# Scenario for the formation of fretting obtained TTS in TA6V from detailed microstructural analysis

Vivien Lefranc, Camille Gandiolle, Maxime Vallet, Julie Bourgon, Eva

Héripré, Siegfried Fouvry, Véronique Aubin

► To cite this version:

Vivien Lefranc, Camille Gandiolle, Maxime Vallet, Julie Bourgon, Eva Héripré, et al. Scenario for the formation of fretting obtained TTS in TA6V from detailed microstructural analysis. Wear, 2025, 571, pp.205834. 10.1016/j.wear.2025.205834. hal-05378039

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https://hal.science/hal-05378039v1

Submitted on 25 Nov 2025

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Distributed under a Creative Commons Attribution 4.0 International License
Scenario for the formation of fretting obtained TTS in TA6V from
detailed microstructural analysis

Vivien Lefranc<fnref n="1" />,<fnref n="2" />, Camille Gandiolle<fnref n="1" />,*, Maxime Vallet<fnref n="1" />,<fnref n="3" />, Julie Bourgon<fnref n="4" />, Eva Heripré<fnref n="1" />,<fnref n="5" />,
Siegfried Fouvry<fnref n="2" />,<fnref n="6" />, Véronique Aubin<fnref n="1" />

*Corresponding Author: camille.gandiolle@centralesupelec.fr

**Affiliations:**

**Keywords:** Fretting, Tribologically Transformed Structure, TEM analysis, Crystal plane indexing, continuous dynamic recrystallization, TA6V.

**Highlight**

- TEM crystal texture analysis of each layer of heterogeneous TTS

- Layers of TTS (I) show strong basal texture

- Layers of TTS are isotropic with a diffuse δ-TiN phase

- TTS is formed by continuous dynamic recrystallization

- Continuous wear explains heterogeneous TTS

**Abstract**

Fretting is defined as surface degradation caused by small oscillatory movements. Fretting in the gross slip regime produces high local stresses at the surface. This results in wear and subsurface plastic deformation, which can create a nanosized-grained structure known as a Tribologically Transformed Structure (TTS). If the condition of TTS formation in the fretted interface is well documented in the literature, the metallurgical transformation leading to this structure is not well known. This article aims to study this phenomenon thanks to an innovative in-depth texture analysis of the TTS which is carried out using a TEM crystal plane indexing

<footnote n="1">¹ Université Paris-Saclay, CentraleSupélec, ENS Paris-Saclay, LMPS - Laboratory of Mechanics Paris-Saclay, France</footnote>

<footnote n="2">² Laboratory of Tribology and System Dynamics, Ecully, France</footnote>

<footnote n="3">³ Laboratory of Structures Properties and Modeling Solids, Gif sur Yvette, France</footnote>

<footnote n="4">⁴ Univ Paris Est Créteil, CNRS, ICMPE, UMR 7182, 2 rue Henri Dunant, 94320 Thiais, France</footnote>

<footnote n="5">⁵ PIMM, Arts et Metiers Institute of Technology, CNRS, Cnam, HESAM Universite, France</footnote>

<footnote n="6">⁶ Mines Paris, PSL University, MAT - Centre des Matériaux, France</footnote>
technique. Several TTS from different steps in the TTS formation are studied. Special care is
taken in determining their texture, and chemical analysis is carried out. Considering the
temperatures involved and the successive structures analyzed, a new model of continuous
dynamic recrystallization via the apparition of shear bands is proposed to explain the
heterogeneity of TTS and its microstructural characteristics.

**1. Introduction**

Fretting occurs when two parts in contact are subjected to slight alternating movements. If the sliding amplitude is such that the entire contact is sliding without the whole surface being exposed, then this is a full-slip fretting regime. This regime is known to generate significant surface wear, which can have a major impact on structural strength [1]. Fretting in the gross slip regime induces high local shear stresses on the top surface. This results in subsurface plastic deformation, which can create a nanosized-grained structure known as a Tribologically Transformed Structure (TTS) [2]. TTS is usually formed from the extreme surface down to a few dozen microns depending on the contact loading and contact size. Indeed, the gross slip fretting stress field is characterized by very severe decreasing stress gradients.

