Exceptional cavitation erosion-corrosion behavior of dual-phase bimodal structure in austenitic stainless steel

Exceptional cavitation erosion-corrosion behavior of dual-phase bimodal structure in austenitic stainless steel

Accepted Manuscript Exceptional cavitation erosion-corrosion behavior of dual-phase bimodal structure in austenitic stainless steel Karthik Selvam, Ja...

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Accepted Manuscript Exceptional cavitation erosion-corrosion behavior of dual-phase bimodal structure in austenitic stainless steel Karthik Selvam, Jaskaran Saini, Gopinath Perumal, Aditya Ayyagari, Riyadh Salloom, Riya Mondal, Sundeep Mukherjee, Harpreet Singh Grewal, Harpreet Singh Arora PII:

S0301-679X(19)30030-1

DOI:

https://doi.org/10.1016/j.triboint.2019.01.018

Reference:

JTRI 5559

To appear in:

Tribology International

Received Date: 11 November 2018 Revised Date:

19 January 2019

Accepted Date: 21 January 2019

Please cite this article as: Selvam K, Saini J, Perumal G, Ayyagari A, Salloom R, Mondal R, Mukherjee S, Grewal HS, Arora HS, Exceptional cavitation erosion-corrosion behavior of dual-phase bimodal structure in austenitic stainless steel, Tribology International (2019), doi: https://doi.org/10.1016/ j.triboint.2019.01.018. 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 proof before it is published in its final 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.

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Exceptional Cavitation Erosion-corrosion Behavior of Dual-phase Bimodal Structure in Austenitic Stainless Steel

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Karthik Selvam1, Jaskaran Saini1, Gopinath Perumal1, Aditya Ayyagari2,3, Riyadh Salloom2, Riya Mondal4, Sundeep Mukherjee2, Harpreet Singh Grewal1, Harpreet Singh Arora1, # 1

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Department of Mechanical Engineering, School of Engineering, Shiv Nadar University, Uttar Pradesh 201314, India 2 Department of Materials Science and Engineering, University of North Texas, Denton, Texas 76203, USA 3 Center for Nanoscale Materials, Argonne National Laboratory, 9700 S. Cass Avenue, Argonne, IL 60439, USA 4 Department of Metallurgical Engineering and Materials Science, Indian Institute of Technology Bombay, Mumbai, India

Corresponding Author E-mail: [email protected]

Abstract

In this study, we demonstrate a novel pathway to engineer the properties of metallic alloys for

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limiting their cavitation erosion-corrosion. A facile single-step processing technique was used to develop bimodal grain structure in stainless steel. The bimodal steel was tested in cavitation erosion and erosion-corrosion conditions. In both the cases, bimodal steel showed exceptionally high degradation resistance, nearly 7 times higher compared to as-received steel. The remarkable

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cavitation erosion resistance demonstrated by bimodal steel is attributed to its high yield strength along with high work-hardening rate. In addition, the bimodal steel showed significantly low

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corrosion rate of 0.001 mm/year in 3.5 wt. % NaCl solution compared to 0.088 mm/year for asreceived stainless steel.

Keywords: Cavitation erosion-corrosion; Surface modification; Bimodal grain structure; Electrochemistry; X-ray photoelectron

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1. Introduction Cavitation erosion, caused by implosion of vapor bubbles, leads to extensive surface damage and fatigue failure of materials and machineries in marine environments [1]. Maritime industries, in particular, are affected by extensive cavitation erosion of marine vessels, ship

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propellers, and control valves. Corrosive environment may further exacerbate material loss by breakdown of surface passive layer [2-6]. Owing to the aggressive nature of cavitation environment, even high strength materials such as alloy steels, nickel and cobalt-based superalloys degrade under typical operating conditions [3, 7-9]. Recently, high entropy alloys

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(HEAs) with attractive mechanical and corrosion properties have shown encouraging cavitation erosion resistance [10, 11]. However, higher manufacturing cost, homogenization and scale-up

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for industrial applications pose significant challenges for their wide spread adoption [12-14]. In contrast, microstructure modification of existing marine-grade materials may be a much more cost-effective path forward.

