Keywords: dual-phase steel, microstructure, gaseous environment, sliding velocity, friction, oxidative wear, oxide layer, scratches, nanoindentation. Resume L'acier biphasé, La microstructure, L'environnement gazeux, La vitesse de glissement, La friction

The tribological behavior of dual-phase steels was investigated at the effects of different martensite volume fraction (MVF), gaseous environments, and sliding velocities, using a ball-on-disk tribometer, under constant load and dry friction sliding conditions. Furthermore, the tribological behavior of oxide layers formed at different temperatures and MVF was investigated by nanoindentation and scratching tests Dry sliding experiments demonstrated that lower MVF, higher sliding velocities, and CO2 environments significantly reduce wear rates due to improved oxidation as well as the formation of protective oxide layers on the worn surfaces. Scratch tests of oxide layers revealed that the wear rate of same composition oxide layers decreased with increasing MVF. The wear resistance of the FeO-dominated oxide layer is the worst. Under the same conditions, the oxide scale formed on martensite has higher specific wear resistance than that formed on ferrite.

For the life of my PhD at Université de Technologie de Compiègne (UTC), I would like to express my gratitude to all who made it possible. I would like to thank the group of Roberval lab to provide the job opportunity from 2018 to 2022. Also, thanks to the China Scholarship Council (CSC) for the scholarship, which allows me to concentrate on scientific research in good living conditions.

First and foremost, my sincere gratitude goes to my supervisor M. Abdeljalil Jourani for offering me the opportunity to study in France and providing me with academic guidance and inspiration throughout the course of this work. He gave me suggestions, encouragement, and support all the time. His patience gave me the power to persevere in the beginning stage of my research. Besides, his guidance and advice helped me to overcome the difficulties in my research process. I was deeply affected by his passion and attitude towards scientific work. His support and encouragement in the realization of this thesis will always be in my heart. I particularly appreciate the time I spent with him in France.

I acknowledge the support from Mme. Isabelle Velluet, M. Loic Escher, and Mme. Caroline Verdari in the operation of the laboratory devices and training. Also, thanks to Mme. Salima Bouvier and M. Ludovic Cauvin for their advice and assistance to me in the thesis stage assessment every year. Without them, this thesis could never have been realized. I would also like to thank my colleague M. Corentin Penfornis, who shares a good cooperation with me and we always have good discussions about the difficulties in the research. I will always remember the friendship. I thank the honorary jury members who have taken time from their busy schedules and agreed to bring their expertise for assessing the present PhD work. I will not forget to thank all my colleagues in my lab, and all my Chinese friends.

Thank you for the happiness we shared together during these wonderful times, and for the help you gave me when I was in trouble, which will be the most unforgettable

IX

List of Figures

Thesis overview

Tribology is the science and technology that investigates relative motion, interacting surfaces, and their related theories, mainly including friction, wear, and lubrication. Tribology is not a directly visible consequence of human activities and is divided into energy consumption due to friction losses and additional resource consumption through wear and failures of machine elements and products. It involves traditional machining, transportation, aerospace, marine, chemical, bioengineering and many other industrial fields. Statistics show that friction consumes approximately 1/3 of the worldwide primary energy, while wear causes about 60% of machine parts failure.

In addition, more than 50% of mechanical accidents result from defective lubrication or excessive wear. In the future, humanity will be under great ecological pressure to efficiently use resources and reduce energy and resource consumption. Tribology can be used to improve the efficiency of mechanical systems, extend their service life, reduce accidents, and provide useful solutions to energy shortages, resource depletion, environmental pollution, and health problems facing the development of human society.

From a socio-ecological perspective, tribology can help to double the utility value while consuming the same resources resulting in an overall reduction in CO2 and greenhouse gases emissions.

Dry friction as a special working formation, due to lots of friction pairs are employed in lubrication conditions, there is still a lack of systematic investigation of it.

This also indirectly impacted the improvement of the fundamental tribological theory.

Recently, with the development of motion parts to high speed, high load, and miniaturization, the traditional lubrication can hardly meet the requirements. Some components that are difficult to lubricate, such as computer magnetic heads and space motion parts, are working under dry friction. Moreover, dirty lubricant oil is a drain on industrial productivity. Meanwhile, environmental pollution from improper handling of lubricants has become a critical issue. In addition, it is also necessary to have a favorable dry tribological property for some friction pairs that work under lubricated conditions. Especially for some friction pairs applied in military or aerospace fields, sufficient dry friction service life must be considered under lubrication failure, in order to achieve the specified running time. Therefore, it can be expected that dry friction will become a research hotspot in the future.

There is a strong system dependence and complex space-time characteristics of dry friction. Tribology behavior is not only governed by the speed, load, temperature, atmosphere, and microstructure, but also by the time and location during rubbing This thesis is divided into 5 chapters:

The first chapter focuses on the research background from the following three aspects: overview of dual-phase low-alloy steels, effect of microstructure on tribological behavior, and investigation of oxidative wear. The first part is devoted to the history and prospects of the investigated material (dual-phase steel) and the related tribological issues encountered in engineering applications. In the second part, the related concepts of dry sliding friction and its main wear mechanisms are explained based on previous scientific works, and the impact of microstructure on the tribological behavior under dry sliding conditions is highlighted. The third part concentrates on the presentation of the factors affecting oxidative wear, oxide formation, oxidative wear mechanism, and tribological properties of the oxide layer. After understanding the background of this thesis, the last part points out the research objectives and significance.

In Chapter 2, the selection of materials and the generation of microstructures are described in detail. First, an introduction to the selected experimental materials is given.

Then, the procedure of heat treatment is illustrated and the evolution of the microstructure achieved after heat treatment is shown. Afterwards, the macro-hardness and microhardness of the specimens with different microstructures are measured and analyzed.

Chapter 3 focuses on the effect of microstructure, i.e., MVF, on oxidative wear of dual-phase steels in air. Meanwhile, the effect of sliding velocity on oxidative wear is also investigated. The investigations concentrate on three aspects, including friction coefficient, wear rate, and wear mechanism. The obtained methods and variation laws are presented for COF and wear rate, respectively. The wear mechanism, wear debris and plastic deformation are also explored.

In Chapter 4, the investigation concentrates on the combined effects of microstructure and gaseous atmosphere on the oxidative wear of dual-phase steels. At first, an experimental method is performed for different gas environments.

Subsequently, the corresponding experimental results are given, i.e., the results of COF, wear rate, wear mechanisms, plastic deformation, and wear debris size variations under different gaseous atmospheres and MVF. Then, the combined effects of gaseous atmosphere and microstructure on oxidative wear are explored in terms of experimental results and electron microscopic analysis.

In Chapter 5, the tribological behavior of the oxide layer after high temperature oxidation is the major point of study. The first part is devoted to the generation of oxide layers and oxide layer characterization (composition and thickness). Later, the nanoindentation and scratch experiments are presented with methods and results, individually. The nanoindentation is aimed at the hardness properties of the oxide layer.

The emphasis of scratch tests is concentrated on the acquisition of the tribological properties for oxide layers, i.e. the COF and the specific wear resistance. Finally, the effects of oxidation temperature and microstructure on the tribological performance of oxide layers are discussed and explained in the light of the experimental data and the analysis of the wear mechanism, thereby the corresponding conclusions.

This thesis ends with general conclusions and perspectives on the investigations performed.

Introduction

Tribological systems are strongly impacted by many parameters, such as working conditions and contact surface properties. As a bibliography, this chapter aims to build the existing knowledge situation in tribology allowing a better understanding of the subject and background of the thesis. The development of dual-phase steels, the effect of microstructure on dry sliding tribology, and oxidative wear mechanisms are introduced via different investigated literatures. Then, the current research status and limitations for investigations about the effects of microstructure, gaseous environment, sliding velocity, and temperature on the oxidative wear under dry sliding conditions are analyzed to demonstrate the research objectives of this thesis.

Development and application of dual-phase steel

As the most important structural material for mankind, steel is widely employed in various fields such as energy, transportation, security, and infrastructure. Since the invention of the automobile, automakers have been committed to finding the right steel for automotive production. In the early 1900s, the emerging automotive industry ushered in the era of alloy steels. During the first years of automotive production, formability and cost were the main drivers for the steel used. In this period, the strength of the steel was of less interest because if additional strength was needed, the thickness would be increased. Steels with two phases were first mentioned in some research and patents in the 1930s and 1940s, but basic research on dual-phase steels was not initiated until the 1960s. In the 1950s and 1960s, high-strength low-alloy (HSLA) steels entered the automotive market, and the popularity of HSLA steels was driven by safety regulations that were becoming more prevalent in the industry. While HSLA steels proved good mechanical properties, their ductility and formability were not as much as necessary to manufacture press-formed car components [START_REF] Farias Moreno | Influence of the Microstructure of the Corrosion Performance of DP Steels[END_REF][START_REF] Tasan | An overview of dual-phase steels: advances in microstructure-oriented processing and micromechanically guided design[END_REF][START_REF] Allain | Dual-Phase Steels: The First Family of Advanced High Strength Steels[END_REF].

The rise of dual-phase steels in 1975 and the initial 25 years of development did not exceed 600-700 MPa tensile strength, especially as summarized in the book "Lowalloy dual-phase steels" published in 1986, while the demands of the major automotive manufacturers are for dual-phase steels with a tensile strength of 1200 MPa and a ductility close to 600-800 MPa [START_REF] Fonstein | Dual-phase steels[END_REF]. Then, in the 1990s, two major contributions to the development of dual-phase steel were generated: the correlation between the strength and MVF was determined, and the properties of ferritic-martensitic steels treated by continuous annealing were discussed. It was at this time that these steels were first officially referred to as "dual-phase" steels [START_REF] Allain | Dual-Phase Steels: The First Family of Advanced High Strength Steels[END_REF][START_REF] Granbom | Structure and mechanical properties of dual phase steels: An experimental and theoretical analysis[END_REF].

Subsequently, advanced high strength steels (AHSS) have been highly regarded in the industrial field [START_REF] Tasan | An overview of dual-phase steels: advances in microstructure-oriented processing and micromechanically guided design[END_REF][START_REF] Fonstein | Dual-phase steels[END_REF][START_REF] Guo | Effect of microstructure quenched around Ac3 point on the damage behavior in 0.087 C-1.35 Mn steel[END_REF][START_REF] Kalashami | The effect of Nb on texture evolutions of the ultrafinegrained dual-phase steels fabricated by cold rolling and intercritical annealing[END_REF][START_REF] Yuan | Effect of warm rolling temperature on the microstructure and texture of microcarbon dual-phase (DP) steel[END_REF]. The initial AHSS steels were classified into four types:

dual phase steels, complex phase steels, transformation induced plasticity steels, and martensitic steels [START_REF] Soleimani | Transformation-induced plasticity (TRIP) in advanced steels: a review[END_REF]. These steels are widely used in the automotive industry owing to the superior balance of strength and ductility [START_REF] Soleimani | Transformation-induced plasticity (TRIP) in advanced steels: a review[END_REF][START_REF] Zhou | The fatigue properties, microstructural evolution and crack behaviors of low-carbon carbide-free bainitic steel during low-cycle fatigue[END_REF][START_REF] Liu | Revisiting the martensite/ferrite interface damage initiation mechanism: The key role of substructure boundary sliding[END_REF]. A major breakthrough in AHSS occurred in the mid-1990s when the steel supplier further refined the construction of equipment such as thermomechanical processing and continuous annealing lines [START_REF] Tamarelli | AHSS 101: the evolving use of advanced high strength steels for automotive applications[END_REF].

In 1994, a consortium of 35 major steel producers worldwide launched the Ultra Lightweight Steel Body (ULSAB) program to design and build lightweight steel body that could meet new safety and performance goals while minimizing fuel consumption [START_REF] Keeler | The proliferation of new types and grades of advanced high strength steels[END_REF][START_REF] Badkoobeh | The bake hardening mechanism of dual-phase silicon steels under high pre-strain[END_REF][START_REF] Badkoobeh | Microstructure and mechanical properties of resistance spot welded dual-phase steels with various silicon contents[END_REF][START_REF] Hou | Investigation of evolving yield surfaces of dual-phase steels[END_REF][START_REF] Avishan | Effect of pre-existing martensite within the microstructure of nano bainitic steel on its mechanical properties[END_REF][START_REF] Kumar | Stability of retained austenite in carbide free bainite during the austempering temperature and its influence on sliding wear of high silicon steel[END_REF]. The usage of AHSS steels in automotive production has steadily increased since the 2000s. Further research has been conducted to create second and third generation AHSS [START_REF] Soleimani | Transformation-induced plasticity (TRIP) in advanced steels: a review[END_REF]. In the last decade, AHSS steels have become the fastest growing material for automotive components.

Figure 1.1: Independent marketing research of material growth in kg for future automotive applications [START_REF] Kupke | Effect of the microstructure on the unloading characteristics of DP steel[END_REF] As one would expect from a material with high tensile strength, dual-phase steels are well suited for automotive components that absorb a lot of energy during crashes.

Increased use of dual-phase steel in automotive components threefold from 2009 to 2020 (Figure 1.1).

The microstructure of dual-phase steels consists of two phases, ferrite and martensite, but small amounts of bainite, pearlite or residual austenite may also be present. The soft ferrite matrix has a low yield strength and fine formability, while the hard martensitic phase provides high tensile strength. All low carbon steels can be processed in a variety way to obtain a dual-phase microstructure, including hot rolling, continuous annealing, or batch annealing. To generate dual-phase steels by continuous annealing, the material is heated above the subcritical temperature (A1) into the austenitic-ferritic two-phase zone [START_REF] Kupke | Effect of the microstructure on the unloading characteristics of DP steel[END_REF][START_REF] Becker | Dualphasen-Stähle mit erhöhter Festigkeit und Verformbarkeit[END_REF]. After a short holding time, the material is rapidly cooled (quenched) below the martensite onset temperature (Ms) shown in Increasing the temperature before quenching leads to an increase in the MVF. However, the carbon concentration in martensite decreases with increasing MVF due to the mixing rule [START_REF] Movahed | The effect of intercritical heat treatment temperature on the tensile properties and work hardening behavior of ferrite-martensite dual phase steel sheets[END_REF][START_REF] De Cooman | Fundamentals of steel product physical metallurgy[END_REF], as the carbon content in the material remains constant. The current research aims to enhance the microstructure of dual-phase steels to improve material properties such as material strength, uniform elongation, wear resistance and material hardness. There are several approaches to modify its microstructure, such as changing the MVF, adjusting the carbon concentration in martensite and grain refinement [START_REF] Kupke | Effect of the microstructure on the unloading characteristics of DP steel[END_REF].

These variations in microstructure provide a wide range of strength and ductility to dual-phase steels. Dual-phase steels can be also produced as Hot-dip Galvanized, Hot-dip Galvannealed and Electro Galvanized. Dual-phase steels can also be fabricated by hot-rolled and cold-rolled depended on materials. Hot rolled and cold rolled dualphase steels both provide an incredibly beneficial combination of low yield, high tensile strength, ease of cold working and weldability because of the ferrite -martensite lattice microstructure. Meanwhile, the high carbon content in the formed martensite during cooling process improves the hardenability of the steel. Typically, higher carbon content promotes greater strength and a higher percentage of martensite [START_REF] Thakur | Processing-microstructure-property correlation for producing stretch-flangeable grade dual-phase steel[END_REF][START_REF] Li | Microstructure-Based Modeling of Mechanical Properties and Deformation Behavior of DP600 Dual Phase Steel[END_REF][START_REF] Hu | Effect of Two Steps Overaging on Mechanical Properties of Tailor Rolled Blank of Dual Phase Steel[END_REF][START_REF] Dai | High-strength dual-phase steel produced through fast-heating annealing method[END_REF][START_REF] Shan | Modeling of Bake Hardening Kinetics and Carbon Redistribution in Dual-Phase Steels[END_REF].

In addition, there is a bake hardening effect in dual-phase steels, which is an important advantage over traditional HSLA materials. Moreover, due to its strainhardening capability from the beginning of deformation, there is an excellent combination of strength and tensile capability in dual-phase steel [START_REF] Badkoobeh | The bake hardening mechanism of dual-phase silicon steels under high pre-strain[END_REF][START_REF] Shan | Modeling of Bake Hardening Kinetics and Carbon Redistribution in Dual-Phase Steels[END_REF].

Dual-phase steel is also widely used in the petrochemical industry. The service environment of petrochemical equipment and materials is not only complex, but also very demanding, facing high load, high impact, high abrasion, high temperature, high pressure, alternating fatigue, ablative abrasion, and other complex service environment, in addition to meet the low cost, easy processing, easy welding and other technical and economic requirements. In addition, dual-phase stainless steel is also extensively applied in various other fields with its unique advantages, such as pulp and paper industry, transportation industry, fertilizer industry [START_REF] Tasan | An overview of dual-phase steels: advances in microstructure-oriented processing and micromechanically guided design[END_REF][START_REF] Fonstein | Dual-phase steels[END_REF][START_REF] Soleimani | Transformation-induced plasticity (TRIP) in advanced steels: a review[END_REF][START_REF] An | A transmission electron microscopy and atom-probe tomography study of martensite morphology and composition in a dual-phase steel[END_REF][START_REF] Chen | A novel technique for developing a dual-phase steel with a lower strength difference between ferrite and martensite[END_REF][START_REF] Nouri | Relationship between microstructure and corrosion behavior in dual-phase steels with various Si content[END_REF].

The relationship between the microstructure and mechanical properties of dualphase steels has been investigated due to their high significance in these industries.

Pierman et al. [START_REF] Pierman | The influence of microstructure and composition on the plastic behaviour of dual-phase steels[END_REF] revealed that the effect of martensite morphology is more important than MVF on the mechanical properties of dual-phase steels. Other investigators demonstrated that the intermediate quenching treatment/water provides the best mechanical properties in dual-phase steels [START_REF] Bag | Influence of martensite content and morphology on tensile and impact properties of high-martensite dual-phase steels[END_REF][START_REF] Deng | Effect of heat-treatment schedule on the microstructure and mechanical properties of cold-rolled dual-phase steels[END_REF]. Several studies focused on the effects of martensite morphology and volume fraction on the COF and wear rate of dual-phase steels. Some of them showed that the wear rate increases with decreasing MVF [START_REF] Abouei | A study on the wear behaviour of dual phase steels[END_REF][START_REF] Modi | High-stress abrasive wear response of 0.2% carbon dual phase steel: Effects of microstructural features and experimental conditions[END_REF][START_REF] Tyagi | Effect of martensite content on friction and oxidative wear behavior of 0.42 Pct carbon dual-phase steel[END_REF][START_REF] Burbure | Studies on tribological wear behaviour for optimization of dual-phase steels[END_REF][START_REF] Hamza | Microstructural and Tribological Characterization of API X52 Dual-Phase Steel[END_REF]. And Wei et al.[40] investigated the impact of ferrite -martensite morphology on the scratch and wear resistance of dual-phase steel.

This thesis focuses on the tribological behavior of dual-phase steels, hoping to investigate their oxidative wear performance under different working conditions via changing the microstructure (MVF) of dual-phase steels. This can provide a great deal of help to researchers to make dual-phase steels with excellent performance and apply them in different working environments.

1.2

The effect of microstructure on friction and wear under dry sliding condition

Dry sliding friction

Generally, friction is the force which resists the relative motion generated between sliding bodies. As a principle, it is possible to distinguish between two types of frictional forces, irrespective of the shape, size, and nature of the sliding body: (a) the force required to hold a solid surface sliding against another is referred to as kinematic or dynamic friction, and (b) the slightly higher force necessitated by initiating motion is called static or limiting friction [START_REF] Dawud | Oxidational wear of low alloy steel in gases other than air[END_REF]. It has been found that there are two fundamental frictional laws for the sliding of metals on metals through experiments:

1) Frictional force Ff is commonly related to the applied load positively. It can be represented as

𝐹 𝑓 = 𝜇𝐹 (1.1)
Here (F) is the applied normal load and μ is the coefficient of friction. It is a dimensionless unit which is constant for a given pair of sliding materials, under the same environmental conditions. There are different coefficients of friction with different materials and conditions.

2) There is no relationship between the friction of two sliding bodies and the surface contact area of the metal. When these two surfaces are in static contact within a surface contact area, they are, in fact, only contacting at the point where the tip corners of the two surfaces touch. The plastic flow at the contact surface leads to the increasing actual area of contact (A), which is suggested by Bowden and Tabor [START_REF] Bowden | The Friction and Lubrication of Solids-Part II[END_REF] and is expressed as:

𝐴 = 𝐹 𝑃 𝑚 (1.2)
where Pm is the flow pressure for the utilized material. Bowden and Tabor [START_REF] Bowden | The Friction and Lubrication of Solids-Part II[END_REF] argued that the friction is primarily dependent on the shear strength of the cold weld junctions generated among the contacting bodies. In general, any assumptions about friction must consider the truth that there are numerous "valleys and hills" included even in the smoothest surfaces.

Blau [START_REF] Blau | Interpretations of the friction and wear break-in behavior of metals in sliding contact[END_REF] identified three stages concerning the friction process: Firstly, the establishment may appear at the beginning of sliding, and is determined by the surface roughness, the surface cleanliness, and the nature and destruction of the oxide film;

Secondly, there is an elastic and plastic deformation of the surface, with a significant increase in contact temperature leading to process hardening of the subsurface layer.

Simultaneously, accompanied by crystallographic reorientation of the subsurface morphology; Finally, there is a balance of constant friction related to the sliding process, which is governed by a temperature-stable equilibrium of oxide generation and breakdown. Wilson,Stott,and Wood[44] presented a friction model predicated on the effect of the compacted oxide on the COF. The COF is correlated with the area of the compacted oxide, oxide, and metal/oxide interface parameters in their model. Stott and Wood[45] investigated friction associated with oxide and junction growth. It was determined that the growth of junctions during sliding is comparatively less since the contact pressure decreases with the growth of junctions, returning the oxide to the brittleness. Therefore, the COF is lower than that obtained in metal-tometal contact.

Main wear mechanisms in dry friction conditions

Wear is the continuous loss of material through mechanical and chemical action during the relative movement of surfaces. As a result of friction, wear is generally formed, but under certain conditions, such as fluid lubrication or the transformation of external friction into internal friction, there is only energy loss but not necessarily material loss. Wear, friction, and lubrication are interrelated.

