Enhancement of wear and corrosion resistance of iron-based hard coatings deposited by high-velocity oxygen fuel (HVOF) thermal spraying

Surface & Coatings Technology 249 (2014) 24–41

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Enhancement of wear and corrosion resistance of iron-based hard coatings deposited by high-velocity oxygen fuel (HVOF) thermal spraying Wu-Han Liu a,b, Fuh-Sheng Shieu a,⁎, Wei-Tien Hsiao b a b

Department of Materials Science and Engineering, National Chung Hsing University, Taichung 402, Taiwan Materials and Chemical Research Laboratories, Industrial Technology Research Institute, Chutung 310, Taiwan

a r t i c l e

i n f o

Article history: Received 27 December 2012 Accepted in revised form 21 March 2014 Available online 29 March 2014 Keywords: High-velocity oxygen fuel (HVOF) thermal spraying Amorphous Iron-based hard coatings

a b s t r a c t Fe-based alloy material is widely used in the corrosive environment of seawater. It is replacing expensive cobalt and nickel-based alloys. Amorphous iron-based (Fe-based) alloy coatings can be deposited on stainless steel by high-velocity oxygen fuel (HVOF) thermal spraying. Following spraying, coatings were heat-treated at 500, 600, 700, 800, and 900 °C for up to three hours in a vacuum furnace. The microstructures of such coatings were analyzed herein using an optical microscope (OM) and scanning electron microscope (SEM) to monitor the morphologies of both powders and coatings of Fe-based alloy. Phase analysis was performed by X-ray diffraction (XRD) and X-ray photoelectron spectroscopy (XPS). The goal of this work on the modification of Fe-based alloy feedstock powder is to enhance the corrosion and wear properties of these coatings. The results of this investigation reveal that adding a cobalt material to Fe-based alloy yields coatings with enhanced corrosion and tribological characteristics. © 2014 Elsevier B.V. All rights reserved.

1. Introduction In recent years, marked progress has been made in iron-based (Febased) wear and corrosion-resistant, amorphous-metal coatings for use by shipping and petrochemical industries in most island nations [1–15]. Such high-performance Fe-based materials might replace more expensive cobalt or nickel-based alloys in corrosion and wear environments, reducing costs in various industry applications [16]. Although Fe-based amorphous-metal ribbons or rods have been studied widely, few researchers have focused on coating films, and contributed their coating formation using thermal spray [17–22]. An amorphous Fe49.7Cr17.7Mn1.9Mo7.4W1.6B15.2C3.8Si2.4 (SAM2X5) patent alloy deposited by HVOF has been applied on the corrosion environment [8,13]. Chokethawai et al. reported that Fe52.2Cr18.9B16.1C4.0Si2.8Mo2.4Mn1.9W1.7 (SHS7170) coatings consisted of amorphous and nanocrystalline phases are well formed by HVOF thermal spraying [10]. The partially amorphous structure of HVOF coated quaternary alloy (Fe–Cr–Si–B) was examined and correlated highly with its hardness and cavitation erosion resistance [11]. Amorphous steel coating associated with yttrium element was developed and fabricated by HVOF spraying due to its alloy having high glass forming ability (GFA) and corrosion resistance [12]. Fe49.5Cr18.3Mo7.7W1.6Mn1.8B14.9C3.6Si2.6 (SHS7574), as-sprayed coating

⁎ Corresponding author. Tel.: +886 4 22840500x414; fax: +886 4 22857017. E-mail address: [email protected] (F.-S. Shieu).

http://dx.doi.org/10.1016/j.surfcoat.2014.03.041 0257-8972/© 2014 Elsevier B.V. All rights reserved.

deposited by HVOF process within partially amorphous phase also has been mentioned by Zois et al. [15,21]. The SHS9172 alloy, which is the focus of this research, is an Fe-based amorphous alloy and is composed of Cr, Mo, W and Nb as well as inexpensive B, which is added to enhance ability of alloy to form a metal glass [23,24]. Many elements and significant variation of atomic sizes are required to satisfy the important conditions for glass formation, described by Inoue's empirical rules [25,26]. The majority of relevant papers have focused on melting mixtures of constituent elements in-situ and then performing conventional rapid quenching [25–27]. However, this work aims to provide an amorphous coating by using a postmixing procedure that involves an electroless method for modification of commercially available powder, followed by thermal spraying. This approach is a good example of the development of modified amorphous materials. Thermally sprayed coatings on a metal substrate have a lower overall manufacturing cost than ferrous bulk metallic glass (BMG) casting because their constituent layers are fabricated in an atmosphere of air. Thermal spray process can cool substrate at a sufficient rate (104 K/s), which is required for fabricating amorphous materials [1]. The conventional high-velocity oxygen fuel (HVOF) thermal spraying process is one of the best thermal spray processes because its gas and particle velocities are three to four times the speed of sound (Mach 3 to 4). In HVOF, both thermal and kinetic energies for heating and accelerating feedstock powder are provided by combusting fuels with oxygen. Additionally, higher in-flight particle velocities and lower melt or semi-melt

W.-H. Liu et al. / Surface & Coatings Technology 249 (2014) 24–41

temperatures than those of flame and plasma spray processes result in higher coating densities, improved adhesion, and more uniformly dispersed coating structures owing to large particle deformations that occur on impact. This process is also ideal for deposition of metal, alloy, and cermet coatings including Ni–Al, steel, Ti, Mo, WC–Co, and chromium carbide with nickel–chromium coatings and similar. These coatings have bond strengths of approximately 60–80 MPa, porosities of under one percent (b1%), and a micro-hardness of 40–68 HRC. Branagan et al. stated that HVOF thermal spraying of an Fe-based coating on stainless steel can also form an amorphous matrix with hard micro/nanoscale borocarbide precipitates [1]. The resulting advantages of improving hardness, density, and mechanical strength cause HVOF thermal spraying to be utilized extensively in formation of amorphous steel coatings to provide resistance against wear, corrosion, and oxidation [1]. In this work, an experiment on HVOF thermal spraying is carried out in a reducing flame. Protective coatings provide resistance to oxidation and corrosion resistance at high temperatures for machine parts in order to improve their service life and performance, even in severe and especially hightemperature environments. High Cr–W–Nb–Mo steel alloy coating is a high-temperature protective coating that can be easily deposited by thermal spraying [23,24]. This alloy has favorable mechanical properties and oxidation as well as corrosion resistance at high temperatures (N800 °C). Co has a few, but very specialized, uses in many grades of steel that are used in high-speed tools [28]. It is also an important constituent of 18% Ni maraging steels and various other ultra-high-strength steels. It is added to one grade of austenitic stainless steel [29]. Co superalloy is used to make heavy industrial gas turbine engines or other machine parts that must resist high-temperature oxidation, wear, and corrosion [30]. An Fe coating that contains Co and Ni is applied to sprocket wheels as a corrosion-retarding material [31]. The purpose for adding a little Co to the alloy feedstock powder in this work is to improve its coating properties. The application of a modified Fe-based alloy as an amorphous coating with a partially crystallized structure can enhance the tribological, corrosion, and high-temperature oxidation resistance properties of these coatings, which increase life cycle of machine parts on which they are utilized. 2. Materials and methods SHS powder (15–53 μm), produced by NanoSteel company, was used in this study. Then, cobalt was added to the SHS powder by an electroless process that was developed by the Surface Treatment and Finishing Laboratory, Feng Chia University in Taiwan. Thermal spraying was carried out to deposit a coating with a thickness of over 100 μm on a 76.2 × 25.4 × 12.7 mm3 304 stainless steel substrate to produce specimens for metallographic and wear testing. Table 1 presents the feedstock materials. SHS9172 and SHS9172 + Co powders were deposited on the substrates by HVOF spraying. These two composite powders were independently fed into an HVOF spraying torch and deposited as coatings. The HVOF-spray-deposited coatings of SHS9172 and SHS9172 + Co composite material are denoted as SH1-0 and SH2-0, respectively. A DJ 2600 DJM gun (Sulzer Metco) was used with 728 SLPM hydrogen, 213 SLPM oxygen and 379 SLPM air as the spraying gases to produce thermally sprayed coatings. Table 2 presents spraying parameters. The samples were subsequently annealed in a vacuum furnace by heat treatment at 500 °C, 600 °C, 700 °C, 800 °C, and 900 °C for three hours. Samples with coatings are again denoted by a series of code, as

25

Table 2 HVOF spraying parameters. Gun type Hydrogen flow rate Oxygen flow rate Air flow rate Powder carrier gas (N2) flow rate Powder feed speed Spray distance Traverse number Transverse speed

DJ-2600 728 SLPM @140 psi 213 SLPM @170 psi 379 SLPM @100 psi 17 SLPM @150 psi 45 g/min 254 mm 4 pass 12,500 mm/s

shown in Table 3. For convenience, as-sprayed samples are hereinafter denoted by samples with annealing at 0 °C. The metallurgical morphologies of the as-spray coatings were analyzed using an optical microscopy (OM). An epoxy resin used to mount those samples of coatings was treated with glass sand to increase hardness. All samples were ground and polished before using OM (Olymplus BH2-UMA with a MacroFire microscope digital camera extension instrument from 50× to 1000 ×). Then, OM image acquisition was performed. With the digital image processing software (Image Pro Plus software package, Media Cybernetics, Inc.), the porosity of each coating was analyzed. The ASTM E2109-01 method was referred to calculate pore area percentages, that is, total pore area versus total coating area. For counting these areas, a digital photo scope was limited in a window fixed size (2048 × 2048 pixels) at OM magnification of 200× according to the standard method. The result was then regarded as porosity of a coating. Random five fields of view were selected in each polished cross-section coating. The mean porosity values (%) for each coating were then computed. Microhardness of samples was also evaluated by Mitutoyo HVK-H2. Test was under 300 g of load for 10 s. Random indentations were made on each cross-section coating using a Vickers diamond point. A minimum of ten measurements were performed for each coating in order to determine mean microhardness values. LEO 1530 FESEM (field emission scanning electron microscope) and EDS (energy dispersive spectrometry) were used to analyze microstructures of thermal spray coatings. SEI (secondary electron imaging) and BEI (backscattered electron imaging) modes were also utilized in SEM (scanning electron microscope). DSC (Differential scanning calorimetry) SDT Q600 including its own software package (TA Instruments Universal Analysis 2000, Waters LLC.) was used to analyze and determine crystallization and glass transition temperature of powder and coating. Samples were measured from 350 °C to 950 °C with a scan rate of 5 °C/min and 20 °C/min in 100 ml/min nitrogen atmosphere. Bond testing of coatings was conducted using the ASTM C 633 standard method, but samples were bonded to a medium carbon steel alloy with a diameter of 8.16 mm. The test was then conducted on an adhesion tester (Elcometer, PAT GM 01/6.3 kN). Plasmatex Klebbi glue was used as bonding adhesive material between coating and dolly. Bond strengths of coatings were averaged over eight positions. The X-ray diffraction (XRD), Phillips PW3710 Vertical Goniometer, was used to characterize structures of coatings that had been sprayed and heat-treated. However, Phillips PW1710 X-ray instrument was used to characterize structures of powders. The crystallinity (%) of powders and coatings was measured by corresponding XRD peak fitting tool (MDI Jade 5.0). XRD profile fitting and overlapped peak separation with Pseudo-Voigt function are described in Jade 5.0 operation manual. The fit report for

Table 1 Feed materials used in experiment. Feed powder no.

SH1-0 SH2-0

Chemical composition (wt.%)

Remark

Fe

Cr

W

Nb

Mo

B

C

Mn

Si

Co

Bal Bal

b25 b25

b15 b15

b12 b12

b6 b6

b5 b5

b4 b4

b3 b3

b2 b2

– 1–2

SHS9172 commercial powder Electroless Co coated SHS9172 powder

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Table 3 Coating type and heat treatment. Material

Samples

Heat treatment

SHS9172

SH1-1 SH1-2 SH1-3 SH1-4 SH1-5 SH1-6 SH2-1 SH2-2 SH2-3 SH2-4 SH2-5 SH2-6

As spray 500 °C, 3 600 °C, 3 700 °C, 3 800 °C, 3 900 °C, 3 As spray 500 °C, 3 600 °C, 3 700 °C, 3 800 °C, 3 900 °C, 3

