Ülkü Anik

Yudum Tepeli

Maher Sayhi

Jihene Nsiri

Mohamed Fethi Diouani

Towards the Electrochemical Diagnostic of Influenza Virus: Development of Graphene-Au Hybrid Nanocomposite Modified Influenza Virus Biosensor Based on Neuraminidase Activity

An effective electrochemical influenza A biosensor based on graphenegold (Au) hybrid nanocomposite modified Au-screen printed electrode has been developed. The working principle of the developed biosensor relies on the measurement of neuraminidase (N) activity. After the optimization of experimental parameters like the effect of bovine serum albumin addition, immobilization time of fetuin A and immobilization time of PNA lectin, the analytical characteristics of the influenza A biosensor were investigated. As a result, a linear range between 10 -8 U/mL and 10 -1 U/mL has been found with relative standard deviation value of 3.23% (for 10 -5 U/mL of N, n:3) and limit of detection value as 10 - 8 U/mL N .The developed biosensor has also been applied for real influenza virus A (H9N2) detection and very successful results were obtained.

Introduction

Influenza virus A can be described as a negative stranded RNA virus which belongs to Orthomyxoviridae family. The virus contains two surface glycoproteins namely hemagglutinin (H) and neuraminidase (N). The classification of the virus subtype can be made according to the antigenic properties of these 18 H (1-18) and 11 N (1-11) glycoproteins [1][2][3][4][5][6][7][8][9][10][11][12][13][14] . Since the influenza virus is very mutagenic, it can easily change the antigenic portions of H and N proteins and as a result very serious antigenic drift has been occurred 15 . Seasonal influenza viruses can easily be affected by these antigenic drift mutations which cause millions of serious infections and approximately 500.000 deaths for every year 15,16 . Also, sometimes two or more influenza A viruses of different origin infects the same cell and as a result, new strains or subtypes are emerged because of the occurrence of genetic reassortment 1

. As an example, the reassortment between human and avian virus strains was the reason of influenza A pandemics in 1957 (H2N2) and in 1968 (H3N2) while in 2009 (H1N1) pandemic virus was found to be a reassortant containing gene segments from human, avian and swine influenza viruses 2,4 . Besides these, highly mortal avian viruses like H5N1 and H7N9 have been discovered to transmit from birds to humans 4,17 . Because of these, some researchers define the virus as "continuously emerging infectious disease" which means it still continues to be a threat for animals and human beings 15 .

Although rapid detection of this virus is mandatory, present influenza diagnosis techniques cannot compete with the mutagenic behavior of the influenza virus. For example, viral culture, which was accepted as a golden standard method, needs two or three days to retrieve the results which is a kind of long time for the diagnosis of influenza virus [18][19][20] . On the other hand, quantitative polymerase chain reaction (qPCR), which is chosen frequently by the health personnel can be described as sensitive but relatively slow (took about 2 h). qPCR is also expensive and can be performed only with the specialists 20-22 . Electrochemical techniques are accepted as practical techniques for many applications 25,26 . The combination of electrochemical techniques with biosensor systems, results with electrochemical biosensors which provide not only practicality but selectivity and sensitivity as well 27,28 . For these reasons, electrochemical biosensors can be accepted as good candidates to be used in point of care systems 29 .

Recently electrochemical biosensors were combined with nanomaterials and as a result, more sensitive and accurate results were obtained 30-32 . Graphene is a kind of twodimensional nanomaterial and provides properties like higher surface area and good conductivity which are very important for the electrochemical biosensors 33-35 . Apart from graphene, recently graphene-metallic nanocomposites were produced and used in these systems. These nanocomposites provide more sensitive and selective results when combined with the biosensor systems [36][37][38][39][40][41][42][43] Considering the influenza biosensors, it can be stated that, up to now almost only H based systems have been developed [7][8][9][10][44][45][46][47][48] . Lately our group managed to develop a N based electrochemical influenza biosensor 14 . In that work, as a preliminary data, only the electrochemical impedance spectroscopy (EIS) diagrams about the fabrication of biosensor were presented 14 . For that earlier work, glassy carbon paste electrode was used as the main electrode. Firstly fetuin A that is a kind of glycoprotein was attached onto the electrode. Then, N was immobilized onto fetuin A. Fetuin A includes terminal 12-14 sialic acid residues per molecule and N clevages fetuin A from this sialic acid ends. Lastly, peanut agglutinin (PNA) lectin was immobilized onto the electrode surface to monitor this cleavage 14 .

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In this work, as a continuation of our previous paper, we improve our system by using gold screen printed electrode (AuSPE) together with graphene-Au hybrid nanocomposite. Also the experimental parameters about the developed system were optimized and analytical characteristics were examined. Lastly, the developed biosensor was applied for real influenza A virus (H9N2) detection.

Experimental

Synthesis procedure of graphene-Au nanocomposite

In order to synthesize graphene-Au nanocomposite, graphene oxide (GO) was used as the starting material. Therefore, first GO was synthesized by using modified Hummers-Offeman method from graphite 38,49 . For this purpose, 1g of graphite powder was added into 23 mL 98% H 2 SO 4 solution and stirred at room temperature for 24 h. Then, 100 mg of NaNO 3 was added into the mixture and stirred for another 30 min. After the mixture was cooled to 5 o C by using an ice-bath, 3 mg of KMnO 4 was added to the medium and then the stirred mixture was heated up to 40 o C. 46 mL of double distilled water was added into above mixture during a period of 25 min. Finally, 140 mL of double distilled water and 10 mL of 30% H 2 O 2 were added into the mixture to stop the reaction. The unexploited graphite in the resulting mixture was removed by centrifugation. Then the mixture was dried in a desiccator at room temperature.

Additionally, for the graphene-Au nanocomposite synthesis step, 10 mg portion of GO powder which was synthesized by modified Hummers-Offeman method was dispersed in 10 mL of water by sonication for 1 h to form stable GO colloid 50, 51 . Then, 20 mL of EG solution and 0.5 mL of 0.01 M HAuCl 4 .H 2 O were added to the GO colloid and stirred for 30 min. After that, the mixture was heated at 100 o C for 6 h by applying the magnetic stirring. Subsequently, the graphene-Au nanocomposites were separated from the EG solution via centrifugation and washed with deionized water for five times. The resulting product was dried in a vacuum oven at 60 o C for 12 h 37 . Finally, the synthesized graphene-Au nanocompsite was dispersed to 10 mg/mL in double distilled water by ultrasonication and stored at 4 o C until it was used 38, 40, 41, 52 .

Preparation of developed electrochemical

ELLA measurements

An Enzyme-Linked Lectin Assay (ELLA) was performed to measure influenza A virus N activity as described by Couzens, L. et al. 3 . Fetuin-coated plate was prepared by dispense 100µl/well of 25µg/ml fetuin working solution into each well. The plate was covered and placed at +4°C for 24hours. Then, the plate was washed 3 times with 200µl/well PBS wash buffer for 3 minutes. 100µl sample diluent buffer was added in column 12 as negative control and 50µl serial dilution antigen (influenza virus, N and non-infected allontoic fluid) was transferred from the dilution plate to each well in column 1-11 of the fetuin-coated plate contain an equal volume of sample diluent buffer. The plate was incubated in humidified incubator with 5% CO 2 for 18hours at 37°C, and then washed 6 times as described before. 100µl/well PNA-peroxidase was dispensed into each well and incubated at room temperature for 2 hours. After this step, 3 times wash was performed before adding of 100µl/well freshly prepared substrate solution for 10 min in dark at room temperature. The color reaction was stopped by the addition of 100µl/well stop solution. The plates were read at 490 nm using 96-well Automatic plate reader.

Results and Discussion

Characterization of synthesized graphene-Au hybrid nanocomposite

Fig. 1 demonstrates the SEM images and EDS results of graphene-Au nanocomposite. It can be seen from the Fig. 1 that AuNps are located on the graphene sheets as white dots (average diameter of 250 nm). EDS results show that atomic dispersion percentages of graphene-Au nanocomposite are 83.73 %, 14.97 % ,1.06 % and mass percentages of these elements 69.29%, 14.66% and 14.32% for C, O and Au, respectively.