During wear, the rate of material removal depends on a great number of variables such as temperature, load, velocity, contact area, sliding distance, atmosphere, material composition, type of lubrication, finish, and type of motion. Initially, Archard, Hirst, Burwell, Kerridge, Lancaster, Welsh and others have done extensive researches in the field of metal-to-metal sliding wear [START_REF] Archard | The wear of metals under unlubricated conditions[END_REF][START_REF] Hirst | Surface film formation and metallic wear[END_REF][START_REF] Hirst | The influence of speed on metallic wear[END_REF][START_REF] Lancaster | The formation of surface films at the transition between mild and severe metallic wear[END_REF][START_REF] Kerridge | The stages in a process of severe metallic wear[END_REF][START_REF] Welsh | The dry wear of steels I. The general pattern of behaviour[END_REF][START_REF] Burwell | Survey of possible wear mechanisms[END_REF]. Archard [START_REF] Archard | The wear of metals under unlubricated conditions[END_REF] revealed the empirical equation for unlubricated sliding mild wear of metals,

𝑉 = 𝑘𝐹𝑆 𝐻 (1.3)
Which means that the volume of wear (V) is proportional to the normal load (F) and sliding distance (S), but inversely to the hardness of the metal (H). Burwell classified the wear mechanisms according to the physical processes [START_REF] Burwell | Survey of possible wear mechanisms[END_REF], namely abrasive wear, adhesive wear, fatigue wear, erosion wear, corrosion wear and fretting wear. As mentioned, any one of these can work alone or in combination with another mechanism and cause wear under lubricated or non-lubricated conditions. Besides, Archard and Hurst [START_REF] Archard | The wear of metals under unlubricated conditions[END_REF] distinguished two types of wear, mild wear and severe wear, based on the degrees of wear. Mild wear is characterized by high contact resistance, small debris particle size, and smooth oxidized wear surfaces. In mild wear, there is slow material removal from the surface. Severe wear is typified by low contact resistance, large particles of metal debris and severe plastic deformation. In some cases, a rough and deeply torn surface would be generated in severe wear. These theories formed the basis for subsequent investigations of wear science.

In the industry, there are many situations where liquid lubrication is impracticable, such as in the nuclear and aerospace industries. Under such circumstances, the moving parts must run successfully in such cases that are difficult to provide any conventional lubrication due to overheating, nuclear radiation, vacuum, etc. The work presented in this investigation is limited to this dry friction, in which the most common wear mechanisms are adhesive wear, abrasive wear, fatigue wear, and corrosion wear.

The theory of adhesive wear was first proposed by Holm [START_REF] Holm | Electric contacts: theory and application[END_REF], which refers to the transfer of material from one surface to another during the relative motion due to welding or local adhesion between the two surfaces. Holm [START_REF] Holm | Electric contacts: theory and application[END_REF][START_REF] Holm | Metal-centered oxygen atom transfer reactions[END_REF] considered adhesion as the capture of atoms from one surface by atoms from the other surface due to frictional effects of sliding surfaces in contact with each other. Subsequently, Bowden

and Tabor [START_REF] Bowden | Friction and lubrication[END_REF], Rabinowicz [START_REF] Rabinowicz | The least wear[END_REF], and Archard [START_REF] Archard | The wear of metals under unlubricated conditions[END_REF] further refined this theory. Bowden

and Tabor [START_REF] Bowden | The friction and lubrication of solids[END_REF] first determined that the actual contact area is only a very small fraction, revealing the role of plastic flow and transient high temperature at the contact asperities for the formation of adhesion junctions. It was concluded that when two surfaces slide against each other and the pressure between the contact surfaces is high enough to cause local plastic deformation, adhesion occurs and results in material transfer. The amount of adhesive transfer depends on the mechanical performance of the interfaces, the nature of the surfaces, and the surrounding environments. It is possible to reduce adhesive wear by any oxides or absorbent films e.g. oxygen and water vapor.

Lubrication is another powerful approach to reduce adhesive wear. Extremely severe adhesive wear causing catastrophic surface damage is termed "scuffing". The general process of adhesive wear shows as fellow:

1) When the friction surfaces are in contact with each other, the higher microconvex body is first in contact, and the contact pressure causes plastic deformation of the micro-convex body, and then, the phenomenon of adhesion.

2) When the relative motion of the friction vice, the adhesion point shear damage, so that the material transferred from the softer surface to the harder surface.

3) When further sliding, under the action of friction, the material transferred to the hard surface is peeled off resulting in the formation of abrasive particles.

Abrasive wear occurs when surfaces slide against each other, and the rough asperities on hard surfaces or hard particles between sliding surfaces penetrate and ploughing out of material from soft surfaces, that is, abrasive wear. Abrasive wear can be classified as two-body abrasive wear and three-body abrasive wear according to the mutual position of the abrasive and the surface. In the case of a rough hard surface sliding against a softer surface, it is referred [START_REF] Avient | An experimental study of friction and wear during abrasion of metals[END_REF] as two body abrasion. When abrasion is caused by loose, hard particles sliding between rubbing surfaces, it is known as 3body abrasion [START_REF] Rabinowicz | A study of abrasive wear under threebody conditions[END_REF]. Burwell [START_REF] Burwell | Survey of possible wear mechanisms[END_REF] first proposed that the greater the deformation of the contact part within the elastic limit, the less mechanical damage will occur. Therefore, the greater the ratio of hardness to longitudinal modulus of elasticity, the better the wear resistance. Rabimowicz [START_REF] Rabinowicz | The least wear[END_REF] developed a simplified micro-cutting model and derived the equation,

𝑉 = 𝑘𝐹𝑆𝑡𝑎𝑛𝛼 𝐻 (1.4)
Which considering that the wear volume (V) is proportional to the normal load (F), the sliding distance (S), and the angle between the abrasive grain and the softer surface (θ). Meanwhile, it is inversely correlated with the hardness (H) of the softer material.

Fatigue wear means that even a slight stress within the elastic limit acts on the interior of the material, it leaves a certain "trace", which accumulates and causes fatigue.

In other words, as a result of repeated frictional stresses working inside the material, part of the material peels off due to fatigue. Unlike adhesive and abrasive wear, fatigue wear occurs in all frictional contacts and is unavoidable even under favorable oil film lubrication conditions [START_REF] Bowden | Friction and lubrication[END_REF]. In some cases, alternating contact stresses might cause extensive surface damage and pitting, followed by surface cracking and the formation of large fragments. Or, in less severe cases, a delamination wear might occur. There are more arguments for fatigue wear, the first and representative one is Suh [START_REF] Suh | The delamination theory of wear[END_REF][START_REF] Suh | An overview of the delamination theory of wear[END_REF], who proposed the theory of delamination wear, which has a significant impact on the study of the sliding wear mechanism. Because of the repeated friction, cavities form below the friction surface and grow into cracks. As the crack continues to propagate at a certain depth, the upper surface flakes off.

Corrosion wear is caused by the formation and peeling off of corrosion products formed by the chemical or electrochemical reaction between the frictional surfaces and the environmental medium during the friction process [START_REF] Bateni | Wear and corrosion wear of medium carbon steel and 304 stainless steel[END_REF]. Oxidation wear is one of the most typical corrosion wear, has been concerned about and research. The surface of any part present in the air always has a layer of oxygen adsorption, and when relative motion occurs, a large contact pressure produces at the projection of the surface, resulting in plastic deformation. Normally, the large plastic deformation accelerates the diffusion of oxygen into the metal, leading to the formation of an oxide layer. The formed oxide layer plays a role in reducing wear due to its higher hardness than the substrate. As the contact surfaces continue to move relative to each other, the formed oxide layer will gradually peel off owing to the low strength of the oxide. Meanwhile, the new metal surface will be exposed to re-oxidation, so that the process of oxidation and peeling alternately, resulting in oxidation wear. Most of the opinions [START_REF] Wilson | The development of wear-protective oxides and their influence on sliding friction[END_REF][START_REF] Archard | The wear of metals under unlubricated conditions[END_REF][START_REF] Tenwick | A simplified theory for the oxidative wear of steels[END_REF][START_REF] Tao | A study of oxidation phenomena in corrosive wear[END_REF][START_REF] Quinn | Application of the oxidational theory of mild wear to the sliding wear of low alloy steel[END_REF] hold that the presence of a certain amount of oxide on the friction surface is beneficial to reduce the direct metal-metal contact, which can reduce the COF and wear rate of the friction pair.

The effect of microstructure on tribological behavior

The microstructure can be changed by altering the chemical composition, alloying elements, carbon content and heat treatment parameters of the steel, all of which have an effect on the wear resistance of the material. Most of the steels, especially alloy steels, are usually utilized after heat treatment. Therefore, the effect of heat treatment and microstructure on tribological properties has attracted the interest of most researchers.

In the last few years, many researchers have contributed to the study of the impact of microstructure on friction and wear mechanisms [START_REF] Karlsson | Influence of tool steel microstructure on friction and initial material transfer[END_REF][START_REF] Trevisiol | Effect of hardness, microstructure, normal load and abrasive size on friction and on wear behaviour of 35NCD16 steel[END_REF][START_REF] Trevisiol | Effect of microstructures with the same chemical composition and similar hardness levels on tribological behavior of a low alloy steel[END_REF][START_REF] Tuckart | Influence of microstructure on tribologically mixed layers[END_REF][START_REF] Zambrano | The sliding wear behaviour of steels with the same hardness[END_REF][START_REF] Cui | Wear characteristics and mechanisms of H13 steel with various tempered structures[END_REF].

Wang [START_REF] Wang | Wear behavior of steel 1080 with different microstructures during dry sliding[END_REF][START_REF] Wang | Tribo-metallographic behavior of high carbon steels in dry sliding: III. Dynamic microstructural changes and wear[END_REF][START_REF] Wang | Tribo-metallographic behavior of high carbon steels in dry sliding: II. Microstructure and wear[END_REF] investigated the wear mechanisms of AISI 52100 and AISI 1080 carbon steels with different microstructures under dry sliding friction. In mild wear, there is no significant difference in the wear volume of various microstructures.

However, in severe wear such as adhesion and delamination, there are considerable differences in the wear volume of different microstructures. The wear resistance increases in the order of martensitic carbide, granular organization, martensite, bainite and lamellar pearlite. The wear resistance of each microstructure depends on the thermal stability, the resistance to plastic deformation, the resistance to crack initiation and expansion, especially on the variation of microstructure of the contact surface and energy consumption during the sliding process.

Meanwhile, Wang [START_REF] Wang | Effect of secondary carbides on oxidation wear of the Cr-Mo-V cast steels[END_REF] et al. investigated the effect of second-phase carbides on the dry tribological behavior of cast hot forging die steel. It was found that there are two aspects to the role of second-phase carbides. When the brittle and hard second phase is continuously distributed in the matrix, it is extremely detrimental to the wear performance. When the carbide in the steel is small and uniformly distributed, the plasticity and toughness of the material is better which resulting in excellent overall performance and higher wear resistance.

Surface strengthening treatment can effectively improve the tribological performance of steel. After surface strengthening treatment, the surface organization is changed, which causes modification of the tribological properties. Totik [START_REF] Totik | The effects of induction hardening on wear properties of AISI 4140 steel in dry sliding conditions[END_REF] et al.

studied the dry tribological behavior of surface high frequency quenched AISI4140. It was noticed that the COF of the surface treated AISI 4140 steel was significantly reduced and also the wear rate. The longer the time of high frequency quenching, the more obvious the reduction of COF and wear rate.

Cui [START_REF] Cui | Research on oxidation wear mechanism of the cast steels[END_REF] performed a unidirectional pin-disk experiment on Cr-Mo-V cast steel to investigate the effects of the morphology, composition, and structure of the oxide layer and the substrate on tribological behavior. During the sliding process, surface oxidation and fatigue peeling of the oxide layer alternate. When there is no coarse second phase, the oxide layer is the main determinant of the wear rate, which is in accordance with Quinn's theory [START_REF] Quinn | Application of the oxidational theory of mild wear to the sliding wear of low alloy steel[END_REF] of oxidative wear. In this case, oxide layer delamination is observed from inside the oxide layer or at the interface between the matrix and the oxide layer, resulting in a lower wear rate. When a coarse second phase is present in the steel, the wear rate is intensely determined by the microstructure of the substrate. The oxide layer is delaminated from inside the substrate, causing high wear rate.

Abouei [START_REF] Abouei | Effect of microstructure on the oxidative wear behavior of plain carbon steel[END_REF] investigated the effect of microstructure on oxidative wear of plain carbon steel. The wear mechanism of each microstructure under all experimental conditions was oxidative wear. As the load increasing, the effect of different microstructures (such as F+P, F+M, M, etc.) on oxidative wear becomes obvious and the difference in wear resistance gets significant.

For dual-phase steel, Davies et al. [START_REF] Davies | Influence of martensite composition and content on the properties of dual phase steels[END_REF] suggest that the strength of the microstructure is related to the grain size of ferrite and the MVF, and irrelevant to the composition and strength of the martensite. Tyagi et al. [START_REF] Tyagi | Modelling of dry sliding oxidation-modified wear in two phase materials[END_REF] showed that as the martensite content increasing, the wear volume decreases linearly, i.e. the wear resistance increases. The hardness of martensite is also a major factor affecting the wear resistance of dual-phase steels. The results of Aksoy et al. [START_REF] Aksoy | An evaluation of the wear behaviour of a dual-phase low-carbon steel[END_REF] showed that the wear resistance of dual-phase steel increases linearly with the hardness of martensite. Some investigations are concentrated on the effect of MVF on tribological behavior [START_REF] Jha | Correlating microstructural features and mechanical properties with abrasion resistance of a high strength low alloy steel[END_REF][START_REF] Nadlene | Study on the effect of volume fraction of dual phase steel to corrosion behaviour and hardness[END_REF][START_REF] Saai | Influence of volume fraction and distribution of martensite phase on the strain localization in dual phase steels[END_REF][START_REF] Saghafian | Correlating microstructural features with wear resistance of dual phase steel[END_REF][START_REF] Trevisiol | Effect of martensite volume fraction and abrasive particles size on friction and wear behaviour of a low alloy steel[END_REF][START_REF] Xu | The effect of martensite volume fraction on the scratch and abrasion resistance of a ferrite-martensite dual phase steel[END_REF][START_REF] Xu | Design of low hardness abrasion resistant steels[END_REF].

Trevisiol et al. [START_REF] Trevisiol | Effect of martensite volume fraction and abrasive particles size on friction and wear behaviour of a low alloy steel[END_REF] show that the COF and wear rate decrease with increasing MVF. For low MVF, the main wear mechanisms are plowing and adhesive wear, while cutting wear becomes predominant instead of plowing and adhesion under high MVF.

Saghafian and Kheirandish [START_REF] Saghafian | Correlating microstructural features with wear resistance of dual phase steel[END_REF] explain that the wear rate of dual-phase steels decreases with increasing MVF, which is probably due to the reduction of the real contact area and the formation of cracks. However, Jha et al. [START_REF] Jha | Correlating microstructural features and mechanical properties with abrasion resistance of a high strength low alloy steel[END_REF] suggest that for dual-phase steel, compared with the full martensitic microstructure, better wear resistance can be provided by a microstructure with a small content of soft ferrite. To a certain extent, although the presence of ferrite reduces the overall strength and hardness of the material, it reduces the brittleness, enhancing the overall fatigue resistance and wear resistance.

In addition, scratch tests on dual-phase steels at low normal load show that the wear depth decreases with increasing MVF [START_REF] Xu | The effect of martensite volume fraction on the scratch and abrasion resistance of a ferrite-martensite dual phase steel[END_REF][START_REF] Xu | Design of low hardness abrasion resistant steels[END_REF]. On the other hand, under high normal load, with increasing MVF, it first decreases and then increases. However, the optimal MVF which displays the best scratch resistance is yet to be determined. Hence, further studies are needed to investigate the effect of MVF on friction and wear.

When considering the wear mechanism, there are fewer investigations about the impact of microstructure such as phase volume fractions on oxidative wear. Previous research of the microstructural factors demonstrates that the wear rate of the dual-phase steel decreases linearly with an increasing MVF for a particular load under a pin-ondisc machine. The main reason for the decline in the coefficient of friction is the decreasing actual contact area caused by the increasing material hardness. Similarly, the average coefficient of friction decreased with an increasing MVF at a given load [START_REF] Tyagi | Effect of martensite content on friction and oxidative wear behavior of 0.42 Pct carbon dual-phase steel[END_REF][START_REF] Abouei | Effect of microstructure on the oxidative wear behavior of plain carbon steel[END_REF][START_REF] Tyagi | Modelling of dry sliding oxidation-modified wear in two phase materials[END_REF]. The above studies have all involved the impact of MVF on the wear rate of dual-phase steel, but there are huge differences in the experimental conditions of friction. The method adopted by Saghfian et al. [START_REF] Saghafian | Correlating microstructural features with wear resistance of dual phase steel[END_REF] was a unidirectional sliding contact of a cylindrical steel pin the steel disc, with a load of 61-83 N, and a relatively high sliding speed (1.2 m/s). The experiments of Trevisiol et al.[86] were focused on the contact of cylindrical pin samples against the abrasive papers. The normal load was also relatively high (50~110 N), and the sliding speed was constant at 0.06 m/s. The test of Jha et al. [START_REF] Jha | Correlating microstructural features and mechanical properties with abrasion resistance of a high strength low alloy steel[END_REF] was performed by a rubber wheel tribometer, in which the rubber wheel was directly rubbed against the flat steel sample in the vertical direction. The experimental load reached 48 N and the sliding velocity was 3.27 m/s. Due to the different tribological systems, the wear mechanism of dual-phase steel is constantly changing. For example, for the frictional system of Trevisiol et al. [START_REF] Trevisiol | Effect of martensite volume fraction and abrasive particles size on friction and wear behaviour of a low alloy steel[END_REF], the main wear mechanism is abrasive wear, while in the tribological system of Saghafian et al. [START_REF] Saghafian | Correlating microstructural features with wear resistance of dual phase steel[END_REF],

the main wear mechanism is delamination wear. Considering that different tribological systems have a great effect on the properties of friction and wear, it is necessary to establish different frictional systems to investigate the relationship between their corresponding tribo-logical behaviors and microstructures. Therefore, using the friction condition established in this paper to investigate the effect of the microstructure of dualphase steel on oxidative wear is a good supplement to the investigations of the tribological behavior of dual-phase steel.

Research on oxidative wear

Impact parameters of oxidative wear

In general, there are many parameters that impact the tribological properties of materials, such as normal load, temperature, sliding velocity, microstructure, gaseous atmosphere, alloy composition, etc. For each wear mechanism, the various influencing parameters play different roles on tribological behavior. In this study, besides the effect of microstructure on oxidative wear, the combined effects of microstructure, gaseous atmosphere, and sliding velocity on oxidative wear are also investigated.

The effect of gaseous atmosphere on oxidative wear

Most of the research on metal wear was performed in air. There is seldom research carried out in the atmosphere other than air.

In 1965, McCoy[90] studied the oxidation behavior of AISI 304 steel in carbon dioxide and found an approximate parabolic regularity. In addition, α-Fe2O3 formed at the contact interface and a spinal at the metal -oxide interface.

Clark et al. [START_REF] Clark | Paper 12: mild wear of unlubricated hard steels in air and carbon dioxide[END_REF] in 1967 studied the mild wear of unlubricated hard steels in air and carbon dioxide at temperatures up to 500 ℃ over a range of loads and speeds. In the CO2, it was found an extremely low wear rate, which was attributed to the formation of protective iron alloy oxides on the surface during sliding. The wear rate increases significantly above a given temperature, probably because the substrate become so soft that no longer support the oxide layer.

Stott and Wood [START_REF] Stott | The influence of oxides on the friction and wear of alloys[END_REF] in 1978 studied the evolution of compact oxide "glazes" for metals on sliding tests in air and at different oxygen partial pressures. It was noted that the formation of a smooth oxide "glaze" with increasing oxygen partial pressure results in a significant decrease in COF and wear rate. Another point to note is that once the oxide is formed at a given oxygen partial pressure, the oxide persists during the whole sliding process, even at low oxygen pressure.

Later, in 1980, Wilson,Stott,and Wood[44] investigated the tribological behavior of unidirectional and reciprocating sliding of Fe and Fe-Cr alloys at 293 K and different oxygen partial pressures. It was noticed that, the COF gradually decreases with the formation of a compact oxide layer on the contact surface, eventually reaching a stable value of about 0.6 when a complete oxide coverage is reached. This is achieved more rapidly at higher oxygen partial pressures.

Sullivan and Granville [START_REF] Sullivan | Reciprocating sliding wear of 9% Cr steel in carbon dioxide at elevated temperatures[END_REF] performed experiments in 1984 on the initial stages of reciprocal sliding wear of 9% Cr steels in a CO2 environment at temperatures ranging from 200 to 550°C. Above 290°C, there is initially a severe wear, followed by an transition to mild oxidative wear.

Tennenhouse et al. [START_REF] Tennenhouse | The effects of oxygen on the wear of Si3N4 against cast iron and steel[END_REF] in 1986 showed that the wear of silicon nitride on cast iron and steel is influenced by the oxygen concentration. In both pin-disk wear experiments and machining tests, the wear of silicon nitride was decreased by reducing the oxygen concentration in the environment. This suggests that there is more than one wear mechanism for silicon nitride. For example, mechanical and chemical factors may be affected the wear process.

In 1986, Newman and Skinner [START_REF] Newman | The high temperature sliding wear of stainless steel in a CO2 atmosphere-The effect of adding low concentrations of O2[END_REF] examined the sliding friction of 316 stainless steel in a carbon dioxide. They revealed that the addition of controlled oxygen to CO2 can generate a transition from severe to mild wear, which confirms the observations of Wallace[95]. This can be explained by the increased rate of substrate hardening due to oxide inclusions into the substrate layer. However, they argued that the increase in hardening rate is due to the acceleration of linear oxidation, contrary to the hypothesis of Wallace [START_REF] Wallace | Oxidational and tribological studies of austenitic stainless steels under CO2-based environments[END_REF], who found that oxidation is essentially logarithmic in growth.