The microstructure of TTS in titanium alloys has been studied in various articles. It consists of α-nanograins smaller than 100 nm [3,4]. TTS is formed through extreme refinement of the grains regardless of their original size, as well as an α+β to α phase transformation even for β alloys [5]. In a previous article, we defined the conditions for the appearance of TTS as a function of the pressure exerted during fretting and the number of cycles for a TA6V alloy under planar conditions [6]. One of the main conclusions was that the TTS was an heterogeneous structure alternating a nanosized and textured TTS (I) with a finer isotropic TTS (II). However, these elements are not sufficient to understand the metallurgical transformations at the origin of TTS formation. For this, it would be necessary to characterize the crystallographic texturing beneath the contact. To date, there is no study looking specifically at this point in fretting TTS, as conventional means (e.g. texture goniometer and EBSD) lack the necessary resolution.

However, some authors have taken an interest in textures on fretted surfaces. For example, V.D. Scott and H. Wilman have quantified the orientation of crystallographic textures by electron diffraction of surfaces obtained for different abrasive wear loads on polycrystalline beryllium and magnesium samples [7], two metals with a compact hexagonal structure and behaviors close to that of α-titanium. They observe an oblique fiber texture of the basal plane with an orientation (0001) inclined at an angle of approximately δ = 21° to the normal to the rubbed surface for beryllium and δ = 22° for magnesium. This angle appears to be related to the coefficient of friction, as tan-1(μ) ≈ δ is observed. They also observe a second extreme surface texturing in which the normal to the basal plane is oriented parallel to the normal of the rubbed surface, allowing alignment of the hexagonal sliding planes with the direction of friction. Swalla et al. using EBSD, observed basal microtexturing at a depth of between 30 µm and 50 µm, which is very pronounced under gross slip fretting conditions on TA6V [8]. In this study, the normal to the basal plane is parallel to the normal of the extreme rubbed surface. It should be noted, however, that these results were obtained on plastically deformed layers and not on TTS. Indeed, since TTS grains are very small (in the nanometer range), it is impossible to characterize them using conventional techniques such as EBSD.

This article aims to fill this gap by completing the microstructural analysis of the nanostructural
TTS structure. A study of the microstructure of TTS under various fretting pressure was carried
out using TEM. Four TEM samples were taken under initial pressure conditions of 200MPa
and 300MPa. Of these, one “punctual” TTS surrounded by a deformed layer and three
“massive” TTS zones were selected and extracted using a SEM-FIB and studied by coupling
various TEM analyses: imaging, Energy Dispersive X-ray Spectroscopy (EDX), Electron
Energy Loss Spectroscopy (EELS) and automated crystal orientation indexing technique under
TEM, called Astar. From this analysis, a possible metallurgical mechanism for the formation of
TA6V TTS generated by fretting is derived.

## 2. Material

The study material is a TA6V alloy. It was held at a temperature in the α-β range under transus-
β and then quenched with water. Finally, it was tempered at a temperature of 700 °C. This α-β
Ti-6Al-4V alloy has been extensively investigated in many fretting wear studies [9,10] and, in
particular, in a study on the conditions of TTS formation as a function of pressure and cycle
number [6]. Its elastic mechanical properties are E= 119GPa, v=0.29, σ<sub>y</sub>= 970MPa.

The microstructure of the initial bulk, α-β Ti-6Al-4V alloy, is characterized by EBSD on a Zeiss SUPRA 55VP SEM-FEG with an EDAX-OIM CMOS EBSD camera. The acquisition step is 0.1 µm. Data are post-processed with EDAX-OIM software. Figure 1 shows a color mapping of the inverted pole figure (IPF//TD). An area of 100µm by 100µm is extracted from the mapping and shown in Figure 2.

<figure><img src="image_2.png" /><figcaption>Figure 1: IPF//TD mapping of the initial microstructure of TA6V. (SD) fretting direction, (ND) direction normal to the rubbed surface, and (TD) transverse direction.</figcaption></figure>

<figure><img src="image_3.png" /><figcaption>Figure 2: IPF extracted from IPF/TD mapping in Figure 1</figcaption></figure>
A bimodal structure can be seen in Figure 2, with alpha nodules and alpha Widmanstätten needles inherited from the beta phase transformed during quenching. The mean equivalent circle diameter of the alpha nodules is 17.15 µm with a standard deviation of 6.39 µm. The pole figures derived from Figure 1 mapping of the α phase are shown in Figure 3, demonstrating that there is no strongly marked crystallographic preferential orientation. The initial material has no texture, even from potential surface machining to generate the samples.

