# Tribological investigation of new self-fluxing nickel alloys for high temperature application: The effect of silicon distribution on glaze layer formation

S. Baydoun, F.Z. Moul-El-Ksour, S. Fouvry, G. Guillonneau, J.C. Pereira, F.

Santos, A. Niklas, P. Lopez-Ruiz, J. Vega, J. Rocchi

► To cite this version:

S. Baydoun, F.Z. Moul-El-Ksour, S. Fouvry, G. Guillonneau, J.C. Pereira, et al. Tribological investigation of new self-fluxing nickel alloys for high temperature application: The effect of silicon distribution on glaze layer formation. Wear, 2024, 564-565, pp.205631. 10.1016/j.wear.2024.205631. hal-05378004

HAL Id: hal-05378004

https://hal.science/hal-05378004v1

Submitted on 22 Nov 2025

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<figure><img src="image_1.png" /></figure>
S. Baydoun, F.Z. Moul-El-Ksour, S. Fouvry, G. Guillonneau, J.C. Pereira, F. Santos, A. Niklas, P. Lopez-Ruiz, J. Vega, J. Rocchi, Tribological investigation of new self-fluxing nickel alloys for high temperature application: The effect of silicon distribution on glaze layer formation, Wear 564–565 (2025) 205631
(https://doi.org/10.1016/j.wear.2024.205631)

Tribological investigation of new self-fluxing nickel alloys for high temperature
application: the effect of silicon distribution on glaze layer formation

S. Baydoun¹*, F.Z. Moul-El-Ksour¹, S. Fouvry¹,⁶*, G. Guillonneau¹, J.C. Pereira², F. Santos³, A. Niklas³, P. Lopez-Ruiz⁴, J. Vega⁵, J. Rocchi⁵

¹Ecole Centrale de Lyon, LTDS Laboratory, 36 av Guy de Collongue, 69130 Ecully, France

²LORTEK Technological Centre, Basque Research and Technology Alliance BRTA, Arranomendia Kalea 4A, Ordizia.
Gipuzkoa, Spain

³AZTERLAN Metallurgy Research Centre, 6, 48200 Durango, Bizkaia, Spain

⁴CIDETEC, Basque Research and Technology Alliance (BRTA), Pº. Miramón 196, 20014 Donostia-San Sebastián,
Spain

⁵Liebherr Aerospace Toulouse SAS, 408, Avenue des Etats-Unis, BP 52010, CEDEX 2, 31016 Toulouse, France

⁶MINES Paris, PSL University, MAT - Centre des Matériaux, CNRS UMR 7633, BP 87 91003 Evry, France

*Corresponding authors email addresses: soha.baydoun@ec-lyon.fr (S. Baydoun) & siegfried.fouvry@minesparis.psl.eu (S.Fouvry)

**Abstract**

The objective of this study is to investigate the tribological performance of new nickel-
based self-fluxing alloys for replacing cobalt-based alloys in aircraft components. Two new
NiCrSiFeB alloys, manufactured by gravity casting and laser metal deposition processes,
having different microstructures, are submitted to fretting loadings and tested at
temperatures ranging from ambient to 650 °C. Their wear behavior is compared to a
reference Stellite 6 cobalt- based alloy. Although the two NiCrSiFeB alloys have the same
chemical composition, the alloy having a coarse microstructure rich in silicates was able
to form a protective glaze layer at high temperature as Stellite 6. On the other hand, the
alloy having a fine microstructure underwent severe adhesive wear. This highlights the
crucial effect of the microstructure and the silicon distribution in activating the formation of
a protective glaze layer at high temperatures.

**Keywords:** Fretting wear; Metal Additive Manufacturing; Gravity Casting; Self-fluxing NiCrSiFeB alloys; Inconel718; Silicon; Sealing rings; High temperatures; Glaze layer.

# 1. Introduction

Cobalt-based alloys are broadly used in wear-related applications involving high temperature, poorly lubricated systems and corrosive environments [1,2]. They are recognized as an attractive material for many engineering applications as aircraft engines, power generation, and automotive industry. They are also characterized by forming a protective glaze layer at high temperature inhibiting wear under sliding conditions [1,3–5]. Due to the adverse health implications associated to cobalt wear particles in different applications and in particular in bleed air valve systems in aircrafts, there is a global tendency to replace cobalt by less contaminant alloys.

Nickel-based self-fluxing alloys are good candidates for replacing cobalt as they exhibit remarkable wear resistance properties at high temperature [6,7]. Ni-Cr-B-Si are among the most widely used alloys in hard-facing applications. Indeed, boron and silicon exhibit an interesting self-fluxing property through reacting with oxides at high temperature forming borosilicate limiting in turn additional oxidation of the main elements of the alloy [8]. The presence of chromium next to boron and carbon ameliorates the wear and corrosion resistance of the material through formation of hard carbides and borides [8].

In most of cases these alloys are used for coatings through different techniques as thermal spraying, yet limited are the studies showing the casting applications of these alloys.
Unfortunately, the effect of the microstructure and the chemical composition of the cast
alloys on wear resistance was poorly investigated in literature especially regarding the
formation of a protective glaze layer. Recent studies conducted by Viat *et al.* [3] and
Dreano *et al.* [4] compared the capacity of different pure metals including nickel, chromium
and cobalt in forming a lubricant glaze layer at high temperature. This was achieved by
doing fretting wear experiments. The authors underlined that among the tested elements
(Co, Cr, Ni), cobalt element has the best capacity in forming a protective glaze layer. They
attributed this behavior to the high cobalt auto-diffusion coefficient in its own oxides [4].
This enhances the tribo-sintering process which is at the origin of glaze layer formation.
Viat *et al.* [3] also suggested that this process is kinetically driven implying that nickel and
chromium alloys can still form glaze layer if adequate test duration, temperature and
debris retention are ensured. This hypothesis is consistent with the fact that glaze layer
was formed in many nickel based alloys as reported by Stott *et al.* [9] and Inman *et al.*
[10].

The objective of this study is to compare the performance of cobalt free Ni-Cr-B-Si alloys
which are usually applied as coatings using thermal spray technologies. One originality of
this research work is using these alloys as bulk metals obtained by advanced
manufacturing process and investigate their tribological behavior. The ultimate purpose
behind this research is to ameliorate cabin air quality by evading potential wear particles
of cobalt alloys. Another goal is fabricating higher wear resistant sealing rings, sliding
bearing seats and leakage-free valves to be installed in new hotter and more pressurized
bleed air for the future aircraft engines involving less bleed Environmental Control
Systems (ECS) [11].
Most of the literature studies suggest that the benefit of silicon is related to an increase of the oxidation reaction and the hardness. However, there is still an open question regarding the formation of a protective glaze layer which was shown to be a driving factor of wear resistance at high temperature. A study reported by Buckley and Johnson [12] revealed that the presence of small amount of silicon in nickel alloys led to a significant reduction and friction and wear at high temperatures. They attributed this to the formation of a protective surface films enhanced by affinity of silicon to oxidants resulting in the formation of silicon dioxide (SiO2) that combines with low metal oxides favoring the formation of glaze layer. This study focused on the silicon content in the nickel alloy with a little focus on the effect on its microstructural distribution in the alloy. Hence, the objective of this study is to assess the effect of the microstructure distribution of silicon element in the constitution of the glaze layer. To achieve this target, a similar Ni-Cr-B-Si composition alloy was fabricated by different manufacturing processes to monitor the formation of Ni-Si phases. By adjusting the proportion of Ni-Si, it is expected to investigate how the silicon can influence the creation of a protective glaze layer.

