Raquel Rodriguez-Lamas

Resistive Switching in LaMnO 3+δ based heterostructures for Valence Change Memories

Keywords: , classification and terminology

This thesis focuses on the study of perovskite oxides for resistive switching based nonvolatile memories (NVM's), their integration in devices and the evaluation of their electrical characteristics. Among the candidates for NVM's, valence change memories (VCM) have been chosen as the subject of this study. VCM's, presenting bipolar resistive switching and based on the drift of oxygen vacancies in an oxide matrix, are of particular interest since these memories can present interface or filamentary-type resistive switching depending on their materials and operation characteristics. In this work, we studied LaMnO 3+δ as a potential candidate for VCM's, based on its flexible oxygen stoichiometry related to the valence changes in the Mn transition metal ion, which maintains the electro-neutrality of the structure. First, the conditions for the growth of LaMnO 3+δ perovskite thin films by pulsed injection metal organic chemical vapor deposition were optimized. The window of parameters that allow the growth of LaMnO 3+ δ (LMO) with different oxygen content (δ) and for a La/Mn ratio close to one were defined. In addition, studies performed on the LMO films grown on different single crystal substrates allowed the measurement of the oxygen mass transport properties of epitaxial LMO films. The very high oxygen diffusion coefficients obtained at intermediate temperatures (500 ºC) and their increase with the presence of structural defects, open the possibility of using LMO as functional oxide in other nanoionic devices, such as for solid oxide fuel cells cathodes. The successful integration of LMO films on platinised silicon substrates was fundamental for the resistive switching studies in devices in top-bottom configuration. We developed several strategies permitting both, the deposition on Pt, and the control of the morpho-structural characteristics of the LMO film. The polycrystalline LMO films grown on platinised silicon were then used for device fabrication in top-bottom configuration and for standard electrical characterisation. Furthermore, we designed and fabricated devices for operando advanced characterisation using Synchrotron-based X-ray photoelectron spectroscopy (HAXPES) combined with resistive switching measurements. TiN was selected as the top electrode of interest in this thesis and TiN/LMO/Pt devices were used for the majority of the electrical characterization studies. The devices were divided in two main families according to the resistance of the pristine state. An extensive study of the activation of the devices and its influence in the device performance was carried out. The best performing devices were then iv evaluated in terms of cycle to cycle variability, cycling sense and opening window as a function of the operation parameters. Finally, the relevance of using an "oxygen active" electrode (such as TiN) was put in evidence, through comparison with tests performed on Au/LMO/Pt devices prepared from the same LMO film.

Truth be told, the path that leads to the obtention of a PhD title is often arduous and not always gratifying. Despite that inevitable truth, I've chosen today to look at this path with the fondness that our past deserves, for it is as dark or bright as we decide to remember it and I'd rather leave a permanent reminder of all the good that happened in the last four years.

First of all I'd like to thank the jury members, for accepting to evaluate my work, without them I wouldn't have fulfilled this task. Thanks for your time, remarks and for the great discussion on the day of the defence.

The work collected in this manuscript was the fruit of a joint effort. It is my belief that research is not to be done alone, and like a big machine each engine does a part of the work. Hence, in these lines I try to acknowledge the efforts of everyone that contributed to the development of the thesis. I would like to thank Michel Boudard, the Director of the thesis, for all the help he offered, from minute one: with the inscription, to face a different bureaucratic system and all of his assistance; with the particularly hard beginning, where I found myself in an unknown environment following new protocols and feeling like I was not working in any particular direction. Thanks for offering your expertise on diffraction, for all the hours of scientific discussion and all the literature recommendations that helped to find a bit of context for the field I was beginning to work with. I need to thank Carmen Jiménez for her direction, the always clear ideas and how she kept us all in track to move forward in a Project in which results were continuously branching out to very promising new ideas and possibilities. I need to thank her for her humanity, and her way to help us all and make us feel "home" at the laboratory, no matter from how far we come. I need to thank her for how she has always listened to all of us, and for everything she has taught me. I do not think I would be as autonomous and independent in a laboratory if not for her. You have been a true role model. I want to thank Mónica Burriel for her amazingness, you´ve been a great team leader, a great director, and I hope I have learnt as much from you as I could in these four years. It is truly a blessing to work with you, and I need to thank you for all the time that you have xxiv List of Tables dedicated to me and the Project, for all the hours of discussion and experiment planning, for all the motivation you have transmitted to me in these years, and for being a great example of researcher.

There is, of course the Post Doc working in ALPS before I arrived, she was the battery bank of the project and I could not believe how many things a single person could do in a day or a week when I met you, but then you taught me how to do that myself, I hope I helped with the workload. Dolors, I really need to thank you for everything you taught me and for always having the good questions to ask (which were of course the ones you would not expect or know how to answer on the spot), for your energy and dynamism, and for all the knowledge you transferred to me in the form of "Know how" or otherwise.

I want to thank all of the collaborators in the ALPS memories project because as I said at the beginning, no machine can work without all the pieces, our collaborators Eugenie Martinez, Olivier Renault and Nicolas Nuns from CEA leti, that provided us with the XPS and C-AFM expertise and equipment, and that carried out the campaigns in SOLEIL and Spring 8 synchrotron facilities.

I need to thank Benjamin Meunier, the Post Doc that took care of "the cream of the crop" of my samples, performing all the XPS and C-AFM characterisation in collaboration with the people form CEA, brought the very best devices to the synchrotron to analyse them properly, and for every other experimental contribution to the project, but at a personal level I'd like to thank him for the moral support he offered to me in very stressful times, and for the occasional chit-chat that made me feel understood at the time. And why not, also for being a great neighbour.

Thanks to Xavier Mescot and Quentin Rafhay from IMEP LaCH, it has been a pleasure to work with you, learn from you and have the opportunity to discuss science with you. I would like to thank Caroline Pirovano and Rose Noelle Vannier, for their support on the oxygen Exchange and Tof-SIMS section of the project, for the enriching discussions and for bearing with my sample pictures and indications of where to "drill the holes", that might have been clear for me but I bet they were not so simple to read. I would like to thank Etienne Pernot for the AFM tutorial in LMGP, and for trusting me with the machine, for being a very nice office partner and for the nice discussions about work and life. I want to thank Odette Chaix, I want to thank her extensively, for the infinite hours we have worked together, for teaching me so much about Raman Spectroscopy, for her kindness and good mood, and for bearing with my first attempts at speaking french when I was still blabbering nonsense. I want to thank her for being one of the people in LMGP that makes it feel like the best place to work in, for the camaraderie and the good vives she transmits to all List of Tables xxv of us. I want to thank Laetitia and Erinie, from that first TEM formation in the lab to the last minute of the thesis, for their availability and all the help they offered, thank you. I must, thank Hervé Roussel for all the time he dedicated to us, all the insight the gave me and his never-ending kindness.

I need to thank the LMGP people, every one of them make possible to work in the amazing environment this lab has, all of them are important and all of them should be thanked for it. This shall be done with a special mention to the administration people, who make everything possible, and a particular thanks to Michèle, J'ai vraiment apprécié ta patience, ta compagnie pendent les repas à la salle café, tu donnes de la vie au labo.

I want to thank also: Sarunas, for you were the first of the team I actually worked with/for, from the lab in Barcelona during my Master, even if I did not know it. It was always great to discuss with you in the mornings.

Klaasjan, there is no way I'll manage to write everything I am thankful for. I'll leave it at a great thank you for being the best team player I've known so far. For working for all the students that follow and for the team, your effort was very much appreciated. Thanks for being an even greater friend than co-worker. I will always treasure the memories we shared.

To the CDI, that appeared one day in the team, H. Jain, thanks for your contribution on the project, and for bearing with me and my hundred samples to measure in the 4 probe system.

Carlos, there goes another infinite thanks, I think we've learnt a lot together, and it was awesome, I wish you the best of lucks with the last year. Thanks not only for how much you helped me at the lab, but also for how much you helped outside. With friends like you one forgets that family is far away.

Alexander Stangl, I am very grateful to have had the chance to meet you, thanks for many interesting discussions, scientific and otherwise, and for the many hours you invested in the last article written before this manuscript was completed. You brough great vibes to the team and the lab. I need to thank my "not a husband" despite IKEA thinking so, I know that I've got a friend in you, and I "blame" you for me recovering my investment on creative activities. I cheer to your future Ragnar, and of course, I thank you for the simulation work I'd have never had the time to do. Khan If you're reading this I thank you for your everlasting good mood, it was very nice to have you around, even if we didn't work together.

To the people forming the LMGP family that accompanied me during this trip: The ones that left and the ones that are still there.

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Joao from the first smile in the staircase to the fastest farewell in Corona times, I have always known you had my back, you've been there for the good the better the bad and the worst, and everytime I knew I was going to make it, partially because of that inexplicable faith you have in me. You've been inspiring and I wish our scientific careers cross again, I know our lives will always be entangled, see you soon in Lisbon or anywhere.

Mirasbek, I miss you, thanks for sharing the "cat-parent" far away from kitten feeling with me, and for everything else.

Antalya, my admiration to you goes beyond words, you're a true fighter, and I am grateful to be counted among your friends, thanks for sharing so much with me during these years.

Morgane, on est arrivé à la même heure au LMGP le premier jour de contrat, et on gardera toutes les deux la mémoire de cet entretien au bureau d'Hubert. C'est un plaisir de partager avec toi, bureau, pause, science, plaintes, la vie. Merci de me faire reprendre les skis. A très bientôt ! Fanny parce que parler avec toi ça m'a toujours remonté la morale.

Adria, si te escribo un libro para agradecerte tu eterna actitud optimista y la amistad que me brindas, igual me llegan las páginas para expresarlo apropiadamente. Eres un rayo de sol. Y nos transmites a todos esa energía y ese positivismo. Gracias.

Cesar, Viet, Getnet, Sara, Evgenii, Carol, Max, Julius, Maxime, Alexandre, Guillaume, Octavio, Fidel, Damir, some of you where there at the beginning and some are still here by the end of this PhD journey of mine, I hope you each know I appreciated, really, having the chance to meet you and share time and food and wine and conversation and films and fun with you.

Pierre ça me manquent les séances de dessin de figures pour articles, merci de m'avoir introduit au "bouldering" et de ta bonne humeur et de ton sourire et d'être venu à Strasbourg! Dorina, I profess you a great professional respect, and you know how much I love you girl, thanks for sharing the journey with me! To my partner in crime, Chiara, for arriving to the lab and being the best company one can hope for, you made the last two years really special, they would not have been the same if you weren't there. Do you imagine a Grenoble without La Choza? Without Cultural lunches? Without Christmas dinners? I don't, not even after the pandemic. José Villafuerte you have been an unexpected constant, and a great surprise, the intern I once met that came back to join the PhD team, your "inner peace" is always welcome, thanks for bearing with us. Monica V. Gracias, gracias, y gracias. Siempre es un placer tenerte cerca, te deseo lo mejor.

Guislain, Monsieur Il faut que je vous remercie de tout coeur, d'abord pour votre amitié, List of Tables xxvii mais aussi pour m'avoir recruté pour les LDD, à la fin de la thèse quand j'avais vraiment besoin de « blow of some steam ».

To the intern I did not have: Claire, I wish the circumstances that made us get that close had been happier, but life hands you lemons and we did the best lemonade ever !.

During the PhD experience I had the chance to visit some other laboratories and facilities, be it for research experiments or just for communications, I'd like to thank particularly:

Santi and Roque in ICN2 for letting me do some PLD in an attempt to grow very fancy films that didn't show up in the manucript.

Fabrice Wilhelm for the help he offered during the Beamtime I could attend to personally in ID12 for our XANES experience in ESRF.

Miguel Manso from UAM that gave us an exceptional reception in Madrid and the possibility to perform RBS experiments in the CEMAM facilities.

Franciso Rivadulla that kindly invited me to share my reaserch through a semminar in CiQUS-USC in Santiago de Compostela.

Of course I could not forget to thank the friends that I have always close to my heart, no matter where we all decide to scattered, in this world that is so big and so small:

To Manel and Jose for starting the Grenoble adventure with me. To Mari and Jess, no need to state the reason. To Carles, Albert, Gerard and Jordi, I might have moved but I can always count on you. To Nil for Joining the LMGP adventure. To Iris, the co-parent of my cat, a great roommate, a great scientist, and a great friend. To Ana for always asking how I am doing, to Paloma for always making me laugh no matter how bad the weather. To Claudia and Andreas, because kilometres are just a number and our friendship is just timeless, for believing in me and being always supportive. To Constance, because some stars aligned for us to meet. To Sergi, "Like ashes in the wind, like spoken words in time" and totally not like us.

A Román y Rubén, por nuestras excursiones al mar, que tanto me faltan ahora que no vivo en la costa y los Ferroleos esporádicos, que tanto ayudan a desconectar.

A David y Rebeca, por recibirme siempre con los brazos abiertos cuando vuelvo a casa. A mi familia, todos ellos, mi hogar, por su apoyo incondicional desde el principio. Por sus oídos atentos, por su paciencia y por su cariño.

To "my roommate", thanks for sharing these years with me and for not letting me drive us both crazy when thesis writing synched with pandemic times.

And finally to you, that are reading this four pages and therefore I owe you a thank you.

This thesis is dedicated to my family

Foreword

Frame and structuration of the thesis

The manuscript you hold in your hands (or your computer screen) contains the collection of results obtained in the frame of a larger project that joined the efforts of the LMGP team and our collaborators. In this 'Foreword' I contextualize the thesis and explain the structure and organisation of its contents.

The ALPS memories project, acronym for 'Alternative perovskite materials for ReRAM memories: understanding and tailoring resistive switching', was an ANR project (ANR-15-CE24-0018) developed in the context of 'Micro and nanotechnologies for information and communication' (ANR Axis 10). The project, coordinated by Dr. Carmen Jiménez (LMGP), gathered teams from four different laboratories: the Institute of Microelectronics, Electromagnetism and Photonics -Laboratory of hyperfrequency and characterization (IMEP-LaHC, Grenoble), the 'Comissariat à l'énergie atomique et aux énergies alternatives: laboratoire d'électronique des technologies de l'information' (CEA-Leti, Grenoble), the 'Unité de catalyse et chimie du solide' (UCCS, Lille) and the 'Laboratoire des matériaux et du génie physique (LMGP, Grenoble), to work towards the common goal of understanding and tailoring resistive switching in perovskite materials, such as the manganite presented in this manuscript.

The project was developed from 04/01/2016 to 03/07/2019, while this PhD thesis started on the 01/10/2016 and extended after the end of the ALPS memories project and focused on the 'resistive switching in lanthanum manganites (LaMnO 3 ), thanks to the presence of mobile oxygen ions'. It aimed to evaluate LaMnO 3 (LMO) as a candidate for resistive switching (RS) based microelectronic devices. For this purpose devices integrating LMO thin films between two metallic electrodes were fabricated for the first time. The proof of RS on this particular LMO based heterostructure was obtained and the study of the physicochemical properties of the material in the different resistance states was carried out to understand the mechanisms implicated in the switching process.

The structure of the thesis and the links between the main chapters and the appendixes are illustrated in Figure 1, with the exception of Appendix C which presents the main techniques xxxii List of Tables and equipments used during the thesis.

Chapter 1 serves as an introduction to non-volatile memories, and particularly to valence change memories (VCMs), and reviews and explains the choice of manganites as candidates for these memories.

Chapters 2 to 4 present the control of the growth and optimization of LMO thin films on Si 3 N 4 , on single crystals and on platinised silicon substrates, respectively, as well as their morpho-structural properties. While Chapter 2 is a very general introduction to growth of LMO on silicon based substrates, the following two chapters diverge defining two different research lines of the thesis, i.e.:

• The main research line, in which the LMO is integrated on platinised silicon for posterior evaluation of the sandwiched memristive devices (metal/LMO/metal), is presented in Chapter 4.

• A second research line developed in Chapter 3, in which first the deposition and then the oxygen ion transport in highly textured LMO grown in single crystal substrates is presented. This chapter studies the oxygen exchange and diffusion in highly textured LMO films, as the switching mechanism in VCMs relies on the movement of oxygen/oxygen vacancies. A fortunate stroke of serendipity led to the discovery of ultra fast oxygen ion diffusion in these films, which led to the potential use of these films for other nanoionic applications such as solid oxide fuel cell cathodes.

Following the main line of the thesis, films with controlled oxygen content were grown, and the related studies proving the changes in oxygen stoichiometry were published in the form of an article, which is integrally reproduced in Chapter 4. After this goal was achieved, the films were used for the fabrication of two types of devices:

• Standard devices for electrical characterization studies, and

• Advanced characterisation-compatible devices for Synchrotron-based measurements.

