Corrosion behavior of Ni3Al-bonded TiC-based cermets in H2SO4 and NaOH solutions

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Corrosion behavior of Ni3Al-bonded TiC-based cermets in H2SO4 and NaOH solutions ⁎

Qiao Mao, Qingqing Yang , Weihao Xiong, Shengtao Li, Man Zhang, Linji Ruan State Key Laboratory of Material Processing and Die & Mould Technology, Huazhong University of Science and Technology, Wuhan 430074, PR China

A R T I C LE I N FO

A B S T R A C T

Keywords: Ti(C,N)-based cermets Ni3Al Corrosion behavior H2SO4 solution NaOH solution

Corrosion behavior of multicomponent Ni3Al-bonded Ti(C,N)-based cermets with various Ni3Al contents and Ti (C0.5,N0.5) addition was investigated in 1 M H2SO4 and 1 M NaOH solutions, using immersion tests, potentiodynamic polarization measurements, XRD, SEM and XPS. In 1 M H2SO4 solution, corrosion process of all cermets consisted mainly of their binder phase dissolution, and during potentiodynamic polarization, their passivation and pseudo-passivation corresponded to passivation of their binder phase and ceramic grains, respectively. Corrosion potential and current density decreased with increasing Ni3Al content, while corrosion potential increased and current density decreased with Ti(C0.5,N0.5) addition. In 1 M NaOH solution, corrosion process of all cermets consisted mainly of their ceramic grain dissolution, and during potentiodynamic polarization, their passivation and pseudo-passivation corresponded to passivation of their ceramic grains and binder phase, respectively. Corrosion potential increased and current density decreased with increasing Ni3Al content, while corrosion potential decreased and current density increased with Ti(C0.5,N0.5) addition.

1. Introduction TiC- and Ti(C,N)-based cermets are wear-resistant materials, and they have been used as cutting tools for high-speed semi-finishing and finishing of carbon steels and stainless steels, due to their excellent resistance to wear [1,2], plastic deformation [3] and oxidation at high temperature [1,4], together with low density and low cost, compared with commonly-used WC–Co hardmetals. In general, Ni or/and Co is used as binder to improve transverse rupture strength and fracture toughness of TiC- and Ti(C,N)-based cermets. Mo2C (or Mo) or/and WC are often indispensable ingredients in TiC- and Ti(C,N)-based cermets, in order to enhance their wettability and sinterability [5–7]. Besides, the addition of Mo2C (or Mo) and WC can significantly inhibit the growth of TiC and Ti(C,N) ceramic grains during liquid-phase sintering [5,6,8]. In order to improve performance in corrosive environments and high-temperature environments, intermetallic compound Ni3Al is considered as potential binder of TiC- and Ti(C,N)-based cermets, because it possesses many attractive properties, for example, high melting point, superior high-temperature mechanical properties [9], and excellent corrosion and oxidation resistance [10–12]. The wetting angle of Ni3Al on TiC and WC substrates at 1400 °C is ≤ 15° and close to 0 °, respectively [13]. At 1390 °C, the wetting angle of Ni3Al on TiC substrate is about 17°, and it decreases to 11° with the addition of about 5 wt% Mo [14]. Consequently, it is possible for TiC- [9,15–18] and Ti(C,N)-based

⁎

[18–20] cermets to use Ni3Al as binder. So far, Ni3Al-bonded TiC-based cermets have been prepared using different methods, for instance, melt-infiltration and sintering [9,15], self-propagating high-temperature synthesis and sintering [16,17], and liquid-phase sintering [18], whilst Ni3Al-bonded Ti(C,N)-based cermets have been prepared using liquid-phase sintering [18–20]. However, it is worth noting that there are few studies on multi-component Ni3Albonded TiC- and Ti(C,N)-based cermets [19,20]. Stewart et al. [19] investigated the effect of Mo2C addition on densification and mechanical properties of as-vacuum-sintered Ti(C0.7,N0.3)–xNi3Al (x = 20, 30, 40 vol%) cermets, as well as sliding wear resistance against WC–6Co (wt%) hardmetals. Huang et al. [20] investigated the effects of Ni3Al content and sintering temperature on microstructure and mechanical properties of as-vacuum-sintered TiC–10TiN–10WC–5Mo–0.8C–xNi3Al (x = 15, 20, 25, 30 wt%) cermets. On the other hand, there are few studies on corrosion behavior of Ni3Al-bonded TiC- and Ti(C,N)-based cermets [21,22]. Memarrashidi et al. [21,22] found that at 20–30 vol% Ni3Al content, Ti(C,N)–Ni3Al cermets exhibited better corrosion resistance in 3.5 wt% NaCl solution than TiC–Ni3Al cermets, due to significant refinement of ceramic grains. Therefore, in the present work, Ni3Al powders are synthesized by mechanical alloying, and then five multicomponent Ni3Al-bonded TiC-based cermets containing 20 wt% WC and 10 wt% Mo are prepared by liquid-phase sintering in vacuum, in order to investigate corrosion behavior of multicomponent Ni3Al-

Corresponding author. E-mail addresses: [email protected], [email protected] (Q. Yang).

https://doi.org/10.1016/j.ceramint.2018.04.161 Received 20 March 2018; Received in revised form 17 April 2018; Accepted 18 April 2018 0272-8842/ © 2018 Published by Elsevier Ltd.