<figure><img src="image_4.png" /><figcaption>Figure 3: Pole figures derived from IPF//TD mapping in Figure 1.</figcaption></figure>
## 3. Methods

### a. TTS from fretting loading

The first step of the study consisted in creating some TTS by fretting loading. A fretting test apparatus mounted on an MTS hydraulic machine was used [11] similar to the one from [6]. Tangential displacement and normal force were controlled throughout the test for constant tangential force and pressure. The tests were performed at ambient temperature (25 °C ± 5 °C) and relative humidity (RH=40 % ± 10 %). The frequency was fixed at 10 Hz, and the sliding amplitude δ₀ at 100 µm to remain in gross slip conditions in all the configurations. It resulted in a sliding speed of 4 mm/s. The bottom and top samples were bars of Ti-6Al-4V, 4 mm wide, with the top sample being set perpendicular to the bottom sample. It resulted in a square contact area of 16 mm² (4x4 mm²). From a previous study of the pressure and number of cycles necessary to form TTS [6], 3 specific loading conditions were selected allowing us to study both some localized and some massive TTS, the second type with varying number of cycles. The loading conditions are given in Table 2.

<table><caption>Table 2: Loading conditions for each observed TEM Sample</caption><thead><tr><th>TEM Sample</th><th>A</th><th>B</th><th>C</th><th>D</th></tr></thead><tbody><tr><td>Pressure (MPa)</td><td>200</td><td>300</td><td>300</td><td>300</td></tr><tr><td>Cycles</td><td>20000</td><td>1000</td><td>1000</td><td>2000</td></tr><tr><td>Observed structure</td><td>Local TTS and GDL</td><td>Massive TTS</td><td>Massive TTS</td><td>Massive TTS</td></tr><tr><td>Sample orientation</td><td>(YZ)</td><td>(XZ)</td><td>(XZ)</td><td>(YZ)</td></tr></tbody></table>

### b. Cross-section analysis

After the fretting tests have been carried out, the samples were cut in the direction of friction and polished to gain access to the center of the trace. Chemical etching with Kroll's reagent ensures that the TTS has been formed. TTS are resistant to etching with Kroll's reagent and appear white under the optical microscope.

Figure 4 illustrates the cross-section after chemical etching of the fretting scar obtained under 300 MPa contact pressure and 2000 fretting cycles. It was taken at the center of the contact parallel to the sliding direction. In the inner part of the contact a uniform white layer, up to 60 µm in thickness, the Tribologically Transformed Structure (TTS), is seen (figure 4b). On the lower part of the Figure, unaffected by the fretting load, the undeformed bulk material is found. Between the TTS and the bulk, we find the General Deformed Layer (GDL) [6] which has a thickness of 20 µm to 30 µm. On the outer part of the surface, a Third Body Layer (TBL) of thickness ranging from 5 to 15 µm is observed and is composed of agglomerated power debris (Figure 4c and 4d).
<figure><img src="image_5.png" /><figcaption>Figure 4: Cross section of the fretting test with a contact pressure of 300 MPa and 2000 cycles. Optical micrograph (a) of the overall contact, (b) with a zoom on TTS. (c) Secondary electron image of the outer contact zone with (d) a zoom on TBL. [6]</figcaption></figure>

c. Extraction of TTS samples for TEM observations

There are several methods for TEM sample extraction and sample preparation. One of them is extraction and refining by Focused Ion Beam (FIB) which is used throughout the study [12]. The area of interest is first covered with a 1 µm layer of platinum. Then trenches are dug by FIB to extract the sample with a micromanipulator. It is then soldered by platinum deposition on a copper TEM grid. Finally, the sample is refined and polished by decreasing the ion acceleration voltage down to 5 kV and 15 pA to achieve an excellent surface finish. This technique allows, among other things, very good localization of the area to be analyzed, as well as good control of the sample manufacturing processes.