## 2. Experimental procedure

### 2.1 Materials and contact geometry properties

The contact geometry consists of crossed-cylinder configurations at 90° (Figure 1). The top and bottom cylindrical specimens have an external radius $R=12.5$ mm, a length $L=30$ mm, and a surface roughness $Ra=0.5$ µm.
<figure><img src="image_2.png" /><figcaption>Figure 1: Description of the crossed cylinders contact geometry</figcaption></figure>

The bottom counterface sample is made up of an Inconel718 casting substrate with 90 µm coating consisting of nickel matrix and chromium carbide applied using High Velocity Oxy Fuel (HVOF) thermal spray coating process.

The top samples consist of three different materials including Stellite 6 and two new nickel self-fluxing alloys denoted by C-LMD and C-GC. C-LMD is fabricated using Laser metal deposition (LMD) which is Directed Energy Deposition Laser Beam (DED-LB) metal Additive Manufacturing (AM) such that a laser beam forms a melt pool on the surface of a metal substrate into which powder is fed using a gas stream. On the other hand, C-GC is fabricated using gravity casting (GC) which is normally used to cast thin-walled cylinders by pouring a molten metal into a casting mold, usually a spinning die. Table 1 shows the chemical composition of each alloy and of the coated Inconel718 (coating and substrate) analyzed using Energy Dispersive Spectroscopy (EDS). The two nickel-based alloys (C-GC and C-LMD) have the same chemical composition but different microstructures as detailed in the following section.
<table><caption>Table 1: Summary of the chemical composition (% weight) for the reference alloy (Stellite 6) the nickel alloys (C-LMD and C-GC) and the Inconel718 (coating + substrate) analyzed using Energy Dispersive Spectroscopy (EDS).</caption><thead><tr><th>Alloy</th><th>Ni (%)</th><th>Cr (%)</th><th>Si (%)</th><th>Fe (%)</th><th>B (%)</th><th>C (%)</th><th>Co (%)</th><th>F (%)</th><th>Others</th></tr></thead><tbody><tr><td>Stellite 6</td><td>1.2</td><td>38.7</td><td>0.8</td><td>1.4</td><td>0.0</td><td>7.7</td><td>44.4</td><td>-</td><td>5.9</td></tr><tr><td>C-LMD</td><td>76.8</td><td>12.5</td><td>3.7</td><td>4.0</td><td>2.4</td><td>0.5</td><td>0.1</td><td>-</td><td>0.0</td></tr><tr><td>C-GC</td><td>76.8</td><td>12.5</td><td>3.7</td><td>4.0</td><td>2.4</td><td>0.5</td><td>0.1</td><td>-</td><td>0.0</td></tr><tr><td>Inconel coating</td><td>36.8</td><td>47.2</td><td>-</td><td>0.0</td><td>-</td><td>10.8</td><td>-</td><td>3.0</td><td>2.5</td></tr><tr><td>Inconel718 substrate</td><td>54.8</td><td>20.5</td><td>-</td><td>19.8</td><td>-</td><td>-</td><td>-</td><td>2.3</td><td>2.6</td></tr></tbody></table>

<figure><img src="image_3.png" /><figcaption>Figure 2: BSE micrographs showing the microstructures of the studied alloys: Stellite 6, C-GC, and C-LMD.</figcaption></figure>

Figure 2 compares the microstructures of the studied alloys. Stellite 6 exhibits a cobalt based matrix (Figure 2, A), along with a chromium-rich secondary phase (Figure 2, B) and light precipitates rich in W and Mo (Figure 2, C).

The characterization of NiCrSiFeB alloys reveals a coarse heterogeneous microstructure with star-shaped precipitates for gravity casting samples (C-GC). This is not the case for the C-LMD alloy fabricated by laser metal deposition (LMD) displaying rather a fine homogeneous microstructure with a dendritic matrix. Concerning the chemical composition of C-GC alloys, the main phases in the matrix (Figure 2, D) is γ-Ni along with
lamellar Ni-Si fine eutectic (Ni₃Si) (Figure 2, E) and a globular Ni-B coarse eutectic (Ni₃B)
(Figure 2, F) in the inner part of the matrix. On the other hand, the precipitates of C-GC
consist of chromium-rich borides displaying dark grey color (Figure 2, G), along with black
chromium-rich carbides (Figure 2, H). As for C-LMD alloys, the matrix phase is composed
mainly of γ-Ni dendrites (Figure 2, I) along with a grey phase of Ni-Si and Ni-B eutectics
(Figure 2, J). In the case of precipitates of C-LMD, chromium-rich borides are detected
with dark grey color (Figure 2, K) along with black chromium-rich carbides (Figure 2, L).

In order to determine that phases that can be formed by the manufacturing process during
solidification, Scheil and equilibrium solidification simulations were performed elsewhere
[13,14]. Since GC process involves a slow cooling rate, a thermal equilibrium can be
assumed hence an equilibrium solidification model has been considered [13] (Figure 3a).
In the case of LMD process involving an out-of-equilibrium, Scheil model has been applied
[14] (Figure 3b). It can be underlined that cooling rate highly influences the microstructural
evolution of this material due to its microstructural sensitivity variations. In the case of the
LMD process, the formation of borides and carbides is favored over Ni₃Si precipitation
due to the different diffusion rates of the elements involved. Silicon (Si) diffuses more
slowly compared to boron (B) and carbon (C). During the rapid solidification of the LMD
process, there is insufficient time for significant diffusion of Si, which restricts the formation
of Ni₃Si and leads to the preferential formation of borides and carbides. Conversely, in the
gravity casting (GC) process, the cooling rate is much slower, allowing for a near-
equilibrium solidification. This slower cooling rate provides ample time for the diffusion of
Si, promoting the formation of Ni₃Si phases. The extended diffusion time in the C-GC alloy
facilitates the growth of a more significant proportion of Ni-Si eutectics, leading to a
microstructure that is rich in Ni₃Si. Additionally, the slower cooling rates in the GC process result in a coarser eutectic structure compared to the finer structure observed in the rapidly cooled C-LMD alloy. This tendency is supported by solidification simulations (Figure 3) suggesting higher proportions of Ni₃Si in C-GC compared to C-LMD.

<figure><img src="image_4.png" /><figcaption>Figure 3: Solidification simulations showing the amounts of phases formed during solidification as a function of temperature. (a) Equilibrium solidification (C-GC) [13] and (b) Scheil solidification (C-LMD) [14]</figcaption></figure>
These computations are qualitatively confirmed by bulk microstructural observations (Figure 4) through Scanning by Electron Microscopy and Energy Dispersive Spectroscopy (SEM-EDS). These observations clearly underline the presence of eutectic structure with high proportion of Ni₃Si phase in the C-GC alloy compared to the C-LMD microstructure.