Please cite this article as: Mao, Q., Ceramics International (2018), https://doi.org/10.1016/j.ceramint.2018.04.161

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Table 1 Nominal composition of five experimental cermets. Cermets

Weight fraction (wt%) TiC

A B C D C1

49.4 44.4 41.4 39.4 34.4

Volume fraction (vol%)

Ti(C0.5,N0.5)

WC

Mo

C

Ni3Al

TiC

7.0

20.0 20.0 20.0 20.0 20.0

10.0 10.0 10.0 10.0 10.0

0.6 0.6 0.6 0.6 0.6

20.0 25.0 28.0 30.0 28.0

65.58 60.31 57.03 54.79 47.50

Ti(C0.5,N0.5)

WC

Mo

C

Ni3Al

9.44

8.37 8.57 8.69 8.77 8.71

6.42 6.57 6.66 6.72 6.67

2.18 2.23 2.26 2.28 2.27

17.45 22.32 25.35 27.43 25.41

2.3. Microstructure characterization

bonded TiC-based cermets with various Ni3Al contents and Ti(C0.5,N0.5) addition in 1 M H2SO4 and 1 M NaOH solutions at room temperature.

Phase identification was carried out using an X′Pert PRO X-ray diffractometer (XRD, PANalytical BV, Netherlands) with Cu Kα radiation (40 kV, 40 mA). Microstructure observation was carried out using a Quanta 200 scanning electron microscope (SEM, FEI Inc., Netherlands) in backscattered electron (BSE) mode or secondary electron (SE) mode.

2. Experimental 2.1. Materials and sample preparation Table 1 lists nominal composition of five experimental multicomponent Ni3Al-bonded TiC-based cermets containing 20 wt% WC and 10 wt% Mo in the present work. Ni3Al powders, i.e. Ni76–Al24 (at %) powders, were synthesized using mechanical alloying by dry high energy ball milling of Ni and Al powders at a ball-to-power weight ratio of 10: 1 and a speed of 350 rpm for 20 h. Available commercial powders of Ni (FSSS, ~ 2.25 µm), Al (~ 80 µm), TiC (~ 2.88 µm), Ti(C0.5,N0.5) (~ 1.50 µm), WC (~ 0.82 µm), Mo (~ 2.80 µm), C (< 30 µm) were used as starting materials. The powder mixtures were milled with alcohol media in a planetary mill at a ball-to-power weight ratio of 7: 1 and a speed of 220 rpm for 48 h. Nylon jars and WC–Co hardmetal balls were used as milling balls. After drying at 80 °C for 12 h and sieving with 100 mesh sieve, the powder mixtures were pressed into green compacts under a uniaxial pressure of 300 MPa for 1 min. The green compacts were liquid-phase sintered in vacuum (about 10−1–10−2 Pa) at 1480 °C for 1 h. After liquid-phase sintering, three-point transverse rupture strength (span 14.5 mm) and hardness were measured at room temperature using a universal testing machine (Zwick/Roll Z020, Ulm, Germany) and a Rockwell hardness tester (HR-150A, Laizhou Huayin Testing Instrument Co., Ltd., China), respectively. Specimens for potentiodynamic polarization measurements were jointed to a conductive copper wire using brazing, and then were sealed with epoxy resin remaining an exposed area of 35 mm2. Prior to immersion tests and potentiodynamic polarization measurements, all specimens were ground and then polished with 1 µm diamond, followed by cleaning with acetone and high-purity deionized water in an ultrasonic bath, and drying in hot air.

2.4. XPS analysis Surface chemical state analysis was carried out using an ESCALAB 250Xi X-ray photoelectron spectroscope (XPS, Thermo Fisher, USA) with monochromatic Al kα X-ray (1486.6 eV) radiation at a power source of 300 W and a pass energy of 25 eV. The binding energy shift due to surface charging was corrected using C 1s peak at 285.0 eV as an internal standard. All high-resolution XPS spectra were fitted using XPSPEAK 4.1 software with Gaussian-Lorentzian line shape (Ti 2p1/2 peaks, 40% Gaussian and 60% Lorentzian; the other, 70% Gaussian and 30% Lorentzian) after subtracting Shirley-type background. 3. Results and discussion 3.1. Microstructure and mechanical properties of materials Fig. 1 shows XRD patterns of Ti(C0.5,N0.5)-free cermets A, B, C, and D, and Ti(C0.5,N0.5)-containing cermet C1, as well as of pure Ni3Al

2.2. Corrosion tests Immersion tests were carried out in 1 M H2SO4 and 1 M NaOH solutions at room temperature (25 ± 1 °C) for 60 h, according to ASTM G31-72 (2004). The ratio of solution volume to specimen surface area was 0.5 ml: 1 mm2. Potentiodynamic polarization measurements were carried out at room temperature (25 ± 1 °C) in a three-electrode electrochemical glass cell containing 500 ml of 1 M H2SO4 or 1 M NaOH solution, which was connected with a CS350 electrochemical workstation (Wuhan CorrTest Instrument Co., Ltd., China). Cermet specimen, platinum electrode and saturated calomel electrode (SCE) were used as working electrode, counter electrode and reference electrode of the three-electrode electrochemical system, respectively. After 1 h immersion for open circuit potential (OCP) to reach a stable value, potentiodynamic polarization was performed from −0.4 V to + 1.8 V with respect to OCP at a scan rate of 1 mV/s. Each measurement was repeated at least two times under the same condition to ensure reproducibility of the results.

