The fretting fatigue performance of Ti–6Al–4V alloy influenced by microstructure of CuNiIn coating prepared via thermal spraying

Tribology International 145 (2020) 106156

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The fretting fatigue performance of Ti–6Al–4V alloy influenced by microstructure of CuNiIn coating prepared via thermal spraying Amin Ma a, Daoxin Liu a, *, Xiaohua Zhang a, Guangyu He b, **, Dan Liu a, Chengsong Liu a, Xingchen Xu a a b

Institute of Corrosion and Protection, School of Aeronautics, Northwestern Polytechnical University, Xi’an, Shaanxi, 710072, China Air Force Engineering University, Science & Technology Plasma Dynamics Laboratory, Xi’an, Shaanxi, 710038, China

A R T I C L E I N F O

A B S T R A C T

Keywords: CuNiIn coating Fretting fatigue Rotary flap peening Ti–6Al–4V alloy

Two types of CuNiIn coatings with different microstructures were successfully prepared on the surface of a Ti–6Al–4V alloy by means of high-velocity oxygen fuel (HVOF) and plasma spray (PS), respectively. The effect of the coating microstructure on the fretting fatigue (FF) resistance of the alloy as well as crack initiation and propagation was investigated. Compared with the HVOF-CuNiIn coating, the PS-CuNiIn coating was composed of a more significant lamellar microstructure and exhibited higher FF resistance. HVOF-CuNiIn possessed low resistance to crack initiation due to its low toughness, thus leading to a short FF life of the titanium alloy. The excellent FF resistance of the rotary flay peening (RFP)-treated sample resulted mainly from the large value of the compressive residual stress.

1. Introduction Ti–6Al–4V alloy is the main structural material of aerospace com­ ponents due to its excellent corrosion resistance, high specific strength, low density, and outstanding oxidation resistance [1–3]. However, this alloy has low hardness, poor thermal conductivity, and high friction coefficient, and is therefore quite sensitive to fretting fatigue (FF) damage [4.5]. Notably, the fretting occurring between the compressor blade dovetail and the fan represents a typical case of FF damage that reduces the safety and life of aero-engines [6.7]. The FF process results from a synergistic effect between fretting wear and fatigue damage [8–11]. Thus, surface engineering technologies for reducing the friction coefficient as well as enhancing the anti-wear and anti-fatigue abilities can control the FF damage to titanium alloys [12–14]. Practice has shown that lubrication reduces the friction coefficient, and shot peening is effective in improving the FF resistance of alloys [15–17]. However, surface-hardening treatments (including coating hardening, plasma nitriding, and plasma carburizing) yield only modest improvements in the FF life of titanium alloys [13.18 21]. This results from the low toughness of the treated layers as well as the high notch sensitivity of these alloys. Most of the previous studies have indicated that, owing to their low

friction coefficient, CuNiIn coatings are effective in retarding the fretting damage of titanium alloys [22–24]. Thus, as an anti-fretting measure, plasma spray (PS)–CuNiIn coatings have been applied to a compressor dovetail structure [25]. However, in another case [26], adhesive wear occurred between the titanium alloy and a CuNiIn coating, thereby accelerating the wear damage of the alloy. This may limit the CuNiIn coating-induced enhancement in the FF resistance of Ti–6Al–4V alloy [19]. Furthermore, previous studies have shown that the FF resistance of single PS-CuNiIn coatings is better than that of the shot-peened and PS composite coatings, owing mainly to the release of compressive residual stress after PS [27]. In addition, the high-velocity oxygen fuel (HVOF) process has been widely used in the preparation of anti-wear and anti-corrosion coatings due to the low porosity, high bonding strength, and strong resistance to delamination [28–30]. A few studies have considered the effect of HVOF-prepared CuNiIn coatings on the FF resistance of titanium alloy. Moreover, a few studies have compared the FF behavior and fretting damage mechanism of PS-coated titanium al­ loys with those of HVOF-coated alloys [25]. The RFP method yields outstanding local strength of components, is economical, easily controlled, and conducive to on-site repair. Therefore, this method has been widely used in the on-site repair of aircraft parts [31]. The influ­ ence of RFP on the FF behavior of titanium alloys and the effect of a

* Corresponding author. ** Corresponding author. E-mail addresses: [email protected], [email protected] (A. Ma). https://doi.org/10.1016/j.triboint.2019.106156 Received 1 November 2019; Received in revised form 6 December 2019; Accepted 30 December 2019 Available online 31 December 2019 0301-679X/© 2020 Elsevier Ltd. All rights reserved.

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rotary flay peening (RFP) and CuNiIn coating composite treatment on this behavior have, however, rarely been investigated. Hence, in this work, the influence of PS-prepared and HVOF-prepared CuNiIn coatings on the FF resistance of a Ti–6Al–4V alloy was investigated. Additionally, the effect of a RFP and PS-CuNiIn coating composite treatment on the FF behavior of this alloy was assessed with the aim of improving the FF resistance of titanium alloy engineering equipment. Determining the mechanism governing the effect of different CuNiIn-coating micro­ structures on the FF behavior of Ti–6Al–4V alloy will provide new insight into the mechanism of fatigue fracture and control of the fretting damage to aviation equipment.

Table 1 Chemical composition of Ti–6Al–4V alloy (wt.%). Elements

Al

V

Fe

C

O

H

Ti

wt%

6.1

3.9

0.2

0.10

0.14

0.003

Balance

(1) BM: Base material; (2) HVOF-P: A CuNiIn coating was prepared (by means of HVOF) on the surface of FF samples, and the same polishing technique was then used to treat the surface of this coating; (3) PS-P: A CuNiIn coating was prepared (via PS) on the surface of FF samples, and the same polishing technique was applied to the surface of this coating; (4) RFP: The surface of FF samples was treated via RFP; (5) RFP-PS-P: A CuNiIn coating was prepared (via PS) on the surface of RFP-treated FF samples, and the same polishing technique was then applied to the surface of this coating.