In 1992, Dawud [START_REF] Dawud | Oxidational wear of low alloy steel in gases other than air[END_REF] examined the oxidative wear of BS970En8 steel in air, argon, carbon dioxide, 100% O2, 20%O2-80% Ar and 2%O2-98% Ar atmospheres by a pin-ondisk tribometer. The curves of the wear rate indicated that the gas composition has a significant influence on the wear behavior. For the Ar-O atmosphere, the wear rate decreases with increasing oxygen concentration and greater oxide coverage. The maximum wear rate occurs in Ar. The wear rate is lower in air compared to the 20%

Chapter 1-Introduction

20

O2-80% Ar atmosphere, probably due to the existence of water vapor in the air. The lowest wear rate appears in the CO2, because the oxidative growth in this environment is controlled by dissociation kinetics rather than diffusion.

There is now widespread agreement that the presence of a gas atmosphere has a significant impact on the friction and wear behavior of metallic materials. This means that when the material is exposed to a gaseous environment, the contact and adsorption capabilities of the newly formed metal surface vary significantly due to the distinct physical and chemical properties of gas molecules [START_REF] Akagak | Sliding friction and wear of metals in vacuum[END_REF][START_REF] Fukuda | Influences of Gases on the Tribological Properties of Pure Iron[END_REF][START_REF] Mishina | Atmospheric characteristics in friction and wear of metals[END_REF][START_REF] Mishina | Chemisorption of diatomic gas molecules and atmospheric characteristics in adhesive wear and friction of metals[END_REF]. In general, when the rubbing pair interacts in the air, a large amount of oxygen increases the surface activity, which is beneficial to the formation of an oxide film, reducing adhesion and providing favorable friction and wear characteristics [START_REF] Mishina | Chemisorption of diatomic gas molecules and atmospheric characteristics in adhesive wear and friction of metals[END_REF]. However, depending on the gas molecules in the environment, these surface layers can differ dramatically [100,101].

In CO2 atmospheres, surveys such as those conducted by Igor Velkavrh[102,103] revealed that outstanding wear reduction was observed. Metal oxides may allow for relatively favorable friction conditions at intermediate concentrations (e.g., air, CO2 atmosphere) while also preventing the formation of Fe-Fe bonds, which leads to severe adhesion.

In a study using AISI 52100 steel ball and medium carbon steel disc under fretting conditions, friction and wear were higher in N2 atmosphere than in air atmosphere [104],

and also in a study using AISI 52100 steel ball and medium carbon steel disc under fretting conditions, friction and wear were higher in N2 atmosphere than in air atmosphere [105]. Although many scientific studies have been conducted on the effect of gaseous environment on material wear resistance, more experiments in different atmospheres are required to thoroughly study oxidative wear.

Gaseous atmospheres have long been known to have a significant impact on the friction and wear behavior of non-lubricated contacts, such as metal-metal adhesion, friction of oxidized metal surfaces, and so on. The findings from these studies, however, do not apply to all gas-lubricated metal/metal pairings due to the complexity and sensitivity of these systems. Adsorption of gas molecules and chemical modifications of contact surfaces, in particular, determine the behavior of each metal differently.

Furthermore, there is still a lack of knowledge on the subject, because the tribological behavior and tribochemistry of gas-lubricated systems are highly sensitive to the choice of operating parameters, and the results from available literature are frequently scattered and, in some cases, contradictory.

This article focuses on different gas environments in determining the friction and wearing mechanisms in non-lubricated dual-phase steel of the technical gasses used at atmospheric gas pressure. As a measure, the atmospheres in Argon (Ar), nitrogen (N2)

and carbon dioxide (CO2) were used. Gas analysis considers their chemical reactivity with steel surfaces and the distinction is made between the properties and the observed friction and wear mechanisms of different tribo-chemical products. This is our main objective in researching the properties of double-phase stainless steel in various gas environments.

The effect of sliding velocity on oxidative wear

Conventionally, normal load and sliding velocity are the most significant parameters in tribological systems. [START_REF] Abouei | Effect of microstructure on the oxidative wear behavior of plain carbon steel[END_REF]113,114]. Inman et al.[115] reported that loose oxide debris are distributed on the worn surface and then form compact oxides at nearly room temperature. Sen et al.[116] examined the tribological behavior of aluminum oxide and AISI 52100 on a ball-on-disc tribometer. The results

showed that both the COF and the wear rate decrease with the increasing sliding speed.

Jiang G. Straffelini[109] investigated the oxidative wear of heat-treated steels at low sliding speeds (between 0.2 and 1 m/s). It is indicated that this type of wear can be analyzed by the formation and aggregation of oxide debris during sliding process.

Moreover, the available equations for two different types of oxidative wear are also presented to simulate the results.

I. Boromeia[118] performed friction experiments with Aluminium bronze and

Cold work tool steel (AISI A2) at sliding speeds of 0.6, 0.9, and 1.2 m/s. It was shown that surface hardening occurred during the friction process and a continuous adhesive oxide layer was generated. At higher sliding velocities, it is possible to have a more stable oxide layer resulting in better wear resistance and less friction.

Tang [119] investigated the effects of sliding velocity and hardness on the tribological performance of hardened tool steel AISI D2. It was demonstrated that the effect of hardness on the COF was more significant at sliding speeds of 0.05m/s and 0.5 m/s than at 0.10 m/s, while the wear rate was greatly sensitive to sliding velocity and hardness.

Jayashree [120] carried out tribological tests on silicon carbide-graphite composites against martensitic stainless steel using a pin-on-disk tribometer at two speeds of 1.57 m/s and 7 m/s. Severe wear was observed at 1.57 m/s, while at 7 m/s, a mild wear was characteristic. The wear is significantly reduced when the tribolayer is smooth, compact, and continuous. Bourithis et al. [121] reported the observations of pin-on-disk experiments performed at different loads from 9.8 to 39.2 N and sliding speeds from 0.15 to 0.75 m/s. As the sliding speed increased, a decrease in the wear rate of the D2 disc was noticed, which was explained by the formation of an oxide layer during the high-speed test.

F.E. Kennedy[122] performed dry sliding pin-on-disc experiments with

Fe30Ni20Mn25Al25 and AISI347 to determine the effect of frictional heating on wear and to investigate the role played by the third body (oxidized wear debris and friction layer)

during wear process. The tests were carried out in air at sliding speeds of 0.1 to 1.0 m/s.

It was found that the wear rate at high sliding velocities (≥0.5 m/s) was significantly lower than that at lower speeds. Moreover, the wear rate was higher at high sliding speeds when wear particles were continuously swept out from the disc surface during the experiment. An analysis of the worn surfaces and wear debris shows that oxidation occurs within the sliding contact process. At high sliding speeds, oxide debris and flaking material form a hard mechanical mixed friction layer which is contributing to lower wear rate.

In addition, there are some conclusions about the effect of sliding velocity on tribological behavior that contradict the above arguments, so it is necessary to examine further.

H. So[123] studied the wear mechanism of solid materials with a pin-on-disk device. When the load applied to the specimen exceeds a certain value, the COF increases with the sliding speed. He also reasonably explained the results by the accurate calculation of the overall temperature and flash temperature in the contact area.

Wang [124] completed dry sliding wear experiments on A52100 steel at different sliding speeds (1.0, 1.5, and 2 m/s) in room temperature. The results revealed that the wear volume increases with sliding speed.

P. Palanikumar[125] studied the effect of sliding velocity on the COF at the contact interface of SS304 alloy. Experiments were run at three different sliding speeds of 1m/s, 2m/s and 3m/s and it was concluded that the COF increased with the sliding speed.

Paul C. Okonkwo[126] applied a pin-on-disk setup to investigate the combined effects of sliding speed (12 mm/s to 760 mm/s) and surface temperature on the wear behavior of an unlubricated mild steel-tool steel contact pair. Adhesive wear of the tool steel is dominant at all sliding velocities. At low speeds, the adhesive volume is maximum and the adhesive wear rate is quite sensitive to the sliding velocity. However, at high speeds, the adhesive wear rate is less sensitive to the sliding velocity. It is well known that dual-phase steels are widely adopted in mechanical engineering, and the investigation of the wear properties has become a topical issue in the research of dual-phase steels. However, so far, there is limited knowledge about the effect of sliding speed on wear characteristics and mechanisms of dual-phase steels.

Therefore, it is urgent to investigate the properties of dual-phase steel in depth for the rapid development in engineering applications.

Formation of oxides

Steel is unstable in the air and when in contact with oxygen undergoes oxidation and forms several types of oxides, including FeO, Fe2O3, and Fe3O4, which adhere to the surface of substrate to form an oxide layer. At high temperatures, the oxidation reaction is more aggressive, and the oxide layer is more visible. This compact oxide layer is like a "glaze", which has a significant effect on the wear resistance of the steel, especially at high temperatures [113,132].

The different natures of oxides formed with temperature as follows: α-Fe2O3 appears below 200°C, Fe3O4 at 200 a 570°C, and FeO above 570°C.

FeO(wüstite) is a P-type ion-deficient semiconductor with a cubic crystal structure.

Since Fe 2+ is easily oxidized to Fe 3+ , this allows a small amount of Fe 2+ to be present in the void of the lattice as Fe 3+ , leading it to be a non-integrable compound. Besides, Fe has a wide quantitative ratio to O, from Fe0.95O to Fe0.84O at 1000ºC. Such FeO has a high vacancy concentration, which allows ions and electrons to migrate outward through the vacancies easily.

Fe3O4 (magnetite) is a mixed oxide with an anti-spinel structure, which can be regarded as FeO•Fe2O3. In other words, Fe3O4 has both Fe 2+ and Fe 3+ , which occupy octahedral vacancies and tetrahedral vacancies, respectively, and the iron ions can diffuse outward through the defects in the vacancies. At the same time, the presence of two valence cations allows electrons to migrate rapidly in both oxidation states, so Fe3O4 has better electrical conductivity.

Fe2O3 (hematite) is an n-type oxide, generally α-Fe2O3, with a rhombic structure.

γ-Fe2O3 is formed after special treatment of Fe3O4, with a cubic structure, and γ-Fe2O3 is a non-stable structure, which is highly susceptible to transform into α-Fe2O3. In the rhombic structure oxygen ions are in the position of dense hexagonal crystal nodes and iron ions are in the interstices. Therefore, the iron ions are movable. The latest results

suggest that Fe2O3 grows through the migration of surface cations. At the same time, O 2-can also migrate through the Fe2O3 layer to the interior of the oxide at the Fe2O3/O interface and the oxidation reaction can continue. At the interface of the Fe/FeO layer (( 1)), iron ions and electrons migrate outward through the FeO layer via iron ion vacancies and electron vacancies, respectively, according to equation (1.5).

𝐹𝑒 = 𝐹𝑒 2+ + 2𝑒 - (1.5)
At the FeO/Fe3O4 interface (( 2)), Fe3O4 is reduced by Fe ions and electrons according to equation (1.6).

𝐹𝑒 2+ + 2𝑒 -+ 𝐹𝑒 3 𝑂 4 = 4𝐹𝑒𝑂 (1.6)
The remaining unreacted ions and electrons continue to diffuse via defects in the tetrahedral and octahedral vacancies in the Fe3O4 layer and react at the Fe3O4/Fe2O3 interface (( 3)) according to equation (1.7).

𝐹𝑒 𝑛+ + 𝑛𝑒 -+ 4𝐹𝑒 2 𝑂 3 = 3𝐹𝑒 3 𝑂 4 (1.7)
There are two modes in the Fe2O3/O2 layer, one is the outward migration of Fe ions and the other is the ionization and inward diffusion of oxygen.

If iron ions are movable in Fe2O3, they will migrate outward via cavities through Fe2O3 to the Fe2O3-air interface, forming new Fe2O3 according to the following equation.

4𝐹𝑒 3+ + 12𝑒 -+ 3𝑂 2 = 2𝐹𝑒 2 𝑂 3 (1.8)
When oxygen ionizes at this interface and diffuses inward:

𝑂 2 + 4𝑒 -= 2𝑂 2- (1.9)
In the Fe2O3 layer, if the oxygen ions are mobile, it is the excess iron ions and electrons that react with the inwardly migrating oxygen ions to form new Fe2O3, according to the following equation.

2𝐹𝑒 3+ + 3𝑂 2-= 𝐹𝑒 2 𝑂 3 (1.10)
The corresponding electrons migrate outward through Fe2O3 to reach the Fe2O3air interface to participate in the ionization reaction of oxygen.

The FeO layer is much thicker than the Fe3O4 and Fe2O3 layers due to the high mobility of ions in the FeO layer with more vacancy defects. In fact, at 1000ºC, the thickness proportion of the three oxide layers is approximately: FeO: Fe3O4: Fe2O3 = 95:4:1 [133,134]. Below 570ºC, the oxidation rate is relatively low due to the absence of FeO. Above 570ºC, Fe has a fast diffusion rate and a thick oxide layer can be formed soon. Despite the relatively high plasticity of the FeO layer, the adhesion of the metal is lost and a porous FeO inner layer is formed. The stresses associated with the fastgrowing oxide layer undoubtedly cause physical defects in the outer oxide layer and the infiltration of gas molecules. In particular, the CO/CO2 and H2/H2O redox systems play a role in oxide formation.

Since FeO dominates in the oxide layer above 570ºC. The growth of this layer controls the overall oxidation rate. However, since the concentration of defects in FeO is determined by the reaction equilibrium at a given temperature, the parabolic rate constant is relatively unaffected by the external oxygen partial pressure. In the gas phase, the increase in oxygen partial pressure theoretically leads to an increase in the relative thickness of the Fe2O3 layer. Nevertheless, because this layer accounts for only 1% of the total thickness, the variation of the oxidation rate with oxygen pressure is difficult to notice.

From the above, it can be found that when the temperature is above 570 ℃, FeO has a high defect concentration, which provides better conditions for diffusion, so that it cannot effectively prevent the rapid occurrence of the reaction, and the oxide layer can be generated and thickened rapidly. At this point, FeO becomes dominant layer and controls the growth of the oxide layer. Below 570 ℃, no FeO layer exists, and the Fe3O4 layer is attached to the substrate, effectively preventing the diffusion of ions and inhibiting the rapid oxidation, so the oxidation is not rapid below 570 ℃, and the thickness of the oxide layer is relatively low.

During the wear process, the oxide formed on the surface plays an important role in the wear resistance. The composition, structure and adhesion of oxide layers have a significant impact on the wear of the alloy, especially on the high temperature wear behavior. Unlike static oxidation, the oxidation reaction on the rubbing surface is influenced not only by the friction conditions (temperature, pressure, speed, etc.) but also by the physicochemical structure and properties of the friction surface [113].

The activation energy required for friction reactions is much smaller than that required for thermochemical reactions, e.g., the activation energy for the friction reaction of 4Fe+3O2=2Fe2O3 requires only 0.7 kJ/mol, while its activation energy required for thermochemical reactions reaches 54 kJ/mol [135]. The rate of oxidation reaction during friction is thousands or even tens of thousands of times faster than the rate of thermochemical reaction. Quinn[136,137], based on diffusion theory, suggested that the rate of oxidation reaction during wear depends on the diffusion rate of oxygen ions into the surface layer or the diffusion rate of metal ions through the oxide layer to the outside. The oxidation follows a parabolic law and the oxidation rate increases exponentially with temperature. However, some experimental results are not in accordance with the parabolic theory proposed by Quinn, such as the study by Smith the oxidation rate depends on the interface between the oxide and the substrate rather than the diffusion through the oxide layer, which is a linear law of oxidation kinetics.

Molgaard [141] studied the oxidation behavior, the thickness of the oxide layer and the transfer during dry sliding wear of steel and found that the oxidation behavior was more in line with the parabolic law at temperatures of 300-800°C but more in line with the linear law when there were a large number of cracks and severe peeling in the oxide layer.

There is a great correlation between the contact temperature of the frictional surface and the formation of oxide layers. Some scholars [142,143] have suggested that oxidation occurs on the friction surface of steel even at very low average surface temperatures. However, others [START_REF] Quinn | Application of the oxidational theory of mild wear to the sliding wear of low alloy steel[END_REF]144,145] have suggested that surface oxidation is related to the flash point temperature of the actual contact area, where oxidation generates initially at these points. For low-alloy steels, the different oxides are produced at different temperatures. Some authors [142] found that the oxide layer formed under wear conditions is dominated by Fe3O4 and Fe2O3, unlike the static oxidation where FeO is dominant. The decrease of FeO is owing to the activation of friction that promotes the oxygen uptake process on the metal surface, and the decrease of the supply of iron ions at the interface because of the detachment of the oxide layer from the metal.

There is also an impact of alloying elements on the formation of oxide layers.

Oxidative wear mechanism

In 1930, Fink[156] first proposed the role of oxidation in the wear of steel. It was found that in the ambient environment, the oxide layer generated on the surface of steel could reduce the metal-metal contact and effectively protect the steel; for the vacuum environment, oxidation could not occur due to the absence of oxygen, and the lack of the protective oxide layer caused more severe wear. Subsequently, Rosenberg and Jordan et al.[157] performed wear experiments with carbon steel in air and hydrogen, respectively. It was found that the wear rate in air was only 1/50th of that in hydrogen, thus the contribution of the oxide layer to wear was inferred, which supports Fink's conclusion.

In 1956, Archard and Hirst [START_REF] Archard | The wear of metals under unlubricated conditions[END_REF] carried out pin-ring experiments on several metals and alloys. Based on the results, they classified the wear behavior and firstly proposed mild wear and severe wear. At low velocity and load, an oxide layer generated and covered on the wear surface due to oxidation, which reduces the wear rate. At this stage, the worn surface is smooth and the wear debris is fine. There is a linear relationship between the wear volume and the sliding distance, and the wear tends to be steady state.

As the load and sliding velocity increasing, a severe wear is observed, resulting in a rough worn surface, and accompanying by the peeling and material transfer from the contact surface. Meanwhile, the initial steady state is broken and the wear rate increases dramatically with sliding distance, which is not in line with the linear relationship under mild wear.

For the mild-severe wear theory, Kerridge and Lancaster [START_REF] Kerridge | The stages in a process of severe metallic wear[END_REF]158] conducted experiments on radioactive brass and carbon steel to explore the process of the severe wear. This is the first time that the concept of a transfer layer has been proposed. During severe wear, the soft material on the contact surface is continuously transferred to the harder material surface and a transfer layer is formed. Also, as the wear continues, the transfer layer delaminates and flakes off, resulting in a dynamic equilibrium. In addition, they also suggested that during both mild and severe wear, an oxide layer would be formed, and there were only differences in the properties of oxide layer and the transfer layer. Combining the results of various investigations, Quinn[140] and Wilson [START_REF] Wilson | The development of wear-protective oxides and their influence on sliding friction[END_REF] presented their respective theories of oxidative wear. An earlier study by Quinn[159] determined that the peeling off of metals causes rapid wear, whereas the peeling off of the reactive layer on the surface causes a much slower wear. The reactive layer is generated in addition to the ambient atmosphere, but also in oil lubrication conditions.

Later, an oxidative wear theory model was proposed under room temperature, high sliding velocity and low load conditions. The schematic diagram of this oxidation wear theory is given in Figure 1.4. With the friction process, the contact temperature increases, which leads to the rapid formation and thickening of the oxide layer, especially in the area of the contact junction, where the oxide layer grows significantly.

The formation of the oxide layer enhances the wear resistance. When the thickness of the oxide layer reaches a critical value, it flakes off and causes wear, but the wear volume is quite slight. The oxidative wear theory model by Wilson et al.[44] is predicated on high temperature, low sliding speed and low load conditions, and a schematic diagram of this oxidative wear theory is given in Figure 1.5. Unlike Quinn's oxidative wear model, Wilson et al. considered that the oxide layer on the wear surface gradually flakes off during the friction process. In the early stage of wear, there is no oxide layer, and it is severe wear. With the accumulation and compaction of oxide particles on the worn surface, a thick oxide 'glaze' is formed which rapidly reduces the wear rate. They believe that the oxide particles formed on the wear surface might be removed to cause wear, and might also form a laminar structure on the wear surface to avoid direct contact between the two surfaces, thus playing a protective role, which is the oxidation-flakingreoxidation mechanism. In other words, there are two relative processes of oxide layer formation, namely, oxide sintering and oxide flaking. If a dense glaze layer formed by oxide particles occurs on the wear surface, which can significantly reduce the wear rate, otherwise it is easy to flake off from the rubbing pair, increasing the wear rate.

In summary, the theories of Quinn and Wilson are based on different working conditions, so the morphology and formation mechanism of the oxide layer during the friction process are different. The former is generated by in-situ reaction, while the latter is formed by sintering of abrasive debris. However, they both agree that the oxide layer plays a decisive part in the oxidative wear, and that the formation of the oxide layer leads to a reduction in the wear rate, exhibiting mild wear.

Subsequently, in 1973, Suh et al.[160] investigated another type of wear. Under dry sliding condition, the generated abrasive debris forms a smooth surface without transfer, i.e., delamination wear. The theory proposes the concept of subsurface crack expansion, which eventually leads to wear surfaces delaminated in flakes and the surface layer can undergo large plastic deformations [START_REF] Suh | The delamination theory of wear[END_REF][START_REF] Suh | An overview of the delamination theory of wear[END_REF]160,161]. When experimented in the absence of oxygen, the wear debris is generated by the fatigue of the surface. Hong et al.[139] proposed a modification of Quinn's oxidative wear theory, which suggests that the wear rate is linearly proportional to oxidation and to the actual contact area of the friction surface. However, based on many wear experiments and microscopic analysis of worn surfaces, there is little evidence to support Quinn's theory that the wear rate is proportional to the oxidation rate[162].

Lim and Ashby[163] distinguished oxidative wear as two states, mild oxidative wear and severe oxidative wear. In this context, mild or severe refers to the degree of oxidation rather than the degree of wear.

In the literature[164], the effect of the formation and morphology of oxide layer on wear under sliding conditions was investigated. It was shown that the formation of a smooth oxide glaze significantly reduces the COF and wear rate. It is also believed that the oxide glaze formed in sliding has a high brittleness, which reduces the adhesive junction growth. As a result, the COF is lower than that of metal-metal contact. These glazes can be formed prior to sliding, as well as through the compaction of oxide abrasive debris during sliding [START_REF] Stott | The influence of oxides on the friction and wear of alloys[END_REF]. Oxide glazes are composed of fine crystalline oxide particles with a diameter of 10 to 50 nm. Ashby et al.[165] studied surface temperatures in dry sliding. A calculation of the average (or "overall") and local (or "flash") temperature of the contact surface was

given. The drawback of this work is that it does not consider the oxide film thickness, which has a considerable effect on the surface temperature of dry sliding. So[162,166] showed that the oxidative wear is determined by the positive pressure and sliding velocity. The wear rate is directly related to the rate of oxide layer peeling off from the surface. However, the flaking of the oxide layer is caused by fatigue, and the oxide that peels off is only part of the whole oxide layer. Therefore, the wear rate is not linearly related to the positive pressure and the actual contact area.