Fig. 1. (a) XRD patterns of experimental cermets. (b) (111) diffraction peak of Ni3Al-based binder phase in experimental cermets and pure Ni3Al (JCPDS file No. 32-1383). 2

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Fig. 2. SEM-BSE images of typical microstructure of experiment cermets A (a), B (b), C (c), C1 (d).

(JCPDS file No. 50-1265). Diffraction peaks of TiC-based carbide or carbonitride ceramic phase with B1-NaCl type structure and Ni3Albased binder phase with L12-Cu3Au type structure were present in XRD patterns of all cermets. On the whole, diffraction peaks of ceramic phase for cermets A and B were similar in position, and those for the other three cermets were similar in position, whose was on the higher angle side of those for cermets A and B. (111) diffraction peak of Ni3Al binder phase in all cermets shifted towards lower angles, compared with that of pure Ni3Al (Fig. 1(b)). To be specific, (111) diffraction peak of Ni3Al binder phase for Ti(C0.5,N0.5)-free cermets shifted further towards lower angles with increasing Ni3Al content, and that for Ti (C0.5,N0.5)-containing cermet C1 shifted less towards lower angles, compared with that for all Ti(C0.5,N0.5)-free cermets. Atomic radii of Ti, Ni, Al, Mo, W, C and N are 0.147, 0.125, 0.143, 0.136, 0.137, 0.077 and 0.071 nm, respectively [23]. Consequently, that diffraction peaks of ceramic phase shifted farther towards higher angle indicates its higher alloying content (W, Mo and Al) [24], and that diffraction peaks of binder phase shifted farther towards lower angle indicates its higher alloying content (Ti, W and Mo). Fig. 2 shows SEM-BSE images of typical microstructure of Ti (C0.5,N0.5)-free cermets A, B and C, and Ti(C0.5,N0.5)-containing cermet C1. Microstructure of as-sintered Ti(C0.5,N0.5)-free cermet D was very similar to that of as-sintered Ti(C0.5,N0.5)-free cermet C, thus it was omitted in the present paper. TiC-based carbide or carbonitride ceramic grains were embedded in Ni3Al binder phase, and moreover, ceramic grains in cermet C1 were smallest in size, indicating that the addition of Ti(C0.5,N0.5) inhibited ceramic grain growth during liquid-phase

sintering. In all cermets, ceramic grains generally exhibited the corerim structure, consisting of black core, white inner rim and grey outer rim, or of white core and grey rim. For cermet A, white inner rims were very thin and incomplete in shape, whilst grey outer rims were thick, and for cermet C1, white inner rims were thick and incomplete in shape, and white cores increased in number. For the other three cermets, white inner rims were thick, whilst grey outer rims were very thin. BSE-SEM images show atomic number contrast (Z-contrast), since heavy atoms scatter more electrons than light atoms. Consequently, white inner rims had higher content of Mo and W than grey outer rims (Z: Ti, 22; Mo, 42; W, 74), and black cores were basically free of Mo and W [24]. White inner rims were expected to form by diffusion at solidstate sintering stage, and grey outer rims formed by dissolution-reprecipitation at liquid-phase sintering and cooling stages [25,26]. In addition, for cermet A, there were many agglomerate white particles. These white particles were undissolved WC particles during liquidphase sintering, perhaps including undissolved Mo2C particles (Mo and C reacted to form Mo2C at high-temperature solid-state sintering stage [27]). This was in good agreement with XRD results (Fig. 1), indicating that alloying content (Mo and W) in ceramic grains and Ni3Al binder phase in cermet A was lower than that in cermets B, C and D. Fig. 3 shows transverse rupture strength (TRS) and hardness of Ti (C0.5,N0.5)-free cermets A, B, C and D, and Ti(C0.5,N0.5)-containing cermet C1. Obviously, for Ti(C0.5,N0.5)-free cermets, TRS increased whilst hardness decreased slightly with increasing Ni3Al content. TRS and hardness were about 1353 MPa and 89.8 HRA for cermet C containing 28.0 wt% Ni3Al, and 1458 MPa and 89.2 HRA for cermet D 3