2. Experimental 2.1. Materials and specimens In this work, the fretting fatigue samples were fabricated from annealed-state Ti–6Al–4V alloy bars (αprimary þ βtransformed, diameter: 16 mm, shown in Fig. 1). The chemical composition of this alloy is shown in Table 1, and the mechanical properties (at room temperature) were as follows: yield strength: 1010 MPa, ultimate tensile strength: 1080 MPa, elongation: 15%, area reduction: 41%, and micro-hardness: 380 HK0.245N. The fretting fatigue pads were fabricated from annealed-state PH13–8Mo stainless steel bars (diameter: 16 mm, shown in Fig. 2), in order to simulate the typical connecting condition of aerospace com­ ponents. This steel was characterized by a relatively high Cr and Ni content and low C content (see Table 2 for the chemical composition); Al and Mo acted as precipitation-hardening elements. The mechanical properties of this alloy were as follows: yield strength: 1307 MPa, ulti­ mate tensile strength: 1356 MPa, elongation: 12%, area reduction: 58%, and micro-hardness: 552 HK0.245N.

2.3. FF testing The FF resistance of the Ti–6Al–4V alloy was determined using a GPS-100 high-frequency fatigue machine at room temperature. Testing was performed under the following conditions: load mode of the fatigue samples: pull-pull, stress ratio: 0.1, frequency: 115.0 Hz, maximal load: 650 MPa. The contact state between the fatigue samples (Ti–6Al–4V alloy) and the fretting pads (PH13–8Mo steel) was cylinder to camber [32]. Rela­ tive fretting displacement between the fatigue specimen and the fretting pads was produced via the differing elastic deformation of the specimen and the pads. The contact stress of the pads with the specimen was 85 MPa, which was applied by a stress ring tester. For each group of specimen, the FF test was repeated three times. The morphologies of fracture surfaces and wear regions were eval­ uated via scanning electron microscopy (SEM). In addition, the fatiguecrack initiation process was investigated by assessing the Ti–6Al–4V alloy after the FF test. The fractured alloy was separated via wire elec­ trical discharge machining (WEDM) performed in the vicinity of the fatigue crack source. Subsequently, this cross-section of the fractured titanium alloy was inlaid, and was mechanically polished using sand­ paper (#300, #400, #600, #800, and #1200) and silicon suspensions (1 μm). The crack-initiation characteristics of the sample were analyzed via SEM.

2.2. FF samples with different treatments For comparison, two types (HVOF and PS) of CuNiIn coatings were prepared on the surface of the Ti–6Al–4V alloy. During thermal spray­ ing, the powder particles were heated until fully melted or partially melted by the high-temperature and high-speed flame or plasma. These particles undergo flattening and splashing when they impact the sub­ strate at a certain speed. The deposition velocity of HVOF differed from that of PS and the surface of each coating was rough (Fig. 3(b), (c), and (e)). Thermal spray powders referred to as Amdry 500 F (Cu36Ni5In; particle size: 10 μm 50 μm) were considered in this work. The coatings (thickness: 160 μm 180 μm) deposited by HVOF and PS on the surface of the FF samples were polished to ~100 μm (Fig. 5). The polishing technique used for the base material subjected to the FF test was employed. Furthermore, the surface roughness (Ra) of the HVOF and PS coatings was ~1.0 μm (Fig. 3(e)). The RFP treatment (shown in Fig. 4) was performed with a rotary blade (14.28 mm � 31.75 mm) bonded with cast steel shot (shot peening strength: 0.22–0.24 mmA, coverage: 200%, motor speed: 6000 rpm). In brief, five types of Ti–6Al–4V alloy FF samples were considered in this study:

2.4. Characterization methods � Microstructure of coatings: The surface morphology and crosssectional microstructure of the HVOF and PS CuNiIn coatings were investigated via SEM. Furthermore, the phase patterns of the coat­ ings were measured by means of X-ray diffraction (XRD). The surface roughness of the samples was measured with a roughness/contour tester (TR300), and characterized by Ra and Rz. � Toughness and bonding strength of coatings: The toughness and bonding strength of the coatings were determined using a static indention tester and an in-house-developed cyclic press-press device

Fig. 1. Schematic drawing of FF specimens. 2

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Fig. 2. Schematic drawing of FF pads.

3. Results

Table 2 Chemical composition of PH13–8Mo steel (wt.%). Elements

C

Mn

Si

Cr

Ni

Mo

Al

Fe

wt%

0.04

0.03

0.03

12.7

8.2

2.2

1.1

Balance

3.1. Microstructural characterization The surface morphologies of Ti–6Al–4V alloy (treated and un-treated by RFP), HVOF-CuNiIn, and PS-CuNiIn coatings (before and after pol­ ishing) are shown in Fig. 3(a)–(d). Additionally, the surface roughness of various surface-modified samples is shown in Fig. 3(e). The original surface of the HVOF-CuNiIn coating was flatter than that of the PSCuNiIn coating (Fig. 3(b) and (c)). Compared with the particles in the PS-CuNiIn coating, many partially melted powder particles were more closely packed on the surface of the HVOF-CuNiIn coating (Fig. 3(b)), owing mainly to the characteristics of HVOF. That is, (i) compared with those of the particles corresponding to PS, the flight speed of the powder particles was higher and the in-flight particles stayed in the hightemperature flame for a shorter time. Thus, most of the powder impacting on the surface of the substrate was partially melted; (ii) during coating formation by means of a thermal spray process, the fully melted, partially melted, and unmelted powder particles are deposited layer by layer. The partially melted CuNiIn powder impacted the sub­ strate at high speed (500–700 m/s). Simultaneously, the partially melted and un-melted powder particles helped to improve the substrate-impact power of the powder and enhance the interlaminar bonding strength of the flat particles. This was beneficial for the formation of a dense coating (porosity < 0.1%). However, the partially melted powder particles could retard the flattening of the fully melted powder [34]. The surface of the PS-CuNiIn coating was characterized by a spindrift-like structure, due to the well-flattened features and splashing of the fully melted powder particles. The flight speed of the powder particles was low during the PS process. This resulted in sufficient time for heat transfer between the CuNiIn powder and the high-temperature plasma. Therefore, most of the powder melted upon impact with the substrate resulting in flattening to a laminar structure. However, the fully melted particles were easily splashed, leading possibly to a decrease in the density of the coating, an increase in the number of pore defects, and a decrease in the interlaminar bonding strength. The surface roughness of both coatings decreased significantly after mechanical polishing. In fact, the polishing marks and roughness