In most cases, the oxide layer may reduce the metal-to-metal contact in the relative sliding, thereby reducing wear. Therefore, oxide layers are of great importance in determining wear rates. Iwabuchi et al.[164] simulated the effect of oxides in dry friction. The transition from severe to mild wear was achieved by adding oxide particles to the wear surface during sliding, which eventually reduced the wear volume by a factor of 10. Moreover, it was found that the variation of the wear rate corresponds to the variation of the oxide composition [167]. During the dry sliding of iron alloys, three types of iron oxides (α-Fe2O3, Fe3O4 and FeO) can be formed. Their composition varies with the contact temperature, depending on the conditions of the wear experiments, especially the environmental temperature and the sliding speed. Other studies [START_REF] Clark | Paper 12: mild wear of unlubricated hard steels in air and carbon dioxide[END_REF][168][169][170][171]] also found that the oxide layer generated on the metal surface during sliding has a significant effect on the tribological behavior, and the composition of the oxide significantly determines the wear rate. Mølgaard[172] pointed out that the oxides of iron formed in low-alloy steel under dry sliding vary with contact temperature, and their main oxides vary with temperature as follows: α-Fe2O3 appears below 200°C, Fe3O4 at 200 a 570°C, and FeO above 570°C.

The type of oxide is influenced by the sliding velocity in addition to the temperature.

Low sliding velocity forms α-Fe2O3 and Fe3O4, while FeO is formed only at higher sliding velocities. Godfrey[171] found that Fe3O4 forms a more protective layer than a α-Fe2O3 on worn surfaces, and that the wear rate is low when Fe3O4 is predominant.

Clark et al. [START_REF] Clark | Paper 12: mild wear of unlubricated hard steels in air and carbon dioxide[END_REF] suggested that the main factor for the decreasing wear rate is the change in the nature of the oxide layer associated with the varying temperature. At 500℃, the main protective oxide is Fe3O4, when the main oxide is Fe2O3, the wear is more severe.

There is still some disagreement about the contribution of FeO. Bisson et al.[173] and Cornelius et al.[174] found that FeO is protective, while Tsuji[175] and Hurricks[176] revealed a significant increase in wear rate between 500 and 600°C when FeO is the dominant oxide.

Most of the previous investigations on oxidative wear have adopted oxide layers as a protective layer to reduce wear. It has been reported that the wear rate can be reduced by two orders of magnitude [177]. Therefore, the oxidative wear is closely associated with the oxide layer and not with the substrate. At most, the substrate needs to be strong enough to suppress additional deformation during wear in order to support the oxide film [162,166]. Otherwise, severe wear can occur [140,159,178,179]. The wear mechanism depends mainly on the test conditions [167]. There is almost no significant impact of changes in the microstructure of the steel on the transformation of the wear mechanism. This agrees with the theory of oxidative wear in the case of mild wear. Therefore, the microstructure does not lead to significant differences in the wear volume.

In recent researches, there is some controversy about the role of the substrate in oxidative wear [START_REF] Wang | Wear behavior of steel 1080 with different microstructures during dry sliding[END_REF]177,180,181]. An increasing number of investigations have shown that oxidative wear does not only depend on the oxides, but that the microstructures and properties of the substrate also have an impact on the wear behavior.

In the literature [180], the effect of oxidation on the wear resistance of ductile iron, gray cast iron and creep cast iron was investigated. The effect of oxidation on wear rate was found to be determined by graphite morphology and microstructure of substrate.

When there is high hardness in cast iron, oxidation reduces the wear rate of gray cast iron, but increases the wear rate of ductile and creep cast iron. When there is a low hardness in cast iron, the rule is exactly the opposite. It can be found that the effect of oxidation on wear is strongly related to the microstructures and properties of the substrate.

Eargles et al. at high loads, the wear is severe, but the wear mechanisms are not clearly distinguished yet, and they are all generally considered to be a mixture of severe wear such as abrasive wear, adhesive wear, and oxidative wear.

From the above analysis, we can determine that there are two different types of oxidative wear, which are not distinguished by friction oxides, but mainly by the degree of softening and deformation of the substrate. When both the oxide layer and the substrate affect the wear behavior, the wear mechanism is different from the theory of mild oxidative wear proposed by Quinn and Wilson et al. [START_REF] Wilson | The development of wear-protective oxides and their influence on sliding friction[END_REF]159]. However, so far, there seems to be no clear correlation between oxidative wear and microstructure. In addition, there are few reports relating the oxidative wear mechanism to the variations in microstructure during sliding process. 1.3.4 The tribological behavior of oxide scales

The study of tribological behavior of oxide layers at the microscopic level is vital As the properties of oxides are highly dependent on their microstructural characteristics and mechanical properties, which results in a more complex tribological behavior of the oxide layers. Yu et al.[191] performed tribological experiments on lowcarbon microalloyed steels and found that the oxide layers formed with different microstructures of Fe3O4 exhibited distinct tribological behavior. The spherical structure may contribute to being a lubricant and therefore resist wear, while the lamellar structure has a stronger effect on adherent wear. The fine lamellar structure leads to a lower COF, the coarse one results in an increased roughness. And the adhesion of oxide layer to substrates varies with different structures. Rojacz et al.[192] investigated the effect of different alloying elements in iron alloys on oxidation kinetics and wear resistance of oxides at 700 °C. Chromium oxide was found to be highly correlated with the frictional stability of the oxide layer. Cr, Ti, and Si showed significant relevance to the oxidation stability and the adhesion of oxide scale. And Ni reduced the brittleness of the oxide scale and enhanced the tribological properties. The scratch test results [193] showed that the oxide scales formed on AISI 445 stainless steel exhibits different adhesion strengths, and the COF increases with loads. The oxide scale caused an increase in the COF compared to that without oxide scale. In the dry sliding pin-on-disk experiment, the COF was slightly lower with an oxide layer on the disk, which unlike the scratch test. The formation of chromium oxide layer reduces adhesive wear, but causes abrasive wear. The chromium-rich oxide layer is well bonded to the substrate, which contributes to wear resistance. Cheng et al.[194] suggested the medium thickness layer formed in static dry air for 120 mins showed the maximum COF and high surface roughness. The thick oxide layer does not always have a positive effect on reducing the COF, which not only depends on the thickness of the oxide scale, but also on its composition and surface morphology. Munther,P.A. and J.G. Lenard[195] revealed that the thickness of oxide scale seems to have a greater impact on friction than oxide composition. And a thicker oxide layer generally generates lower friction values. the same results were confirmed by Luong,L. and T. Heijkoop[196].

In this work, investigations of oxide layer properties have focused on medium-to long-term oxidation (>4 h) at high temperatures of 600-1200°C. For short-term oxidation (<4 h) at lower temperature ranges (<600°C), the wear and friction behavior of the oxide layer is still lacking. Accordingly, the materials investigated for oxide layers are mostly focused on stainless steels, hot-rolled steels, and high-speed steels.

Few researches have been reported for low-alloy steels, especially for dual-phase steels.

Therefore, in this paper, the effects of nature, thickness, and matrix microstructure of oxide layers on friction and wear of oxide scales are discussed in detail.

Research objective and value

As we known, during the friction process between the surfaces, even at room temperature and low speed, there is oxidation on the contact surface due to the heat flux.

In this subject, we can increase the oxidation in the following ways:

(a) Increasing sliding speed and contact pressure (b) Injecting more oxygen: O2, CO2

For(a), owing to the frictional heat flux is related to COF, pressure, and sliding speed [125,197], experiments were done with velocity 0.1m/s, 0.6m/s, and 1m/s to study the oxidation. For (b), four different gaseous atmospheres are used for the experiments: air, argon, nitrogen, and carbon dioxide. The air atmosphere is used as a reference. And the oxygen concentration in carbon dioxide is higher than that in air, while the oxygen concentration in nitrogen and argon is lower than that in air.

For the investigation of oxide scales, since the temperature of oxide formation:

Fe2O3 (T <200°C), Fe3O4(T> 400°C), and FeO (T> 570°C), oxidative experiments were made at 200℃, 400℃, and 600℃ to acquire different types and thicknesses of oxide

Conclusion

This literature review shows that the microstructure is a fundamental factor in determining the tribological behavior. Under dry sliding conditions, there is an essential role for microstructures on friction and wear of the material. In fact, most of the literatures are carried out for various tribological systems that result in different wear mechanisms. It is meaningless to compare these results quantitatively. Even under the same wear mechanism, there is a lack of literature to determine an optimal microstructure that provides the favorable friction and wear resistance for the given material in a systematic way. In addition, there is still a relative absence of studies on the effect of microstructure on the oxidative wear. Some of the work in this bibliography overview revealed the effects of gaseous environment, sliding velocity, and temperature on the oxidative wear, respectively. However, the combined effects of the microstructure with the above factors on the oxidative wear is unexplored. Finally, investigations on the tribological behavior of the oxide layers with different compositions are still missing.

In this case, the tribological behavior of the same low alloy steel (25CD4) with the same chemical composition, especially oxidative wear, is investigated in this thesis.

The different microstructures were achieved by variations of the heat treatment. Ballon-disk dry sliding experiments were performed with samples of different microstructures at the same surface roughness and normal load. The tribological system parameters can be modified by varying the gaseous environment (Air, Ar, N2, and CO2)

and the sliding speed (0.1, 0.6, and 1m/s) to modulate the factors affecting oxidative wear. These tests permit to study the separate impact of microstructure of the same steel on the tribological behavior under the given tribological system. For the same microstructure, the effects of changes in tribological system parameters on friction and wear behavior can also be compared. Meanwhile, the combined effects of microstructure and tribological system factors (sliding velocity and gas environment)

on the tribological behavior were also analyzed. Besides, in this thesis, scratch experiments on different oxide layers were tested by a nano-indenter to examine the tribological behavior of the oxide layers formed at varying microstructures and oxidation temperatures.

Material and microstructure

This chapter is devoted to the selected materials and the generated microstructures in order to better understand the contribution of microstructures to the tribological behavior of low-alloy steels. First, the reasons for the chosen steels and their properties are presented. Secondly, the methods of generation for different microstructures are introduced. The investigation of microstructure in this thesis concentrates on the variation of MVF. Since hardness can be changed by a variation of microstructural characteristics, the third part of this chapter is devoted to the measurement of macrohardness and microhardness.

Experimental materials

The main objective of this chapter is to develop identical microstructures of the same steel with different microscopic characteristics in order to compare their tribological behavior. The microstructure is obtained by different heat treatments for selected low-alloy steel. Therefore, the steel is selected by analyzing CCT (continuous cooling transformation) and TTT (time-temperature-transformation) diagrams of different low-alloy steels. After austenitization, these diagrams can describe the specific phase transformation of the steel. The CCT chart predicts the phase transformation after continuous cooling as well as the final macro-hardness, while the TTT diagram depicts the phase transformation during isothermal holding. Thus, with these charts, it is possible to predict the final generated microstructure, as well as the final macrohardness, as a function of the heat treatment for a given steel.

For a given microstructure, the variation in microstructural characteristics can be modulated by the grain size, the morphology of the phases, and the volume fraction of these phases. Therefore, the selected steel should have the ability to produce the same microstructure, while providing different tribological properties due to different microstructural characteristics. The chemical composition of 25CD4 is shown in Table 2.1. The diameter and the thickness of each sample were 55 mm and 3 mm, respectively.

Heat treatments

The specimens were subjected to step quench heat treatments in order to obtain dual phase microstructures with martensite and ferrite. As shown in Figure 2.2, this is done in two stages: first, a full austenization was performed at 900 °C for 30 minutes;

then, to achieve the dual phase region consisting of austenite and ferrite, the specimens were kept at an intercritical temperature of 725 °C. The established microstructures at 725 °C were composed of remained austenite and ferrite, which transformed from austenite. Finally, the specimens were quenched in water to achieve the transformation from austenite to martensite, while ferrite remained unchanged. Nonetheless, during the intercritical heat treatment at 725 °C, six samples were held for six different times ranging from 0 to 7 minutes to vary the volume fraction of martensite and ferrite. Three samples were prepared for each constant holding time. The purpose of austenitization is to dissolve all the carbon existed in the alloy. As

shown in Figure 2.3, the solubility of carbon in the ferritic phase is 0.0218%, while that in the austenitic phase is 2%. Therefore, the redissolution of carbon is achieved by increasing the temperature during the formation of γ-austenite. To make austenite transformation occur, the temperature of the sample must be raised higher than the transformation temperature A3. As shown in Figure 2. Besides, the carbon content directly affects the hardness of the martensite. 2.3 Microstructure characterization

Method of obtaining martensite volume fraction

The six specimens with unequal MVF were mechanically ground on SiC grinding papers of 80 grit to remove 0.5 mm of material from the surface prior to the friction tests. The goal of this procedure was to obtain the bulk material's structure. Then, using

SiC grinding paper with grits ranging from 180 to 4000, mechanical grinding was used.

The surfaces were then polished separately with 3 µm and 1 µm diamond paste. The polished surfaces were then cleaned in ethanol for 10 minutes using ultrasound.

The specimens were then etched with a 1 % Nital solution for 3 seconds before being cleaned with ultrasound and ethanol for 10 minutes. The microstructures of each sample were examined using a ZEISS Sigma Field Emission Electronic Gun-Scanning

Electron Microscope (FEG-SEM) at 20 keV to obtain the MVF. The volume fraction of martensite and ferrite was calculated using ImageJ image processing.

Generated microstructure

The volume fraction of ferrite and martensite at a given critical annealing temperature is known to be dependent on the holding time [202,203]. SEM images of dual phase microstructures composed of martensite and ferrite at various holding times are shown in Figure 2.7.

To the beginning, a complete austenitic microstructure is formed during the austenitization process. Then, as the temperature drops from 900 °C to the critical temperature of 725 °C, ferrite nucleates at the austenite grain boundaries. Finally, the austenite undergoes a water quenching transformation to martensite, producing very little retained austenite while the ferrite formed at the critical temperature remains unchanged. It is widely acknowledged that the amount of retained austenite is proportional to the carbon content of the steel. The residual austenite in steels with less than 0.5 % carbon is less than 2% [204]. As a result, for 25 CD4, only martensite and ferrite can be seen in the micrographs, with very little retained austenite (less than 2%) [START_REF] Xu | The effect of ferrite-martensite morphology on the scratch and abrasive wear behaviour of a dual phase construction steel[END_REF]. Micro-hardness measurements were performed with a Vickers Zwick ZHVµs(H115) under a load of 0.01 kg. The micro-hardness values reported were an average of three measurements taken at the same time in different locations. At each phase, two positions at the edge and one position in the center were chosen for each measurement.

Result of macro-hardness

Figure 2.9 depicts the macro hardness as a function of MVF. The specimen with the greatest macro hardness is the one that has been directly water-quenched after austenitizing, and its macro hardness exceeds 600 Hv. The macro hardness of the samples gradually decreases as the MVF decreases, i.e. as the holding time increases.

When the holding time is 7 minutes, that is, when the volume fraction of martensite is 48%, the minimum value of macro hardness is around 300 Hv. Many previous studies [START_REF] Nadlene | Study on the effect of volume fraction of dual phase steel to corrosion behaviour and hardness[END_REF]205,206] have found that increasing macro hardness is caused by an increase in the hard MVF and a decrease in the soft ferrite volume fraction. It has been established that the carbon content of a phase determines its hardness. Martensite, as a metastable iron phase, has a much higher amount of carbon content than in solid solution. Martensite has a very high hardness because of the transformation from austenite to martensite. Furthermore, the maximum carbon content dissolved in ferrite is only 0.02 %, resulting in ferrite having a much lower hardness than martensite. According to some studies [START_REF] Modi | High-stress abrasive wear response of 0.2% carbon dual phase steel: Effects of microstructural features and experimental conditions[END_REF][START_REF] Tyagi | Effect of martensite content on friction and oxidative wear behavior of 0.42 Pct carbon dual-phase steel[END_REF][START_REF] Xu | The effect of ferrite-martensite morphology on the scratch and abrasive wear behaviour of a dual phase construction steel[END_REF], even if the hardness of the individual phase follows the opposite trend, the overall macro-hardness of the dual phase steel increases with the MVF. Furthermore, as described in some studies [START_REF] Movahed | The effect of intercritical heat treatment temperature on the tensile properties and work hardening behavior of ferrite-martensite dual phase steel sheets[END_REF]207], the carbon concentration in separated martensite increases with decreasing MVF, resulting in an increase in microhardness.

The carbon concentration of the dual phase steel can be calculated by the following formula:

𝐶 = 𝐶 𝑓 𝑉 𝑓 + 𝐶 𝑚 𝑉 𝑚 = 𝐶 𝑓 (1 -𝑉 𝑚 ) + 𝐶 𝑚 𝑉 𝑚 (2.1)
Further, the carbon concentration of martensite can be expressed as:

𝐶 𝑚 = 𝐶 𝑓 + (𝐶 -𝐶 𝑓 )/𝑉 𝑚 (2.2)
Among them, Cm is the carbon concentration of isolated martensite, Cf is the carbon concentration of ferrite, C is the carbon concentration of dual phase steel, Vm and Vf are respectively the MVF and ferrite volume fraction. According to the above expression, the carbon concentration of martensite decreases with the increasing MVF.

For lower MVF, the carbon content of this phase is high, in contrast, for phases with high MVF, the carbon content is relatively low. Therefore, the main reason for the macro hardness increasing may be the increase of the volume fraction of the hard martensite phase.

Result of microhardness

The results of microhardness experiments, as shown in Figure 2.10, also support the preceding conclusion. Furthermore, the microhardness measurement points chosen for the experiment are located at the edge and center of each phase (Figure 2.11).

Further experiments show that the microhardness in the center of martensite is significantly higher than near the edge. The carbon at the austenite grain boundary is transferred from the edge to the center during the nucleation of ferrite in the austenite grain boundary, which is more beneficial to ferrite nucleation and growth due to the lower amount of dissolved carbon in the ferrite. Because the carbon content in the center of the formed martensite is higher than in the edge after water quenching due to the transformation from austenite to martensite, the microhardness at the center is greater than that of the edge. Because the amount of carbon dissolved in ferrite is so small, there is no discernible difference in microhardness between the center and the edge of ferrite.

Conclusion

In this chapter, the generation of different microstructures is described in detail to understand better the role of microstructure in tribological behavior. The low-alloy steel 25CD4 was chosen since they are capable to generate these distinct microstructures with the same chemical composition by heat treatment. Different dual-phase microstructures (consisting of ferrite and martensite) were obtained with the martensite volume fraction as a characteristic. Meanwhile, for the same steel, the increase in macro-hardness of the dual-phase microstructure is more related to the proportion of phases than to the local hardness of the phases. Thereby, in the following chapters, the tribological behavior of these samples will be investigated more specifically. The aim is to analyze the effect of martensite volume fraction on a dual-phase microstructure composed of ferrite and martensite. Afterwards, other factors influencing oxidative wear are coupled for a comprehensive analysis.

Effect of martensite volume fraction on oxidative wear in air

This chapter consists of two main parts, mainly focusing on dry sliding experiments in an air environment. The first part (section 3.1-3.4) analyzes the effect of microstructure on oxidative wear in terms of friction coefficient, wear rate, and wear mechanism. Among them, the method of dry sliding experiments and the corresponding variation result of the friction coefficient are described in section 3.1. The section 3.2

presents the calculation method and the results for the wear rate, respectively.

Regarding the wear mechanism, the microscopic analysis, plastic deformation, and wear debris analysis are studied in the section 3.3. Then, the section 3.4 is a discussion of the effect of microstructure on oxidative wear. The second part (section 3.5) focuses on the effect of sliding velocity on oxidative wear. Similarly, three aspects of friction coefficient, wear volume, and wear mechanism are analyzed and discussed. In addition, the combined effects of sliding velocity and MVF on the tribological properties of dualphase steels are investigated.

Friction coefficient

Experimental method

Dry sliding tribological tests were carried out on a ball-on-disc rotating test rig

Result of friction coefficient

Wear rate

Calculation of wear rate

To calculate the wear rate, the topography of the worn samples was analyzed with a 3D optical profilometer (Sensofar Metrology). Figure 3.4 depicts the methods and steps. Three different locations were chosen as measurement samples for each scar using the SensoSCAN software. Similarly, the SensoMAP application was used to extract three profiles from each surface to determine the area of the peak and the hole.

Finally, the wear track's final wear accumulation amount was represented by the average of the nine measurements. Where 𝑆 ℎ𝑜𝑙𝑒 and 𝑆 𝑝𝑒𝑎𝑘 are the area of hole and peak, 𝐿 is the length of the scar and 𝐷 is the total sliding distance.

The final wear rate under the same experimental parameter was determined by the average of at least three tests. The trend of wear rate with relation to MVF is shown in Figure 3.5. There is low wear on the sample with low MVF (59% or less) compared with high MVF (93% or more). In addition, it can be seen that the minimum wear occurs on the sample with the MVF of 59% and the maximum occurs on the sample with 93% MVF. Meanwhile, for the samples with MVF between 59% and 93%, the wear rate increases with the increasing MVF.

Result of wear rate

Wear mechanisms

Microanalysis

After tribological tests, wear scars were observed by a Field Emission Electronic Gun-Scanning Electron Microscope (FEG-SEM) ZEISS Sigma. Besides, EDX analysis was performed for the wear debris and scars to investigate the wear mechanisms. The results of EDX analysis on the points of the wear track surface are shown in Figure 3.7 and Table 3.2. It can be seen that the oxygen content of the wear debris was high, which also indicates that significant oxidative wear occurred during the experiment. It can be found that the dark or black areas of the wear track have higher oxygen concentrations, while the light gray areas have low oxygen concentration (3%).

