Properties of HVOF-sprayed TiC-FeCrAl coatings

Author’s Accepted Manuscript Properties of HVOF-sprayed TiC-FeCrAl coatings Giovanni Bolelli, Alberto Colella, Luca Lusvarghi, Pietro Puddu, Rinaldo Rigon, Paolo Sassatelli, Veronica Testa www.elsevier.com/locate/wear

PII: DOI: Reference:

S0043-1648(18)30921-9 https://doi.org/10.1016/j.wear.2018.11.002 WEA102537

To appear in: Wear Received date: 31 July 2018 Revised date: 22 October 2018 Accepted date: 4 November 2018 Cite this article as: Giovanni Bolelli, Alberto Colella, Luca Lusvarghi, Pietro Puddu, Rinaldo Rigon, Paolo Sassatelli and Veronica Testa, Properties of HVOF-sprayed TiC-FeCrAl coatings, Wear, https://doi.org/10.1016/j.wear.2018.11.002 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Properties of HVOF-sprayed TiC-FeCrAl coatings Giovanni Bolellia,*, Alberto Colellab, Luca Lusvarghia,c, Pietro Puddua, Rinaldo Rigond, Paolo Sassatellia,1, Veronica Testaa

a

Department of Engineering “Enzo Ferrari”, University of Modena and Reggio Emilia, Via Pietro Vivarelli 10/1, 41125 Modena (MO), Italy b

c

MBN Nanomaterialia S.p.A., 31050 Vascon di Carbonera (TV), Italy

Consorzio Interuniversitario Nazionale per la Scienza e Tecnologia dei Materiali (INSTM), Local

Unit Università di Modena e Reggio Emilia, Via Pietro Vivarelli 10/1, 41125 Modena (MO), Italy d

Ecor International S.p.A., via Friuli 11, 36015 Schio (VI), Italy

*

Corresponding author: Tel.: +39 0592056233; fax: +39 0592056243; [email protected]

Abstract As an alternative to WC-CoCr and Cr3C2-NiCr coatings for wear and corrosion protection, a TiC – 25vol.% (Fe-20wt.%Cr-5wt.%Al) powder, free from hazardous and/or supply-critical elements (Ni, Co, W), was produced by high-energy ball-milling and processed by High Velocity Oxygen-Fuel (HVOF) spraying, obtaining dense (<1 vol.% porosity), hard (HIT > 12 GPa) layers with reasonably good deposition efficiency of ≈54%. Tribological testing revealed that the TiC-FeCrAl coatings are particularly promising for sliding contacts, as their ball-on-disc wear rates against an Al2O3 counterpart were lower than those of an HVOF-sprayed Cr3C2-NiCr reference, both at room temperature and at 400 °C, although they could

1

Present address: Il Sentiero International Campus S.r.l., Via Friuli 11, 36015 Schio (VI), Italy

not match the performance of WC-CoCr. At room temperature, brittle fracture along oxidized lamellar boundaries caused localized spallation, releasing debris in the contact region, but, in the incubation period before spallation cracks could propagate, remarkably low friction (≈0.27) was recorded. At 400 °C, spallation was largely suppressed by thermal softening, whilst coarser abrasive grooving became the dominant wear mechanism. TiC-FeCrAl coatings appeared less suited to high-stress abrasion, since extensive brittle fracture resulted in higher wear rates than HVOF-sprayed Cr3C2-NiCr, and to (acidic) corrosive environments. Electrochemical polarisation tests in 0.1 M HCl indeed revealed limited corrosion resistance of the FeCrAl matrix.

Keywords: Thermal spray coatings; Cermets; Sliding wear; High temperature; Three-body abrasion

1. Introduction A significant challenge for materials science and engineering research is nowadays posed by the need to replace critical materials, where the “criticality” can stem from two distinct, but equally important factors: (1) Supply scarcity and insecurity, which are especially an issue when a material is also essential to one or more key industrial technologies. Such considerations have prompted many industrialized countries to develop lists of “critical raw materials” and related policies aiming to secure their supply, enhance their recycling, and/or promote their substitution with more abundant and easily sourced ones [1]. The European Union has been particularly active on this subject lately, releasing comprehensive documents on critical raw materials, their profiles and their substitutionability [2,3]. Among these materials, cobalt and tungsten stand out for their remarkable importance in the field of protective coatings, being the main constituents of thermally sprayed WC-Co-based hardmetals [4,5] and of WC-based metal-matrix composite layers manufacturing by cladding [6,7]. These coatings find numerous applications for their wear and corrosion resistance under particularly severe conditions. Chromium, which was added to the list of critical raw materials for the EU in 2014 [8], has been removed from it in the 2017 update [3]; however, it still remains very close to the supply risk threshold in addition to its key economic importance [9,10]. It is therefore important to try reducing its consumption. In its most significant application as the fundamental alloying element of stainless steels, its complete replacement is deemed unfeasible, but approaches have been suggested to reduce its usage [1]. Analogously, it would be desirable to reduce the consumption of Cr in Cr3C2-NiCr hardmetal coatings, frequently used for their high-temperature wear resistance [11] and better thermal expansion compatibility to steel substrates than WC-Cobased ones [12].

(2) Health hazardousness, which raises safety issues particularly during those stages of the manufacturing and service lifecycle of a component, when direct contact with the hazardous material(s) is involved. Ni, Co and WC-Co are frequently used in inhalable, powdered form in the production of the above-mentioned wear- and corrosion-resistant coatings. Inhalation of these powders produces a variety of adverse acute and long-term reactions in the respiratory system, as well as potential carcinogenesis [13–17], which is one of the main risk factors in the thermal spray industry [18]. Based on the above considerations, an obvious need emerges to find replacements for conventional wear- and corrosion-resistant coating materials. With specific regard to the thermal spray field, such replacement has been identified as a way to further enhance the occupational and environmental safety advantages of this family of techniques [19] over competitive processes such as hard chromium electroplating [20]. TiC-based hardmetals and metal-matrix composites have long been proposed as thermal spray coatings, drawing on extensive research concerning their use in bulk form, as an alternative to WC-based cutting tools [21]. Compositions such as (Ti,Mo)(C,N)-Ni(Co) have been proposed [22–26]. A common finding inferable from all of the mentioned studies is that the wear resistance of TiC-based coatings is somewhat intermediate between that of Cr3C2-NiCr and WC-Co-type hardmetals under various wear testing conditions, including dry sliding, rubberwheel abrasion, and dry erosion. The related wear mechanism were reportedly quite typical of thermal spray coatings in general, often involving a certain degree of brittleness as evidenced by splat removal in ball-on-disc tests against an alumina counterbody and by brittle fracturing when eroded by hard particles impinging at 90° [22]. Under high-temperature sliding wear conditions, the coatings were reported to maintain good wear resistance up to at least 600 °C; moreover, the coatings develop a thin oxide scale starting from 400 °C [26]. Irrespective of technical considerations, a Fe-based matrix would be preferable [19] to avoid or minimize the content of Co, Ni and Cr. TiC-Fe-based feedstock powders can be produced by a number of processes, including conventional powder metallurgy, carbothermic or aluminothermic