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respectively. Fig. 6(a) shows the whole XPS spectra obtained from surfaces of Ti (C0.5,N0.5)-free cermet C and Ti(C0.5,N0.5)-containing cermet C1 with the same Ni3Al content before and after 60 h immersion in 1 M H2SO4 and 1 M NaOH solutions at room temperature without sputtering, and Fig. 6(b) to (f), in turn, show experimental and fitted high-resolution Ni 2p3/2, Al 2p, Ti 2p, Ti 3p and W 4f, Mo 3d spectra. Note that the intensity ratio was 2: 1 for Ti 2p3/2 to Ti 2p1/2, 3: 2 for Mo 3d5/2 to Mo 3d3/2, and 4: 3 for W 4f7/2 to W 4f5/2, respectively [28], and the intensity ratio was fixed to 0.12 for Ti 3p to Ti 2p [29]. Table 2 lists binding energies of all deconvoluted XPS peaks. Before and after 60 h immersion in both solutions, Ni 2p3/2 XPS spectra of both cermets exhibited a complex structure with intense satellite signals (Fig. 6(b)). Prior to immersion, there were two main peaks in Ni 2p3/2 spectra of both cermets, corresponding to Ni° [30–33] in Ni3Al binder phase of cermet substrate and Ni(OH)2 [30,32,34], and after 60 h immersion in 1 M H2SO4 solution, there were two main peaks corresponding to Ni° [30–33] and Ni(OH)x(SO4)y [31,35,36], and after 60 h immersion in 1 M NaOH solution, there were three main peaks corresponding to Ni° [30–33] in Ni3Al binder phase, Ni(OH)2 [30,32,34] and NiOOH [32–34]. Before and after immersion, there were three peaks in Al 2p spectra of both cermets, corresponding to Al° [37–39] in Ni3Al binder phase of cermet substrate, AlOx [37,39,40] and Al2O3/Al(OH)3 [39–41] (Fig. 6(c)). There were three 2p3/2–2p1/2 spin-split doublets in Ti 2p spectra for both cermets, corresponding to Ti–(C,N) bonds [42–44] in ceramic grains of cermet substrate, Ti2O3 [43,45–47], and TiO2 [43,46–48] (Fig. 6(d)). There were three peaks in Ti 3p spectra for both cermets, corresponding to Ti–(C,N) bonds [42], Ti2O3 [47] and TiO2 [48] (Fig. 6(e)). There were two 4f7/2–4f5/2 spin-split doublets in W 4f spectra of both cermets, corresponding to W–(C,N) bonds [49–51] in ceramic grains and WO3 [51–53] (Fig. 6(e)). There were four 3d5/ 2–3d3/2 spin-split doublets in Mo 3d spectra of both cermets, corresponding to Mo–(C,N) bonds [54–56] in ceramic grains, MoO2 [57–59], MoO(OH)2 [57,58,60] and MoO3 [57–59] (Fig. 6(f)). In addition, some amount of TiO and Mo2O3 might be present on the surface of both cermets, since the first Ti 2p3/2–2p1/2 spin-split doublet (454.85–455.16 and 460.75–461.06 eV) and the first Mo 3d5/2–3d3/2 spin-split doublet (228.49–228.69 and 231.69–231.89 eV) was consistent in position with TiO [45–47], and Mo2O3 [59,60], respectively. The oxide films were very thin on the surface of cermets C and C1 before and after 60 h immersion in 1 M H2SO4 and 1 M NaOH solutions at room temperature, since cermet substrate (Ni° and Al° in Ni3Al binder phase, and (Ti,W,Mo)–C or (Ti,W,Mo)–(C,N) bonds in ceramic grains) were clearly detected by XPS (Fig. 6(b) to (f)). Based on XPS results,

Fig. 3. Transverse rupture strength (TRS) and hardness of experimental cermets.

containing 30.0 wt% Ni3Al, respectively. TRS and hardness of Ti (C0.5,N0.5)-containing cermet C1 was lower about 235 MPa and 0.1 HRA than those of Ti(C0.5,N0.5)-free cermet C with the same Ni3Al content, respectively. High TRS of cermets C and D was expected to be mainly attributed to the synergistic effect of the following two factors: (1) higher content of Ni3Al binder phase; (2) the good wettability of Ni3Al binder phase on ceramic grains, resulting from higher alloying content (Mo and W) in Ni3Al binder phase and ceramic grains, according to XRD results and SEM observation (Figs. 1 and 2).

3.2. Immersion corrosion behavior After 60 h immersion in 1 M H2SO4 solution at room temperature, for Ti(C0.5,N0.5)-free cermets A, B, C and D, and Ti(C0.5,N0.5)-containing cermet C1, Ni3Al binder phase was preferentially dissolved, thus leaving TiC-based carbide or carbonitride ceramic skeleton in the surface layer. Fig. 4(a), (b) and (c) show surface SEM-SE images of cermets A, C and C1 after 60 h immersion in 1 M H2SO4 solution, respectively. After 60 h immersion in 1 M NaOH solution at room temperature, no significant dissolution was observed for Ti(C0.5,N0.5)-free cermets A, B, C and D, and Ti(C0.5,N0.5)-containing cermet C1, except that some local areas were attacked. Fig. 5(a), (b) and (c) show surface SEM-SE images of cermets A, C and C1 after 60 h immersion in 1 M NaOH solution,

Fig. 4. Surface SEM-SE images of experimental cermets A (a), C (b) and C1 (c) after 60 h immersion in 1 M H2SO4 solution at room temperature. 4

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Fig. 5. Surface SEM-SE images of experimental cermets A (a), C (b) and C1 (c) after 60 h immersion in 1 M NaOH solution at room temperature.