Electrochemical characterization of developed electrochemical Influenza A biosensor

The fabrication of developed influenza A biosensor was monitored via EIS in the presence of 10 mM K 3 [Fe(CN) 6 ]/K 4 [Fe(CN) 6 ]. From the Nyquist plots of EIS (Fig. 2); it can be seen that bare AuSPE has the small semi-circle domain because of high conductivity of Au surface (curve a). After the electrode modification with graphene-Au hybrid nanocomposite, semi-circle domain increases a little bit (curve Please do not adjust margins Please do not adjust margins b). After that, when the electrode surface was covered with fetuin A, the resistance of the electrode surface increases since active surface area was blocked (curve c). Then, N was immobilized on to fetuin A from the sialic acid sides. Therefore, electrode surface was covered with N and charge transfer ability of the electrode surface was reduced resulting with the increment in the semicircle diameter (curve d). Lastly, PNA lectin which has specificity to galactose molecules that appears after the cleavage of fetuin A with N, was used. When the PNA lectin binds onto the galactose molecules, the electron transfer was getting even harder and semi-circle diameter also increases due to the resistance increase on the electrode surface (curve e).

Optimization of Experimental Parameters

The experimental parameters consisting of the effect of BSA addition on developed system, immobilization time of fetuin A, immobilization time of PNA lectin were investigated by EIS. Since developed system is considered to be applied for human samples like throat swabs in future, physiological human pH and body temperature were chosen as the working conditions so that all the experiments were carried out at pH 7.4 and incubation step of N was done at 37 o C .

The effect of BSA on the developed system

As it is well known, BSA is generally used for the blockage of unspecific interaction on the electrode surface, after the immobilization of the analyte. However, even for this step, it has to be waited for a period of time for the completeness of the attachment of BSA. In order to observe this effect, developed influenza A biosensor was fabricated in the presence of BSA, in the absence of BSA and with the simultaneous immobilization of fetuin A at the same time. All these steps were monitored by EIS under the previously mentioned working conditions. The obtained EIS diagrams (Supporting Information, Fig. S1.) demonstrate that for our system there isn't any necessity for BSA addition which means that, there is no unspecific binding during the fabrication of the developed influenza A biosensor. Also by eliminating the addition of BSA, the preparation procedure of biosensor is shortened and as a result, the practicality of the fabrication procedure has been increased. In conclusion, no BSA attachment onto the electrode surface was done for the future experiments.

Optimization of fetuin A immobilization time

Fetuin A immobilization time experiments were carried out for 15 min, 30 min, 45 min, 60 min and 90 min incubation times and the results were monitored by using EIS. The obtained EIS diagrams were presented at the supporting information (Supporting Information, Fig. S2) and the excel plots in it. It can be seen from the Fig. S2. that the best result was obtained at 30 min. Incubation times less than 30 min. results with lower difference in electrode resistance demonstrating that more times are needed probably for the exchange reactions between fetuin A amino groups and NHS to be completed. On the other hand, incubation times more than 30 min, resulted with decrease in resistance difference that might be attributed to the separation of non-specific bonding from the electrode surface. As a result, further experiments were conducted by using 30 min. as the optimum incubation time for fetuin A.

The effect of PNA Lectin immobilization time

PNA lectin immobilization time experiments were carried out for 30 min, 45 min, 60 min and 90 min immobilization times by using EIS. The obtained EIS diagrams were presented at supporting information (Supporting Information, Fig. S3) and the excel plots were demonstrated in Fig. S3. also. It can be seen from the Fig. S3. that the best result was obtained at 60 min. Immobilization times, less than 60 min. results with lower difference in electrode resistance demonstrating that more times are needed probably for the specific binding reactions between PNA lectin and galactose molecules to be completed. Immobilization times that lasted longer than 60 min, resulted with decrease in resistance difference that might be attributed to the falling out of PNA-galactose structure from the electrode surface 11 . Please do not adjust margins Please do not adjust margins

ARTICLE

Analytical Characteristics

After the optimization of experimental parameters, the analytical characteristics were examined. For obtaining a calibration graph, the specific interactions between PNA lectin and galactose molecules that occur after the cleavages of N from the fetuin A were monitored. When the concentration of N is increased, since more fetuin A will be cleavage, more galactose molecules will appear and as a result, more PNA lectin will link to the galactose ends. All these changes were monitored via resistance values that were obtained from EIS measurements (Fig. 3.) As you can see from the Fig. 3., the linear range and limit of detection (LOD) values which based on N concentration was obtained between 10 -8 U/mL and 10 -1

U/mL with the equations of y=2746x+31.69 (R 2 =0.99) and 1 x 10 -8 U/mL, (stated by brand that 7.9 units/mg solid) respectively. Relative standard deviation (R.S.D) value was calculated for 10 -5 U/mL N (n=3) and found as 3.23%. For the comparison, obtained analytical characteristics values of developed electrochemical biosensor were compared with influenza A virus enzyme-linked immunosorbent (ELISA) assay results in literature. These ELISA assays relied on the N activity or antibody-antigen interaction. From these ELISA assays LOD values such as 0.7 ng/mL antigen 53 , 33 ng/mL antigen 53 , 0.5 ng/mL protein 54 and 10 2 -10 3 Tissue Culture Infectious Dose 50% (TCID 50 )

55 can be seen. From these results it is obvious that developed sensor is sensitive enough. On the other hand considering rapidity and practicality, it can be said that developed electrochemical biosensor could definitely offer better performances. -2 U/mL and g)10 -1 U/mL; which were used for calibration graph and h)the calibration graph of developed electrochemical influenza virus A biosensor.

Sample application and control studies

Evaluation of ELLA results

We developed the electrochemical influenza virus A biosensor which based on N activity. Therefore, before we applied developed electrochemical influenza virus A biosensor for the real H9N2 influenza type A virus detection, ELLA experiment was carried out to demonstrate that isolated H9N2 influenza type A virus is active according to N (Supporting Information, Fig. S4). Moreover, ELLA measurements were carried out with Please do not adjust margins control group to demonstrate that control group is not active according to N. For this purpose, the ELLA results of H9N2 virus and control group were compared with the ELLA results of standard N. According ELLA results, LOD value of N was found as 6.52x10 -6 U/mL. As a result, we also supported the accuracy and sensitivity of the developed influenza virus A biosensors working principle with the ELLA using the same method.

Evaluation of EIS results

On the other hand, developed and optimized influenza virus A biosensor was applied for real virus detection, H9N2, that was prepared according to the procedure given in experimental part. For this purpose, during the fabrication of electrochemical influenza virus A biosensor, H9N2 influenza virus A sample was used instead of N. After that, in order to determine the N activity of H9N2 influenza virus, PNA lectin immobilization step was done.

The control study was also carried out by using uninfected egg sample (allantoic fluid). It can be seen from the Fig. 4. that the developed electrochemical influenza biosensor detects influenza A virus selectively.

Conclusions

Another novel and effective electrochemical influenza A biosensor based on graphene-Au modified Au-SPE has been developed. This work is the continuation of our previous work where working principle relies on the measurement of N activity 14 . From the LOD values, it is obvious that with developed electrochemical influenza A biosensor more sensitive results were obtained compared to ELLA assay. We believe that, the wide linear range and sensitive results from real influenza A sample (H9N2) analysis, increase the potential of usage of our system as a point of care tool for the diagnosis of influenza virus A in the future 56 .

Scheme 1 .

 1 Scheme 1. Schematic diagram depicting the steps of developed electrochemical influenza A biosensor.

Figure 1 .

 1 Figure 1. A), B) SEM images of graphene-Au nanocomposite and C) EDS results of graphene-Au nanocomposite.

Figure 2 .

 2 Figure 2. Nyquist plots of developed influenza A viral biosensor. a.Plain AuSPE, b. AuSPE/graphene-AuNp, c. AuSPE/graphene-AuNp/fetuin A, d. AuSPE/graphene-AuNp/fetuin A/N, e. AuSPE/graphene-AuNp/fetuin A/N/PNA lectin. The EIS procedure was set to measure the electron transfer resistance in the frequency range of 0.1 Hz -10 kHz at potential of 0.1 V and 10 mM K 3 [Fe(CN) 6 ]/K 4 [Fe(CN) 6 ] (in pH:7.4, 0.1 M PBS) was used as a redox probe.

Journal Name 6 |

 6 J. Name., 2012, 00, 1-3 This journal is © The Royal Society of Chemistry 20xx

Figure 3 . 10 - 8 U

 3108 Figure 3. Nyquist plots of different concentration of N; a)10 -8 U/mL, b)2x10-8 U/mL, c)10 -6 U/mL, d)10 -5 U/mL, e)10 -3 U/mL, f)10 -2 U/mL and g)10 -1 U/mL; which were used for calibration graph and h)the calibration graph of developed electrochemical influenza virus A biosensor.