reduction, and combustion synthesis, as reviewed by Das et al. [27]. Carbothermic reduction of oxide precursors, such as ilmenite, has indeed been proposed in [28,29]. In-situ combustion synthesis has been exploited by using the heat of the thermal spray jet to trigger the self-propagating reaction in powders consisting of agglomerated mixtures of ferrotitanium alloys with C sources such as graphite [30,31], asphalt or sucrose [32–36], with optional addition of a Ni source (e.g. carbonyl Ni) for a Fe-Ni matrix. Alternatively, self-propagating high-temperature combustion synthesis (SHS) was employed to pre-synthesize a TiC-containing powder, which was subsequently sprayed, as in [37,38]. Conventional powder metallurgy techniques such as agglomeration and sintering [31,39,40] (with optional, subsequent plasma densification [41]) were also used to produce feedstock powders from fine TiC and Fe-alloy particles, which were dispersed in water, spray-dried, and consolidated at high temperature. However, in spite of extensive research, these materials have not yet made their way into industrial production. Moreover, less information is available on the tribological behaviour of TiC-Fe-type coatings, although a tendency to brittle behaviour, similar to TiC-Ni/Co compositions, can be inferred by the occurrence of surface delamination during abrasion by angular alumina particles and of sub-surface cracking when abraded by rounded silica sand [37]. The present paper aims to characterize the properties of coatings obtained by High-Velocity Oxygen Fuel (HVOF) spraying of a TiC-based hardmetal powder with Fe-based alloy matrix, produced through an industrial high-energy ball milling process already in use for large-scale powder production, as previously illustrated in [42,43]. The production of the feedstock presented in this study is therefore amenable for direct, cost-effective scale-up. TiC was synthesized from precursors (elemental Ti and C) during the milling process itself, which implies multiple advantages, including: (i)

excellent bonding between the hard phase and the surrounding matrix;

(ii)

no need for a costly high-temperature post-treatment, as it occurred for the SHS powders described in the previously mentioned refs. [37,38];

(iii)

better control over the chemical and phase composition, in comparison to in-situ synthesis, which led to large splat-to-splat variability in the coatings described in [32–34,36].

A FeCrAl alloy, similar to Kanthal-type alloys [44], was chosen as the matrix for its potentially good corrosion and oxidation resistance. It has indeed been suggested in [1] that FeCrAl alloys can potentially yield stainless Fe-based alloys with reduced consumption of Cr. The material is also free from W, Co and Ni, with considerable advantages in terms of occupational safety and of reduced critical raw materials usage. The paper particularly focuses on room- and high-temperature dry sliding wear behaviour, highstress particle abrasion and electrochemical corrosion in aqueous environment, comparing the coating properties to those of reference, WC- and Cr3C2-based ones obtained by the same HVOF deposition process.

2. Experimental 2.1

Feedstock powders and coating manufacturing

The feedstock powder was produced by MBN Nanomaterialia (Vascon – TV, Italy) using a proprietary, industrial high-energy ball milling process starting from elemental Ti, C, and a FeCrAl alloy, mixed in suitable amounts to achieve an eventual composition of 75 vol.% TiC – 25 vol.% (Fe-20wt.%Cr-5wt.%Al). The milled powder was air classified and sieved between 10 µm and 38 µm, vacuum packed to prevent oxidation, and eventually pre-heated to 60 °C shortly before spraying, in order to enhance flowability by avoiding humidity uptake in the feeding system. The powder was sprayed using a Diamond Jet 2600 (Oerlikon-Metco, Wholen, CH) HVOF torch onto AISI 304 stainless steel plates of 60 mm length × 25 mm width × 3 mm thickness. The latter were manually grit-blasted using 36 mesh brown alumina abrasive at a pressure of 5.5 bar up to a surface roughness Ra ≈ 6 µm, ultrasonically cleaned in acetone, and pre-heated to 60 °C. The torch was scanned linearly in front of the substrates, mounted on a rotating turntable, in order to achieve a

linear torch-substrate speed of 0.75 m/s and a pass spacing (pitch) of 4 mm. Compressed air jets were used to cool the substrates and blow any overspray off their surface. Two distinct sets of process parameters were chosen (Table 1) based on previous experience with HVOF processing of hardmetal compositions: higher oxygen and fuel (H2) flow rates, with an overall reducing flame stoichiometry (equivalence ratio λ = 0.86), were employed to create a hot and fast jet, whilst lower flow rates with almost neutral flame stoichiometry (λ = 0.98) produced a somewhat slower, colder flame.

Table 1: HVOF process parameters. TiC-FeCrAl

WC-10Co4Cr

Cr3C2-NiCr

Set 1

Set 2

O2 flow rate [SLPM]

214

188

214

234

H2 flow rate [SLPM]

635

553

635

615

Air flow rate [SLPM]

274

274

274

344

Equivalence ratio λ

0.86

0.98

0.86

1

O2 pressure [psi/kPa]

170 / 1172

170 / 1172

170 / 1172

H2 pressure [psi/kPa]

140 / 965

140 / 965

140 / 965

Air pressure [psi/kPa]

100 / 689

100 / 689

100 / 689

Powder feed rate [g/min]

20

70

75

Stand-off distance [mm]

240

240

250

Torch-to-substrate relative velocity [m/s]

0.75

0.75

1.00

Track pitch [mm]

4

6

4

N° of cycles

24

40

30

The deposition efficiency was calculated as the ratio

⁄

between the actually

deposited coating mass and the overall mass of powder sprayed towards the substrates. The former was measured by weighing the plates before and after spraying, using an electronic scale with ±0.01 g accuracy. The latter was calculated based on the powder feed rate ( ̇ ), the length (l) and width (w) of the substrate, the torch/substrate relative traverse speed (v), the track pitch (p) and the overall number of torch cycles (n), according to equation (1): (⁄ ) ( ⁄ ) ̇

(1)

As a term of comparison, commercially available WC-10Co-4Cr (WOKA 3652) and Cr3C2-25(Ni20Cr) (WOKA 7102) powders (both from Oerlikon Metco WOKA GmbH, Barchfeld, Germany) were also sprayed by the same HVOF torch, using previously optimized process parameters listed in Table 1.