[33], respectively. A four prismatic diamond indenter was used for both devices. The following experimental parameters were employed: for the indention tester, maximum loading, 20 N; loading rate, 20 N/min; and dwell time, 60 s. For the press-press device, loading, 40 N; number of cycles, 50,000; and frequency, 20 Hz. After the static indention and cyclic press-press test, the indentations were evaluated by using SEM to compare the toughness and bonding strength of the HVOF-CuNiIn and PS-CuNiIn coatings. � Microhardness: The variation in surface hardness of the samples with increasing load was evaluated using a micro-hardness tester (HV1000) equipped with a Knoop diamond indenter. Each sample was subjected to 20-s loading under 10 g, 25 g, 50 g, 100 g, 200 g, 250 g, and 500 g. � Compressive residual stress: The characteristics of the residual stress along the axial direction and the effect of the PS coating on the re­ sidual stress of the Ti–6Al–4V alloy (treated by RFP) were evaluated via XRD (LXRD-MG2000). The residual stress on the surface of the CuNiIn coatings was measured using Cr–K radiation at the {h k l311} plane of the Cu phase (Bragg angle: 125.28� , irradiated area: 3.14 mm2). The residual stress along the axial direction of Ti–6Al–4V alloy was evaluated using Cu-Kα radiation at the {h k l-213} plane of αTi (Bragg angle: 139.31� , irradiated area: 3.14 mm2, titanium alloy was successfully removed by a mixture of HF and HNO3 solution to determine the in-depth residual stress of this alloy). The effect of the PS-CuNiIn coating on the residual stress of the RFP-treated titanium alloy was determined via the aforementioned method after complete removal of the CuNiIn coating by a mixture Fe(NO3) and HCl solution.

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Fig. 3. Surface morphologies and roughness of treated, un-treated Ti–6Al–4V alloy and two kinds of coatings:(a) Ti–6Al–4V alloy base material; (b) original HVOF coating; (b-1) polished HVOF coating; (c) original PS coating; (c-1) polished PS coating; (d) RFP; (e) surface roughness of treated, un-treated Ti–6Al–4V alloy and polished, unpolished coatings.

observed on the surface of the HVOF-CuNiIn coating were similar to those observed for the substrate (Fig. 3(a), (b-1), and (e)). However, in addition to the longitudinal polishing marks, many micro-pits were observed on the surface of the polished PS-CuNiIn coating (see Figs. 3(c1)). Additionally, the surface roughness value was higher than that of the HVOF-CuNiIn coating, indicating that the HVOF-CuNiIn coating was denser than the PS-CuNiIn coating. Most of the polishing marks dis­ appeared from the surface of the Ti–6Al–6V sample after the RFP treatment, and the sample surface, which was slightly damaged, was evenly covered with spot peening pits. The PS-CuNiIn coatings depos­ ited on the surface of the Ti–6Al–4V alloy substrate and the RFP-treated samples were characterized by similar surface roughness values, and

shot peening had only a slight effect on the roughness. The cross-sectional morphologies corresponding to the two types of polished CuNiIn coatings are shown in Fig. 5. The lamellar microstruc­ ture of the PS coating was more pronounced than that of the HVOF coating (thickness: ~100 μm). Additionally, many pore defects and inter-splat cracks occurred in this coating (see Fig. 5(b)), owing to the low flight speed of the powder during the PS process. This low speed resulted in sufficient heat transfer between the powder and the hightemperature plasma and, hence, many powder particles could be fully melted before arriving at the substrate (Fig. 3(c)). In the case of fully melted particles, with a certain flight speed, the in-pressure of the impacting molten powder particles induced the mass flow and a well4

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Fig. 4. Sketch of RFP equipment.

Fig. 6. XRD analysis of CuNiIn coatings prepared on the surface of Ti–6Al–4V alloy by HVOF and PS.

3.2. Residual stress and hardness

Fig. 5. Cross-section of the polished CuNiIn coatings: (a) the coating prepared by HVOF; (b) the coating prepared by PS.

The compressive residual stress on the surface of the PS-CuNiIn and HVOF-CuNiIn coatings, and the effect of the PS technique on the dis­ tribution of the RFP-induced compressive residual stress are indicated in Fig. 7. As shown in Fig. 7(a), the compressive residual stress on the surface of the HVOF coating was slightly larger than that of the PS coating. As shown in Fig. 7(b), RFP generated a large-magnitude deep residual stress field in the surface and subsurface of the Ti–6Al–4V alloy, thereby improving the fatigue resistance of the material. The afore­ mentioned surface and subsurface distribution of the stress field changed only slightly after the plasma spray. Many experimental studies have considered the compressive residual stress field induced on the surface of titanium alloys via surface deformation strengthening tech­ nology, such as traditional shot peening [35] or an ultrasonic surface rolling process (USRP) [36.37]. The results revealed that this field was the dominant factor in improving the plain fatigue (PF) resistance and FF resistance of these alloys. That is, the residual stress of the RFP-treated layer was only moderately attenuated after PS deposition of the CuNiIn coating, thereby leading to improvement in the FF resistance of the substrate. Fig. 8 shows the microhardness determined for the surface of the PSCuNiIn coating, HVOF-CuNiIn coating, RFP-treated Ti–6Al–4V alloy,