SHS9172 + Co

h h h h h h h h h h

amorphous content of powders and coatings is provided in the next Section (to be described later in Table 5). The diffraction patterns of annealing coating samples at 900 °C were also simulated by the Rietveld full pattern refinement method [32]. This is a profile refinement method for nuclear and magnetic structures [32]. Maud software, version 2.33, is then used. According to Maud software, amorphous/disordered phase can be defined by its grain size less than or equal to 10 nm. A very good agreement between the experimented and simulated XRD patterns can be obtained if their weighted residual error (Rwp) and goodness values are given by less than 15% and approximate 1, respectively. To analyze phase chemical composition of coatings, X-ray photoelectron spectroscopy (XPS) was utilized. XPS was carried out using a Thermo VG ESCAlab 250 (ThermoVG Co.) with an Al (Kα) X-ray source. The X-ray photoelectron spectra of samples were calibrated using carbon 1s (284.6 eV). The surface of each coating was cleaned using Ar+ ion sputtering for 300 s. The abrasion properties of sprayed coatings were characterized by dry sand wear testing (Yen Wu Mechanical Co., Ltd.) using the ASTM G65 standard method. The wear load and wear rotation speed during dry sand wear testing were 130 N and 200 rpm, respectively. The diameter of abrasion wheel was 228.6 mm (9.00 in.). The rate of sand flow through nozzles was controlled between 300 and 400 g/min. Three samples of each coating were tested. The duration of test was 30 min. The material loss was measured by weighting samples before and after testing. A theoretical density of 7.59 g/cm3 [24] was used to transform mass into volume to express wear rate in terms of material removal. The corrosion behavior of stainless steel with different coatings was evaluated by potentiodynamic tests at 25 °C in 5% NaCl solution. Tests were made using CHI potentiostat/galvanostat model 440 in a three electrode cell without aerated. Coating with defined sample area of 1.5 cm2 was used as working electrode. A platinum plate served as counter electrode; reference electrode consisted of Ag/AgCl (3 N KCl). A scan rate of 10 mV/s was used in polarization experiment. In this paper, we presented results of metallographic examinations, tensile bond strength tests, and wear tests, only on as-sprayed specimens and those that had been annealed at 500 °C, 700 °C, and 900 °C for three hours. Electrochemical tests were carried out on as-sprayed specimens and those that had been annealed at 900 °C for three hours. 3. Results and discussion 3.1. SEM images of feedstock material The SEM photograph in Fig. 1 reveals that the microstructures of the powder with added Co (SH2-0) and that of powder without added Co (SH1-0). At region A of Fig. 1(b) nano particles (arrows B) are seen and some pores are on the surface of micro powders (SH1-0). The morphology of SH1-0 and SH2-0 powders is spherical and crystalline, as shown in Figs. 1(a) and (m), respectively. The size of powder particles is measured to be in the range of 15–53 μm. According to SEM and EDS

mapping images of polished cross-section of SH2-0 powder, these powder particles are coated with cobalt at outer layers (Figs. 1(d)–(m)). SH2-0 powder, prepared by an electroless process, comprises aggregation particles ranged from hundred nanometers to one micrometer in diameter (Fig. 1(c)). The thickness of cobalt film is in the range of 1–2 μm (arrows C in Fig. 1(m)). The cobalt content of SH2-0 powder is determined by EDS to be 1–2 wt.%, and confirmed by subsequent XPS analysis. The feature of SH2-0 powder which is similar to core–shell microstructure brings about the corresponding coating characteristics, and will be mentioned later. 3.2. DSC analysis Table 4 lists the peak crystallization temperatures, Tx, of SHS9172 and SHS9172 + Co powders and coatings, as estimated from DSC curves in Fig. 2. Notably, SH2-0 and SH2-1 have a lower Tx than SH1-0 or SH1-1. The glass transition temperature, Tg, can be determined only for assprayed SH1-1 and SH2-1 specimens. The Tg of SH1-1 is slightly lower than that of SH2-1. Adding Co to SHS9172 alloy coating reduces its crystallization temperature, but increases glass transition temperature. It also reduces enthalpy of crystallization of SHS9172 + Co alloy coating. Moreover, a partial explanation for this finding may lie in the following fact. While alloy is with higher Tg temperature, the less difference between Tg and Tx is, the more amorphous phase tendency during the cooling process of manufacture is (according to reduced glass transition temperature, Trg is) [33,34]. Namely, SH2-1 is more thermally stable against crystallization than SH1-1. Because SH2-1 has higher Tg temperature and lower temperature difference between Tg and Tx. SH2-1 coating yields a slightly lower onset crystallization temperature (Tx = ~ 627 °C), comparing with onset temperature of Fe48Cr15Mo14C15B6Y2 amorphous coating (Tx = ~637 °C) [12,35]. The other possible reason for Tg difference will be discussed in the following Subsections 3.6 and 3.7. Adding Co to an Fe-based alloy coating can improve its glassforming ability (GFA), as also reported elsewhere [36,37]. 3.3. Metallographic analysis of coating The SHS9172 (from samples SH1-1 to SH1-6) and SHS9172 + Co (from samples SH2-1 to SH2-6) coating were prepared from spherical powders of high Cr–W–Nb–Mo steel alloy with and without added Co, respectively. Fig. 3 shows that the thickness of SH2-1 coating is thicker than that of SH1-1, under the same operation parameters for HVOF process. OM observations reveal that deposition rates of coating with added Co exceed those of coating without added Co (Fig. 3). Of cross-sections of hard coatings for HVOF spraying, Fig. 3 also presents the typical regions. For instance, it shows that there are some vertical cracks within microstructure of SH1-1 coating (Fig. 1(a)), whereas none of them takes place in SH2-1 coating (Fig. 1(b)). Same results will be described in Figs. 5 and 6. The main cracks are typically perpendicular to substrate, while shorter horizontal cracks branch out from these perpendicular vertical cracks in Fig. 3(a) (SH1-1). Moreover, Fig. 3 shows that individual splat is seldom rounded as opposed to splat-shaped particles in as sprayed coatings. This microstructure is owing to partially or fully melted particles while it was transported in flame stream. Effective mechanical anchorage of coatings onto substrate irregularities is also observed (Figs. 3 and 4). Figs. 4–8 illustrate SEM and EDS mapping images of microstructural features of SHS9172 and SHS9172 + Co coating. To describe the positive effect of adding Co in SHS9172 + Co, Fig. 4(c) illustrates that a little Co comes with Fe and distributes almost along Fe. And, Co in SH2-6 coating diffuses through coating/substrate interface into AISI 304 stainless steel substrate. The finding, the distribution of Co over coating/substrate interface, is expected to improve bonding strength with its substrate. Comparing with upper right zones of each image in, it is also clearly found that some Co are located around splat particles. However, other Co are located at inner splat particle, as illustrated in other zones of

W.-H. Liu et al. / Surface & Coatings Technology 249 (2014) 24–41

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Fig. 1. SEM and EDS images of powders: (a) SEI SH1-0, (b) higher magnification image of (a) corresponding to region A, (c) SEI SH2-0, (d)-(l) EDS mapping images of SH2-0: (d) Fe; (e) Cr; (f) W; (g) Nb; (h) Mo; (i) C; (j) Mn; (k) Si; (l) Co element, (m) SEM of SH2-0. Arrows B indicate nano particles grew on the surface. Arrows C indicate thickness of the cobalt film in the range of 1–2 μm.

the same figure (Fig. 4), for example the zone near sand blasting grit. Thus, the existence of cobalt for SHS9172 + Co coatings is also expected to improve cohesion between splat particles. According from above observations on SHS9172 + Co coatings (Fig. 4), the results imply that a small amount of Co dissolved in Fe-rich matrix and present in/beside splat particles can act as a binder for SHS9172 + Co coatings effectively. Co in coatings is closely related to the similar core–shell microstructure of SHS9172 + Co feedstock powder. The results will be further discussed in Subsection 3.7. The following will describe the defect features of SEM of some representative samples. Some pores and microcracks appear as dark regions in SEI images (Figs. 5 and 6). The pore size for SHS9172 and SHS9172 + Co coatings relatively increases with increasing of heat treatment temperature until up to about 700 °C for 3 h, then decreases with further

increasing in annealing temperature up to 900 °C for 3 h (Figs. 5(a), (c), (e), and (g); and 6(a), (c), (e), and (g)). Some researchers reported that large pores (e.g., see in Figs. 5(e), (f), 6(c) and (d)) between flattened droplets mostly come from loose packing of layered structure or are formed by atmospheric gas, whereas small pores (e.g., Figs. 5(a) and 6(a)) within flattened particles are formed by shrinkage of these particles [38]. To highlight differences between the two coating SHS9172 and SHS9172 + Co, quantitative data analyses of OM cross-section microstructure was performed (Fig. 9). Despite the presence of defects (e.g., cracks), the coatings have a low porosity about 0.3–3.5% which is the typical characteristic of HVOF-sprayed coatings. For SHS9172, similar to pore size, average area percentage of pores (that is, porosity) relatively increases with increasing of heat treatment temperature until up

Table 4 DSC experiments. Samples no.

SH1-0 SH2-0 SH1-1 SH2-1 SH1-0 SH2-0 SH1-1 SH2-1

Glass phase transition

Crystallization phase transition

Heat rate (°C/min)

Glass transition temperature, Tg (°C)a

Specific heat Capacity at glass transition, ΔCp (J/(g · °C)

Peak crystallization temperature, Tx (°C)

Enthalpy at crystallization, ΔHx (J/g)

5

– – 587.4 ± 1.6 590.7 ± 1.6 – – – –

– – 0.9 ± 0.0 1.4 ± 0.2 – – – –

647.8 644.0 627.8 627.6 667.4 662.7 648.3 637.2

6.9 6.1 25.4 16.9 7.3 5.0 49.7 31.5

20

Value given as mean ± SD within three repetitional times. a Onset temperature.

± ± ± ± ± ± ± ±

0.9 0.1 0.5 0.2 0.1 0.1 0.1 0.3

± ± ± ± ± ± ± ±

1.4 0.4 0.3 0.5 0.6 0.7 0.4 2.4

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Fig. 2. DSC curves of SHS9172 and SHS9172 + Co powders and coatings.

to about 700 °C for 3 h, then decreases with further increasing in annealing temperature up to 900 °C for 3 h (Fig. 9). By comparison, the maximum average value of porosity (2.7%) for SHS9172 + Co is obtained after heat treatment at low temperature, that is, about 500 °C, and then porosity also decreases with further increasing in annealing temperature up to 900 °C for 3 h (Fig. 9). In other words, maximum and minimum average values of porosity for SHS9172 are 2.4% at 700 °C and 0.8% at 900 °C respectively, whereas corresponding values for SHS9172 + Co are 2.7% at 500 °C and 0.4% at 900 °C, respectively. It is also found that the microcrack defects in SHS9172 coatings are clearer than those in SHS9172 + Co coatings (Figs. 5 and 6). Specially, vertical microcracks only exist in SHS9172 coatings (e.g., Fig. 5(c) and (e)). It seems to be able to develop the defect more easily for SHS9172 than SHS9172 + Co, this finding implies that deposition of SHS9172 onto AISI 304 substrate may be on fully melting condition, overheating state during HVOF process. Therefore, onto the relatively cold substrate, the corresponding coatings produce microcrack defects in order to balance off by raising heat accumulation. It is also found that the number of coarse partial melted/unmelted particles in SH2-1 is greater than that in SH1-1 (Figs. 3(b) and 6(a)). However, the SH2-6 is denser than that of SH1-6 (Figs. 5(g) and 6(g)). This observed evidence suggests that SHS9172 + Co feedstock powder does not have enough time to be fully melted by means of HVOF spraying. This phenomenon should be related to the similar core–shell microstructure of feedstock powder itself. The generalized results may also help to explain some interesting differences highlighted in DSC, SEM and OM analyses as mentioned previously. According to the report, self-fluxing alloys (such as FeCrBSiCNi or SHS9172 alloy system) contain Si and B, which act as deoxidizers and have been used for vitrification, and HVOF or flame spraying can be used to deposit such alloys, causing some self-fluxing reaction during the process, itself, or subsequent heat treatment [8,39–42]. Thus, the above-mentioned evidence also suggests that this self-fluxing reaction has stronger effect on densification and make amorphous/disordered structure of SHS9172 + Co coatings stable during subsequent heat treatment. That is, self-fluxing reaction results in relatively high glass forming temperature for SHS9172 + Co (Tablet 4). The self-fluxing reaction also may result in relatively low porosity for SH2-4 and SH2-6, as shown in Fig. 9. Wu et al. [11] stated that the partial melted/unmelted particles involve in semi- or un-molten borides, if in the melting states during the HVOF process, it could be easy to transfer to the formation of amorphous phase and then produce solid solution strengthening in Fe matrix (e.g., α-Fe) solution. This is the kind of self-fluxing reaction for the Fe based alloy (47.52 wt.% Fe-44.7 wt.% Cr-5.8 wt.% B-1.98 wt.% Si) existing high B during the HVOF and/or its following high temperature

Fig. 3. OM cross-section images of (a) SH1-1 and (b) SH2-1 for SHS9172 as-sprayed coating and SHS9172 + Co as-sprayed coating, respectively.

treatment process. The self-fluxing reaction may be accompanied with crystal grain nucleation and growth during annealing in the alloy. Accordingly, SHS9172 (this article used), SHS8000 [14], SHS7574 [15], and SHS7170 [7,10,21] alloys may also have this kind of behavior. Branagan et al. [13] found that the decreasing crystallinity of Fe–Cr– Mn–Mo–W–B–C–Si as-sprayed coating indeed results from un-melted particles in which un-melted particle boride phase is present. However, the effect of this reaction for the partial melted/unmelted particles in SH2-2 is not helpful for decreasing its porosity due to heat treatment temperature below its own Tg temperature. It means that there is not enough driving force to re-start self-fluxing reaction at such low annealing temperature. Another reason is variation of porosity related to coatings. This variation may result from phase growth and phase transformation, is further discussed in the following Subsection 3.6. The following will describe oxide features of SEM with some of representative samples. Figs. 5 and 6 show oxide discontinuous layers of

W.-H. Liu et al. / Surface & Coatings Technology 249 (2014) 24–41

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Fig. 4. SEM cross-section images of SH2-6 located at interface between coating and AISI 304 stainless steel substrate: (a) SEI image, (b) composite element map of C, Fe, Mn, Cr, Mo, Nb, W, Si, B and Co elements, (c) composite element map of Fe and Co elements.

Fig. 5. SEM cross-section images of SHS9172 coatings with various heat treatments: (a) SEI of SH1-1, (b) BEI of SH1-1, (c) SEI of SH1-2, (d) BEI of SH1-2, (e) SEI of SH1-4, (f) BEI of SH1-4, (g) SEI of SH1-6, and (h) BEI of SH1-6.

Fig. 6. SEM cross-section images of SHS9172 + Co coatings with various heat treatments: (a) SEI of SH2-1,(b) BEI of SH2-1, (c) SEI of SH2-2, (d) BEI of SH2-2, (e) SEI of SH2-4, (f) BEI of SH2-4, (g) SEI of SH2-6, and (h) BEI of SH2-6.

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Fig. 8. SEM cross-section and EDS mapping images of SHS9172 and SHS9172 + Co samples at the interface between coating and its outside surface near epoxy resin zone: (a) SEI of SH1-6, (b) composite element map of Fe, C, Mn, Cr, Mo, Nb, W, Si, B and O elements of (a), (c) composite element map of Fe, Mn, Cr and Nb elements of (a), (d) composite element map of Fe, Mn, Cr, Nb and Si elements of (a), (e) SEI of SH2-6, (f) composite element map of Fe, C, Mn, Cr, Mo, Nb, W, Si, B, Co and O elements of (e), (g) composite element map of Fe, Mn, Cr and Nb elements of (e), (h) composite element map of Fe, Mn, Cr, Nb, Si and Co elements of (e).