Primary failure during cavitation occurs through propagation of fatigue cracks from cyclic stresses caused by micro-jet and shock waves [6, 13, 15-17]. Increasing grain boundary area enhances a material’s yield strength and may help arrest the propagation of these fatigue

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cracks. This makes ultra-fine grain structures effective in enhancing cavitation resistance [6, 1822]. Further, the strain hardenability of a material also plays a vital role in circumventing degradation by cavitation [23, 24]. A recent study on Al0.1CoCrFeNi single phase HEA with grain size of nearly 2 mm showed significantly higher cavitation resistance compared to stainless

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steel (SS316L) [10]. Exceptional cavitation resistance of the HEA was attributed to its high strain-hardening index. Thus, a microstructure that results in a combination of high yield strength

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and large strain-hardening rate could be promising for the cavitation environment. Typically, ultrafine grain (UFG) structure lacks work-hardening ability due to limited accommodation of dislocations [25, 26]. Bimodal grain structures comprising of fine grains in the matrix of coarse grains are promising in this regard. Fine grains in the bimodal microstructure increase the yield strength while coarse grains provide high ductility and strain-hardening ability. However, the coarser grains in bimodal structure lowers the yield strength compared to UFG microstructure of the same alloy composition [27-29]. Therefore, materials with a bimodal grain structure with yield strength comparable or better than the UFG microstructure could be potentially transformative for applications demanding high cavitation resistance. Realization of a bimodal 2

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grain structure is challenging and is limited by the multiple complex processing steps. These may involve a combination of severe plastic deformation processing, heat treatment and/or cryogenic processing [29-31]. Here, we demonstrate a facile processing route for achieving dual-phase bimodal grain structure in marine grade stainless steel with exceptional resistance against

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cavitation erosion. The bimodal steel also showed superior corrosion and erosion-corrosion resistance.

2. Experimental procedure

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2.1. Material selection

Commercially available stainless steel plates of SS316L, annealed and air cooled, with

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dimensions 75 cm × 50 cm × 0.5 cm was used for the current study. The measured chemical composition and mechanical properties of the as-received material are listed in Table 1 and Table 2 respectively.

Table 1: Measured chemical composition of austenite stainless steel (SS316L) C

Mn

Si

Wt. %

0.024

1.43

0.57

P

S

Cr

Mo

Ni

Fe

0.032

0.01

16.8

2.1

10.1

Balance

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Elements

Table 2: Experimentally obtained mechanical properties of austenite stainless steel (SS316L) Ultimate Strength

Tensile strength

580 (MPa) ± 25

Yield strength

290 (MPa) ± 20

Percentage elongation

45 ± 3

Hardness (Vickers)

226 HV0.025 ± 10

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Mechanical properties

2.2. Experimental method Surface properties of the stainless steel were tailored using submerged friction stir

processing performed using a pin-less tool made of tungsten carbide with 12 mm shoulder diameter. The following two different processing conditions were used: (1) Submerged friction stir processing (SFSP): here, the tool was traversed along the longitudinal direction of the workpiece at a rotational speed of 1800 rpm, while submerged in a pool of liquid (mixture of distilled water and ethanol in equal proportion) at 0° C; (2) Stationary submerged friction stir processing 3

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(SSFSP): in this case, the tool was rotated at 1800 rpm at a particular location of the work-piece for a period of 15 minutes while submerged in the pool of liquid. The tool plunge depth used in both cases was 400 µm. The FSP fixture used for holding the work-piece was connected to an external chiller through inlet and outlet ports. As-received and processed specimens were

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polished down to 3000 grit followed by electro-polishing in 10% oxalic acid solution at 650 mV for 2 minutes. The grain size and phase distribution were obtained using electron back scatter diffraction (EBSD) and X-ray diffraction. For transmission electron microscopy (TEM) studies, samples were prepared using FEI Nova NanoLab 200TM focused ion beam (FIB). FEI Tecnai

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F20 field emission gun TEM was used to probe the microstructure of bimodal and UFG steel. For tensile testing, dog-bone shaped mini-tensile specimens with dimensions of 5 mm x 1.25 mm

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x 0.4 mm were prepared and tested at room temperature at a strain rate of 10-3 s-1. Cavitation erosion tests were carried out using an ultrasonic vibratory oscillator (Sonics: VCX750) at 20 ± 0.5 KHz frequency and peak amplitude of 50 µm ± 5 %, as per ASTM standard G32. Cavitation setup used in the current study is an indirect test method, where samples were placed at a standoff distance of 500 µm from a replaceable titanium probe. Samples were mounted on a tool fixture and immersed in distilled water for cavitation erosion

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testing, while the medium used for cavitation erosion-corrosion was 3.5 wt. % NaCl solution. The temperature of the distilled water was maintained at 25 ± 5º C using a copper coil connected to the chiller. Figure 1 shows a schematic representation of a vibratory apparatus used in the current study. Mass loss measurements during cavitation erosion tests were done using high