<figure><img src="image_5.png" /><figcaption>Figure 4: SEM-EDS analysis revealing the presence of eutectic structures with high proportion of Ni₃Si phase in the C-GC alloy compared to the C-LMD</figcaption></figure>

Figure 5 compares the Vickers hardness of the Stellite 6 and the nickel alloys at T= 25°C and 650°C. It can be noticed that the three alloys display similar hardness values especially at T=650°C. However, alloy C-LMD alloy has the highest mean values at both ambient and high temperature with lowest scatter. This is attributed to the fact that C-LMD
has a homogeneous microstructure which is not the case of C-GC having heterogeneous phases and consequently high dispersion in hardness values. This stems from disparity between the high hardness values of the secondary phases being equal to 1899 ± 153 HV for chromium carbides and 2257 ± 44 HV for borides compared to the low hardness (398 ± 54 HV) of the nickel matrix.

<figure><img src="image_6.png" /><figcaption>Figure 5: Comparison of the Vickers hardness at T=25°C and T=650°C for the reference alloy (Stellite 6) and NiCrSiFeB nickel-based alloys (C-GC and C-LMD).</figcaption></figure>

## 2.2 Experimental setup and data acquisition

Reciprocating wear tests were conducted on a high temperature tribometer to conduct high temperature wear tests (Figure 6a). The coated Inconel counterface was fixed on the bottom sample holder while the new tested alloys are fixed on the top sample holder where a displacement $\delta$ with a frequency f is imposed using an electromagnetic shaker. The normal load P is kept constant by applying dead loads. High temperatures up to 700°C are applied thanks to a heating system consisting of five heating rods placed in the bottom sample holder as well as an insulated chamber allowing the confinement and the homogeneity of the sample temperature during the test. The sample temperature is recorded during the test using a thermocouple. Note that temperature monitoring was previously optimized by inserting a thermocouple very close to the contact surface without
fretting motion. A piezoelectric sensor records the tangential force (Q) induced by the
friction and an optical laser sensor measures the displacement δ. Sliding loops, Q versus
δ, are plotted for each cycle of the test (Figure 6b). The friction coefficient μ is defined as
the ratio of the maximum tangential force Q* to normal force P. To guarantee a given
sliding amplitude independent of the compliance of the tribometer, gross slip sliding tests
were performed by monitoring the sliding amplitude δ₀, defined as the residual
displacement when Q=0. On the other hand, the displacement amplitude δ* is
continuously adjusted to keep a constant δ₀ at a given required value.

<figure><img src="image_7.png" /><figcaption>Figure 6: Schematic illustration of the: (a) fretting experimental set-up; (b) sliding log with δ₀ maintained constant by adjusting δ*.</figcaption></figure>

2.3 Test conditions

To assess the performance of the new alloys with respect to the reference Stellite 6 alloy,
a crossed-test strategy varying the temperature and the test duration was applied. Figure
7 illustrates the applied methodology: two key reference test conditions (repeated three
times each) corresponding to 25°C and 650°C sample temperatures keeping constant the
test duration at N=200k cycles, normal force P=23 N corresponding to maximum Hertzian
contact pressure p=700 MPa, sliding amplitude δ₀= ±450 µm, and a sliding frequency f=20
Hz (Figure 7). A second set of experiments is then applied where the sample temperature is varied from T=25°C to 650°C while fixing the number of fretting cycles at 200k cycles. Then, for each reference temperature (i.e. T=25°C and 650°C), the number of cycles is varied from 1k to 200k cycles.

<figure><img src="image_8.png" /><figcaption>Figure 7: Multi-scale experimental strategy to assess the evolution of wear volume as a function of sample temperature and number of cycles.</figcaption></figure>

## 2.4 Surface analysis

To estimate the wear volume, the top and bottom samples are cleaned in ultrasonic bath with ethanol for 20 minutes to eliminate the remaining debris particles. Then, optical observations are done followed by 3D wear scars acquired using interferometric surface profilometry (Figure 8).
<figure><img src="image_9.png" /><figcaption>Figure 8: (a) Optical observation; (b) 3D interferometric profilometry for wear volume estimation (after [15]).</figcaption></figure>

By defining a reference surface, the wear volume is calculated from the difference between the volumes below and above the reference plane denoted by V⁻ and V⁺ respectively (Eq. 1).

$$V_t = V^{-} - V^{+} \qquad (1)$$

By summing the wear volumes of the top (Stellite 6, C-GC, or C-LMD alloy) and bottom (associated Inconel counterpart) samples, the total wear volume V can be calculated (Eq. 2).

$$V = V_t (\text{alloy}) + V_t (\text{Inconel counterpart}) \qquad (2)$$

### 3. Results

#### 3.1 Fretting wear behavior versus sample temperature

Figure 9 compares the evolution of the wear volume and the friction coefficient of the reference alloy (Stellite 6) and the nickel-based alloys (C-LMD and C-GC) by varying the sample temperature from 25°C to 650°C. It is interesting to note that Stellite 6 and C-GC exhibit similar behaviors.
<figure><img src="image_10.png" /><figcaption>Figure 9: Evolution of the (a) total wear volume (Equation 2, i.e. the sum of the wear volumes of the C-GC, C-LMD and Stellite 6 alloys and their corresponding coated Inconel counterparts). (b) Evolution of the friction coefficient as a function of the sample temperature for the reference alloy (Stellite 6) and the nickel-based alloys (C-GC, and C-LMD) tested against coated Inconel alloy (P=23 N, f=20 Hz, $\delta_0$= ±450 µm, and N=200k cycles).</figcaption></figure>

Three domains can be discerned [4]. Domain I ($T < T_1$) characterized by a continuous increase in the wear volume up to maximum value reached at $T_1=300^\circ C$. The friction coefficient is quite high in this domain. Domain II ($T_1 < T < T_2$) displays a sharp decrease
in the wear volume and the friction coefficient until reaching very low wear rate above
500°C (domain II) although the friction coefficient remains quite stable around 0.4. This
signifies that the production of wear debris is drastically reduced at the interface due to
the formation of a lubricant protective glaze layer above a threshold temperature as
revealed by [3,4,16–18]. Finally, domain III (T > T₂) is detected such that a very low wear
volume and a relatively low coefficient of friction marking the creation of glaze layer.
Hence, glaze layer formation, which is classically observed in cobalt alloy is also detected
in this study for C-GC nickel based alloy.