From the EDX results and SEM micrograph analysis, a great number of tribo-oxides appeared on worn surfaces of dual-phase steel as MVF decreased. The morphology of worn surfaces gradually presented smooth un-delaminated regions which were typical tracks of oxidative wear. Therefore, the previously mentioned compression film formed by wear debris can be determined as an oxide film. Due to the existence of the oxide film, it effectively prevents direct contact with the frictional matrix. As the friction continues, the debris of oxides that remains in the wear track is better compacted on the contact surface. This also promotes the formation and maintenance of the oxide layer.

Due to the existence of the hard oxide layer, the wear rate gradually decreases, especially the adhesive wear, which was also observed in this experiment.

Plastic deformation

Calculation of plastic deformation

The calculation about the volume of plastic deformation during the rubbing process is shown in Figure 3.8. The cross-sectional profile of the track was achieved by a 3D optical profilometer in three different positions of the wear track (Figure . 3.8(a)).

The area of material accumulation due to extrusion on both sides of the wear track is defined as the area of plastic deformation at this cross-sectional location, as indicated in Figure 3.8(b) and 3.8(c). The values are computed by the software SensoMAP. The total volume of plastic deformation for the worn track is determined by the average cross-sectional plastic deformation area and the track length, which can be calculated as follow.

𝑉 𝑝𝑑 = 𝐴 𝑝𝑑 × 𝐿 (3.2)
Where, 𝐴 𝑝𝑑 is the average area caused by plastic deformation for nine measurements after obtaining the cross section. L is the length of the worn track. specimen with a MVF of 59%, which is also the specimen that exhibits the lowest wear rate. Consequently, this also provides a supplement to explain the variation in COF and wear rate in these experiments.

Result of plastic deformation

Wear debris

Each of the various processes by which material can be lost from a surface in service leaves its fingerprint both in the topography of the worn surface and in the size, shape and number of the particles which make up the wear debris. Since debris concentration and size both increase by increasing of wear degree, they can reflect wear severity and wear rate [211,212]. In some literatures [113,213], the wear of alloys is affected by the existence of debris, especially when the debris can be retained on the contact surface under reciprocating conditions. During sliding, the wear debris kept within the track is crushed by repeated plastic deformation and oxidation. However, once they have been reduced to a sufficiently small size, they are agglomerated at certain locations, particularly in grooves, due to adhesion forces between solid surfaces arising from surface energy[214], and develop compact layers which eventually become load-bearing. The wear is also reduced to lower levels.

In this experiment, only EDX was used for the elemental analysis of the wear debris, as the number of wear debris produced was small and difficult to collect. It can be found that the wear debris consists mainly of metal oxides, the results of which are also confirmed in Figure 3.10. And it also shows the evolution of abrasive particle size as a function of MVF, which can be concluded that the size of debris decreases with the MVF. For the specimens with high martensite volume fraction, the size of the wear debris is relatively large, with an average diameter around 8 µm. Meanwhile, an aggregation of small wear debris was observed on the larger wear particles. As the martensite volume fraction decreases, a decrease in the overall wear debris size can be noted, despite the uneven size of the generated debris. For samples with lower martensite volume fractions, a large number of wear debris with diameters less than 1 µm are independently existed. The shape of the wear debris is smoother and more rounded.

Discussion of the effect of martensite volume fraction on friction coefficient and wear rate

Combining SEM images and EDX analysis results, the tribological behaviors of dual-phase steel are analyzed and discussed as the MVF varies.

As we all know, metal oxides generally provide relatively favorable friction properties [215,216], and direct contact of metallic Fe causes severe adhesion [209,210].

The higher adhesion caused by metal-to-metal contact dominates the tribological behavior of the contact points. Previous studies have shown that the increased COF is caused by dominant adhesive wear [START_REF] Trevisiol | Effect of martensite volume fraction and abrasive particles size on friction and wear behaviour of a low alloy steel[END_REF]. In addition, the delamination of the compacted transfer layer and the formation of oxide can also be a foundation. Compared with adhesive wear, the compacted oxide layer can effectively reduce the real contact area and junctions, which may bring down the COF. The results of EDX analysis show that the oxidative wear increases with decreasing MVF, which explains the phenomenon that COF decreases with decreasing MVF.

The main wear of samples with low MVF is oxidative wear. The relatively high oxygen concentration of the wear track confirms this point. The existence of a composite oxide film is like a protective layer on the metal surface [141,166,[217][218][219],

which reasonably explains why the wear of low MVF is lower than that of high MVF.

The increase in plastic deformation represents an increase in oxidative wear. Under the conditions of mild oxidative wear theory, the increase in oxidative wear means a decrease in wear rate.

During the sliding process, adhesion and junctions occur only on the real contact areas, where high contact pressure causes the plastic deformation of the surface. There is a higher temperature at plastically deformed contacting areas. Thus, these plastically deformed areas are preferential locations for oxidation [START_REF] Abouei | Effect of microstructure on the oxidative wear behavior of plain carbon steel[END_REF]166]. As the plastic deformation of the substrate increases, more cracks would be formed and propagated [220,221]. This could easily lead to oxygen entering these cracks and reacting with the steel substrate to form oxides. A large plastic deformation would occur in a softer matrix under the same normal load and sliding speed, and so produce more tribo-oxides. Due to the low normal load and sliding velocity, the plastic deformation generated during wear is helpful to retain more tribo-oxides. The model proposed by Quinn[137,222] can also be used to show that mild oxidative wear is usually generated under low normal load and velocity, and wear property is only related to tribo-oxides.

As the low-speed friction proceeds, the oxide remaining in the wear track is compacted into a hard tribo-oxide layer to prevent metal-metal contact and thus reduce wear.

According to the results of the macro-hardness test and the plastic deformation measurement, it can be determined that the macro-hardness of the dual-phase steel matrix decreases with decreasing MVF, and the plastic deformation increases with decreasing MVF. Meanwhile, owing to the oxide increases during plastic deformation, oxidative wear increases with decreasing MVF. This reasonably explains the relationship between the microstructure of dual-phase steel and oxidative wear.

The effect of sliding velocity on oxidative wear in air

3.5.1 Friction coefficient

Experimental method

The experiments regarding the investigation of sliding velocity were performed on the same experimental device, but in unidirectional rotation mode, as shown in Figure 3.11. This is because the reciprocating rotation mode does not provide a larger velocity variation compared to the unidirectional rotation. The sliding velocities were 0.1m/s, 0.6m/s, and 1m/s, respectively, with a constant load of 15N and ambient environment.

For different sliding speeds, the number of rotating revolutions is 10000. The additional experimental preparation and parameter settings are described in Section 3.1.1. A challenge for this type of wear experiment is the rapidly changing conditions, especially contact pressure, when the ball enters and rubs the disc. The literature [126] demonstrates that the approach using concentric wear tracks of various radius offers identical friction and wear outcomes on the disc surface for tracks worn at different radius [129]. This provides a reliable justification for this thesis to investigate the effect of sliding velocity on the tribological behavior of dual-phase steels at the same MVF. Nevertheless, it is more distinct on the specimens with higher MVF. It can be seen that, both the sliding velocity and the MVF have an effect on the COF. The review [223] indicated that the COF increased significantly with increasing sliding velocity as Si4N3 steel balls rubbing against stainless steel, which is not in consistent with the results of this paper. [224] and the trend shown by So[162]. Besides, it seems that the effect of MVF on the average steady-state COF does not show a distinct regularity under the working conditions of this experiment.

Wear rate

Calculation of wear rate

The wear rate is measured as described as figure 3.14. Since the experiment was performed with a constant number of revolutions instead of the same sliding distance, the wear degree was expressed by calculating the wear rate at a given length in the selected areas. The value of wear rate for the selected area is obtained by the average of 8 locations (shown as figure 4.14(a)). Afterwards, as introduced in section 3. Where, Shole and Speak are the areas of the "peak" and "hole" (Shown as Figure 3.14(c)). V is volume loss, S is total sliding distance, R is radius of the track, and N is number of revolutions. Overall, the wear rate decreases with increasing sliding velocity. The minimum wear rate occurs at a sliding speed of 1 m/s. Compared to specimens with low MVF (MVF 48% and MVF 59%), the wear rate is larger on specimens with higher MVF (MVF ≥ 68%). It becomes more and more obvious as the sliding speed decreases. Especially at a sliding velocity of 0.1 m/s, the wear rate on specimens with high MVF is 4-8 times larger than that of low MVF. This means that, in addition to the sliding velocity, the variation of the MVF also impacts the wear resistance of dual-phase steels. As can be seen from the figure, the diameter of the worn Al2O3 ball decreases with increasing sliding velocity. The maximum worn diameter of Al2O3 balls is observed at a sliding velocity of 0.1 m/s. Moreover, it can be observed that the area covered by black material on the worn surface of Al2O3 ball increases with the sliding velocity. The black adhering material is difficult to be removed from the ball surface. Also, it is hard to directly determine the composition under electron microscopy due to the poorer electrical conductivity of the ceramic alumina balls. However, it can be presumed from the literature [238] that the black material might be an accumulation of oxide wear debris generated during the sliding process.

The average diameter of the worn Al2O3 ball as function of sliding speed in all experiments is given in Figure 3.17. It can be found that the wear trend of Al2O3 balls is consistent with that of the dual-phase steel discs, i.e., the wear decreases with the increasing sliding velocity. Meanwhile, the worn diameter of the ball was lower when rubbed against dual-phase steel with low MVF (MVF 48% and MVF 59%) than against with high MVF (MVF ≥ 68%). 3.3 show the results of the EDX elemental composition analysis on the worn surface points. It is clear that there is a higher oxygen content in wear debris and dark or black regions of the track, indicating that oxidation was generated during the sliding process at the tested positions. There is low oxygen concentration in the light gray zone, which is essentially the same as the oxygen content in the unworn surface, and can be identified as an unoxidized surface. Generally, when sliding occurs between two mating surfaces, frictional heating can raise the contact temperature. If the sliding speed is adequately large, the flash temperature at the contact asperities might result in oxidation. Once the oxide achieves a critical thickness, it may be fractured and wear debris is generated. This is the oxidative wear mechanism, which was firstly modeled by Quinn et al.[145]. Oxidative wear, one of the two basic categories of non-lubricated sliding wear [START_REF] Archard | The wear of metals under unlubricated conditions[END_REF], is characterized by the generation of oxidative wear debris and smooth oxidative worn surfaces. This type of wear evidently concerns reactions with the surroundings, especially with oxygen [137]. In addition, when sliding surfaces are submitted to thermal oxidation, whether the heat is provided by an external heat source or internally by frictional heat generated from sliding contact, an oxidative wear mechanism is induced. In steels, this wear mechanism normally occurs at sliding speed larger than 1 m/s, where the flash temperature is greater than 400°C [109,163]. However, experience demonstrates that oxidative wear might also occur at lower sliding speeds for heattreated high-strength steels, although the thermal effect is not significant enough to cause direct asperity oxidation. This is of great significance since the oxidation wear that occurs at low sliding speeds is commonly related to the achievement of a mild wear. Indeed, in this case, the compaction and sintering of the iron oxide forms a particularly compact and smooth friction layer, also called 'glaze', a name suggested by Stott[225]. The existence of the 'glaze' leads to a reduction in the adhesion and prevents metal-to-metal contact, which avoids severe adhesive wear. This is known as "low-sliding velocity oxidative wear" [226].

In this experiment, as the sliding speed increased from 0.1m/s to 1m/s, it was observed that both the COF and wear volume decreased. When the sliding speed is low (0.1m/s), the frictional heating is not evident and the oxidation of the contact surface is relatively light. At this point, the main wear mechanism is plowing, deep grooves, dominated by abrasive wear, accompanied by material peeling and less oxidation. This is the reason for the higher COF and wear rate at 0.1m/s. As the sliding speed increases to 0.6m/s, oxidative wear is enhanced and a great number of oxidative debris appear on the worn surface. The appearance of oxidized worn debris leads to variations in the COF [227]. Generally, wear debris acts as a roller or ball at the contact interface, reinforcing the movement between sliding parts and reducing the COF [125,228]. For wear, wear particles are typically critical in dry sliding process as a type of "third body" [122,229]. In some situations, if oxide particles are harder than the underlying substrate, they can behave as a third wear body, leading to an increased wear rate, such as in the case of NiAl [230]. On the other hand, oxides formed on iron-based materials are generally softer and more adhesive than other metal oxides so that lead to a slight wear formation called mild oxidative wear [162,163]. Besides, oxides may bear part of the normal load so that the underlying metal matrix is protected from wear [162].

Studies of wear on several high carbon steels showed that there is a transition to lower wear rates as oxide particles form on the sliding surface, whether the oxide particles are formed in situ [231] or are supplied to the asperity from the outside source [164]. For a larger sliding speed (1 m/s), the flash temperature is further increased and the wear debris melted or micro-welded to the contact interface is gradually compacted as the sliding process proceeds, resulting in a smooth oxide 'glaze'. The relatively continuous, smooth, and compact oxide 'glaze' affects the direct metal-metal contact, causing a poorer strength of the coordinated covalent bonds among the adhesion effects.

Meanwhile, the harder oxide layer which fits well with the substrate also reduces the abrasive wear to some extent. At this time, oxidation wear progressively replaces abrasive wear as the main wear mechanism, which explains the lowest COF and wear rate obtained at 1m/s.

In summary, it can be concluded that as the sliding speed increases, high temperatures generated by friction lead to certain chemical reactions between the contact surface of specimens, resulting in a series of wear effects. During the sliding process, although only a small amount of frictional oxide is present, it modifies the wear behavior even the wear mechanism. As the sliding speed increases, the wear mechanism transforms from predominantly abrasive wear at low velocities to oxidative wear at high velocities. The formation of a relatively continuous, smooth, compacted oxide layer effectively limits the COF and wear of the material.

Combining the COF, wear volume, average oxygen content, and SEM images of specimens with different MVF, it was found that specimens with lower MVF (MVF 48%

and MVF 59%) exhibited lower wear volume compared to those with high MVF (MVF ≥ 68%). The reason could be explained that specimens with lower MVF are more prone to oxidative wear during sliding wear. Both the SEM micrograph and the oxygen content of the worn track clearly demonstrate this point. Nevertheless, this observation is more significant at low sliding velocities. With the increase of sliding speed, the effect of MVF on the wear of dual-phase steel decreases progressively, and the sliding velocity then becomes the main factor affecting the wear.

Conclusion

The study in this chapter focuses on the effect of microstructure (MVF) on the tribological behavior of dual-phase steels. In addition, the effect of sliding velocity on oxidative wear was investigated. The dry sliding experiments were all carried out on a ball-disk tribometer, under a constant load (15 N) and an air environment. In order to change the microstructure of the same steel, 25CD4 steel was heat treated in different procedure. The effect of sliding velocity on tribological behavior was investigated by varying the sliding velocity of the tribometer. Next, the experimental findings are summarized in terms of MVF and sliding velocity, respectively.

For the effect of microstructure on the tribological behavior of dual-phase steels, the COF decreases with decreasing MVF. The COF, as the most direct response to the interfacial state during the friction process, has a variation pattern that proves the inevitable link between the MVF and the tribological performance of the dual-phase steel. For the wear rate, the specimens with lower MVF (≤59%) have lower wear rate compared to high MVF under the present experimental conditions. In this investigation, the variation of COF and wear rate were attributed to the evolution of the oxidative wear. Microscopic observation and analysis of the elemental content in the wear tracks reveals that the oxygen content in the wear tracks increases as the MVF decreases, leading to increased oxidative wear. The formation of dense oxide film (glaze layer) can effectively reduce wear, especially adhesive wear. For low MVF, the dense oxide layer is more easily formed and retained on the worn surface, which can be clearly observed in the micrographs. For lower MVF, the main wear mechanism is oxidative wear. However, for higher MVF, adhesive wear and ploughing predominate. Besides, both plastic deformation and wear debris size decrease with increasing MVF. This also provides a reason to believe that oxidation wear increases with decreasing MVF.

For the effect of sliding speed on the tribological performance of dual-phase steels, both the COF and wear decrease with increasing sliding velocity. This is owing to the increase in oxidative wear as the sliding speed increases, which can be justified by microscopic observation of wear tracks and elemental content analysis. The frictional oxide layer induced by the sliding velocity plays a protective role for the metal surface.

The transformation of the wear-affected zone at different sliding velocities leads to a transformation in the wear mechanism and eventually to a modification of the wear behavior. At 0.1m/s, the major wear mechanism is abrasive wear, accompanied by partial material flaking and less oxidation wear; at 0.6m/s, the wear mechanism is in the transition between abrasive wear and oxidation wear; at 1m/s, oxidation wear becomes the dominant wear mechanism. For the same sliding velocity, specimens with lower MVF exhibited lower wear rates compared to specimens with higher MVF. It is especially evident at low sliding velocities. As the sliding velocity increases, the effect of MVF on the wear decreases and the sliding velocity becomes the dominant factor. presents the calculation method and results of the wear rate. The third part (section 4.3)

analyzes the wear mechanism from three aspects: microscopic analysis, plastic deformation, and wear debris, respectively. Finally, the combined effects of gaseous environment and microstructure on oxidative wear of dual-phase steels are discussed and the corresponding conclusions are summarized.

Friction coefficient

Experimental method

For experiments in non-ambient atmospheres, there is an internal cover ( Nitrogen and argon are the shortest transition in all gasses, and the longest transition is the carbon-dioxide gas. The transition time in argon and nitrogen are both between in air and carbon dioxide. However, the transition in the argon environment is slightly longer than that in nitrogen. For Figure 4.2(a), it is interesting that at the beginning of the friction curve, the lowest coefficient of friction was measured in the atmosphere of carbon dioxide; during some tests, the coefficient of friction remained stable over the whole test.

However, after 100s of testing, a steep increase in friction was observed, with the COF exceeding that of other gaseous atmospheres ( In a steady state, the average COF is the average value of the instantaneous COF (i.e. from 100 seconds to the end of the test). The average COF as a function of MVF is shown in Figure 4.7 for various gas environments. It is worth noting that the COF is affected by both the gas environment and the MVF. Overall, the minimum friction coefficient is found in the air environment with a value of 0.49. The maximum friction coefficient occurs in the CO2 environment with a value of 0.71. Furthermore, when the volume fraction of martensite is less than 80%, the influence of the gas environment dominates the effect of the MVF on the coefficient of friction.

Wear rate

Calculation of wear rate

The wear rate can be calculated using the method previously described in Section 3.2.1, which was shown as Figure 3.4. When the volume fraction of martensite is high (93% or more), nitrogen has the highest wear rate, followed by air, argon, and carbon dioxide; The volume fraction of martensite is between 68 and 80%, the argon atmosphere has the highest wear rate, followed by nitrogen, air, and carbon dioxide.

Result of wear rate

The wear rate changing curves in argon and carbon dioxide differ slightly from those in air and nitrogen. The maximum wear rate in an argon environment occurs on specimens with a MVF of 80%. In the air and nitrogen environments, the maximum wear rate occurs on specimens with a MVF of 93%. For the carbon dioxide environment, the wear rate of the dual-phase steel is less affected by the variation of the MVF. In comparison with the other gas environments, the lowest wear rates were obtained in the CO2 environment, regardless of the MVF.

Wear mechanisms

Microanalysis

SEM images of wear tracks in air, argon, nitrogen, and carbon dioxide environments are shown in Figure 4.9. The material's performance is relatively close in the experimental environment of argon and nitrogen. In comparison to other gas environments, nitrogen and argon exhibit more adhesive wear, with visible adhesive scuffing on the track surface. In contrast, in a carbon dioxide environment, adhesive wear and material removal are significantly reduced, and abrasive wear is very low.

Furthermore, when compared to the wear debris from the argon and nitrogen atmospheres, the wear debris from air and carbon dioxide appears to be well-adhered to the wear track, even forming a layer of a compressed thin film that smeared across the wear track. This phenomenon is more visible in the presence of carbon dioxide. The oxygen concentration of the wear track surface was determined using EDX analysis in the air, argon, nitrogen, and carbon dioxide atmospheres. It is obvious that the oxygen concentration in air and carbon dioxide gas environments (18.49% and 21.79%, respectively) is much higher than in argon and nitrogen gas environments (5.08% and 7.49 %) from Table 4.1. This indicates that the surface of the dual phase steel produces more oxidative wear in the air and carbon dioxide environments than in nitrogen and argon.

It is known that, when compared to non-oxidized metal surface areas, iron oxides (such as Fe2O3 and Fe3O4) formed in the air and carbon dioxide environments [208,209,232] reduce adhesion. This explains why adhesive wear is less in the air and carbon dioxide environment than in argon and nitrogen. 4.2 show the results of EDX analysis on the wear track surface points. The oxygen content of the wear debris can clearly be seen to be extremely high, indicating that significant oxidative wear occurred during friction.

Furthermore, the dark or black areas of the wear track have higher oxygen concentrations, whereas the light gray areas have the same oxygen concentration (3%) as non-wear areas outside the wear track. The SEM pictures in Figure 4.11 and the results of EDX in Table 4.3 show that the oxidative wear increases with the decreasing MVF under the same gaseous atmosphere.

The morphology of worn surfaces revealed smooth undelaminated regions, indicating oxidative wear. As a result, the previously mentioned compression film formed by wear debris can be categorize as an oxide film. The oxide film formed on the wear track can reduce wear, particularly adhesive wear, as observed in this experiment.

Calculation of plastic deformation

The plastic deformation can be calculated by the method described in Section 3.3.2, which was shown as Figure 3.8.

Result of plastic deformation

Figure 4.12 depicts the plastic deformation as function of MVF under different gaseous atmospheres. It can be observed that the trend of plastic deformation in carbon dioxide is similar to that in air, with the maximum plastic deformation occurring in the specimen with a MVF of 59%. Meanwhile, this trend is opposite to the wear rate curve in air and carbon dioxide rubbing environments. In addition, the maximum volume of plastic deformation is found in the carbon dioxide rubbing environment, regardless of the MVF. It seems to provide one explanation why the material exhibits an extremely low wear rate in the carbon dioxide atmosphere versus the other gaseous environments.