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precision weighing balance with an accuracy of 0.01 mg. Cavitation erosion test was conducted for a duration of 20 hours and mass loss was measured for every hour. All the experiments were

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done three times to ensure the repeatability. The surface of all eroded samples was examined using a scanning electron microscope (SEM) to understand cavitation erosion mechanisms. Hardness measurements of samples before and after cavitation testing was done across the specimen cross-section using micro-hardness tester at 25 gm load (Hv0.025). Corrosion behaviour of as-received and processed samples was investigated by open

circuit potential, potentiodynamic polarization and electrochemical impedance spectroscopy (EIS) using Gamry Interface 1000E electrochemical setup. Each test was repeated at least three times to ensure repeatability of the results. A standard three electrode cell configuration was used with saturated calomel electrode (SCE) as reference, high density graphite rod as counter, and 4

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sample as working electrode. 3.5 wt. % NaCl solution was used as the electrolyte for corrosion studies. Before running the corrosion tests, the open circuit potential (EOCP) for each sample was measured for a period of 3600 seconds to keep fluctuations within 0.1 mV. Potentiodynamic polarization was done in the voltage range of -0.25 V vs EOCP to +0.4 V vs EOCP with forward

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and reverse scan rate of 0.166 mV/s. EIS measurements were obtained at EOCP over a frequency range of 0.01 Hz to 100 kHz with a set AC voltage amplitude of 10 mV. The Gamry Echem analyst 7.03 was used to model the electrical equivalent circuit (EEC) and fit to Nyquist and Bode plots using simplex algorithm. The thin oxide film formed on as-received and processed

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samples were characterized by X-ray photoelectron spectroscopy (XPS). XPS spectra of Cr, Fe and O were recorded on the surface of all the samples using monochromatic Al-Kα X-ray source

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(1.486 keV, Scienta Omicron Nanotechnology). XPS peak fitting, elemental identification, and

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atomic percentage were obtained by Casa XPS software (V2.3.16) and XPSPEAK 4.1.

Figure 1: Schematic of cavitation erosion and cavitation erosion-corrosion set-up with ultra-

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sonicator and sample holder. Specimens with dimensions 10 mm × 10 mm × 5 mm were used for cavitation erosion and cavitation erosion-corrosion.

3. Results and Discussion

3.1 Microstructure and Mechanical Properties Figure 2 shows a schematic of SFSP and SSFSP used in this study for obtaining two different microstructures in SS316L: (1) ultra-fine grain structure obtained using SFSP (hereafter referred as UFG) and (2) bimodal grain structure obtained using SSFSP (hereafter referred as 5

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BM). Severe plastic deformation during linear traversing on the surface concurrent with forced cooling from the submerging liquid resulted in restricted grain growth of recrystallized grains resulting in ultra-fine grain structure. In contrast, prolonged rotation of the tool at a particular location on the work-piece resulted in bimodal structure, likely due to partial recrystallization.

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Figure 2(a) and Figure 2(b) shows the EBSD map of BM and UFG specimen respectively. The bimodal specimen had fine grains embedded in the matrix of coarse grains (~ 10 µm) with an average grain size of nearly 3.5 µm. In contrast, UFG had fine equiaxed grain structure with an average grain size of nearly 0.9 µm compared to 22 µm for the as-received steel. The processing

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depth was found to be nearly 250-300 µm in both the cases. TEM images of the bimodal specimen (Figure 3(a) and (b)) showed fine grains of the order of 200-400 nm, fine deformation

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bands and sub-grains. EBSD phase map and XRD analysis for all the specimens are given in

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Figure S1

Figure 2: Schematic representation of friction stir processing for developing bimodal grain structure and ultra-fine grain structure in stainless steel. Traversing the rotating tool resulted in ultra-fine grain structure while stationary processing produced bimodal grain structure; Electron

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back scatter diffraction (EBSD) image for as-received [18] (a) bimodal specimen (BM) and (b) ultra-fine grain (UFG) structure. (supplementary information). As-received alloy primarily contained austenite as shown by both

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XRD and phase map. EBSD phase map as well as XRD analysis for UFG indicated small fraction (~ 8 %) of martensite phase formation. EBSD phase map for BM (Figure S1 (c)) indicated fine martensite grains with nearly 30 % volume fraction embedded in the matrix of

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coarse austenite grains, resulting in a dual- phase micro-structure.