On the other hand, alloy C-LMD manifests a very different behavior although its chemical
composition is similar to C-GC. C-LMD alloy displays the same tendency in domain I and
partly in domain II until 400°C above which a very disturbed evolution is observed. Such
behavior can be explained by the absence of the glaze layer at high temperature. Hence,
very high wear volume is observed in C-LMD above 500°C. Similarly, the coefficient of
friction is relatively high at ambient temperature, then it decreases and remains stable till
T<400°C. After that, the coefficient of friction increases gradually with the increase of
temperature. Additionally, by observing the wear scars of C-GC and C-LMD after contact
opening at T=650°C (Figure 10), it appears that no wear particles are detected on C-GC
which is not the case for C-LMD where a considerable amount of wear particles induced
by abrasive/adhesive wear processes was observed. This confirms that C-LMD fails to
form a stable protective glaze layer at high temperature generating large quantity of wear
particles.
<figure><img src="image_11.png" /><figcaption>Figure 10: Comparison of the wear scars of alloys C-GC and C-LMD tested against coated Inconel alloy directly after opening the contact at 650°C tested against coated Inconel alloy (P=23 N, f=20 Hz, δ₀= ±450 µm, and N=200k cycles).</figcaption></figure>

To better clarify the surface damage between samples, Figure 11a & b picture the evolution of the wear volume and the maximum wear depths of Stellite 6, C-LMD and C-GC as a function of the sample temperature. The wear volume of Stellite 6 and the maximum wear depth increase until T=150°, then decrease progressively until reaching very low values starting from T=500°C. On the other hand, the wear volume of C-GC and its wear depth remain relatively low compared to C-LMD and Stellite 6 at all temperatures.
<figure><img src="image_12.png" /></figure>

<figure><img src="image_13.png" /></figure>

<figure><img src="image_14.png" /></figure>

<figure><img src="image_15.png" /><figcaption>Figure 11: Evolution of the wear volume of (a) alloys (Stellite 6, C-GC, and C-LMD) and the (b) maximum wear depth of these alloys; (c) wear volume of coated Inconel and the (d) maximum wear depth of the coated Inconel as a function of the sample temperature (P=23 N, f=20 Hz, $\delta_0$= ±450 µm, and N=200k cycles).</figcaption></figure>

C-LMD alloy, however, has the highest wear volume and wear depth at high temperatures exceeding 300°C confirming that this alloys fails to maintain or create an anti-wear glaze layer. In contrast, the coated Inconel counterpart of C-LMD displays a surprising low wear volume at high temperature which will be shown later as an artifact related to material transfer from C-LMD to the counterface (Figure 11c). Figure 11d indicates that the
substrate is reached for temperature between T=100°C and T=400°C. This indirectly infers that high wear volume of C-LMD alloy is independent of the presence of coating on the Inconel counterpart. Besides, for the high temperature situations (i.e. T>500°C) being the domain of interest for this study, the substrate is not reached, hence, the coated interface can be considered for tribological analysis.

## 3.2 Fretting wear behavior

### 3.2.1 Wear and friction response at T=25°C

Figure 12 displays the wear behavior for the alloys (Stellite 6, C-LMD and C-GC) for different number of cycles at T=25°C. At ambient temperature, the results show that Stellite 6 and C-GC exhibit the same bilinear behavior suggesting a severe wear at the beginning followed by a mild wear once a stable third body layer is formed (Figure 12a). The decrease in the wear rate is, however, smaller for C-LMD which still maintains a high wear rate even after 200k loading cycles. On the other hand, these three alloys have equivalent coefficient of friction which is increases significantly until 50k cycles then stabilizes or increases slightly after this value (Figure 12b).
<figure><img src="image_16.png" /></figure>

<figure><img src="image_17.png" /><figcaption>Figure 12: Evolution of the (a) total wear volume and the (b) friction coefficient at T=25°C for the alloys (Stellite 6, C-GC, and C-LMD) tested against coated Inconel 718 (P=23 N, f=20 Hz, δ₀= ±450 µm, and T=25°C).</figcaption></figure>

Figure 13 compares the evolution of the wear scars of the tested alloys and their corresponding coated Inconel counterpart at ambient temperature. From optical images and 3D profilometry for all alloys, we can notice the presence of abrasive grooves indicating that abrasive-oxidational wear is dominant at low temperature.
<table><caption>Figure 13: Comparison of the wear tracks of the alloys (Stellite 6, C-GC, and C-LMD) and their coated Inconel counterpart using optical images and 3D profilometry at ambient temperature (P=23 N, f=20 Hz, δ₀= ±450 µm and N=200000 cycles).</caption><thead><tr><th></th><th>T=25°C</th><th>Stellite 6 / coated Inconel</th><th>C-GC / coated Inconel</th><th>C-LMD / coated Inconel</th></tr></thead><tbody><tr><th rowspan="2">optical images</th><th>counter body</th><td>Image of a counter body coated with Stellite 6 / coated Inconel</td><td>Image of a counter body coated with C-GC / coated Inconel</td><td>Image of a counter body coated with C-LMD / coated Inconel</td></tr><tr><th>coated Inconel</th><td>Image of a coated Inconel</td><td>Image of a coated Inconel</td><td>Image of a coated Inconel</td></tr><tr><th rowspan="2">3D Profilometry</th><th>counter body</th><td>Image of a counter body coated with Stellite 6 / coated Inconel</td><td>Image of a counter body coated with C-GC / coated Inconel</td><td>Image of a counter body coated with C-LMD / coated Inconel</td></tr><tr><th>coated Inconel</th><td>Image of a coated Inconel</td><td>Image of a coated Inconel</td><td>Image of a coated Inconel</td></tr></tbody></table>

Examining the 3D surface profiles, it can be concluded that the largest wear value is observed in fretting surface of C-LMD fretting followed by that of Stellite 6 and finally C-GC. Alternatively, the fretting scar volume of counter surfaces (i.e. coated Inconel) is the largest when fretted against C-GC, then C-LMD and finally Stellite 6. Hence, C-GC alloy, displaying large carbide and boride structures tend to limit its own surface wear but causes high wear damage on the coated Inconel counterpart.
### 3.2.2 Wear and friction response at T=650°C

At T=650°C (Figure 14a), Stellite 6 and C-GC confirm a very low wear volume evolution compared to C-LMD.

<figure><img src="image_18.png" /><figcaption>Figure 14: Evolution of the (a) total wear volume and the (b) friction coefficient at T=650°C for the alloys (Stellite 6, C-LMD, C-GC) tested against coated Inconel 718 (P=23 N, f=20 Hz, $\delta_0$ = ±450 µm, and T=650°C).</figcaption></figure>
This suggests that, by contrast to Stellite 6 and C-GC, C-LMD alloy is unable to form or maintain a glaze at high temperature. The is also reflected by the friction coefficient (Figure 14b) which remains relatively high for C-LMD and equivalent to that obtained at ambient temperature. Yet, Stellite 6 and C-GC have relatively low friction coefficient at T=650°C compared to that obtained at ambient temperature marking the creation of a lubricant glaze layer.