prior to immersion, the thin oxide films on the surface of both cermets consisted of Ni(OH)2, AlOx, Al2O3/Al(OH)3, Ti2O3, TiO2, WO3, MoO2, MoO(OH)2 and MoO3, perhaps including TiO and Mo2O3. After 60 h immersion in 1 M H2SO4 solution, there was not Ni(OH)2 but Ni (OH)x(SO4)y in the thin oxide films, and after 60 h immersion in 1 NaOH solution, there was extra NiOOH in the thin oxide films, compared with that prior to immersion. Prior to immersion, there were no other nickel oxides besides Ni (OH)2 on the surface of cermets C and C1 (Fig. 6(b)), which was in agreement with 718 Ni-based alloy [61] and Co–Ni–Cr–Mo alloy [62], and there were AlOx and Al2O3/Al(OH)3 (Fig. 6(c)), which was in agreement with pure Al [39] and Ni3Al foil [40]. During immersion in 1 M H2SO4 solution, Ni(OH)x(SO4)y formed on the surface of both cermets. Nakamura [63] et al. detected NiSO4-like species on the surface of as-active Ni electrode in pH 3 H2SO4 solution, using in situ infrared reflection adsorption spectroscopy (IRAS). Al2O3/Al(OH)3 content increased on the surface of both cermets, compared with that prior to immersion. However, the oxide films containing Ni(OH)x(SO4)y, AlOx and Al2O3/Al(OH)3 were not enough to protect Ni3Al binder phase in both cermets, thus leading to active dissolution of Ni3Al binder phase (Fig. 4). Only the formation of NiO can effectively suppress active dissolution of Ni-based alloys in H2SO4 solutions, according to the previous studies [64,65]. During immersion in 1 M NaOH solution, NiOOH formed on the surface of both cermets, and moreover, AlOx and Al2O3/Al(OH)3 contents did not significantly change, compared with those prior to immersion. It is well known that NiOOH forms on the surface of Ni electrodes in NaOH solutions under higher potentials, due to the electrochemical oxidation of Ni(OH)2 to NiOOH [64,66]. Consequently, the formation of NiOOH on the surface of both cermets indicated the occurrence of considerable micro-galvanic coupling between Ni3Al binder phase and ceramic grains during immersion in 1 M NaOH solution. For both cermets, the XPS intensities of Ti 2p3/2–2p1/2, W 4f7/2–4f5/ 2 and Mo 3d5/2–3d3/2 spin-split doublets corresponding to Ti–(C,N), W–(C,N) and Mo–(C,N) bonds in ceramic grains of cermet substrate did not significantly change after 60 h immersion in 1 M H2SO4 solution, compared with those prior to immersion, indicating that ceramic grains were not significantly attacked, while those significantly decreased after 60 h immersion in 1 M NaOH solution, indicating preferential dissolution of ceramic grains, which was partially attributed to the occurrence of considerable micro-galvanic coupling between Ni3Al binder phase and ceramic grains (Fig. 6(d) to (f)). Before and after 60 h immersion in both solutions, Ti2O3, TiO2, WO3, MoO2, MoO(OH)2 and MoO3 were present on the surface of both cermets, perhaps including

TiO and Mo2O3. Ti2O3 and TiO2, including TiO, were generally present in the air-formed films on the surface of pure Ti [46,67] and Ti-containing alloys [68]. No W-containing oxides excerpt WO3 was present on the surface of WC–MgO composites after corrosive wear tests [51]. Meanwhile, MoO2, MoO(OH)2 and MoO3 were detected on the surface air-formed films of pure Mo [57]. Compared with those prior to immersion, after 60 h immersion in 1 M H2SO4 solution, Ti2O3, WO3 and Mo-oxides contents did not significantly change on the surface of both cermets, and TiO2 content decreased, indicating ceramic grains did not exhibit passivation characteristic. The cause why TiO2 content decreases is so far not clear, and it needs further investigation. After 60 h immersion in 1 M NaOH solution, TiO2 and Ti2O3 contents did not significantly change whilst WO3 and molybdenum oxide contents decreased on the surface of both cermets, indicating that air-formed TiO2 and Ti2O3 were not nearly dissolved whereas air-formed WO3 and Mooxides were partially dissolved during immersion in 1 M NaOH solution. There was not obvious difference in corrosion mechanism between Ti(C0.5,N0.5)-free cermet C and Ti(C0.5,N0.5)-containing cermet C1 during immersion in 1 M H2SO4 and 1 M NaOH solution, based on surface SEM-SE images and XPS results. No significant micro-galvanic coupling occurred between Ni3Al binder phase and ceramic grains during immersion in 1 M H2SO4 solution, and however, considerable micro-galvanic coupling occurred during immersion in 1 M NaOH solution. 3.3. Electrochemical corrosion behavior Fig. 7 shows typical potentiodynamic polarization curves of as-polished Ti(C0.5,N0.5)-free cermets A, B, C and D, and Ti(C0.5,N0.5)-containing cermet C1, after 1 h immersion at OCP in 1 M H2SO4 solution at room temperature. Obviously, all potentiodynamic polarization curves were similar in shape. There were a passive region in the range from about 38–125 mV vs. SCE to about 643–693 mV vs. SCE, and a pseudopassive region from about 817–863 mV vs. SCE to about 1700–1750 mV vs. SCE. Table 3 lists some characteristic electrochemical parameters of all cermets, including corrosion potential and current density (Ecorr and icorr), passivation potential and critical current density (Epp and icrit), transpassive potential (Etp), and pseudo-passivation potential and critical current density (Epsp and ipsp). Note that icorr was determined by extrapolating the cathodic Tafel line to Ecorr, because there was not an apparent anodic Tafel region. All cermets mainly consisted of Ni3Al binder phase and TiC or Ti (C,N) ceramic phase (Figs. 1 and 2). The active-passive behavior of 5

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Fig. 6. Whole XPS spectra (a) and high-resolution Ni 2p3/2 (b), Al 2p (c), Ti 2p (d), Ti 3p and W4f (e) and Mo3d (f) XPS spectra obtained from surfaces of cermets C (1, 2, 3) and C1 (4, 5, 6) before immersion (1, 4) and after 60 h immersion in 1 M H2SO4 (2, 5) and 1 M NaOH (3, 6) solutions (experimental spectra: black lines; fitted spectra: colored lines). Note: Ni° and Al° come from Ni3Al binder phase in cermet substrate, and (Ti, W,Mo)–(C,N) bonds come from ceramic grains in cermet substrate.