Figure 3: Transmission electron microscopy (TEM) image of bimodal steel showing (a) fine

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deformation bands and fine grains, (b) sub-grain structure [32].

Austenite to martensite transformation during processing is attributed to high strain-rate

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deformation. The fine slip bands, their intersections, dislocation bundles and other defects act as a nucleation site for martensite grains [33-35]. Further, deformation induced nano-twins and sub grain structures with high stacking faults increases the inelastic strain energy. After a critical strain, such dislocation clusters tend to transform into martensite. However, the accumulation of internal energy through dislocations is limited by the effect of recovery, thus constraining the continual growth of martensite grains [36, 37]. Continuous growth of martensite grains is also hindered by the stabilized austenite, resulting in a dual phase bimodal structure. The true stress-strain plot for all the three specimens is shown in Figure 4(b). As-received steel showed a yield strength of 300 MPa with 40 % true strain. As expected, UFG specimen 7

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demonstrated a higher yield strength of 420 MPa with reduced true strain of 30 %. Interestingly, the bimodal specimen showed highest yield strength of 620 MPa along with 32 % true strain. Thus, bimodal specimen demonstrated superior mechanical properties, even better than the UFG structure. The work-hardening rate as a function of strain for all cases is shown in Figure 4(c).

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All specimen showed nearly similar variation of work-hardening rate with increasing strain. The initial high work-hardening rate is followed by a lower and nearly steady rate with continuous decrease towards the end. Here also, bimodal steel showed highest work-hardening rate followed by UFG steel and the as-received steel. The dual-phase bimodal grain structure can be

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considered as a dispersion strengthened system [38]. For dispersion strengthened alloys, the yield strength primarily depends on the pinning of dislocations around the particles and consequently

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delay the onset of plastic deformation. Mathematically, yield strength for dual-phase alloys is given as:

0

where



0.41

[39]

is the initial flow stress constant,

is the true stress,

is the true strain, is the shear

modulus, f is the volume fraction of second phase, d is the particle size, and b is the burgers /

/

. An increase in

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vector. The equation suggests a direct correlation between σ and

volume fraction of secondary phase can increase the yield strength, while decreasing the particle size can enhance the work-hardening rate [38]. The dual-phase bimodal grain structure has large volume fraction of fine martensite grains which contributes towards enhancing the yield strenth

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and work-hardening rate simultaneously. Thus, the unique properties of bimodal specimen are

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attributed to its dual-phase microstructure.

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Figure 4: (a) Engineering stress-strain curve, (b) True stress-strain curve and (c) work hardening

3.2 Cavitation Erosion and Erosion-Corrosion

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rate for as-received, bimodal and ultra-fine grain steel specimen [32].

The results of cavitation erosion for all samples are shown as cumulative mass loss in Figure 5(a). The time required for mass loss to reach 1 mg (minimum measurable mass) was considered for calculating the incubation period which represents the period of insignificant

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material loss. As-received steel showed the least incubation period of about 2 hours while both UFG and BM specimen showed significantly higher incubation period of 7 hours and 11 hours

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respectively. Substantially higher incubation period for BM specimen indicates that bimodal grain structure was able to limit the damage initiation for a significantly longer time compared to the as-received and ultra-fine grain steel. This is likely attributed to higher yield strength as well as high work-hardening rate of bimodal grain structure (Figure 4(b) and (c)), both of which enhances the resistance to plastic deformation. For the same alloy system, linear dependence of incubation period on yield strength has been shown in previous studies as well [6, 40-42].

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Following the initial incubation period, the cumulative mass loss (CML) for all specimen showed an acceleration period of rapid material loss which finally ends in a steady state condition. The slope of acceleration part of CML plot is very steep for the as-received steel compared to both the processed specimen. BM steel showed the highest resistance in the acceleration stage as well,

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characterized by the lowest value of slope. The CML after 20 hrs of pure cavitation erosion was 24 mg for the as-received stainless steel, 6 mg for UFG and 3.5 mg for BM specimen. Compared

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to as-received steel, ultra-fine grain steel showed nearly 4 times higher erosion resistance, while it is nearly 7 times higher for the bimodal steel. This is again attributed to superior mechanical properties of bimodal steel, especially high yield strength and work-hardening rate, which significantly enhanced its degradation resistance. The CML during cavitation erosion-corrosion for all samples is shown in Figure 5(b). Compared to pure erosion, all samples showed lower incubation period and increase in overall mass loss. As-received steel showed an incubation period of about 1 hour while it was nearly 5 hours for both UFG and BM specimen. Similarly, the mass loss increased to nearly 32 mg for the as-received steel, 10 mg and 8 mg for UFG and BM specimen respectively. Thus, during cavitation erosion-corrosion also, the bimodal steel 9

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showed nearly 4 times higher resistance compared to the as-received steel. The reduction in degradation resistance during cavitation erosion-corrosion is likely due to rapid removal of the work-hardened layer by metal removal through passivation. The continuous removal of work-

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material removal during cavitation corrosion.