flattened lamellar microstructure was produced by the rapid cooling. However, the low impact energy of the melted particles was unfavorable to the tight bonding of inter-splat, thus pores were easily formed be­ tween the well-flattened lamellas. The HVOF-CuNiIn coating, with a low density of pit defects, was denser than the PS-CuNiIn coating. This resulted mainly from the high flight speed of the partially melted pow­ der particles during the HVOF process. These particles can enhance the effect of kinetic energy on the substrate, thereby increasing the compactness of the coating. The phase analysis of the two coatings (Fig. 6) revealed that, compared with the HVOF-CuNiIn coating, the PS-CuNiIn coating con­ tained more of the Cu0.2Ni0.8O phase. This resulted from the fact that the flight speed of the powder particles was lower than that of the HVOF process, and these particles were exposed for longer times to the hightemperature air plasma (leading to oxidation). Additionally, the (111) preferred orientation of the PS-CuNiIn coating was more pronounced (than that of the HVOF-CuNiIn coating), consistent with the lamellar microstructure of the coating.

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the PS-CuNiIn coating. This resulted mainly from the fact that the HVOFCuNiIn coating was denser than the PS coating and confirmed the use­ fulness of HVOF technology in preparing wear-resistant and corrosionresistant coatings. The RFP method can yield significant improvement in the surface microhardness of the titanium alloy substrate. Under small loads (10 g and 25 g), the microhardness value of the RFP-treated tita­ nium alloy was similar to the value obtained for the fretting pad of the PH13–8Mo steel. The mechanism of RFP method was similar to shot peening and other surface mechanical strengthening treatments. Addi­ tionally, the RFP-treated Ti–6Al–4V alloy substrate was characterized by (~58%) higher microhardness and better loading-capability retention (with increasing load) than the untreated substrate, which might be ascribed to severe plastic deformation and high compressive residual stress induced by RFP. The severe plastic deformation and high compressive residual stress-produced on the substrate surface via these mechanical strengthening treatments were beneficial to improve the microhardness, loading-capability, anti-wear property and fatigue resistance of metallic materials [35,36]. 3.3. Toughness and bonding strength Fig. 9 shows SEM images of the HVOF and PS cross-sections before and after the static indentation test. The regions denoted by the blue dotted lines in Fig. 9(a) and (d) are the same as the indented regions in Fig. 9(b) and (e), respectively. The shape of the indentation on the crosssection of the HVOF coating was smaller than that of the PS coating. This indicated that the HVOF coating exhibited a better loading capability than the PS coating in this cross-section direction (Fig. 8), owing possibly to the high density of the HVOF coating. However, new cracks were induced in the cross-section of this coating (see high-magnification SEM image shown in Fig. 9(c)) during the static indentation test. These cracks may be attributed to the following: (i) during the process of HVOF, most of the CuNiIn powder particles were partially melted, and impacted the substrate and precoating at high flight speed and with high impact kinetic energy. This led to excellent compaction, high loading capability, but poor cohesive strength, low fracture toughness, and easy cracking of the HVOF coating under stress conditions; (ii) micropores were present in the HVOF coating (Fig. 5(a)), although this coating was denser than the PS coating. These micropores played a key role in the generation of stress concentrations for crack initiation and propagation under stress conditions. Simultaneously, cracks formed between the coating and the substrate during diamond indentation (see Fig. 9(a) and (b)) indicating that the bonding strength of the coating on the substrate surface was only moderately high. However, the PS coating had low loading capability. Therefore, the occurrence of crack defects between the original lamellar structure (quite flat melted powder particles) had a certain buffering and releasing effect on the static indentation process. This prevented the formation of new cracks in the cross-section of the PS coating (Fig. 9(e)) subjected to static indentation. The indentation led instead to the closing and opening of crack defects originally parallel to the interface between the coating and the substrate. Furthermore, many pores were generated at the interface between the PS coating and the titanium alloy substrate (see Fig. 9(d) and (e)). That is, the bonding strength between the coating and the substrate was low, owing to the low flight speed of the powder particles during the PS process. Fig. 10 shows the morphologies of the cycle press-press indentation on the surface of the HVOF and PS coatings. For the same test condition, the shape on the surface of the HVOF coating was smaller than that on the PS coating, indicating that the loading capability of the HVOF coating was higher (Fig. 8). This was consistent with the results of the static indentation tests, where smaller indents were obtained due to the better compactness of the HVOF coating, compared with that of the PS coating (Fig. 5). After 50000 press-press cycles, layered debris produced by pressing on the edge of the indentation (see zone indicated by blue rectangle in Fig. 10(a)) revealed that the cohesive strength of the HVOF coating was low. Simultaneously, the crack induced in the corner of the

Fig. 7. Residual Stress on the surface of PS and HVOF CuNiIn coatings (a) and residual stress at various depth from the surface of treated Ti–6Al–4V alloy by RFP before and after PS coating (b).

Fig. 8. The surface hardness of samples various with loading increasing.

PH13–8Mo steel substrate, and Ti–6Al–4V alloy substrate subjected to increasing indentation load. For a load of 10 g, the microhardness of the HVOF-CuNiIn coating is 27% higher and 19% higher than the hardness values of the PS-CuNiIn coating and the titanium substrate, respectively; moreover, the loading capability (with increasing loading) of the HVOFCuNiIn coating was similar to that of the substrate and better than that of 6

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Fig. 9. Static indentation morphology of CuNiIn coatings: (a) CuNiIn coating prepared by HVOF before indent; (b) CuNiIn coating prepared by HVOF after indent; (c) the enlarged image of yellow rectangle marked with dot line in (b); (d) CuNiIn coating prepared by PS before indent; (e) CuNiIn coating prepared by PS after indent. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)

Fig. 10. Vickers indentation surface morphology of CuNiIn coatings prepared by HVOF ((a), (c)) and PS ((b), (d)).

indentation indicated that the toughness of the coating was low (consistent with the results of static indentation tests). This was quite different from the PS-CuNiIn coating, where considerable interlayer bulging (due to extrusion) was observed at the indentation edge of the coating. That is, the bonding strength of interlayers in the PS-CuNiIn coating was lower than that of the HVOF-CuNiIn coating (Fig. 5: cross-sectional morphology shows the transverse crack parallel to the interface of the coating and the substrate). This may be related to the relatively low velocity of CuNiIn powder particles during plasma spraying (fully melted powder particles could be obtained via sufficient heat exchange with a high-temperature flame). The collision between the molten particles led to a strong metallurgical bond and, hence, the toughness of the single-layer coating was better than that of the HVOF coating.