Fig. 7. SEM cross-section and EDS mapping images of SHS9172 and SHS9172 + Co at the region of coating: (a) SEI of SH1-6, (b) composite element map of Fe, C, Mn, Cr, Mo, Nb, W, Si and B elements of (a), (c) composite element map of Fe, C and B elements of (a), (d) composite element map of Fe, C, Cr and B elements of (a), (e) SEI of SH2-6, (f) composite element map of Fe, C, Mn, Cr, Mo, Nb, W, Si, B and Co elements of (e), (g) composite element map of Fe, C and B elements of (e), (h) composite element map of Fe, C, Cr and B elements of (e), (i) composite element map of Fe and Co elements of (e).

remain in a vacuum furnace because the annealing was carried out under ~10−2 Torr and oxides of Fe and Cr are slightly increased with increase of heating temperature. These oxides can be further identified in XPS analysis as shown in the following Subsection 3.7.

4.0 SH1-1 SH2-1 SH1-2 SH2-2

SH1-4 SH2-4

SH1-6 SH2-6

700

900

3.5 3.0

Porosity (%)

SHS9172 and SHS9172 + Co at 900 °C clearly. Fig. 8 shows that structure of oxides at outer surface of SH1-6 coating is looser than that of SH2-6 coating. SH1-6 oxides, detected by X-ray mapping, are involved in Fe, Cr, Nb, Mn and Si composite oxides along with trace elements B, W, and Mo. However, SH2-6 oxides still contain trace element Co except for the former elements. Main oxide is Cr oxide for SH1-6 (Fig. 8(b)–(d)). Moreover, the concentration of chromium oxide at region of composite oxide film in SH2-6 (Fig. 8(f)–(h)) is less than that in SH1-6 (Fig. 8(b)–(d)). Besides, the concentration of Si in oxide film of SH1-6 is less than that in oxide film of SH2-6 (comparison between Fig. 8(d) and (h)). The corresponding results imply that SHS9172 oxide film having primary Cr oxide transforms to SHS9172 + Co oxide film having Fe, Cr and Si whose concentrations are high due to added Co. During HVOF process in a reducing flame, a few Fe and Cr oxides also probably remain in SHS9172 and SHS9172 + Co specimens. Oxides of Fe and Cr will still

2.5 2.0 1.5 1.0 0.5 0.0

as spray

500

Annealing Temperature (oC) Fig. 9. Bar plot of porosity of as-sprayed coating and coatings annealed at 500 °C, 700 °C and 900 °C, respectively.

W.-H. Liu et al. / Surface & Coatings Technology 249 (2014) 24–41

For precipitation and grain features of SEM of the representative samples, significant lighter and well-proportioned precipitation phase in coatings increases with an increase in annealing temperature (Fig. 7(a) and (e)). Composite element mapping analysis identifies increased W and Mo elemental contents in these lighter regions (Fig. 7(b) and (f)). No significant difference exists among morphology of precipitation particles in SH1-6 (Fig. 7(a)) and SH2-6 (Fig. 7(e)). The size of precipitation particles is ranged from 80 to 200 nm. The number of precipitation particles is increased with increasing heat treatment temperature for either SH1-6 (Fig. 5(b), (d), (f) and (h)) or SH2-6 (Fig. 6(b), (d), (f) and (h)). However, Cr elemental content in SH1-6 is more than that in SH2-6 (Fig. 7(d) and (h)). And, the finding also shows that Co element in SH2-6 prefers staying at the regions of Fe matrix to that of white precipitation particles. On the other hand, SEM micrographs for SH1-6 and SH2-6 again indicate that for two regions in Fig. 7(a) and (e), marked regions A and B respectively, fine grain size below 100 nm is seen.

3.4. Microhardness testing of coating The hardness test shows that the hardness firstly increases greatly and then decreases as annealing temperature is increased (Fig. 10). The microhardness of coatings increases with annealing temperature up to 700 °C when they are annealed for 3 h, but decreases with a further increase in annealing temperature up to 900 °C, at which annealing temperature, the value of microhardness is close to that of as-spraycoated coatings. In other words, the maximum average values of hardness for SHS9172 and SHS9172 + Co specimens are: Hv300 = 1216 and 1125 at 700 °C for 3 h, respectively. However, the minimum average values of hardness are: Hv300 = 745 at 900 °C (SH1-6) as compared with SHS9172 and SHS9172 + Co. For improving hardness or mechanic properties, this result reveals that hardness of SHS9172 or SHS9172 + Co coating is just high enough at operating temperature up to about 700 °C. Furthermore, the analyzed data indicates surprising trends that microhardness of coating with added Co is slightly higher than that of any coatings without added Co and the extent of difference increases with annealing temperature, except for annealed at 700 °C. A previous study found that adding Co improves microhardness of Fe48Cr15Mo14C15B6Y2 amorphous bulk alloy [43]. The difference between microhardness values of SH1-4 and SH2-4 may result from growth of alloy phase with Co, which is further discussed in the following Subsections 3.6 and 3.7. Besides, the mean hardness value of SH1-1 (Hv300 = 850) is higher than that of SH1-6 (Hv300 = 745); however, the corresponding value of SH2-1 (Hv300 = 853), as well as SH1-1, is approximately equal to that of SH2-6 (Hv300 = 850).

1400

3.5. Bonding strength of coating The SH1-2 and SH2-2 have the maximum average values of tensile bonding strength (91 and 96 MPa), whereas SH1-1 and SH2-1 have the minimum average values of their strength (59 and 73 MPa), respectively, as shown in Fig. 11. Moreover, the tensile bond strength of SH1-1 and SH2-1 coatings increases greatly with annealing temperature up to 500 °C but then slightly drops down as annealing temperature is increased (Fig. 11). However, the coating samples with annealing at any temperature up to 900 °C yield higher tensile bond strength than those for as-spray coating samples. Comparing with porosity and hardness values (Figs. 9 and 10), a similar tendency toward the increasing– decreasing curves can be found as annealing temperature is increased. In this study, adding Co is also observed to increase coating strength, except for SH2-4. And both SHS9172 and SH9172 + Co alloys with annealing is observed to enhance tensile bonding strength. Tensile bonding strength of thermal spray coating is distinguished into two different strength modes: cohesive (intrasplats and intersplats in the coating) and adhesive strength (between the coating and its substrate) [44–47]. Furthermore, the cohesive (intrasplats and intersplats in the coating) strength of the coating depends on a base metal matrix, and precipitates besides pores and cracks in its own coating structure. Since the vertical cracks are caused by residual stresses in the coating (Fig. 3), bond strength for SH2-1 is higher than that of SH1-1 due to absence of vertical cracks. And, diffusion phenomenon with element species is also helpful for increasing tensile bonding strength with increased annealing temperature. The diffusion phenomenon between coating/substrate interface and intermetallic compound through interaction with base metal atoms (substrate), is probably responsible for the stronger bond, e.g. metallurgical bond [44–47]. By comparison the annealing samples with their own as-sprayed samples, it can be explained that remained residual stresses are further released resulting from mismatch between thermal expansion coefficient of coating and substrate. Moreover, it seems to be explained as there are some precipitates, for instance the nanocrystal dispersoids (Figs. 5–7), in the amorphous/disordered alloy coatings. Those precipitates mainly improve their tensile strength with heat treatment, as compared with assprayed samples. This result for amorphous/disordered with nanocrystalline is consistent with those of Inoue et al. [48] and Frankwicz et al. [49]. Besides, for the case with annealing samples at 700–900 °C, tensile strength is lower than that for the annealing samples at 500 °C, it is possible to provide valid reason that a few oxide precipitates (Figs. 5(g) and 6(g)) in coating or between coating and substrate also have negative influence on bonding strength. By comparing SH1-1 and SH2-1 with SH14 and SH2-4, an exothermic crystallization reaction can improve mechanical properties (strength) of coating, as reported earlier [50].

120 SH1-1 SH2-1 SH1-2 SH2-2

SH1-4 SH2-4

31

SH1-6 SH2-6

SH1-1 SH2-1

SH1-2 SH2-2

SH1-4 SH2-4

SH1-6 SH2-6

1200

Bonding Strength (Mpa)

Micro Vickers (Hv300)

pure glue via blast substrate

1000 800 600 400 200 0

as spray

500

700

Annealing Temperature

900

(oC)

Fig. 10. Bar plot of microhardness of as-sprayed coating and coatings annealed at 500 °C, 700 °C and 900 °C, respectively.

100 80 60 40 20 0

as spray

Annealing Temperature (oC) Fig. 11. Bar plot of bonding strength of as-sprayed coating and coatings annealed at 500 °C, 700 °C and 900 °C, respectively.

32

W.-H. Liu et al. / Surface & Coatings Technology 249 (2014) 24–41

Wear data provide further evidence of predicted tendency for cohesive strength as shown in Subsections 3.6–3.8. 3.6. Phase analysis using X-ray diffraction The XRD results are presented in Fig. 12 and summarized in Tables 5 and 6. The analyses are shown in the following: (i) The as-sprayed and annealed specimens contain no less than six solid solutions (α-Fe) as the major probable phases. They are: JCPDS 06-0696 Fe (2θ = 44.67°, 65.02°, 82.33° and 98.95°), 65-9130 Fe0.90Si0.1 (2θ = 44.78°, 65.19°, 82.56° and 99.24°), 65-6323 Fe0.90Si0.095 (2θ = 44.77°, 65.185°, 82.56° and 99.24°), 65-4664 CrFe4 (2θ = 44.67°, 65.02°, 82.33° and 98.94°), and 65-7775 Fe0.905Si0.1 (2θ = 44.63°, 64.97°, 82.24° and 98.82°) for both SHS9172 and SHS9172 + Co with as-sprayed and annealed at temperature up to 900 °C, whereas 41-1466 (Fe,Cr) (2θ = 31.14°, 44.83°, 55.44°, 65.19°, 74.00°, 82.56° and 99.40°) only for SHS9172 + Co with annealed at 800 °C and 900 °C. Because of XRD patterns with complex multi-phases, the α-Fe planes (110), (200), and (211) are selected to estimate the size of its mean grain by using Sherrer equation within build-in Jade 5.0 package (Table 5). The α-Fe matrix grain with nano size crystallites for both SHS9172 and SHS9172 + Co coatings shows that the utmost 900 °C for 3 h heat treatment may be not able to grow enough into microcrystalline level. Multi-components (at least 9 elements) and over twenty probable phases are likely to be the two primary causes of forming α-Fe nanocrystalline in Fe-based alloys. On the other hand, the HVOF process provides rapid quenching, and multi-components provide conditions for active mixing with a negative enthalpy of mixing. The production of nanocrystals is thought to have been caused mainly by heat that accumulated within the as-sprayed coatings. To produce coatings of a certain thickness, spray gun must travel over deposited coatings. This operation causes localized reheating directly, owing to both gas jet and evolution of latent heat as successive layers of melted splats solidify. Therefore, it is responsible for formation of nanocrystalline grains in amorphous/disordered matrix. This result provides a detailed account for α-Fe matrix grain with Co added into SHS9172. Namely, the size of α-Fe nanocrystalline for SHS9172 + Co is always less than that for SHS9172. Notably, the grains in the sample that is annealed even at 900 °C are not larger than those in the feedstock powder (SH1-0 and SH2-0). Although the feedstock powders of both SHS9172 and SHS9172 + Co are polycrystalline (not shown here), the as-sprayed coatings of these alloys are nearly amorphous/disordered. This fact is related to HVOF process and various components of metal glass alloy. The XRD patterns of both SHS9172 and SHS9172 + Co coatings have a partially amorphous/disordered structure, since sharp peaks associated with crystallization are observed on the broad halo peak. Fig. 12 reveals that the coatings typically are composed of amorphous/ disordered phases. For comparing crystalline α-Fe phase content in its amorphous/disorder, we set a series of peaks in selected parameter range at 2θ = 44.0°–44.7°. And their corresponding fullwidth-half-maximum (FWHM) values of XRD profiles, which are more than or equal to 3°, are used as a fitting condition with main amorphous/disordered peak to fit all XRD patterns at 2θ = 20°– 100° in both SHS9172 and SHS9172. Then, the crystallinity (%) can be estimated after fitting the XRD patterns. The finding is justified by a fact that samples of SHS9172 + Co are always with higher amorphous/disordered content than samples of SHS9172 (Table 5). At the utmost 900 °C for 3 h heat treatment, the amorphous/disordered content of near 12% for SHS9172 + Co is present, whereas none of that for SHS9172. The trend of crystallinity shows that the residual content of amorphous/

Fig. 12. X-ray diffraction patterns of (a) SHS9172 and (b) SHS9172 + Co. (a) SH1-1, SH1-2, SH1-3, SH1-4, SH1-5, SH1-6 (b) SH2-1, SH2-2, SH2-3, SH2-4, SH2-5, SH2-6 as-spray coating and coatings annealed at 500 °C, 600 °C, 700 °C, 800 °C, and 900 °C, respectively. The patterns of SH1-6 and SH2-6 include both experiment and calculated data, respectively.

W.-H. Liu et al. / Surface & Coatings Technology 249 (2014) 24–41 Table 5 Grain size and crystallinity (%) for XRD experiments. Samples no.