2.2

Characterization of microstructure, phase composition, and hardness

The particle size distribution of the feedstock powder was measured by laser diffraction using a wet dispersion technique (Mastersizer with Hydro-2000 wet dispersion unit: Malvern Instruments, Malvern, UK). Its oxidation behaviour in static air was characterized by simultaneous thermogravimetry (TG) / differential thermal analysis (DTA) using an STA409 apparatus (NETZSCH-Gerätebau GmbH, Selb, Germany): ≈100 mg of feedstock powder were heated in an alumina crucible at a rate of 20 °C/min up to 1200 °C The phase composition of the TiC-FeCrAl powder and coatings was assessed by X-ray diffraction (XRD: X’Pert PRO diffractometer, PANAlytical, Almelo, NL) using Cu-Kα radiation emitted from a source tube operated at 40 kV, 40 mA and collected by a 1D array of solid-state detectors (X’Celerator). Patterns were acquired over a 20° < 2θ < 85° angular range, with a scan step of 0.017° and a counting time of 2 s/step.

The microstructure of the powder and of the corresponding HVOF-sprayed coatings was observed on polished cross-sections using scanning electron microscopes (Quanta-200 and Nova NanoSEM 450 with FEG source: FEI – Thermo Fisher Scientific, Hillsboro, Oregon, USA) equipped with energy-dispersive X-ray (EDX) detectors (INCA: Oxford Instruments Analytical; Quantax-200: Bruker). Powders were cold-mounted in epoxy resin, whereas coated plates were cut with a metallographic saw and hot-mounted in phenol resin. Grinding and polishing were carried out in both cases using SiC papers (up to 2500 mesh size), 3 µm diamond slurry and colloidal silica suspension. Samples were attached to aluminium stubs and their surface was coated with a ≈10 nmthick layer of gold by DC sputtering before SEM observation, in order to ensure proper electrical conductivity and avoid electronic charge accumulation. Thickness and porosity were measured by image analysis (ImageJ: NIH, Bethesda, MA, USA [45]) on SEM micrographs of polished cross-sections: 10 images at 2000× magnification were employed for porosity and 2 images at 200× magnification were employed for thickness. Microhardness was measured at the mid-thickness location on polished cross-sections using a depth-sensing microindenter (Micro-Combi Tester: Anton-Paar TriTec, Peseux, CH) equipped with a Berkovich diamond tip, operating with a maximum applied load of 1 N, holding time of 15 s, and loading/unloading rate of 1.5 N/min. Calibration and operation of the instrument and data analysis were carried out according to the ISO 14577 standard. Results are expressed in terms of indentation hardness (HIT) as defined in ISO 14577-1 and are also converted to Vickers hardness (HV1N), using the geometric relation between the projected and actual contact area of a Berkovich indenter.

2.3

Wear testing

Dry sliding wear tests were performed on the TiC-FeCrAl coatings and on the WC-CoCr and Cr3C2-NiCr references in rotating ball-on-disc configuration (High-Temperature Tribometer: Anton-Paar TriTec). A 6 mm-diameter Al2O3 sphere with average surface roughness Ra < 0.1 μm

was used as counterpart, because this setup is capable of simulating the recurring practical event of sliding wear caused by hard asperities, either embedded in the mating counterpart (e.g. particlefilled polymer gaskets or particle-reinforced metal-matrix composites) or coming from trapped, foreign debris, as explained in [46]. Tests were run at room temperature and at 400 °C. In the latter case, the temperature of the sample, attached to the rotating disc, was measured on its back face with a thermocouple; heating lasted 1 h and an additional waiting time of 40 min ensured homogenisation of the system temperature. All samples were ground and polished to Ra ≈ 0.02 µm using diamond papers and 3 µm-size diamond slurry, and ultrasonically degreased in acetone prior to testing. Test conditions included a relative sliding speed of 0.10 m/s, an overall sliding distance of 5000 m, a wear track radius of 7 mm and a normal load of 10 N, to yield an initial average Hertzian contact pressure of approximately 1.3 GPa. It should be noted that contact pressures decrease during the test, as ball wear proceeds, eventually reaching minima of ≈15 MPa at room temperature and ≈35 MPa at 400 °C. These conditions were chosen to match prior tests by the same authors, in order to allow comparison to an extensive database of reference values [12,47,48]. Two samples were tested for each coating type. The friction coefficient was monitored by measuring the tangential force acting on the counterpart with a load cell. High-stress three-body abrasion tests were performed on as-deposited TiC-FeCrAl and Cr3C2-NiCr coatings using a dry sand-steel wheel apparatus equipped with a 200 mm-diameter, 10 mm-wide Fe360A steel wheel rotating at 85 rpm. FEPA-80 Al2O3 abrasive was fed tangentially at 80.2 g/min while the sample was pressed by a 40.2 N normal load against the wheel. The test lasted 50 laps (≈31 m); 4 abrasion tracks were performed on each sample. In both sliding and abrasive wear tests, the volume loss of the sample was assessed using a scanning, white-light confocal profilometer (Conscan: Stil, Aix-en-Provence, FR) and was converted into specific wear rate by dividing over the sliding distance and the applied normal load.

Moreover, in sliding wear tests, the wear loss of the spherical counterpart was assessed by measuring the diameter of the worn cap using an optical microscope. Worn surface morphologies were observed by SEM; moreover, cross-sections were obtained by metallographic cutting of vacuum cold-mounted worn samples (in order to avoid any cuttinginduced pull-outs), which were then ground and polished using SiC papers, diamond slurry and colloidal silica suspension as mentioned above. Micro-Raman spectra were acquired on surfaces and wear debris using a 532.81 nm-wavelength solid-state laser (LabRam spectrometer: Jobin-Yvon, Longjumeau, FR) focused through a 100× objective.