3.4. FF resistance The FF life associated with different surface modification treatments of the Ti–6Al–4V alloy is shown in Fig. 11. The lowest FF life, which was ~1/2 that of the untreated substrate, was obtained for the alloy with the HVOF-CuNiIn coating, indicating that this coating was detrimental to the FF resistance of the alloy. However, the FF life of the substrate increased by 35%, owing to the PS-CuNiIn coating. The difference be­ tween the effects of the HVOF and PS coatings on the FF resistance of the Ti–6Al–4V alloy was strongly correlated with the mechanical properties and microstructure of the coatings. The FF life of the RFP-treated alloy was almost two times higher than that of the untreated alloy, owing to the good surface integrity introduced during shot peening [38.39]. The PS-CuNiIn coating prepared on the RFP-treated surface led to a two-fold 7

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(f)). However, for the HVOF-CuNiIn coating, fatigue peel off and sec­ ondary cracks protruding in the fretting wear area were observed (as indicated by the blue arrow in Fig. 12(c)). This may have resulted from the low toughness of the coating. In conclusion, the reduction in the FF resistance of the alloy substrate, owing to the HVOF-CuNiIn coating, may have resulted from the coating-induced acceleration of fatiguecrack initiation and reduction in the crack-initiation life. Moreover, as shown in Fig. 12(d), a crack formed in the interface between the coating and the titanium alloy substrate. This crack was induced by multi-axial loading, and could be attributed to the low coating/substrate bonding strength (see Fig. 9(b)). The coating and the substrate were bonded via mechanical bonding [23.25]. Previous studies have illustrated that, owing to their low friction coefficient, CuNiIn coatings are beneficial for retarding the fretting damage of the alloys [22.23]. Furthermore, the low hardness of these coatings is beneficial for accommodating the plastic deformation and alleviating the stress concentration of the contact zone between the fa­ tigue sample and fretting pads. This may lead to significant improve­ ment in the FF life of metal materials [27.40.44]. Furthermore, the present research indicated that the good toughness of the coatings is beneficial for retarding fatigue crack initiation in the fretting area [28]. The secondary crack induced on the contact surface of the HVOF-CuNiIn coating (as shown in Fig. 12(c)) revealed that the toughness of this coating is low. The reduction in the FF life of the BM may be attributed to this low toughness. Additionally, the thickness of the coating contributed significantly increase in durability under fretting conditions. A previous study [13] has illustrated that a thermal spray CuNiIn coating (thickness: 30 μm) was worn away after the fretting wear experiment and, hence provided the substrate no protection from fretting damage. Moreover, Rajase­ karan et al. [47] reported that the fatigue life of Al–Mg–Si alloy gun sprayed with a CuNiIn coating (thickness: 40 μm) is superior to that of the uncoated and 100-μm-thick coated specimens. This superior life resulted from the fact that the surface compressive residual stress in the former was higher than that of the latter. Zhai et al. [48,49] gave a comprehensive review on wear performance of the additives-containing nanocomposite coating (fullerenes, CNTs, graphene and nano­ diamonds). These composite coatings performed more dramatic anti-wear property, compared with the pure metal coating, which was ascribed to their lower friction coefficient. Thus, a suitable thickness, a high compressive residual stress and a low friction coefficient of the coatings are important for enhancing the FF life of an alloy [50]. Simultaneously, the large number of pore-defects occurring in the PS-CuNiIn coating had a negative effect on the FF-life improvement of this alloy, and the number of these pores would increase with increasing coating thickness. Thus, the PS-CuNiIn coating yielded only modest enhancement in the FF resistance of the substrate. The fretting wear occurring in the contact zone between fretting pads and the RFP-treated titanium alloy resulted mainly from fatigue wear and abrasive wear (Fig. 12(i)). The FF crack of RFP-treated sample initiated from the bottom of fretting wear pits in the contact zone (Fig. 12(j)). The fracture surface characteristic of both the BM and RFPtreated samples was multiple crack source. However, the FF crack growth of the BM was instable when the crack initiated (Fig. 12(b)). Moreover, severe delamination occurred in the fretting wear region (Fig. 12(i)), owing to the long FF life of the sample and the severe cu­ mulative damage. Fatigue wear occurred on the surface of the fretting region in the Ti–6Al–4V alloy subjected to the RFP and PS combined treatment. The characteristics of the fracture (Fig. 12(h)) in the coating area and the titanium alloy area were similar to those of the coating (Fig. 12(f)) and the RFP-treated alloy (Fig. 12(j)), respectively. Therefore, the enhancement in the FF resistance of the alloy via the compound treat­ ment resulted from the effect of the RFP method on the fatigue behavior of the BM. That is, owing to the treatment, fatigue-crack initiation in BM was impeded, and the crack-propagation resistance as well as the crack-

Fig. 11. Fretting fatigue life of treated and un-treated Ti–6Al–4V samples.