Grain size (nm)

Crystallinity(%)

SH1-0 SH1-1 SH1-2 SH1-3 SH1-4 SH1-5 SH1-6 SH2-0 SH2-1 SH2-2 SH2-3 SH2-4 SH2-5 SH2-6

43.0 10.6 19.7 18.4 25.2 24.6 26.8 23.2 b3 4.1 14.1 12.2 13.9 11.6

80.72 29.56 36.07 64.39 94.52 97.68 100.00 94.89 24.73 30.88 60.70 75.27 82.22 88.93

(ii)

(iii)

(iv)

(v)

disorder for SHS9172 + Co at 900 °C annealing may result from self refluxing with unmelted particles. The traces of intermetallic phases exist probably in as-sprayed and/or annealing coatings, including of borides (JCPDS 34-1198 Fe3B, 51-1410 CrFeB, 65-5159 Mn2B, 47-1629 B2Fe15Si3, 351340 BFe2), carbides (JCPDS 20-0508 Fe5C2, 72-1988 Fe6W6C, 23-1127 Fe6W6C, 05-0721 (Cr,Fe,W,Mo)23Fe21(W,Mo)2C12, 340001 Fe3C, 73-1341 Fe2.7Mn0.3C, 78-1500 Cr21.34Fe1.66C6, 781499 Cr22.23Fe0.77C6, 35-0783 Cr23C6, 14-0407 Cr23C6, 65-3132 Cr23C6, 65-8781 NbC, 65-8789 Nb4WC5) and of other compounds (JCPDS 65-8992 Cr6Fe18Mo5, 48-1816 Co3Fe7). In SHS9172, Cr23C6 phases (JCPDS 65-3132 in SH1-1) start growing at annealing and the content gradually increase as a function of increased heat treatment temperature. At 700 °C or above, SH1-4 has three kinds of Cr23C6 phases (JCPDS 35-0783, 140407 and 65-3132 Cr23C6). However, in contrast, SH2-1, SH2-2 and SH2-3 have none of Cr23C6 phases. Only SH2-4, SH2-5 and SH2-6 have these three Cr23C6 phases, with the distinct annealing temperature at the same 700 °C or above. The existence of Fe3B phase starts to grow for both SHS9172 and SH9172 + Co at 600 °C or above. The Mn2B phase is only presented for both SHS9172 and SH9172 + Co until 900 °C or above. The Cr6Fe18Mo5 phase exists for both SH9172 and SH9172 + Co at 700 °C or above, however, the Co3Fe7 phase only for SH9172 + Co at 600 °C or above.

Table 6 Composition of possible phase of samples at 900 °C for XRD experiments. Data calculated using Rietveld full pattern refinement method. Possible phase

Wt.% (SH1-6)

Wt.% (SH2-6)

48-1816 Co3Fe7 65-8781 NbC 35-1340 BFe2 65-5159 Mn2B 65-3132 Cr23C6 14-0407 Cr23C6 35-0783 Cr23C6 78-1499 Cr22.23Fe0.77C6 78-1500 Cr21.34Fe1.66C6 65-8992 Cr6Fe18Mo5 51-1410 CrFeB 73-1341 Fe2.7Mn0.3C 34-1198 Fe3B 65-7775 Cr1.07Fe18.93 65-4664 CrFe4 23-1127 Fe6W6C 20-0508 Fe5C2 65-6323 Fe0.905Si0.095 41-1466 (Fe,Cr) 65-3190 Fe0.9Si0.1 06-0696 Fe

– 1.12 5.10 2.26 2.71 2.45 2.35 6.27 0.75 4.84 2.34 0.54 1.06 0.28 62.56 0.43 0.07 2.14 – 0.65 2.08

0.02 1.82 5.94 2.94 1.05 1.08 1.14 5.66 0.59 0.02 2.77 1.09 1.40 11.38 53.17 0.57 5.95 0.13 2.87 0.40 0.01

33

(vi) The existence of (Cr,Fe,W,Mo)23Fe21(W,Mo)2C12, Cr21.34Fe1.66C6 and Cr22.23Fe0.77C6 phases starts to grow for SHS9172 at 600 °C or above, whereas those phases for SHS9172 + Co start to grow up at 700 °C or above. The reason is that the shielding effect of Co results from the similar core–shell structure of SHS9172 + Co feedstock powder can have an impact on the phase formation rate of the corresponding coatings (Fig. 1). The existence of borides B2Fe15Si3 and BFe2, and carbides Fe5C2, Fe6W6C, Fe3C, NbC and Nb4WC5 is presented at 800 °C, but only B2Fe15Si3, Fe3C and Nb4WC5 phases disappear at 900 °C for SHS9172. By comparison, the same precipitates of SHS9172 + Co are presented, only the existence of borides B2Fe15Si3 at earlier temperature of 700 °C is seen, and then the same disappearance at such 900 °C is seen, too. The differences (e.g., in SHS9172 + Co: Cr23C6 phases are absent at 600 °C or below but Co3Fe7 starts to grow at 600 °C; (Cr,Fe,W,Mo)23Fe21(W,Mo)2C12, Cr21.34Fe1.66C6, Cr22.23Fe0.77C6 and B2Fe15Si3 phases grow at 700 °C; (Fe,Cr) phase is produced at 800 °C) result from Co added to SHS9172 alloys and may lead to formation of coarse unmelted particles in coatings (Figs. 3 and 6) to produce diffusion and self-flux reaction during heat treatment. The primary reason of the difference may result from the similar core–shell structure of SHS9172 + Co feedstock powder. It is noted that B2Fe15Si3 phase in both cases is stable at higher temperature (700–900 °C, completely disappeared at 900 °C) than that of Mini´c et al. study (527–650 °C, completely disappeared at about 650 °C) [51]. B2Fe15Si3 phase is denoted by one type of Fe3B with stabilized Si [15,21]. (vii) Due to Co3Fe7 phase at such low temperature, that is, 600 °C (cf. Fig. 2), it suggests that SHS9712 + Co may be alloyed with a slight amount of Co element during HVOF spraying (cf. Fig. 7). The melting temperature of Co3Fe7 phase (about 1493 °C [52]) is lower than that of α-Fe phase (about 1538 °C); thus the result implies that melting temperature of SHS9172 + Co coating is lower than that of SHS9172 coating. Namely, this is one of more amorphous phase tendency condition for alloy having lower liquidus temperature Tl or melting temperature Tm according to reduced glass transition temperature Trg [33,34]. Combined with conclusion of Subsection 3.2 (the higher Tg), it indicates that SHS9172 + Co coating has higher GFA in comparison with SHS9172 coating. Adding Co to SHS9172 coating can further produce a disordered atomic configuration, suppressing crystallization with the multiphase and promoting GFA, similar to reference [36]. Moreover, based on XRD analyses, intensity of α-Fe (2θ = 44.5–44.7°) peaks dramatically increases at temperatures ranged from 600 °C to 700 °C (Fig. 12), and those amorphous/disordered phases yield crystallization peaks at 620–650 °C; thus these crystallization temperatures of SHS9172 and SHS9172 + Co coatings indicate that they are the crystal transformation temperatures of amorphous/ disorder into α-Fe primary phase (Fig. 2 and Table 4). (viii) The increase of porosity with increasing annealing temperature (Fig. 9) is probably related to crystalline growth, phase transformation and precipitation. However, porosity for SHS9172 decreases until above 700 °C, whereas that for SHS9172 + Co decreases at the early annealing temperature, namely at 700 °C; that is to say, a wide temperature range of those decreasing data is correlated to the precipitation of B2Fe15Si3 besides Co3Fe7 phase. The result implies that B2Fe15Si3 metastable phase precipitate in coatings may result from the self-flux reaction. Moreover, self-flux reaction containing more B and Si concentration can reduce its porosity in coating [39]. The boron–iron–silicon phase for SHS9172 + Co coatings is presented at such low 700 °C, it is likely that Tg temperature is increased by addition of Co (Fig. 2). The Co3Fe7 and (Fe,Cr) phases grown at 900 °C may decrease the porosity of SHS9172 + Co coatings further.

34

W.-H. Liu et al. / Surface & Coatings Technology 249 (2014) 24–41

(ix) Generally, among these precipitates, borides are the most hardness phases, the next are carbides, and intermetallic compounds without metalloids are rather tough. The relative hardness of corresponding borides and carbides can be found elsewhere [53–66]. Due to solid solution strengthening (by dissolution of B, Si, Cr, etc. in α-Fe) and grain boundary strengthening (by α-Fe nanosized grains), partially amorphous/disordered structure with α-Fe nanocrystalline may mainly have an effect on increase of microhardness for both SHS9172 and SHS9172 + Co coatings when they are annealed at 0–600 °C temperature for 3 h. The CrFeB is with high hardness phase, belonging to M2B-type boride (M stands for metal) [63]. Moreover, the evidence with 700 °C annealing condition suggests that precipitate of carbides (i.e., (Cr,Fe,W,Mo)23 Fe 21 (W,Mo)2 C 12 and 14-0407 Cr23C 6 ) and growth of carbides (i.e., Cr21.34Fe1.66C6, Cr22.23Fe0.77C6, and 35-0783 and 65-3132 Cr23C6), and growth of Fe3B as well as CrFeB borides are responsible for the highest microhardness of both coatings. The microhardness of SH2-4 is relatively low in comparison to SH1-4 (Fig. 10), the finding is probably related to the B2Fe15Si3 metastable phase precipitate and the growth of Co3Fe7. Although at 900 °C annealing process, some precipitates (e.g., carbides and borides) can further help to improve the microhardness of Febased coatings, on the other hand, primary intensity ferrite and trace intermetallic compound without metalloids seem not to be clearly pronounced. The existence of Cr6Fe18Mo5 may be, for example, responsible for decreasing microhardness with both coatings at 900 °C annealing. In fact, the hardness of carbides decreases with increasing heat treatment temperature, especially at such high temperature 900 °C, as shown in other references [53,55]. The reason for higher hardness of SH2-6 as compared with SH1-6 will be discussed in next Subsection, even though the (Fe,Cr) solid solution and Co3Fe7 intermetallic compound without metalloids are only presented in SH2-6 coating. The previous research indicates three beneficial influences on the microhardness as follows: grain boundary strengthening (by α-Fe nanosized grains), dispersion strengthening (by borides) and solid solution strengthening (by dissolution of B, Si, Cr, etc. in α-Fe) have strong effect on the highest microhardness of a as-sprayed Fe based coating [11]. Furthermore, during annealing process (b 800 °C), some precipitate (e.g., carbides, borides) can help to improve the microhardness of Fe-based coatings [15]. These issues of the microhardness are highlighted in our research. (x) Due to overlapping effect of Co layer on X-ray light, grain size of SHS9172 + Co powder is smaller than its certain size as well as that of SHS9172 powder. Adding Co reduces both grain size and crystallinity of SHS9172 coating at various annealing temperatures. (xi) Fe23B6 and Fe23(C,B)6 is hardly presented due to weaker intensity in XRD spectra. Fe23(C,B)6 and Fe3B phases are less stable than Fe2B and FeB phases [67,68]. Fortunately, borides of Nb, Mo, and W cannot be found after annealing. Silicides of Mo and W are also not detected by XRD. The tendency of these phases to precipitate out depends on the degree of segregation of solid solution in Fe-based alloy. (xii) The phases of oxides are difficult to identify in XRD analysis due to the low content in both coating types and the formation of borides and carbides which are used to against high-temperature oxidation of coatings. (xiii) Composition of complex multi-phase for SH1-6 and SH2-6 is calculated by Maud software according to Rietveld analysis of XRD patterns (Table 6). For simplified calculation, the 72-1988 Fe6W6C, 05-0721(Cr,Fe,W,Mo)23Fe21(W,Mo)2C12, and 65-8789 Nb4WC5 are not added to simulate the XRD patterns of both SH1-6 and SH2-6. The 65-3190 Fe0.9Si0.1 is set as an

amorphous/disordered phase in SH1-6; however, the 65-7775 Cr1.07Fe18.93, 65-3190 Fe0.9Si0.1, 65-6323 Fe0.905Si0.095, and 060696 Fe are set as amorphous/disordered phases in SH2-6. Fitting plots are shown in Fig. 12. Weighted residual error (Rwp) and goodness values of fitting profile are given by about 14.59%, 1.38 and 12.50%, 1.2 for SH1-6 and SH2-6, respectively. The fitting works are thus considered as in a very good agreement. Similar results of amorphous/disordered content can be presented in Table 6. The grain size of 65-7775 Cr1.07Fe18.93, 654664 CrFe4, 65-6323 Fe0.905Si0.095, and 06-0696 Fe in SH1-6 is 26.8, 53.4, 26.8, and 26.8 nm, respectively. The grain size of 657775 Cr1.07Fe18.93 and 65-4664 CrFe4 in SH2-6 is both with 11.6 nm. These results with α-Fe in nanosized grain size are highlighted again (comparison to the grain size shown in Table 5). The iron and manganese boride, and iron and niobium carbide phase contents of SH1-6 are lower than those of SH2-6. However, the chromium intermetallic compound phase content of SH1-6 is higher than that of SH2-6. These results suggest that the hardness of SH2-6 is higher than that of SH1-6 because the total sum of boride and carbide phase contents is higher. On the other hand, α-Fe solid solution, some of borides, carbides with metalloids, and intermetallic compounds without metalloids in annealing specimen may have help for characteristic of the bonding strength compared with as-sprayed specimen (see in Fig. 11 and Table 6). These borides and carbides are also metallic [55,59,61]; hence improvement of the bonding strength is obtained. Sá Brito et al. stated that the brides formed in metallic coating onto carbon steel can improve the adhesion strength of the coating [69]. The reason causing the highest bonding strength is for the case at 500 °C, not for as-sprayed case, probably due to the precipitate of Cr23C6. However, formation of B2Fe15Si3 metastable phase may considerably cause the relative low bonding strength of SH2-4 coating in comparison with SH1-4. The existence of (Fe,Cr) and Co3Fe7 for SHS9172 + Co materials at 900 °C annealing may improve bonding strength as compared with SHS9172 at same annealing condition. 3.7. XPS analysis of surface The precipitated phase in coatings upon annealing at higher temperature is studied by XPS. Added cobalt in 1–2 wt.% range in SHS9172 + Co powder and coating specimens. Fig. 13 presents the XPS peaks of Si 2p, O 1s, and Co 2p. The data of XPS including Fe 2p3/2, Cr 2p3/2, Nb 3d, Mo 3d, B 1s, C 1s, Mn 2p3/2, Si 2p, O 1s, and Co 2p is shown in Table 7 for SHS9172 and SHS9172 + Co following annealing at 900 °C for 3 h. This Table presents the probable associated phases and other contents (Auger, multiplet splitting and shake-up lines) of SHS9172 and SHS9172 + Co following annealing at 900 °C for 3 h based on the XPS spectra. The XPS peaks are separated based on the reference [70] and using Pauling's electro-negativity scale [71]. The core level fitted peaks correspond to their probable associated phases with SH1-6 and SH2-6 are discussed further. 3.7.1. XPS peaks of Fe 2p3/2 Firstly, the deconvolution of Fe 2p3/2 energy level for SH1-6 and SH26 yields four peaks, respectively. Primary peak corresponds to Fe0/ Fe2.7Mn0.3C/Fe3Si/FeB/Fe2B (707.78 ev). The XPS peaks of Fe 2p3/2 for Fe0 are related to α-Fe, which is also identified from the XRD profiles. Peaks associated Fe2.7Mn0.3C/Fe3Si/FeB/Fe2B are also found; their peaks overlap each other at 707.78 eV and correspond to Fe0. Fe2B can be BFe2 (35-1340) and CrFeB (51-1410) phases in XRD. FeB usually grows up on the outer surface of iron borides at high temperature [68]. The Fe3Si may indicate that it is metastable Fe3B and/or Fe3(Si,B) [15], as compared with XRD analysis. These results suggest that the devitrification of coatings with increasing temperature causes transformation of metastable Fe3B into stable Fe2B [67]. On the other hand, M23C6 (i.e., (Cr,Fe,