2.4

Electrochemical corrosion testing

Electrochemical polarisation tests were performed on the TiC-FeCrAl coatings and on the Cr3C2NiCr and WC-CoCr references using a three-electrode cell (K0235 flat cell: Ametek Princeton Applied Research) filled with 300 mL of a 0.1 M HCl aerated aqueous solution at room temperature. The exposed surface of the sample (ground and polished as described in Section 2.3) was 1 cm2; the counter-electrode was a Pt mesh and the reference electrode was an Ag/AgCl/KCl(3M) electrode. The tests started after 1 h of free corrosion, to ensure sufficient stabilization of the open-circuit potential (OCP); a -400 mV / +1400 mV overpotential range was scanned at a 0.5 mV/s range. Corrosion current density and corrosion potential were extracted by Tafel analysis [49]. Two samples were tested for each coating type. Corroded surfaces and crosssections (prepared as for worn samples, see Section 2.3) were observed by SEM.

3. Results and discussion 3.1

Characterization of the feedstock powder

Figure 1. SEM micrographs of the high-energy ball milled TiC-FeCrAl feedstock powder: overview (A) and high magnification details of particles containing different amounts of TiC grains of micrometric (B), sub-micrometric (C) and nanometric (D) size.

Powder particles are dense with somewhat irregular but, in most cases, rather equiaxed morphology (Figure 1A); only few of them have an elongated shape. As a result, the powder was experimentally found to produce a stable, homogeneous flow in the HVOF feeding system. The particle size distribution (Figure 2), characterized by d10 = 10 µm, d50 = 22 µm, d90 = 42 µm, is also quite consistent with the sieving and classification operation mentioned in Section 2.1 and the morphology and flowability are suitable for gas-fuelled HVOF processes [50].

Figure 2. Particle size distribution of the high-energy ball milled TiC-FeCrAl feedstock powder, as determined by laser diffraction.

In detail, the hard phase exhibits a quite wide range of grain sizes, from micrometric (Figure 1B) to sub-micrometric (Figure 1C) and nanometric ones (Figure 1D), which might result from a somewhat variable intensity of the milling action on distinct particles. In any case, the hard phase consists solely of TiC, as revealed by both EDX analyses (Figure 3: spctr 2) and XRD patterns (Figure 4). The matrix is an α-Fe (b.c.c.) solid solution (Figure 4) containing Fe, Al and Cr (Figure 3: spctr 1), which means that Ti and C reacted extensively during the milling process, leaving no free Ti in the matrix within the detection limits of the EDX technique. In the overall spectrum acquired on the powder, no impurities are detected within measurable amounts (Figure 3: overall spectrum): the strong O and C signals originate from the surrounding resin, whereas the Au peak present in all spectra comes from the sputtered gold layer deposited onto the surface of resinmounted samples as described in Section 2.2.

Figure 3. EDX spectra of the feedstock powder, acquired by FEG-SEM using an electron beam energy of 5 keV: “Spctr 1” and “Spctr 2” were acquired on the locations marked in Figure 1B, the “Overall” spectrum was obtained by a large-area scan over a low magnification image such as that seen in Figure 1A.

The thermogravimetric analysis (Figure 5) shows that the powder starts oxidizing perceivably at ≈350 °C. At the 400 °C, the temperature chosen for the high-temperature sliding wear test, the mass gain is still very limited. Mass gain continues up to ≈1050 °C, with an overall increase of approximately 15%, which also suggests that the material is quite resistant to oxidation, at least within the studied temperature range. The corresponding DTA signal is characterized by a broad but rather weak peak, indicative of a limited oxidation rate.

Figure 4. XRD patterns of the TiC-FeCrAl feedstock powder and of the two HVOF-sprayed coatings.

Figure 5. Simultaneous TG/DTA analysis of the TiC-FeCrAl feedstock powder.

3.2

Microstructure of the coatings

Figure 6. SEM micrographs of the HVOF-sprayed coatings deposited using the first (A,C,E) and second (B,D,F) set of process parameters listed in Table 1: low-magnification overviews (A,B), intermediate- (C,D) and high-magnification (E,F) details.

The coatings are macroscopically dense and free of major defects (Figure 6A,B); accordingly, their porosity, as determined by image analysis (Table 2) is <1 vol.%. Thicknesses, however, differ greatly (Figure 6A,B and Table 2), which, given that both coatings were deposited with an identical number of torch cycles, indicates a significantly different deposition efficiency, as confirmed by the

corresponding, quantitative measurement (Table 2). Specifically, spraying with high gas flow rates and reducing flame stoichiometry resulted in remarkably low deposition efficiency of about 33%, which may tentatively be ascribed to frequent rebounding of particles. Hardmetal powders usually employed for HVOF spraying, obtained by such techniques as agglomeration and sintering [50], possess intrinsic porosity which has been shown i.a. by Kamins et al. [51] to favour deformation upon impact in a semi-solid condition. By contrast, the dense particles obtained by mechanical milling (Section 3.1 and Figure 1A) presumably have lower deformability; hence, they are likely to bounce off when their impact velocity is too high. Adjusting the process parameters by lowering the gas flow rates accordingly resulted in a satisfactory deposition efficiency of about 54%, comparable to the values reported for gas-fuelled HVOF spraying of commercially available hardmetal feedstock [52]. On the other hand, no major microstructural differences appear between samples obtained using the two parameter sets. Detailed views (Figure 6C-F) show that the wide range of TiC grain sizes existing in the feedstock (Section 3.1) was retained in the coatings, which means that most particles were not excessively heated, so that extensive dissolution (particularly of the nanometric grains) did not occur. In some lamellae, however, TiC grains did disappear in the matrix (see for instance Figure 6D), which is a consequence of the stochastic nature of the HVOF spraying process, where distinct particles are likely to experience different degrees of heating and melting. XRD patterns (Figure 4) accordingly indicate the formation of small amounts of decarburized TiC1-x and of Fe2C, the latter resulting from the interaction of the matrix with some dissolved TiC.