improvement in the FF life compared with that of the untreated titanium alloy. Previous studies illustrated that the PS-CuNiIn coating led to a reduction in the friction coefficient and resulted in absorption of the fretting energy [27.40.41]. Simultaneously, the compressive residual stress was only modestly reduced by the plasma spray (Fig. 7) and, hence, the FF-life dispersion after the compound treatment was lower than that of the RFP-only treatment. Shot peening, a surface strength­ ening technique, can induce a deep and large compressive residual stress, and the residual stress field can cancel some of the fatigue tensile stress. This can inhibit the initiation and early growth of fatigue cracks, thereby leading to significant improvement in the FF resistance of the titanium alloy. Furthermore, the compressive residual stress has a sig­ nificant effect on the FF life of the alloy [5]. 4. Discussion 4.1. Characteristics of friction and fracture Fig. 12 shows the morphologies of the fracture surface and fretting wear damage on the contact region of the titanium alloy under various surface conditions. The alloy has poor thermal conductivity, a high friction coefficient, and low microhardness and is, thus quite susceptible to FF damage [8.42]. The surface fretting wear occurring in the contact region of the fatigue sample (Ti–6Al–4V base material) with the hard fretting pads (PH13–8Mo steel) resulted mainly from fatigue wear and abrasive wear (Fig. 12(a)). The accumulated fretting wear led to severe damage in this area. Moreover, the fatigue cracks were initiated from the bottom of the fretting wear pits (Fig. 12(b)), owing to the multi-axial load and the severe stress concentration on the bottom of the pit. The fretting wear damage area is a quite strong source for crack initiation under FF conditions [24.35]. Due to the effect of the normal load applied between the fretting pads and fretting fatigue sample via the stress ring, the coatings prepared through HVOF and PS were characterized by a high bonding strength with the Ti–6Al–4V base material. The process of FF fracture began from the surface of the coatings, and propagated to the substrate (Fig. 12(d), (f), and (h)), consistent with previously reported results [25]. The FF sample with the HVOF-CuNiIn coating and PS-CuNiIn coating on the fretting wear region underwent mainly fatigue wear (Fig. 12(c) and (e)). The fatigue crack was initiated from the surface of the coatings, and propagated perpendicular to the direction of external alternating load (the source of FF cracks indicated by yellow arrows in Fig. 12(d) and (f)). The equivalent crack effect might promote fatigue-crack propagation to the interior of the titanium alloy until fracture occurs when the crack reaches the substrate/coating interface (Fig. 12(d) and 8

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Fig. 12. Fretting wear region ((a), (c), (e), (g), and (i)) and fracture surface ((b), (d), (f), (h), and (j)) of treated and un-treated Ti–6Al–4V samples: (a) (b) Ti–6Al–4V alloy base material; (c) (d) HVOF coating; (e) (f) PS coating; (g) (h) RFP-PS coating; (i) (j) RFP.

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zone. As shown in Fig. 13(a), the FF source on the surface of the Ti–6Al–4V alloy underwent mainly peel-off wear. Furthermore, the leading crack was initiated from the bottom of the fretting wear pit and propagated along the direction of the maximum shear force. The secondary-cracks were either initiated from the bottom of the pits or from the subsurface of the severe wear region. These cracks then grew along the direction of the maximum shear force (~30� from the surface). This is the general rule of FF damage in metallic materials, such as ti­ tanium alloys [5.43]. The FF crack of the titanium alloy coated with the HVOF-CuNiIn coating was initiated from the surface of the coating. The internal crack developed rapidly along the direction perpendicular to the surface

initiation life were improved [2.15.43]. The RFP-induced compressive residual stress was, however, only partly reduced by the PS treatment. Moreover, compared with the pure RFP treatment, the combined treatment yielded only modest improvement in the FF life because, the: (i) RFP-induced compressive residual stress was slightly reduced by the PS process; and (ii) PS-CuNiIn coating resulted in only slight improve­ ment in the FF resistance of the BM. 4.2. Fracture behavior Fig. 13 shows the cross-sectional morphologies of the titanium alloy fracture under various surface conditions near the fatigue crack resource

Fig. 13. Cross-sectional morphology of FF specimen after failure with different modification: (a) Ti–6Al–4V alloy base material; (b) HVOF coating; (c) PS coating; (d) (e) RFP-PS coating; (f) cracks propagation from coating to substrate; (g) RFP. 10

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(indicated by the blue bot line in Fig. 13(b)), consistent with the results of Fig. 12(c) and (d). This low resistance to crack development may be correlated with the poor toughness of the HVOF-CuNiIn coating. The decrease in the FF resistance of the Ti–6Al–4V alloy resulted mainly from the HVOF-CuNiIn coating-induced reduction in the fatigue crack initiation life. Fig. 13(c) shows that the fatigue crack of the substrate coated with the PS-CuNiIn coating was initiated from the bottom of the fretting wear pit. The crack propagated along the direction of maximum shear force until arriving at the inherent inter-plate cracks and then either branched or terminated. This branching or termination played a key role in improving the FF resistance of the BM via the PS-CuNiIn coating. Additionally, the low hardness of this coating was beneficial for absorbing the fretting energy, accommodating the plastic deformation, and alleviating the stress concentration of the contact zone [27.44]. Therefore, the rate of crack propagation in this coating was slow and the FF life of the BM under this condition was high, i.e., the FF resistance of the titanium alloy was improved by the PS-CuNiIn coating. Different microstructures were obtained for the coatings prepared via different thermal spray technologies. However, the effect of the coating microstructure on the FF resistance of the substrate has only been considered in a few study [45]. In this work, the FF life of the HVOF-CuNiIn coated Ti–6Al–4V alloy was lower than that of the BM, whereas the FF life of the PS-coated alloy was higher. The results might be related to the effect of the hardness, thickness, toughness, and surface compressive residual stress of the HVOF and the PS coatings, as well as the different microstructure of the coatings. As shown in Fig. 13(b)–(e), for the typical lamellar microstructure of the PS-CuNiIn coating, the interface of the plates can increase the fatigue crack-growth resistance (hinder fatigue-crack propagation to the interior of the substrate).This increased resistance and the zigzag nature of the growth path lead to good FF resistance of the coating. Additionally, during the fatigue-crack propagation process, low amounts of crack branching occurred in the dense HVOF coating, indicating that the resistance to crack development was low. The significant lamellar microstructure of a PS-CuNiIn coating may therefore be beneficial for controlling the propagation of fatigue cracks. The results obtained for the influence of the coating micro­ structure on fatigue-crack propagation provide new insight for further study of the fatigue mechanism characterizing these coatings. In brief, the FF-life improvement of the titanium alloy via the PSCuNiIn coating was mainly correlated with the enhancement of the initiation life and early growth life of the fatigue crack. The PS-CuNiIn coating yielded only modest enhancement in the FF resistance of the substrate, owing to the fact that, the: (i) hardness of the PH13–8Mo steel pads was dramatically higher than that of the coating; (ii) contact stress of the contact region between the substrate and fretting pads was high; and (iii) pore defects occurring in the coating had a negative effect on the FF-life improvement of this alloy. The RFP method led to significant enhancement in the fatigue wear resistance and abrasive wear resistance of the titanium alloy substrate, thereby limiting nucleation and initiation of the FF crack. Thus, fewer secondary cracks were induced in the RFP-treated sample than in the BM (Fig. 13(g)). Limiting the formation of these cracks plays a key role in improving the FF resistance of titanium alloys by means of the RFP technique [15.38.46]. The FF crack is usually initiated from the bottom of fretting wear pits on the surface of the PS-CuNiIn coating prepared on the RFP-treated titanium alloy. The FF crack-propagation process of this coating pro­ ceeded as follows (see Fig. 13(d) and (e)). At the beginning of this stage, the crack would propagate along the direction of maximum shear stress. Afterward, the crack would either propagate along the inter-splat crack (Fig. 13(d)) or terminate. Energy accumulation for the subsequent initiation of this crack (indicated by red arrows in Fig. 13(e)) would occur when the crack arrives at inherent inter-splat cracks (perpendic­ ular to the propagation direction). Additionally, when the fatigue crack arrived at the coating/substrate interface, the crack would penetrate the