The fabrication procedure for both types of devices is presented in Chapter 5. Regarding the standard devices several batches were produced using LMO thin films with different oxygen content. In this thesis only a small fraction of the fabricated devices are analysed. Most of the effort was directed to the analysis of the TiN/LMO/Pt heterostructures, while other heterostructures (Au/LMO/Pt, Pd/LMO/Pt and Ti/LMO/Pt) were either analysed by co-workers and collaborators or left for future work.

As shown in the thesis structuration Figure 1, Chapter 6, completes the main research line of the thesis and deals with resistive switching in the TiN/LMO/Pt heterostructure. It presents 1.1 Towards the smallest library ever built

From the first Palaeolithic cave paintings, passing by the lithographs to scroll writing, and from scrolls to libraries and databases, humanity has always tried to store information as a means of preserving and transmitting knowledge, among other purposes.

In the same way it has been possible to compress the conventional paper-based libraries by storing virtual data in form of bits, a revolution on data storage technology is now required to keep increasing the storage density to further miniaturize devices.

The miniaturisation concept accompanied the evolution of electronics, starting from the vacuum triode [1], used in old radios and TVs, passing by the transistor [2][3][4] and the integrated circuits (IC) [5][6][7]. Since the 70's it was observed that the number of transistors in an IC Introduction doubled every 18 to 24 months. This observation is known as Moore's law. The transistors have effectively miniaturised from the MOSFET process nodes of 10 µm in 1971 to 7 nm in 2018. The doubling of transistors in an IC seems to have 'slowed down' as the field effect transistors (FET) reaches the limits of scaling [8]. Nevertheless the semiconductor manufacturers seem to have found their way around the issue, envisioning the production of 3 nm semiconductor nodes. This can be accomplished through changes in the transistor architecture, such as the gate all around (GAA) FET [9] or the fin field effect transistor (FiNFETs). Another possibility to keep working towards the miniaturisation would be to turn to the alternatives to the FET. While Moore's law is a trend that responds to the goal of a 'minimisation of the space required for a function', a new strategy for miniaturisation known as 'more than Moore' ('MtM') passes by 'the provision of several functions on the same amount of space'conferring the capability to interact with the outside world to the compacts systems. While CMOS (complementary metal-oxide-semiconductor) technology is used for constructing the classic IC's, 'Beyond CMOS' devices, are used for 'MtM' applications. 'Beyond CMOS', one of the focus groups in the 'International roadmap for devices and systems', thrives to overcome the CMOS scaling limitations of device density and speed [10]. In the scope of this focus group, emerging memory devices, based on different 'strategies' are studied. Among them, architectures based on 'memristive elements' are a promising candidate to overcome the scaling limitations in the memory field. The devices based on these elements offer the possibility to store more information in less physical space, as they allow building denser memories.

Memristive devices as candidate for new generation memories

Before going deeper into the description of possible alternatives for new generation memories, it is convenient to introduce the different types of memories. To begin with, we can make a trip through the standard 'start-up' of a computer. The first thing that happens is the reading of the 'read only memory' (ROM), which is a non-volatile memory (NVM) containing the selfdiagnosis instruction to check the proper functioning of the system and the instructions to load the operating system. This operating system is stored in a massive memory device, generally a hard disk drive (but it can be also a CD, DVD, USB-key or other devices) which is as well nonvolatile, has a large storage capacity, and must be as cheap as possible (in the case of the USB key, the storage capacity is reduced and the price constraint removed in favour of a fastest operation velocity). At last, a part of the operating system is loaded in the random access memory (RAM). This allows the microprocessor to execute the basic instructions enabling the The different types of existing memories are presented according to memory hierarchy in Figure 1.1. Among the memories previously described, the only volatile memory is the RAM (Figure 1.1 dynamic random access memory DRAM and static random access memory SRAM). RAM needs energy to keep its content, but the write and read operations are very fast. Nowadays, the minimum operation speed for SRAM is around 5 ns (depending on the operation 5 to 40 ns), and for DRAM around 10 ns (depending on the operation 10 to 100 ns, 60 ns in average) [11]. Here it is where resistive random access memories (RRAM) could be of most advantage, by bringing the 'non-volatility' character to the RAM, while presenting reported speed operations comparable with those of SRAM and DRAM (speed operations range of 5 ns to 20 ns) [12][13][14].

RRAM can be considered as a NVM memory that works by changing the resistance of Introduction an active material or interface. RRAMs can be classified as memristive devices, given that the latter are 'electrical resistance switches capable of retaining a resistance state dependent on its electric history' [15]. They are capacitor-like structures, as generally they are built in a sandwich-like metal-insulator-metal [16] (or metal-insulator-semiconductor [17]) fashion. They integrate a metal-oxide or organic compound as sandwiched material ('top-bottom' configuration) presenting resistive switching behaviour.

These devices offer the possibility to store more information in less physical space as they allow building denser memories. On top of this, the fact that the memory effect is related to a change of resistance of the material allows great scalability. Furthermore, they enable reducing memories power consumption during the write-read processes. For RRAM to be competitive with Flash memories the write voltages (voltage required for engraving information on the memory) need to be compatible with CMOS technology in the range from some mV to V. The resistance ratio between different states can be of 1.2 to 1.3 but it will be preferably larger than 10. Also, write cyclability larger than 10 3 times must be provided. Moreover to be competitive with DRAM and best Flash memories, the read pulses need to be faster than 100 ns. To preserve the resistance state upon reading, the read voltages must not exceed 1/10 of the write voltage value. And finally the minimum read currents should be of 1 µA to facilitate its detection. [18] In a nutshell, resistive switching based memories provide the possibility to have highdensity, high-speed and low-power NVMs.

Resistive switching: concept, classification and terminology

Resistive switching (RS) refers to the reversible phenomena occurring in two terminal circuit elements, which resistance changes under the effect of electrical stimuli (applied electric field) in a non-volatile manner, i.e. the different resistance states (two to more) are preserved in absence of electric field. [19] Resistive switching can be classified according to the switching conditions. In order to proceed to the classification some terms will be defined:

• HRS: High resistance state • Current compliance (Icc): Limit of the current set on the measurement apparatus, mainly to protect the device from breakdown, it can be used to regulate the switching procedure.

• Polarity: Refers to the direction of the applied electric field relative to the grounded electrode. Hence positive polarity will refer to a defined V > 0, and negative polarity to a defined V < 0.

• Volatility: The lack of preservation of the resistance state in absence of electric field.

If the switching occurs for the same polarity (applying a V SET and V RESET voltage of the same sign) the operation mode is referred to as unipolar resistive switching. In Figure 1.2 a, a switch operation in the positive polarity has been taken as descriptive example. Starting in a HRS (blue line in first quadrant) when a certain voltage is reached (V SET > 0) there is a non-volatile resistance change from HRS to LRS. Then, when a V RESET voltage (V RESET > 0) is applied, the state of the memory reverts to the precedent resistance state (HRS). When changing from HRS to LRS there is an imposed current limit (current compliance), while when changing from LRS to HRS there is no current limitation. Both resistance states can be read at small voltages such that V read << V RESET and V read <<V SET .

If the switching is given from a HRS to a LRS in one polarity and reversed from LRS to HRS in the opposite polarity, it is known as bipolar resistive switching (Figure 1.2 b). This mechanism can be non volatile and suitable for memory devices. In Figure 1.2 b the bipolar counter-clockwise (CCW, referred to as the cycling sense at the positive IV quadrant) mode is depicted (where the I-V sweep in positive goes from HRS to LRS). Bipolar resistive switching can also be given clockwise (CW) (from LRS to HRS).

If both resistance changes (SET and RESET) occur for the same polarity when doing a voltage sweep such as 0 V → V max → 0 V, and there is only one stable state at 0 V, the mechanism is classified as threshold switching (Figure 1.2 c). This is a volatile resistance change between two states, and cannot be used as a memory, but can have other potential applications such as selectors to avoid crossbar array sneak path problems [20].

Complementary resistive switching is similar to threshold switching in appearance (Figure 1.2 d), however the HRS and LRS are both stable non-volatile states. Figure 1.2 d shows that between V th3 and V th1 only HRS is available to read. In first instance V > V th1 reads a LRS. Once V th2 is overcome, a jump in resistance state is given and the device is in HRS. This is mirrored in the opposite polarity, i.e. if a sweep operation is performed at negative voltage, a jump in resistance will be given when overcoming V th3 to a LRS and back to HRS for V > V th4 . The memory state can only be read between the V th1 (V th3 ) -V th2 (V th4 ) [21]. This read operation is destructive, which means that the memory state is changed after the read and it is necessary to rewrite it.

It should be noted that the LRS and HRS for any of the described switching types can be ohmic for both states, non-ohmic for both states, ohmic and non-ohmic, respectively, or vice-versa, depending on the electrodes and oxide materials of the device.

The switching mechanisms can be originated by different physical and/or chemical phenomena. It can be derived from nano-mechanical [22] or magnetoresistive phenomena [23], linked to electrical effects [19,24], given by a phase change [25] or by nanoionic redox phenomena [18]. In the last case, ionic motion in the nanoscale leads to redox reactions that affect the resistance of the elements of the two terminal type element. RAM memories based on these phenomena are known as ReRAM (redox based resistive switching random access memories) [19]. Ion induced RS can be classified depending on the mobile species as: cation driven, anion driven and cation-anion jointly driven [26]. The classification of ReRAM according to the nanoionic mechanism implicated separates ReRAMs in: Electrochemical memory (ECM): Where the mobile species are cations from electrochemically active electrodes, such as Ag (cation driven ReRAM) and the switching is caused by electrochemical metallization. The migration of species back and forth from one electrode to the other is controlled by reversing the polarity of the applied voltage [27]. This leads to bipolar resistive switching.

Valence Change Memory (VCM): Where the mobile species are oxygen anions (anion driven mechanism) presents bipolar resistive switching.

Thermochemical memory (TCM) : based on thermal effects. The material across the 2 terminal device is modified due to Joule heating. It may be introduction of carbon from residual organics, decomposed insulator or part of the electrode metal that gets transported into the insulator. [27] This mechanism leads to unipolar resistive switching.

After classifying the memory according to the phenomena implicated, a second classification can be made according to the location of the switching in the two terminal device. Taking the cross section of the electrode/insulator/electrode sandwich, it can be described as filamentary, interfacial or bulk type. Nevertheless, the existence of intermediate scenarios is possible [28,29], as well as the coexistence of several mechanisms.

According to our previous classification, ECMs VCMs and TCMs can all be filamentary, while VCMs can also present interfacial or even bulk type switching. Filamentary RS occurs when a filament of metal or oxygen vacancies is formed through the active material, leading to a change in resistance value. This type of RS normally requires a forming step at high voltage, Introduction and presents a large variability given stochastic nature of the filament formation. In the case of interfacial resistive switching, the switch is related to physical or chemical phenomena taking place at the heterostructure interface (i.e. for an oxide based heterostructures electrode/oxide interface). Interface type switching is based on homogeneous interfacial processes and generally does not require a high-voltage forming step. Among the interface type memories, the switching can be also related to electronic and electrostatic interactions, or ferroelectric polarization. In the bulk type switching, the valence change is given in the total volume of sandwiched active material, it can be the case of memories where a phase change is driven by the variation on oxygen content in the active material [30].

ReRAM have been reported to be scalable down to 10 x 10 nm 2 [19] and their resistance ratios are usually in the 10 -1000 margin (while for Magnetic RAM the ratios are smaller than 5) giving a large sensing margin. The possibility to access to intermediate resistance states between HRS and LRS conducting to multi level resistance potentially provides a large increase in storage density. Writing speeds have been reported at the regime of ns or even 100 ps which makes them faster than Flash memories. The memory readout is non destructive so no restoring cycle is required (as in phase change memories such as PCRAM). The operation energy per bit, i.e. the energy needed to change resistance state, has been reported to be smaller than pJ/bit for TaOx based VCMs [19,31]. And the endurance in term of cycles is typically 10 6 -10 8 , larger than the range for NAND Flash memories (10 3 -10 5 ). In this thesis we will focus on ReRAMs based on the VCM mechanism as they may not require a high-voltage forming process to be activated, and have the potential to be filamentary or interfacial.

Valence change memories (VCM)

The first spectroscopic evidence of VCMs was reported in 2007 by Janousch et. al. in Cr-Doped SrTiO 3 , the authors measured the Cr and Ti absorption K-edges by XANES. The Cr K-edges spectra collected for anode and cathode matched Cr 4+ and Cr 3+ respectively, while Ti K-edges 'are identical within the experimental accuracy' [32]. In 2013, Sandisk and Toshiba presented a VCM type 32GB ReRAM-based chip, developed with cross-point architecture in 24 nm technology. [33][19]; this was the proof of the potential of VCMs to become new generation NVMs.

As indicated above, this work focuses on VCMs, which present bipolar resistive switching, based on the drift of oxygen vacancies in an oxide matrix and the concomitant changes in oxidation state of the cation sublattice. The two terminal element is generally a metal oxide sandwiched between two inert metal electrodes (such as Pt or TiN) of different oxygen affinity and/or different workfunction.

In an oxide, oxygen vacancies have higher mobility than cations and therefore are expected Several examples of VCM have been reported in binary and ternary oxides in the amorphous and crystalline phases. In this work we will focus on the case of a crystalline ternary oxide.

As an example we consider a model material such as SrTiO 3 (STO), an ABO 3 cubic perovskite where the oxygen vacancies can easily reduce the TiO sublattice. In the particular case of a single-crystal, STO is highly homogeneous, with no grain boundaries, and has a very well known defect chemistry, oxygen diffusion and cation diffusion, extendedly reported in literature [35]. STO shows the typical characteristics of a VCM. When it is placed as the 'insulator' in the two terminal element for RS, the point defects, such as oxygen vacancies, electrons or holes, are expected to play a mayor role. In the same manner, out of the single crystal scenario, it is important to notice that interfaces such as grain boundaries, strongly affect the local defect concentration. If we focus on the role of oxygen vacancies implicated in the valence change phenomenon, resistive switching can be given in a filamentary manner, by the creation of bundles of oxygen vacancy chains showing higher electric conductivity than the STO matrix. While when the valence change mechanism is given in an interfacial manner, the oxygen vacancies accumulate in the cathodic region, generating an n-conducting section in the STO matrix [18]. Intermediate scenarios where both mechanisms coexist are also possible.

Introduction

Oxygen vacancies in a certain material can move driven by a gradient in chemical potential (oxygen concentration), by a gradient in electrical potential, and/or by a gradient in temperature. The 'oxygen vacancy movement' in each case is more precisely called diffusion, drift and thermodiffusion, respectively. [35] It should be noted that in a memristive device under operation (applied voltage) the oxygen vacancies/ions can be submitted to simultaneous thermal and electrochemical gradients. In all cases the flux of vacancies/ions will be proportional to the so-called diffusion coefficient ( D v for the vacancy diffusion coefficient and D O for the oxygen diffusion coefficient). Thus the diffusion coefficient of the material (either pristine or electroformed) will determine the switching kinetics in VCMs. Of course, the distance that the species has to travel is also of major importance, and will depend on the geometrical size of the device, but also in the dimensions of the filament and of the gap between filament and electrode in the case of filamentary-type switching.

Research on fast ion transport phenomena and their applications is covered by the field known as Solid State Ionics, and more recently by the subfield of Nanoionics when nanosized dimensions and effects are involved. Since the early 2000s the research on nanoionic devices for energy storage and conversion and for information storage attracted considerable attention. The oxygen diffusion in oxide materials has been widely studied for oxygen ion conductors (OIC) [37] mixed ionic-electronic conductors (MIEC) [38] and composites of both (OIC +MIEC) for their use as electrolytes and electrodes for solid oxide fuel cells, and oxygen permeation membranes, among other applications.

Understanding the oxygen diffusion processes is fundamental for all theses applications, but also the oxygen exchange kinetics with the surrounding gas atmosphere. Although the oxygen exchange is not expected to occur in embedded memories, it can nevertheless play a very important role in laboratory studies in which the memristive devices are directly in contact with the atmosphere and where they could exchange oxygen in localized hot regions during device operation. For applications in which fast kinetics are desirable, this process can be enhanced in films by improving the surface/volume rate of the electrodes [39][40][41], through control of film thickness and decoration, or porosity of the films. Other studies proved that, in epitaxial films, strain variations in the lattice improved performance regarding diffusion [42], surface exchange or both [43,44] of certain materials such as YSZ [42], LaCoO 3 [45], or LaSrCoO 3-δ [46]. In several perovskite films, the role of extended defects such as dislocations in the oxygen diffusion have been studied. On one hand, for SrTiO 3 dislocations normal to the oxygen diffusion direction slow down transport, and dislocations parallel to the oxygen diffusion direction seem to neither enhance nor slow diffusion [47,48]. On the other hand, for manganites such as La 0.8 Sr 0.2 MnO 3 (LSMO) extended defects such as dislocations parallel to the oxygen diffusion direction have been claimed to accelerate 1.3 Resistive switching: concept, classification and terminology oxygen transport [49]. The same behaviour has been observed for several polycrystalline perovskites (such as: La 0.8 Sr 0.2 Mn O 3 , (La 0.8 Sr 0.2 ) 0.95 MnO 3 , La 0.8 Sr 0.2 MnO 3+ δ or La 0.8 Sr 0.2 (Mn 1-x Co x ) 0.85 O 3±δ ) through their grain boundaries (GBs) [50][51][52].