TEM samples were taken from the cross-section of the fretted sample, after polishing and etching with Kroll's reagent using a HELIOS 660 Dual-Beam SEM-FIB. Four FIB-machined TEM samples, labelled A to D, were cut in several orientations and on three different fretting scars. Sample fretting parameters, sampling location and plane parallel to the sliding plane are summarized in Table 2. The orientation indicated in the table is explained in Figure 5 in relation to the fretted surface and the sliding direction.
<figure><img src="image_6.png" /><figcaption>Figure 5: Representation of the orientation of the TEM samples in relation to the fretting wear scars.</figcaption></figure>

The “local” TTS extraction zone of sample A was chosen to provide an analysis of an
intermediate state of TTS formation as defined in [6]. The other samples (B to D) were
extracted at 300 MPa and a low number of cycles to study a “massive” TTS that is still largely
free of cracks.

d. Analyses

Bright-field observations in TEM mode were carried out on a JEOL JEM-2100Plus TEM with an accelerating voltage of 200 kV and equipped with a Gatan CCD CMOS RIO 16 4k camera. Diffraction patterns to identify the phases present were produced on the same instrument and post-processed on Gatan's DigitalMicrograph software with the DiffTools plugin [13]. In addition to the above analyses, chemical EDXS analyses were carried out on a TEM TITAN³ G2 80-300 (FEI-ThermoFisher) with a Super-X detector. Post-processing was carried out using Bruker's Esprit software and the Cliff and Lorimer method [14].

Deconvolution of Vanadium peak in EDXS analysis is not easy for TA6V. Indeed, the main energy loss peak of Titanium is at the same value as that of Vanadium. Quantifying nitrogen is also difficult for similar reasons. To determine whether vanadium and nitrogen are present in TTS, complementary electron energy loss spectroscopy (EELS) analyses were carried out on the Titan 300 kV TEM equipped with the Gatan Enfinium spectrometer. EELS enables more reliable quantification of light elements such as nitrogen. The positions and intensities of the energy loss peaks are taken from [14].
For crystallographic texturing analysis, a TEM-based crystal plane indexing technique known
as ASTAR, developed by E. Rauch et al. [15], was used. The principle of Nanomegas' ASTAR
technique involves a device mounted on a TEM, that precisely controls the electron beam in
order to scan a TEM sample and make a grid of diffraction patterns in precession mode,
enhancing resolution. These patterns are then compared, using software, with simulated
diffraction patterns from known phases of the material, analyzed in different orientations. The
most probable orientation is selected and linked to a grid point. By integrating the orientations
of the points, a grain orientation map similar to SEM EBSD, but now transposed to a TEM high
resolution, can be reconstructed. This technique requires the TEM diffraction images to
correspond to a single crystal. Data are post-processed using OIM software. Astar analyses of
our samples were carried out on a FEI Tecnai F20 TEM fitted with Astar equipment at ICMPME
Thiais.

4. Results of the analysis of microstructure and chemistry with TEM

Samples were imaged by TEM and analyzed by diffraction pattern. Samples B, C and D were
further analyzed chemically.

a. TEM microstructure analyses of samples

TEM sample A was taken from the punctual TTS formed at 200 MPa and 20,000 cycles. TEM imaging of this sample has already been presented in [6] but is re-submitted here for ease of analysis later. Its sampling location is shown in Figure 6a, the TTS is below a thick third body layer and at a depth of 15 µm from the surface. Figure 6b shows the overall appearance of the TEM sample. Two very distinct zones are observed: the general deformed layer (GDL) with larger grains, between 70 nm and 150 nm in diameter, and the localized TTS. The latter is around 2 µm thick and consists of grains between 20 and 50 nm in diameter.