Table 2: deposition efficiency, porosity and instrumented hardness for the tested TiC-FeCrAl coatings

TiC-FeCrAl: Set 1 TiC-FeCrAl: Set 2

Deposition efficiency (%)

Porosity (vol. %)

HIT (GPa)

HV1N

Thickness (µm)

33.1

0.82 ± 0.29

13.8 ± 1.0

1281 ± 99

307 ± 26

54.2

0.96 ± 0.32

12.2 ± 1.4

1129 ± 131

520 ± 9

Figure 7. EDX spectra acquired at the locations marked in Figure 6E using an electron beam energy of 5 keV.

Many lamellae exhibit oxide stringers rich in Al and Ti along their boundary (e.g. see Figure 6D – spctr. 2; corresponding EDX spectrum in Figure 7). Al was accordingly depleted from the matrix, as can be seen by comparing the rather weak Al peak in the corresponding EDX spectrum (Figure 7: spctr. 4) to the more intense peak appearing in the matrix phase of the feedstock powder (Figure 3: spctr. 1). Such oxidation did not occur during the flight stage, as this would have also produced intra-lamellar oxide inclusions through turbulent flow within the melt [53] and would have involved all of the alloy elements, including Fe and Cr, which are instead very scarce in the present oxide stringers. To the contrary, oxidation mainly occurred in the solid state, immediately after deposition, when the lamella was still hot enough to react with the surrounding environment, according to the mechanism laid out in [53]. The TG/DTA analyses in Section 3.1 indicated that the material is quite resistant to oxidation up to 1300 °C; hence, it is likely than such oxidation occurred while the lamellae were at higher temperature, closer to the matrix solidification point, when the

oxidation rate was likely much larger. In these conditions, the more oxidizable constituents (Al and Ti) diffused towards the surface of the lamella to generate an oxide scale. In a few lamellae, however, oxidation was more extensive: the whole matrix phase was replaced by a dark-contrast oxide (Figure 6D: spctr. 3) containing Fe (Figure 7: spctr. 3) along with Al, Ti and, possibly, Cr (whose EDX signal is covered by that of oxygen). These lamellae probably resulted from overheating and in-flight oxidation of particles that were especially rich in TiC grains.

The coatings exhibit high indentation hardness in excess of 12 GPa, corresponding to a Vickers hardness above 1100 HV0.1 (Table 2), which far exceeds the typical values for Cr3C2-NiCr coatings [12] and is comparable to those for WC-CoCr ones [47]. However, if the test load is increased from 1 N (Section 2.2) to 3 N (≈ 300 gf), the usual value for hardness measurement on thermal spray coatings [54], the material is frequently uplifted around indentation marks (Figure 8). This, together with the presence of microcracks within individual TiC grains (Figure 6F) or even across entire lamellae (Figure 6D), suggests a certain brittleness of the coating material at both the intra-and inter-lamellar level.

Figure 8. SEM micrograph showing a Berkovich indentation mark produced at an applied load of 3 N on the polished cross-section of the “Set 1” coating.

3.3

Dry sliding wear behaviour

3.3.1 Wear behaviour at room temperature

Figure 9. Results of ball-on-disc dry sliding wear tests performed on the TiC-FeCrAl coatings and on the reference materials: wear rates of coatings (A) and counterparts (B), and friction coefficients (C).

Figure 10. SEM micrographs of TiC-FeCrAl coatings after ball-on-disc testing at room temperature: worn surface of sample “Set-2” (A,B), cross-sectional overview of sample “Set-1” (C) with nearsurface delaminations (circled), and details of the same (D-F). Labels 1 = shallow abrasive grooves; 2 = pits (partly filled with debris); 3 = surface microcracks; 4 = intra-lamellar sub-surface microcracks; 5 = interlamellar sub-surface microcracks; 6 = carbide pull-out. At room temperature, the TiC-FeCrAl coatings exhibit slightly lower wear rates (≈2.5×10-6 mm3/(N∙m): Figure 9A) than the Cr3C2-NiCr reference (3.7×10-6 mm3/(N∙m)) and also inflict slightly lower wear to the Al2O3 counterpart (Figure 9B), although they produce higher friction (Figure 9C). Other Cr3C2-NiCr coatings tested in [12] under identical conditions also exhibited the

same wear rates (≈4×10-6 mm3/(N·m)) as the present reference sample, corroborating to the conclusion that the TiC-FeCrAl samples have somewhat better sliding wear resistance. Their performance is similar to that of a HVOF-sprayed FeVCrC alloy (rich in VC hard phase) tested in [55] (≈2×10-6 mm3/(N·m)), and much superior to that of electroplated hard chromium, which, under similar [56] or identical [47] testing conditions returned wear rates of ≈1×10-4 mm3/(N·m). The WC-CoCr reference, on the other hand, significantly outperforms the TiC-FeCrAl coatings: its wear rate is lower by more than one order of magnitude; friction coefficient and counterpart wear are also slightly lower (Figure 9). On the TiC-FeCrAl coatings, shallow abrasive grooves are produced by contact with the hard asperities on the counterpart and/or with wear debris (Figure 10A,B: label 1). The very small groove size is consistent with the high hardness of the coatings (Section 3.2 and Table 2). The wear mechanisms of Cr3C2-NiCr coatings under the same test conditions (which have been studied in great detail in [12]; hence, the analysis need not to be repeated here) involve comparatively more severe abrasive grooving, which is consistent with their lower hardness and explains the difference in wear rates. Had grooving been the only wear mechanism of the TiC-FeCrAl coatings, their wear rates would have likely been much lower than are shown in Figure 9A. However, cracks (Figure 10B: label 3) also appear, which lead to frequent near-surface delamination (Figure 10A,B: label 2). Delamination pits and microcracks are very shallow (Figure 10C – circles) and do not extend down along the section. In detail, delamination cracks immediately below the surface run both between lamellae, mostly following oxide inclusions (Figure 10D,E – label 5), and inside lamellae (Figure 10D – label 4). In addition, occasional pull-out of carbide particles occurs (Figure 10F – label 6), albeit to a very minor extent. Cracking and delaminations are probably produced through a surface fatigue process, due to the cyclic contact stresses applied at every disc revolution. Nucleation and propagation of fatigue cracks are most likely favoured by the presence of interlamellar oxide inclusions and by the overall brittleness of the coatings, both at the inter- and intra-lamellar levels, as reported in Section 3.2.

Figure 11. Friction curve recorded during a room-temperature ball-on-disc test on the TiC-FeCrAl “Set 1” sample. The inset shows a detail of the initial stages of the test.