W.-H. Liu et al. / Surface & Coatings Technology 249 (2014) 24–41 Cr 2p3/2

SH1-6 peak sum SH1-6 background SH1-6 peak 707.78 ev SH1-6 peak 708.40 ev SH1-6 peak 709.60 ev SH1-6 peak 711.00 ev

SH2-6 Coating

SH2-6 peak sum SH2-6 background SH2-6 peak 707.78 ev SH2-6 peak 708.50 ev SH2-6 peak 709.90 ev SH2-6 peak 711.90 ev

SH1-6 Coating

720

715

710

705

700

582

SH1-6 Coating

206

204

202

200

SH1-6 raw data SH1-6 peak sum SH1-6 background SH1-6 peak 203.70 ev SH1-6 peak 204.01 ev SH1-6 peak 205.43 ev SH1-6 peak 206.42 ev SH1-6 peak 206.73 ev SH1-6 peak 207.50 ev SH1-6 peak 208.15 ev SH1-6 peak 210.22 ev SH2-6 raw data SH2-6 peak sum SH2-6 background SH2-6 peak 203.70 ev SH2-6 peak 203.95 ev SH2-6 peak 205.43 ev SH2-6 peak 206.42 ev SH2-6 peak 206.67 ev SH2-6 peak 207.50 ev SH2-6 peak 208.15 ev SH2-6 peak 210.22 ev

578

576

574

572

570

W4f

Intensity (a.u.)

Intensity (a.u.)

SH2-6 Coating

208

580

B.E. (eV)

Nb 3d

210

SH1-6 Coating

695

B.E. (eV)

212

SH1-6 raw data SH1-6 peak sum SH1-6 background SH1-6 peak 574.75 ev SH1-6 peak 575.10 ev SH1-6 peak 575.50 ev SH1-6 peak 576.30 ev SH1-6 peak 577.10 ev SH2-6 raw data SH2-6 peak sum SH2-6 background SH2-6 peak 574.70 ev SH2-6 peak 575.00 ev SH2-6 peak 575.50 ev SH2-6 peak 576.25 ev SH2-6 peak 578.00 ev

SH2-6 Coating

Intensity (a.u.)

Intensity (a.u.)

Fe 2p3/2

725

35

SH1-6 peak 31.47 ev SH1-6 peak 32.25 ev SH1-6 peak 33.65 ev SH1-6 peak 34.43 ev

SH2-6 Coating

SH2-6 background SH2-6 peak 31.36 ev SH2-6 peak 32.17 ev SH2-6 peak 33.54 ev SH2-6 peak 34.35 ev

SH1-6 Coating

36

B.E. (eV)

35

34

33

32

31

30

B.E. (eV) SH1-6 raw data SH1-6 peak sum SH1-6 background SH1-6 peak 188.18 ev SH1-6 peak 188.80 ev SH1-6 peak 189.65 ev SH2-6 raw data SH2-6 peak sum SH2-6 background SH2-6 peak 188.10 ev SH2-6 peak 189.12 ev

Intensitry (a.u.)

SH1-6 raw data SH1-6 peak sum SH1-6 background SH1-6 peak 228.37 ev SH1-6 peak 229.10 ev SH1-6 peak 229.50 ev SH1-6 peak 231.55 ev SH1-6 peak 232.28 ev SH1-6 peak 232.68 ev SH2-6 raw data SH2-6 peak sum SH2-6 background SH2-6 peak 228.29 ev SH2-6 peak 229.10 ev SH2-6 peak 229.90 ev SH2-6 peak 231.47 ev SH2-6 peak 232.28 ev SH2-6 peak 233.08 ev

SH2-6 Coating

SH1-6 Coating

237

235

233

231

229

227

225

B.E. (eV)

Intensity (a.u.)

B1s Mo 3d

SH2-6 Coating

SH1-6 Coating

193.4

191.4

189.4

187.4

185.4

B.E. (eV)

Fig. 13. XPS patterns of SH1-6, and SH2-6 coatings. Peak of Si 2p cannot be separated because it overlaps with peak of Fe 3s or Co 3s.

W,Mo)23Fe21(W,Mo)2C12, Cr21.34Fe1.66C6 and Cr22.23Fe0.77C6) and M12C (i.e., Fe6W6C) are not mentioned in current work due to no study reported yet. In fact, SHS9172 alloy belongs to an Fe–C–Si–B quaternary alloy and thus M23C6 intermetallic phase can be gradually grown within it by annealing procedure [72,73]. The information, the intensity with main peak of Fe2.7Mn0.3C/Fe3Si/FeB/Fe2B peaks for SH2-6 is relatively high, implies that the concentration of Fe carbides, borides and silicides for SH2-6 is slightly higher than that for SH1-6. The trace of boron–iron– silicon phase (e.g., B2Fe15Si3 metastable phase in XRD analysis) onto the external surface of SHS9172 + Co coating may be still presented at such a high temperature of 900 °C.

Fe oxides are also the reason for the M23C6 and M12C species hard to identify, these oxides which formed on both SH1-6 and SH2-6 alloy surfaces in Fe 2p3/2 XPS signals (Fig. 13) and the result agrees with SEM observation (Figs. 5, 6, 8). This fact is consistent with the literatures reported [74–76]. Although they exist in both coatings, there are no continuous oxide films onto the coatings (Fig. 3) therefore the composition of their coatings is detected. The intensity of peak at 707.78 ev for SH2-6 is slightly higher than that for SH1-6, it means that the relative high content of Fe is presented, which is consistent with the result of SEM (Fig. 7). The information, the position of oxide peaks for SH2-6 shifted to higher binding energy, represents that iron concentration of oxide

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W.-H. Liu et al. / Surface & Coatings Technology 249 (2014) 24–41

SH1-6 peak sum SH1-6 background SH1-6 peak 282.20 ev SH1-6 peak 283.43 ev SH1-6 peak 283.90 ev SH1-6 peak 285.04 ev SH1-6 peak 286.53 ev

SH2-6 Coating

SH2-6 peak 282.25 ev SH2-6 peak 283.36 ev SH2-6 peak 283.90 ev SH2-6 peak 285.04 ev

Mn 2p3/2

SH2-6 Coating

Intensity (a.u.)

Intensity (a.u.)

C1s

SH1-6 peak 639.50 ev SH1-6 peak 640.10 ev SH1-6 peak 640.60 ev SH1-6 peak 641.60 ev SH1-6 peak 642.30 ev SH1-6 peak 643.00 ev SH1-6 peak 646.30 ev SH2-6 peak sum

SH1-6 Coating

SH2-6 peak 639.60 ev SH2-6 peak 640.10 ev SH2-6 peak 641.60 ev SH2-6 peak 643.00 ev SH2-6 peak 644.30 ev SH2-6 peak 646.30 ev

SH1-6 Coating

305

300

295

290

285

280

275

665

270

660

655

B.E. (eV)

650

645 640 B.E. (eV)

635

630

SH1-6 raw data SH1-6 peak sum SH1-6 background SH1-6 peak 530.68 ev SH1-6 peak 531.60 ev SH1-6 peak 533.20 ev SH2-6 raw data SH2-6 peak sum SH2-6 background SH2-6 peak 531.15 ev SH2-6 peak 531.60 ev SH2-6 peak 533.20 ev

O 1s SH1-0 raw data SH2-0 raw data SH1-6 raw data SH2-6 raw data

SH2-6 Coating

Intensity (a.u.)

Intensity (a.u.)

Si 2p

SH2-6 Coating

SH1-6 Coating

SH2-0 Powder

625

SH1-6 Coating

SH1-0 Powder

114 112 110 108 106 104 102 99.8 97.8 95.8 93.8 91.8 89.8 87.8

B.E. (eV)

550

Intensity (a.u.)

Co 2p

545

SH1-6 raw data SH1-6 peak sum SH1-6 background SH1-6 peak 784.30 ev SH1-6 peak 789.38 ev SH2-6 raw data SH2-6 peak sum SH2-6 background SH2-6 peak 779.29 ev SH2-6 peak 784.40 ev SH2-6 peak 789.00 ev SH2-6 peak 794.25 ev

SH2-6 Coating

540

535 530 B.E. (eV)

525

520

515

SH1-6 Coating

830

820

810

800 790 B.E. (eV)

780

770

760

Fig. 13 (continued).

is slightly higher than those for SH1-6. The fact is also compatible with the result of SEM (Fig. 8). 3.7.2. XPS peaks of Cr 2p3/2 Deconvolution of Cr 2p3/2 energy level of SH1-6 and SH2-6 yields five peaks, respectively. The main peaks of Cr 2p3/2 core level are associated with CrB2 or Cr–C compounds. Here, Cr–C may be assigned to (Cr, Fe,W,Mo)23Fe21(W,Mo)2C12, Cr21.34Fe1.66C6 Cr22.23Fe0.77C6 and Cr23C6, as compared with XRD analysis. XPS peaks at 574.70, 574.75, 575.00 and 575.10 eV for Cr 2p3/2 band are similar to those from carbides (Fig. 13). The Cr0 metallic-state is also observed in the 574.75– 574.7 eV and thus indicates that Cr metal is presented in the Fe-rich matrix. The coexistence of Cr0 and Fe0 (shown in Subsection 3.7.1.) may be assigned to (Fe,Cr), CrFe4, Cr1.07Fe18.93 and/or Cr6Fe18Mo6, as compared with XRD analysis. CrO2, Cr2O3, Cr(OH)3, and CrO3 are predicted to

present on the surfaces of several iron–chromium-based alloys [77, 78]. The metal related signals reveal that thickness of oxide layers is less than that of a few tens of Angstroms, explaining why photoelectrons that escaped from the alloy coatings can pass through oxide film. The information, the intensity with main peak and position of oxide peaks for SH1-6 is relatively high and shifted to higher binding energy, implies that the concentration of Cr and its oxide for SH1-6 is slightly higher than that for SH2-6. The fact is also compatible with the result of SEM (Figs. 7 and 8). Notably, the HVOF was performed in a reducing flame environment (Table 2) with a fuel to oxygen ratio of 2.48 (neutral fuel/ oxygen ratio of close to 2.1). Therefore, the coatings adsorbed hydroxyl species and water species and these contaminants are formed at their surfaces probably. The binding energies of the Fe 2p3/2 (see the previous Subsection) and Cr 2p3/2 electrons are 711.90 eV and 577.10 eV for trace FeOOH (711.5 ev) and Cr(OH)2 (577.1 ev), respectively [79].

W.-H. Liu et al. / Surface & Coatings Technology 249 (2014) 24–41

37

Table 7 XPS at heat treatment of 900 °C for 3 h. Core level

Fe 2p3/2

Cr 2p3/2

Nb 3d

W 4f Mo 3d

B 1s

C 1s

Mn 2p3/2

O 1s

Co 2p a

Probability phase or characteristic peak

0

Fe /Fe2.7Mn0.3C/Fe3Si/FeB/Fe2B Fe3O4 FeO Fe2O3 FeOOH Cr0/CrB2 CrB2/Cr–C CrO2 Cr2O3 Cr(OH)3 CrO3 NbC NbO Nb2O3 Nb2O5 W0 W–C Mo0/Mo–C/Mo–B MoO2 Mo2O5 Fe–B/Co–B/Cr–B/Mn–B/W–B/Mo–B Fe–B/Cr–B/Mn–B/W–B/Mo–B MoB2 Co2B B0 NbC Cr–C/W–C/Mo–C Fe2.7Mn0.3C C (graphite) C–O Mn–Si/Mn0 Fe2.7Mn0.3C MnCr2O4/MnFe2O4/MnO Mn2B/Mn2O3/(CoMn2O4)a MnSiO3 MnO2 Shake up Nb2O5/MoO3/WO3 Cr2O3/MoO3/(Co2SiO4)a SiO2 CoO Fe (LMM)

Bind energy (ev) SH1-6 coating

SH2-6 coating

707.78 708.40 709.60 711.00 – 574.75 575.10 575.50 576.30 577.10 – 203.70, 206.42 204.01, 206.73 205.43, 208.15 207.50, 210.22 31.47, 33.65 32.25, 34.43 228.37, 231.55 229.10, 232.28 229.50, 232.68 – 188.18 188.80 – 189.65 282.20 283.43 283.90 285.04 286.53 639.50 640.10 640.60 641.60 642.30 643.00 646.30 530.68 531.60 533.20 – 784.30,789.38

707.78 708.50 709.90 – 711.90 574.70 575.00 575.50 576.25 – 578.00 203.70, 206.42 203.95, 206.67 205.43, 208.15 207.50, 210.22 31.36, 33.54 32.17, 34.35 228.29, 231.47 229.10, 232.28 229.90, 233.08 188.10 – – 189.12 – 282.25 283.36 283.90 285.04 – 639.60 640.10 – 641.60 – 643.00 644.30, 646.30 531.15 531.60 533.20 779.29, 794.25 784.40, 789.00

(CoMn2O4) and (Co2SiO4) are only for SHS9172 + Co specimen.