Due to the direct relationship between oxygen transport and switching characteristics of memristive devices, the oxygen diffusion and exchange in manganite thin films were studied in this thesis. The measurements of the mass transport coefficients were carried out at intermediate temperature (500-600 ºC), and thus the results obtained are of large interest not only for VCMs, but also for micro solid oxide fuel cells (µ-SOFC). For this reason the results detailed in Chapter 3 are presented with a focus in this last application.

Resistive switching in manganites

In the 1990's the topic of resistive switching re-emerged (after a first wave of research in the 1960's). It was while investigating the magnetoresistive properties of Pr x Ca 1-x MnO 3 (PCMO) that the Tokura group found electrically triggered RS [53]. Initial publications reported resistance changes under the effect of an applied electric field (in absence of external applied magnetic field) presented by magnetoresistive manganites [53,54].PCMO was used to fabricate the first integrated ReRAM by Zhuang et. al. [55] [19], taking advantage of the colossal magnetorestive films. A large number of articles on resistive switching in manganites have been published since 2000 (127 according to the WOS by 01/07/2020 containing the keywords 'Resistive switching' and 'Manganite').

Since then, the materials from the manganite family (manganites) have proven to present high HRS/LRS ratios (10 3 ) [56,57] fast operation speed (~10 ns write) [58], multi level resistance [59,[START_REF] Lau | Nonvolatile multilevel memory effect by resistive switching in manganite thin films[END_REF], and no need for electroforming [START_REF] Stoliar | Nonvolatile Multilevel Resistive Switching Memory Cell: A Transition Metal Oxide-Based Circuit[END_REF]. They are thus considered suitable for new generation RRAM memory devices.

Particular attention has been paid to La 1-x Sr x MnO 3 (LSMO) [30,[START_REF] Lau | Nonvolatile multilevel memory effect by resistive switching in manganite thin films[END_REF][START_REF] Huang | Study on the resistive switching properties of epitaxial La0.67Sr0.33MnO3 films[END_REF][START_REF] Xu | Unusual resistive switching induced by voltage in La0.7Sr0.3MnO3 thin films[END_REF][START_REF] Lee | A Study on the Resistive Switching of La0.7Sr0.3MnO3 Film Using Spectromicroscopy[END_REF][START_REF] Manca | Reversible oxygen vacancies doping in (La 0.7 ,Sr 0.3 )MnO 3 microbridges by combined self-heating and electromigration[END_REF][START_REF] Kumari | Structural and resistive switching behaviour in lanthanum strontium manganite -Reduced graphene oxide nanocomposite system[END_REF][START_REF] Gonzalez-Rosillo | Nanoscale Correlations between Metal-Insulator Transition and Resistive Switching Effect in Metallic Perovskite Oxides[END_REF], La 1-x Ca x MnO 3 (LCMO) [START_REF] Duhalde | Electric field-driven resistive switching in magnetoresistive La0.5Ca0.5MnO3 thin films[END_REF][START_REF] Shang | Effect of carrier trapping on the hysteretic current-voltage characteristics in Effect of carrier trapping on the hysteretic current-voltage characteristics in Ag La0.7 Ca0.3 MnO3 Pt heterostructures[END_REF][START_REF] Villafuerte | Resistance switching induced by electric fields in manganite thin films[END_REF][START_REF] Sheng | Current-induced resistive switching effect in oxygen-deficient La0.8Ca0.2MnO3-delta films[END_REF][START_REF] Belogolovskii | Interface resistive switching effects in bulk manganites[END_REF][START_REF] Liu | Improved resistive switching properties in stacked structures[END_REF][START_REF] Rubi | Two resistive switching regimes in thin film manganite memory devices on silicon[END_REF][START_REF] Acevedo | Manganite-based three level memristive devices with self-healing capability[END_REF] and PCMO [29,[START_REF] Chen | Perovskite RRAM devices with metal/insulator/PCMO/metal heterostructures[END_REF][START_REF] Kim | Growth and characterization of Pr0.7Ca0.3MnO3 thin films for resistance random access memory[END_REF][START_REF] Liu | Low programming voltage resistive switching in reactive metal/polycrystalline Pr0.7Ca0.3MnO3 devices[END_REF][START_REF] Liu | Low-power and controllable memory window in Pt/Pr 0.7Ca 0.3MnO 3/yttria-stabilized zirconia/w resistive random-access memory devices[END_REF][START_REF] Norpoth | In situ TEM analysis of resistive switching in manganite based thin-film heterostructures[END_REF][START_REF] Lee | Manganite-based memristive heterojunction with tunable non-linear I-V characteristics[END_REF][START_REF] Arndt | Spectroscopic Indications of Tunnel Barrier Charging as the Switching Mechanism in Memristive Devices[END_REF] manganites, for which a valence change mechanism has been identified, although the coexistence with other mechanisms is possible. The VCM cells can be filamentary [START_REF] Kumari | Structural and resistive switching behaviour in lanthanum strontium manganite -Reduced graphene oxide nanocomposite system[END_REF] or interfacial [29] depending on the electrodes of choice [START_REF] Lee | Manganite-based memristive heterojunction with tunable non-linear I-V characteristics[END_REF], the electroforming process [START_REF] Acevedo | Manganite-based three level memristive devices with self-healing capability[END_REF] and potentially other factors, such as trap rich interfaces [START_REF] Shang | Effect of carrier trapping on the hysteretic current-voltage characteristics in Effect of carrier trapping on the hysteretic current-voltage characteristics in Ag La0.7 Ca0.3 MnO3 Pt heterostructures[END_REF]. A selection of studies from 2004 is presented in Tables 1.1 and 1.2 at the end of this chapter, showing the manganite composition and type of electrodes used in each case.

Evidence of the oxygen role in manganites has been shown by different methods, among which the simplest is by performing current voltage sweeps on films grown under different oxygen partial pressure [START_REF] Kim | Growth and characterization of Pr0.7Ca0.3MnO3 thin films for resistance random access memory[END_REF]. In their work, Kim et. al. observed resistive switching in the I-V curves of PCMO films with different oxygen content. The HRS/LRS ratio showed dependence with the oxygen content, it was larger for the films grown under larger oxygen partial pressure.

Introduction

Other studies provided prove of the role of oxygen in resistive switching by using spectroscopic techniques such as electron energy loss spectroscopy (EELS) [START_REF] Arndt | Spectroscopic Indications of Tunnel Barrier Charging as the Switching Mechanism in Memristive Devices[END_REF], X-ray photoelectron spectroscopy (XPS) [START_REF] Kumari | Structural and resistive switching behaviour in lanthanum strontium manganite -Reduced graphene oxide nanocomposite system[END_REF][START_REF] Meunier | Microscopic Mechanisms of Local Interfacial Resistive Switching in LaMnO 3+δ[END_REF] or even synchrotron radiation techniques; such as hard X ray photoelectron spectroscopy (HAXPES) [START_REF] Arndt | Spectroscopic Indications of Tunnel Barrier Charging as the Switching Mechanism in Memristive Devices[END_REF][84] or X ray absorption spectroscopy (XAS) [START_REF] Rathod | Investigations on the Electronic Excitations through Spectroscopic Measures for Resistive Switching Character of Manganite Thin Films[END_REF]. Arndt et. al. observed through HAXPES PCMO based devices incorporating yttria stabilized zirconia (YSZ). Clear changes in the binding energy of Zr 3d orbital of YSZ are observed in the different resistance states. They complemented this study with EELS, which provided proof of the reduction of Mn in the YSZ/PCMO interface and hence of an oxygen exchange between PCMO and YSZ, that was fundamental for the RS shown by their devices [START_REF] Arndt | Spectroscopic Indications of Tunnel Barrier Charging as the Switching Mechanism in Memristive Devices[END_REF]. Kumari et. al. used XPS measurements of the binding energy of Sr 3d peak (in LSMO based devices) to support the observation of conducting filament formation induced by oxygen vacancies [START_REF] Kumari | Structural and resistive switching behaviour in lanthanum strontium manganite -Reduced graphene oxide nanocomposite system[END_REF]. XAS experiments, as the ones performed by Rathod et. al., confirm changes on the oxidation state of Mn through near edge x ray absorption fine structure (NEXAFS). In their work, they study RS in Y 0.95 Sr 0.05 MnO 3 (YSMO) through NEXFAS. They observe the changes in the energy at the Mn L 3,2 -edges related to a change in Mn valence from 4+ to 3+ given by the formation of oxygen vacancies [START_REF] Rathod | Investigations on the Electronic Excitations through Spectroscopic Measures for Resistive Switching Character of Manganite Thin Films[END_REF].

Other techniques, such as in-situ transmission electron microscopy (TEM) have also been fruitful to provide direct observation of changes in distribution or movement of oxygen vacancies movement related to RS [30,[START_REF] Norpoth | In situ TEM analysis of resistive switching in manganite based thin-film heterostructures[END_REF]. Yao et. al. observed the migration of oxygen due to the combined effects of an applied electric field and the Joule heating. They studied LSMO through in-situ TEM combined with EELS, applying a voltage to the sample during observation. Images were obtained in which a structural transition from perovskite to brownmillerite is observed. The oxygen vacancies lateral migration in the LSMO appears to trigger the RS. [30]

LaMnO 3±δ

In LaMnO 3±δ lanthanum usually presents an oxidation state of +3 while the Mn presents a mixed oxidation state Mn 3+ /Mn +4 depending on the oxygen stoichiometry (3±δ) to maintain charge neutrality. This capability to accommodate the valence and oxygen content suggest that LaMnO 3±δ has the potential to be integrated in VCMs.

Lanthanum manganites belong to the the perovskite family. The prototype or ideal perovskite structure has ABO 3 chemical stoichiometric formula (Figure 1.3 a) and space group Pm-3m. The cubic unit cell can be described with the A cations are at the centre of the unit cell with O anions forming an AO 12 coordination cuboctahedron, the B cations are placed on the vertices with O anions forming a BO 6 coordination octahedron as shown in Figure 1.3 a. The position of the ions in the unit cell correspond to the centre of the cell A (1/2, 1/2, 1/2), Distortions from the ideal perovskite can be quantified through the Goldschmidt tolerance factor (t), described by the formula:

t = (r A + r O ) √ 2(r B + r O )
Where r A and r B stand for the cation radii and r O for the oxygen radius.The perovskite is cubic as long as t is in the range [0.89<t<1.02] [START_REF] Coey | Mixed-valence manganites[END_REF]. In the same manner t in the range [0.7<t<0.9] correspond to orthorhombic or rhombohedral perovskites. LaMnO 3±δ , presents a distorted perovskite structure and can be orthorhombic (Figure 1.3 a) or rhombohedral (Figure 1.3 c) depending on the oxygen stoichiometry (Figure 1.4 b). According to Van Roosmalen et. al. [START_REF] Van Roosmalen | Phases in the Perovskite-Type LaMnO3+δ Solid Solution and the La2O3-Mn2O3 Phase Diagram[END_REF], as published in the the La 2 O 3 -Mn 2 O 3 phase diagram (Figure 1.4 b), the perovskite LaMnO 3±δ structure can accommodate La/Mn ratios in the range from of 0.908-1.202 at 850 °C in air. This range can be larger, as proposed by Bertrand et. al. [START_REF] Bertrand | Low-pressure-MOCVD LaMnO3 very thin films on YSZ (100) optimized for studies of the triple phase boundary[END_REF], if the temperature is lower.

Stoichiometric LaMnO 3 is orthorhombic with a Goldschmidt factor of 0.905. It presents, at room temperature, orthorhombic Pbnm structure (Pnma by permutation) with cell parameters of: a = 5.533 Å, b = 5.727 Å and c = 7.668 Å [START_REF] Ritter | Influence of oxygen content on the structural, magnetotransport, and magnetic properties of LaMnO(3+δ )[END_REF].

LaMnO 3 shows a cooperative Jahn-Teller (JT) effect, due to the distortions of the anionic octahedra. This distortion in the structure is spontaneously given in LaMnO 3 by the Mn 3+ being surrounded by a symmetric environment (oxygen octahedra). Variation on the distortion allow us to follow the structure evolution by Raman spectroscopy associated to the evolution of the JT active modes, as will be detailed in chapters 3 and 4. In Figure 1.5, the reference Introduction Raman modes from orthorhombic LMO are presented as published by Iliev et. al. [START_REF] Iliev | Raman phonons and Raman Jahn-Teller bands in perovskite-like manganites[END_REF].

It is important to notice that the structure of this perovskite varies with non-stoichiometry not only regarding the oxygen content but also regarding the La/Mn ratio i.e. the full formula for the structure taking into account the lanthanum and manganese vacancies would be La 1-x Mn 1-y O 3±δ ( x, y>0, δ > 0 or δ < 0). In fact, oxygen cannot occupy interstitial sites in LaMnO 3 . Hence the δ > 0 formulation, often used in the literature for simplicity, corresponds in reality to La and Mn vacancies as described by the complete formula: La 1-ε Mn 1-ε O 3 where ε = δ /(3 + δ ) for a La/Mn ratio of 1. Along the manuscript the notation LaMnO 3+δ is employed imposing a La/Mn ratio of 1. As as shown in Figure 1.4 (reported by J. Töpfer and J.B.Goodenough [START_REF] Töpfer | LaMnO(3+δ ) Revisited[END_REF]) for LaMnO 3+δ the increase of δ leads to a rhombohedral stoichiometry at room temperature, while δ close to zero leads to an orthorhombic structure . For simplicity the formula is expressed as oxygen excess LaMnO 3+δ, which from now on will appear as LMO throughout the thesis.

The conductivity of LMO varies with the crystal structure, while orthorhombic phase has been reported as insulating phase, rhombohedral LMO shows better conductivity. Several studies reported the phase transition for stoichiometric LMO given at the T JT = 750 K (477 °C) [START_REF] Rodríguez-Carvajal | Neutron-diffraction study of the Jahn-Teller transition in stoichiometric LaMnO3[END_REF] where Jahn-Teller distortion is lost. Resistivity measurements in a temperature range from room temperature to above the T JT have been reported in literature [START_REF] Zhou | Paramagnetic phase in single-crystal LaMnO3[END_REF][START_REF] Souza | Magnetic susceptibility and electrical resistivity of LaMnO3, CaMnO3 and La1-xSrxMnO3 (0.13<x<0.45) in the temperature range 300-900K[END_REF], showing a drop of resistivity at this temperature. The structure changes and resistivity changes in temperature are reproduced in Figure 1.6. Few studies in literature refer to resistive switching of the parent compound for the lanthanum manganite family, and until 2016 RS had only been reported for LMO grown on crystal substrates [START_REF] Xu | Evidence for a Crucial Role Played by Oxygen Vacancies in LaMnO3 Resistive Switching Memories[END_REF][START_REF] Jin | Mechanism study on oxygen vacancy induced resistance switching in Au / LaMnO 3 / SrNb 0[END_REF].In 2012 Xu et. al. [START_REF] Xu | Evidence for a Crucial Role Played by Oxygen Vacancies in LaMnO3 Resistive Switching Memories[END_REF] provided proof of the oxygen vacancy role on resistive switching in LMO films grown on Nb:SrTiO 3 (0.8 wt%). In that work, they presented a comparative study where LMO films were grown under different oxygen pressures, producing LMO films with different content of oxygen vacancies (Vö). They showed evidence of the variation in Vö content obtained by X-Ray diffraction measurements using synchrotron radiation and scanning transmission electron microscopy (STEM). The bipolar resistive switching measurements, performed in Pt/LMO/Nb:SrTiO 3 devices, depicted larger variations in resistance for the films with larger oxygen deficiency. Voltages of ± 5 V where applied in pulsed measurements using 1 ms width write pulses and the read currents were of 0.1 µA. Also in that work, a model based on variation of a Schottky barrier at the Pt/LMO interface was proposed. In 2013, Jin et. al. [START_REF] Jin | Mechanism study on oxygen vacancy induced resistance switching in Au / LaMnO 3 / SrNb 0[END_REF] (from the same group) elaborated a study on the oxygen vacancy induced switching in Au/LMO/Nb:SrTiO 3 devices. In this case, the applied voltages upon I-V voltage sweeps were applied in sequences of -1.5 V 0 V 3 V 0 V -1.5 V. Here again the RS is attributed to the variation of the oxygen vacancy concentration of the Au/LMO interface modulating the Schottky barrier thickness.