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hardened layer exposes the underneath surface to impact of micro-jets which increases the

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Figure 5: (a) Cumulative mass loss for all samples subjected to cavitation erosion for 20 hours, (b) Cumulative mass loss for all samples subjected to cavitation erosion-corrosion for 20 hours, (c) sub-surface hardness of post-cavitation (pure erosion) tested samples and (d) XRD analysis of post-cavitation tested samples.

The work-hardening by implosion of bubbles during cavitation is evident from the sub-surface hardness of post cavitation tested samples, shown in Figure 5(c).

All samples showed a

significant increase in hardness post cavitation which was largest for the bimodal specimen (~ 1.5 times), followed by as-received and ultra-fine grained specimen. This is primarily attributed 10

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to grain refinement by high strain deformation of the surface during cavitation as well as dislocation accumulation and strain induced phase transformation [43-45]. The XRD analysis of the post cavitation samples, shown in Figure 5(d), indicates the formation of martensite [46]. In addition, there was significant peak broadening which supports grain refinement during

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cavitation [47].

Figure 6: Comparison of mean depth of erosion during cavitation erosion and erosion-corrosion

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for bimodal steel (current study) with different high entropy alloys (HEA) and amorphous coatings and bulk materials. Bimodal stainless steel showed significantly higher erosion

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resistance compared to HEA and amorphous coatings. The mean depth of erosion for bimodal steel is comparable to bulk HEA in both cavitation erosion and erosion-corrosion conditions. The performance of the developed bimodal stainless steel was also compared with other

advanced materials (bulk/coatings) and is shown in Figure 6. Interestingly, the cavitation erosion and erosion-corrosion resistance of bimodal steel is substantially higher compared to all the investigated high entropy alloy coatings. Further, the MDE are comparable to recently investigated bulk high-entropy alloy, Al0.1CoCrFeNi [10, 11]. In addition, the bimodal steel was far superior compared to iron-based amorphous coating [48]. High entropy alloy and amorphous metals are multi-component/multi-principle alloy systems with exceptional mechanical 11

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properties. In particular, HEAs have been reported to exhibit remarkable degradation resistance which is attributed to their unusually high work-hardening ability [12, 49, 50]. The demonstration of better/similar performance by bimodal grain structure in stainless steel is of

3.3 Electrochemical Corrosion Behavior 3.3.1

Potentiodynamic polarization

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particular significance which is attributed to its unique dual-phase bimodal microstructure.

The results of potentiodynamic polarization are shown in Figure 7(a). As-received steel

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showed active-passive behavior during anodic polarization followed by rapid surge in the current density. The anodic polarization curve for both the processed alloys showed rapid passivation

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with subsequent jump in current density similar to as-received alloy. BM showed higher Ecorr compared to UFG specimen, indicating nobler characteristic for the former. Further, the passive current density (ipass) for as-received steel was significantly higher compared to UFG steel, while ipass was not clearly discernable for the bimodal specimen. Tafel fitting was used to obtain corrosion current density (icorr) and the values are summarized in Table 3. Icorr for BM was the lowest and an order of magnitude smaller compared to the as-received alloy. The corrosion

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current translated to corrosion penetration rate of 0.001 mm/year for the bimodal specimen compared to 0.017 mm/year and 0.088 mm/year for UFG and the as-received specimen. Thus, bimodal specimen demonstrated exceptionally low corrosion rate, nearly 8-10 times lower compared to ultrafine grain stainless steel. Lower corrosion rate for bimodal alloy compared to

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ultrafine grain structure in a passive electrolyte is rather unusual. UFG structure is likely to form a more uniform and intact passive layer due to smaller standard deviation in the grain size

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distribution [51, 52]. This unusual behavior is discussed further in the subsequent section (section 3.4). Low magnification scanning electron microscopy (SEM) images of the corroded surfaces for all specimens are shown in Figure 7(b) to 7(d). While as-received steel showed extensive pitting, both UFG and BM showed very few and smaller pits. The pit depth was measured in each case and was found to be nearly 200 µm for as-received steel, 115 µm for UFG and nearly 85 µm for the bimodal specimen (Figure S2).