The TEM B sample was extracted from the massive TTS formed at 300 MPa after 1000 cycles, which is around 60 µm thick at this point in the trace. Its sampling location, directly at the rubbed surface, is shown in Figures 7a and 7b. Figure 7c shows the overall appearance of the sample. Two types of alternating microstructures are observed, each about 1µm thick: TTS (I) and TTS (II). One difference between TTS (I) and TTS (II) concerns the grain size: TTS (I) grain size ranges from 20 to 50nm, while TTS (II) displays grain size systematically smaller than 20nm.
<figure><img src="image_7.png" /><figcaption>Figure 6: Sample A. a) BSE image of the extraction zone of sample A after chemical etching, 15 microns from the rubbed surface. The location of the sample is represented by the red rectangle. b) Bright-field TEM image of the entire sample. [6]</figcaption></figure>

The diffraction patterns associated with each of the 4 zones: GDL and TTS of sample A and TTS (I) and TTS (II) of sample B are shown in Figure 8. In all these structures, only the diffraction circles corresponding to the α phase appear. There is no β-phase grain. The GDL and the TTS consist only of α grains. The diffraction circles of GDL (Figure 8a), TTS (Figure 8b) and TTS (I) (Figure 8c) are discontinuous, suggesting the presence of crystallographic texturing. By contrast, this texturing is not observed in TTS (II) since continuous circles, representative of untextured material can be observed (Figure 8d). Similar results are observed in samples C and D. Based on these observations, the localized TTS of sample A is considered to be similar to the TTS (I), which is why this sample was not analyzed further.
<figure><img src="image_8.png" /><figcaption>Figure 7: Sample B. a) Optical image and b) BSE image of the extraction zone of sample B after chemical etching. The location of the sample is represented by the red rectangle. c) Bright-field TEM image of the entire sample. [6]</figcaption></figure>
<figure><img src="image_9.png" /><figcaption>Figure 8: Diffraction patterns of a) GDL in sample A, b) TTS in sample A, c) TTS (I) in sample B and d) TTS (II) in sample B.</figcaption></figure>

Besides, TTS (II) also reveals a diffraction circle between circles (1010) and (0002) that does not seem to appear in the other diffraction patterns and does not correspond to a β-phase. It was also observed in the other samples presenting TTS (II), that is samples C and D.

To better observe this anomaly, Figure 9 plots the diffraction spectrum obtained from a
rotational averaging of the intensity around a central point (SAED) taken from the local TTS
(A-TTS) and TTS (II) (B - TTS (II)) diffraction patterns (obtained using Difftools [13]). The position of the α-phase peaks is indicated by black dotted lines. Red dotted lines indicate peaks corresponding to a potential TiN-δ phase, with a face-centered cubic structure. The additional contribution between circles (1010) and (0002) of the α-phase, observed for TTS (II), appears to coincide well with a fcc TiN-δ phase. This contribution is particularly visible for the circle located at the position of the (200) TiN-δ peak (Figure 9).

This peak is not present in the local TTS of the TEM A sample. In the transformation of the base metal into TTS, we thus have, in addition to the transformation of β into α, the appearance of δ-TiN precipitates in the TEM B sample within TTS (II). It would appear from further study of all diffraction spectra, that TTS (I) is also partly made up of TiN, but this seems less significant than for TTS (II), as the TiN circles are less intense when viewed in Figure 8b and c. It, therefore, seems important to complete this investigation, hence the EDXS and EELS chemical analyses carried out.

<figure><img src="image_10.png" /><figcaption>Figure 9: Diffraction pattern of (a) local TTS formed at 200 MPa, (b) massive TTS (II) formed at 300 MPa, and their diffraction circle intensity spectra (c) and (d) respectively. The location of the Ti-α and TiN-δ phase peaks are shown in dotted black and red respectively.</figcaption></figure>

b. Chemical analysis

Results in this paragraph come from sample C which was extracted from the same fretting scar as sample B and presents similar observations to samples B and D.

EDXS and EELS chemical analyses were carried out to confirm the presence of TiN. The area analyzed by EDXS is shown in Figure 10. It includes both TTS (I) and (II). Figure 10(e) shows
the EDXS map of nitrogen concentration. The entire observed zone shows nitrogen, but there
is inhomogeneity of nitrogen concentration in TTS (I) and TTS (II). From the color intensity, it
appears that TTS (II) has a higher average nitrogen content than TTS (I), even though TTS (I)
appears to contain a nitrogen-rich precipitate (which is not observable in ADF images). While
these images seem to confirm the hypothesis of the presence of TiN precipitates in TTS (I) and
(II), with a higher distribution in TTS (II) than in TTS (I), it is difficult to obtain a reliable
quantification of nitrogen in this material using this method. A high-resolution EELS analysis
was carried out on the same zone and is shown in Figure 11. In the titanium and nitrogen
concentration maps, the whiter the pixel, the higher the element concentration. The same trend
can be observed as with EDXS: TTS (II) is richer in nitrogen than TTS (I). The latter, however,
shows a nitrogen-rich precipitate in the observed zone.