In 2013, Nath et. al. reported a study on the destabilisation of the insulating state of LMO, hinting towards potential, reversible, non volatile resistance changes [START_REF] Nath | Electric field driven destabilization of the insulating state in nominally pure LaMnO 3[END_REF]. The resistance change was justified as a percolation type transition, where two phases play a role: an insulating phase and a high resistance bad metallic phase, the resistance of the bad metallic phase is almost temperature independent. These experiments were carried out in LMO single crystals and the changes from LRS to HRS and vice-versa observed during the I-V sweeps are very abrupt. In this study, evidence of a certain temperature dependence for the V SET and V RESET is given.

In addition some other works on RS of LMO-based devices have been reported after he beginning of my PhD thesis, showing the growing interest in the topic. In 2017 Gadani et. al. published results on electric pulse induced resistance (EPIR) in nanostructured LMO in absence of magnetic field and also with a 5 kOe applied magnetic field for spintronic based devices [START_REF] Gadani | Low field magnetoelectric and magnetotransport properties of sol-gel grown nanostructured LaMnO3 manganites[END_REF]. Very recently in 2020 Yin et. al. reported (in their to be published work) RS in Pt/LMO/Pt/Ti/TiO 2 /SiO 2 /Si heterostructures showing the coexistence of unipolar and threshold-type switching [START_REF] Yin | Coexistence of Nonvolatile Unipolar Memory and Volatile Threshold Resistance Switching In LaMnO3 Thin Films[END_REF]. They showed endurances larger than 500 I-V cycles for unipolar switching obtained at voltages of V SET ≈ -7V and V RESET ≈ -2V. The read operation was carried out at -0.3V. The changes in resistance are attributed to variations on oxygen vacancy content at the top Pt/LMO interface, leading to variations on the width and height of the 1.4 Scope of the thesis conduction barrier.

In this work LMO was selected as active memristive material of study for RS applications, which constitutes the core subject of the thesis. LMO has been grown by pulsed injection metal organic chemical vapour deposition in the form of thin films on platinised silicon substrates, with Pt used as bottom electrode. To the best of our knowledge there were no publications related to such systems in the literature previous to 2016.

Scope of the thesis

The goal of this PhD thesis was to work towards a better understanding of the physicochemical mechanisms involved in the resistive switching process for valence change memories, and more specifically in LaMnO 3±δ -based devices. As explained in Section1.3.5 the parent compound LaMnO 3±δ was chosen in this thesis as a model material to study the phenomena leading to resistive switching in this memories. To achieve this goal, a series of steps were required. In first instance the growth of LaMnO 3±δ thin films on Si-based and on single crystal substrates (Chapters 2 and 3) had to be optimised. The optimised films were employed in a study dedicated to measure and understand oxygen ion transport in epitaxial LaMnO 3±δ films (Chapter 3). The results obtained in this chapter are relevant, not only for resistive switching, but also for other nanoionics applications such as micro Solid Oxide Fuel Cells. LMO films growth on platinised Si needed to be re-optimised for RS studies Chapter 4. The microfabrication of devices (Chapter 5) was fundamental to perform electrical measurements on standard sandwich like devices and advanced characterisation of the material on heterostructures adapted for operando measurements. Finally, different types of electrical measurements were performed in order to assess the RS behaviour of the fabricated heterostructures and to analyse the mechanisms playing a role in RS (Chapter 6). This study covered the performance of LMO-based devices in terms of cyclability, retention time, device operation parameters (V SET , V RESET ), programming window and device to device variability.

Additional analysis of LMO heterostructures performed in a collaborative manner during the duration of 'ALPS memories' project, which offers further insight in the phenomena implicated in RS in LMO heterostructures is presented in Appendix A and B in the form of co-authored articles published in the last 3 years. This chapter is dedicated to the growth of LaMnO 3+δ (LMO) thin films on Si based substrates. It is divided in three main sections.

Section 2.1 describes the pulsed injection metal organic chemical vapour deposition technique (PI-MOCVD), the reactor and the tunable parameters used in the procedure to optimize the growth of LMO thin films suitable for resistive switching applications.

Section 2.2 is dedicated to the growth of polycrystalline LMO on stress-free Si 3 N 4 /SiO 2 /Si insulating substrates.

This optimisation preceeded, the integration of LMO films on platinised silicon which was then used to study resistive switching of LMO in Top-Bottom configuration. This specific Growth of polycrystalline LaMnO 3+δ thin films by Pulsed Injection Metal Organic Chemical Vapour Deposition.

optimisation on platinised silicon substrates is presented separately in chapter 4. On the other hand, the optimisation of LMO films on Si 3 N 4 /SiO 2 /Si insulating substrates was performed simultaneously with the deposition of highly textured films on single crystals. The specific details to the growth of LMO on single crystals is further developed in Chapter 3.

Pulsed Injection Metal Organic Chemical Vapour Deposition System

The films presented in this work were grown in a Pulsed Injection Metal Organic Chemical Vapour Deposition (PI-MOCVD) JIPELEC reactor. Aside from the common features of the CVD technique, such as large area of homogeneous composition deposition and uniform thickness, the pulsed injection allows the use of liquid precursors at room temperature and its preservation in an inert atmosphere. Chemical vapor deposition techniques (CVD) such as conventional metal-organic (CVD) MOCVD [START_REF] Bertrand | Low-pressure-MOCVD LaMnO3 very thin films on YSZ (100) optimized for studies of the triple phase boundary[END_REF][START_REF] Bosak | Cation and oxygen nonstoichiometry in RâMnâO (R=La, Nd) bulk samples and thin films[END_REF][109][START_REF] Pripko | Growth and characterization of LaMnO 3 /SrTiO 3 bi-layer[END_REF][START_REF] Gorbenko | MOCVD of LaMnO3 on Biaxially Textured Ni-Based Substrates in a Reducing Atmosphere[END_REF] and pulsed injection MOCVD (PI-MOCVD) [START_REF] Vincent | Crystal Structure Transformations of a Magnetoresistive La0.8MnO3-δ Thin Film[END_REF][START_REF] Pignard | Chemical vapor deposition and structural study of La0.8 MnO3âδ thin films[END_REF] have been used for the deposition of LMO and substituted LMO in the form of thin films. These previous studies proved the possibility to control the cation and oxygen stoichiometry for the deposition of polycrystalline and epitaxial thin films [START_REF] Pripko | Growth and characterization of LaMnO 3 /SrTiO 3 bi-layer[END_REF][START_REF] Pignard | Chemical vapor deposition and structural study of La0.8 MnO3âδ thin films[END_REF].

The JIPELEC reactor presented in Figure 2.1 allows the growth of perovskite films, in our case the LMO manganite at high temperatures and controlled pressure environment. Further details about the reactor and growth conditions can be found in the experimental section (4.3.1) of Chapter 4 . As shown in Figure 2.1b, the temperature can be controlled and regulated section per section (in the Figure 2.1b the fixed set temperatures are given). The real temperature for the evaporator has been measured to be between 230 and 260 °C, which is well adapted for the evaporation of the precursor solution used of Tris(2,2,6,6-tetramethyl-3,5heptanedionato) lanthanum(III) (La(thd) 3 ) and Tris(2,2,6,6-tetramethyl-3,5-heptanedionato) manganese(III) (Mn(thd) 3 ) in m-Xylene. The commercial metal-organic precursors La(thd) 3 and Mn(thd) 3 were povided by Strem chemicals (CAS: 14319-13-2 and 14324-99-3 respectively), the m-xylene solvent was purchased from Alfa Aesar (CAS: 108-38-3). The total concentration of species in this precursor solution was kept to 22.5 mM, modifying the relative concentration of La(thd) 3 and Mn(thd) 3 during the film composition optimisation procedure.

A sketch of the reactor is presented in Figure 2.1a, which can be divided in 3 sections. Section 1, named evaporator, corresponds to the precursor solution storing system, the injection line and the evaporation system. The injection line is formed by a glass bottle containing the precursor solution that is kept at room temperature and under inert gas atmosphere. Under these conditions, the precursor is stable for a long time. The evaporation is based in the in- Growth of polycrystalline LaMnO 3+δ thin films by Pulsed Injection Metal Organic Chemical Vapour Deposition.

jection of the solution containing the precursors from the pressurised line to a heated chamber under vacuum. This method provokes a sudden evaporation of precursor even for molecules with a low vapour pressure. The pulsed injection mode uses a car injector, which works as a high speed electro-valve, as liquid flow controller. The amount of solution injected is a function of: i) the pressure used to push the solution to the injector, ii) the valve opening time and iii) the injection frequency. The injector plays a double role: it controls the size of the injected volume and it provokes a pulverisation into microdroplets, improving the evaporation rate of the solvent and the precursors. After evaporation, the precursor gases are transported to the reaction chamber. In order to avoid the condensation of the vapours, the line has to be heated respecting an increasing thermal gradient. This section is identified as 2 in Figure 2.1a. The evaporated product is transported by a controlled Ar flow towards the reactor chamber, identified as 3. The chamber contains the substrate that is placed at the centre of a furnace, to be heated to the selected deposition temperature. The system used is a hot-wall quartz reactor. The pressure in the chamber is regulated by vacuum pumping through a butterfly valve. In this system, the flux of these gases can be sent from top to bottom, during the preparation step to prevent the deposition of contaminating particles, and it is possible to switch to a flow from bottom to top during the deposition. The by-products of the reaction flow out of the chamber towards an outer collector cooled with liquid N 2 . In this work, the solution container was pressurized with argon at 1.5 bar, the injection of the liquid precursors was performed using a frequency of 2.5 Hz and an opening time of 2 ms with a solution feeding rate of 0.35 mL/min. The evaporator was thermalized at 250 °C and an Ar flow of 0.1 slm was used as carrier gas. The total pressure in the chamber was fixed at 5 Torr and oxygen gas was added directly in the reaction chamber to obtain an oxygen partial pressure of 50%. The deposition temperature ranged from 500 to 750 °C. A thermally activated chemical reaction takes place at the surface of the heated substrates. The steps involved in the MOCVD reaction are schematized in Figure 2.1a: adsorption, gas-solid reaction, nucleation and growth of LMO, and release of the by-products formed.

Silicon-based Si 3 N 4 /SiO 2 /Si substrates

The goal of this series of depositions was to obtain polycrystalline LMO films with perovskite crystal structure, a La/Mn ratio of 1, which are highly dense and homogeneous, and presenting a surface morphology suitable for device fabrication. Hence, the film properties needed to be known in terms of composition (ratio La/Mn and oxygen content), crystallographic structure and morphology (thickness, roughness and grain size).

The choice of Si based substrate, in this case Si 3 N 4 /SiO 2 /Si, facilitated the compositional analysis of the films, as the substrate does not contain heavy elements nor shares any element with the film. This allowed the use of techniques such as Rutherford backscattering spectrometry (RBS) to determine the La/Mn ratio of the films. The specifics of the techniques are explained in Appendix C. RBS has been used in order to characterise the film composition, in terms of the relation between La/Mn ratio introduced in the precursor solution and incorporated in the films. The results obtained from these two techniques allowed to obtain a calibration curve for energy dispersive x ray spectroscopy (EDX) measurements. In this way the La/Mn ratio can be estimated for any sample grown on Si based substrates using EDX data. Preferrably the samples should be calibrated with a series grown under the same temperature conditions and with the same thickness to reduce the error on the estimation of the La/Mn ratio.

The LMO films were measured by grazing incidence X ray diffraction (GIXRD). This technique is suitable to study the diffraction pattern obtained for crystallites of different orientation, and establish a comparison with different crystal structures of LMO, orthorhombic and rhombohedral. This study was combined with Raman spectroscopy, which is also sensitive to the crystal structure. Scanning electron microscopy (SEM) was used to determine the morphology of the sample in top view (surface), and atomic force microscopy (AFM) to study the topology and roughness. The specification of the equipements used for the characterisation are described in Appendix C. Initial test thin film samples were grown on Si 3 N 4 /Si substrates. These substrates were prepared in the PTA by deposition of a Si 3 N 4 film of 300 nm on a Si-wafer by plasma enhanced chemical vapor deposition (PECVD). This test thin films showed that the LMO films tended to crack due to thermal instability (probably related to differences in the thermal expansion coefficient between Si and Si 3 N 4 at high temperature [START_REF] Habermehl | Coefficient of thermal expansion and biaxial Young's modulus in Si-rich silicon nitride thin films[END_REF]). It must be noted that the Si 3 N 4 film had an expected nitride stress of 130±20 MPa after deposition on Si.

A second 'stress free' substrate consisting of 300 nm Si 3 N 4 over 200 nm oxidised Si on a Si (001) oriented wafer (Si 3 N 4 /SiO 2 /Si) was purchased to the 'Institute of Microelectronics of Barcelona -Centro Nacional de Microelectronica' (IMB -CNM). These substrates withstanded better the deposition conditions, no cracks were observed in the films after the deposition. The roughness of these substrates measured by AFM on an area of 1 µm 2 is of 0.45 nm ± 0.03 nm.

Effect of the solution composition in the growth of polycrystalline LaMnO 3+δ films

This section presents the substrates used for the growth of LMO films, and the consecutive optimisation steps followed to obtain polycrystalline films suitable for memristive devices in Top-Top configuration. In particular, the discussion on the morpho-structural characteristics related to the variation of the film thickness, the concentration of the precursor species in the solution are presented here, as well as the discussion on the morpho-structural characteristics related to the variation of the growth temperature and the cooling down atmosphere.

Previous to this work, in the context of 'Alps Memories' project, Dr. D. Pla had performed a first optimisation of LMO films on Si based substrates using an injection frequency of 5 Hz and 2 ms as electrovalve opening time for the injection, a deposition temperature of 700 °C and 5 Torr pressure in reactor chamber, while the relation Ar/O 2 was kept at ~50%. These films presented a semi-quantitative La/Mn ratio of 1.07 estimated by EDX, with an average grain size of 25 -50 nm and the estimated film thickness of 150 nm (measured from the SEM cross section) for a sample grown by injecting 15000 pulses, Figure 2.2, the deposition yield was of 0.01 nm/pulse and the growth rate of 3 nm/min. Nevertheless, the columnar opening was too broad for electrode fabrication, given that the separation between colums could easily lead to the formation of percolation paths fo the metallic electrodes.

In order to optimize the LMO films in terms of film density, homogeneity, stoichiometry and roughness, new conditions were selected for the deposition by PI-MOCVD. On the one 2.3 Effect of the solution composition in the growth of polycrystalline LaMnO 3+δ films 27 hand, an injection frequency of 2.5 Hz was set to decrease the growth rate and improve the film density. On the other hand the deposition temperature was increased to 750 °C to improve the crystallisation of LMO. Given the variation in the deposition parameters, the analysis of the influence of the precursor solution cationic ratio in the composition and thickness was required. Besides, as the new deposition temperature was higher, the composition of the film could be also affected. With this purpose, two new series of samples were prepared, keeping the optimised oxygen partial pressure and opening time from the precedent study. Two sets of different thicknesses corresponding to 10000 and 15000 pulses were grown using a ratio La/Mn = 2 in the precursor solution, as this was the precedent optimal precursor ratio. Three different La/Mn ratios of precursor species (1.5, 2 and 2.5) were selected for a second series grown by the injection of 10000 pulses. In order to ease the identification of these samples they are presented in table 2.1 with their respective identifiers.

In first instance, a series of 3 samples with different precursor solution ratios was prepared. The variation of composition in the films was analysed by combining EDX and RBS (see Appendix C). RBS, a highly precise quantification technique, will be used as reference for the measured samples. Additionally, the RBS technique allows to estimate the thickness of the samples after having calibrated it with a measured film of known thickness. EDX is faster than RBS, it is more readily available and easy to use, but it is a semi-quantitative technique at best. Nevertheless, a RBS/EDX calibration can be used to estimate the La/Mn ratio from EDX data.