6

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Table 2 Binding energies of the deconvoluted XPS peaks in Fig. 6. XPS spectra

Peaks

Species

Binding energy (eV) Cermet C

Ni 2p3/2

Main

Satellite

Al 2p

2p

Ti 2p

2p3/2

2p1/2

Ti 3p

3p

W 4f

4f7/2 4f5/2

Mo 3d

3d5/2

3d3/2

Ni° (Ni3Al) Ni(OH)2 Ni(OH)x(SO4)y NiOOH Ni° (Ni3Al) Ni(OH)2 Ni(OH)x(SO4)y NiOOH Al° (Ni3Al) AlOx Al2O3/Al(OH)3 Ti–(C,N) Ti2O3 TiO2 Ti–(C,N) Ti2O3 TiO2 Ti–(C,N) Ti2O3 TiO2 W–(C,N) WO3 W–(C,N) WO3 Mo–(C,N) MoO2 MoO(OH)2 MoO3 Mo–(C,N) MoO2 MoO(OH)2 MoO3

Cermet C1

Before immersion

H2SO4

NaOH

Before immersion

H2SO4

NaOH

852.65 856.00

852.75

852.84 856.04

852.61 855.96

852.91

852.84 855.94

856.25 858.45 862.00

858.55

856.36 856.69 858.64 862.04

858.41 861.96

861.95 72.30 74.40 75.35 454.85 455.95 458.45 460.75 461.45 464.15 33.95 35.15 37.05 32.30 35.65 34.40 37.85 228.60 229.35 230.15 232.35 231.80 232.55 233.30 235.50

72.55 74.35 75.40 454.95 456.05 458.35 460.85 461.55 464.05 34.00 35.20 37.10 32.35 35.70 34.45 37.90 228.65 229.42 230.40 232.45 231.85 232.62 233.55 235.60

858.71

856.64 858.64 861.94

862.06 863.29 72.44 74.39 75.54 455.09 456.14 458.89 460.99 461.64 464.59 33.94 35.24 37.09 32.39 35.64 34.49 37.84 228.69 229.34 230.25 232.54 231.89 232.54 233.40 235.69

72.41 74.36 75.36 454.91 456.06 458.46 460.81 461.56 464.16 33.81 35.21 37.16 32.26 35.66 34.36 37.86 228.56 229.16 230.26 232.36 231.76 232.36 233.41 235.51

72.51 74.46 75.46 455.16 456.26 458.56 460.1.06 461.76 464.26 33.86 35.16 36.91 32.31 35.46 34.41 37.66 228.61 229.11 230.11 232.26 231.81 232.31 233.26 235.41

863.24 72.43 74.38 75.38 454.99 455.99 459.04 460.89 461.49 464.74 34.04 35.24 37.09 32.29 35.64 34.39 37.84 228.49 229.14 230.14 232.44 231.69 232.34 233.29 235.59

about 359–459 mV vs. SCE, and passivation took place in the range of about 859–1459 mV vs. SCE, followed by rapid dissolution [72]. For TiC coatings prepared by low-energy ion implantation (LEII) [44] and chemical vapour deposition (CVD) [73], similar active-passive behavior was observed in 0.5 or 1 M H2SO4 solution. Passivation of TiC was found to result from the formation of TiO2 [73]. According to electrochemical corrosion behavior of Ni3Al, Ni and TiC in H2SO4 solutions as described above, in 1 M H2SO4 solution, active dissolution and passivation of all cermets were associated with active dissolution and passivation of their Ni3Al binder phase, respectively, and their transpassive dissolution and pseudo-passivation corresponded to active dissolution and passivation of their TiC or Ti(C,N) ceramic grains, respectively. The pseudo-passive plateau with a width of about 700–800 mV resulted from the interaction between passivation of ceramic grains and transpassive dissolution of binder phase. Above about 1700 mV vs. SCE, the steep increase of current density was attributed to rapid dissolution of ceramic grains. Passivation of all cermets, i.e. passivation of their binder phase, mainly resulted from the formation of NiO, and their pseudopassivation, i.e. passivation of their ceramic grains, mainly resulted from the formation of TiO2. In 1 M H2SO4 solution, for Ti(C0.5,N0.5)-free cermets A, B, C and D with different Ni3Al content, Ecorr and icorr decreased with increasing Ni3Al content (Fig. 7 and Table 3). Ecorr indicates the thermodynamic characteristics of cermets to corrosion in H2SO4 solution [70]. When corrosion took place, the kinetics characteristics of cermets determine the corrosion rate [74]. The decrease of Ecorr was attributed to that alloying content (Ti, W and Mo) in Ni3Al binder phase increased with increasing Ni3Al content, according to XRD results (Fig. 1), and the slight decrease of icorr might be attributed to the increase of Al2O3 content in the oxide films with increasing Ni3Al content. Ecorr and icorr

Fig. 7. Typical potentiodynamic polarization curves of experimental cermets after 1 h immersion at open circuit potential in 1 M H2SO4 solution at room temperature and a sweep rate of 1 mV/s.