Figure 4 .

 4 Figure 4. Nyquist plots for A) 64 UHA H9N2 influenza virus, B) control group by using developed electrochemical influenza virus A.

Analyst Accepted ManuscriptPublished on 26 October 2017. Downloaded by University of Reading on 26/10/2017 23:05:28.View Article Online

Published on 26 October 2017. Downloaded by University of Reading on 26/10/2017 23:05:28.View Article Online

Analyst Accepted ManuscriptPublished on 26 October 2017. Downloaded by University of Reading on 26/10/2017 23:05:28. View Article Online DOI: 10.1039/C7AN01537B

Acknowledgements

The grant from The Technical and Scientific Council of Turkey (TUBITAK) Project no: 114Z654 is gratefully acknowledged. The Grant from Tunisian Ministry of Higher Education and Scientific Research (MHESRT): Projet Tuniso-Turque 2015-2017.

Conflicts of interest

There are no conflicts to declare. Please do not adjust margins Please do not adjust margins

ARTICLE


M Salem

S Le Roux

G Dour

P Lamesle

K Choquet

F Rézaï-Aria

Effect of aluminizing and oxidation on the thermal fatigue damage of hot work tool steels for high pressure die casting applications

Keywords:

The interaction between thermal fatigue and aluminizing and/or oxidation is investigated using an experimental approach based on decoupling of mechanisms. Virgin and pre-aluminized steel specimens are tested in air and nitrogen between 100 and 650 °C. Homogeneous uniaxial micro-crack network forms on the oxidised or prealuminized surface in air, with a better resistance to micro-cracking for the intermetallic coating. The propagation of the micro-cracks is delayed in nitrogen, whilst no evidence of micro-cracking is observed on the virgin specimen. The premature cracking of the steel depends on the formation of the superficial micro-crack network, and the crack propagation is assisted by oxidation.

Introduction

In the High Pressure Die Casting (HPDC) process, molten aluminium is injected between 650 and 700 °C in the cavities of the mould, at a speed higher than 30 m/s and under high pressure (up to 1300 bar) [START_REF] Kaufman | Aluminum alloy castings: properties, processes, and applications[END_REF]. The surface of the mould is thus exposed to strong thermo-mechanical loadings and severe corrosive attacks. It is damaged by complex mechanisms combining thermal fatigue (TF), wear (by shock or friction) and corrosion (by aluminizing and oxidation) [START_REF] Manson | Thermal fatigue of die casting dies -A Review and Proposed Future Program[END_REF][START_REF] Sjostrom | Thermal fatigue in hot-working tools[END_REF], which impact the industrial process in different ways [START_REF] Graham | Thermal Processing Structure and Thermal Fatigue Relation for Die Steel[END_REF].

After several thousand cycles, the mould cavities (especially the gate areas) may suffer a substantial and irreversible loss of materialcalled "washout" -, which decreases their ability to meet the dimensional tolerance of cast parts [START_REF] Norström | General aspects on "washout" mechanisms in aluminium die casting dies[END_REF][START_REF] Chen | Die erosion and its effect on soldering formation in high pressure die casting of aluminium alloys[END_REF]. This phenomenon is attributed to the interaction between erosion and corrosion, which leads eventually to the dissolution of the die [START_REF] Chen | Die erosion and its effect on soldering formation in high pressure die casting of aluminium alloys[END_REF][START_REF] Shankar | Die soldering -A metallurgical analysis of the molten aluminium/die interface reactions[END_REF][START_REF] Gopal | Soldering in die casting: aluminum alloy and die Steel interactions[END_REF]. In service, the surface of the die keeps its initial quality as long as it remains protected against oxidation or aluminizing by the lubricant coating [START_REF] Gopal | Soldering in die casting: aluminum alloy and die Steel interactions[END_REF][START_REF] Salem | endommagement par fatigue thermique des moules de fonderie sous pression d'aluminium: effet de l'interaction avec l'aluminisation et l'oxydation[END_REF], as shown in Fig. 1a. By stripping the die surface in areas of high flow (e.g. gate) or with direct impingement (e.g. pins in front of the gate), erosion enables the molten aluminium to come into contact with the surface of the steel mould. The inter-diffusion of aluminium and iron atoms results in the formation of intermetallic compounds at the interface between cast alloy and steel [START_REF] Salem | endommagement par fatigue thermique des moules de fonderie sous pression d'aluminium: effet de l'interaction avec l'aluminisation et l'oxydation[END_REF][START_REF] Chacrone | Kinetic and metallographic study of the formation of the alloy layer during the coating of steel with an aluminium-silicon mixture[END_REF]. Post-mortem metallographic analyses of industrial dies [START_REF] Salem | endommagement par fatigue thermique des moules de fonderie sous pression d'aluminium: effet de l'interaction avec l'aluminisation et l'oxydation[END_REF][START_REF] Mitterer | Application of hard coatings in aluminium die castingsoldering, erosion and thermal fatigue behaviour[END_REF][START_REF] Kosec | Analysis of casting die failures[END_REF] have indeed highlighted the presence of oxide (Fig. 1b) and/ or intermetallic phases (Fig. 1c), both in surface and inside cracks. Oxide and intermetallic compounds contribute to modify the chemical composition of the steel sub-surface, and consequently its mechanical behaviour [START_REF] Sundqvist | Effects of liquid aluminium on hot-work tool steel[END_REF]. Among the phases composing the complex aluminide layer, the Fe 2 Al 5 (Si) phase formed just above the mould surface is particularly compact and brittle. In addition, sticking problemsalso known as "soldering" -manifests by a total or partial attachment of the cast part on the die. This adhesion may be caused by chemical interaction between aluminium and iron [START_REF] Shankar | Die soldering -A metallurgical analysis of the molten aluminium/die interface reactions[END_REF][START_REF] Gopal | Soldering in die casting: aluminum alloy and die Steel interactions[END_REF][START_REF] Joshi | Intermetallic formation and its relation to interface mass loss and tribology in die casting dies[END_REF][START_REF] Zhu | Evaluation of soldering, washout and thermal fatigue resistance of advanced metal materials for aluminum die-casting dies[END_REF][START_REF] Chen | Formation and progression of die soldering during high pressure die casting[END_REF] or by mechanical bonding, due to the infiltration of the aluminium through the cracks from the "heat-checking" which is the main damage mode experienced by mould cavities after a large number of cycles [START_REF] Kosec | Analysis of casting die failures[END_REF][START_REF] Persson | Thermal fatigue cracking of surface engineered hot work tool steels[END_REF][START_REF] Miquel | Heat-checking of hot work tool steels[END_REF]. The soldering issues require stopping the production to remove the part manually and rectifying the tool surface by mechanical processes (machining, grinding, etc.) [START_REF] Chen | Die erosion and its effect on soldering formation in high pressure die casting of aluminium alloys[END_REF][START_REF] Gopal | Soldering in die casting: aluminum alloy and die Steel interactions[END_REF]. In service, the die surface undergoes strong alternate heating-cooling cycles, respectively in the injection phase and the opening step during which the part is ejected and a lubricant is sprayed onto the mould surface. As a result of these high transient thermal gradients, thermal stresses are induced at the surface of the mould cavity where cracks are initiated by anisothermal low cycle fatigue process. The micro-cracks grow to form a superficial network, either multi-axial on plane surfaces (Fig. 1d) or uniaxial in geometrical singularities (Fig. 1e) [START_REF] Medjedoub | Effect of local stress on the heatchecking morphology in high temperature tool steels under thermal fatigue: Transition from multi-axiality to uniaxiality[END_REF]. The formation of a macroscopic heat-checking network on the mould significantly reduces the quality of produced parts, requiring subsequent finishing operations and the scrapping of the die when the defects are too large. Moreover, macrocracks can propagate in-depth in the tool, leading to its premature retirement if not rupture [START_REF] Graham | Thermal Processing Structure and Thermal Fatigue Relation for Die Steel[END_REF][START_REF] Danzer | Thermal fatigue cracks in pressure die casting dies[END_REF][START_REF] Berns | Beispiele zur schädigung von warmarbeitswerkzeugen[END_REF].