The three LMO films on Si 3 N 4 /SiO 2 /Si were measured by these two techniques and the La/Mn ratios obtained are presented in Table 2.2. All the data show an excess of La in the films with the exception of 1.5LMO10kp. RBS results for thickness estimation show that the films with less La in the precursor solution have a smaller growth rate than those in wich La is twice as abundant as Mn. However, the growth rate is comparable for both films with a ratio La/Mn≥2 (0.011 nm/pulse and 0.012 nm/pulse). The La/Mn ratio in the film obtained by RBS for a given La/Mn ratio in the precursor solution. b) Calibration of the La/Mn ratio in the films. La/Mn ratio extracted from RBS data vs the La/Mn ratio obtained from EDX at 15 kV. The linear relation between both that allows correcting the EDX La/Mn ratios is presented in the red box. The green line stands for a slope s = 1. The light green highlighted region stands for the tolerable deviation from ratio 1 to obtain a sample in the 'perovskite-type phase' according to Van Roosmalen et.al [START_REF] Van Roosmalen | Phases in the Perovskite-Type LaMnO3+δ Solid Solution and the La2O3-Mn2O3 Phase Diagram[END_REF].

the linearity of the incorporation of the elements to the film when increasing the La/Mn ratio. The green line corresponds to a slope s = 1. The light green highlighted region corresponds to the tolerable deviation from ratio 1 to obtain sample growth in the 'perovskite-type phase' according to Van Roosmalen et.al [START_REF] Van Roosmalen | Phases in the Perovskite-Type LaMnO3+δ Solid Solution and the La2O3-Mn2O3 Phase Diagram[END_REF]. From these results it can be concluded that the the 1.5 La/Mn ratio in the solution results in stochiometric LMO films. The RBS data were then used to obtain the relation with the EDX data presented in Figure 2.3 b. The relation between RBS and EDX data is established to be 1.23 from this calibration. The surface morphology of this composition series was analysed by SEM. In Figure 2.4, the surface morphology of different LMO films/Si 3 N 4 /SiO 2 /Si is presented. The grain size, calculated from 100 measured grains per sample, has decreased 50% in size compared to the precedent samples (Dr. D. Pla Figure 2.2) In addition, the grain size is more homogenous, being the largest standard deviation of 6 nm. An increase in grain size with La/Mn ratio is observed and can be related to an increase in the thickness of the film (see Table 2.2). Although in Figure 2.4 a the image quality permited only to estimate an average grain size value taking into account the bigger grains, there is a remarkable difference in grain size bewteen this sample (1.5LMO10kp) and the other two.

To complete the analysis of this series, the crystal structure of the films was characterised by XRD to evaluate if the difference in stoichiometry leads to the presence of impurity phases. XRD θ -2θ (°2θ ) diffractograms obtained from the LMO/Si 3 N 4 /SiO 2 /Si samples showed mainly the substrate peaks, while the Bragg peaks corresponding to LMO presented intensities at a level comparable to noise. Thus the polycrystalline films were then studied by GIXRD, using a grazing angle between 0.3°and 0.6°which was optimised for each sample. This techniques allow to increase the intenstity of the diffraction of the LMO film. as presented in Figure 2.5. The GIXRD diffraction patterns obtained for the LMO films prepared from different composition precursor solutions are shown along with a Pnma orthorhombic and a R-3c rhombohedral patterns in Figure 2.5. Only one diffraction peak from the substrate is detected, labelled with an S. As presented in Figure 2.5, the broad peaks could fit both references, the orthorhombic and the rhombohedral patterns. A broad signal around 2θ 28°, as well as a shift towards smaller angles can be observed with increasing La excess. All the other visible peaks (except from that of the substrate) can be attributed to LMO even when the composition is farthest from the stoichiometric value. As explained by Van Roosmalen et. al. [START_REF] Van Roosmalen | Phases in the Perovskite-Type LaMnO3+δ Solid Solution and the La2O3-Mn2O3 Phase Diagram[END_REF] there is a range for which the LaMnO 3 perovskite can be grown without parasitic phases. The La 2 O 3 -Mn 2 O 3 phase diagram was presented in the introduction (1.3.4). In air at 850 °C this La/Mn ratio range goes from 0.908-1.202, but it can be extended at lower temperatures in agreement with Bertrand et. al. [START_REF] Bertrand | Low-pressure-MOCVD LaMnO3 very thin films on YSZ (100) optimized for studies of the triple phase boundary[END_REF]. In our case the perovskite films grown at 750 °C seem to be free of crystalline impurity phases up to La/Mn ≈ 1.5. However, the presence of amorphous impurity phases cannot be fully discarded given the large La/Mn ratio. The reduction in intensity and the broadeding of the peaks for the film with largest La excess can be explained by the formation of a semi-amorphous phase of La 2 O 3 [START_REF] Zhang | Mechanochemical synthesis of LaMnO3 from La2O3 and Mn2O3 powders[END_REF] upon increase of La excess in the film.

In summary, for the La/Mn composition ratio series, the standard characterization has shown that:

• It had been previously found that a solution with a La/Mn ratio of 2 leads to the stoichiometric composition of La/Mn = 1 for a deposition temperature of 700 °C. • The increase of La in the solution provokes a diminution of the growth rate, probably due a lower deposition yield of managanese.

• Larger thicknesses of the films lead to bigger grains.

• All films present La excess with the exception of 1.5LMO10kp. The excess of La probably provokes the formation of an amorphous La 2 O 3 phase and could correspond to the broad signal detected in GIXRD for La/Mn ratio in the film close to 2.

Effect of the temperature in the growth of polycrystalline LaMnO 3+δ films

With the first optimisation we determined a precursor La/Mn ratio range to obtain the desired La/Mn composition in the film without secondary phases. Now the goal is to determine the optimal thickness and master the control of the phases (orthorhombic or rhombohedral). In this second optimisation the influence of the temperature on the features of the LMO films was studied. The sample identifiers and deposition conditions are presented in Table 2.3.

Reducing the deposition temperature is convenient to avoid substrate instability. The challenge of the temperature reduction comes from finding a compromise between substrate stability and good crystallisation of the LMO film. It is also known that differences in temperature in the growth process affect the species incorporation to the film [START_REF] Pignard | Chemical vapor deposition and structural study of La0.8 MnO3âδ thin films[END_REF] as shown by Pignard et. al. (Figure 2.6). In their study, Pignard et. al. [START_REF] Pignard | Chemical vapor deposition and structural study of La0.8 MnO3âδ thin films[END_REF] grew LMO by PI-MOCVD in horitzontal configuration. They determined that the La/Mn ratio varies with the deposition temperature for temperatures above 675 °C. The variations they observed consisted on a small increase in the La/Mn ratio between 650 °C and 700 °C and considerably large increase in La content above 700°C. Here we varied the deposition temperature between 600 °C and 750°C (Table 2.3). Given the tendency presented in Figure 2.6, we assume that the temperature decrease will lead to a reduction of incorporated La to the film. Hence, the precursor solution ratio for this series was fixed to La/Mn = 2.

Since the objective is to obtain flat, dense and as homogeneous as possible films for the latter device fabrication, and since, as we have seen, thinner samples present smaller grain sizes, the thickness of the films was reduced by decreasing the number of pulses injected.

To study the effect of thickness the films were grown by injecting 6000 and 10000 pulses (Table 2.3).

The RBS results presented in Table 2.4 show that the La/Mn ratio of the LMO films increases with the deposition temperature. As previously explained, this tendency is expected for temperature variations (See Figure 2.6). This can also be observed in Figure 2.7, where the La/Mn ratios obtained by RBS are plotted for the sample series grown with 6000 pulses and 10000 pulses. The same trend was found for both thicknesses. The composition of the films does not vary largely with temperature below 675 °C but the ratio of La incorporated to the film increases considerably for temperatures above this value.

Regarding the films thicknesses, they are comparable below 675 °C , hinting towards the growth being controlled by the precursors diffusion [START_REF] Pignard | Chemical vapor deposition and structural study of La0.8 MnO3âδ thin films[END_REF]. Above 675 °C several variations are observed. First a small increase in thickness is given better observed in the 6000 pulses series between 660 °C and 700 °C, which occurs simultaneously to an increase of the La incorporation to the film. Since there is La excess, we assume that La is accomodated in the perovskite structure and the film becomes Mn vacancy rich. In the series at 10000 pulses, between 700 °C and 750 °C the thickness decreases as the temperature increases, in agreement with the tendency presented by Pignard et. al. [START_REF] Pignard | Chemical vapor deposition and structural study of La0.8 MnO3âδ thin films[END_REF]. This decrease in film thickness with temperatures higher than 700 °C can be related to the decomposition of the Mn precursor before reaching the substrate, since its decomposition temperature is lower than that of the La precursor. This would explain the fastest incorporation of La in films grown above 750 °C observed in section 2.3.

Through the RBS technique the oxygen quantification in the films was possible (Appendix C). The atomic percentage (at.%) oxygen reference for a stoichiometric LaMnO 3 is of 0.6 (La 0.2, Mn 0.2, O 0.6). In Table 2.4, the oxygen atomic percentage is always above 0.6. This implicates that δ > 0 in LaMnO 3+δ . As previosuly explained LMO cannot accomodate oxygen in interstitial sites [START_REF] Töpfer | LaMnO(3+δ ) Revisited[END_REF], hence the nominal oxygen excess is accomodated by the formation of cationic vacancies of La and Mn. The incorporation of oxygen seems fairly independent of the growth temperature, although the largest concentrations were found for the highest temperatures in each series.

The surface morphology of this series was studied by SEM and the surface roughness was measured by AFM (Figure 2.8). Figure 2.8 shows that the grain size slightly increases from 13 to 17 nm whith the increase in La excess (corresponding to a thickness increase from 65 to 112 nm). It can be observed in Figure 2.8 for 700 °C that the SEM shows white and black particles which could correspond to larger grains in the cross section direction and to missing grains or grain separation. This contrast matches well with the height scale showing differences of up to 20 nm in height for the 700°C sample. The roughness follows the same tendency increasing from 1.45 to 1.79 nm. This could hint towards a larger separation between the grains, which would be preferrably to avoid for device fabrication. Hence, temperatures below 700 °C seem to fulfill the composition and morphology requirements better.

Nevertheless the crystal structure is of fundamental importance, hence to validate the advantages of lower temperature deposition, Raman spectroscopy was used to study the phase variations on the films grown at different temperatures. LMO is a well know material and the Raman lines for both crystal structures have been reported in the literature [START_REF] Iliev | Raman phonons and Raman Jahn-Teller bands in perovskite-like manganites[END_REF][START_REF] Iliev | Role of Jahn-Teller disorder in Raman scattering of mixed-valence manganites[END_REF] as previously presented in the introduction (1.3.4). The reference Raman lines extracted from the literature are indicated as guide for the eyes. In general, the film seems to be a mixture of the two phases as it presentes peaks corresponding to both the orthorhombic (marked as an olive green O) and rhombohedral (represented by a navy blue R) phases. In Figure 2.9 a for the thinnest films a clear increase of the rhombohedral contribution with temperature can be observed, proved by the intensity increase of the band at 671 cm -1 . The biggest change is produced from 675 °C to 700 °C, while between 650 and 675 °C no differences are observed, as expected since the change in composition is given at 700 °C. In Figure 2.9 b the second series in temperature (10000 pulses) is shown. Since the films are thicker there is a larger peak to background signal. The same tendency (as for the 6000 pulses series) can be observed, only this time with a very small shift of the band at lower wavenumber from 278 cm -1 towards 217 cm -1 . This observation reinforces the idea of rhombohedral structure gaining importance in the mixture in films deposited at higher temperature. Both Figures 2.9 a and b show the presence of the orthorhombic mode at 482 cm -1 overlapped by the tail of the main Si peak, supporting that in all cases there is an orthorhombic-rhombohedral phase mixture.

The range of temperatures validated to grow LMO is 650 °C -700 °C. In this range, the limit ratio of precursor solution La/Mn = 2 has been validated to obtain films of perovskite

Effect of the cooling atmosphere in the deposition of polycrystalline LaMnO 3+δ films

Once the window of La/Mn ratio in the precursor solution and the temperature to grow LMO perovskite have been determined, it is of interest to tune the oxygen content for a fixed cationic ratio since, as explained in the introduction (Subsection 1.3.5), oxygen is expected to play a role in resistive switching. For that, LMO films of different thicknesses were deposited at 700 °C, and cooled immediately after deposition in different atmospheres. To tune the oxygen content of the films, these were cooled at 5 Torr with either 100% of Ar or a 50% Ar/O 2 mixture. This should not affect the La/Mn ratio incorporated to the films, since it does not affect the growth procedure, but it may induce changes in surface morphology and film structure.

In Figure 2.10 the surface morphology of films with two different thicknesses with different cool down atmospheres is shown. The cool down environment appears to affect only slightly the grain size, within the error bars of the measurements. In order to study the morphology changes on the thinnest films AFM was used . The topography of samples obtained by the injection of 2500 pulses (15 nm thick) for the two different cooling environments is presented in Figure 2.11. The rms roughness was 0.75 nm for the film cooled in oxygen and 0.73 nm for the one cooled in argon. This roughness is smaller than the ones measured for thicker samples. The minimisation of roughness is critical for devices fabrication. The roughness, as is the case of the grain size, seems unaffected by the cool down environment.

To verify if a variation in oxygen content was obtained by changing the cool down conditions, the samples were analysed by Raman spectroscopy, as shown in Figure 2.12 .For all the samples presented were (grown at 700 °C) the rhombohedral component is present in all the films, as indicated by the presence of the Raman mode at 671 cm -1 . Despite having always a mixture of phases, the spectra show an increase of the rhombohedral signal and a decrease of the orthorhombic signal when increasing the oxygen content. The composition in terms of La/Mn ratio has been verified to be close to the stoichiometry for the samples grown at temperatures lower than 700 °C (La/Mn close to 1). The increase in oxygen stoichometry δ (LaMnO 3+δ ) leads to a mainly rhombohedral structure, as reported in literature [START_REF] Töpfer | LaMnO(3+δ ) Revisited[END_REF][START_REF] Töpfer | Charge transport and magnetic properties in perovskites of the system La-Mn-O[END_REF].

Conclusions

• The composition study led to the determination of the range of the La/Mn ratio in precursor solution that leads to the growth of perovskite LMO without impurity phases. The La/Mn ratio in precursor solution should be limited to La/Mn = 1.5 if the growth is performed at 750°C, while a ratio of La/Mn = 2 allows growing stoichiometric LMO Growth of polycrystalline LaMnO 3+δ thin films by Pulsed Injection Metal Organic Chemical Vapour Deposition.

below 700 °C. The composition study provided insight on how La is the limiting species for the growth, and how some La excess is tolerated by the structure before La 2 O 3 phases appear (as shown by GIXRD) .

• The temperature study permitted finding a compromise to lower the growth temperature preserving substrate stability, preserving the La/Mn close to 1 ratio in the films. A clear dependence of La/Mn ratio with deposition temperature was identified. Growth temperatures lower than 700 °C led to LMO films that fulfill better the morpho-structural requirements.

• A nominal oxygen excess was identified in all the grown samples (LaMnO 3+δ with δ > 0). The oxygen excess implicates the presence of La and Mn vacancies in the films.

• The study of oxygen incorporation in the film by changing the cool down atmosphere proved that it is possible to tune the oxygen content, as verified by the change in structure measured by Raman spectroscopy.

All these results allowed to estabilish the optimised depositon conditions to grow LaMnO 3+δ films integrated on platinised silicon, as presented in chapter 4.

Chapter 3

LaMnO 3+δ growth on single crystal substrates and oxygen ion transport studies. Most of the optimisation of the growth of the LMO films on single crystal substrates was performed simultaneously to that of the films grown on Si based substrates. In this chapter, an overview of the morpho-structural characterisation of the deposited thin film samples is presented following the same scheme as in chapter 2. The chapter is divided in two main sections. The first section consists of a recapitulation of the optimisation regarding film thickness and cool down environment. The second section presents the oxygen ion transport studies that were performed on the LMO films grown on single crystals. These results are written in the form of an article. Although the article's motivation and discussion were written with a focus on micro SOFCs applications, as previosyly explained in the introduction section 1.3.2, the finidings related to the diffusion and surface exchange coefficients are also of great importance for the application of LMO as memristive material, as they can be directly related to the switching kinetics of the devices. .

Optimisation of the growth of epitaxial LaMnO 3+δ thin films

The optimisation of the growth of LMO epitaxial films on the substrates of choice, namely SrTiO 3 (001) (STO) and LaAlO 3 (001) (LAO) purchased from CrysTech GmbH (Germany), was performed simultaneously with the composition optimisation of LMO/Si-based substrates. Hence the growth conditions are as described in Chapter 2 in Section 2.3.

The LMO films morphology was analysed by SEM and AFM, which also allowed to study the roughness of the sample surface. X ray reflectometry (XRR) studies allowed to estimate the thickness, roughness and density of the films. The crystal structure analysis was performed by X ray diffraction (XRD) in θ -2θ configuration, and complemented by Raman spectroscopy, which permited the observation of structural changes on the LMO/LAO films.

LaMnO 3+δ epitaxial films: thickness and oxygen content tunability

This section focuses on the optimisation of two features of LMO films on single crystal substrates. On the one hand, the film changes related to differences in thickness are studied to evaluate and optimize the film thickness for device fabrication. On the other hand, the oxygen content of the films is varied by changing the cool down environment in the reactor chamber after film growth. For this new series the temperature was fixed to 700 °C and the La/Mn ratio in the precursor solution was fixed as La/Mn = 2. These conditions were selected to perfom the deposition of epitaxial films simultanesously with that on Si-based substrates. The general conditions indicated in chapter 2 (injection frequency of 2.5 Hz, electro-valve opening time of 2 ms and a ratio of Ar/O 2 of 50% in the reactor chamber for a fixed P of 5 Torr) were kept for all the films deposited. The thickness was tuned by varying the number of pulses between 2500, 6000 and 10000 pulses. Different cool down environments (argon or oxygen) were selected to tune the oxygen content.