Ni3Al was found to be similar to that of pure Ni in H2SO4 solutions [69]. For Ni3Al [69] and pure Ni [70,71], active dissolution, passivation and transpassive dissolution in 0.1–1 M H2SO4 solutions started below −200 mV vs. SCE, at about 40–100 mV vs. SCE and above about 800–900 mV vs. SCE, respectively, and moreover, there was a passive plateau with a width of about 800–900 mV. Passivation of pure Ni was well known to be attributed to the formation of NiO [30,64,65]. For TiC single crystals, in pH 1–3H2SO4 solutions, active dissolution started at 7

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Table 3 Characteristic electrochemical parameters of five experimental cermets in 1 M H2SO4 solution. Experimental cermet

A

B

C

D

C1

Ecorr (mV vs. SCE) icorr (μA/cm2) Epp (mV vs. SCE) icrit (μA/cm2) Etp (mV vs. SCE) Epsp (mV vs. SCE) ips (μA/cm2)

38 ± 0.8 3.59 ± 0.009 69 ± 0.4 33.62 ± 0.008 682 ± 1.4 823 ± 1.8 158.16 ± 0.048

–4 ± 0.6 3.34 ± 0.008 56 ± 0.8 23.81 ± 0.013 671 ± 1.0 860 ± 1.2 95.12 ± 0.041

–29 ± 0.6 2.03 ± 0.005 80 ± 0.5 22.25 ± 0.011 693 ± 0.9 860 ± 0.7 74.55 ± 0.022

–52 ± 0.9 1.64 ± 0.004 65 ± 0.3 16.18 ± 0.004 679 ± 1.1 863 ± 1.1 113.43 ± 0.058

60 ± 0.5 1.88 ± 0.006 125 ± 0.5 3.57 ± 0.005 643 ± 0.8 817 ± 0.8 67.09 ± 0.021

Fig. 8. (a) Typical potentiodynamic polarization curves of experimental cermets after 1 h immersion at open circuit potential in 1 M NaOH solution at room temperature and a sweep rate of 1 mV/s. (b) magnification of narrow passive regions in (a).

behavior is very similar to that to pure Ni [69]. Ni is usually passivated but Al usually dissolves rapidly as aluminate ion AlO2- [69]. The electrochemical oxidation of Ni(OH)2 to NiOOH starts at about 360 mV vs. SCE, forming a pseudo-passive region, prior to the oxygen evolution reaction. According to electrochemical corrosion behavior of Ni3Al, Ni and TiC in NaOH solutions as described above, in 1 M NaOH solution, active dissolution and passivation of all experimental cermets corresponded to active dissolution and passivation of their ceramic grains, respectively, and their pseudo-passivation was mainly attributed to the electrochemical oxidation of Ni(OH)2 to NiOOH. In 1 M NaOH solution, for Ti(C0.5,N0.5)-free cermets A, B, C and D with different Ni3Al content, Ecorr increased and icorr decreased with increasing Ni3Al content (Fig. 8 and Table 4). The increase of Ecorr was attributed to the increase of the total content of alloying elements W and Mo in ceramic phase with increasing Ni3Al content, according to XRD results (Fig. 1), and the decrease of icorr was mainly attributed to the decrease of ceramic grain content with increasing Ni3Al content (Fig. 2). Low corrosion resistance of cermet A might also be associated with the presence of many white particles (Fig. 2(a)). Lower Ecorr and higher icorr of Ti(C0.5,N0.5)-containing cermet C1, compared with those of Ti(C0.5,N0.5)-free cermet cermet C with the same Ni3Al content, indicates that the addition of Ti(C0.5,N0.5) decreased corrosion resistance of Ni3Al bonded TiC-based cermets in 1 M NaOH solution. This might be mainly attributed to that ceramic grains with white core had worse corrosion resistance in 1 M NaOH solution than those with black core,

of Ti(C0.5,N0.5)-containing cermet C1 was higher and lower than those of Ti(C0.5,N0.5)-free cermet cermet C with the same Ni3Al content, respectively, indicating that the addition of Ti(C0.5,N0.5) improved corrosion resistance of Ni3Al-bonded TiC-based cermets in 1 M H2SO4 solution. Fig. 8 shows typical potentiodynamic polarization curves of as-polished Ti(C0.5,N0.5)-free cermets A, B, C and D, and Ti(C0.5,N0.5)-containing cermet C1, after 1 h immersion at OCP in 1 M NaOH solution at room temperature. Obviously, all potentiodynamic polarization curves were very similar in shape. There were a very narrow passive region in the range from about 202–239 mV vs. SCE to about 206–244 mV vs. SCE and a wide pseudo-passive region in the range from about 373–448 mV vs. SCE to about 657–781 mV vs. SCE. Table 4 lists some characteristic electrochemical parameters of all cermets, including corrosion potential and current density (Ecorr and icorr), passivation potential (Epp), and pseudo-passivation potential (Epsp). Note that icorr was determined by Tafel extrapolation method. TiC exhibits low stability in NaOH solutions, according to literature available on its corrosion behavior [72,75]. For TiC single crystals, during potentiodynamic polarization in 1 M NaOH solution, Ecorr was about −400 mV vs. SCE, and current density increased with increasing potential, and however, it is worth noting that current density increased slowly in the range of about –200 to −120 mV vs. SCE, after active dissolution is over, according to the study by Payer et al. [72]. Ni3Al exhibits good stability in NaOH solutions [76], and its electrochemical

