. 1990, 31, 179–190. [CrossRef] 47. Lu, Y.; Ives, M.; Clayton, C. Synergism of alloying elements and pitting corrosion resistance of stainless steels. Corros. Sci. 1993, 35, 89–96. [CrossRef] 48. Cavanaugh, M.; Buchheit, R.; Birbilis, N. Modeling the environmental dependence of pit growth using neural network approaches. Corros. Sci. 2010, 52, 3070–3077. [CrossRef]

© 2017 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).