The code names of the samples and their specific growth parameters are presented in Table 3.1. Raman could only be used to analyse the LMO/LAO films (STO substrates have a very intense signal in Raman and the film/substrate signal ratio is very poor and most lines overlap) and EDX allowed the semi-quantitative compositional analysis of the films only on STO (LMO/LAO films composition analysis is difficult given the presence of La in the substrate).

Effect of the thickness in the growth of epitaxial LaMnO 3+δ films

The morphology of the samples was studied by SEM and AFM. Figure 3.1 presents the evolution of surface morphology with the thickness for LMO films on LAO (Figure 3.1 a) in terms of grain size and surface roughness. As precedently observed in the study of the LMO growth on Si based substrates, the grain size slightly increases with film thickness. The values are within the same range, from 11 nm to 30 nm. In Figure 3.1 b the same tendency can be appreciated on LMO films on STO, although the thickness difference is in the range of the error. Comparing the LMO films grown on both substrates, an increase in roughness (from 0.5 to 1.6 nm) is appreciated for the samples grown on LAO. In the thicker film, (B-LMO10kp-O2) it can be appreciated that the grains have a squared shape and are arranged in two perpendicular directions.

XRR was performed to measure the thickness of the films, as the film fringes allow to extract the film thickness from the simulation performed fitting the reflectometry data. The thickness values are presented in table 3.2 and range from 10 to 20 nm depending on the number of pulses, accompanied by the estimated density values obtained from the simulations performed in LEPTOS software. According to Van Roosmalen et. al. [START_REF] Van Roosmalen | The Defect Chemistry of LaMnO3+δ[END_REF] the theroetical density of LaMnO 3+δ is of (6.57 ± 0.02) g • cm -3 and can be considered independent of δ in the range [0,0.18]. The values obtained through XRR simulation are in good agreement and within the error of the expected density. An example of fitting simultation is shown in Figure 3.2. Sample B-LMO10kp-O2 was not meassured by XRR since, due to the large thickness, previous measurements on samples grown at 750°C for the same number of pulses presented no oscillations.

The structural study was performed by XRD in θ -2θ configuration. The obtained diffractograms showed textured films in the (00l) direction (for the pseudo-cubic structure) as presented in Figure 3.3. For simplicity, the peaks are indexed in the pseudocubic system for LMO all along section 3.1.The observation of only one family of peaks is consistent with highly textured or epitaxial films. Figure 3.3b shows the LMO films on STO single crystal substrate. The overlap between the LMO film peaks and the substrate peaks is also indicative of highly textured films. In Figure 3.3a doublets can be observed in the (002) and (003) peaks for the LMO/LAO films in B-LMO10kp-O2 and B-LMO6kp-O2. The different out of plane parameters (calculated as pseudocubic parameters for simplicity and previously presented in table 3.2 ) could indicate distinct domains with different in plane orientations of the LMO cell.

In order to obtain information on the exact crystallographic phase and its evolution with the film thickness Raman spectroscopy was performed on the LMO/LAO films, as shown in The shoulder at 671 cm -1 is related to the rhombohedral structure and appears in all films. The contribution of the orthorhombic line at 611 cm -1 increases with the thickness up to 6000 pulses. This peak seems to displace towards the rhombohedral position for the thickest films.

Cool down environment effect

The samples were either cooled in a 5 Torr of Ar environment, or in a 50% Ar/O 2 environment. We showed in the previous chapter that the crystallographic phase of the LMO can be tuned with the incorporation of oxygen through this means. Here we show its effect in films grown on single crystal substrates. As this variation in oxygen content did not induce great changes in morphology, we focus in the crystal structure changes. In this section only films grown on LAO are shown. The cell parameters calculated in pseudocubic system are shown in Table 3.3. Figure 3.5 shows the θ -2θ diffractograms obtained from three different thicknesses in both cool down conditions. A displacement of the main peaks towards higher 2θ of the samples cooled in O 2 rich atmosphere is given (lower main cell parameter in Table 3.3) . Triplets are visible in the 002 and 003 reflexions in the difractograms of the samples cooled in Ar.

The same samples cooled in different atmospheres were also analysed by Raman spectroscopy in Figure 3.6. Comparing the samples cooled in Ar with those cooled with 50% O 2 , there is a remarkable increase in the orthorhombic contribution for all films when they are cooled in Ar environment. The rhombohedral contribution from the 671 cm -1 line is observable in all films and does not disappear with the reduction of oxygen content. In the B-LMO6kp pair, it becomes clear that the orthorhombic contribution is stronger for B-LMO6kp-Ar, as the lines at low wavenumber corresponding to the orthorhombic structure are observable (the triplet at 257 cm -1 , 284 cm -1 , 308 cm -1 ).

Characterisation Summary

The film thickness affects the morphology of the films. The thicker the film the larger the roughness and grain size, further characterisation by TEM in cross section is convenient to analyse the changes in structure with thickness of the samples, it will be presented in section 3.2.4.

The oxygen content does not affect notably the surface morphology features of the films, but it does affect the structural features. Changes in the out-of-plane cell parameter have been observed (smaller parameter for samples cooled in oxygen), and the mixture of phases that composes the film varies from mainly orthorhombic (cool down in Ar) to mainly rhombohedral (cool down in O 2 ).

The two thinnest films (2500 and 6000 pulses) were further analysed to determine their tracer oxygen surface exchange and diffusion coefficients by 18 O exchange experiments coupled with time-of-flight secondary ion mass spectrometry (Tof-SIMS) analysis. This led to the study presented in section 3.2 in the form of an article.

Oxygen transport in nano-architectured thin films

The contents presented in this section is reproduced here in the form of a draft of an article, which is planned to be submitted soon to a scientific journal. The samples used for this study resulted from the optimisation procedure described in the previous sections. They correspond to the A(B)-LMO6kp-O2 samples and a new batch of A(B)-LMO2.5kp-Ar.

Design of nano-architectured LaMnO 3 films with fast oxygen transport properties at low temperature

Introduction

In the early 2000s with the need for new ionic solid-state electrochemical micro devices for energy conversion and storage devices (such as micro batteries and micro fuel cells) and for information storage (such as redox-based resistance random access memories, ReRAM), research on fast ion transport phenomena and their applications became a hot topic. Amidst the materials required for these applications, oxygen ion conductors (OIC), mixed ionic-electronic conductors (MIEC) [37] and composites of both (OIC +MIEC) [38] are of particular interest.

In this field, micro solid oxide fuel cells (µ-SOFCs) are one of the most promising candidates as power sources for portable electronic devices, as they have been predicted to have the highest energy density and specific energy [START_REF] Evans | Review on microfabricated micro-solid oxide fuel cell membranes[END_REF][START_REF] Chiabrera | Micro solid oxide fuel cells: a new generation of micro-power sources for portable applications[END_REF][START_REF] Bieberle-Hütter | Syngas generation from n-butane with an integrated MEMS assembly for gas processing in micro-solid oxide fuel cell systems[END_REF][122]. In these devices the catalytic reduction of oxygen ions takes place in the cathode, which requires porous materials with both high oxygen surface exchange and diffusion coefficients at operating temperatures, typically above 400 °C. [START_REF] Evans | Review on microfabricated micro-solid oxide fuel cell membranes[END_REF][122][START_REF] Adler | Factors Governing Oxygen Reduction in Solid Oxide Fuel Cell Cathodes â[END_REF][START_REF] Carter | Factors limiting the low temperature operation of SOFCs[END_REF] Since the oxygen reduction reaction (ORR) is a limiting factor for lowering the operation temperature, its improvement through different strategies has been broadly studied in the last times [START_REF] Sun | Cathode materials for solid oxide fuel cells: a review[END_REF].

The surface exchange kinetics and ionic conductivity can be improved by introducing strain in the lattice [42][43][44], as reported for a large number of materials [42][43][44][45][46][START_REF] Wen | Interfacial lattice-strain effects on improving the overall performance of micro-solid oxide fuel cells[END_REF], including 3.2 Oxygen transport in nano-architectured thin films 51 many perovskite oxides. In addition, the role of dislocations in the ion diffusion has also been studied for several epitaxial perovskite films [47][48][49]. For SrTiO 3 , dislocations normal to the oxygen diffusion direction slow down the ion transport, while those parallel to the diffusion direction do not accelerate it [47,48]. In contrast, in the case of La 0.8 Sr 0.2 MnO 3 (LSM) manganites, it has been claimed that dislocations in the form of half loops starting parallel to the oxygen diffusion direction, accelerate oxygen transport [49]. Recently, it has also been reported for several polycrystalline perovskite compositions [50][51][52] that grain boundaries (GBs) highly enhance oxygen exchange and diffusion. Furthermore, the optimization of electrode nano-architectures [39][40][41] through the control of the film porosity, thickness and the "decoration" of the electrode surface lead to a maximization of the surface-volume relation which can also favor gas diffusion [41]. Complementary, strategies such as catalyst infiltration [START_REF] Ding | Enhancing SOFC cathode performance by surface modification through infiltration[END_REF] also lead to the enhancement of the catalytic activity of the electrodes.

In this work, we designed nano-architectured films and investigated their ionic mass transport properties using the parent compound LaMnO 3 (LMO) perovskite as a promising oxygen electrode for µ-SOFCs. LMO possesses a distorted perovskite structure, consisting of corner sharing MnO 6 oxygen octahedra with a central Mn atom, and La cations occupying the 12fold coordination site formed in the middle of the cube of eight such octahedra. The crystal structure of the stoichiometric LMO undergoes a phase transition at a temperature close to ~477 ºC [START_REF] Rodríguez-Carvajal | Neutron-diffraction study of the Jahn-Teller transition in stoichiometric LaMnO3[END_REF][START_REF] Rudskaya | Phase transitions in manganese-containing perovskites[END_REF] from a rhombohedral (R-3c) symmetry (higher temperature phase) to an orthorhombic (Pnma) symmetry (lower temperature phase). Furthermore, upon stoichiometry changes (La, Mn and/or O content variations) this orthorhombic phase can transform at room temperature to a rhombohedral one, concomitant to a Mn 3+ to Mn 4+ valence change (particularly Mn 4+ percentage of 20% or larger [START_REF] Rodríguez-Carvajal | Neutron-diffraction study of the Jahn-Teller transition in stoichiometric LaMnO3[END_REF][START_REF] Wold | Preparation and crystallographic properties of the systems LaMn1-xMnxO3+λ and LaMn1-xNixO3+λ (JA-1106)[END_REF]).

In this line, in a previous study we developed different deposition strategies which allow to tune the oxygen content in LMO polycrystalline films [START_REF] Rodriguez-Lamas | Integration of LaMnO 3+δ films on platinized silicon substrates for resistive switching applications by PI-MOCVD[END_REF]. We proved that the evolution of the crystal structure of the symmetry from orthorhombic to rhombohedral was bound to the increase in Mn +4 content [START_REF] Rodriguez-Lamas | Integration of LaMnO 3+δ films on platinized silicon substrates for resistive switching applications by PI-MOCVD[END_REF].

Here, we combined strain, defects and film nano-architecture effects of epitaxial LMO films to enhance their ionic conduction and oxygen exchange kinetics for µ-SOFC cathodes. Defect-free and defect-rich epitaxial LMO films were obtained by growing by pulsed injection metal organic chemical vapor deposition (PI-MOCVD) on two different substrates: SrTiO 3 (STO) and LaAlO 3 (LAO). In addition, by tuning the deposition parameters we prepared "finger-like" nano-architectured films (thicker than 15 nm) to obtain a large exposed surface area, highly desirable for cathodes. Their oxygen exchange and diffusion properties close to 500 and 600 °C were evaluated by the Isotope Exchange Depth Profile (IEDP) methodology combined with Time-of-Flight Secondary Ion Mass Spectrometry (ToF-SIMS). The influence of strain, defects and dislocations on the oxygen mass transport properties was evaluated. Combining both, strain and nano-architecture engineering, we show how the incorporation of oxygen into the LMO film and the oxygen ion transport can be smartly maximized.

Experimental

Thin film deposition

LMO thin films were grown by PI-MOCVD in a JIPELEC PI-MOCVD reactor. This technique allows to grow uniform and conformal films over large areas. Here, we deposited LMO films with a controlled thickness and composition on multiple 10x10 cm 2 single crystal STO (001) and LAO (001) substrates (from CrysTech) within a circular area of 4 cm diameter. The film growth method is described in the previous work by Rodriguez-Lamas et al. [START_REF] Rodriguez-Lamas | Integration of LaMnO 3+δ films on platinized silicon substrates for resistive switching applications by PI-MOCVD[END_REF], while the specific deposition conditions are given as follows: LMO films were grown at 700 °C in a 50% Ar and 50% O 2 atmosphere with a total pressure of 5 Torr using a precursor solution of La(thd) 3 and Mn(thd) 3 in m-xylene (where La/Mn = 2) with a total metallic precursor concentration of 0.0225 M. The optimized pulse injection parameters were a 2.5 Hz frequency and a 2 ms electrovalve opening time. After growth, the samples were cooled down to room temperature in an Ar atmosphere for the dense thin films and in an O 2 atmosphere for the nano-architectured films with controlled temperature ramps.

Structural characterisation

The structural characterisation of the LMO film was performed by X-Ray diffraction (XRD) to know the phase structure and orientation of the films, Raman spectroscopy to complement the phase structure study, scanning electron microscopy (SEM) to reveal the surface morphology and homogeneity, atomic force microscopy (AFM) to measure the roughness of the films, and transmission electron microscopy (TEM) and high resolution TEM (HRTEM) to observe the crystal structure and microstructure of the film cross-section.

XRD characterization was performed in Bragg-Brentano θ-2θ configuration in a Bruker D8 Advance serie II diffractometer, with monochromatic CuKα1 radiation (λ = 1.5406 Å) and 1D detector (Bruker Lynxeye) with a 0.06°resolution in 2θ. The diffractogram was collected from 20 to 110°(2θ) with a step of 0.011°and 2 s by step counting time. The sample was spinning during the measurement.

Raman spectra were recorded using a Jobin Yvon/Horiba Labram spectrometer equipped with a liquid nitrogen cooled CCD detector. The 488 nm excitation line of an Ar + laser was focused to a 1 µm 2 spot on the sample using an objective x100. The laser power on the sample surface was close to 1 mW. Spectra were calibrated at room temperature using a silicon reference sample with a theoretical position of 520.7 cm -1 .

A SEM FEG Gemini SEM 500 was used for sample morphology characterization. Images were acquired with an energy of 3 kV in high vacuum mode and at a working distance of 2.6 mm.

AFM topography scan in tapping mode was performed in an AFM (Veeco) Digital Instruments Dimension 3100 through the Nanoscope interface, and the Gwyddion software was used for data treatment (extraction of the root mean square, rms, roughness data).

Cross-sections of the samples were prepared by the semi-automated polishing tripod technique with the MultiPrepTM system (Allied High Tech Products, Inc.). The PIPS II from GATAN system was used for the final polishing. TEM and HRTEM images were recorded with a JEOL JEM 2010 LaB6 microscope operating at 200 KV with a 0.19 nm point-to-point resolution. In order to analyze the strain induced in the epilayer by the lattice mismatch from the substrate, high resolution TEM observations were carried out with using the Geometric Phase Analysis (GPA) method. [START_REF] Hÿtch | Quantitative measurement of displacement and strain fields from HREM micrographs[END_REF]

Isotope exchange depth profiling technique

For 18 O/ 16 O isotope exchange experiments, the samples were placed in an exchange rig equipped with a quartz tube. They were pre-annealed in pure oxygen of natural isotopic abundance (0.20 % of 18 O) at ~200 mbar for 1 h at the same temperature as the subsequent exchange experiment. In order to avoid crack formation in the samples, the furnace was heated and cooled using a 5 °C/min rate. Once back to room temperature, natural oxygen was pumped out. Then, 18 O enriched oxygen (P total ~200 mbar) was introduced into the tube and 18 O/ 16 O exchanges were carried out with the same heating/cooling rate. Two different intermediate temperatures in the range of µ-SOFCs operation were selected, i.e.504 °C and 609 °C with 8 h and 1 h plateaux, respectively. For the sake of simplicity, in the future description we will consider 500 and 600 °C as the exchange temperatures. Note that the effective 18 O exchange time was calculated taking into account the finite increase/decrease in the sample temperature. [START_REF] Killoran | The Effective Duration of a Linear Slow-Cool[END_REF] To quantify the 18 O isotope content in the gas phase, prior to the experiments, a piece of silicon wafer was annealed at 1000 °C for 3 h to grow a SiO 2 layer, which was analysed by SIMS, obtaining an 18 O fraction of 0.895. For further structural analysis, some pristine pieces of the same samples were annealed in air using an equivalent thermal history to that of the isotopic exchanged samples (adding the times corresponding to the pre-anneal in 16 O and to the exchange in 18

O).