3.2.2. Electrochemical Impendence Spectroscopy

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The results of EIS testing are shown as Nyquist and Bode plot in Figure 8. The impedance was significantly higher for both the processed samples compared to the as-received alloy (Figure 8(a)). The zoomed in region of the impedance spectra is shown in Figure 8(b). The phase angle vs frequency plot for all samples, given in Figure 8(c), showed a broad peak in the

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mid-frequency range which is indicative of the capacitive characteristic of the system. The

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capacitive

Figure 7: (a) Potentiodynamic polarization curves for as-received steel, bimodal and ultra-fine grain specimen; Scanning electron microscope images of the corroded surfaces for (b) asreceived steel (c) ultra-fine grain (UFG) and (d) bimodal (BM) specimen.

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Table 3: Open-circuit potential (OCP), corrosion current (Icorr), corrosion potential (Ecorr) and corrosion rate per year for the as-received steel, UFG and BM specimens in 3.5 % NaCl solution

As-received Steel UFG BM

-321 -198 -127

Icorr ×10-6 (A/cm2) 7.7 1.5 0.11

Ecorr (V) -0.359 -0.200 -0.164

Corrosion rate (mm/year) 0.088 0.017 0.001

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OCP (mV)

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Sample

Figure 8: Electrochemical Impedance Spectroscopy (EIS) results showing (a) Nyquist plot for as-received, ultra-fine grain (UFG) and bimodal (BM) specimen; (b) zoon-in region of Nyquist plot shown in (a); (c) and (d) Bode plot for all the three specimen. The electrochemical equivalent circuit (EEC) is shown as an inset in (c) and (d).

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Table 4: Electrochemical parameters obtained after fitting the EIS Data of the as-received steel, UFG, and BM specimens in 3.5 % NaCl solution R2 (k ohm cm2)

RP (k ohm cm2)

CPE1 (µ S sn cm-2)

n1

CPE 2 (µ S sn cm-2)

n2

Total CPE (µ S sn cm-2

RS (k ohm cm2)

Goodnes s of Fit

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Sample

R1 (k ohm cm2)

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1.50 3.98 5.5 280.4 0.71 943.40 0.58 1223.80 6.23 0.81x10-3 As0 received 49.05 49.05 0.81 117.60 7.76 7.23x10-3 UFG 87.71 7.54 2.38x10-3 128.60 128.60 0.86 BM characteristic is the result of charge transfer and formation of double layer capacitance at metal/electrolyte interface where, the diameter of the capacitive loop indicates the charge transfer resistance (Rct). Higher Rct signifies lower dissolution rate and correspondingly higher

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corrosion resistance. The Nyquist plot for the BM specimen showed the largest radius signifying highest corrosion resistance amongst all specimen. The Bode plot for the as-received alloy suggests two-time constants (Figure 8(c) and 8(d)), whereas both the processed specimens showed a single time constant. Therefore, a two-time-constant EEC for the as-received alloy and a single-time-constant EEC for both the processed specimens were used to model the

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electrochemical system. The EEC for as-received steel and both the processed specimens are shown as insets in Figure 8(c) and Figure 8(d) respectively. This electrical circuit has been previously shown to give a best fit for austenitic stainless steels [53]. Here, Rs is the solution or electrolyte resistance, CPE1 and R1 are the electrochemical double layer capacitance and charge

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transfer resistance, respectively, for the high-frequency part of the spectrum, whereas CPE2 and R2 are the corresponding elements for the low-frequency spectrum. CPE, a frequency-dependent constant phase element with exponent n, was used instead of pure capacitance to account for

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surface in-homogeneities, such as roughness, adsorption, and diffusion. The values of different elements of the electrochemical equivalent circuit are summarized in Table 4. The total resistance for UFG and BM specimen is given by R1 while it is the summation of R1 and R2 for the as-received alloy. The total resistance was found to be highest for BM followed by UFG and the as-received alloy. Charge transfer resistance shown by BM was noticeably higher compared to other two specimens. This is in line with Tafel results, supporting highest corrosion resistance for the bimodal specimen followed by UFG and the as-received alloy. In addition, the n value for both the processed samples was higher, indicating better homogeneity of the passive layer. The 15

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likely cause of superior corrosion behavior of processed samples is discussed in the forthcoming section (section 3.4).