Understanding the interactions between the different damage mechanisms involved in HPDC is therefore essential for the casting industry, in order to optimize the service conditions and to better design the material solutions to extend the tool life. The effect of the environment, especially oxidation, is well known in low cycle fatigue (LCF) [START_REF] Tsujii | Effect of testing atmosphere on low cycle fatigue of hot work tool steel at elevated temperature[END_REF], but not yet clearly identified in TF, despite a large number of reported studies [START_REF] Sjostrom | Thermal fatigue in hot-working tools[END_REF][START_REF] Miquel | Heat-checking of hot work tool steels[END_REF][START_REF] Howes | A Study of Thermal Fatigue Mechanisms[END_REF][START_REF] Malm | Material-related model for thermal fatigue applied to tool steels in hot-work applications[END_REF][START_REF] Wallace | Effect of composition and processing on the thermal fatigue and toughness of high performance die steel[END_REF][START_REF] Ebner | Thermal fatigue behaviour of hot-work tool steels: heat check nucleation and growth[END_REF][START_REF] Medjedoub | Experimental conditions and environment effects on thermal fatigue damage accumulation and life of die-casting steel X38CrMoV5 (AISI H11)[END_REF][START_REF] Klobčar | Thermal fatigue of materials for die-casting tooling[END_REF]. All these works emphasize the impact of the oxidation or/and aluminizing on TF damage of tool steels. But unlike isothermal fatigue tests, the TF test is still not standardized and the applied solicitations vary depending on the heating and cooling devices and the specimen geometry chosen by each author. This diversity of test conditions complicates the comparison of the different TF results in order to quantify the environment effects objectively.

This contribution is an attempt to better understand the influence of the corrosive environmentboth oxidation and aluminizingon the TF behaviour of a hot work tool steel X38CrMoV5 (AISI H11). The role of the Fe-Al intermetallic compounds, resulting from the steel aluminizing, is investigated based on TF experiments performed on coated specimens. Several techniques can be used to prepare an aluminium coating, such as Physical Vapour Deposition (PVD), Chemical Vapour Deposition (CVD), Hot-Dipping (HD), Electro-Chemical Deposition (ECD), Plasma Sputtering (PS), cladding, spraying, etc. [START_REF] Nicholls | The production and properties of aluminium coatings[END_REF][START_REF] Drewett | Diffusion coatings for the protection of iron and steel: Part 1: Aluminium coatings[END_REF][START_REF] Majumdar | Formation of Al 2 O 3 /Fe-Al layers on SS 316 surface by pack aluminizing and heat treatment[END_REF]. In our case, a low-temperature process allowing the formation of a homogeneous layer on a complex geometry was required, in order to avoid microstructural changes of the heat-treated steel, which would result in the degradation of its mechanical properties by thermal softening [START_REF] Medvedeva | High-temperature properties and microstructural stability of hot-work tool steels[END_REF]. Among processes meeting these constraints, the low-temperature CVD process using Pack Cementation (PC) was selected for its simplicity of use and because it allows a fine surface finish [START_REF] Salem | endommagement par fatigue thermique des moules de fonderie sous pression d'aluminium: effet de l'interaction avec l'aluminisation et l'oxydation[END_REF]. To investigate the role of oxidation, TF fatigue experiments are performed in air and nitrogen, and compared to the same tests on uncoated specimens (virgin steel). The superficial damage and microstructural evolutions are analysed using Scanning Electron Microscopy (SEM) and X-Ray Diffraction (XRD). The effects of the intermetallic and oxide layers on crack initiation and lifetime of the steel are presented and discussed.

Experimental procedure

Material and specimens

A low silicon and impurity grade of X38CrMoV5 tool steel (AISI H11), known as ADC3 and provided by Aubert & Duval, was investigated. Its chemical composition is given in Table 1. The steel was heat-treated by austenitisation at 980 °C for 1 h, then quenching in air, and double tempering at 550 and 605 °C for 2 h at each temperature, resulting in a martensitic microstructure with a final hardness of about 47 HRC [START_REF] Salem | endommagement par fatigue thermique des moules de fonderie sous pression d'aluminium: effet de l'interaction avec l'aluminisation et l'oxydation[END_REF].

Disk shaped axi-symmetric specimens with an outer diameter of 30 mm and 1 mm thickness were machined from the heat-treated bars. An axial hole of 10 mm diameter was drilled in the specimen for internal water-cooling (Fig. 2). The surface of the specimen at tips was polished along the circumferential direction (θθ) to achieve a mean roughness Ra of about 0.05 μm.

Thermal fatigue experiments

A new TF rig, equipped with a vacuum chamber with an Al-double water-cooled wall, was developed in order to perform TF experiments in ambient laboratory air or inert atmosphere (nitrogen N 2 or argon Ar). The heating facility consisted of a 25 kW-Hüttinger high frequency induction generator, coupled with a solenoid inductor coil placed inside the chamber (Fig. 3a). The induction frequency was about 2-4 MHz, resulting in a rapid heating of the outer radius (tip) of the specimen. The depth of the induction "skin effect" was estimated to be a few tens of micrometres. The specimen was carefully positioned inside the induction coil using a movable sample holder (Fig. 3b). A temperature variation less than 5 °C (at 650 °C) was achieved along the circumferential direction. The inner wall of the specimen was continuously cooled by water at 20 °C, with a flow rate of about 10 l/min. The external surface of the specimen was naturally cooled by convection and radiative losses. Severe thermal gradients were induced inwardly from the specimen tip. The temperature was measured by a K-type thermocouple with a wire diameter of 80 μm, spot-welded at the tip of the specimen. The variation of temperature versus time was collected with data acquisition system developed in-house, using National Instrument® hardware and Labview® software. For experiments in N 2 or Ar, the gas was introduced after a primary (4.10 -2 mbar) and secondary (5.10 - 5 mbar) vacuum respectively. The oxygen and water vapour contents in the inert gas were very low (< 3 ppm, see Table 2). Several purges were performed before each test. During the test, the chamber was put under slight overpressure (around 1.2 bar) and the gas was stirred using a ventilator [START_REF] Salem | endommagement par fatigue thermique des moules de fonderie sous pression d'aluminium: effet de l'interaction avec l'aluminisation et l'oxydation[END_REF]. The thermal cycle imposed at the edge of the specimen varied between a minimum temperature (T min ) of 100 °C and a maximum temperature (T max ) of 650 °C. The prescribed T max was achieved by controlling the power of the induction generator, using a PID regulator (Eurotherm®). The heating time (ht) was then adjusted to 1.25 s, resulting in a mean heating rate of about 420 °C s -1 . The cooling time (ct) was set to 7.4 s, so as to achieve T min .

Microstructural characterisation

The TF experiments were regularly interrupted to observe the evolution of the surface damage (especially the oxidation and micro-crack initiation) by means of optical microscopy and SEM, in order to establish the evolution of both the distance between the micro-cracks (d) and the crack depth (a) as a function of the number of TF cycles (N). The micro-crack initiation was determined by direct observation of the specimen surface, while the lifetime for crack growth was given by the crack propagation curves (a = f(N)). TF crack lifetime was divided in two stages: an initiation stage that involves the growth of superficial small cracks called "micro-cracks" (with a depth smaller than about 50 µm), and a propagation stage that involves the growth of the steel macro-cracks reaching an "engineering" size (up to 0.5 mm in depth).

In order to identify the oxide and intermetallic phases, XRD analyses were performed from the outer surface of TF coated specimens before and after TF cycling. A Panalytical X'Pert Pro diffractometer was used, with monochromatic CoKα radiation (λ = 1.78 Ǻ) and parallel beam optics (better suited for thin film characterization). The parameters used were an angular step size of 0.01°an integration time of 5 s, a voltage of 40 kV and a current of 40 mA. In order to reduce the penetration depth and thus the relative peak intensity of iron, incident angle (ω) was fixed to 5.

The cracking mechanisms and the microstructure of the oxide/intermetallic/steel interfaces were investigated by post-mortem metallographic observations and analyses. At the end of the TF test, each specimen was cut and several cross-sectional samples were prepared (different cross section planes). Samples were first mechanically polished and then mirror finished using colloidal silica. Some samples were electrochemically etched using a 10% chromic acid solution and setting parameters (voltage of 5 V, intensity of 0,8 A and etching time of 20 s). This attack aims to reveal the initial austenitic structure of the steel and the oxide microstructure, by dissolving compounds containing chromium (as the chromium carbides and spinel oxides). A FEI Nova NanoSEM 450 Scanning Electron Microscope was used to observe the interface cross-sections. The chemical composition of coated specimens was analysed by Energy Dispersive Spectroscopy (EDS), using an acceleration voltage of 25 kV before the TF tests (Fe/Al binary system), and of 15 kV after the TF tests (O/Cr/Fe/Al system).