A ToF-SIMS V instrument (ION-TOF GmbH Germany) equipped with a Bi liquid metal ion gun (LMIG) for analysis and a caesium gun for sputtering was used for the oxygen isotope analyses. Negative secondary ions were collected and the data were obtained by 6 pulses burst mode operation, applying Poisson correction. Charge effects were compensated by means of a 20 eV pulsed electron flood gun. A 300x300 µm² surface area was sputtered with the Cs + ion beam (1 keV, 85 nA) before analysis with the Bi 3+ primary ion beam (25 keV, 0.25 pA) rastered over a 100x100 µm² surface area. Several SIMS depth profiles were measured for each of the exchanged samples to ensure reproducibility.

In order to estimate the diffusion coefficients (D*) and exchange coefficients (k*) of the dense LMO thin films, the oxygen concentration profiles were analysed using the analytical solutions to the diffusion equation for a plane sheet model developed by Crank [START_REF] Crank | The Mathematics of Diffusion[END_REF] (see Supplementary Information for further details) and fitted by least squares. In addition, simulations using a finite element modelling (FEM) were used to simulate the oxygen concentration profile when the assumption of a single homogeneous film was not valid. The geometry of the model consists of a columnar cell made in agreement with the average grain size values and thickness measured of the LMO films by AFM and HRTEM, respectively (see Fig.

Results and Discussion

Structural characterisation of the LMO films

A comprehensive microstructural characterization of the LMO grown by PI-MOCVD at 700 °C on single crystal STO and LAO substrates is presented in this section.

Epitaxial dense films

First, films were deposited by the injection of a small number of pulses of precursor solution (2500) to obtain very thin dense LMO layers. While the thin film deposition occurs at a temperature above the phase transition (T JT ≈ 477 °C) [START_REF] Rodríguez-Carvajal | Neutron-diffraction study of the Jahn-Teller transition in stoichiometric LaMnO3[END_REF][START_REF] Rudskaya | Phase transitions in manganese-containing perovskites[END_REF], during the cooling of the sample a cooperative Jahn-Teller (JT) transition is expected to take place in LMO, resulting in an orthorhombic crystallographic structure with highly distorted MnO 6 octahedra.

In the diffractograms of the as-deposited LMO films grown on STO (Fig. 3.7a) and on LAO (Fig. 3.7b) only the 00l reflections of pseudocubic LMO can be observed in addition to the substrate reflections, confirming the good crystalline quality, single phase and single orientation of the film. In the case of LMO/STO the film and substrate reflections strongly overlap due to the matching cell parameters between the film and the substrate. For simplicity throughout the article the LMO films have been indexed using the pseudocubic system, although the Pnma space group characteristic of stoichiometric LMO bulk sample (International Centre for Diffraction Data (ICDD) pattern 01-072-0840) and equivalent to the Pbnm space group also used in some cases in the literature is used as the pattern of reference. However, as will be later explained, in some cases the LMO films could be composed of a mixture of orthorhombic (Pnma) and rhombohedral (R-3c) phases. As schematically represented and explained in Pomar et al. [START_REF] Pomar | Formation of Self-Organized Mn3O4 Nanoinclusions in LaMnO3 Films[END_REF], the Pnma orthorhombic LMO cell parameters a = 5.743 Å, b = 7.695 Å and c = 5.537 Å correspond to Mn-Mn distances of a pc = c pc = 3.989 Å and b pc = b/2 = 3.847 Å in pseudocubic cell notation. When grown on single crystal perovskite substrates, in order to accommodate the JT structural distortion the LMO can present 2 different orientations at room temperature, which correspond to the orthorhombic Pnma b parameter in-plane or out-of plane [START_REF] Pomar | Formation of Self-Organized Mn3O4 Nanoinclusions in LaMnO3 Films[END_REF].

The thin dense LMO thin films were then exchanged in 18 O at 500 °C and 600 °C to evaluate the oxygen surface exchange and transport across the LMO films. The LMO/STO and LMO/LAO films were measured in top view by SEM and AFM after the 18 O exchange at 600 °C (Fig. 3.7c-d). Both analyses reveal grain sizes in the nanoscale ( 10-12 nm) and a very flat and homogeneous surface morphology with a film rms roughness of only 0.33 ± 0.03 nm for LMO/STO and 0.40 ± 0.03 for LMO/LAO. The cross-section TEM images presented in Fig. 3.7e-f demonstrate full dense and homogeneous LMO films with thicknesses of 12.3 ± 0.3 nm and 11.2 ± 0.4 nm for STO (Fig. 3.7e) and LAO (Fig. 3.7f), respectively. However, in contrast to the high quality films grown on STO, the films on LAO exhibit a large amount of defects at this low magnification scale.

Nano-architectured thin films

The active surface area of the LMO films was maximized by growing thicker films (6000 pulses) with a "finger-like" microstructure, while keeping a dense continuous film at the bottom part in contact with the substrate. This nano-architecture aims to improve the catalytic properties of the LMO film (cathode) by largely increasing the surface/volume relation. Fig. 3.8 presents the structural analysis performed on the nano-architectured thick films. As for dense thin films, the diffractograms of LMO/STO (Fig. 3.8a) and LMO/LAO (Fig. 3.8b) are only composed of the 00l reflections of pseudocubic LMO and the correspondent substrate reflections. Similarly, the film and substrate diffraction peaks of LMO/STO strongly overlap , which is not the case for the LMO/LAO films, allowing a more accurate peak deconvolution. The 002 LMO diffraction peak is fully asymmetric and can be fitted with two components. Hence, two different pseudocubic cell parameters (Mn-Mn distances) can be extracted from the patterns leading to 3.89 and 3.92 Å for LMO/STO and 3.91 and 3.95 Å for LMO/LAO. These nano-architectured LMO films were exchanged in 18 O at the same time as the epi- taxial dense thin films, and thus were submitted to exactly the same temperature/time/pO 2 history. SEM secondary electron microscopy images after the 600 °C exchange of both nanoarchitectured thin films can be seen in Fig. 3.8c (LMO/STO) and Fig. 3.8d (LMO/LAO), where the top view shows the presence of individual non-attached columnar squared grains. Both films present homogeneous features in large areas with comparable grain dimensions (13-14 nm for the squares sides). The topography extracted by AFM, seen on the right side of the images, indicates that the roughness is much higher in these nano-architectured films (Fig. 3.8c-d) comped with the dense thin films (Fig. 3.8c-d). The rms roughness values obtained are 1.3 ± 0.1 nm for LMO/STO and 1.8 ± 0.8 nm for LMO/LAO. However, it should be pointed out that, the real roughness is indeed larger than these values. As will be shown in cross-section TEM analysis, the gap between the columns is too small to allow the AFM tip to reach the whole depth and thus the obtained rms values obtained from the AFM measurements are underestimated.

TEM cross section images of the thick nano-architectured LMO/STO (Fig. 3.8e) and LMO/LAO (Fig. 3.8f) films revealed the existence of a remarkable "finger-like" microstructure which consists of a continuous dense LMO film at the bottom covered by the growth of columns (the so-called "fingers") separated by an intercolumnar gap of 1-2 nm. The LMO bright field images (Fig. 3.8e-f) show films composed of columns with rounded top surfaces, presenting a homogenous density and microstructure over large areas. The exact thickness of the continuous dense part of the LMO film varies depending on each two consecutive columns considered. In average, a larger thickness of around 23 nm is observed in average for the LMO/STO film (Fig. 3.8e), while this average thickness is close to 15 nm for the LMO/LAO films (Fig. 3.8f). This observation can be easily explained by the smaller mismatch between the LMO film and the substrate for the STO substrate, compared to the LAO, allowing for a more continuous and continuous growth until the strain is released by the formation of separated columns. It can also be appreciated that in both cases not all the grains reach the maximum height, being some of the "fingers" shorter than the rest of the LMO columns. Finally, the maximum height of the columns of the LMO films grown on LAO (~66 nm) are in average longer those of the films grown on STO (~54 nm).

Oxygen mass transport properties of the LMO films

As previously explained the LMO films were exchanged in 18 O at 500 °C and 600 °C to evaluate the oxygen surface exchange and transport across the LMO layer (see sketch of the isotope exchange process in Fig. 3.9a and d). These temperatures allowed us to investigate the LMO thin films in plausible operating conditions that are commonly used for µ-SOFC 18 O isotopic fraction for LMO films grown on STO and LAO: dense thin films exchanged b) at 500 °C for 8 h and c) at 600 °C for 1 h; "finger-like nano-architectured" films exchanged e) at 500 °C for 8 h and d) at 600 °C for 1 h. The magenta diamond pattern indicates the film/substrate interface region which is not taken into account for the fitting due to the intermixing of ion species from the LMO film and the substrate. The blue pattern stands for the dense part of the film (either totality of thin dense film, or continuous part of nanoarchitectured films).

applications. [122,[START_REF] Pla | Is it possible to design a portable power generator based on micro-solid oxide fuel cells? A finite volume analysis[END_REF] It is worth noting that for each temperature, all four samples were exchanged simultaneously, i.e. the dense thin and the nano-architectured thick LMO films, on both STO and LAO substrates. After the isotopic exchange, the 18 O concentration profiles through the LMO/STO and LMO/LAO films were measured by ToF-SIMS. Two depth profiles showing the secondary ion species of interest (LaO - 2 , MnO -, TiO -, 18 O -and 16 O) obtained for LMO/STO and LMO/LAO at 500 °C for 8 h are shown in Fig. SI 3.17 in the Supplementary information as representative examples.

The normalized 18 O isotopic fraction C' depth profiles calculated from the SIMS measurements for the LMO films exchanged at 500 °C are plotted in Fig. 3.9b and Fig. 3.9e (for the dense thin and nano-architectured thick films, respectively), and those for the films exchanged at 600 °C in Fig. 3.9c and Fig. 3.9f (for the dense thin and nano-architectured thick films, LaMnO 3+δ growth on single crystal substrates and oxygen ion transport studies. respectively). C'(x) is calculated according to Equations 1 and 2:

C(x,t) = ( 18O

18 O+ 16 O ) Equation ( 1) where C(x, t) is the intensity ratio between the 18 O -signal measured by SIMS and the total ( 18 O -+ 16 O -) oxygen ion intensity in the material, x is the distance from the gas/LMO thin film interface and t is the time of isotopic exchange. 2) where C bg corresponds to the background 18 O isotope fraction (0.002 for natural abundance) and C g stands for the 18 O concentration in the exchange gas.

C ′ (x,t) = C(x,t)-C bg (C g -C bg ) Equation (
Oxygen tracer diffusion (D * ) and surface exchange (k * ) coefficients were obtained by nonlinear least squares regression curve fitting of the measured diffusion profiles to the solution of Fick's second law solved for diffusion in an infinite solid slab by Crank [START_REF] Crank | The Mathematics of Diffusion[END_REF]. The corresponding equations and the boundary conditions for the plane sheet model are detailed in the Supplementary information (schematics of the model shown in Fig. SI 3). Note that as the penetration of the Bi 3+ ion beam reaches several nanometers, a mixture of oxygen ions from the film and the substrate is measured when the interfacial zone is reached. Accordingly, the region of 1 nm closest to the substrate was not taken into account for the fit as its profile is not meaningful of the diffusion process.

To validate the profiles achieved, several concentration profiles were measured in different regions of the sample obtaining always similar results.

Epitaxial dense films

The normalized concentration profiles for the LMO/STO dense thin films (red squares in Fig. 3.9b and Fig. 3.9c) show curved profiles, with a decrease in concentration from the top surface towards the bottom interface, as expected.

In the case of LMO dense thin films grown on STO and exchanged at 500 °C (red squares in Fig. 3.9b), the values obtained from the fit are: D* = 3.8 ± 0.2 x 10 -17 cm 2 /s and k* = 1.8 ± 0.4 x 10 -11 cm/s. In the case of LMO dense thin films grown on LAO and exchanged at 500 °C (blue circles in Fig. 3.9b), the profile shows a high normalized isotopic concentration close to 0.78. The profile is rather flat throughout the majority of the film thickness, with a decrease in 18 O concentration at the top surface and at the bottom interface. The much higher isotopic concentration reached, together with a much flatter profile when compared to the film grown on STO substrate, indicate that both the oxygen surface exchange and diffusion coefficients are much higher for the film grown on LAO. The effective surface exchange coefficient obtained from the fit of the concentration profiles is k* = 5.8 ± 0.5 x 10 -11 cm/s, a value 3 times larger than for the LMO film grown on STO. Although it was not possible to obtain an exact D* coefficient due to the shape of the concentration profile, the fitting was also carried out taking into account only the film region with a slight curvature (bottom part of the film) in order to obtain the minimum D* value (Dmin*) in that region. This led to a Dmin* of 1.2 ± 0.1 x 10 -16 cm 2 /s, already four times larger than that measured for LMO/STO at the same temperature. However, the real diffusion coefficient of the film is expected to be larger, but would require much thicker dense films or extremely short exchange times to be precisely measured.

Complementary, the values fitted from the normalized concentration profiles of the LMO films exchanged at 600 °C (curves shown in Fig. 3.9c) are D* = 6.2 ± 0.4 x 10 -17 cm 2 /s and k* = 2.6 ± 0.1 x 10 -10 cm/s, and D* = 7.8 ± 0.5 x 10 -17 cm 2 /s and k* = 3.1 ± 0.1 x 10 -10 cm/s for LMO/STO and LMO/LAO, respectively. Hence, the coefficients for LMO/LAO are slightly higher than those extracted from the fit of the LMO/STO profiles at the same temperature. It should be pointed out that for the LMO/STO dense thin films, the fitted curve slightly deviates from the experimental data, as shown for the 600 °C exchange in Fig. SI 3.18. Deviations from the ideal concentration profile have also been previously reported for other epitaxial films such as Sr-doped LMO (LSM) [49,50] or La 2 NiO 4 [START_REF] Burriel | Anisotropic oxygen diffusion properties in epitaxial thin films of La2NiO4+δ[END_REF]. These deviations can be related to the presence of adsorbed oxygen at the top surface, strain-related differences as a function of depth, fast diffusion through dislocations or to the influence of the exchange/diffusion process during the cooling of the sample. However, as a first approximation, this simple model has the advantage of allowing the extraction of the effective oxygen transport parameters and the direct comparison between films grown on different substrates. Going a step further, an important difference in the shape of the two curves can be noticed close to the interface with the substrate (see blue and red curves in Fig. 3.9c), with a marked change in slope and steeper profile observed at about 3 nm from the LMO/LAO film interface. To take into account this behavior, additional FEM simulations were carried out including two regions: a 3 nm thick region corresponding to the "interface LMO" region close to the interface with the LAO substrate and a "bulk LMO" region far from the interface. The continuous black line in Fig. SI 3.18 represents the result of the finite element simulation with the two layers. This results in a first Di* = 3.16 x 10 -17 cm 2 /s for the interfacial region and a second Dbulk* = 8.47 x 10 -17 cm 2 /s, (more than 2.5 times larger than Di*). The origin of regions with different diffusion coefficients will be discussed in the following section.

The average k* and D* values obtained for the diffusion profiles measured for the dense thin films at 500 and 600 °C are presented in Arrhenius plots in Fig. 3.10, together with values reported in literature for LMO pellets with different La/Mn stoichiometry. In the case of the surface exchange coefficient k* (Fig. 3.10a), the values obtained for the LMO/STO films fit perfectly the tendency shown for LaMnO 3 and La 0.9 MnO 3 pellets [START_REF] Berenov | Oxygen tracer diffusion in undoped lanthanum manganites[END_REF] (by low temperature extrapolation). However, one value stands out in the plot. As previously mentioned, the k* measured for LMO/LAO at 500 °C is almost three times higher than that measured for LMO/STO and it is also higher than that expected for the polycrystalline bulk samples as shown in Fig. 3.10a. Fig. 3.10b shows that the diffusion coefficients at 600 °C have (similar) slightly higher values than those expected from LaMnO 3 and La 0.9 MnO 3 pellets [START_REF] Berenov | Oxygen tracer diffusion in undoped lanthanum manganites[END_REF] by low temperature extrapolation following the same slope (activation energy). However, at 500 °C extraordinarily fast D* values have been measured for both films, larger by at least one order of magnitude than those expected by simple low temperature extrapolation. For LMO/STO, the values are close to those measured at 600 °C, while they are even larger for LMO/LAO. It should be noted that in the case of the LMO/LAO, the plotted D* values correspond to a minimum D* in the bottom part of the film, while the average D* value throughout the film could not be obtained in this case as the profile is too flat to be measurable, due to a predictable high D* value. The origin of the very high diffusion coefficients measured at 500 °C will be discussed in the following section.