Spectra, representative of each TTS, are plotted in Figure 11. The spectrum for TTS (I)
corresponds to the zone within the blue rectangular, while the spectrum for TTS (II)
corresponds to the red rectangular.

The nitrogen K peak is located at 402eV, while the titanium L<sub>2,3</sub> peaks are at 456 eV and 462 eV. Thus, unlike EDS, the titanium and nitrogen peaks are separated in EELS, enabling precise quantification of the species present. The peak corresponding to nitrogen is not present in TTS (I) at around 400 eV, whereas it is well-defined in TTS (II). Quantification of Ti and N elements gave 1.7%at nitrogen in TTS (I) and 17%at in TTS (II), confirming the apparent heterogeneity of concentration. The value of 1.7%at should be treated with caution, as it is close to the resolution limit of the method but seems to confirm the low nitrogen concentration of TTS (I) excluding the nitrogen-rich precipitate.

<figure><img src="image_11.png" /><figcaption>Figure 10: Sample C. (a) Bright-field image of MET sample. (b) ADF image of the white-framed area in Figure (a), with the EDS analysis zone framed in orange. (c) ADF image, along with EDS mappings of (d) titanium and (e) nitrogen of the analysis zone. (I) indicates TTS (I) layer and (II) indicates TTS (II) layer.</figcaption></figure>
<figure><img src="image_12.png" /></figure>

Figure 11: Sample C. Center: ADF image of the orange-framed zone in Figure 10b associated with Titanium and Nitrogen EELS mapping. The electron energy loss spectra corresponding to the blue-framed zone (I) of TTS (I) and the red-framed zone (II) of TTS (II) are on the left and right respectively. The EELS quantification of the elements titanium and nitrogen can be found in the table on each spectrum.

TTS (I) therefore appears to be a structure made up of 20-50 nm α-titanium grains with large TiN precipitates. TTS (II) is a structure with α grains of less than 20 nm and probably the same proportion of TiN precipitates, but the latter are much smaller and diffusely distributed throughout the material, resulting in this impression of a nitrogen-rich structure. However, the density of TiN-δ precipitates needs to be put into perspective. Nitrogen is already highly soluble in the α phase of titanium, so this δ phase will always be in the minority compared with the α phase.

### c. Discussion on chemical aspects

This is not the first time that nitrides have been seen in TTS. The observation of the TiN component in TTS generated by fretting in large Ti6-Al-64V contact was previously documented by Mary et al. [16]. The precipitation of this face-centered cubic δ-TiN phase in massive TTS suggests that the material reacts with nitrogen from the air during fretting. Indeed, the TA6V alloy initially has a nitrogen composition of less than 0.05% by mass. From Mary et al. [17], assuming an "air distilling" process, di-oxygen is consumed to form oxides on the outer part of the contact and is less abundant in the inner part of the contact. There, the nitrogen partial pressure is higher and di-nitrogen molecules interact with the reactive titanium metal resulting in the formation of TiN and other nitride components. This hypothesis was later demonstrated by Baydoun et al. by coupling an Advection Dispersion Reaction modelling of the gaz atmosphere operating within a fretting interface [18,19]. Shipway et al. [20] explain that TTS will only form in an oxygen-deprived environment conducive to adhesive phenomena. In the case of the present study, a nitrogen-rich local atmosphere can also be observed at the center of the contact. This also explains why TTS does not form on contact edges, where oxidized debris can form and accommodate stress [6]. It is highly likely that this TiN precipitation will have an impact on the properties of massive TTS. Indeed, TiN is a refractory ceramic often used in the form of coatings on high-speed tools. Its high hardness, around 2400 HV, makes it highly resistant to abrasive wear. TiN is also very brittle. In addition to the presence of TiN precipitate in TTS, it can be assumed that nitrogen is present as a solid solution. In titanium, nitrogen is an alphagenic alloying element that is located on the octahedral interstitial sites of the α-hexagonal phase. It therefore increases the stability of the α phase. Nitrogen has a hardening effect on titanium: it causes a decrease in ductility and an increase in yield strength, as well as embrittlement even at low nitrogen levels. To summarize, the addition of nitrogen to TTS will contribute to its high hardness and brittleness by solid solution and structural hardening through the formation of TiN precipitates.