Coating process

Steel specimens were coated using a CVD process called "AlCl 3 Pack Cementation" (PC), to form Fe-Al intermetallic phases in a large temperature range (500-1000 °C) [START_REF] Mevrel | Pack cementation processes[END_REF][START_REF] Xiang | Relationship between pack chemistry and aluminide coating formation for low-temperature aluminisation of alloy steels[END_REF][START_REF] Rohr | Development of novel protective high temperature coatings on heat exchanger steels and their corrosion resistance in simulated coal firing environment[END_REF]. In order to avoid the thermal softening and keep the initial microstructure and hardness of the steel, the coating temperature was selected to be below the second tempering temperature of the steel (605 °C). The TF specimens were buried in a powder mixture (8Al-2AlCl 3 -90Al 2 O 3 ) consisting of three components: Al as the metallic source (element to be deposited on the surface), AlCl 3 as halide salt activator, and Al 2 O 3 as inert filler preventing powder sintering. The average particle size of the powder mixture was about 70 μm. The sealed pack was loaded into a furnace, and cementation was performed for 8 h at 500 °C under an inert atmosphere (N 2 with 5% H 2 ). The furnace, controlled by a programmable temperature controller, was heated at a rate of 16 °C min -1 to the required temperature and cooled to room temperature at the same rate [START_REF] Salem | endommagement par fatigue thermique des moules de fonderie sous pression d'aluminium: effet de l'interaction avec l'aluminisation et l'oxydation[END_REF]. As shown by SEM observations with a Back Scattered Electron (BSE) detector, a compact and homogeneous coating of less than 5 μm was formed with a columnar grain structure (Fig. 4a). EDS and XRD analyses revealed that the coating consisted mainly of Fe 2 Al 5 (Fig. 4b andc).

Results

Superficial damage 3.1.1. Virgin steel specimens

In air, a uniform and homogeneous oxide layer forms on the surface of the steel specimens in the early hundreds of TF cycles. Under thermal cycling, the oxide layer begins to crack within the first hundred cycles (600 cycles). Then, a homogeneous and uniaxial network of parallel and regular micro-cracks (uniaxial heat-checking) forms on the surface of the steel from 1500 cycles, with locally a few spalled areas (Fig. 5a). These micro-cracks are oriented perpendicular to the orthoradial axis of the specimen (i.e. to the circumferential stress σ θθ ). They propagate through the oxide layer in depth towards the steel and along "zz" direction, reaching the lateral faces of the disc crossing the specimen edge. Beyond 5000 cycles, some micro-cracks (called "secondary") get closed in outer surface and disappear (Fig. 5b). Therefore, the mean apparent inter-crack distance d tends to increase with the number of TF cycles N, as shown in Fig. 6.

In cross-section, it can be observed that the oxide layer formed in air on the virgin steel after 30,000 TF cycles is thick, compact and adherent to the steel. In addition, the oxide scale is multi-layered with inner Crrich and outer Cr-poor oxides. Fig. 7 also shows that the micro-cracks propagate from the superficial "heat-checking" into the base steel through the oxide scale. All micro-cracks are filled by inner Cr-poor oxide. The propagation of secondary micro-cracks seems to be limited in the oxide scale. Few micro-cracks propagate in depth, marking thus the initiation phase of steel cracking. The cracks propagating into the steel follow a straight path perpendicular to the "θθ" axis (Fig. 7a). The crack surfaces are oxidized, as the outer surface of the steel, but with decreased thickness towards the crack tip (Fig. 7a).

When the TF test is performed in nitrogen, both the kinetics and the morphology of the oxides are modified. The oxidation is logically reduced since the oxygen partial pressure in the test atmosphere is lowered. Furthermore, a thin, laminated and porous oxide scale is observed, which shows a strong tendency to spalling during the early stages of its formation (Fig. 8a). Even if the acicular structure of the superficial oxide differs from that observed in air, the same layers are found. No evidence of heat-checking network is observed and cracking is strictly limited to few areas where the oxide layer breaks markedly (Fig. 8b). In cross-section, most of the cracks appear thin, with a slightly opened crack tip, and are filled with oxide up to a very shallow depth, indicating that a few ppm of oxygen is sufficient to enable a crack propagation mechanism assisted by oxidation. Although the propagation paths are meandering (Fig. 9), it remains perpendicular to the "θθ" axis at macroscopic scale.

Coated steel specimens

In air, the surface of the coated specimen is first damaged by a uniaxial "heat-checking" (Fig. 10a), which is quite similar to that the one observed on virgin steel tested in air (Fig. 5a). However, it should be noted that the micro-crack network appears distinctly later, i.e. after 2500 cycles instead of 600 cycles for the virgin steel tested in air. Then, the oxidized areas (appearing in white on the micrographs) emerge through the micro-cracks and tend to completely fill the cracks at higher number of TF cycles N (Fig. 10b). Some micro-cracks open while others seem completely closed. As previously observed for virgin steel in air, the mean inter-crack distance d increases by increasing N (Fig. 11). As can be seen in cross-section, most of the micro-cracks propagate in the steel, to a depth of a few tens of micrometres (Fig. 12a andb). These cracks, which are entirely filled with oxide, are V-shaped. It is also observed that the external surface is particularly rough, due to oxide protrusions through the micro-cracks (Fig. 12b). Some microcracks stop to a few micrometres in depth and seem to be closed, thus explaining the increase of d with N.

Cross sectional observations reveal that the microstructure of the Fe/Al intermetallic coating changes under the effect of thermal and thermo-mechanical loadings. The evolution of the interface and the intermetallic coating (that is transformed into a multilayer) is particularly well highlighted by EDS analyses (Fig. 13) and the microstructure (Fig. 14a). XRD analyses (Fig. 14b) still reveal the presence of Under nitrogen, the superficial "heat-checking" network is heterogonous with shorter and less opened micro-cracks, leading to a lower apparent surface density than in air (Fig. 15a). Oxide still grows inside the micro-cracks (appearing in white in Fig. 15b), but in smaller quantities than in air. Cross-sectional observations reveal that these micro-cracks propagate in depth following a sinuous path (Fig. 16a), and that the micro-crack tips are less opened than in air (Fig. 16b). In addition, some thin and short cracks initiate into the steel and propagate towards the surface and the bulk along planes perpendicular to the Fig. 10. SEM (BSE) micrographs of the surface of the coated specimen tested in air after 5000 TF cycles (a), and detailed view of the homogenous uniaxial "heatchecking" network after 30,000 TF cycles (b). "θθ" axis. It must be emphasized that this was never observed for TF tests carried out in air.

TF life and crack propagation mechanisms

The lifetime data corresponding to the required numbers of TF cycles for micro-crack initiation and macro-crack growth (using two different criteria of crack depth, as described in Section 2.3) are reported in Table 3. It indicates that in air, the micro-crack initiation is strongly delayed (by almost a factor 4) in the intermetallic layer compared to the oxide scale. However, in nitrogen, the crack initiation starts later in the intermetallic layer, with a difference of 2500 cycles compared to the same test in air.

The propagation curves plotted in Fig. 17a andb represent the evolution of the depth of the main crack for the virgin and coated steels, respectively. In all cases, the TF crack growth follows a classical sigmoidal trend, with acceleration just after initiation followed by a stabilization of the propagation rate around of 0.35 µm/cycle, and then arrest for a crack depth greater than 2 mm. On the other hand, the propagation rate of cracks between 0.5 and 2 mm seems to be rather independent from the test atmospheres. The depth of macro-cracks varies depending on the test conditions and the number of macro-cracks propagating in each specimen. In fact, for the virgin steel tested in air and coated specimens tested in both atmospheres, two diametrically opposed macro-cracks (main cracks) propagate in the disc. After 20,000 TF cycles, other secondary macro-cracks propagate in each specimen. However, both the main and secondary macro-cracks started simultaneously in the virgin steel tested in nitrogen.