substrate due to its large accumulated energy (indicated by yellow arrow in Fig. 13(f)). In conclusion, the high FF life of the Ti–6Al–4V alloy treated via RFP and the PS-CuNiIn coating could be attributed to the surface work hardening and the RFP-induced large-magnitude deep compressive residual stress. The RFP treatment inhibited fatigue-crack initiation and propagation [16,31]. The high FF also resulted from the lamellar microstructure of the coating, which led to fatigue crack branching and deflection, thereby controlling fatigue-crack propagation in the coating. 4.3. FF mechanism Based on the above results and analyses, the effects of the various surface conditions (BM, polished PS-CuNiIn coating, polished HVOFCuNiIn coating, RFP, and polished RFP-PS coating) on the FF damage mechanism of the Ti–6Al–4V alloy are schematically presented in Fig. 14. In the beginning of the FF test performed on the Ti–6Al–4V substrate, wear pits were rapidly generated in the mixed slip wear region between the fatigue sample and the fretting pads (Fig. 14I–B). This resulted from the fact that the fretting damage occurring in the mixed slip region was worse than that in the slip region and the adhesive region. After these pits were generated, fatigue cracks were initiated from the bottom of the pits (caused by severe local stress concentration [24]) and then propa­ gated inside along the direction of the large shear stress (Fig. 14I–C). Subsequently, as the depth of the crack propagation increased, the leading fatigue crack propagated along the direction perpendicular to the contact surface (Fig. 14I–D). This was attributed to the reduced ef­ fect of the shear stress and contact stress in the contact region [5]. A schematic of the FF evolution process occurring in the BM is shown in Fig. 14 I. The surface of the HVOF-CuNiIn coating was smooth after the me­ chanical polishing process, and the coating was characterized by high compactness and high hardness. However, the low toughness and low cohesive strength led to a low FF resistance. Additionally, a microcrack perpendicular to the fretting direction was easily initiated in the contact region of the coating (shown in Fig. 14 II-B and Fig. 12(c)). This resulted from the low toughness and the high shear stress in the mixed slip region caused by the fretting pads. The microcrack promoted rapid fatiguecrack initiation and propagation into the depth of the coating sub­ jected to an alternating load (Fig. 14 II-C). Afterward, the high accu­ mulated energy in the fatigue crack-tip resulted in continuous crack growth into the substrate when the fatigue crack reached the coating/ substrate interface. In contrast to that of the PS-CuNiIn coating, the FFlife reduction of the substrate with the HVOF-CuNiIn coating resulted mainly from the low resistance to fatigue-crack initiation and develop­ ment in this coating. A schematic of the FF evolution process occurring in the BM with the HVOF-CuNiIn coating is shown in Fig. 14 II. The surface of the PS-CuNiIn coating, which was generally smooth after the mechanical polishing process, was characterized by low porosity, low compactness, and low hardness. The toughness and fatigue wear resistance of this coating were higher than those of the HVOF coating. In the beginning of the FF test, the high toughness and low hardness of this coating could enhance the resistance to fatigue-crack initiation and alleviate the stress concentration of the contact zone, respectively [27]. However, the pore defects would promote fatigue-crack initiation. The mechanism of fatigue-crack initiation from the mixed slip wear region on the coating surface could therefore be attributed to the aforementioned factors (Fig. 14 III-B and C). Moreover, fatigue crack branching and deflection during the propagation process, resulted mainly from the quite lamellar microstructure (Fig. 14 III-D and Fig. 13(c)–(e)). After this branching and deflection, the high accumu­ lated energy in the fatigue crack-tip led to continuous crack growth into the substrate when the fatigue crack reached the coating/substrate interface (Fig. 13(f)). A schematic of the FF evolution process occurring in the BM with the HVOF-CuNiIn coating is shown in Fig. 14 III. 11

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Fig. 14. Schematic illustration of fretting fatigue damage of treated and un-treated Ti–6Al–4V alloy: (I) Ti–6Al–4V alloy base material; (II) HVOF coating; (III) PS coating; (IV) RFP; (V) RFP-PS coating; (A) surface morphology; (B) beginning of FF; (C) initiation of fatigue crack; (D) cracks propagation.