Apart from that, the variation pattern of plastic deformation in argon and nitrogen atmosphere is not quite noticeable. Comparing the different gaseous environments, it can be found that on the specimens with higher MVF, the size of the debris generated in the air environment is much larger than that in the carbon dioxide environment. As the MVF decreases, the size of the debris formed in air becomes progressively smaller and there is less difference with the size of the debris generated in the CO2 environment. This corresponds well to the wear rate curve. In the air environment, the wear rate decreases as the MVF decreases. In the CO2 environment, the wear rate remains at a lower value, regardless of the MVF. For these two gaseous environments, when the difference in the size of the generated debris is larger, i.e., on specimens with a higher MVF, a large difference in the wear rate can also be found; when the difference in debris size is smaller, i.e., on specimens with a lower MVF, the wear rate can also be found to be approaching gradually in both environments. Therefore, the variation of the wear debris size reflects the evolution of the wear rate considerably. In the nitrogen and argon environment, it was not available for analysis and comparison due to the limited amount of wear debris produced.

Discussion of the combined effects of gaseous atmosphere and microstructure on friction coefficient and wear rate

Combining SEM images and EDX analysis results, the COF, wear rate and wear mechanism are analyzed and discussed as the MVF varies in different gas environments.

The main reason for the lower coefficient of friction in air compared to argon and nitrogen environments is a change in the wear mechanism. Metal oxides, as we all know, have relatively favorable friction properties [215,216], and direct contact between metallic materials causes severe adhesion [209,232]. Because the argon and nitrogen atmospheres lack oxygen and other reactive environmental molecules, metal-to-metal contact has a higher adhesion, which dominates the tribological behavior of the contacting surface. Previous research has suggested that the increase in COF could be explained by predominant adhesive wear [START_REF] Trevisiol | Effect of martensite volume fraction and abrasive particles size on friction and wear behaviour of a low alloy steel[END_REF]. Furthermore, the delamination of the compacted transfer layer and the formation of oxide during the friction test can be used as a foundation. When compared to adhesive wear in argon and nitrogen environments, the compacted oxide layer formed in the air atmosphere can produce a lower real contact area and fewer junctions, potentially lowering the coefficient of friction. The EDX analysis results show that oxidative wear increases with decreasing MVF, which explains why the COF decreases in the air environment with decreasing MVF.

Although oxidation wear increases with decreasing MVF in argon and nitrogen environments, there is significant adhesive wear, which affects the COF. In the case of a carbon dioxide environment, it should be noted that the increase in friction in the CO2 atmosphere was usually related to the thick wear debris tribofilm formed during contact.

Igor Velkavrh[102] obtained similar results, which are explained by the negative effect on COF caused by abrasion debris and tribofilm accumulation during sliding contact.

Analyzing the results of plastic deformation yields another explanation (Figure 4.12).

The material exhibited significant plastic deformation in the presence of carbon dioxide gas. The real contact area increases as a result of the massive plastic deformation, resulting in an increased COF [233,234]. The phenomenon that the COF increases with decreasing MVF in a carbon dioxide environment is caused more by the negative influence of wear debris accumulation, oxidative wear, and plastic deformation.

Because of the relatively high oxygen concentration, the main type of wear in air and carbon dioxide environments is oxidative wear. In contrast, the main wear mechanism in the nitrogen and argon environments is adhesive wear. The presence of a composite oxide film acts as a protective layer on the metal surface [141,166,218,219,235], which explains why wear is lower in air and carbon dioxide environments than in argon and nitrogen environments. The metal oxide glaze layer forms more quickly and easily in the carbon dioxide environment because the oxygen content exceeds that of the air.

Further explanation and analysis of the effect of MVF on wear in the same gas environment can be made using plastic deformation and wear debris size. When compared to other regions, the temperature at plastically-deformed contacting areas is higher. As a result, these plastically deformed areas are preferential oxidation sites [START_REF] Abouei | Effect of microstructure on the oxidative wear behavior of plain carbon steel[END_REF]166]. Cracks would form and spread as the substrate's plastic deformation increased [217,220,221]. This allows oxygen to enter the cracks and react with the steel substrate to form oxides inside the matrix. Under the same normal load and sliding speed, it is obvious that a large plastic deformation would occur in a softer matrix, producing more tribo-oxides. Because of the low normal load and sliding velocity, the plastic deformation produced during wear aids in the retention of more tribo-oxides. The oxide remaining in the wear track is compacted into a hard tribo-oxide layer as the low-speed friction continues, preventing metal-metal contact and thus reducing wear. Because the hardness of the dual-phase steel matrix decreases as the MVF decreases, the plastic deformation increases as the MVF decreases. Meanwhile, because the oxide increases with plastic deformation, oxidative wear increases as the MVF decreases.

The evolution of debris size as a function of MVF in CO2 and air environments is depicted in Figure 4.13. It is possible to conclude that the size of the debris decreases as the MVF increases. This is more evident in the air environment than in the CO2 environment. It is well understood that oxidation prevents rubbing surfaces from mutual material transfer and the subsequent pile up to form larger particles, resulting in finer wear particles. The experimental results show that the wear rate decreases as the wear particles are refined. Meanwhile, the variation of the wear debris size corresponds to the experimental results of wear rate. In air, larger debris sizes correlate to higher wear rates. As the size of the debris generated in the air approximates that in CO2, the wear rate also gradually approaches. The similar and small size of the debris generated in the CO2 environment for the different MVF specimens corresponds well to that the low and stable wear rates in CO2. Thus, the wear debris size can be considered as a direct response to the changing wear rate.

The increased oxidation refines the wear particles, which accounts for the reduction in wear. Oxide and oxidized metal debris can be retained or compacted on the contacting surfaces under low speed and reciprocating sliding conditions.

Furthermore, the oxidized wear-particles that remained between contacting surfaces acted as lubricants, reducing wear. The smooth oxide 'glaze' formed by compacted particles can provide even better anti-wear protection. As a result, as the MVF decreases, so does the size of the debris. This adequately explains how a low MVF contributes to the acceleration of oxidation at a rubbing surface, resulting in a decrease in wear rate.

Conclusion

The main purpose of this chapter is to determine the combined effects of gaseous environment and microstructure on the tribological behavior of dual-phase steels. Dry sliding experiments on specimens with different microstructures were performed on a ball-disc tribometer under different gaseous environments (air, nitrogen, argon, and carbon dioxide) with constant normal load(15N) and sliding speed. From the obtained experimental results, the following conclusions can be found.

For the COF, it can be analyzed in two parts: the average COF at steady state and the initial-to-steady state COF. For the average COF in steady state, the minimum COF appears in the air environment regardless of the MVF. Moreover, the average COF decreases with decreasing MVF. The maximum COF appears in the carbon dioxide environment and its value increases with decreasing MVF. The average COF in argon and nitrogen environments are in between that in air and carbon dioxide. The COF in argon is slightly greater than that in nitrogen. However, the variation of the COF in argon and nitrogen environments is independent of the MVF. It is obvious that there is different reaction occurring on the rubbing surface for each gaseous environment. It is worth explaining that in the CO2 gaseous environment, unlike the steady state where the maximum average COF appears, the longest transition time of the COF also occurs at this environment. It means that CO2 is indeed effective in reducing the COF during the initial rubbing period. There is no relevant explanation in this work for the highest average COF occurred in CO2 after reaching steady state.

The wear in the CO2 environment is lower compared to other gaseous environments. Moreover, the microstructure (MVF) seems not to be a major factor affecting the wear rate in the carbon dioxide gaseous environment. However, in air, nitrogen, and argon environments, relatively lower wear rates can still be found for specimens with lower MVF.

Microscopic observation and the elemental composition analysis show that the lowest wear rate in CO2 environment is due to the generation of a relatively complete smooth compacted oxide layer on the wear tracks. This can be observed clearly in the micrographs. Meanwhile, the wear tracks in the CO2 environment exhibit the highest oxygen content, which also indicates that the oxidation is more easily produced in this environment. Generally, the compacted smooth oxide layer effectively protects the contact surfaces from metal-metal contact, which reduces adhesive wear and significantly decreases the wear rate. This could explain the extremely low wear rate in a CO2 environment. Besides, the results of plastic deformation also suggest that the largest plastic deformation occurs in carbon dioxide, which providing another evidence for the occurrence of larger oxidation in CO2.

Oxidative wear is the dominant wear mechanism in the CO2 environment, regardless of the MVF. For air environment, oxidative wear dominates for specimens with lower MVF (≤59%), while ploughing and adhesive wear predominate for higher MVF. In the nitrogen and argon environments, the wear mechanisms are similar, both dominated by adhesive and ploughing at higher MVF, while lower MVF (≤59%) is characterized by adhesive wear and accompanied by severe material spalling. Finally, the wear mechanisms of different oxide layers were analyzed and the effects of microstructure and oxidation temperature on the friction and wear of the oxide layers were discussed.

Generation and characterization of oxide scales

Preparation of oxide scales

Samples with different MVF were obtained using the methods in sections 2.2 and 2.3. Prior to the sample oxidation, the prepared disc-shaped sample was cut into four equal-sized parts using a Struers Accutom-50 cutter. This ensures the same MVF at different oxidation temperatures, which facilitates independent studies about the effect of oxidation temperature on the behavior of oxide layer.

The various oxide scales were obtained by the static oxidation to keep the cut samples in an experimental oven (air atmosphere) at 200 °C, 400 °C, and 600 °C for 2 hours and then cool to room temperature in air.

Characterization method of oxide layer

To evaluate the thickness of the oxide layer, the oxidized samples were resinmounted and grinded on 1000-4000 grit SiC paper. The cross-sections were examined by SEM to determine the potential oxidation mechanism and scale thickness. For statistical reasons, the thickness was measured at five points across the cross-section.

The analysis of the nature of the various oxide scales is achieved by the physical From literatures [134,236,237,240,241], it can be seen that the oxides formed on low-alloy steel at 200°C is mainly Fe2O3, at 400°C is Fe3O4, and when the temperature is above 570°C, FeO is the dominant oxide component. Considering the obvious color characteristics of the oxide layer at 200°C, it can be determined that the main oxide at 200°C in this experiment is Fe2O3.

Because the oxide layers generated at 400°C and 600°C have quite similar colors, it is difficult to determine the main constituents of the oxide layers simply from their physical characteristics. The main investigation in this paper is the tribological behavior of oxide layers composed of different types of oxides. Thus, it becomes particularly important to identify the exact composition of the oxide layers. The relevant literatures [134,172,176,237,240] have provided references about the main compositions of the oxide layers formed in the low-alloy steel at 400 °C and 600 °C as Fe3O4 and FeO, respectively. The EBSD physical phase identification was performed to determine the main compositions formed at the current conditions more precisely. In summary, from the physical characteristics, the EBSD results, and the references, we can determine that Fe2O3 is the main component of the oxide layer generated at 200°C, at 400°C is Fe3O4, and at 600°C is FeO. Since the minimum thickness observed in this experiment is around 1 μm, the hardness value at the 300 nm was chosen so that minimize the impact of the substrate on the hardness of the oxide layer and to avoid excessive mechanical errors due to the small indentation depth selected. In fact, at this depth, the imprinted area is of the order of 1 μm 2 and should be rarely influenced by the substrate. Figure 5.6(b) shows an example of the imprint after the maximum load, and also indicates the size of the indentation at an indentation depth of 300 nm. All hardness profiles for the MVF 100% sample are shown in Figure 5.6(c), where the thick black curve is the mean of all hardness. The hardness of the oxide layer is the averaged over all relevant curves at a depth of 300 nm. The measurements for the hardness of the substrate after oxidative process were carried out by a Vickers Zwick ZHV-µs(H115) at a load of 10g. All surfaces of the sample are covered with oxide layer after the oxidation process in the oven. The cutting cross-section of the sample is selected as the region for the oxidized substrate hardness measurement. The specimen is mounted and the cross section is polished. This removes the oxide layer adhering to the substrate on the cross section. Thus, the hardness of the oxidized substrate can be measured directly. The values reported for hardness were an average for five measurements from different locations. A measurement of the hardness of the substrate after oxidation is made to contrast with the hardness of the oxide layer.

Result of oxide layer microhardness

To characterize the mechanical properties the oxide layers formed by the different oxides, nanoindentation has been realized on each oxide layer and microhardness measurements have been performed on the corresponding oxide substrates. show an increasing tendency of microhardness with the increasing MVF. The magnitude of increase with temperature for the oxide layer is slightly larger than that for the oxidized substrate. Besides, the microhardness of the oxide layer is basically greater than that of the corresponding oxidized substrate, especially at 200 °C and 400 °C, where there is a significant difference in microhardness. That is in accordance with the results reported in other studies [134,194]. Contrary to the maximum microhardness of the oxide layer achieved at 400°C, the microhardness of the oxide substrate gradually decreases with the increasing oxidation temperature. It indicates that there are different mechanical property changes in the oxide scale and the oxide substrate as the oxidation temperature increasing. This also demonstrates the necessity of separating the oxide layer from the substrate and investigating their respective properties individually.

Acquisition of friction coefficient

Since the Nanoindenter Agilent G200 is equipped with a device to measure the horizontal friction force Ft in the scratch direction, the corresponding instantaneous COF (Ft=µFn) can be calculated from the measured tangential and normal forces during the scratch experiments. The average value of the COF was calculated from the results of three experiments with 500 instantaneous COF values each time. Combining the curves of the COF during scratch test and the mean value of COF, it can be found that, no matter the MVF and the normal load, the maximum COF occurs at the oxide scale produced at 400°C, which corresponds to the main oxide as Fe3O4.

Result of friction coefficient

Next, the COF of the oxide layer increases with normal load regardless of the microstructure of matrix and oxidation temperature. Besides, the MVF is insensitive to the COF of the oxide layer and there is no obvious regularity to be found. The large fluctuations in COF shown may be due to changes in penetration depth and fracture patterns occurring in the contact area.

Wear rate

Calculation of wear rate

The wear rate for the scratch test is calculated by the cross-section method. As shown in Figure 5.13, As is the area of the scratch section, A1 and A2 are the areas of the piled-up material on both sides of the groove, l is the width of the scratch, d is the depth of the scratch, and α is the attack angle of the indenter tip.

Here, it can be found:

𝑡𝑎𝑛𝛼 = 𝑑 1 2 • 𝑙 (5.1) 𝑙 = 2𝑑 𝑡𝑎𝑛𝛼 (5.2)
The wear area of the profile Ar:

𝐴 𝑟 = 𝑓 𝑎𝑏 × 𝐴 𝑠 (5.3)
Where, 𝑓 𝑎𝑏 = 𝐴 𝑠 -(𝐴 1 + 𝐴 2 ) 𝐴 𝑠 (5.4)

𝐴 𝑠 = 1 2 • 𝑙 • 𝑑 = 𝑑 2 𝑡𝑎𝑛𝛼
(5.5) Based on Zum Gahr[243], the fab value is a means of quantifying the number of microchipping (fab=1) and microplowing (fab=0). However, in this scale, an accurate measurement of A1 and A2 could not be performed, we assumed that fab=1. Therefore, the Ar can be written:

𝐴 𝑟 = 𝐴 𝑠 = 𝑑 2 𝑡𝑎𝑛𝛼 (5.6)
The volume loss:

𝑉 = 𝐴 𝑟 × 𝐿 = 𝐿 • 𝑑 2 𝑡𝑎𝑛𝛼 (5.7)
Therefore, the wear rate can be calculated as follow:

𝑊 𝑟 = 𝐴 𝑟 𝐿 = 𝑑 2 𝑡𝑎𝑛𝛼 (5.8)
Since all the tests were carried out with a same indenter tip, tanα is constant. The relationship between Wr and d can be represented as:

𝑊 𝑟 ∝ 𝑑 (5.9)

Therefore, the variation of scratch depth can be used to express the trend of wear rate in this experiment. In other words, the deeper the scratch, the higher the wear rate. From the above results of scratch depth, it is observed that regardless of the load and MVF, the maximum scratch depth happens on the oxide layer generated at 600°C, correspondingly, the main oxide is FeO. Further, the scratch depth increases with normal load for any condition, which is easy to understand. Finally, for the microstructure of the substrate, there is an overall decreasing tendency of the scratch depth as the increasing MVF. Owing to the large surface roughness of the oxide layer, it leads to a greater impact of the system error at low load(20mN). Therefore, this changing is more pronounced at relatively high loads (50mN and 100mN).

Result of wear rate

Figure 5.17 shows the results of scratch width of oxide scale on martensite and ferrite under different conditions. It is clear the scratch width of oxide scale formed on the martensite is significantly less than that on the ferrite. Since this experiment was performed with the same indenter tip, the specific wear resistance of the scratch test was negatively related to the scratch depth. Meanwhile, due to the scratch width is positively to the scratch depth, the scratch width is likewise inversely related to the specific wear resistance. That means the oxide layer formed on the martensite has a greater wear resistance than that on the ferrite. experiments shows that there is no significant delamination of the oxide layer throughout the scratch grooves, which suggests the quality of bonding to the substrate.

Wear mechanisms

In summary, it is concluded that when the main component of the oxide layer is Fe3O4, the dominant wear mechanism is micro-cutting; Ploughing is the main wear mechanism that occurs when the main component of the oxide layer is FeO. Whereas, in the case of Fe2O3, the wear mechanism is intermediate between micro-cutting and ploughing.

Discussion of friction coefficient and wear resistance of oxides scales

As is known, the tribological behaviors of oxide layers depend on many parameters such as thickness, chemical composition, natures, and adherence oxide /substrate. In some cases [134,185,[244][245][246], a three-layer oxide structure is obtained, while in others [134,210] a two-layer oxide structure is obtained, as previously described in the literature. For low-alloyed steels [134,192,210] it is common to obtain a single oxide layer at relatively low temperatures, which corresponds to the oxide layer obtained in this experiment. Therefore, more attention has been concentrated on the impact of nature and thickness on tribological properties of oxide scale with single layer.

Besides friction and wear, the hardness of the oxide layer was identified as an auxiliary mechanical performance for proper correlation.

Although there are extensive previous literatures on the effect of frictional oxides and thickness of oxide layer on tribological behavior by using pin-on-disk experiments, friction and wear in the scratch test are completely different from that in the pin-ondisk test. For the scratch test, only one abrasive grain (the indenter) makes contact with the materials, whereas for the pin-on-disk experiment, many abrasive grains are in contact. This is largely due to the different contact geometries between the scratch test and the pin-on-disk test. Moreover, compared to the scratch test, the pin-on-disk experiment is more like a combination of multiple scratch tests. Hence, for the scratch tests, the measured COF and wear rate mainly provide the basis for the abrasive wear of the oxide layer [190]. The existing studies on the effect of oxide layers or oxidation products on the tribological behavior are more concerned with oxidative wear but not abrasive wear. This is what makes this experiment clearly different from other investigations [172,191,194].

For the COF, a fact seems to stand: when the thickness of the oxide scale decreases, the COF increases. However, the explanation for this is still controversial [247]. The results of the present experiments under different thickness of oxide layers but the same oxide composition show that the effect of thickness on the COF is not regular.

Compared to the impact of thickness, the variations in hardness due to different oxide nature seem to have a greater effect on the COF [194]. Generally, it is accepted that higher hardness correlates with lower COF and the overall trend is valid. However, the exception is the COF in the scratch test [190], where the average COF of the oxide layer increases with hardness. In this experiment, it appears that the hardness of the oxide layer plays a considerable part. Since the Fe3O4-dominated oxide layer formed at 400°C has the maximum hardness, which leads to the wear mechanism transition to microcutting, resulting in the maximum COF occurring in this case.

For the wear rate, it is generally assumed that the oxide composition is the critical factor affecting oxidative wear [210]. As reported, the change in wear rate corresponds to the variation in oxide composition [172], which leads to a reducing than two orders of magnitude alteration in wear rate [177]. At the same time, the different nature of the oxides impact the hardness and structure of the oxide scale, which consequently impacts the variation of the wear rate in the scratch test. It is common that the wear rate increases with decreasing hardness, i.e., the higher hardness signifies better wear resistance [190].

However, there are exceptions. Khruschov[248] showed that the wear resistance decreases with increasing hardness. The same author developed a relation among wear resistance and Young's modulus to indicate that wear resistance increases with the increasing Young's modulus of the material [249].The results of other investigators have also revealed that the dependence of wear resistance on overall hardness is much weaker [250]. Hence, hardness must be cautiously utilized as a parameter to predict wear resistance. For example, the wear mechanism of coatings represents a more complex form owing to the laminar structure. Straffelini et al.[251] identified that the hardness of the reinforcing phase in the material has a stronger impact on wearability than the overall hardness, and that the material with the lowest hardness of the substrate displayed the minimum wear rate. The results of this scratch experiment show that the maximum wear rate appears on the FeO-dominated oxide layer formed at 600°C, and the lowest wear rate occurs on the Fe2O3-dominated oxide layer formed at 200°C.Indeed, it seems that the hardness of oxide layer is not a decisive factor for the wear rate in the scratch tests, despite the fact that the Fe3O4-dominated oxide layer formed at 400°C has the highest hardness,it does not show the lowest wear rate at the mentioned loads (Figure 5.16). Instead, a different tendency was observed between the COF and the hardness of the oxide layer, i.e., the oxide layer with highest hardness showed the largest COF at the loads used [START_REF] Becker | Dualphasen-Stähle mit erhöhter Festigkeit und Verformbarkeit[END_REF][START_REF] Kerridge | The stages in a process of severe metallic wear[END_REF], and 100 mN) (Figure 5.12). There is no brief explanation for this tribological behavior of the Fe3O4-dominated oxide layer formed at 400°C. The opposing behavior of hard coatings was also noted by Gee[252].

He demonstrated that there was almost no correlation between wear resistance and specimen hardness in scratch tests at low loads (less than 4 N). It may be owing to that the oxide layer has a weaker sublayer cohesion strength and is more ductile, but it is also likely to be associated with the microstructure of the oxide layer. For example, Yu et al. suggested that the friction and wear properties is inconsistent for oxide layer formed with different microstructures, even the same composition of Fe3O4 [191]. In addition, the analysis of the wear mechanism suggests that the main wear mechanism of the Fe2O3-based oxide scale is in a transition between ploughing and micro-cutting, while the Fe3O4-dominated oxide layer is micro-cutting, thus leading to a greater wear rate. In summary, it is evident that the variation of wear rate is more attributed to the different oxide nature rather than the hardness.

For the wear rate in the scratch test, in addition to the oxide composition, the thickness of the oxide scale also has a major impact on the wear rate. Consequentially, it is found that the wear resistance decreases with increasing thickness of the oxide layer in this experiment. This is confirmed by the experimental results obtained under different MVF but same oxidation temperature (same major oxide composition), i.e., as the MVF decreases, the oxide layer thickness increases, leading to a higher wear rate.