Detailed analysis of the friction curves (Figure 11) shows that, in the first few tens of metres, friction tends to stabilize at a rather low value of ≈0.27, which corresponds to the incubation stage where fatigue damage has not yet led to the formation of pits. At this stage, the high hardness of the coating (Section 3.2) minimizes both the adhesive and abrasive contributions to friction; indeed, the surface morphology obtained by stopping a wear test after a sliding distance of 50 m is free of any sign of plastic shearing and confirms the mentioned shallowness of abrasive grooves (Figure 12).

Figure 12. SEM micrograph of the worn surface of the TiC-FeCrAl “Set 1” coating after a roomtemperature ball-on-disc test over a sliding distance of 50 m.

Once surface delamination begins, wear debris is released in the contact area and forms clusters on the surfaces of the coating (especially filling in part of the pits themselves: Figure 10A,B) and of the counterpart (Figure 13). Interlocking between the clusters causes friction to rise, as seen in Figure 11. Friction is also unstable as clusters are continuously removed and re-formed, and as new pits open. Friction minima probably correspond to instants when debris clusters are detached from the mating surfaces.

Figure 13. Optical micrograph of the worn surface of the Al2O3 counterbody after ball-on-disc testing at room temperature against the TiC-FeCrAl “Set 2” coating.

Notably, the appearance of debris in the contact area reportedly caused a similar increase in the friction coefficient (up to ≈0.8) of a (Ti,Ta)C-NiCr coating sliding against an alumina ball [41]. Numerous studies have indeed recently confirmed that, during sliding contacts involving thermal spray hardmetal and/or metal alloy coatings, the formation of wear debris is related to increased friction if such debris consists of poorly crystalline or amorphous oxides [47,57–59]. EDX (Figure 14A) and Raman (Figure 14B) spectra confirm that this is the case for the present debris. Namely, the debris contains significant amounts of O (Figure 14A) together with all of the coating constituents (Ti, Fe, Cr, Al), which suggests complete oxidation of coating fragments. Moreover, the very broad peaks of its Raman spectrum (Figure 14B) reveal poor crystallinity. By comparison with open Raman databases [60], the presence of anatase can be confirmed primarily through its

characteristic peak at about 150 cm-1; other peaks, though difficult to identify due to their breadth and low intensity, suggest the presence of rutile, of spinel oxides (such as magnetite and/or chromite) and Fe aluminate [51], and of titanate compounds [62]. Poorly crystalline oxides present no shear plane that would enable solid lubrication activity [57,58]; by contrast, it can be hypothesized that interaction with environmental humidity can favour a sticky behaviour.

Figure 14. (A) EDX spectrum acquired at 12 keV electron beam energy on the oxide clusters seen in Figure 10A,B, and (B) micro-Raman spectrum acquired on loose wear debris collected on the sample surface after the ball-on-disc wear tests at room temperature and at 400 °C.

3.3.2 Wear behaviour at 400 °C At 400 °C, the wear rates of the TiC-FeCrAl coatings increase to ≈1×10-5 mm3/(N·m) (Figure 9A). SEM micrographs of their worn surface (Figure 15A) and cross-section (Figure 15C,D) accordingly show that abrasive grooving has become more severe than it was at room temperature (compare to Figure 10). Grooves are indeed tens of micrometres wide (Figure 15A) and some micrometres deep (Figure 15D). On the other hand, the occurrence of brittle fracture is reduced, as revealed by magnified views of the top surface (Figure 15B) and of the cross-section: very few sub-surface cracks are indeed seen in Figure 15D. This presumably contributed to avoiding an even greater increase of the wear rate at 400 °C. Both the increased severity of abrasive grooving and the reduced incidence of brittle fracture suggest that the material experienced thermal softening at 400 °C.

Figure 15. SEM micrographs of the TiC-FeCrAl “Set 2” coating after ball-on-disc testing at 400 °C: top surface (A: secondary electrons overview; B: backscattered electrons detail) and cross-section (A: overview, B: detail). Arrows in panel B indicate some oxidized clusters.

Figure 16. Backscattered electron SEM micrographs showing details of the cross-section of the oxide scale developed onto the unworn surface of the TiC-FeCrAl “Set 2” coating after ball-on-disc testing at 400 °C (A-C), and corresponding EDX spectra acquired with an electron beam acceleration voltage of 10 kV (D).

Figure 17. (A) secondary electron SEM micrograph of the unworn, oxidized surface of the TiCFeCrAl “Set 2” coating after ball-on-disc testing at 400 °C, and (B) representative Raman spectra acquired on protrusions (1) and smooth regions (2). In panel B: * = rutile; all other peaks = anatase [63].

Detailed SEM micrographs also show that a thin oxide scale (Figure 16A,B: see arrows) formed outside the wear track, after the TiC-FeCrAl coatings had been kept at 400 °C for the overall duration of the ball-on-disc test (≈14.5 h, as inferable from the test parameters listed in Section 2.3). The scale is rich in titanium (EDX spectrum 1 in Figure 16), which is indeed the most oxidizable element in the coating. The limited thickness (<500 nm) of the scale indicates a very low oxidation rate at 400 °C, consistent with the TG/DTA analyses of the feedstock powder discussed in Section 3.1. It is therefore inferred that the scale was readily removed at the beginning of the wear test and could not be re-formed between successive passes of the counterbody, thus playing no further role on the wear process. However, the thickness of the scale is not fully homogeneous, as seen e.g. by comparing the ≈50 nm-thick scale in Figure 16A to the ≈400 nm-thick one in Figure 16B. More specifically, top surface morphologies suggest a slightly thicker oxide scale on TiC grains (Figure 17A), although its composition does not change: the scale consists in all cases of a mixture of anatase and rutile (Figure 17B: peaks marked by an asterisk are assigned to rutile, all other peaks to anatase, according to [63]). During the earliest stages of the wear test, the oxidized TiC grains located on the surface presumably have reduced load-carrying ability, which might favour the initial onset of the coarse abrasive grooves seen in Figure 15A.

The difference between the wear rates of the TiC-FeCrAl coatings and that of the Cr3C2-NiCr reference (≈4.5×10-5 mm3/(N·m), Figure 9A) has anyway become even larger than it was at room temperature. This value is consistent with the wear rates of other Cr3C2-NiCr coatings tested in [12], ranging from ≈4.5×10-5 to ≈1×10-4 mm3/(N·m). It has been reported in [12] that Cr3C2-NiCr also experiences quite severe abrasive grooving at 400 °C [12]. It is likely that Cr3C2-NiCr softens more than does the TiC-FeCrAl composition.