Figure 15 compares the evolution of the wear tracks of the tested alloys and their corresponding coated Inconel counterpart at T=650°C. Optical images and 3D profilometry reveal that the wear scars and the wear volume are relatively small in Stellite 6 and C-GC compared to C-LMD due to the presence of glaze layer in Stellite 6 and C-GC. This latter is more visible in the coated Inconel counterparts of Stellite 6 and C-GC where a shiny layer is observed which is a characteristic feature of glaze layer. In contrast, transfer of material is detected from C-LMD alloy toward the coated Inconel counterpart. This transfer phenomenon induces an "apparent" low wear volume in the coated Inconel counterpart. Note that the surface texture of the coated Inconel fretting wear scar is typical for adhesive wear process. This material transfer phenomenon causes a very high wear volume on C-LMD. What seems quite surprising in these results is that C-GC and C-LMD have the same chemical composition, yet C-GC is capable of creating a glaze layer whereas C-LMD could not. This suggests that the chemical composition is not the sole parameter controlling the genesis of glaze layer as there is a prominent effect of the microstructure which is the only difference between these alloys.
<table><caption>Figure 15: Comparison of the wear tracks of the alloys (Stellite 6, C-GC, and C-LMD) and their coated Inconel counterpart using optical images and 3D profilometry at T=650°C (P=23 N, f=20 Hz, δ₀= ±450 µm and N=200000 cycles).</caption><thead><tr><th></th><th>T=650°C</th><th>Stellite 6 / coated<br>Inconel</th><th>C-GC / coated<br>Inconel</th><th>C-LMD /<br>coated Inconel</th></tr></thead><tbody><tr><th rowspan="2">optical images</th><th>counter<br>body</th><td><img src="optical_image1.png" alt="Optical image of a counter body coated with Stellite 6 / coated Inconel"></td><td><img src="optical_image2.png" alt="Optical image of a counter body coated with C-GC / coated Inconel"></td><td><img src="optical_image3.png" alt="Optical image of a counter body coated with C-LMD / coated Inconel"></td></tr><tr><th>coated<br>Inconel</th><td><img src="optical_image4.png" alt="Optical image of a coated Inconel with Stellite 6 / coated Inconel"></td><td><img src="optical_image5.png" alt="Optical image of a coated Inconel with C-GC / coated Inconel"></td><td><img src="optical_image6.png" alt="Optical image of a coated Inconel with C-LMD / coated Inconel"></td></tr><tr><th rowspan="2">3D Profilometry</th><th>counter<br>body</th><td><img src="3d Profilometry_image1.png" alt="3D Profilometry image of a counter body coated with Stellite 6 / coated Inconel"></td><td><img src="3d Profilometry_image2.png" alt="3D Profilometry image of a counter body coated with C-GC / coated Inconel"></td><td><img src="3d Profilometry_image3.png" alt="3D Profilometry image of a counter body coated with C-LMD / coated Inconel"></td></tr><tr><th>coated<br>Inconel</th><td><img src="3d Profilometry_image4.png" alt="3D Profilometry image of a coated Inconel with Stellite 6 / coated Inconel"></td><td><img src="3d Profilometry_image5.png" alt="3D Profilometry image of a coated Inconel with C-GC / coated Inconel"></td><td><img src="3d Profilometry_image6.png" alt="3D Profilometry image of a coated Inconel with C-LMD / coated Inconel"></td></tr></tbody></table>

To better assess the C-LMD adhesive wear process at T=650°C, the wear tracks of this alloy and its counterpart are compared to C-GC for different test durations at T=650°C. Figure 16 clearly underlines that material transfer from C-LMD to the coated Inconel is activated at early stages of the tribological process at least before 1000 cycles. It becomes more pronounced with the increase in the test duration. This leads to large wear volume on C-LMD due to the considerable material transfer toward the coated Inconel counter
face accompanied by abrasive-oxidational wear phenomenon causing large wear particles ejection. Alternatively, C-GC alloy fretting wear scar as well as its corresponding coated counterface display very low wear damage. The wear damage is slightly more pronounced after few thousands of fretting cycles on the counterface which can be attributed to the initial phase preceding the creation of glaze layer where abrasive-oxidational wear occurs. Yet, once the glaze layer is created, no further wear would be detected as revealed by Dreano et al. [4]. Again, these results confirm the fast and efficient effect of the glaze layer formed with the C-GC alloy which cannot be achieved using the microstructure related to the C-LMD alloy.

<table><caption>Figure 16: Evolution of the 3D wear profiles of the alloys (C-GC and C-LMD) and their coated Inconel718 counterpart as a function of the number of cycles (P=23 N, f=20 Hz, δ₀= ±450 µm, and T=650°C).</caption><thead><tr><th>number of cycles</th><th>1000</th><th>100 000</th><th>200 000</th></tr></thead><tbody><tr><td>C-GC</td><td></td><td></td><td></td></tr><tr><td>coated Inconel of C-GC</td><td></td><td></td><td></td></tr><tr><td>C-LMD</td><td></td><td></td><td></td></tr><tr><td>coated Inconel of C-LMD</td><td></td><td></td><td></td></tr></tbody></table>
In order to better describe the transfer process generated by C-LMD surface fretted against coated Inconel at 650°C, SEM and EDS analyses of C-LMD and C-GC fretting wear scars have been performed (Figure 17).

<table><caption>Figure 17: EDS of analysis of the coated Inconel718 counterpart of the alloys (C-GC and C-LMD) as a function of the number of cycles (P=23 N, f=20 Hz, δ₀= ±450 µm, and T=650°C).</caption><thead><tr><th rowspan="2">number of cycles</th><th colspan="3">coated Inconel against C-GC</th><th colspan="3">coated Inconel against C-LMD</th></tr><tr><th>El.</th><th>%wt</th><th></th><th>El.</th><th>%wt</th><th></th></tr></thead><tbody><tr><td rowspan="4">1000</td><td>Ni</td><td>23.7</td><td></td><td>Ni</td><td>46.9</td><td></td></tr><tr><td>Cr</td><td>55.7</td><td></td><td>Cr</td><td>9.3</td><td></td></tr><tr><td>O</td><td>9.3</td><td></td><td>O</td><td>23.2</td><td></td></tr><tr><td>C</td><td>11.3</td><td></td><td>C</td><td>8.8</td><td></td></tr><tr><td rowspan="4">100 000</td><td>Ni</td><td>30.1</td><td></td><td>Ni</td><td>61.9</td><td></td></tr><tr><td>Cr</td><td>32.6</td><td></td><td>Cr</td><td>10.4</td><td></td></tr><tr><td>O</td><td>26.8</td><td></td><td>O</td><td>14.3</td><td></td></tr><tr><td>C</td><td>3.3</td><td></td><td>C</td><td>2.7</td><td></td></tr><tr><td rowspan="4">200 000</td><td>Ni</td><td>26.3</td><td></td><td>Ni</td><td>62.3</td><td></td></tr><tr><td>Cr</td><td>39</td><td></td><td>Cr</td><td>10.6</td><td></td></tr><tr><td>O</td><td>23.2</td><td></td><td>O</td><td>16.3</td><td></td></tr><tr><td>C</td><td>5.3</td><td></td><td>C</td><td>4</td><td></td></tr></tbody></table>

Concerning the coated Inconel of C-GC, it can be discerned that the chromium content is higher than the nickel content at all number of cycles such that the average ratio of nickel to chromium is equal to 0.67. This proportion is equivalent to nickel to chromium ratio in the Inconel coating which is equal to 0.78. This means that the EDS analysis is capturing the chemical composition of the Inconel coating suggesting that the thickness of the glaze layer is relatively small as will be shown later in the discussion. In contrary, the coated Inconel sliding against C-LMD manifests a lower chromium content compared to nickel at all number of cycles such that the average ratio of nickel to chromium is equal to 5.63.
This proportion has the same order of length to the nickel to chromium ratio in C-LMD which is equal to 6.14. This confirms the transfer of material from C-LMD to the coated Inconel counterpart forming a thick adhesion zone on the surface.