3.4.XPS Analysis

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High resolution XPS analysis of the thin oxide layer for all samples are shown in Figure 9. Figure 9(a) shows the XPS wide scan for as-received steel, ultra-fine grain (UFG) and bimodal (BM) specimen while Figure 9(b) gives the quantification of the oxide layer in all the three specimen. Figure 9(c) to 9(e) shows the XPS scans of Cr-2P, Fe-2P and O-1s at the surface of all

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specimen. It is seen that the composition of oxide layer across all samples was similar, primarily

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Figure 9: (a) XPS wide scan for as-received steel, ultra-fine grain (UFG) and bimodal (BM) specimen, (b) Quantification of the composition of oxide layer, High resolution XPS scans of (c) Cr-2P (d) Fe-2P and (e) O-1s at the surface in all the three specimen.

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consisting of oxides of Fe and Cr. Chromium oxide (Cr2O3) and ferric oxide (Fe2O3) were the major oxides formed in all the cases (Figure 9(a) and 9(b)). Further, oxygen was mainly present in the form of O2- and OH- and no metal hydroxides were found (Figure 9(c)). OH- peaks likely result from the adsorbed moisture on the oxide layer surface. XPS results indicate that processing

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did not substantially change the nature of oxide layer formed on stainless steel. Figure 9(d) summarizes the atomic percent of ionic forms of iron and chromium along with oxygen. The atomic fraction for all elements was highest for the bimodal specimen, followed by ultra-finegrained steel and lowest for the as-received steel. Higher atomic fraction of elements in the oxide

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indicates higher diffusivities of these elements in the bimodal steel compared to other two specimen. Typically, fine grain structure enhances the atomic diffusivity due to availability of

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more grain boundary area. The fine martensite grains in the bimodal specimen (200-400 nm, Figure 2) may have resulted in higher diffusivity. Besides grain size, the diffusivity significantly depends on crystal structure and packing fraction [54]. Martensite has body-centered tetragonal structure for which the atomic packing fraction is dependent on the ratio of its lattice parameters i.e. ⁄ . Further, the ⁄ ratio for the martensite phase is a function of carbon fraction and is typically equal to 1 for C wt. % < 0.5 [55]. The carbon fraction in stainless steel SS 316L is

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nearly 0.03% which indicates that martensite in SS316L forms with

⁄ ratio of ≤ 1. Also, the

presence of BCC phase in the selected area diffraction pattern (SADP) analysis of bimodal specimen (Figure S3) supports the presence of martensite phase with ⁄ ratio = 1. The atomic packing fraction for martensite with ⁄

1 is 68 %, equivalent to the BCC crystal. Typically,

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diffusivity is higher for BCC crystal compared to FCC due to its more open structure. Therefore, both finer grain size and lower atomic packing fraction for martensite phase in dual-phase

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bimodal steel likely resulted in faster diffusion of oxygen, chromium and iron.

3.5 Synergy

Cavitation erosion-corrosion synergy has a substantial influence on the degradation

behavior. Material could exhibit either positive or negative synergy depending on the passivating medium. Generally, positive synergy causes excessive material degradation and is typically observed in SS316L. Positive synergy is primarily through the combined influence of erosion enhanced corrosion (EIC) and corrosion enhanced erosion (CIE). Hence, assessing their

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individual roles could provide further insight into their degradation behavior. Synergy could be empirically expressed through the equation [56]:

where, VT is total volume loss during cavitation erosion-corrosion, VE is the volume loss from

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pure erosion, VEIC is measured from polarization curve using Faraday’s law in pure erosion for 20 hours, VC is the volume loss calculated from icorr under static condition and VCIE is measured from the above empirical equation. Figure 10(a) shows the calculated volume loss from individual components and Figure 10(b) shows percentage contribution of each factor towards

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the total volume loss. The results showed a positive synergy for all the tested samples. Further material loss is aggravated mostly through mechanical and corrosion enhanced erosion while

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erosion enhanced corrosion is negligible in all the cases. This is attributed to the continuous removal of passive layer, leading to the exposure of underneath material to corrosion and erosion. Both the processed samples show lower synergy compared to the as-received steel and is the lowest for BM specimen. This is likely due to superior cavitation erosion as well as

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corrosion resistance of processed specimen.