3.7.3. XPS peaks of Nb 3d The Nb 3d spectra are deconvoluted with four peaks for both SH1-6 and SH2-6. The main Nb 3d band XPS peaks, corresponding to NbC and Nb oxides. They are explained by the fact that Nb in coatings interacts with C and O elements rather than B element. Since Nb is a strongly carbide-forming transition metal element and insoluble in Fe3C, it forms independent carbide phases (such as NbC) easily in Fe-based alloy that contains Nb metal. The result is compatible with the result of XRD (Fig. 12). No Nb0 peak is observed in Nb 3d region and the growth of NbO, Nb2O3 and Nb2O5 oxides on the surface of both coatings is easily seen as shown in previous literature [80,81]. Moreover, the highlighted results show that the NbC content of SH2-6 is higher than that of SH1-6 (compared with the intensity of peaks at Nb 3d5/2 of 203.70 eV), whereas the NbO content of SH2-6 is lower than that of SH1-6 (compared with the intensity of peaks at Nb 3d5/2 of 203.95 and 204.01 eV). 3.7.4. XPS peaks of W 4f and Mo 3d Two states of W0 and W–C, are formed in surfaces of both SH1-6 and SH2-6 and identified by the deconvolution of W 4f energy region (Fig. 13 and Table 7). The W0 metal state is with high intensity in W 4f band energy. This state shows the presence of one solute with an α-Fe matrix. W–C (e.g., WC = 32.17–32.25 eV in W 4f7/2) is applied

to a domain of high hardness and wear resistance with both SHS9172 and SHS9172 + Co coatings. Adding Co to SH2-6 makes the peak associated with the W 4f band shifts to a lower binding energy. This indicates that less W oxide is formed. The existence of W–C in W 4f region may also represent Fe6W6C and/or (Cr,Fe,W,Mo)23Fe21(W, Mo)2C12, as compared with XRD analysis. The deconvolution of Mo 3d energy region reveals that three Mo oxidation states in the SH1-6 and SH2-6 with Mo0/Mo–C/Mo–B, MoO2 and Mo2O5 (Fig. 13 and Table 7). Mo0/Mo–C/Mo–B yields the most intense Mo 3d energy peak for all species. Mo carbide and boride (e.g., Mo2C = 227.5–227.8 eV and MoB2 = 227.5–227.9 eV in Mo 3d5/2 [70,82,83]) are identified in both SH1-6 and SH2-6. Moreover, the highlighted result shows that Mo0/Mo–C/Mo–B content of SH2-6 is lower than that of SH1-6 (compared with the intensity of peaks at Mo 3d5/2 of 228.29 and 228.37 eV). The corresponding peak of SH2-6 shifts to lower binding energy. Therefore peak at 228.29 eV for SH2-6 shifts to corresponding Mo compounds (Mo–C/Mo–B) more closely than that for SH1-6, whereas the trend of peak at 228.37 eV for SH1-6 is toward corresponding Mo0 metal (Mo0 = 228.0 eV [70]). The interesting difference of the trend may result from unmelted particles in SHS9172 + Co coating, the particles further cause the self-fluxing reaction with varying heat treatment. The phenomenon is described by results of SEM mapping in Subsection 3.3 (Figs. 3 and 6).

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W.-H. Liu et al. / Surface & Coatings Technology 249 (2014) 24–41

The peaks at 229.1 eV and 232.28 eV are attributed to MoO2 typically [84]. Very small traces of Mo2O5 at 229.9–230.9 eV are also observed, as also reported by Damyanova et al. [85] The coexistence of Mo0 and Fe0 and Cr0 may form iron solid solution or Cr6Fe18Mo6 phase. And, the information, Mo–C in Mo 3d region, W–C in W 4f region and Cr–C in Cr 2p3/2 region, again shows that it is likely to be (Cr,Fe,W, Mo)23Fe21(W,Mo)2C12 phase corresponding to XRD analysis. 3.7.5. XPS peaks of Mn 2p3/2 Multi-peaks of Mn fitted in Mn 2p3/2 XPS band indicate presence of complex compounds. The main phase of Mn–Si/Mn0 is detected by XPS, but analysis of XRD results is difficult. The existence of Mn0 may form iron solid solution phase that has already been mentioned. Mn peaks in Mn 2p region are much weaker than other peaks associated with Fe 2p3/2, Cr 2p3/2, Nb 3d, Mo 3d, W4f, B 1s, C 1s, O 1s and Co 2p core levels. Data in the literature [86,87] reveal two satellite peaks (Mn sharke up) in Mn 2p3/2 region at approximately 644.30 eV and 646.30 eV. Mn3C might appear in both SH1-6 and SH2-6 coatings as Fe2.7Mn0.3C due to their lattice structure of Fe3C and Mn3C is the same isomorphic [88]. Mn2B may exist in both SH1-6 and SH2-6 coatings [89], the phase is compatible with the result of XRD (Fig. 12). Additionally, Mn can fully dissolve in Cr2O3 during annealing at high temperature [90] and form MnCr2O4 spinel at outer surface of SHS9172 coatings. The explanation of this observation is similar to that of Mn oxide containing Fe and Co, with weak peaks of MnFe2O4 in SH1-6 and CoMn2O4 in SH2-6. MnSiO3 has also been identified in SH1-6 coating in very small amounts and with Gibbs free energy lower than that of MnO or SiO2 in Ellingham diagram [91]. The content of Mn2O3 (641.60 ev) and MnO2 (643.00 ev) in Mn 2p3/2 region for SH2-6 is relatively high as illustrated in Fig. 13. Mn concentration of outer surface for SH1-6 is also lower than those for SH2-6 present in SEM mapping (Fig. 8). 3.7.6. XPS peaks of B 1s The deconvolution of B 1s energy level of SH1-6 and SH2-6 yields three and two peaks, respectively. The main peaks at binding energies of 188.10 eV and 188.18 eV are close to that reported in literature for metal borides (187.7 eV), B\C, and B\B bonds [70,92]. Interestingly, the peak at B 1s, 189.2 eV, corresponds to Co2B. This finding shows that added Co not only mixes with the original SHS9172 feedpowder but also reacts with other elements in SHS9172 + Co coating in HVOF process. This material participates with crystallization reaction of coating as presented in Fig. 2. The interdiffusion of B with W and Mo atoms is slower than that of with Fe and Cr atoms [93,94]. It is thus difficult to identify Mo 3d and W 4f peaks for Mo and W simple borides. Through deconvolution of B 1s region corresponding to peaks at 188.10 eV and 188.18 eV, the following: FeB, Fe2B, CoB, CrB2, W2B5, Mn2B and Mo2B5 transition metal borides are probably presented [70, 82]. Some transition metal boride phases (e.g., Fe2B and Mn2B) are identified and are also agreed well with XRD analysis (Fig. 12 and Table 6). As regards Cr, Mn, Mo and W borides, they combine with Fe borides (Fe 2p3/2 energy level) easily, they may exist in both coatings, as (Fe, M)B, (Fe,M)2B or (Fe,M)3B, where M denotes those transition metals [15]. MoB2 (188.80 eV) is identified in SH1-6, whereas Co2B (189.12 eV) in SH2-6. The presence of CoB and Co2B shows that adding cobalt can be dissolved in SH2-6 coating, the fact which is consistent with the results of SEM mapping in Subsection 3.3 (Figs. 4 and 7). Besides, the full width at half maximum (FWHM) of B1s peak for SH26 (FWHM = 0.87 at 188.10 eV) is broader than that for SH1-6 (FWHM = 0.85 at 188.18 eV). Some of this further evidence depicts that coarser unmelted particles in SHS9172 + Co may produce selffluxing reaction during varying heat treatment, as illustrated SEM observation in Subsection 3.3 (Figs. 3 and 6). The more boron in coatings could enable glass formation through high temperature process [8, 39–42]. Namely, SHS9172 + Co feedstock powder with similar core– shell structure having significant effect on coatings is highlighted again.

The results having more borides cause the higher hardness in SH2-6, as compared with SH1-6 (Fig. 10). The precipitation of transition metal borides may contribute positively to hardness of both coating annealed at 700 °C compared with those as sprayed coatings. Transition metal borides have been reported their hardness and wear properties to a greater extent than those properties for carbides in alloy [53–66,72, 95]. XRD data analysis demonstrates that B0 may be dissolved into Fe matrix (α-Fe). XRD and XPS analyses (Figs. 12 and 13) show that adding more than 0.01 wt.% B to steel formed various borides [96]. Neither B nor Si is volatilized out of spray stream as described elsewhere [97]. It is likely that the temperature (typically below 3000 °C) in HVOF flame is low enough that B and Si in alloys are not easy to volatilize. 3.7.7. XPS peaks of C 1s and O 1s Deconvolution of C 1s energy level for SH1-6 and SH2-6 yields five and four peaks, respectively. The peaks at 282.25 eV, 283.36 eV, and 283.90 eV in C 1s energy region are typically owing to NbC, Cr–C/W– C/Mo–C, and Fe2.7Mn0.3C, respectively. NbC for SH1-6 and SH2-6 is also found in XRD data (Fig. 12). Cr carbides (Cr\C bond at 283.36– 283.43 eV) including Cr23C6, Cr21.34Fe1.66C6, and Cr22.23Fe0.77C6 can be identified from the main peaks at binding energies of 282.8–283.1 eV [83]. However, these binding energies are also close to those of W–C (e.g., WC = 281–283 eV) and Mo–C (e.g., Mo2C = 282.7–282.9 eV) [83,98]. The Fe6W6C species found in XRD might be also seen on W–C overlapped peaks in C 1s binding energy spectrum. The Cr, Fe, W, and Mo with C are likely to form complex carbide in both SH1-6 and SH26 coatings at such high temperature, such as the presence of (Cr,Fe,W, Mo)23Fe21(W,Mo)2C12 phase corresponding to XRD analysis when coexistence of Cr–C, W–C, Mo–C and Fe–C species in C 1s energy region is shown. The O 1s band in XPS profiles shows that most oxide compounds close to outer surface of SH1-6 were Cr-, Nb-, Si-, W-, and Mo-oxides. However, O 1s band in XPS profiles reveals that Co2SiO4 may be present in SH2-6. In thin layer of oxides in SH1-6 and SH2-6 coatings, metal atoms diffuse toward outer surface while oxygen diffuses into them. The oxygen diffuses into coating following the thickness direction and thus the concentration distributions of oxygen vary in the thickness direction. C (graphite) and C–O species may come from the reducing flame atmosphere with high C concentration during HVOF spraying (neutral fuel/oxygen ratio of close to 2.1). The intensity of O 1s band (i.e., peak sum) in XPS profiles for SH1-6 is higher than that for SH2-6, the fact means the relatively high oxidation content in SH1-6 coating. 3.7.8. XPS peaks of Co 2p and Si 2p Although Co 2p3/2 peak at 784.30–784.4 eV overlaps Fe (LMM) Auger lines, intensity of peak from SH2-6 specimen is stronger than that from the corresponding SH1-6. Some Co compounds are present in SH2-6 but Co 2p peak shows the presence of CoO only (at 779.24 eV in Co 2p3/2). Confirming the presence of other Co compounds is rather difficult. There is a relative high intensity of Fe0 peak (Fe 2p3/2 energy level) and Co 2p peak at 784.4 eV in SH2-6 coating. We thus suggest that only a slight amount of Co is dissolved in Fe-rich matrix and accompanied by growth of Fe species. Namely, Co–Fe species (e.g., Co3Fe7 phase in XRD analysis) may be present in SH2-6. The relative intensity of Co 2p peak at 784.4 eV in SH2-6 coating also implies that formation of Co grows along Fe species, the finding agrees with SEM observation (i.e., a little Co come with Fe and is almost along Fe) (Fig. 4(c)). This metal–metal bond acts as a binder for SH2-6 coating. Specially, Co is presented at inner splats and between coating and substrate. This suggestion is based in part on SEM and XRD (Figs. 4, 7 and 12). Thus, the existence of Co at inner splats thus plays a dominant role in coating to improve mechanical properties of coating, such as by bonding well with themselves and its substrate (Fig. 11). Overlapping of binding energies of Si with Fe 3s or Co 3s core level makes the separation from peak of Si 2p difficult. The SHS9172 and SHS9172 + Co coatings yield broader XPS Si 2p core level peaks in the