The absence of pits on the worn surface of the TiC-FeCrAl coating, and consequently the absence of their highly abrasive, sharp edges, can also explain the decreased wear rate of the counterbody at 400 °C (Figure 9B). Although the worn TiC-FeCrAl surface does show some oxidized debris clusters (Figure 15B – see arrows), whose composition is identical to that observed at room temperature (Figure 14B), no clusters accumulate on the counterbody (Figure 18), minimizing the chance for interlocking as described for the room-temperature tests. The friction coefficient produced by TiC-FeCrAl coatings therefore decreases, compared to the room-temperature values (Figure 9C), though it is still higher than that caused by Cr3C2-NiCr.

Figure 18. Optical micrograph of the worn surface of the Al2O3 counterbody after ball-on-disc testing at 400 °C against the TiC-FeCrAl “Set 2” coating.

WC-CoCr still outperforms all other coatings with orders-of-magnitude lower wear rate (Figure 9A); indeed, the abrasive grooves formed on WC-CoCr, as seen in detail in [47], are much shallower than they are on TiC-FeCrAl, even at 400 °C. However, at this temperature, the WCCoCr reference produces higher friction (Figure 9C) and much higher counterpart damage (Figure 9B), consistent with data in [47], since adhesive wear phenomena were shown to occur and to cause

material transfer to the counterbody surface. Moreover, it should be remarked that the applicability of WC-CoCr coatings onto steel substrates at 400 °C is at least partly impaired by the significant mismatch in thermal expansion coefficients [47]. Under specific conditions, this can indeed result in severe cracking of the coating layer [47].

3.4

High-stress three-body abrasion

Figure 19. SEM overviews of the alumina abrasive before (A) and after (B) dry sand-steel wheel testing of the TiC-FeCrAl “Set-2” coating, and backscattered electron details of fractured coating fragments detectable amidst the debris (C,D).

The abrasive particles experienced considerable fragmentation after the dry sand-steel wheel wear test (Figure 19B), as compared to their original size and shape (Figure 19A). With reference to the distinction between low- and high-load abrasion given in [64], the present test conditions can clearly be classified in the latter category. Note that some of the unbroken abrasive particles in Figure 19B might have not entered the contact zone between the sample and the steel disc.

Figure 20. SEM micrographs (A,B: overviews in secondary and backscattered electrons imaging modes respectively, C: backscattered electrons detail) of the TiC-FeCrAl “Set 2” coating after dry sand-steel wheel abrasion testing. Labels 1 = particle indentation; 2 = abrasive groove; 3 = brittle fracture; 4 = embedded alumina particles.

Some debris particles from the worn TiC-FeCrAl coating can also be detected, revealing extensive fracturing (Figure 19C). Sometimes, coating debris was plastically deformed and mixed with alumina fragments (Figure 19D). This finding corroborates the assumption of high contact stresses

and also suggests that brittle fracture is a significant wear mechanism for TiC-FeCrAl coatings during dry particles' abrasion. Accordingly, top surface views reveal three main wear mechanisms: 1) Particles' indentation (Figure 20A – label 1), sometimes resulting in the embedment of alumina fragments on the coating surface. These fragments are clearly recognizable through their dark contrast in backscattered electron micrographs, both on top surface views (Figure 20B) and on the cross-section (Figure 21E and corresponding EDX spectra in Figure 21G); 2) Abrasive grooving (Figure 20A – label 2), which occurred as abrasive particles slid onto the coating surface; 3) Brittle fracture, recognisable on both surfaces (Figure 20A – label 3 and detail in Figure 20C) and cross-sections (Figure 21A-D). In the centre of the abrasion scar, where greatest reduction in thickness occurred (Figure 21A: see arrows), damage is not limited to the near-surface area: extensive micro-cracking can be seen up to a depth of ≈100 μm. Closer to the top and bottom edges of the scar, cracks affect a shallower region (Figure 21B). It should be noted that the seemingly high porosity of the microcracked areas in Figure 21A,B is probably due to pull-out of unsupported, fractured fragments during polishing, as better seen in the details of Figure 21C,D, since the mounting resin could not impregnate closed sub-surface cracks. More specifically, cracks propagated both along interlamellar boundaries and within individual lamellae (Figure 21C,D), leading, on a larger scale, to the removal of entire lamellae or portions thereof (Figure 20A – label 3). In a few cases, plastically deformed coating debris, mixed with fragmented abrasive particles as previously seen in Figure 19D, could stick back onto the worn surface (Figure 21F and EDX spectrum in Figure 21H). Moreover, at intra-lamellar level, fracturing (Figure 20C – circles) and pull-out (Figure 20C – arrow) of individual TiC grains took place, particularly affecting the largest ones.

Figure 21. Cross-sectional views of the TiC-FeCrAl “Set 2” coating after dry sand-steel wheel abrasion testing: overviews of the centre (A) and side (B) of the abrasion scar (arrow indicate the extension of microcracked areas), details of microcracks next to (C) and below the surface (D), embedded alumina residuals (E) and mixed debris lumps (F), and corresponding EDX spectra (G,H).

Brittle fracture is likely the most severe of the three wear mechanisms: in accordance with previous considerations in Section 3.2, it would seem that both inter-lamellar brittleness, probably promoted

by oxide stringers, and intra-lamellar brittleness are impairing the high-stress abrasive wear resistance of TiC-FeCrAl coatings. It is indeed because of brittle fracture that the TiC-FeCrAl coatings exhibit poorer performance the Cr3C2-NiCr reference under high-stress particle abrasion conditions (Table 3). The mentioned brittleness of the TiC-FeCrAl coatings would therefore seem to exert an even more negative influence under these conditions than it did in room-temperature sliding wear tests (Section 3.3).

Table 3: wear rates measured after dry sand-steel wheel high-stress abrasion test. Sample

TiC-FeCrAl: Set 1

TiC-FeCrAl: Set 2

Cr3C2-NiCr

Wear rate [×10-4 mm3/(N·m)]

41.7 ± 4.0

55.4 ± 9.1

3.7 ± 0.4

3.5

Electrochemical polarisation testing

Figure 22. Corrosion current densities (ICorr) and corrosion potentials (ECorr) obtained by electrochemical polarisation tests.