It is evident that the thickness of the oxide layer seems to have a greater impact on the wear resistance compared to the nature of oxide. Meanwhile, our previous researches [253,254] regarding oxidative wear of dual-phase steels with various MVF proved that oxidative wear increases with decreasing MVF, which is a good proof of the impact of MVF on the oxidizability of the material. It is well verified by the results of thickness of oxide layer at different MVF, i.e., samples with low MVF have greater oxidizability.

Thus, the microstructure before oxidation has an effect on the thickness of the oxide scale, which in turn impacts the tribological behavior of the oxide layer. It is more general for the impact of load on the friction and wear of the oxide layer, where both the COF and the wear rate increase with the increasing load in the scratch experiment.

Currently, most of the studies on oxide layers are performed on materials with relatively simple and homogeneous microstructures [177,210]. Tyagi [START_REF] Tyagi | Modelling of dry sliding oxidation-modified wear in two phase materials[END_REF] Furthermore, the high difference in hardness between the substrate and the oxide layer also leads to higher wear [256].This also explains well the poorer wear resistance of the oxide layer formed on the substrate with lower MVF under the same oxidation temperature and load.

Conclusion

The investigation of this chapter focused on the tribological behavior of the oxide layers obtained for low-alloy steels (dual-phase steels) at different oxidation Comparing the hardness of the oxide layer and the corresponding oxidative substrate, it can be observed that the oxide layer generally has a higher hardness than the oxidative substrate.

The tribological behavior of the oxide layer was tested by unidirectional scratching experiments on a nanoindentation device. The tribological properties of the oxide layers were analyzed in terms of friction coefficient, wear rate, and wear mechanism, respectively.

For COF, the maximum COF occurs on the Fe3O4-dominated oxide layer generated at 400°C. The difference in composition of the oxide layer caused by the oxidation temperature is the main factor affecting the COF, while the MVF of the substrate and the thickness of the oxide layer are not sensitive to the change in the COF of the oxide layer. For normal load, the COF of all oxide layers increases with increasing normal load. It is due to the increase in the contact area between the nanoindenter and the oxide layer caused by the increase in normal load.

For the wear rate, both the composition and thickness of the oxide layer are the main factors affecting the wear rate. The maximum wear rate occurs on the FeO. The wear rate of the oxide layer decreases with increasing MVF and decreasing oxide layer thickness. In contrast to the nature and thickness of the oxide layer, hardness is not a critical factor in determining the wear rate for scratch testing. In addition, the wear rate of all the different oxide layers increased with increasing normal load.

For the wear mechanism, microscopic analysis of the scratch tests showed that the main wear mechanisms are in the transition between ploughing and micro-cutting on the Fe2O3-dominated oxide layer produced at 200°C, and ploughing on the FeOdominated oxide layer produced at 600°C. Micro-cutting mainly occurs on the Fe3O4dominated oxide layer formed at 400 °C. Besides, under the same conditions, the oxidation properties of different microstructures on dual-phase steel are different.

Compared to ferrite, the oxide layer formed on martensite has better wear resistance.

6 Conclusions and perspectives

Conclusions

This research focused on the effects of microstructure, gaseous atmosphere, and sliding velocity on the tribological behavior of dual-phase steels. Furthermore, the tribological properties of the oxide layers formed at different oxidation temperatures and MVF were investigated.

The first chapter is devoted to the general introduction of the research background.

The development of dual-phase steels, the impact of microstructure on tribological behavior, and oxidative wear are introduced, respectively. Later, the research objectives of this thesis are explained. The macro-hardness increases with the increase of MVF. The microhardness of martensite decreases with increasing MVF, while the variation of microhardness for ferrite is not significant. Therefore, for the same steel, the increase in macro-hardness of the dual-phase microstructure is more related to the proportion of phases than to the local hardness of the phases. In addition, for a single martensite island, the microhardness in the center is usually greater than that at the edges. Compared to ferrite, the oxide layer formed on martensite has a better wear resistance.

Perspectives

The investigation of oxidative wear of dual-phase steel might also be explored more extensively by combining the microstructure with other tribological parameters, which allows the dual-phase steel in an optimal performance under different service conditions. This will be the main research direction for further work, which will contribute to a better and more comprehensive understanding of the frictional oxidation.

➢ High-temperature sliding experiments will be performed to investigate the combined effects of microstructure and temperature on oxidative wear of dualphase steel. The corresponding results will be compared with the tribological properties of oxide layers obtained by static oxidation to analyze the differences between frictional oxides and static oxides, which provide a deeper knowledge of the role of oxides in the sliding friction.

➢ Experiments of oxygen partial pressure will be carried out to analyze the combined effects of oxygen partial pressure and microstructure on the oxidative wear of dual-phase steels. Thus, it is possible to determine at what oxygen content the dual-phase steel offers the greatest wear resistance, thereby providing an optimized solution for practical engineering applications.

➢ In this paper, the temperature of the contact surface was not measured while investigating the combined effects of sliding velocity and microstructure on oxidative wear of dual-phase steel. It is necessary to improve this experiment and build the relationship between sliding velocity, microstructure, and contact surface temperature to better understand the essence of oxidative wear caused by sliding velocity.

➢ Simulations and theoretical modeling are necessary. Based on the relevant data obtained from this experiment, simulation and theoretical modeling should be created, which enable preliminary prediction of the tribological behavior of dual-phase steel.

Figure 1 .

 1 Figure 1.1: Independent marketing research of material growth in kg for future automotive applications ....................................................................................................................... 6 Figure 1.2: Annealing cycle for dual-phase steel consisting of ferrite and martensite .............. 7 Figure 1.3: Oxidative mechanism of steel ............................................................................... 26 Figure 1.4: Diagram of Quinn's theory of oxidative wear: (1) Fracture of oxide layer;(2) Formation of thick oxide layer;(3) The initial generation of thick oxide layer at contact asperities;(4) Without contact ......................................................................................... 31 Figure 1.5: Schematic diagram of oxidative wear by Wilson ................................................. 32

1 : 6 Figure 1 . 2 : 7 Figure 1 . 3 : 26 Figure 1 . 4 : 31 Figure 1 . 5 : 32 Figure 2 . 1 : 44 Figure 2 . 2 : 45 Figure 2 . 3 : 49 Figure 2 . 8 : 50 Figure 2 . 9 : 53 Figure 3 . 1 : 58 Figure 3 . 5 : 61 Figure 3 . 8 : 63 Figure 3 . 9 : 68 XFigure 3 . 12 : 77 Figure 4 . 1 : 86 Figure 4 . 3 : 86 Figure 4 . 4 : 87 Figure 4 . 5 : 97 Figure 5 . 1 :XIFigure 5 . 2 :

 161271326143115322144224523492850295331583561386339683127741864386448745975152 Figure 1.1: Independent marketing research of material growth in kg for future automotive applications ....................................................................................................................... 6 Figure 1.2: Annealing cycle for dual-phase steel consisting of ferrite and martensite .............. 7 Figure 1.3: Oxidative mechanism of steel ............................................................................... 26 Figure 1.4: Diagram of Quinn's theory of oxidative wear: (1) Fracture of oxide layer;(2) Formation of thick oxide layer;(3) The initial generation of thick oxide layer at contact asperities;(4) Without contact ......................................................................................... 31 Figure 1.5: Schematic diagram of oxidative wear by Wilson ................................................. 32

  process. The varied tribological systems present complex tribological laws. For the previous investigation in the Roberval laboratory at Université de Technologie de Compiègne(UTC), Trevisiol and Penfornis concentrated more on the effect of microstructure on abrasive wear of dual-phase steel under dry friction. Based on their work, this thesis focuses on the combined effects of microstructure, sliding velocity, and gaseous atmosphere on oxidative wear of dual-phase steel at lower surface roughness. The correlation between microstructure and oxidative wear was developed by performing sliding tribological tests and analyzing macro-hardness, microhardness, plastic deformation, and wear debris size. In addition, the tribological behavior of the oxide layer generated at different microstructures and oxidation temperatures was also investigated. Therefore, the research work is intended to enrich the tribological models and to contribute for industrial applications of dual-phase steel, especially for oxidative wear.

Figure 1 . 2 .

 12 Figure 1.2. As the temperature reaches Ms during cooling process, austenite transforms to martensite, while the ferrite phase remains almost unchanged. The amount of martensite in the dual-phase microstructure can be adjusted by modifying A1.

Figure 1 . 2 :

 12 Figure 1.2: Annealing cycle for dual-phase steel consisting of ferrite and martensite[START_REF] Bleck | Microalloying of cold-formable multi phase steel grades[END_REF]

  [117] performed dry sliding experiments with AISI H13 steel of different hardness at different velocities in air. Due to the evolution of the frictional oxide layer, the sliding speed significantly affects the wear behavior and mechanism. As the sliding speed increases, more oxides are produced during the frictional process. The stability of the oxide layer becomes a critical issue in determining the wear rate. In addition, three wear mechanisms were identified as dominant. At relatively low sliding speeds, adhesive wear and abrasive wear dominated, accompanied by mild oxidative wear. As the sliding speed increases, mild oxidative wear dominates. With a further increase in sliding velocity, wear rate decreases in the transition from mild oxidation to severe wear.

  Gåård et al.[127] recently evaluated the dry sliding wear behavior of hardened AISI D2 tool steel against martensitic carbon steel. The experiments were run on a sliding plane tribometer with sliding velocities between 281-409 mm/s and contact pressure of 2000 MPa. The results suggested that the adhesive wear increased as the sliding velocity increased. Gåård explained it as an increase in surface temperature due to the friction heating, resulting in an increased adhesion between surfaces [128]. A similar trend[123] was also observed in metal-metal contact with sliding speeds in the range of 200-8000 mm/s. However, a research[129] showed that adhesive wear decreases as the temperature increased from 20°C to 150°C for unlubricated steel-tool steel contact pairs with pin-on-disc sliding conditions under low speeds. When sliding speed and overall temperature were increased separately, it showed different wear behavior. The literature[130, 131] also demonstrates that the magnitude of adhesive wear reduces with an increase in overall temperature. These are contrary to the suggestion that an increase in temperature at the contact surface (e.g. by frictional heating) results in an increasing adhesive wear.For dual-phase steels, M. Ruiz-Andres[107] investigated friction and wear of DP600 at sliding speed (0.05-3.75 m/s) and contact frequency (0.6-16.5 Hz) under unidirectional ball-on-disc dry sliding conditions. The wear mechanism of DP600 is mainly oxidative wear. The nature of oxide is determined by the sliding velocity and the contact frequency. With the increasing sliding velocity, the COF first increases and then gradually decreases.

Figure 1 . 3 :

 13 Figure 1.3: Oxidative mechanism of steel[132]

[

  138] on the oxidation behavior in dry sliding wear of Fe-12Cr steel. Therefore, Hong et al.[139] and Sullivian et al.[140] proposed another linear theory, which believed that the oxidation can be accelerated by the flaking of the oxide layer during sliding, and

  Fontalvo[146] et al. studied the effect of elements Si and Al on the wear resistance of X38CrMoV5-1. The oxides formed by elements Si and Al were segregated on the worn surface, which slowed down the oxidation of the steel and reduced the thickness of the oxide layer. The thin oxide layer is not sufficient to protect the material and increases the wear. Stott[START_REF] Stott | The influence of oxides on the friction and wear of alloys[END_REF] et al. found that oxide formation is faster in iron-rich alloys than in nickel-rich alloys, and in chromium-poor alloys than in chromium-rich alloys. Radu and showed that the addition of Y as an oxidation-active element to Stellite 6 and 21 alloys significantly improves the mechanical properties and adhesion of oxide layer, resulting in high wear resistance in air.In addition, the growth of oxide layer depends on the contact surface and oxidation conditions, such as epitaxial growth, crystal state of the oxide layer, crystal state of the metal surface, mechanical activation, presence of liquid on the surface, and chemical activity of the oxidizer[150][151][152][153][154][155].

Figure 1 . 4 :

 14 Figure 1.4: Diagram of Quinn's theory of oxidative wear: (1) Fracture of oxide layer;(2) Formation of thick oxide layer;(3) The initial generation of thick oxide layer at contact asperities;(4) Without contact[159]

Figure 1 . 5 :

 15 Figure 1.5: Schematic diagram of oxidative wear byWilson[44]

  [170] studied the dry sliding tribological behavior of AISI 1213, AISI 1020 steels, and AISI Nimonic 75 Nickel Alloys. It was concluded that the change in wear behavior depends on the properties of the substrate and the oxide layer formed at the junction. This research clearly indicates that the substrate and the oxide layer jointly influence the wear behavior.Starffelini[182] studied dry sliding wear of austenitic and austenitic-ferritic stainless steels. It was argued that the plastic deformation of the subsurface layer contributes to the formation of the wear layer. The wear behavior is determined by the combined effects of the oxide and subsurface deformation.Rai et al.[183] andPellizzari et al.[184] investigated the wear behavior of a variety of microstructures of steel at room and high temperatures, respectively. Under air environment, oxidative wear plays an important role at both room temperature and high temperature to reduce the wear rate. At low loads, mild oxidative wear is predominant;

  for a further understanding of high-temperature oxidative wear and the role of oxide nature in the frictional wear process at the macroscopic level[113]. The conventional methods of studying oxide scales by metallographic cross-section preparation normally cause damage to brittle oxide scales. Thus, it is difficult to investigate the tribological properties of the extremely thin oxide layers in detail by the conventional approaches[185]. In recent years, with the advancement of microscopic characterization techniques and mechanical experimental techniques in the micro-and nano-scale, sufficient capabilities have been available to perform tribological experiments in the micro-and nano-scale to independently analyze the tribological behavior of oxide layers. However, existing studies on oxide scales have focused on the effects of oxidation temperature, cooling rate, oxygen partial pressure, and alloy composition on the composition, microstructure, surface morphology, and adhesion of the oxide scales. Limited number of articles have been issued on the tribological properties of oxide scales themselves.Since the thickness of oxide scale is extremely thin and it can be considered as a special coating, the tests of the tribological and mechanical performance of the oxide scale can be explored by attempting nanoindentation and scratch tests[186]. Lately,Zamri et al.[187, 188] successfully acquired the mechanical performance of oxide layers on HSS materials with the assistance of nanoindentation instead of conventional experiments. However, the oxide layers they tested were too thick (more than 20 μm) and only the cross-section of the oxide scales was tested, not on the surface of the oxide layers. As a common non-standard wear resistance test in laboratories, the scratch test has been utilized for years to assess the resistance of scratch damage. It allows relatively simple and fast comparison of abrasive wear of different materials under several loads and provides basic information about the wear mechanism[188, 189], with well reproducible results. Despite the wear rate from the scratch test might not definitely portray that of practical applications due to a variety of intrinsic and extrinsic factors, it offers some fundamental indications for abrasive wear because of its simplicity and idealization. As a semi-quantitative test of the abrasive wear behavior and general mechanical stability, it is of great significance[190].

Figure 2 .

 2 Figure 2.1 shows the CCT and TTT curves for the 25CD4 steel. It concluded that this steel can produce quenched martensitic microstructures after rapid continuous cooling. In addition, the continuous cooling or isothermal holding time required to reach the non-martensitic phase after austenitization is relatively short. This facilitates the obtainment of identical microstructures (dual-phase structures) with various microstructural characteristics (martensite volume fraction). This is an important reason for choosing 25CD4 as the experimental material in this thesis.

Figure 2 . 1 :

 21 Figure 2.1: Diagrams (a) CCT and (b) TTT of 25CD4 low alloy steel[198]

Figure 2 . 2 :

 22 Figure 2.2: Schematic drawing of the step quenched heat treatment for specimens

3 ,

 3 A3 corresponds to the transformation temperature located between the domain composed of γ-austenite only and the domain composed of γ-austenite and α-ferrite[199]. Moreover, this transformation temperature A3 is closely related to the carbon content of the considered material. Generally, for a given alloy, the relevant austenitizing temperature is chosen to be A3 + (25 to 50°C). Meanwhile, the sample must be held at a given austenitization temperature for a sufficient time to allow complete carbide dissolution and austenite grain homogenization. However, the duration of austenitization should not be too long to avoid the generation of excessive coarsened austenite grains, which could worsen the final mechanical properties of the material[200].

Figure 2 . 3 :

 23 Figure 2.3: Phase diagram of iron-carbon alloy[198]

Figure 2 . 4 :

 24 Figure 2.4: TTT diagram of 25CD4 steel depicting the range of isothermal holding durations[198]

Figure 2 . 5 :Figure 2 . 6 :

 2526 Figure 2.5: Transformation of austenitic grains to martensitic grains[201]

Figure 2 . 7 :

 27 Figure 2.7: SEM images of the developed microstructures for an intercritical holding time of: (a) 0 mins; (b) 1 min; (c) 2 mins; (d) 3 mins; (e) 6 mins; (f) 7 mins

Figure 2 . 8 :

 28 Figure 2.8: Evolution of MVF as function of intercritical holding time at 725 °C

Figure 2 . 9 :

 29 Figure 2.9: Evolution of macro-hardness as function of MVF

Figure 2 . 10 :Figure 2 . 11 :

 210211 Figure 2.10: Evolution of microhardness as function of volume fraction of: (a)martensite; (b)ferrite

(

  CSM Tribometer) which shown as Figure 3.1. The 6 mm diameter ball was made of ceramic Al2O3 and was a commercially available standard sample. The average surface roughness of the disc samples is Ra 0.3. The contact pressure is 1.27 Gpa. Under a normal load of 15 N and an average sliding speed of 0.013 m/s, all samples were subjected to reciprocating circular motion in air environment to investigate the effects of microstructure on tribological behavior. This low speed was chosen to avoid high temperatures during the friction test. With a single stroke length of 10 mm, the total sliding distance was 7 m.

Figure 3 . 1 :

 31 Figure 3.1: CSM Tribometer

Figure 3 .

 3 Figure 3.2 shows the variation of the COF with MVF. The average value of the COF at a steady state is acquired and calculated by the experimental data from 100 to 546 s. Figure 3.3 shows the average value of the COF at a steady state as a function of MVF and it remains 0.49~0.60. It is worth noticing that the minimum COF appears on the sample with an MVF of 48%, with a value of 0.49. The maximum COF also appeared on the specimen with an MVF of 100%, and its value reached 0.60. As can be seen in Figures 3.2 and 3.3, the coefficient of friction declines with decreasing MVF.

Figure 3 . 2 :Figure 3 . 3 :

 3233 Figure 3.2: Variation of the COF with time under different MVF

Figure 3 . 4 :

 34 Figure 3.4: The method and steps to calculate the wear rate: (a) and (b) 3D morphology acquisition by Sensofar Metrology; (c) The selection of profiles by SensoMAP application to obtain the area of the peaks and holes

Figure 3 . 5 :

 35 Figure 3.5: The wear rate as a function of MVF

Figure 3 . 6 :

 36 Figure 3.6: SEM images of wear tracks at MVF (a) 100%; (b) 93%; (c) 80%; (d) 68%; (e) 59%; (f) 48%

Figure 3 . 7 :

 37 Figure 3.7: EDX analysis of the points at MVF 68%

Figure 3 . 8 :

 38 Figure 3.8: The plastic deformation measured by SensoMAP: (a) The wear groove;(b) The crosssectional view of the wear groove; (c) Schematic drawing of the cross-sectional area of the wear groove and definition of the plastic deformation

Figure 3 . 9 :

 39 Figure 3.9: The volume of plastic deformation as a function of MVF

Figure 3 . 10 :

 310 Figure 3.10: The variation of wear debris size with martensite volume fraction and the corresponding EDX results: (a) 100%; (b) 93%; (c) 80%; (d) 68%; (e) 59%; (f) 48%

Figure 3 Figure 3 . 12 :Figure 3 .

 33123 Figure 3.11: Schematic diagram of unidirectional rotation mode

Figure 3 . 13 :

 313 Figure 3.13: The average COF as function of sliding velocity under different MVF

  2.1, obtain the value of Shole and Speak (shown as figure 3.14(b)(c)). The equation for calculating the wear rate is as follows:

Figure 3 . 14 :Figure 3 .

 3143 Figure 3.14: The measurement of wear rate: (a) Position selection; (b) Selected location of crosssection profile for wear track; (c) Area calculation for section profile

Figure 3 .

 3 Figure 3.16 exhibits the worn surface of Al2O3 balls at different sliding velocities.

Figure 3 . 16 :Figure 3 . 17 :Figure 3 .

 3163173 Figure 3.16: Variation of worn diameter of Al2O3 ball with sliding velocity

Figure 3 .

 3 Figure 3.18: SEM images of wear tracks under 0.1m/s,0.6m/s, and 1m/s sliding speed at MVF of 93%,68% and 48%

Figure 3 . 20 :

 320 Figure 3.20: Oxygen concentration (in at%.) as function of sliding velocity under different MVF

Figure 3 .

 3 Figure 3.21 shows further micrographs and local enlargements of the worn tracks formed in different MVF and sliding speeds to facilitate more accurate observation and analysis of the wear mechanism.Combined with Figure3.21(a)(d) and the related enlarged views, it can be concluded that at a sliding velocity of 0.1m/s, a large number of deep grooves appear on the worn surface, which is a relatively obvious characteristic of abrasive wear, and some of the worn surfaces show material peeling. The surface flaking exhibits a high oxygen content, as the enlarged view ofFigure 3.21(a). In this case, abrasive wear is predominant and accompanied by less oxidative wear, which results in a larger volume of material removal. As the sliding speed increases to 0.6m/s, a significant amount of grinding snowball-shaped oxidized wear particles appears on the worn surface, and the oxidized region gradually increases. It contrasts with the worn surface generated at a

Figure 3 .

 3 21(c) shows, representatively, that as the number of compacted smooth oxide islands increases, those initially in isolation gradually connect and may eventually cover the entire contact surface. It is quite different from the microscopic morphology of the worn tracks formed at 0.1m/s and 0.6m/s. Meanwhile, the number and depth of grooves on the worn surface are further minimized. At this point, it seems that oxidation wear gradually replaces abrasive wear in a dominant position.