Fractographic analyses show that crack surfaces are oxidized and rough. Different shapes are observed for short cracks (a < 0.5 mm) propagating in the edge specimen and long cracks (a > 0.5 mm) propagating in the disc. The front of short cracks is longer in surface than in depth (Fig. 18a), whereas the long crack front is deeper inside the disc (Fig. 18b). On the non-oxidized areas of the steel surface, crack propagation follows a "zig-zag" path along symmetrical planes inclined at 45°to the loading axis (Fig. 19a). At the macroscopic scale, cracks propagate in depth in mode I, along a plane perpendicular to the circumferential "θθ" axis. The crack propagation usually follows a transgranular mechanism (Fig. 19b), but evidence of inter-granular propagation can also be observed occasionally (Fig. 19c).

Discussion

During the TF test, the specimen is self-constrained since the "free" thermal expansion of the surface is hindered by the bulk, which is at lower temperature. Therefore, each element of the specimen undergoes cyclic strains and associated stresses resulting from the thermal transients. The external surface of the steel is thus subjected to a thermomechanical stress-strain hysteresis loop, and goes into compression during attempted expansion (heating phase) and into tension during attempted contraction (cooling phase) [START_REF] Qayyum | Numerical simulation of thermal fatigue behavior in a cracked disc of AISI H-11 tool steel[END_REF][START_REF] Salem | Fatigue life modelling of hot forming tools[END_REF]. Since the oxide or intermetallic layer is strongly bonded to the substrate, it undergoes the same thermo-mechanical loadings imposed to the steel surface. Other stresses (related to geometric factors, material heterogeneity or layer growth, etc.) are superposed to thermal stresses induced by the substrate [START_REF] Schütze | Protective oxide scales and their breakdown[END_REF]. The local mechanical incompatibility, generated by the thermal mismatch between the substrate and the multilayered coating of different coefficients of thermal expansion [START_REF] Spera | What is thermal fatigue? Thermal fatigue of materials and components[END_REF][START_REF] Roux | Role of heat-flux density and mechanical loading on the microscopic heat-checking of high temperature tool steels under thermal fatigue experiments[END_REF], significantly impacts the thermo-mechanical loading on the surface. The sub-surface also undergoes metallurgical modifications, along the thickness and in the composition of the layer. In addition, the thermo-mechanical loading induces cyclic softening of the steel, phenomenon that has been widely reported for tool steels [START_REF] Miquel | Heat-checking of hot work tool steels[END_REF][START_REF] Klobčar | Thermal fatigue of materials for die-casting tooling[END_REF][START_REF] Majumdar | Formation of Al 2 O 3 /Fe-Al layers on SS 316 surface by pack aluminizing and heat treatment[END_REF][START_REF] Velay | Cyclic behavior modeling of a tempered martensitic hot work tool steel[END_REF][START_REF] Grüning | Cyclic stress-strain behavior and damage of tool steel AISI H11 under isothermal and thermal fatigue conditions[END_REF]. Whether or not the specimen is coated, the damage usually initiates first in the superficial layer (composed of the intermetallic layer and/or the steel surface). After a critical number of cycles (depending on the amplitude and rate of the thermal cycle and environmental conditions), cracks initiate and propagate under mode I along the circumferential "θθ" axis [START_REF] Qayyum | Numerical simulation of thermal fatigue behavior in a cracked disc of AISI H-11 tool steel[END_REF].

Various damage mechanisms are involved depending on the nature of the specimen (virgin or coated steel) and the TF test environment (air or nitrogen), as detailed below.

Damage mechanisms

It was shown that the oxide layer formed on the virgin steel is always composed, of superficial hematite (Fe 2 O 3 ), magnetite (Fe 3 O 4 ) and inner spinel (Fe,Cr) 3 O 4 [START_REF] Salem | endommagement par fatigue thermique des moules de fonderie sous pression d'aluminium: effet de l'interaction avec l'aluminisation et l'oxydation[END_REF][START_REF] Sundqvist | Effects of liquid aluminium on hot-work tool steel[END_REF][START_REF] Miquel | Heat-checking of hot work tool steels[END_REF][START_REF] Medjedoub | Experimental conditions and environment effects on thermal fatigue damage accumulation and life of die-casting steel X38CrMoV5 (AISI H11)[END_REF], whatever the test atmosphere. The thickness proportions of these different oxide layers were estimated at 20, 40 and 40%, respectively. In air, the oxide was compact and homogeneous, and its surface was damaged very early by a "heatchecking" network. This mode of cracking, by initiation and propagation of through-scale micro-cracks perpendicular to the steel/ coating interface, is generally attributed to the hard coating failure under tension [START_REF] Schütze | Deformation and cracking behavior of protective oxide scales on heatresistant steels under tensile strain[END_REF][START_REF] Nagl | The mechanical failure of oxide scales under tensile or compressive load[END_REF][START_REF] Strawbridge | Mechanical failure of thin brittle coatings[END_REF][START_REF] Teixeira | Residual stress and cracking in thin PVD coatings[END_REF]. Schütze et al [START_REF] Schütze | Development of a comprehensive oxide scale failure diagram[END_REF] proposed an oxide scale failure mechanism depending on applied strain, scale thickness and defect sizes. Accordingly, the first through-scale cracks developed from pre-existing defects in regions with high stress concentration, when critical thickness and tensile strain are reached. At low temperature (corresponding in our TF tests to the maximum stress σ max ), the accommodation of tensile stress was achieved by oxide cracking when this stress exceeded, locally, the elastic limit of the oxide. The brittle breakdown of the oxide into a finite number of segments led in surface to the formation of homogenous and dense "heat-checking" network, propagating on the surface along "zz" axis and in depth towards the steel. The growth of micro-cracks through the inner oxide reduced the oxide scale cross section (distance between the free surface and the steel). Oxidation was thus accelerated through the micro-cracks, which represent preferential diffusion paths. The outward migration of iron cations, resulting in the formation of magnetite on the micro-crack surfaces, partially filled the micro-cracks. The inward diffusion of oxygen caused the penetration of oxide into the steel around the crack tip, leading to the formation of magnetite, and then spinel when the micro-crack propagated beyond a Cr-depleted zone (Fig. 7b). In fact, some micro-cracks stopped because screened by the deeper neighbouring micro-cracks, which locally deformed the orthoradial stress fields at the blunted tips. These arrested micro-cracks disappeared from the surface because covered by hematite, and seemed completely filled by magnetite in depth. Indeed, the internal growth of the oxide reduced the distance between the free surface and the steel surface normalizing thus the oxidation kinetics of the steel in the cracked and uncracked oxide zones. The inside protrusion of oxide was prevented or stopped, since the preferential diffusion was decelerated. Therefore, the healing oxide partially restored the protective action of the layer, resulting in a decrease in micro-crack density (Fig. 6). Meanwhile, in-depth propagating micro-cracks amplified local stress fields and, in conjunction with local thermo-mechanical softening and oxidation, caused nucleation and propagation of macro-cracks (especially on the first hundred micrometres). Previous investigations [START_REF] Medjedoub | Experimental conditions and environment effects on thermal fatigue damage accumulation and life of die-casting steel X38CrMoV5 (AISI H11)[END_REF] suggested a link between superficial micro-cracks and macro-cracks propagating in depth. When the crack propagated in the steel, its fresh surfaces in contact with air oxidized in the same way as the free surface of the steel. A compact multilayer oxide grew on each crack lip, with a decreasing thickness towards the crack tip where the oxidation kinetics was less important than at the top of the crack (decrease in depth of oxidation temperature and time). Under nitrogen, the oxide layer formed after a large number of cycles was thinner and more porous than in air. The deformation of the superficial layer, caused by the thermo-mechanical stresses, led to development of surface wrinkling, but without the formation of a "heatchecking" network (even after 10,000 cycles). This should be mainly related to the mechanical behaviour of the porous oxide layer and its weak bound with substrate. Accordingly, oxide layer accommodated the thermo-mechanical loadings by deforming rather than cracking, unlike the compact oxide or intermetallic layer.