Figure 10: (a) Total volume loss (VT) in 3.5% NaCl, volume loss for pure erosion (VE), volume loss from pure corrosion (VC), volume loss from erosion enhanced corrosion (VEIC) and volume loss from corrosion enhanced erosion (VCIE) for as-received, ultra-fine grain (UFG) and bimodal (BM) specimen (b) percentage contribution of individual components: erosion, corrosion and synergy in total volume loss. 3.6 Damage mechanism 19

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SEM images showing topography of eroded samples are shown in Figure 11(a) to 11(d). The analysis of the figures show a significant difference in the degree of degradation. Compared to BM and UFG samples, the surface of the as-received steel was covered by significantly deep and larger craters and cracks. Further, as-received and UFG specimens primarily exhibited

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almost similar degradation mechanism comprising of pits, micro-cracks, visible striations and craters. Striations indicate plastic deformation of the material under fatigue loading through slip band formation. Prolonged exposure to cyclic strain causes micro-cracks to nucleate at the slip bands which further propagates with further deformation. Generally, grain boundaries act as a

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barrier for the crack propagation, thus increase in grain boundary area can restraint the crack propagation locally. This is consistent with the degradation mechanism observed for the UFG

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specimen (Figure 11(b)). Compared to the as-received steel, UFG show shallow pits, cracks and lesser striations which is the consequence of finer grain structure. In contrast, BM specimen showed significantly different degradation mechanism comprising of micro-scale ductile deformation, fine dimples and micro-plastic tearing (Figure 11(c)), typically referred as tearing topography surface (TTS). The zoomed in image of the region marked in Figure 11(c) is shown in Figure 11(d). TTS has been commonly observed in martensitic steels when exposed to cyclic

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stresses, featured by extrusion at localized regions with micro plastic deformation[57]. The fine dimples in BM specimen indicate localized extrusion through pining during plastic deformation, thus confirming the dispersion strengthening effect in dual-phase bimodal steel. The damage mechanism was found to be similar for both cavitation erosion and erosion-corrosion conditions.

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The results of current study suggest that bimodal grain structure could be highly promising for applications demanding high cavitation and cavitation-corrosion resistance, surpassing the

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performance of ultra-fine grain structures.

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Figure 11: Scanning electron microscope (SEM) images for (a) as-received steel, (b) ultra-fine

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grain (UFG) specimen, (c) bimodal (BM) specimen, (d) zoomed in image of the region marked

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in (c) after 20 hours of cavitation erosion.

Conclusion

A facile single step processing technique was utilized for developing bimodal grain structure in austenitic stainless steel, SS316L, and its cavitation erosion and erosion-corrosion behavior was evaluated. 1. The microstructure of bimodal specimen comprise of fine martensite grains (< 500 nm) embedded in coarse austenite grains (~ 10 µm). The bimodal specimen 21

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demonstrated remarkable mechanical properties with higher yield strength compared to ultra-fine grained steel. 2.

The bimodal steel showed exceptional cavitation erosion and cavitation erosioncorrosion resistance, ranging between 4 to 7 times the as-received steel.

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3. The superior performance of bimodal steel is attributed to its higher yield strength and work-hardening rate which significantly delayed the damage initiation and the overall material loss during cavitation. 4.

In addition, the bimodal steel showed substantially lower corrosion rate which was

5.

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attributed to higher stability of the passive layer with higher chromium fraction. The reduction in degradation resistance during cavitation erosion-corrosion was

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attributed to the removal of work-hardened layer during cavitation through passivation and consequent breakdown.

Declarations of interest: none Authors Contribution:

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HSA and HSG plan the whole study; KS and JS carried out processing of stainless steel to develop bimodal structure; KS and JS performed cavitation and cavitation corrosion studies; GP performed all corrosion studies on as-cast and processed samples; KS and JS performed XPS studies on all samples; AA and RS did sample preparation using EDM and performed all tensile testings on miniature tensile tester under the supervision of SM; AA and RS also did SEM and TEM studies on all samples under the supervision of SM; RM performed XRD and EBSD studies of all samples. KS and HSA wrote different sections of the manuscript with the help of HSG and SM. All authors reviewed the manuscript.

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Highlights

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Bimodal grain structure was developed in austenitic stainless-steel. Bimodal structure showed exceptional cavitation erosion-corrosion resistance Bimodal structure showed substantially lower corrosion rate in 3.5 % NaCl solution Unique dual-phase bimodal grain structure contributed to remarkable performance Performance of bimodal steel was found to be similar to high entropy alloys

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o o o o o