Figure 3 . 21 :

 321 Figure 3.21: Micrographs and local enlargements of the worn tracks formed in different MVF and sliding speeds

4

  Combined effects of gaseous atmosphere and microstructure on oxidative wearThis chapter explores the combined effects of gaseous environment and microstructure on oxidative wear of dual-phase steels. Specimens with different martensite volume fractions were exposed to dry sliding experiments in different gaseous environments (air, argon, nitrogen, and carbon dioxide). The first part (section 4.1) describes the method of dry sliding friction experiment and the variation of the friction coefficient under different test conditions. The second part (section 4.2)

Figure 4 .

 4 1(a)) and an outer chamber(Figure 4.1(b)) exist for the equipment as a secondary sealed chamber. Gases were continuously supplied directly from the gas cylinder into the inner gas chamber. Before beginning the experiment, the required gas must be supplied for at least 20 minutes to fill the inner and outer chambers, preventing other gases from affecting the experimental environment. The tribological experiments used four different gases: air, Ar, N2, and CO2. All non-air gases, namely Ar, N2, and CO2, were commercially available technical gases. The purity of all gases is 99.99999%. The experimental method and parameter settings are described in Section 3.1.1.

Figure 4 . 1 :

 41 Figure 4.1: Tribological test setup in gas environment:(a) CSM Tribometer (with gas cover); (b) Schematic diagram of equipment operation in gas environment

Figure 4 . 2 :

 42 Figure 4.2: COF as function of gaseous environment for samples of (a) 100% martensite; (b) 93% martensite; (c) 80% martensite; (d) 68% martensite; (e) 59% martensite; (f) 48% martensite

Figure 4 . 3 :Figure 4 . 4 :Figure 4 . 5 :Figure 4 . 6 :

 43444546 Figure 4.3: The definition of transition time

Figure 4 .

 4 Figure 4.4 depicts the time required to transition from mild to severe wear as a function of MVF in various gaseous atmospheres.

Figure 4 .

 4 2(b),(d)). Furthermore, the average COF (Figure 4.5) for the first 100 seconds clearly shows the same results as the transition, indicating that the carbon dioxide environment reduced the COF more effectively in the early stages of the experiment. The evolution of the COF as a function of MVF is depicted in Figure 4.6. In the air, the COF decreases as the volume fraction of martensite decreases. Oppositely, the COF increases as the MVF increases in carbon dioxide. Meanwhile, in nitrogen and argon environments, the COF is slightly affected by the volume fraction of martensite, but its average value remains constant, ranging from 0.59 to 0.62, as shown in Figure 4.7.

Figure 4 . 7 :

 47 Figure 4.7: The average COF as function of MVF under different gaseous atmospheres

Figure 4 .

 4 Figure 4.8 depicts the evolution of wear rate as a function of MVF in various gas environments. Obviously, there is extremely low wear in the carbon dioxide environment, which is much smaller than that in other gas environments, regardless of how the volume fraction of martensite changes. The wear rate tends to decrease as the MVF decreases in specimens with a low MVF (80% or less).

Figure 4 . 8 :

 48 Figure 4.8: The wear rate as function of MVF under different gaseous atmospheres

Figure 4 . 9 :

 49 Figure 4.9: SEM images of wear tracks from (a) air, (b) N2, (c) Ar, and (d) CO2 atmosphere at MVF 48%

Figure 4 .

 4 Figure 4.10 and Table 4.2 show the results of EDX analysis on the wear track

Figure 4 .

 4 Figure 4.11: SEM images of wear tracks from air, N2, Ar, and CO2 atmosphere at MVF 93%,68% and 48%

Figure 4 . 12 :

 412 Figure 4.12: The plastic deformation as function of MVF under different gaseous atmospheres

Figure 4 . 13 :

 413 Figure 4.13: The size of wear debris under air and CO2 environments for different MVF samples: (a) MVF93% in air; (b) MVF93% in CO2; (c) MVF68% in air; (d) MVF68% in CO2; (e) MVF48% in air; (f) MVF48% in CO2.

5

  Friction and wear of oxides scale after high-temperature oxidationIn this chapter, the effects of microstructure and oxidation temperature on the tribological behavior of the generated oxide layers are investigated. First, the method of obtaining oxide layers is described. Specimens with different martensite volume fraction (MVF) were kept at different oxidation temperatures to generate oxide layers with different compositions and thicknesses. The hardness of the generated oxide layers was measured by nanoindentation experiments. Afterward, the tribological behaviors of the oxide layers were investigated. The generated oxide layers were exposed to unidirectional scratching tests on a nano-indenter under various normal loads to obtain the friction coefficients and calculate the corresponding specific wear resistance.

Figure 5 . 1 :

 51 Figure 5.1: Oxide layers formed under different experimental conditions

Figure 5 .

 5 Figure 5.2 exhibits the cross-sectional micrographs, Band Contrast, and Phase Color on the oxide layers generated at 400 °C and 600 °C. In Figure 5.2(b), it is observed that the oxide layer consists mainly of equiaxed grains with a grain size of about 250 nm.The result inFigure 5.2(c) shows that the main phase of the oxide at this condition is

Figure 5 . 2 :

 52 Figure 5.2: Micrographs of oxide layers formed at 400°C and 600°C and EBSD phase analysis: (a) Cross-section of the oxide layer formed at 400°C; ( b)Band contrast of oxide formed at 400°C; (c)Phase color of the oxide formed at 400°C; (d) Cross-section of the oxide layer formed at 600°C; (e) Band contrast of oxide formed at 600°C; (f) Phase color of the oxide formed at 600°C

Figure 5 .

 5 [START_REF] Allain | Dual-Phase Steels: The First Family of Advanced High Strength Steels[END_REF] shows micrographs of the cross-sections of the oxide layers obtained at different MVF and temperatures.

Figure 5 . 3 :Figure 5 . 4 :

 5354 Figure 5.3: SEM images of the oxide layer cross-section under different oxidation temperatures and MVF

Figure 5 . 5 :

 55 Figure 5.5: Nanoindenter Agilent G200

Figure 5 . 6 :

 56 Figure 5.6: Nanoindentation micrographs of the oxide layer formed at 400°C on MVF 100% and the selection of oxide scale hardness: (a) Micrograph of the indentations;(b) Enlarged view of indentation and size at 300nm depth; (c) the selection of oxide scale hardness

Figure 5 .

 5 Figure 5.7 shows the variation of the oxide layer microhardness with temperature

Figure 5 . 7 :

 57 Figure 5.7: The microhardness variation of oxide layer with oxidation temperature for different MVF

Figure 5 . 8 :

 58 Figure 5.8: Comparison of microhardness between the oxide layer and the corresponding substrate

Figure 5 . 9 :

 59 Figure 5.9: Micrograph of the scratch experiments under different conditions

Figure 5 . 10 :Figure 5 . 11 :Figure 5 . 12 :Figure 5 .

 5105115125 Figure 5.10: Variation of COF of oxide scales formed at different temperature during the scratch test under different MVF and 50mN

Figure 5 . 13 :

 513 Figure 5.13: The cross-section area of scratch

Figure 5 .

 5 Figure 5.14 displays the variation of the scratch depth of the oxide layer formed at different MVF and oxidation temperatures under a constant 50 mN. The variation of scratch depth with load for oxide scales produced at different oxidation temperatures on specimens with a constant MVF of 80% is shown in Figure 5.15. Besides, Figure 5.16 exhibits the average value of scratch depth for all scratch experiments in relation

Figure 5 . 14 :Figure 5 . 15 :Figure 5 . 16 :

 514515516 Figure 5.14: Variation of scratch depth of oxide layer formed at different temperature during the scratch test at different MVF under 50mN

Figure 5 .Figure 5 .

 55 Figure 5.18: The SEM micrographs of scratch tests under different experimental conditions: (a)20mN; (b)50mN; (c)100mN

Figure 5 . 19 :

 519 Figure 5.19: Microscopic comparison of the bottom of scratches for oxide layers formed at different temperatures on MVF 100% specimens: (a) Fe2O3-dominated formed at 200°C; (b) Fe3O4-dominated formed at 400°C; (c) FeO-dominated formed at 600°C

  proposed a model of oxidation-modified wear resistance in two phases by improving the homogeneous material used by Archard and Quinn for predicting oxidative wear. Three distinct cases are regarded:(a) both phases have the uniform oxidation behavior and critical thickness; (b) there is the same oxidation behavior of the two phases but significantly distinct critical thicknesses; (c) both phases present distinct oxidationbehaviors along with various critical thicknesses. Where, the critical thickness is the maximum thickness to ensure the oxide layer does not peel off or remove from the substrate. However, based on their conclusions, there is no direct evidence available to determine which of the case of dual-phase steel belongs to. The results of this paper show that the oxidation capacity of martensite and ferrite seems to be different under the same conditions, resulting in a lower wear rate of the oxide layer formed on martensite and higher on ferrite.Yanagisawa et al. [255] have done passivation experiments on dual-phase steels and noticed different thickness of oxide scales on ferrite and martensite. The effect of oxide scale thickness on the wear rate in scratch tests has been demonstrated above and this can provide a reason for interpretation.

  temperatures and MVF. The composition of the oxide layer depends only on the oxidation temperature, while the thickness of the oxide layer is determined by both the MVF and the oxidation temperature. The differences in oxide layer composition, thickness, and hardness result in different tribological behavior. In this experiment, the hardness of the oxide layer was examined by nanoindentation. Also, the friction and wear of oxide layers with different natures and thicknesses were investigated by scratch tests under different normal loads. Besides, the scratch observations on different microstructures were made to compare the wear resistance of the oxide layer formed on the martensite and ferrite. The detailed conclusions are described as follows.The composition of the oxide layer is mainly influenced by the oxidation temperature. At 200°C, the main oxide is Fe2O3; at 400°C, it is Fe3O4; at 600°C, the dominant is FeO. Both the oxidation temperature and the microstructure of the initial material have an effect on the thickness of the resulting oxide layer. The thickness of oxide scales increases with the increase of oxidation temperature. This leads to increase the activation energy of the oxidation and promotes the diffusion ability of the oxygen into the steel substrate, which increases the depth of oxidative region. As the MVF increases, the thickness of the oxide layer decreases. This is due to the varying oxidation properties of the different microstructures. The experiments in Chapters 3 and 4 have demonstrated that specimens with high MVF have stronger oxidation resistance compared to those with low MVF. The tests about oxide layer hardness were achieved on a nanoindenter. The nanoindentation results showed that the maximum hardness appeared on the oxide layer generated at 400 °C, with Fe3O4 as the dominant oxide.

Chapter 2

 2 describes the selected low alloy steel (25CD4) and the method of altering the microstructure (MVF). For the same steel, it has the same chemical composition. The obtained dual-phase microstructure are all consisted of ferrite and martensite. Samples with different MVF are generated by changing the heat treatment procedure. In this way, it is ensured that only one parameter (MVF) changes while the others remain constant during the investigation. The variation of the macro-hardness and microhardness with MVF are studied separately. The relevant conclusions are as follows:

Chapters 3

 3 and 4 are dedicated to the dry sliding tribological behavior under different experimental parameters. Friction coefficients, wear rates, and wear mechanisms are analyzed by varying the martensite volume fraction, sliding velocity, and gas environment, particularly focusing on oxidative wear. The purpose of these chapters is to establish the link between the influencing parameters and the tribological behavior of low-alloy steels. The main findings are described below:➢ The effect of microstructure on tribological behavior For COF, it decreases with decreasing MVF. For wear rate, samples with low MVF (≤ 59%) exhibits less wear rate compared to high MVF. the variation of COF and wear rate are attributed to the evolution of the oxidative wear. It increases with decreasing MVF. The formation of a compact oxide film (glazed layer) can effectively reduce wear, especially adhesion wear. For lower MVF, the main wear mechanism is oxidative wear. However, for higher MVF, adhesive wear and ploughing predominate. Besides, both plastic deformation and wear debris size decrease with increasing MVF.➢ The effect of sliding velocity on tribological behaviorBoth the COF and wear decrease with increasing sliding velocity. This is owing to the increase in oxidative wear as the sliding speed increases. The frictional oxide layer induced by the sliding velocity plays a protective role for the metal surface. As the sliding velocity increases, the effect of MVF on the wear decreases and the sliding velocity becomes the dominant factor. At 0.1m/s, the major wear mechanism is abrasive wear, accompanied by partial material flaking and less oxidation wear; at 0.6m/s, the wear mechanism is in the transition between abrasive wear and oxidation wear; at 1m/s, oxidation wear becomes the dominant wear mechanism.➢ The effect of gaseous environment on tribological behaviorFor the mean COF at steady state, the minimum COF occurs in the air environment regardless of the MVF. The maximum value is found in the carbon dioxide environment, where the value increases with decreasing MVF. The COF in argon and nitrogen environments are between that in air and carbon dioxide.The COF in argon is slightly greater than that in nitrogen. However, the variation of the COF in argon and nitrogen environments is independent of the MVF. It is worth noting that the longest transition time for COF occurs in the CO2 environment. It means that CO2 is indeed effective in reducing the COF during the initial rubbing period. For the wear rate, it is much lower in the CO2 environment compared to other gaseous environments, which is due to the generation of a relatively complete smooth compacted oxide layer on the wear tracks. In addition, the largest plastic deformation occurs in CO2, as well as the smaller debris size. For the wear mechanism, oxidative wear is the predominant in CO2 environments. In nitrogen and argon environments, more adhesive wear is observed accompanied by severe material spalling. Chapter 5 is focused on the effects of microstructure and oxidation temperature on the tribological behavior of the generated oxide layers. The COF of the different oxide layers was obtained by scratch tests and specific wear resistance was calculated. The microscopic observation was carried out to analyze the varying wear mechanisms. The related conclusions are as follows: Both the oxidation temperature and the microstructure of the initial material have an effect on the thickness of oxide scale. It increases with increasing oxidation temperature and decreases with increasing MVF. The hardness has a stronger effect on the COF in the scratch test, while the MVF of substrate and oxide layer thickness are insensitive to the COF of oxide layer. The maximum wear rate occurred on the FeOdominated oxide layer exposed at 600°C. The wear rate decreased as the increasing MVF and the decreasing thickness. Both the COF and the wear rate of the different oxide layers increased with increasing normal load. The main wear mechanisms are in the transition between ploughing and micro-cutting on the Fe2O3-dominated oxide layer, and ploughing on the FeO-dominated oxide layer. Micro-cutting mainly occurs on the Fe3O4-dominated oxide layer formed at 400 °C. Under the same conditions, the oxidation properties of the different microstructures on dual-phase steels are different.

Table 2 .

 2 1: The chemical composition of 25CD4 steel (weight %) .......................................... 44

Table 3 .

 3 1: Oxygen concentration (in at%) of wear tracks at different MVF ........................... 60 Table 3.2: Elemental composition (in at%) of wear tracks and debris at MVF 59% by EDX analysis ............................................................................................................................ 61 Table 3.3: Elemental composition (in at%.) of wear tracks and debris at 1m/s and MVF 68% by EDX ........................................................................................................................... 75

Table 4 .

 4 1: Oxygen concentration (in at%.) of wear tracks from air, Ar, N2, and CO2 atmospheres at MVF 80% .................................................................................................................... 93 Table 4.2: Elemental composition (in at%.) of wear tracks and debris from CO2 atmospheres at MVF 48% by EDX ...................................................................................................... 94 Table 4.3: Oxygen concentration (in at%.) of wear tracks from air, N2, Ar, and CO2 atmospheres at MVF 93%,68 and 48% ................................................................................................ 94

Table 2 . 1 :

 21 The chemical composition of 25CD4 steel (weight %)

	C	Mn	Si	S	P	Cr	Ni	Mo
	0.25	≤0.90	≤0.40	≤0.035 ≤0.035	≤1.20	-	≤0.25

Table 3 .

 3

		1: Oxygen concentration (in at%) of wear tracks at different MVF
	MVF	Test 1 (%)	Test 2 (%)	Test 3 (%)	Average (%)
	MVF100%	7.57	5.23	7.53	6.78
	MVF93%	6.88	12.28	13.13	10.76
	MVF80%	18.22	19.51	19.12	18.95
	MVF68%	21.18	18.58	21.98	20.58
	MVF59%	29.33	26.64	27.78	27.92
	MVF48%	31.68	28.07	27.2	28.98

Table 3 .

 3

		2: Elemental composition (in at%) of wear tracks and debris at MVF 59% by EDX
			analysis
	Label	W%(C)	W%(O

) W%(Al) W%(Si) W%(Cr) W%(Mn) W%(Fe) W%(Mo)

	1	0.45	6.77	0.05	0.24	0.97	0.85	90.22	0.46
	2	1.00	2.84	0.00	0.15	0.97	0.83	94.00	0.22
	3	0.88	35.02	0.32	0.18	0.58	0.30	62.34	0.37
	4	0.80	35.39	0.40	0.20	0.60	0.49	61.79	0.33
	5	1.44	44.00	0.61	0.22	0.41	0.42	52.81	0.08
	6	1.24	3.09	0.05	0.20	1.12	0.76	93.23	0.30
	7	1.43	23.33	0.21	0.23	0.80	0.56	73.30	0.14
	8	1.02	3.97	0.00	0.18	0.94	0.76	92.78	0.34

Table 3 . 3

 33

: Elemental composition (in at%.) of wear tracks and debris at 1m/s and MVF 68% by EDX

Label W%(C ) W%(O ) W%(Al) W%(Si) W%(Cr) W%(Mn) W%(Fe) W%(Co)

	1	4.022	4.5926	0	0.3071	0.7711	0.6053	87.9653	1.7365
	2	2.7241	2.5699	0	0.2941	0.8323	0.7048	91.313	1.5617
	3	3.3139	2.7304	0.0586	0.3922	0.8202	0.6627	90.2602	1.7617
	4	3.3462	23.3195	0.1299	0.3232	0.7502	0.5561	70.2369	1.3379
	5	4.5221	10.6651	0.2546	0.5331	0.7478	0.5687	81.1351	1.5734
	6	5.8012	11.36	0.1266	0.3991	0.688	0.5637	79.2962	1.7652
	7	3.3081	2.6749	0.0265	0.3606	0.8199	0.6019	90.6245	1.5837
	8	2.9506	2.8476	0.0177	0.3469	0.7849	0.6783	90.5717	1.8024

Table 4 .

 4 1: Oxygen concentration (in at%.) of wear tracks from air, Ar, N2, and CO2 atmospheres at MVF 80% Figure4.10: EDX analysis on the points for CO2 atmosphere at MVF 48%

Table 4 .

 4 2: Elemental composition (in at%.) of wear tracks and debris from CO2 atmospheres at MVF 48% by EDX

Table 4 . 3

 43

: Oxygen concentration (in at%.) of wear tracks from air, N2, Ar, and CO2 atmospheres at MVF 93%,68 and 48%

259(1-6): p

  behaviour ofFe-12Cr steel. Tribology international, 1989. 22(4): p. 243-252. 139. Hong, H., R. Hochman, and T. Quinn, A new approach to the oxidational theory of mild wear. STLE TRANSACTIONS, 1988. 31(1): p. 71-75. 140. Sullivan, J., T. Quinn, and D. Rowson, Developments in the oxidational theory of mild wear. Tribology International, 1980. 13(4): p. 153-158. 141. Mølgaard, J., A discussion of oxidation, oxide thickness and oxide transfer in wear. Wear, 1976. 40(3): p. 277-291. 142. Sullivan, J. and S. Hodgson, A study of mild oxidational wear for conditions of low load and speed. Wear, 1988. 121(1): p. 95-106. 143. Sexton, M. and T. Fischer, The mild wear of 52100 steel. Wear, 1984. 96(1): p. Investigation of the role of oxide scale on Stellite 21 modified with yttrium in resisting wear at elevated temperatures. Wear, 2005. . 453-458. 148. Wang, L. and D. Li, Effects of yttrium on microstructure, mechanical properties and high-temperature wear behavior of cast Stellite 6 alloy. Wear, 2003. 255(1-6): p. 535-544. 149. Radu, I., D. Li, and R. Llewellyn, Tribological behavior of Stellite 21 modified with yttrium. Wear, 2004. 257(11): p. 1154-1166. 150. Fehlner, F.P. and N.F. Mott, Low-temperature oxidation. Oxidation of metals,

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144. Lim, S., et al., The unlubricated wear of sintered steels. Tribology international, 1987. 20(3): p. 144-149. 145. Quinn, T., J. Sullivan, and D. Rowson, Origins and development of oxidational wear at low ambient temperatures. Wear, 1984. 94(2): p. 175-191. 146. Fontalvo, G. and C. Mitterer, The effect of oxide-forming alloying elements on the high temperature wear of a hot work steel. Wear, 2005. 258(10): p. 1491-1499. 147. Radu, I. and D. Li, 1970. 2(1): p. 59-99. 151. Evans, H., D. Norfolk, and T. Swan, Perturbation of parabolic kinetics resulting

Acknowledgements

II

Acknowledgements

Acknowledgements III

layers.

The subject is focused on the effect of microstructure on oxidative wear by changing the gas environment and sliding speed under low surface roughness. In addition, the tribological properties of oxide layers formed at different temperatures was also investigated for low alloy steel at microscopic scale. Dry sliding friction tests were performed by using a ball-on-disk tribometer (CSM Tribometer). The wear rate calculation of ball/disk sliding experiments was realized on a 3D optical profilometer.

In order to investigate the mechanical and tribological properties of the different oxide layers, the experiments of nanoindentation and the scratch tests were carried out on the nano-indenter (Nano-indenter Agilent G200) over a wide loading range. The composition of the oxide layer is determined by a combination of physical characterization and electron backscatter diffraction (EBSD) phase identification.

Scanning Electron Microscope (SEM) and Energy-dispersive X-ray spectroscopy (EDX) were used as complementary techniques to characterize the and wear mechanisms. The research objectives are as follows.

For (a) and (b):

✓ How does the microstructure impact the formation of oxides during friction? ✓ How do sliding velocity and gaseous environment affect oxidative wear? ✓ How do the oxides impact the tribological behavior of the contact surfaces? ✓ Whether the formed oxide protective layer can reduce wear.

For the properties of oxide scales:

✓ Does the microstructure impact the formation of different types of oxides at high temperatures?

✓ How does the nature and thickness of these oxide layers impact tribological behavior?

✓ Whether the adhesion between the oxide layer and substrate is favorable, and whether three-body wear occurs?