These results can be further confirmed by comparing the wear scars of C-GC and C-LMD alloys at T=650°C as pictured in Figure 18. After 1000 cycles, the post-mortem BSE analysis of C-GC alloy reveals an oxide layer all over the wear track residing over the hard phases of C-GC. In contrary, very low oxide content is detected in C-LMD validating the premise that material transfer takes place from C-LMD to its coated Inconel counterpart. By passing to 200000 fretting cycles, a continuous glaze layer is formed all over the wear scar in C-GC. In contrary, the C-LMD fretting scars display powdery oxides in some regions underlying that C-LMD alloy is unable of forming a continuous glaze layer which in turn favours the activation of critical adhesive wear phenomenon.
<table><caption>Figure 18: Comparison of the BSE images of the wear scars of the alloys (C-GC and C-LMD) as a function of the number of cycles (P=23 N, f=20 Hz, δ₀= ±450 µm, and T=650°C).</caption><thead><tr><th>number of cycles</th><th>C-GC</th><th>C-LMD</th></tr></thead><tbody><tr><td>1000</td><td></td><td></td></tr><tr><td>100 000</td><td></td><td></td></tr><tr><td>200 000</td><td></td><td></td></tr></tbody></table>

### 3.3 Influence of CrC-Ni coating

The current study was done using a CrC-Ni coated Inconel counterface. There is a key aspect to evaluate to what extent the presence of CrC-Ni coating can influence the wear and friction response of the studied alloys. To assess this aspect, some tests were performed without the CrC-Ni coating using directly Inconel718 alloy machined to provide similar surface roughness (Ra<0.5 µm). Figure 19 compares the friction coefficient at ambient and high temperature using coated and uncoated Inconel counterfaces. Similar tendencies are observed suggesting that the presence of CrC-Ni coating does not influence significantly the tribological behavior which appears mainly driven by the studied material. Both Stellite 6 and C-GC display equivalent response whereas higher friction
coefficient is detected for C-LMD at high temperature due to its inability to form a protective glaze layer.

<figure><img src="image_19.png" /><figcaption>Figure 19: Comparison of the friction coefficient of the alloys (Stellite 6, C-GC, and C-LMD) tested against coated and uncoated Inconel 718 at (a) T=25°C and (b) T=650°C (P=23 N, f=20 Hz, δ0= ±450 µm, and N=200k cycles).</figcaption></figure>

A similar tendency is observed regarding the total wear volume evolution (Figure 20).

Under low temperature (Figure 20a), equivalent wear volumes are observed whether the
counterface is coated or not. However, at 650°C, C-LMD alloy displays a very high wear
volume compared to Stellite 6 and C-GC especially in the absence of coating on the
Inconel counterpart (Figure 20b). Hence, the presence of CrC-Ni coating does not
influence the wear mechanism which is driven by the capacity of the studied alloy to form
or not a protective glaze layer.
<figure><img src="image_20.png" /></figure>

<figure><img src="image_21.png" /><figcaption>Figure 20: Comparison of the total wear volume of the alloys (Stellite 6, C-GC, and C-LMD) tested against coated and uncoated Inconel 718 at (a) T=25°C and (b) T=650°C (P=23 N, f=20 Hz, δ0= ±450 µm, and N=200k cycles).</figcaption></figure>

The objective of these additional tests is to show that the CrC-Ni coating does not modify the wear mechanisms so that the given results could be extended to describe the tribological behavior of the given alloys against uncoated Inconel. However, this analysis also suggests that the presence of CrC-Ni coating tends to reduce the dramatic transfer of C-LMD to the counterface. This is confirmed in Figure 21 where the wear volumes
measured on the counterface are compared. Without the CrC-Ni coating, significant negative wear volumes are measured on the Inconel fretted against C-LMD marking huge material transfer process. This can be very detrimental for the application (i.e. catastrophic seizure). Hence, the application of the CrC-Ni coating on the Inconel surface tends to limit this detrimental effect reducing the tribological damage when the glaze layer is not formed.

<figure><img src="image_22.png" /><figcaption>Figure 21: Comparison of the friction coefficient of the alloys (Stellite 6, C-GC, and C-LMD) tested against coated and uncoated Inconel718 at T=650°C (P=23 N, f=20 Hz, δ₀= ±450 µm, and N=200k cycles).</figcaption></figure>

## 3.4 Hardness influence

The wear resistance increases with the material Hardness according to Archard wear model [19]. To quantify the potential effect of hardness, the Archard wear coefficient K (mm³/J), defined as the ratio between the total wear volume (V) and the accumulated Archard work (ΣW) with W=4.δ₀.P [19], is computed for each temperature test condition.

$$V = K \times \Sigma W \qquad (2)$$

Figure 22 compares the evolution of Archard wear coefficient K of the alloys Stellite 6, C-GC and C-LMD as a function of the hardness H depending on the sample temperature.
Two different domains can be discerned for C-GC and Stellite 6 (Figure 22a & b). For T<300°C, the wear rate increases when the hardness decreases, in other words for this temperature range the wear volume is inversely proportional to the hardness which is consistent with Archard wear theory. When the temperature goes from 300°C to 650°C, a decrease in the wear volume is noticed, yet the hardness keeps decreasing slowly. This behavior can be attributed to the fact that when the temperature is below 300°C, the glaze layer is not formed yet so the Archard law is still valid where abrasive-oxidational wear occurs. But when the temperature surpasses 300°C, glaze layer is formed. The formation of this cohesive glaze layer highly modifies the tribological behavior. The wear rate strongly decreases with the reduction of the hardness. The Archard theory is no longer valid in this case as it does not take into account the tribo-chemical process resulting in the genesis of glaze layer. For C-LMD, an oxidational-abrasive wear is still occurring over the whole temperature domain. Hence, the Archard wear coefficient in case of C-LMD decreases with the increase of the Vickers hardness. Therefore, the Archard theory is still applicable over the entire temperature domain. Hence, Archard theory appears only valid for oxidational-abrasive wear process and is no longer satisfied when the glaze layer is formed since a major part of the friction work is consumed by the debris layer.
<figure><img src="image_23.png" /><figcaption>Figure 22: Evolution of the Archard wear coefficient K as a function of the hardness for the reference alloy (Stellite 6) and the nickel-based alloys (C-LMD and C-GC).</figcaption></figure>

**4. Discussion**

In this study, the objective is to compare the tribological behavior of two nickel-based alloys (C-GC and C-LMD) compared to a reference cobalt alloy Stellite 6 sliding against coated Inconel718 alloy. Results show that the two alloys (C-GC and C-LMD), having the same chemical composition but different microstructures, behaved differently especially at high temperatures. Indeed, C-GC had a low wear volume at all temperatures and favored the creation of a protective glaze layer at high temperature as Stellite 6 alloy. On the other hand, C-LMD has higher wear volume compared to C-GC and Stellite 6 especially at high temperature where severe adhesive wear and material transfer was detected. Hence, the purpose is to investigate the origins behind the difference in wear behavior of these alloys.