### 5. Crystallographic texture analysis
a. Astar results

An area containing TTS (I) and TTS (II) from sample B was mapped as an inverse pole figure using AstaR (Figure 12a). A grey level indicating the confidence of measurement is superimposed on the color image. This index defines the accuracy of the match between simulated and measured orientations. The first thing we notice is that TTS (II) has a low confidence index, indicating that this area has been incorrectly indexed. The same issue can be seen in Figure 12b, which shows the analysis of the TEM C sample. It seems that the grain size in this layer is too fine to be indexed correctly, even with Astar. Several grains diffract at each measurement point, blurring the identification of orientation.
<figure><img src="image_13.png" /><figcaption>Figure 12: Sample B. Inverse pole figure (from high-angle annular dark-field (HAADF) image) of the red-framed zone on the BF TEM image of the TEM Sample on the left. The IPF confidence index is given by the grey level. (a) Sample B, (b) Sample C.</figcaption></figure>

TTS (I), with its larger grains, is indexed with good confidence. Some grains are elongated in the sliding direction (SD) direction (Figure 5). The equivalent grain size circle diameter is 27
nm with a standard deviation of 22 nm, indicating strong grain size heterogeneity. The pole figures associated with TTS (I), for each sample, are shown in Figure 13. A strong basal fiber texture is observed, confirming the texture observed by diffraction. A very pronounced texture is found in the (0002) plane figure.

<figure><img src="image_14.png" /><figcaption>Figure 13: Pole Figure of textures from TTS (I) from Figure 12 in (a) Sample B and (b) Sample C</figcaption></figure>

The normal to the basal plane has a preferential angle of around 30° to the surface normal for both samples. Texturing due to the azimuthal rotation of the hexagonal mesh is less pronounced. A schematic representation of the preferential orientation found using pole figures is shown in Figure 14.

<figure><img src="image_15.png" /></figure>
Figure 14: Schematic of the preferential orientation of the α-cell of titanium in massive TTS with respect to fretting solicitation (Figure 5).

b. Discussion

Metals with an hexagonal structure, such as α titanium, have 5 sliding systems that describe the sliding of dislocations. For titanium, the preferred system is the prismatic plane. Basal fiber texture is a type of texture also found during simple shear loading [21,22]. It therefore seems consistent to link the observed texture to the tangential fretting force applied to the surface. Moreover, this texture is also in line with the proposal of Scott et al. [7], who introduced the formula tan-1(μ) = δ, where δ is the angle of the normal to the basal plane with respect to the normal to the surface being fretted, and μ is the coefficient of friction. The stabilized coefficient of friction μ, for the fretting tests carried out in this study, is around 0.55. Thus, the above formula gives δ = 29°, which is consistent with the orientation of the texturing found in TTS (I). This texturing is undeniably the result of major plastic deformations due to the high shear stress generated by fretting sliding. One of the main differences between fretting and simple shearing is that fretting is associated with a strong gradient. It is therefore surprising to find a layered structure within massive TTS, with alternating TTS (I) and TTS (II).

Sauger et al. [2] proposed a model for the formation of TTS based on recrystallization. The severe plastic deformations of fretting create high dislocation densities at the surface, which would allow TTS nuclei to recrystallize. Sauger describes TTS seeds growing progressively from the surface to form grains of a different phase from the base material. However, this model would give a homogeneous structure and does not seem consistent with the microstructure encountered in our study. Furthermore, nucleation and growth of grains under mechanical loading are defined by the community as dynamic discontinuous recrystallization. Such a phenomenon requires a temperature relatively high (above half the melting point) which is not the case in this study due to the low sliding speed. Even if the temperature rises during fretting, it would not be enough for dynamic discontinuous recrystallization to happen [23,24].