The mechanism of surface micro-crack formation observed for the coated steel was close to that found for the compact oxide grown in air, with however some changes in the damage sequence. First, the microstructure of the initial intermetallic coating evolved to a complex multilayered intermetallic and oxide, due to the thermally activated diffusion of iron, aluminium and oxygen at high temperature. The new phases identified in this multilayer were in agreement with the Fe-Al equilibrium diagram and the literature [START_REF] Potesser | The characterization of the intermetallic Fe-Al layer of steel-aluminum weldings[END_REF][START_REF] Deqing | Phase evolution of an aluminized steel by oxidation treatment[END_REF][START_REF] Novák | On the formation of intermetallics in Fe-Al system-An in situ XRD study[END_REF][START_REF] Zamanzade | A review on the properties of iron aluminide intermetallics[END_REF]. Then, the micro-cracks propagated through the modified intermetallic layer both toward the steel (in depth) and along "zz" axis (on the surface), as observed in the compact oxide scale formed on virgin specimens tested in air. However, a difference in the morphology of micro-cracks was observed depending on the test atmosphere (air or nitrogen). In air, greater oxidation increased the propagation rate of the superficial micro-cracks along the "zz" axis to achieve a uniform and dense heat-checking network. In both atmospheres, the outside protrusions of the iron oxides through the micro-cracks caused the uplift of the intermetallic layer without inducing breaking, testifying to the ductility of some of iron aluminium intermetallic phases, especially at high temperature [START_REF] Novák | On the formation of intermetallics in Fe-Al system-An in situ XRD study[END_REF].

Effect of the heat-checking network and oxidation on the TF lifetime

In air, both virgin and coated TF specimens showed premature cracking with the growth of two main macro-cracks in the bulk steel, while under nitrogen the virgin steel simultaneous exhibited multicracking at higher number of cycles. The premature cracking of steel could be initiated, by the concentration of stress fields at the tip of the micro-cracks, from the "heat-checking" network that develops on the oxide growing in air and in the intermetallic coating.

Moreover, the intermetallic coating increased the TF lifespan by about 30% (for a crack length of 1 mm), compared to the virgin steel tested in air. This difference would come from the characteristics of the superficial layer, and in particular the intrinsic properties of Fe-Al intermetallics showing a better TF durability than the Fe-Cr oxides. This was further corroborated by measurements of the "heat-checking" network density (inversely proportional to the distance between the micro-cracks), which was twice as small on the coated specimens and even less by increasing the number of cycles. Therefore, a better resistance to TF "heat-checking" of the superficial layer would increase the lifetime of the steel. However, when the surface was free off a hard coating (such as an oxide or intermetallic layer), the lifetime was much longer because of the high ductility of the steel. Indeed, for the virgin specimen tested under nitrogen, the surface element could be considered composed by the steel only (by neglecting the mechanical role of the oxide, here as a porous and unbound layer).

On the other hand, the TF lifetime of the coated specimen was slightly increased under nitrogen, due to the slower growth of the surface "heat-checking" network. This clearly shows the predominant role of oxidation in both the development of the surface "heatchecking" network and the propagation of micro-cracks through the intermetallic coating. However, oxidation seemed less significant in the premature cracking of the coated steel, with a difference of only 12% on the number of cycles to initiation between TF tests in air and nitrogen.

For all TF tests, the crack propagation curves showed a fairly similar kinetics of the main macro-cracks (for a between 0.5 and 2 mm). The effect of oxidation on the crack propagation rate in the base steel was thus less obvious, although it was significant for propagation mechanisms of short cracks. Oxidation should however influence the crack propagation kinetics, especially at the top of the specimen exposed to maximum temperatures. Indeed, it was shown that the cracks were more opened in air than under nitrogen, suggesting a higher crack propagation rate in air. Unfortunately, this could not be highlighted for short cracks (a < 0.5 mm) in the current propagation curves, due to the limitations of the measurement method (projections along "rr" axis of the crack trace appearing on the lateral faces). The criterion of crack sizing in the disc part should be different than in the specimen edge, where the crack shape was elliptical with two axes (Fig. 18a). In fact, due to the semi-cylindrical geometry of the specimen edge, crack propagation was faster on the surface than in depth. Once the cracks reached the lateral faces of the specimen (for a > 0.5 mm), the crack propagation rates on surface and in depth became uniform and the measured crack (a s ) were similar in both lateral sides (Fig. 18b). However, the crack front remained slightly curved inside the specimen where the crack propagated under plane strain, while it propagated under plane stress on the surface. The temperature drops along the depth where the oxidation became less important and its effects on the crack propagation were negligible. In both atmospheres, the propagation kinetics were thus comparable for a between 0.5 and 2 mm, as previously seen in propagation curves. Nevertheless, it was clearly demonstrated that both the crack arrest and the total number of macrocracks propagating in the specimen depended on the test conditions. Thermo-mechanical simulations showed that by increasing inward crack length, the J-integral drops (a > 2 mm) because of the attenuation of thermal gradients and thermal stresses in depth [START_REF] Salem | endommagement par fatigue thermique des moules de fonderie sous pression d'aluminium: effet de l'interaction avec l'aluminisation et l'oxydation[END_REF][START_REF] Qayyum | Numerical simulation of thermal fatigue behavior in a cracked disc of AISI H-11 tool steel[END_REF]. This explains the S-shape of the TF crack propagation curves. These simulations also led to investigate the interaction between the primary and secondary cracks, and showed that the J-integral decreases by increasing the number of cracks. This can explain the crack arrest at shallower depth for the virgin steel tested in nitrogen, where more than several cracks propagated simultaneously in the disc.

Tool damage in HPDC

In HPDC, the moulds are subjected to fast cyclic thermal transients in a particularly corrosive environment (contact with molten aluminium and air). Damage mechanisms are involved by a complex interaction between TF and corrosion (Fig. 1b-e), close to that observed from TF tests of intermetallic coated and virgin specimens in air. All observations of cracks in HPDC dies reveal such patterns [START_REF] Salem | endommagement par fatigue thermique des moules de fonderie sous pression d'aluminium: effet de l'interaction avec l'aluminisation et l'oxydation[END_REF][START_REF] Chen | The effect of temperature on soldering and the sequence of formation of the soldered layer during high pressure die casting of Al-11Si-3Cu alloy[END_REF][START_REF] Chen | Structures of intermetallic phases formed during immersion of H13 tool steel in an Al-11Si-3Cu die casting alloy melt[END_REF][START_REF] Persson | Failure modes in field-tested brass die casting dies[END_REF][START_REF] Srivastava | Computer modeling and prediction of thermal fatigue cracking in die-casting tooling[END_REF][START_REF] Muhič | Thermal fatigue cracking of diecasting dies[END_REF]. In die areas prone to aluminizing, TF damage should thus be controlled by: -the ability of the lubricant to remain on the die surface, in order to fully play its protective role as a chemical and physical barrier against oxidation, aluminizing (Fig. 1a) and wear; -aluminizing kinetics that is mainly controlled by the die contact temperature at the location where aluminium impinges and attacks the die, and by shot number (contact duration); -the intermetallic cracking and the nucleation of the crack in the die surface.

One question remains regarding a potential critical thickness and the composition of the intermetallic layer, before any cracking might occur. This could be investigated through TF tests on specimens with different coating thicknesses, or by providing an aluminium-rich test environment allowing a continuous growth of the intermetallic layer (in the same way as the oxide increases on a virgin specimen during a test in air). The growth of both the oxide and the intermetallic layers is controlled by diffusion, which is a thermally activated phenomenon. At 650 °C, the rate constants (k p ) for the oxidation of X38CrMoV5 (AISI H11) in air and its aluminizing in a liquid AlSi9Cu3 (A380) bath are equal to 9.14 10 -2 and 2.0 µm s -0.5 , respectively [START_REF] Salem | endommagement par fatigue thermique des moules de fonderie sous pression d'aluminium: effet de l'interaction avec l'aluminisation et l'oxydation[END_REF]. Besides, it was shown that the intermetallic layer acts as a chemical barrier to oxidation (and vice versa), and that the oxide promotes a denser network of micro-cracks, which leads to premature cracking of the steel. As the aluminizing kinetics is ten times higher than the oxidation rate (Fig. 1b andc), it would therefore be detrimental to the dies even if the intermetallic layer showed better FT behaviour. Using HPDC mould under an inert atmosphere being hardly conceivable for practical reasons, hard coating could be a good solution to protect the mould surface, as evidenced by many studies on this topic [START_REF] Mitterer | Application of hard coatings in aluminium die castingsoldering, erosion and thermal fatigue behaviour[END_REF][START_REF] Navinšek | Industrial applications of CrN (PVD) coatings, deposited at high and low temperatures[END_REF][START_REF] Starling | Thermal fatigue of hot work tool steel with hard coatings[END_REF][START_REF] Wang | A study of PVD coatings and die materials for extended die-casting die life[END_REF][START_REF] Gulizia | Performance evaluation of PVD coatings for high pressure die casting[END_REF][START_REF] Srivastava | A multilayer coating architecture to reduce heat checking of die surfaces[END_REF][START_REF] Peter | Study of protective coatings for aluminum die casting molds[END_REF]. To be effective, the coating should meet three criteria: be compatible with the substrate and be ductile enough to avert TF cracking, be chemically stable to prevent corrosion, and be hard enough to avoid damage by wear.