As previously reported in the literature [20], the glaze layer generated by cobalt-based alloys as Stellite 6 at high temperature mainly consists of nanograin-spinel of Cr, Ni and Co oxides as $CoCr_2O_4$ combined with Co binder matrix. This leads to homogeneous 2 to 5 µm thin layer which covers the entire rubbed surfaces of the two counterfaces (Figure 23).
<figure><img src="image_24.png" /><figcaption>Figure 23: SEM and EDS cross-section analysis of Stellite 6 revealing the formation of glaze layer at T=650°C (P=23 N, f=20 Hz, δ₀= ±450 µm, and N=200k cycles).</figcaption></figure>

SEM and EDS analysis of C-GC alloy displays a glaze layer structure similar to that detected in Stellite 6 (Figure 24). EDS mapping underlines a high concentration of Cr in the glaze layer and low concentrations Ni element relative to the bulk although this latter element is the dominant in studied materials (Figure 24).
<figure><img src="image_25.png" /><figcaption>Figure 24: SEM and EDS cross-section analysis of C-GC revealing the formation of glaze layer at T=650°C (P=23 N, f=20 Hz, δ₀= ±450 µm, and N=200k cycles).</figcaption></figure>

The SEM-EDS cross-section analysis of C-LMD wear scar displays a discontinuous oxide debris instead of a homogeneous protective glaze layer (Figure 25). This confirms the material transfer process from C-LMD alloy toward the coated Inconel counterpart where thick third body layer having a thickness of 5-10 µm is transferred to coated Inconel counterpart. This thick third body layer has a Ni/Cr ratio around 5 equivalent to the Ni/Cr ratio in C-LMD alloy, confirming hence the EDS surface analysis of Figure 17. By contrast to C-GC, this non protective third body displays a lower concentration of Cr oxides comparatively to Ni ones.
C-LMD

<figure><img src="image_26.png" /></figure>

coated Inconel against C-LMD

<figure><img src="image_27.png" /><figcaption>Figure 25: SEM and EDS cross-section analysis of C-LMD and its coated Inconel counterface revealing the significant transfer process at T=650°C (P=23 N, f=20 Hz, δ₀= ±450 µm, and N=200k cycles).</figcaption></figure>
Micro-Raman analyses of C-GC and its coated Inconel counterface reveal the presence of NiCr₂O₄ spinel in the glaze layer structure with traces of silicon oxides (Figure 26). On the other hand, the discontinuous oxide debris structure of C-LMD and its coated Inconel counterface mainly consists of abrasive NiO and Cr₂O₃ oxides (Figure 26) with traces of NiCr₂O₄ spinel in isolated regions where glaze layer or compact debris bed was able to form (Figure 25). Indeed, several authors in literature argue that nickel-chromium alloys with Cr₂O₃ and NiO oxides, could not generate glaze layer concluding that those latter oxides are inadequately adherent and poorly sinterable [3,21]. However, several studies [16,22,23] reported that when a glaze is created on nickel-based superalloys rich in chromium, it is essentially formed of NiCr₂O₄ spinel due to the reaction between Cr₂O₃ and NiO formed at low temperatures. The driving mechanism leading to the formation and sintering of NiCr₂O₄ spinel out of Cr₂O₃ and NiO oxides is still unclear and is mainly attributed to the presence of other alloying elements in the superalloy as silicon, cobalt and iron for instance facilitating this tribo-chemical interaction. Nevertheless, in the current study the two alloys have the same chemical composition which means that there is an additional factor other than the chemical composition that leads to formation of NiCr₂O₄-rich glaze layer in C-GC and its absence in C-LMD.
<table><caption>Figure 26: Comparison between mirco-Raman analysis of C-GC glaze layer structure and C-LMD non glaze layer structure and their corresponding coated Inconel counterparts (T=650°C, P=23 N, f=20 Hz, δ₀= ±450 µm, and N=200k cycles) [24–27].</caption><thead><tr><th colspan="2">T=650°C</th><th>Point analysis</th><th>Positions on 3D profiles</th><th>Raman spectra</th></tr></thead><tbody><tr><td rowspan="2">C-GC/Inconel interface</td><td>counterface</td><td></td><td></td><td></td></tr><tr><td>C-GC</td><td></td><td></td><td></td></tr><tr><td rowspan="2">C-LMD/Inconel interface</td><td>counterface</td><td></td><td></td><td></td></tr><tr><td>C-LMD</td><td></td><td></td><td></td></tr></tbody></table>

To better assess the origin behind the creation of glaze layer in C-GC, SEM and EDS analyses of the surface and cross-section were performed (Figure 27) to compare the microstructure of C-LMD and C-GC being the only difference between these alloys. Figure 27 exhibits a silicate “tortoiseshell” deposit on C-GC alloy corresponding to the Ni₃Si phase distribution in the alloys, which successively combines with Ni and above all Cr oxides generated in the fretted interface inducing glaze layer formation. Such silicate deposits are absent on the C-LMD alloy specimen which has homogeneous distribution of phases. Silicates, through their viscous and cohesive properties at high temperatures,
have in fact been known for many years as compounds improving the tribological
performance at high temperatures [12,28]. In this case, they appear as an interesting
alternative to replace cobalt as binders for glaze layer activation. This suggests that the
C-LMD through its particular fine microstructure limits the diffusion on silicon on the top
surface. A lower concentration of silicate on the top surface reduces the binder effect so
that the protective glaze layer cannot be formed leading to a catastrophic wear involving
adhesive and abrasive wear phenomena exhibited by aggressive NiO oxides. Besides,
the silicate protrusions formed at surface of the C-GC along with the coarse star-shaped
borides and carbides generate a special surface texture that can facilitate trapping and
cohesion between third body particles which facilitates the formation of glaze layer as
reported by Kong et al. [29]. These two features (i.e. silicate binder effect and the surface
texture) increase sinterability of wear debris, thus promoting the formation of a protective
glaze layer on the contact surface.

<figure><img src="image_28.png" /><figcaption>Figure 27: SEM and EDS analyses of the C-GC and C-LMD surface heated at 650°C (outside the fretting scar).</figcaption></figure>
From this analysis, it can be concluded that protective glaze layers are formed through
the combination of spinel oxides and the presence of a binder which can be cobalt for
Stellite 6 or amorphous silicates for C-GC. When the presence of silicate on the top
surface is not sufficient (i.e. C-LMD), the glaze layer formation is no more possible and
the interface evolves toward a high wear NiO abrasive/adhesive wear process. This
finding is in agreement with Buckley and Johnson [12] results who investigated wear
behavior of binary nickel-silicon alloys by changing the silicon content from 0% to 10%
prepared at NASA Lewis Research Center (now Glenn Research Center). They noticed
that in the absence of silicon, the wear damage of nickel was very high and favored
adhesive wear especially in reducing atmospheres for temperatures exceeding 500°C.
This is not the case when the silicon content is increased from 2.5% to 10% leading to
significant reduction in the friction coefficient and wear rates. They attributed this response
to the vitrification effect of the silicon which is used in alloys as an oxygen "getter" that is
extremely reactive with oxidizers favoring the creation of glaze layer. Similar results were
obtained by Viat et al. [28] who revealed that silicates behave as a binding agent
promoting sintering and adhesion between debris particles through their viscous flow
within the granular media of debris bed.