W.-H. Liu et al. / Surface & Coatings Technology 249 (2014) 24–41

range of 99–104 eV than the corresponding powders. This observation is explained by growth of some silicon oxides on surfaces of coatings. This is also supported by the related peak of silicon oxide with coatings at 533.20 eV in O 1s region. Notably, XRD and XPS results (Figs. 12–13 and Tables 6–7) reveal that precipitation of boride and carbide phases and growth of α-Fe phase may be the main causes of the variation in hardness and bonding strength of coatings (Figs. 10–11) at different temperatures. 3.8. Wear testing The dry sand–rubber wheel abrasion test reveals that the significant variation in wear resistance of the coatings depends upon annealing temperature (Fig. 14). The porosity and the microhardness increase first and decrease then. The wear volume loss decreases first and increases then (Figs. 9, 10 and 14). Firstly, the wear volume loss for both SHS9172 and SHS9172 + Co coatings decreases at less than 500 °C suggests that this is closely related to precipitate content of α-Fe primary phase in amorphous/disordered phase (13.08 ± 1.14, 8.62 ± 0.37, 12.77 ± 0.07, 8.31 ± 0.31 mm3 for SH1-1, SH1-2, SH2-1, SH2-2, respectively). Although Fe2.7Mn0.3C, CrFeB and Cr23C6 phases have help for resistance of wear, they seem not to be major effect on the difference between wear volume loss of both coatings. The result is consistent with grain size and crystallinity of XRD calculated data (Table 5). By comparing SHS9172 with SHS9172 + Co, the interesting notice presents that the increasing trend of wear volume loss is probably related to B2Fe15Si3 metastable phase precipitate at more than 500 °C (6.27 ± 0.09, 10.73 ± 1.35, 5.44 ± 0.17, 9.22 ± 0.72 mm3 for SH1-4, SH1-6, SH2-4, SH2-6, respectively). In other words, the precipitate of B2Fe15Si3 metastable phase for SHS9172 + Co is earlier than that for SHS9172 by about 100 °C, the fact which may result from self fluxing reaction of unmelted particles in SHS9172 + Co (Subsections 3.3 (Figs. 3 and 6), 3.6 and 3.7). The value of wear volume loss for SH2-6 thus is lower than for SH1-6 due to more carbides, borides and silicides (i.e., Fe3Si) in SH2-6 (Subsection 3.7). The precipitate strengthening of complex intermetallic phases (e.g., (Cr,Fe,W,Mo)23Fe21(W,Mo)2C12, Cr23C6 and Fe3B) may lead to the minimum value of wear volume loss for both SH1-4 and SH2-4 coatings. The phase growth of these carbides and borides may then lead to the worse wear properties at such high temperature of 900 °C for both SH1-6 and SH2-6 coatings. The formation of many loose and discontinuous oxides (Figs. 5, 6 and 8) in coatings upon annealing at 900 °C also causes a slightly higher wear volume. However, it is nevertheless still lower than those of as-sprayed coatings. Adding Co improves anti-wear properties of coating. Whang and Giessen stated that a lower friction coefficient of amorphous alloys of Fe, Co, and Ni was associated with a lower wear volume [99]. Furthermore, Co coated on SHS9172 feedstock powder may form a solid solution of Fe-based system in SHS9172 + Co coatings during HVOF flame

16.0

SH1-1 SH2-1

SH2-1 SH2-2 SH1-4 SH2-4

39

spraying and subsequent annealing, subsequently improving the cohesion between splats in coating and adhesion between coating and substrate. 3.9. Electrochemical testing Fig. 15 plots potentiodynamic polarization curves of SHS9172 + Co coatings and the SHS9172 coating in 5% NaCl solution. The corrosion current density is inversely proportional to corrosion resistance. According to Tafel plot, corrosion current density for SH1-1, SH1-6, SH2-6, and SH2-1 can be determined to be 8.09, 2.49, 0.90 and 0.32 μA/cm2, respectively. Clearly, the amorphous/disordered SHS9172 + Co coatings exhibit a better corrosion resistance than the amorphous/disordered SHS9172 coatings. This result is attributed to lower crystallinity (%) and fine crystalline phase of SHS9172 + Co coating, and is also consistent with anticorrosion properties of amorphous Fe alloy [21]. Through microstructure observation with SHS9172 and SHS9172 + Co coatings (Figs. 5 and 6), unmelted particles in SH2-1 cannot help for corrosion resistance properties of its coating. Chokethawai et al. stated that the region of partially unmelted and/or unmelted particles is etched more easily than that of amorphous in the thermal sprayed coatings [10]. Additionally, when small amounts of Co replace Fe in Fe–Co–Cr–B–Si alloys, their corrosion resistance is increased [100]. Three factors including amorphous/ disordered, intermetallic phase content of coatings, and porosity of coatings, have effects on corrosion behavior. However, SH2-1 coating had a slightly better corrosion resistance than SH2-6 coating, the finding suggests that the relative high amorphous/disordered content in SH2-1 plays as the primary role in NaCl solution for corrosion resistance. The size of grains of intermetallic compound does not have significant difference between SH2-1 and SH2-6 coatings (Figs. 6, 12). The research indicates that intermetallic phase content of both coatings SHS9172 and SHS9172 + Co (belong to iron based coatings) is the secondary role for corrosion resistance (Fig. 12). These intermetallic phases, especially boride and carbide interstitial compounds, can decrease corrosion rate as reported in the literature [8,68,69,101]. Besides (Fe,Cr) and Co3Fe7 (see in Subsection 3.6), SH2-6 coating containing more boride intermetallic phases than other coatings (see in Subsection 3.7) thus exhibits greater resistance to corrosion in NaCl solution than SH1-6 and SH1-1. The minimum porosity for SH2-6 also contributes to resistance to corrosion (Fig. 9). Coarse nano-scale grains in SH1-6 coating also slightly reduce the area of grain boundaries because grain boundary attack may occur (Table 5). The formation of oxides including spinel (MnCr2O4 and MnFe2O4) and rhodonite (MnSiO3) may improve corrosion resistance due to the extremely low Gibbs free energy as discussed in Subsection 3.7 on XPS analysis. In contrast, the higher porosity for SH1-1 seems to be helpful to the corrosion (see in Fig. 9).

SH1-6 SH2-6

Volum Loss (mm3)

14.0 12.0 10.0 8.0 6.0 4.0 2.0 0.0

as spray

Annealing Temperature (oC) Fig. 14. Volume loss of as-sprayed SHS9172 and SHS9172 + Co coatings with postannealed at 500, 700, and 900 °C in dry sand wear testing.

Fig. 15. Potentiodynamic polarization of as-sprayed and post-annealed (at 900 °C) SHS9172 and SHS9172 + Co coatings.

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4. Conclusions This investigation examined the structural evolution of an Fe-based (SHS9172) alloy coating with the addition of cobalt and deposition by HVOF spraying. The main results of this investigation are as follows. 1. SHS9172 powders can be covered in a cobalt film by an electroless process and co-deposited with original SHS9172 powder as a modification from the film and form a novel composite amorphous coating. 2. OM observation of cross-section coatings depicts that the thickness of SHS9172 + Co coating onto 304 stainless steel substrate is thicker than that of SHS9172 coating onto such a substrate, under the same operation parameters for HVOF process. It suggests that adding Co to the Fe-based (SHS9172) alloy coating can improve the deposition rate. The average value of porosity for SH2-1 is slightly lower than that for SH1-1. Moreover, the minimum average value of porosity for SHS9172 + Co is obtained after annealing up to 900 °C for 3 h, as compared with other conditions for SHS9172 and SHS9172 + Co. 3. SEM observation of cross-section coatings depicts that a small amount of Co dissolved in the Fe-rich matrix and presented in/ besides the splat particles can act as a binder for the SHS9172 + Co coatings (Fig. 4). Based on the corresponding XRD analysis, the trace of Co3Fe7 in SH2-6 coating can be found (Table 6). Through XPS analysis, Co–Fe species is also found since the relative high intensity of Fe0 peak in Fe 2p3/2 energy level and the Co 2p peak at 784.4 eV. The evidence with the peak of B 1s also supports that Co may react with B to form CoB and/or Co2B. These results aim to provide an adequate explanation of Co effect on the SHS9172 + Co coating, that is, it can be dissolved in the Fe-rich matrix, accompanied by the growing iron species and formed itself intermetallic compounds when heat treatment temperature is increased. 4. The SHS9172 + Co has higher GFA (glass forming ability) than the SHS9172 due to the less difference between Tg (glass transition temperature) and Tx (peak crystallization temperature); and low liquidus or melting point since the existence of Co3Fe7. 5. As compared with the difference between SHS9172 and SHS9172 + Co respectively, the SH1-2 and SH2-2 have the maximum average values of tensile bonding strength, whereas the samples SH1-1 and SH2-1 have the minimum average values. The maximum average values of hardness for SHS9172 (Hv300 = 1216 of SH1-4) is higher than that for SHS9172 + Co samples (Hv300 = 1125 of SH2-4). However, the minimum average values of hardness (Hv300 = 745) are found at 900 °C for SH1-6. 6. Based on the corresponding main peak fitting with α-Fe for XRD calculated data and SEM high magnification images, nano-size grains of the primary phase in HVOF Fe-based coating may be not destroyed by annealing even up to 900 °C. 7. Comparing crystalline α-Fe phase content with its amorphous/ disorder in the coatings, the crystallinity (%) can be estimated after fitting the XRD profiles. The result is justified by a fact that the amorphous/disordered content of samples SHS9172 + Co is always higher than those of samples SHS9172. The coarse unmelted particles suggest that they can be considered as a domain role stabilizing amorphous/disordered phases in SHS917 + Co coatings. It also suggests that the self-fluxing reaction for those unmelted particles aims to investigate the cause and effect of Co added to HVOF SHS9172 coating. The precipitate of B2Fe15Si3 metastable phase for SHS9172 + Co (at 700 °C) is earlier than that for SHS9172 (at 800 °C). The findings may also correlate positively with the self fluxing reaction of unmelted particles in SHS9172 + Co. 8. The existence of (Fe,Cr) and Co3Fe7 for SHS9172 + Co at 900 °C annealing may improve the bonding strength as compared to SHS9172 at same annealing condition. 9. Besides the precipitate of B2Fe15Si3 borides, this research also indicates that the precipitate of carbides (i.e., (Cr,Fe,W,Mo)23Fe21(W,

Mo) 2 C 12 and 14-0407 Cr 23 C 6 ) and the growth of carbides (i.e., Cr21.34Fe1.66C6, Cr22.23Fe0.77C6, and 35-0783 and 65-3132 Cr23C6), and of Fe3B (34-1198 Fe3B, B2Fe15Si3) as well as CrFeB borides are responsible for the behaviors of porosity, microhardness, bonding strength and wear with both coatings at or above 700 °C annealing temperature. 10. The similar content tendency of significantly lighter and wellproportioned precipitation phase in SHS9172 and SHS9172 + Co coatings increases with an increasing annealing temperature (cross section of SEM study). Composite element mapping analysis identifies that W and Mo elemental contents are increased in these lighter regions. The precipitation particles with 80–200 nm embedded in the coatings are obtained. 11. Based on dry-sand testing, the best anti-wear properties are obtained when cobalt is added to HVOF SHS9172 coatings and annealed at a high temperature of 700 °C. Firstly, the wear volume loss for both SHS9172 and SHS9172 + Co coatings decreases at less than 500 °C. It suggests that there are precipitates of α-Fe in amorphous/ disordered phase. Secondly, the increasing trend of the wear volume loss probably is related to the B2Fe15Si3 metastable phase precipitate at more than 500 °C. The value of the wear volume loss for SH2-6 thus is lower than for SH1-6 due to more carbides, borides and silicides (i.e., Fe3Si) in SH2-6 (Subsection 3.7). 12. Based on electrochemical testing, the best anti-corrosion properties in 5% NaCl solution are obtained when cobalt is added to HVOF SHS9172 as-sprayed coatings. The properties are better than those of SH2-6, SH1-6 and SH1-1 coatings. Namely, the amorphous/ disordered phase content for SH2-1 is the highest, leads to the best corrosion resistance, and then the intermetallic phase content of annealed SHS9172 + Co is more than that of annealed SHS9172, the corrosion resistance of SH2-6 is thus better than SH1-6. Finally, the relatively high porosity of SH1-1 leads to the worst corrosion in contrast to other coatings. Conflict of interest statement There is no conflict of interest for our paper. References [1] D.J. Branagan, W.D. Swank, D.C. Haggard, J.R. Fincke, Metall. Mater. Trans. A 32 (2001) 2615–2621. [2] V. Ponnambalam, S.J. Poon, G.J. Shiflet, V.M. Keppens, R. Taylor, G. Petculescu, Appl. Phys. Lett. 83 (2003) 1131–1133. [3] N. Yang, J. Farmer, G. Lucadamo, J. Haslam, High-Performance Corrosion-Resistant Materials (HPCRM) Annual Report, Lawrence Livermore National Laboratory, Sandia National Laboratory, Livermore, California, 2004. [4] D.J. Branagan, M.C. Marshall, Be.E. Meacham, E.J. Buffa, High-Performance Corrosion-Resistant Materials (HPCRM) Annual Report, Lawrence Livermore National Laboratory, NanoSteel Company, Idaho Falls, Idaho, 2004. [5] J.C. Farmer, J.J. Haslam, S.D. Day, D.J. Branagan, M.C. Marshall, B.E. Mecham, E.J. Buffa, C.A. Bue, J.D.K. Rivard, D.C. Harper, M.B. Beardsley, D.T. Weaver, L.F. Aprigliano, L. Kohler, R. Bayles, E.J. Lemieux, T.M. Wolejsza, N. Yang, G. Lucadamo, J.H. Perepezko, K. Hildal, L. Kaufman, A.H. Heuere, F. Ernst, G.M. Michal, H. Kahn, E.J. Lavernia, High-Performance Corrosion-Resistant Materials (HPCRM) Annual Report, UCRL-TR-206717, Lawrence Livermore National Laboratory, Livermore, California, 2004. [6] J.C. Farmer, J.J. Haslam, S.D. Day, D.J. Branagan, C.A. Blue, J.D.K. Rivard, L.F. Aprigliano, N. Yang, J.H. Perepezko, M.B. Beardsley, Pressure Vessels and Piping Division (Publication) PVP, 7, American Society of Mechanical Engineers, New York, NY 10016–5990, 2005. [7] D.J. Branagan, M. Breitsameter, B.E. Meacham, V. Belashchenko, J. Therm Spray Techn. 14 (2005) 196–204. [8] J.C. Farmer, J.J. Haslam, S.D. Day, T. Lian, C.K. Saw, P.D. Hailey, J.-S. Choi, R.B. Rebak, N. Yang, J.H. Payer, J.H. Perepezko, K. Hildal, E.J. Lavernia, L. Ajdelsztajn, D.J. Branagan, E.J. Buffa, L.F. Aprigliano, J. Mater. Res. 22 (2007) 2297–2317. [9] J. Blink, J. Farmer, J. Choi, C. Saw, Metall. Mater. Trans. A 40 (2009) 1344–1354. [10] K. Chokethawai, D.G. McCartney, P.H. Shipway, J. Alloys Compd. 480 (2009) 351–359. [11] Y. Wu, P. Lin, Z. Wang, G. Li, J. Alloys Compd. 481 (2009) 719–724. [12] H.S. Ni, X.H. Liu, X.C. Chang, W.L. Hou, W. Liu, J.Q. Wang, J. Alloys Compd. 467 (2009) 163–167. [13] D.J. Branagan, W.D. Swank, B.E. Meacham, Metall. Mater. Trans. A 40 (2009) 1306–1313.