Table 4 Characteristic electrochemical parameters of five experimental cermets in 1 M NaOH solution. Experimental cermet

A

B

C

D

C1

Ecorr (mV vs. SCE) icorr (μA/cm2) Epp (mV vs. SCE) Epsp (mV vs. SCE)

–530 ± 4.3 5.25 ± 0.028 202 ± 0.9 448 ± 1.3

–458 ± 7.0 3.25 ± 0.013 231 ± 0.4 447 ± 1.5

–298 ± 5.8 2.33 ± 0.009 228 ± 0.5 447 ± 1.1

–249 ± 4.0 2.45 ± 0.011 239 ± 0.8 414 ± 1.5

–395 ± 4.2 3.65 ± 0.018 230 ± 0.4 373 ± 1.8

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thus the number increase of ceramic grains with white cores led to the decrease of corrosion resistance of cermet C1 (Table 4).

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4. Conclusion The present work investigated corrosion behavior of Ni3Al-bonded TiC-based cermets containing 20 wt% WC and 10 wt% Mo with various Ni3Al contents and Ti(C0.5,N0.5) addition in 1 M H2SO4 and 1 M NaOH solutions at room temperature, using immersion tests, potentiodynamic polarization measurements, XRD, SEM and XPS. The main conclusions are as follows: 1. In 1 M H2SO4 solution, corrosion process of all cermets mainly consisted of Ni3Al binder phase dissolution. During immersion, their binder phase exhibited active dissolution, since Ni(OH)x(SO4)y, AlOx, Al2O3/Al(OH)3, etc., were not enough to protect binder phase. During potentiodynamic polarization, their passivation corresponded to passivation of their binder phase, mainly due to the formation of NiO, and their pseudo-passivation corresponded to passivation of their ceramic grains, mainly due to the formation of TiO2. Corrosion potential and current density decreased with increasing Ni3Al content, while corrosion potential decreased and current density increased with Ti (C0.5,N0.5) addition, indicating better corrosion resistance with Ti (C0.5,N0.5) addition. 2. In 1 M NaOH solution, corrosion process of all cermets mainly consisted of ceramic phase dissolution. During immersion, their ceramic grains were preferentially dissolved, partially due to considerable micro-galvanic coupling between their binder phase and ceramic grains, in accompany with the formation of NiOOH. During potentiodynamic polarization, the narrow passive region of about 40–50 mV corresponded to passivation of their ceramic grains, and their pseudo-passivation corresponded to passivation of their binder phase, mainly due to the formation of NiOOH. Corrosion potential increased and current density decreased with increasing Ni3Al content, indicating better corrosion resistance with increasing Ni3Al binder phase, while corrosion potential decreased and current density increased with Ti(C0.5,N0.5) addition. Acknowledgements The present work was supported by Program for Innovative Research Team by the Ministry of Education of PRC under Grant No. IRT1244, and Major Innovation Program of Hubei Province of China under Grant No. 2016AAA067, and Project supported by State Key Laboratory of Powder Metallurgy, Central South University, Changsha, China. The authors would like to thank Analytical and Testing Center at Huazhong University of Science and Technology for providing experimental facilities. References [1] E.B. Clark, B. Roebuck, Extending the application areas for titanium carbonitride cermets, Int. J. Refract. Met. Hard Mater. 11 (1) (1992) 23–33. [2] G. Schmolz, F.J. Momper, Cermets–cutting materials with a future, Prod. Eng. 3067 (6) (1988) 36–40. [3] G. Östberg, K. Buss, M. Christensen, S. Norgren, H.-O. Andrén, D. Mari, G. Wahnström, I. Reineck, Mechanisms of plastic deformation of WC–Co and Ti (C,N)–WC–Co, Int. J. Refract. Met. Hard Mater. 24 (1–2) (2006) 135–144. [4] Q.Q. Yang, W.H. Xiong, S.Q. Li, H.X. Dai, J. Li, Characterization of oxide scales to evaluate high temperature oxidation of Ti(C,N)–based cermets in static air, J. Alloy. Compd. 506 (1) (2010) 461–467. [5] S.Y. Zhang, Titanium carbonitride–based cermets: processes and properties, Mater. Sci. Eng. A 163 (1) (1993) 141–148. [6] S.Q. Zhou, W. Zhao, W.H. Xiong, Y.N. Zhou, Effect of Mo and Mo2C on the microstructure and properties of the cermets based on Ti(C,N), Acta Metall. Sin. (Engl. Lett.) 21 (3) (2008) 211–219. [7] J. Jung, S. Kang, Effect of ultra–fine powders on the microstructure of Ti (CN)–xWC–Ni cermets, Acta Mater. 52 (6) (2004) 1379–1386. [8] H. Matsubara, S. Shin, T. Sakuma, Grain growth of TiC and Ti(C,N)–based cermets during liquid–phase sintering, Solid State Phenom. 25–26 (1992) 551–558. [9] P.F. Becher, K.P. Plucknett, Properties of Ni3Al-bonded titanium carbide ceramics,

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