c. New hypothesis for TTS formation

A mechanism that seems more consistent with our analyses for the formation of TTS is
continuous dynamic recrystallization. It consists of the formation of new grains by the
progressive disorientation of grain sub-joints resulting from an accumulation of dislocations
during deformation [23]. Continuous dynamic recrystallization has been observed during cold
loading (as fretting) and for metals with high stacking fault energy, such as TA6V [23–26].
Deformation induces progressive grain refinement by continuous dynamic recrystallization until
a plateau is reached where the grain size saturates. This plateau is reached at very high plastic
deformations (ε = 30). This could explain why the structure formed at 200 MPa consists only
of TTS (I), as the strains are not sufficient to reach the grain size of TTS (II). The grain size
saturation found in the TTS (II) zone is also in line with this model. However, if we limit
ourselves to this explanation, it is not possible to accurately represent the heterogeneous
nature of TTS layers (I) and (II) of the massive TTS. For this reason, we propose that
deformation is localized in the form of shear bands. There are studies showing the appearance
of shear bands at low sliding speeds similar to those in our test, leading to the refinement of
the microstructure localized below the surface [27].

The following scenario (shown in Figure 15) is then considered for the formation of TTS from
the generally deformed layer (GDL):

- Plastic deformation induced by fretting stresses leads to an increase in dislocation density and induces weak local disorientation. These dislocations form dislocation cells. The stress
also allows the grains to orient themselves to form a fiber texture.

- Deformation localizes and forms shear bands at a shallow depth below the contact.

- Disorientation within the grains increases because the localized shear increases the disorientation between dislocation cells formed within the grains.

- The formation of sub-joints with high disorientation leads to grain fragmentation and microstructure refinement (continuous dynamic recrystallization).

This process explains the heterogeneous structure of TTS (I) and (II) encountered during our study with TTS (II) being formed in the shear bands. One of the limitations of this explanation is related to texturing. Indeed, in continuous dynamic recrystallization models, texture intensity increases with plastic deformation, whereas untextured TTS (II) is observed at the more advanced recrystallization stage. Moreover, the heterogeneous microstructure of the massive TTS means that the observed state is a transient one. If strain is further increased, deformation should localize in TTS (I) due to the difference in grain size between TTS (II) and TTS (I). TTS (I) should then be refined until the grain size of TTS (II). However, this phenomenon is not observed. One possible explanation for the absence of a more homogeneous microstructure is that it doesn't have time to form. Indeed, it is possible that the TTS in this heterogeneous state becomes brittle enough to be destroyed during fretting stresses, allowing debris to form. It destroys itself as it forms, and so does not reach this final state. The low thickness of the TTS, which reaches a plateau as a function of the number of cycles [5], seems to confirm this hypothesis. Finally, a last open question concerns the regular thickness from 1 to 3 µm of the alternated TTS (I) and (II) layers (Figure 12) which is still unexplained. Further investigations need to deepen this typical "composite TTS structure" which perhaps may be related to the original Ti-6Al-4V microstructure and grain size.

<figure><img src="image_16.png" /><figcaption>Figure 15: Descriptive diagram of TTS formation phenomena: (a) initial GDL grains are deformed and develop a fibrous texture; (b) deformation is localized by shear bands; (c) localized shear increases grain subjoin disorientation; (d) significant subjoin disorientation allows new grains to form in the TTS (II). The direction of friction is shown in red.</figcaption></figure>

## 6. Conclusion

During fretting, depending on the level of pressure applied, different stages of microstructure transformation are observed. The first transformation stage results in a plastically deformed layer (GDL) with grain sizes ranging from 70 nm to 150 nm, which is already composed entirely of α-phase. As plastic deformation increases, shear bands form. When the rate of dislocation is sufficient, TTS is formed by continuous dynamic recrystallization, first locally, then massively. This mechanism results in extreme grain refinement coupled with an heterogeneous microstructure. Two distinct zones are highlighted: one called TTS (I) with a grain size ranging
from 20 nm to 50 nm, large nitrogen-rich precipitates and a strong basal texture. The other
structure, TTS (II), is made of equiaxed nanometric grains with no texture and grain sizes below
20 nm. The diffuse presence of precipitates of a δ-TiN phase was revealed in TTS (II)
accompanied by a significant increase in nitrogen concentration, with a portion of the nitrogen
in solid solution form in the α-phase. The hypothesis of continuous wear is put forward to
explain why TTS never reaches a stabilized homogeneous state at minimum grain refinement.

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