W.-H. Liu et al. / Surface & Coatings Technology 249 (2014) 24–41 [14] J. Zhou, J.K. Walleser, B.E. Meacham, D.J. Branagan, J. Therm. Spray Techn. 19 (2010) 950–957. [15] D. Zois, A. Lekatou, M. Vardavoulias, T. Vaimakis, A.E. Karantzalis, Surf. Coat. Technol. 206 (2011) 1469–1483. [16] W.H. Wang, C. Dong, C.H. Shek, Mater. Sci. Eng. R 44 (2004) 45–89. [17] A. Inoue, B.L. Shen, A. Takeuchi, Mater. Trans. JIM 47 (2006) 1275–1285. [18] B.L. Shen, A. Inoue, C.T. Chang, Appl. Phys. Lett. 85 (2004) 4911–4913. [19] J. Zhang, H. Tan, Y.P. Feng, Y. Li, Scr. Mater. 53 (2005) 183–187. [20] Y. Takahara, H. Matsuda, Mater. Trans. JIM 36 (1995) 903–908. [21] D. Zois, A. Lekatou, M. Vardavoulias, J. Alloy. Compd. 504 (2010) 283–287. [22] S. Sampath, Mater. Sci. Eng. A 167 (1993) 1–10. [23] O. Racek, High-Performance Corrosion-Resistant Materials (HPCRM) Annual Report, LLNL-JRNL-402714, Lawrence Livermore National Laboratory, Livermore, California, 2008. [24] SHS Product Technical Data, The NanoSteel, Co, Web site http://www.nanosteelco. com/products/tech_data_pdf/9172_hvof_tech_data.pdf (Accessed April 27, 2012). [25] A. Inoue, Acta Mat. 48 (2000) 279–306. [26] A. Inoue, Mater. Sci. Eng. A 304–306 (2008) 1–4. [27] A. Makino, T. Kubot, M. Makabe, C.T. Chang, A. Inoue, Mater. Sci. Eng. B 148 (2008) 166–170. [28] A.M. Bayer, B.A. Becherer, T. Vasco, in: ASM International (Ed.), 9th ed., High Speed Tool Steels ASM Handbook, vol. 16, ASM International, 1989, pp. 51–59, (Metals Park, OH). [29] R.F. Decker, S. Floreen, in: R.K. Wilson (Ed.), Maraging Steels, Proc. Maraging Steel Symposium, TMS, 1988, pp. 1–38. [30] P. Crook, A.I. Asphahani, S.J. Matthews, US Patent No. 5,002,731 (1991). [31] T. Zhang, D. Tang, Recent Patents on Corrosion Science 1 (2009) 1–5. [32] H.M. Rietveld, J. Appl. Crystallogr. 2 (1969) 65–71. [33] Z.P. Lu, H. Tan, Y. Li, S.C. Ng, Scr. Mater. 42 (2000) 667–673. [34] Z.P. Lu, Y. Li, S.C. Ng, J. Non-Cryst. Solids 270 (2000) 103–114. [35] Z. Zhou, L. Wang, F.C. Wang, Y.B. Liu, Trans. Nonferrous Metals Soc. China 19 (2009) 636–638. [36] P.G. Tao, Y.Z. Yang, Z.J. Dong, Y.P. Ren, Rare Metal Mater. Eng. 40 (2011) 1625–1628. [37] K. Amiya, A. Urata, N. Nishiyama, A. Inoue, J. Appl. Phys. 97 (2005) (10 F91310F913-3). [38] V.V. Sobolev, J.M. Guilemany, Mater. Lett. 18 (1994) 304–308. [39] L. Lin, K. Han, Surf. Coat. Technol. 106 (1998) 100–105. [40] Y.H. Shieh, J.T. Wang, H.C. Shih, S.T. Wu, Surf. Coat. Technol. 58 (1993) 73–77. [41] I.C. Grigorescu, C. Di Rauso, R. Drira-Halouani, B. Lavelle, R. Di Giampaolo, J. Lira, Surf. Coat. Technol. 76 (1995) 494–498. [42] H. Sarjas, D. Goljandin, P. Kulu, V. Mikli, A. Surženkov, P. Vuoristo, Mater. Sci. 18 (2012) 34–39 (Medžiagotyra). [43] H.F. Li, H.B. Zhao, X.W. Geng, Z.J. Fan, China Build. Mater. Sci. Technol. 17 (2008) 68–71. [44] R. Mcpherson, Thin Solid Films 83 (1981) 297–310. [45] S. Dallaire, Thin Solid Films 95 (1982) 237–244. [46] M. Parco, L. Zhao, J. Zwick, K. Bobzin, E. Lugscheider, Surf. Coat. Technol. 201 (2007) 6290–6296. [47] Z.X. Ding, W. Chen, Q. Wang, Trans. Nonferrous Metals Soc. China 21 (2011) 2231–2236. [48] A. Inoue, T. Zhang, Y.H. Kim, Mater. Trans. JIM 38 (1997) 749–755. [49] P.S. Frankwicz, S. Ram, H.-J. Fecht, Appl. Phys. Lett. 68 (1996) 2825–2827. [50] H.S. Ingham, A.P. Shepard, Metco Flame Spray Handbook, Metco Inc., New York, 1965. [51] D.M. Mini´c, A. Gavrilovi´, P. Angerer, D.G. Mini´c, A. Mariˇci´c, J. Alloy. Compd. 476 (2009) 705–709. [52] D. Kodama, K. Shinoda, K. Sato, Y. Konno, R.J. Joseyphus, K. Motomiya, H. Takahashi, T. Matsumoto, Y. Sato, K. Tohji, Adv. Mater. 18 (2006) 3154–3159. [53] A. Miyoshi, A. Hara, J. Jpn. Soc. Powder Metall. 12 (1965) 78–84. [54] G. Roberts, G. Krauss, R. Kennedy, Tool Steels, 5th ed. ASM International, Materials Park, OH, 1998. 45–65. [55] Y. Kumashiro, Y. Nagai, H. Katō, J. Mater. Sci. Lett. 1 (1982) 49–52. [56] M. Anis, W.M. Rainforth, H.A. Davies, Wear 172 (1994) 135–145. [57] N. Parkansky, I.I. Beilis, L. Rapoport, R.L. Boxman, S. Goldsmith, Yu. Rosenberg, Surf. Coat. Technol. 105 (1998) 130–134. [58] A. K¨uper, X. Qiao, H.R. Stock, P. Mayr, Surf. Coat. Technol. 130 (2000) 87–94.

41

[59] P. Vajeeston, P. Ravindran, C. Ravi, R. Asokamani, Phys. Rev. B 63 (2001) (045115045115-13). [60] D.S. Wang, E.J. Liang, M.J. Chao, B. Yuan, Surf. Coat. Technol. 202 (2008) 1371–1378. [61] B. Xiao, J. Feng, C.T. Zhou, J.D. Xing, X.J. Xie, Y.H. Cheng, R. Zhou, Physica B 405 (2010) 1274–1278. [62] T. Liyanage, G. Fisher, A.P. Gerlich, Surf. Coat. Technol. 205 (2010) 759–765. [63] S.Q. Ma, J.d. Xing, D.W. Yi, H.G. Fu, J.J. Zhang, Y.F. Li, Z.Y. Zhang, Surf. Coat. Technol. 205 (2011) 4902–4909. [64] J. Greim, K.A. Schwetz, Ullmann's 6 (2012) 220–236, http://dx.doi.org/10.1002/ 14356007.a04_295.pub2. [65] R. Păunescu, I. Gherghescu, U.P.B., Sci. Bull. Ser. B 74 (2012) 225–234. [66] S. Yin, D. He, C. Xu, W. Wang, H. Wang, L. Li, L. Zhang, F. Liu, P. Liu, Z. Wang, C. Meng, W. Zhu, Int. J. High Pres. Res. 33 (2013) 409–417. [67] T. Koziel, A. Zielinska-Lipiec, Z. Kedzierski, T. Czeppe, J. Microsc. 224 (2006) 27–29. [68] I. Campos, M. Palomar, A. Amador, R. Ganem, J. Martinez, Surf. Coat. Technol. 201 (2006) 2438–2442. [69] V.R.S. Sá Brito, I.N. Bastos, H.R.M. Costa, Mater. Design. 41 (2012) 282–288. [70] J.F. Moulder, W.F. Stickle, P.E. Sobol, K.D. Bomben, in: J. Chastain (Ed.), Handbook of XRay Photoelectron Spectroscopy, 2nd ed., Perkin-Elmer, Corp., Eden Prairie, MN, 1992. [71] L. Pauling, In the Nature of the Chemical Bond, 3rd ed. Wiley, New York, 1960. [72] J.H. Kim, K.H. Ko, S.D. Noh, G.G. Kim, S.J. Kim, Wear 267 (2009) 1415–1419. [73] P.S. Bouaricha, J.-G. Legoux, C. Moreau, P. Ctibor, S. Nourouzi, A. Vardelle, Advances in Technology and Application: Proceedings of the International Thermal Spray Conference, 10–12 May, 2004, Osaka, Japan, 2004, (CD-ROM, file 025.pdf). [74] I. Olefjord, H. Fischmeister, Corros. Sci. 15 (1975) 697–707. [75] S.K. Sharma, S. Hofmann, Appl. Surf. Sci. 51 (1991) 139–155. [76] M. Detroye, F. Reniers, C. Buess-Herman, J. Vereecken, Appl. Surf. Sci. 144–145 (1999) 78–82. [77] R. Natarajan, N. Palaniswamy, M. Natesan, V.S. Muralidharan, Open Corros. J. 2 (2009) 114–124. [78] S. Bera, S. Rangarajan, S.V. Narasimhan, Corros. Sci. 42 (2000) 1709–1724. [79] K. Asami, K. Hashimoto, S. Shimodaira, Corros. Sci. 17 (1977) 713–723. [80] M.D. Novak, C.G. Levi, Proceedings of the ASME International Mechanical Engineering Congress and Exposition, Washington, IMECE2007-42908, 2007, pp. 1–7. [81] R.A. Perkins, G.H. Meier, J. Min. Met. Mater. Soc. 42 (1990) 17–21. [82] G. Mavel, J. Escard, P. Costa, J. Castaing, Surf. Sci. 35 (1973) 109–116. [83] U. Muller, C.V. Falub, G. Thorwarth, C. Voisard, R. Hauert, Acta Mater. 59 (2011) 1150–1161. [84] D.O. Scanlon, G.W. Watson, D.J. Payne, G.R. Atkinson, R.G. Egdell, D.S.L. Law, J. Phys. Chem. C 114 (2010) 4636–4645. [85] S. Damyanova, L. Petrov, P. Grange, Appl. Catal. A General 239 (2003) 241–252. [86] C.J. Cong, L. Liao, Q.Y. Liu, J.C. Li, K.L. Zhang, Nanotechnology 17 (2006) 1520–1526. [87] A.A. El-Moneim, E. Akiyama, H. Habazaki, A. Kawashima, K. Asami, K. Hashimoto, Corros. Sci. 40 (1998) 235–250. [88] J. Briki, S. Ben Slima, J. Metall. 841901 (2012) 1–4. [89] V.V. Nemoshkalenko, T.B. Shashkina, V.G. Aljeshin, A.I. Senkevich, J. Phys. Chem. Solids 36 (1975) 37–40. [90] S. Ranganathan, J.P. Hajpa, Bull. Mater. Sci. 9 (1987) 149–158. [91] K. Yanagihara, S. Yamazaki, Nippon Steel Technical Report, UDC 543. 5: 669. 15Õ782, 100, 2011. 27–32. [92] M. Büttner, H.-J. Wang, A.M. Dongare, G.J. Shiflet, P. Reinke, P. Oelhafen, B.S. Mun, X.J. Gu, S.J. Poon, J. Appl. Phys. 102 (2007) (033501-033501-7). [93] S.P. Lua, O.Y. Kwon, Surf. Coat. Technol. 153 (2002) 40–48. [94] C.V. Cooper, G.G. Peterson, Surf. Coat. Technol. 43–44 (1990) 756–767. [95] V.I. Dybkov, L.V. Goncharuk, V.G. Khoruzha, A.V. Samelyuk, V.R. Sidorko, Mater. Sci. Technol. 27 (2011) 1502–1512. [96] J. Lorinczi, G. Kralik, M. Kovacs, A. Horvath, Mater. Sci. Forum 414–415 (2003) 267–274. [97] P. Niranatlumpong, H. Koiprasert, Surf. Coat. Technol. 206 (2011) 440–445. [98] E. Cappelli, S. Orlando, F. Pinzari, A. Napoli, S. Kaciulis, Appl. Surf. Sci. 138–139 (1999) 376–382. [99] S.H. Whang, in: T. Masumoto, K. Suzuki (Eds.), Wear Behavior of an Amorphous Alloy, Proc. 4th Znt. Conf. on Rapidly Quenched Metals, Sendai, August 24–28, 1981, Japan Institute of Metals, Tokyo, 1982, pp. 1403–1406. [100] I. Solomom, N. Solomon, Rev. Metal. Madrid 46 (2010) 389–395. [101] V. Cremaschi, I. Avram, T. Pérez, H. Sirkin, Scr. Mater. 46 (2002) 95–100.