Conclusion

The effect of the interaction between aluminium and HPDC diesresulting in the formation of intermetallics -on the thermal fatigue (TF) behaviour of X38CrMoV5 (AISI H11) hot work tool steel is investigated. A new experimental approach is proposed to study the interaction between these different damage mechanisms, using virgin and coated specimens by CVD pack-cementation at 500 °C.

The microstructure of the superficial layer undergoes chemical and microstructural evolution during TF cycles, which activates the diffusion of chemical elements (O, Al, Fe, Cr) through the different interfaces. The oxide layer which forms on virgin steel specimens is much thicker, homogenous and compact in air than in nitrogen. On coated specimens, the intermetallic layer, initially composed of Fe 2 Al 5 , evolves to a complex multi-phased layer consisting of new intermetallic phases (Fe 3 Al, FeAl, Fe(Al), FeAl 2 Fe 2 Al 5 and FeAl 3 ) and oxide compounds ( The mechanisms of the crack initiation depend on oxidation and the surface layer properties. In air, a micro-crack network forms on the surface of both virgin and coated specimens after a few hundred to thousand TF cycles. The heat-checking network is less dense on the coated specimen, while it is non-existent on the virgin steel since the oxide layer is too thin and porous to initiate micro-cracks. Preferential oxidation by short-circuit diffusion of the base steel leads to the propagation of the heat-checking network in the steel. In addition, the micro-crack propagation through the intermetallic layer is assisted by oxidation. The intermetallic coating delays the steel cracking and increases the TF life of the steel, due to its better TF durability. Under nitrogen, oxidation is greatly reduced and micro-crack initiation is delayed on both virgin and coated specimens. It is clearly shown that the microscopic "heat-checking" is the main factor determining the premature cracking of the steel and that the crack propagation mechanisms are assisted by oxidation at high temperature.

Fig. 1 .

 1 Fig. 1. SEM observations on X38CrMoV5 (AISI H11) pin of AlSi9Cu3 high pressure die casting mould: undamaged (a), oxidized (b) and aluminising (c) areas in cross section; multi-axial (d) and uniaxial (e) "heat-checking" network in surface [9].

Fig. 2 .

 2 Fig. 2. Geometry and dimensions of the thermal fatigue specimens.

Fig. 3 .

 3 Fig. 3. (a) General view of the thermal fatigue rig, (b) detailed schema of heating and specimen holder devices.

Fig. 4 .

 4 Fig. 4. Coated X38CrMoV5 (AISI H11) steel by AlCl 3 Pack Cementation at 500 °C for 8 h: cross sectional SEM (BSE) micrograph (a), EDS spectrum performed at 25 kV (b), and XRD pattern from the outer surface, using Co-Kα radiation with ω = 5°, V = 40 kV and I = 40 mA (c).

Fig. 5 .

 5 Fig. 5. SEM (BSE) micrographs of the surface of the specimen tested in air after 1500 TF cycles (a), and detailed view of the homogenous uniaxial "heat-checking" network after 5000 TF cycles (b).

Fig. 6 .

 6 Fig. 6. Evolution versus number of cycles of the mean distance between the micro-cracks (d) formed on the oxidized surface of the virgin steel specimen tested in air.

Fig. 7 .

 7 Fig. 7. SEM (BSE) micrographs in cross section of the virgin specimen tested in air after 30,000 TF cycles, before (a) and after electrolytic etching in chromic acid (b).

Fig. 8 .

 8 Fig. 8. SEM (BSE) micrographs of the outer surface of the specimen tested in nitrogen, showing a porous and thin oxide layer after 10,000 TF cycles (a), and detailed view of the isolated cracks that develop after 25,000 TF cycles (b).

Fig. 9 .

 9 Fig. 9. SEM (BSE) micrographs in cross section of the virgin specimen tested in nitrogen after 40,000 TF cycles (a) and detailed view showing the superficial oxide layer (b).

  Fe 2 Al 5 , but also Fe(Al), FeAl and traces of Fe 3 Al, FeAl 3 and FeAl 2 . The interdiffusion of chemical species on the coating free surface and through the micro-cracks (into the steel) induces the formation of Al 2 O 3 , and Fe 2 O 3 , Fe 3 O 4 and (Fe,Cr) 3 O 4 , respectively. After 30,000 TF cycles, the total thickness of the modified coating and steel surface can reach 10 µm.

Fig. 11 .

 11 Fig. 11. Evolution versus number of cycles of the mean distance (d) between the micro-cracks formed on the coated specimen tested in air.

Fig. 12 .

 12 Fig. 12. SEM (BSE) micrographs in cross-section of the coated specimen tested in air after 30,000 TF cycles (a), and enlarged view of the specimen surface (b).

Fig. 13 .

 13 Fig. 13. SEM (BSE) micrograph in cross section of a micro-crack formed on the coated specimen tested in air after 30,000 cycles, and its associated ESD elementary maps showing the distribution of Al, Fe, O and Cr (40 frames, V = 15 kV, time constant = 13 µs).

Fig. 14 .

 14 Fig. 14. Cross-sectional SEM (BSE) micrograph (a) and XRD pattern (b) of the coated specimen tested in air after 30,000 TF cycles. The XRD analysis was performed from the outer surface of the specimen (using Co-Kα radiation with ω = 5°, V = 40 kV and I = 40 mA).

Fig. 15 .

 15 Fig. 15. SEM (BSE) micrographs of the surface of the coated specimen tested under nitrogen after 5000 TF cycles (a), and detailed view of the non-uniform uniaxial "heat-checking" network after 30,000 TF cycles (b).

Fig. 16 .

 16 Fig.[START_REF] Chen | Formation and progression of die soldering during high pressure die casting[END_REF]. SEM (BSE) micrographs in cross-section of the coated specimen tested under nitrogen after 30,000 TF cycles (a), and enlarged view of the micro-cracks (b).

Fig. 17 .

 17 Fig. 17. Propagation curves of the main cracks in air and nitrogen, respectively for virgin (a) and coated (b) steel specimens.

Fig. 18 .

 18 Fig. 18. Post-mortem SEM observations showing the elliptical shape of the short cracks propagating in the specimen edge (a), and the curved front of the crack propagating in the disc (b) (a s and a i are the size of the crack, respectively on surface and in-depth).

Fig. 19 .

 19 Fig.19. Different mechanisms of crack propagation observed in the specimen tested under inert atmosphere at T max = 650 °C: "zig-zag" propagation path observed on the non-oxidized areas of the steel surface (a); trans-granular (b) and inter-granular in depth propagation (c) of a short crack.

Fe 2 O 3 ,

 23 Fe 3 O 4 , (Fe,Cr) 3 O 4 and Al 2 O 3 ).

Table 1

 1 Chemical composition of the X38CrMoV5 (AISI H11) tool steel.

	Element	C	Cr	Mo	V	Si	Mn	S	P	Fe
	Weight%	0.36	5.06	1.25	0.49	0.35	0.36	0.0011	0.006	bal.

Table 2

 2 Impurity content in the argon and nitrogen gases used for the TF tests.

		H 2 O	O 2	C n H n
	Argon (αAr)	< 3 ppm	< 2 ppm	< 0.5 ppm
	Nitrogen (αN 2 )	< 3 ppm	< 2 ppm	< 0.5 ppm

Table 3

 3 Comparison of the lifetime of virgin and coated X38CrMoV5 (AISI H11) steel under thermal fatigue at T max = 650 °C, in air and nitrogen.

	Specimen	Test	Number of	Number of cycles	Number of cycles
		atmosphere	cycles for	(crack	(crack
			micro-	depth = 1 mm)	depth = 2 mm)
			crack
			initiation
	Virgin steel	Air	600	11,000	13,500
		N 2	-	25,000	28,000
	Coated steel Air	2500	14,000	18,500
		N 2	5000	16,000	18,500