Among the reference samples, the corrosion current density of Cr3C2-NiCr (Figure 22) is consistent with previous data acquired on similar coatings tested in the same environment [65], whilst that of WC-CoCr is remarkably low, with a correspondingly noble corrosion potential, which suggests a particularly dense layer, completely free of open porosity. In comparison to those values, the corrosion current density of the TiC-FeCrAl coated samples is higher by one or more orders of magnitude (Figure 22). The whole polarization curves (Figure 23) of the TiC-FeCrAl coatings accordingly lie at lower potentials and higher current densities than those of both references. The corrosion current densities of TiC-FeCrAl coatings are also higher than those typical of electroplated hard chromium layers, which usually exhibit ICorr ≈ 10-6 – 10-7 A/cm2 in a 0.1 M HCl aqueous solution [66] when free of through-thickness microcracks.

Figure 23. Electrochemical polarisation curves of a TiC-FeCrAl “Set 2” coating and of the Cr3C2NiCr and WC-CoCr references.

The poorer performance of the HVOF-sprayed TiC-FeCrAl layers seems not to be due to their defectiveness, as no corrosion is visible at the interface with the substrate (Figure 24D), but to an

intrinsically limited corrosion resistance of the material itself, and particularly of the metal matrix, in comparison to the CoCr- and NiCr-based alloy matrices of the reference coatings. Indeed, the matrix seems unable to passivate, as shown by the polarisation curves (Figure 23), and it suffers selective corrosion in lamellae containing the largest amounts of TiC grains. This occurs both on the coating surface (Figure 24A - arrow) and at various depths along the cross-section (Figure 24B,C arrows), approximately affecting half of the coating thickness (although local variability exists). Some paths may indeed have been opened up for the electrolyte to penetrate towards the interior of the coating, though not across its entire thickness. Galvanic coupling between the nobler hard phase and the less noble matrix is presumably maximized in those lamellae. Aluminium depletion from the alloy due to oxidation during the deposition process (as discussed in Section 3.2) can further impair the corrosion resistance of the matrix.

Figure 24. SEM micrographs of the TiC-FeCrAl “Set 2” coating after electrochemical polarisation tests: surface (A); cross-section at different depths below the surface (B,C) and at the substrate interface (D). The arrows indicate selective corrosion of the matrix in TiC-rich lamellae.

4. Conclusions A TiC – 25vol.% (Fe-20wt.%Cr-5wt.%Al) experimental powder was produced by high-energy ballmilling, and classified and sieved to a -38+10 µm size distribution. Particles have irregular but mostly equiaxed shape. They consist of a TiC hard phase in an α-Fe alloy matrix; however, a certain variability is seen in the amount and size of TiC grains. Using a suitable set of process parameters, gas-fuelled HVOF-spraying of this powder yielded dense (<1 vol.% porosity), hard (HIT > 12 GPa) coatings, with a deposition efficiency of approximately 54%. The tribological characterization indicates that the TiC-FeCrAl coatings are particularly promising for dry sliding wear applications. Under ball-on-disc test conditions at room temperature, the TiCFeCrAl coatings' wear rate is indeed somewhat lower than that of HVOF-sprayed Cr3C2-NiCr and much lower than that reported in the literature for electroplated hard chromium layers, although HVOF-sprayed WC-CoCr outperforms the TiC-FeCrAl coatings by more than one order of magnitude. Remarkably, in the early stages of the sliding wear test, the TiC-FeCrAl coatings exhibit low (≈0.27) dry sliding friction coefficient in contact with the Al2O3 counterpart, with no signs of adhesive wear and very limited abrasive grooving. As the sliding distance increases, damage accumulation results in near-surface spallation by brittle fracture. Oxidized wear debris released in the contact region increases friction to more than 0.8. At 400 °C, though the occurrence of brittle fracture is reduced, abrasive grooving of the TiCFeCrAl coatings becomes more severe. Wear rates therefore increase; nonetheless, the TiC-FeCrAl coatings outperform the Cr3C2-NiCr reference even more than they did at room-temperature. WCCoCr still exhibits much lower wear rate, although the applicability of this coating composition at 400 °C is known to be impaired by technical issues. As a downside, TiC-FeCrAl coatings exhibit brittle behaviour, particularly evidenced by cracking and uplifting of the material during Berkovich indentation under a 3 N load. This seems to impair

the response to high-stress three-body abrasion conditions, as compared to a HVOF-sprayed Cr3C2NiCr reference. Both inter- and intra-lamellar brittle fracture seems to contribute to the higher abrasive wear loss. The electrochemical corrosion resistance of the TiC-FeCrAl coatings is somewhat limited, as compared to Cr3C2-NiCr and WC-CoCr references and to literature data for electroplated hard chromium. This is not due to defects in the coatings, but to selective corrosion of the FeCrAl matrix. Overall, the present research has shown that TiC-FeCrAl coatings can be an alternative especially to Cr3C2-NiCr for sliding wear protection at least up to 400 °C, with the advantage of being free from critical raw materials and health-hazardous elements, such as Ni, Co and W. They could be promising for a range of applications in the food processing and packaging industry (for instance, as counterface for reinforced polymer gaskets in rotary seals), in hydraulic and pneumatic pistons (wear protection of sliding faces), in journal bearings (including high-temperature machinery), etc. Future developments shall include the use of different deposition processes, such as High-Velocity Air Fuel (HVAF) spraying, in order to minimize the oxide content in the coatings, as well as refinements to the matrix alloy formulation, in order to overcome its intrinsic brittleness and improve the corrosion resistance.

Acknowledgements The authors are grateful to Dr. Magdalena Lassinantti Gualtieri (University of Modena and Reggio Emilia) for performing the laser-diffraction particle size analysis of the feedstock powder.

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Highlights

A TiC – 25vol.% (Fe-20wt.%Cr-5wt.%Al) powder was produced by high-energy ball-milling

HVOF-sprayed coatings have <1% porosity, HIT > 12 GPa and deposition efficiency ≈54%

The coatings resist ball-on-disk sliding wear better than Cr3C2-NiCr at ≈25 °C and 400 °C

Brittle cracking reduces their resistance to high-stress particle abrasion conditions

Degradation of the FeCrAl matrix limits their electrochemical corrosion resistance