Another interesting finding is that the higher the concentration of Ni₃Si phase in the alloy,
the higher the silicate concentration on top surface and the easier the formation of the
protective glaze layer. The Ni₃Si phase formation in nickel alloys is a function of the Si
concentration but above all on the Si diffusion in the matrix during the cooling stage of the
manufacturing process. Solidification simulations have been performed using the classical
Scheil's model (which considers segregation during solidification without considering the
cooling rate) for C-LMD manufacturing and a thermal equilibrium simulation for the gravity
casting (C-GC) process (Figure 3). In both cases, it was concluded that the Ni₃Si phase
formation starts at a rather low temperature (below 800°C) comparatively to the other
phases. These computations reveal that the rapid cooling of the LMD process (C-LMD
alloy) tends to freeze the formation of Ni-Si phase in contrast to the lower cooling rate
involved in the gravity casting (C-GC alloy) which extends the formation of eutectics.
These simulations confirm former microstructural expertizes underlying a significantly
higher concentration of Ni₃Si phase in C-GC comparatively to C-LMD (Figure 27). Hence,
by adjusting the cooling rate, it is possible to optimize the formation of Ni₃Si phase which
itself allows silicate diffusion on the top surface leading to the formation of a protective
glaze layer when this interface is fretted under high temperature. If the concentration of
Ni₃Si phase is too low due an insufficient concentration of Si element or a too fast cooling
rate (i.e. C-LMD), the surface concentration of silicate at high temperature becomes too
low and the protective glaze layer cannot be formed (C-LMD).

It should be noted that C-GC is also characterized by coarse Ni₃B phase compared to C-LMD. However, several studies revealed that Ni-B alloys do not show satisfactory wear behavior at high temperature [30,31]. Mukhopadhyay *et al.* [30,31] tested at temperatures ranging from ambient to 500°C Ni-B coating containing 5-5.5% of boron, with the presence of Ni₃B and Ni₂B microcrystalline phases. They revealed that the tribological behavior of the Ni-B coating worsens at 100°C compared to room temperature, the it displays better wear resistance at 300°C which again deteriorates at 500°C due to severe oxidation and softening of the deposits. Additionally, they showed that boron is not detected in the worn surface of the coatings at 500°C since at elevated temperatures, the boron atoms diffuse
to the surface and evaporate after reacting with oxygen due to the creation of boron oxides that evaporate above 450°C [30,31]. This is not the case of silicon oxides having much higher melting and boiling points. Another study [30] revealed that the addition of molybdenum or tungsten to Ni-B coatings (i.e. Ni-B-W or Ni-B-Mo) enhances the tribological behavior of the alloy especially for Ni-B-Mo due to the formation of a protective glaze layers at 500°C. This indicates that boron alone cannot lead to enhanced tribological performance of nickel alloys at high temperature and it needs additional agent that favours the creation of glaze layer which is silicon in the light of this study. Additionally, Yoo *et al.* [32] investigated the effect of boron content on the high-temperature wear resistance of FeCrCSi alloys. They noticed that alloys with more than 1 wt% boron showed significant wear loss than alloys with lower boron content at high temperatures due to the presence of brittle FeB particles in the microstructure that facilitate crack initiation leading to spalling during sliding.

**5. Conclusions**

This study focuses on the wear behavior of the cobalt alloys (Stellite 6) and two NiCrSiFeB alloys (C-LMD and C-GC) fretted against a coated Inconel718. The effect of the temperature and the number of cycles were investigated and the following points are outlined:

* The wear mechanics are strongly influenced by the temperature (from 20°C to 650°C). At low temperature, the wear mechanism combined an abrasive and an oxidative process, leading to a severe wear. When the temperature exceeds a specific
transitional temperature, wear decreases with the formation of a protective glaze layer
that leads to a no-wear domain.

* Stellite 6 and C-GC are characterized by an abrasive oxidational wear at temperature lower that 300°C where Archard wear law is valid and the wear rate and the friction coefficients are high in this domain. For T>300°C, glaze layer starts to form in these alloys leading to decrease in the wear rate and the friction coefficient until no wear is detected at 650°C. Due to the formation of this tribofilm, Archard law is no longer valid in these alloys for T>300°C.

* C-LMD has the worst wear behavior at both ambient and high temperature as it is unable to form a protective glaze layer, so it wears a lot compared to the other alloys at high temperature. Since oxidational-abrasive wear prevails at all tested temperatures (25°C<T<650°C), Archard law remains valid for C-LMD. This is explained by the important material transfer from alloy C-LMD to the coated Inconel presenting the creation of a protective glaze layer. This result becomes worse when C-LMD is tested against uncoated Inconel where the wear rate becomes very high and a negative wear volume is obtained in the uncoated Inconel counterpart marking material transfer and adhesive wear.

* The protective glaze layer is made of nano-grain spinel oxides agglomerated thanks to a binder. For Stellite 6 the binder is cobalt, for C-GC the binder is a silicate structure. Whatever the composition of the new NiCrSiFeB alloys, if the presence of silicates on the top surface is not sufficient, the glaze layer cannot be formed and the wear rate under high temperature becomes catastrophic (C-LMD).
* The presence of silicates results from an oxidation-diffusion process of the eutectic Ni₃Si phase embedded in the microstructure. Whatever the NiCrSiFeB alloys, the higher the proportion of Ni₃Si phase the higher the silicate concentration on the top surface. The proportion of Ni₃Si phase for a given [Si] concentration depends on the cooling rate which must be slow enough (i.e. GC process) to allow the Si diffusion within the matrix. If it is too fast (i.e. LMD process) the Ni₃Si phase proportion is too low, the silicate surface deposit cannot be formed and finally the protective glaze layer cannot be generated at 500°C.

* This investigation also underlines that the correlation between hardness and wear resistance is no more satisfied as soon as the third body layer plays a key influence in the friction dissipation process so that above 300 °C. Hence, hardness cannot be considered to extrapolate the wear performance of a NiCrSiFeB friction interface for high temperature applications.

* Ni₃Si phase has a beneficial effect on wear performance at high temperature thanks to the formation silicate binder. For each application, a balance must be established between tribological and mechanical properties to optimize the concentration of Ni₃Si phase.

* The exceptional performance of Stellite 6 alloy at high temperature can be explained by the fact that Co metal is an optimal binder to form protective wear glaze layer and in the meantime provides high toughness properties preventing fracture phenomena. This explains why Stellite 6 cobalt alloys are so difficult to replace in industrial high temperature applications.
## 6. Acknowledgement

The authors would like to thank Olivier Graton at Laboratory of Tribology and System Dynamics (LTDS) for his help in the instrumentation part of the experiments. Besides, the authors extend their gratitude to Louis Grillot, Evan Russo, Mohammed Fartas, Tabasser Hossain, and Laulade Quesne being the interns working on this project at LTDS Laboratory. This European research project called NEMARCO (http://nemarco.eu/) received funding from the Clean Sky 2 Joint Undertaking (JU) under the grant agreement No. 101007948. The JU receives support from the European Union's Horizon 2020 Research and Innovation Program, and the Clean Sky 2 JU members other than the European Union.

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