For the RFP-treated substrate, the RFP-induced compressive residual stress and surface work-hardening layer inhibited fatigue-crack initia­ tion. At the beginning of the FF test, microcracks were initiated and then propagated along the direction approximately parallel to the contact surface, resulting in large-flake wear (Fig. 14 IV-B and Fig. 13(g)). During continuous wear, the size of the large flake debris decreased and micro pits were induced by the severe fatigue wear. The local stress concentration in the bottom of the pits resulted in fatigue-crack initia­ tion in the mixed slip wear region (Fig. 14 IV-C). Moreover, the fatiguecrack propagation was similar to that occurring in the BM (Fig. 14 IV-D). A schematic of the FF evolution process associated with the RFP-treated sample is shown in Fig. 14 IV. The mechanism of the FF damage occurring in the titanium alloy treated via the RFP and PS-CuNiIn coating compound method was similar to that of the titanium alloy treated with the PS-CuNiIn coating only. However, the RFP-induced compressive residual stress and surface work-hardening layer could prevent fatigue-crack propagation to the interior of the alloy. Thus, the high FF life of the alloy obtained via the compound treatment was attributed to two main factors: (i) retardation of fatigue crack development, owing to the quite lamellar microstruc­ ture of the PS-CuNiIn coating; and (ii) inhibition of fatigue-crack initi­ ation and propagation in the coating/substrate interface, owing to the RFP-induced compressive residual stress. Additionally, the RFP-only

method yielded high dispersibility, which resulted mainly from the high initiation life of the fatigue crack on the surface of the substrate [36.38]. The dispersibility of the FF life associated with the compound treatment condition was lower than that of the RFP-only treatment (Fig. 11). This dispersibility represented a very important feature of the compound treatment. A schematic of the FF evolution process occurring in the BM with the PS-CuNiIn coating and the RFP compound treatment is shown in Fig. 14 V. 5. Conclusion (1) A CuNiIn coating prepared on the surface of Ti–6Al–4V alloy via HVOF was characterized by excellent compactness, high micro­ hardness, and loading capability similar to that of the titanium alloy substrate. The cohesive strength and toughness of the coating were, however, both low. The PS coating exhibited excellent splat toughness, but (unlike the HVOF coating) this coating consisted of many pores and inter-splat crack defects, and was characterized by low microhardness and lower loading capability than the BM. However, the microhardness of the RFPtreated titanium alloy was similar to that of the PH13–8Mo steel (fretting pads), owing possibly to the RFP-induced surface work hardening. The RFP treatment induced a compressive residual 12

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(3)

(4)

(5)

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stress field (maximum value, ~820 MPa; distribution depth, ~200 μm) on the surface of the titanium alloy. Furthermore, for the combined RFP and PS treatment, the compressive residual stress was only modestly reduced after PS. The FF life of the Ti–6Al–4V alloy decreased by ~50% through the use of the HVOF-CuNiIn coating. This decrease could be attributed to the reduction in the fatigue-crack initiation life due to the inherent toughness of the HVOF coating. Through the use of the PS-CuNiIn coating, the FF life of the Ti–6Al–4V alloy increased by ~35%. This may have resulted mainly from the synergistic effect of the high toughness and low microhardness that inhibited fatigue-crack initiation. The low hardness of the coating could effectively alleviate the stress concentration of the contact zone. Additionally, the direction of the inter-splat cracks in the PS coating was approximately par­ allel to the surface of the samples. That is, the inter-splats cracks were approximately perpendicular to the propagation direction of the fatigue cracks, which may have resulted in enhanced FF resistance of the BM. The surface work hardening of the Ti–6Al–4V alloy during RFP resulted in improved resistance to surface wear and fatigue-crack initiation and, consequently, an almost two-fold improvement in the FF life of the alloy, due to this treatment. At the same time, the compressive residual stress field led to significant inhibition of fatigue-crack initiation and propagation. Through the combined RFP and PS technique, the FF life of the BM was improved almost two-fold, owing to the synergistic effect of the PS coating and RFP on enhancing the FF resistance. Additionally, compared with the RFP-only treatment, this com­ pound treatment led to lower dispersibility of the FF life, result­ ing from alleviation of the stress concentration in the contact zone via the PS coating.

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Author contributions section Amin Ma: designed and performed the experiments, performed the data analysis and mechanism analysis, and wrote the paper. Daoxin Liu: designed the experiments, and helped Amin Ma to write the paper and analyze the mechanism. Xiaohua Zhang: helped Amin Ma to design the experiments. Guangyu He: performed the data analysis. Dan Liu: per­ formed the data analysis. Chengsong Liu: contributed to the mechanism analysis. Xingchen Xu: contributed to the mechanism analysis. All au­ thors contributed to the general discussion. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgments This work was supported by the National Natural Science Foundation of China (51771155), the National Science and Technology Major Project (2017-VII-0012-0107) and the Equipment pre-research field fund (61409220202). References [1] Carroll BE, Palmer TA, Beese AM. Anisotropic tensile behavior of Ti-6Al-4V components fabricated with directed energy deposition additive manufacturing. Acta Mater 2015;87:309–20. https://doi.org/10.1016/j.actamat.2014.12.054. [2] Tang CB, Liu DX, Tang B, Zhang XH, Qin L, Liu CS. Influence of plasma molybdenizing and shot-peening on fretting damage behavior of titanium alloy. Appl Surf Sci 2016;390:946–58. https://doi.org/10.1016/j.apsusc.2016.08.146. [3] Lu JX, Chang L, Wang J, Sang LJ, Wu SK, Zhang YF. In-situ investigation of the anisotropic mechanical properties of laser direct metal deposition Ti6Al4V alloy. Mater Sci Eng A 2018;712:199–205. https://doi.org/10.1016/j.msea.2017.11.106.

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[36] [37]

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