Nano-architectured films

The thick nano-architectured LMO films were submitted to the same exchange protocols at 500 °C and 600 °C as those previously presented for thin dense layers. The 18 O normalized isotopic fraction profiles measured by SIMS are shown in Fig. 3.9e (exchange at 500 °C) and Fig. 3.9f (exchange at 600 °C). The curves present a humped shape in all cases. The maximum concentration is found at certain distance from the bottom interface which is somewhat larger for the LMO/STO films, i.e. of around 18 nm for LMO/STO and 16 nm for LMO/LAO for the samples annealed at 500 °C, and of 28 nm for LMO/STO and 22 nm for LMO/LAO for the samples annealed 600 °C.

In this case, the exchange and diffusion processes cannot be explained by a simple geometrical model. As we have previously seen, the films are composed of a thick dense region covered by columnar "fingers" in which both thicknesses vary for different grains. As shown in Fig. 3.9d, the exchange process will not only take place on the top surface of the film, but also at the lateral side of all the columnar grains, as well as at regions at the "feet" of the grains between the columns. In addition, importantly, as previously observed in the TEM images of the films (Fig. 3.8e-f) the columns are not all of the same height and therefore, when sputtering through the film the contribution from columns of different height will successfully appear. Each of the columns from top to bottom is expected to have first region ("finger") governed by the surface exchange in the lateral sides of the column, in which the concentration profile is expected to be quite flat. Next, once the dense part of the film is reached, the concentration profile is expected to be governed by the oxygen diffusion, as in the dense films (as in Fig. 3.9b-c). Globally, when a large area composed by columns of different height with different basal dense thickness is sputtered and analysed jointly, a complex humped profile is expected. In addition, it should be noted that due to the inter-columnar porosity, the SIMS sputtering process will not be homogenous throughout the whole surface area, and thus the thickness calibration of the concentration profiles in Fig. 3.9e-f are just an average calibration. It should be noted that humped profiles have been previously reported for epitaxial LSM thin films grown on LAO. [49] In that case, the shape of the profiles was explained and fitted by a model that included half-loop dislocations with fast oxygen diffusion across and parallel to the film (U-shape).

It is important to notice that the oxygen concentration profiles of the dense thin (blue background region in Fig. 3.9b-c) and nano-architectured thick LMO films in the bottom 12 nm (blue background region in Fig. 3.9e-f) practically overlap for both temperatures and both substrates, confirming the oxygen diffusion governing in this region of the film.

In order to compare the total amount of 18 O introduced into the LMO films, the areas beneath the concentration profiles have been integrated, as they represent the oxygen tracer amount incorporated into the film (see Table SI 3.5). As expected, these values are larger for the thicker films, showing that a larger active surface area obtained from the "finger-like" morphology, allows for a much larger total oxygen content to be introduced into the films.

Understanding enhanced oxygen mass transport properties

We have shown that the oxygen concentration profiles and the amount of oxygen incorporated into the LMO films depend strongly on the substrate used as template for the LMO deposition.

While the shape of the profiles and oxygen transport coefficients of the films grown on STO at 600 °C are close to the values extrapolated from high temperature measurements for polycrystalline samples, the LMO/LAO films presented more complex profiles composed of different regions, and slightly higher k* and D* values. However, the most remarkable observation is the ultra-fast oxygen diffusion measured at 500 °C, being at least one order of magnitude higher than expected for films on both substrates, and considerably higher for LMO/LAO. At 500 °C, the LMO films on this substrate also present a remarkably high surface exchange coefficient.

In contrast to the numerous studies related to the oxygen transport properties in polycrystalline [50,51] and epitaxial [49] LSM thin films, to the best of our knowledge these properties had never been studied previously in LMO thin films. In both polycrystalline and epitaxial LSM films, structural defects parallel to the diffusion direction are considered as the origin of the enhanced k* and D* values, and of the appearance of two distinct slopes in the diffusion profile. The profile fitting was carried out using a model with two parallel diffusion paths: bulk and grain boundaries for the polycrystalline films [50,51], and bulk and dislocations for the LSM epitaxial films grown on LAO [49]. In the latter case, an additional dislocation-free LSM interfacial region with a different diffusion coefficient had to be added in the model to properly fit the profiles.

With the aim of explaining the very high diffusion coefficients measured at 500 °C, and the differences found between LMO films grown on different substrates, an in-depth structural comparative analysis of the films before and after the annealing treatments was carried out.

Epitaxial dense films

The HRTEM images (Fig. 3.11) confirm this remarkable difference in the films nanostructure depending on the substrate. Fig. 3.11a presents a HRTEM bright field image and the selected area electron diffraction (SAED) pattern of the LMO film on STO, showing a very quality and defect-free film with a perfect cube on cube (perovskite on perovskite) growth. Contrary, the HRTEM bright field image in Fig. 3.11c confirms the elevated density of defects in the LMO/LAO thin films. In this case a defective interface between the dense LMO film and the LAO substrate can be observed at this high magnification level.

First, we try to explain the origin of the differences between the theoretical and real oxygen concentration profiles for LMO/STO (see red squares and red dotted line as example of fit in Fig. SI 3.18). Indeed, the oxygen concentration profiles measured for the LMO/STO films (Fig. 3.9b and c) can be well fitted by FEM using models with different regions: either parallel diffusion paths corresponding to dislocations/subgrain boundaries and bulk, or two regions in series with different D* values. However, for the thin dense LMO/STO films the selective area electron diffraction (SAED) pattern along the [001] STO zone axis (inset in Fig. 3.11a) shows that the film is clamped on the substrate with Qx,STO = Qx,LMO (and compensated by Qy,STO < Qy,LMO according to a negative Poisson coefficient see also Fig. 3.7a for XRD). The inset in Fig. 3.11b shows the Fourier transform extracted from Fig. 3.11a in which a mask is applied to the diffraction spots corresponding to the 010 STO , 101 LMO , 0-10 STO and -10-1 LMO reflections. The inverse Fourier transform of the masked region (Fig. 3.11b) shows only the diffraction planes normal to the interface and proves the very good parallelism of the planes across the interface and throughout the whole film, and thus confirms the absence of defects such as dislocations or subgrain boundaries. As the LMO thin film deposited on STO can be considered as a dense defect-free film showing epitaxial cube on cube growth on top of the STO substrate, we conclude that we have not found any experimental structural evidence to support the FEM model with parallel diffusion pathways for the LMO/STO films.

Next, in order to verify if the LMO/STO film consists of two regions in series with different D* values related to different strain states (close and far from the bottom interface), the GPA method has been carried out in the HRTEM images of the LMO/STO dense thin films (Fig. 3.12). The corresponding GPA strain (ε) images shown in Fig. 3.12b and c are calculated based on the following equations, respectively and by using a Gaussian mask of g/4 on the 100 spatial frequency of STO:

ε GPA zz = d 10 1 LMO -d 100 ST O d 100 ST O Equation (3) ε GPA xx = d 101 LMO -d 010 ST O d 010 ST O Equation (4)
, where dhkl are the d-spacings of the related planes in LMO and STO.

The average calculated GPA out-of-plane strain is found equal to 2.2 ± 1% (Fig. 3.12b). The theoretical corresponding GPA strain (relative to STO) of relaxed LMO is expected to be equal to 1.5%. In Fig. 3.12c, the GPA strain in the basal plane of LMO on STO is found to be negligible (zero in average). This means that the {101} of LMO are elastically strained (compressed) by 1.5% in order to adapt the d-spacing of {010} of STO. Since the basal-plane lattice parameter of LMO is compressed by 1.5% in order to adapt the corresponding lattice constant of STO, hence, following the bi-axial stress formulation [START_REF] Hirth | Theory of Dislocations (2nd ed.)[END_REF], the perpendicular lattice constant is expanded by 0.7%. Within the experimental error, both the out-of-plane and in-plane strains are constant throughout the LMO film thickness and thus a FEM model with two distinct regions in depth with different D* in depth does not represent the structural observations obtained by GPA analysis.

Thus we propose that the observed small deviations from an ideal diffusion profile, which mainly occur in the first few nanometers of the LMO/STO film (see Fig. On the one hand, a surface space-charge layer depleted of oxygen vacancies would lead to a lower diffusion coefficient due to a lower amount of mobile defects (lower concentration of oxygen vacancies) [START_REF] De Souza | Probing Diffusion Kinetics with Secondary Ion Mass Spectrometry[END_REF]. On the other hand, the presence of La and Mn cation vacancies in LMO has been calculated by molecular dynamic simulations, which strongly inhibits the longrange diffusion of oxygen vacancies (D v ). [START_REF] Börgers | The surprisingly high activation barrier for oxygen-vacancy migration in oxygen-excess manganite perovskites[END_REF] Finally, we cannot discard a small change in the oxygen concentration profile shape due to the influence of the exchange/diffusion process during the cooling of the sample.

Next, to understand the large differences found at 500 °C for the films grown on different substrates, we focus on the HRTEM bright field image of the LMO film on LAO films shown in Fig. 3.11c. The Fourier transform analysis of this image, shown in Fig. 3.11d, confirms the presence of a large number of defects in this film in the form of slinky lines with occasional discontinuities corresponding to defective crystallographic planes perpendicular to the surface. These defects correspond generally to collective shifting of planes. Although defects are present throughout the whole film thickness, few of them appear in the first 3 bottom nanometers of the observed region, while their density is higher through the rest of the LMO film. We propose that the presence of these structural defects has a positive effect with higher oxygen concentration in the films (see Table ?? and Fig. 3.9 b and c) resulting from higher surface exchange and diffusion coefficients for the LMO films grown on LAO (Fig. 3.10). It should be noted that the shape of the LMO/LAO profiles are qualitatively similar to those measured by Navickas et al. for LSM epitaxial films [49], explained in their case by a model including the presence of half-loop dislocations. This can be justified considering the model reported by Bagués et al. [START_REF] Bagués | The Misfit Dislocation Core Phase in Complex Oxide Heteroepitaxy[END_REF] for LSM/LAO, where the formation and accumulation of oxygen vacancies is promoted around the dislocation core area. Moreover, according to Navickas et. al. [49], the increase of oxygen vacancy concentration leads to faster diffusion through the dislocations. In the same work [49] the authors measured an increase in the "bulk" diffusivity, together with an enhanced diffusivity along the dislocations, both of which increased with film thickness.

In the case of the LMO/LAO concentration profile for 600 °C, the change in slope occurring around 3 nm (blue circles in Fig. 3.9c), qualitatively represented and simulated by two regions with different diffusion coefficients (Fig. SI 3.18), would be related to a change from a defect-rich region (Dbulk*, top of the film) to a defect-poor strained region (Di*, interface, bottom of the film). The strain imposed by the substrate would not be totally relaxed by the formation of defects in the first 3 bottom nanometers, which would result in a lower diffusion coefficient in this interface region. On the other hand, the slightly "humped" shape of the LMO/LAO concentration profile for the 500 °C exchange (blue circles in Fig. 3.9b) can be qualitatively explained by the existence of regions with fast diffusion (defect-rich), which do not reach the bottom interface, parallel to other regions with lower diffusivity. The tracer oxygen would reach the center of the LMO film through the fast-diffusion regions, resulting in a high normalized isotopic concentration in the central region, and would then diffuse back towards the top, resulting in a lower 18 O concentration at the top surface. Although a FEM including the presence of half-loop dislocations would also explain these findings, in our case we do not have the proof of this type of defects in the LMO/LAO films, and thus here we have not applied this model. By the structural and functional comparison of the properties of both dense thin films, we confirm that the presence of defects present in the LMO films grown on a substrate with a large mismatch (LAO) enhances the oxygen mass transport properties of the LMO films. This effect is particularly large at 500 °C, as will be discussed later.

Nano-architectured Films

Next, we will focus on the HRTEM images of the "finger-like" nanostructured LMO films grown on STO and on LAO, which are shown in Fig. 3.13a and Fig. 3.14a, respectively, and in which the crystallographic planes of the perovskite structure can be observed. As already shown in Fig. 3.8e-f, the LMO films grown on both substrates are formed by a dense continuous film covered by columnar "finger-like" grains. Occasional changes in cell orientation in the columns are observed, supporting the existence of multiple orientations of the orthorhombic cell . Fig. 3.13b and Fig. 3.14b present the LMO/STO and LMO/LAO film SAED obtained along the [001] STO and [001] LAO zone axis, respectively. While the LMO and STO diffraction spots are almost superimposed in Fig. 3.13b, the diffraction array in Fig. 3.14b gives evidence of the orthorhombic Pnma structure for LMO as multiple orienta-tions of the orthorhombic LMO b parameter can be observed in the case of LAO substrate. The Pnma structure can also be identified in the HRTEM images in different regions of the columnar features. The fact that different orientations were observed in LMO/STO images without altering the superposition of their diffraction spots is in favour of a larger misorientation in LMO/LAO. Fig. 3.13c and Fig. 3.14c show high resolution bright field images of the LMO films on STO and LAO, respectively, where the continuous film region and the base of the columns can be differentiated. Similarly to Fig. 3.11b and Fig. 3.11c for the dense thin films, Fig. 3.13d and Fig. 3.14d show the Fourier transform of Fig. 3.13c and Fig. 3.14c, and the reciprocal space spots selected to obtain the inverse Fourier transform in which the planes normal to the interface can be observed as well as the dislocations.

A significant conclusion that can be extracted from the combination of the structural and oxygen transport properties of these thick films results is the importance of the combination of nano-architecture and defect engineering to optimize the functional properties of LMO. It is clear that the augmentation of the active surface area by the creation of a "finger-like" microstructure is a very effective way of increasing the total tracer amount of oxygen which is incorporated into the film (see Table SI 3.5). However, for the nano-architectured LMO films grown on STO, although a high concentration of oxygen enters the pillars of the film, only a small amount effectively reaches the substrate (red squares in the blue background region in Fig. 3.9e-f), as the transport is diffusion-limited in the dense region close to the interface. On the other hand, in the case of LMO/LAO, thanks to the presence of a high density of defects, the oxygen reaches more effectively the bottom interface by a faster diffusion, particularly high at 500°C (blue circles in the blue background region in Fig. 3.9e-f). These results are of major importance and have to be carefully considered for possible applications such as micro-SOFC or oxygen permeation membranes, in which the LMO film would be in direct contact with an electrolyte or a support material.

Origin of the high oxygen diffusion at 500 °C

As previously mentioned, at 500 °C the D* values measured for the LMO dense thin films are at least one order of magnitude higher than expected by the high-temperature activation energy (Fig. 3.10b). In order to understand the origin of the fast diffusion, Raman spectroscopy was carried out on the LMO/LAO films before and after the isotopic exchange treatments (Fig. 3.15). The as-deposited LMO films (Fig. 3.15a for the thin dense film and Fig 3 .15b for the thick nano-architectured film) present a main sharp mode centered at 612 cm -1 , characteristic of the LMO orthorhombic structure, related to a symmetric stretching vibration of the oxygen ions in the MnO 6 octahedra. This in-plane stretching vibration is determined by the Mn-O 2 distances and thus activated by Jahn-Teller (JT) distortions [START_REF] Iliev | Distortion-dependent Raman spectra and mode mixing in R MnO3 perovskites (R=La[END_REF]. It is over- lapped with a broad mode centered at 660 cm -1 , probably related to the LMO rhombohedral structure and also activated by JT distortions [START_REF] Iliev | Raman phonons and Raman Jahn-Teller bands in perovskite-like manganites[END_REF]. This signature seems to attest a mixed orthorhombic/rhombohedral structure at the local scale which cannot be observed using XRD.

After performing the exchanges in an 18 O-enriched atmosphere, a shift of the stretching mode towards smaller wavenumbers occurs. This shift is due to the replacement of 16 O atoms with 18 O atoms of larger mass. The peak * marked with an asterisk is not assigned. It might correspond to the Mn 2 O 3 phase as reported in literature [START_REF] Gao | A comparison study on Raman scattering properties of α-and β -MnO2[END_REF]. However the presence of Mn 2 O 3 has not been observed by any of the other characterization techniques neither in the annealed nor in the exchanged samples.

In addition to this, a clear evolution in the profile of the Raman lines is visible. The observed broadening and intensity decrease can originate from an increase in lattice disorder due to both the introduction of 18 O atoms in the oxygen lattice and to an increase in the oxygen content in the LMO lattice contributing to an evolution towards a more rhombohedral structure. Such oxygen excess would increase the concentration of Mn 4+ ions inside the lattice and thus decrease JT distortions and the intensity of the JT activated modes in Raman spectra. [START_REF] Dubey | Signature of JahnâTeller distortion and oxygen stoichiometry in Raman spectra of epitaxial LaMnO3+δ thin films[END_REF] Dubey et al. observed the described Raman spectra changes in epitaxial thin films of LMO grown by Pulsed Laser Deposition and annealed in different oxygen environment. In their work the increase in oxygen stoichiometry observed by Raman was also confirmed by a decrease in unit cell volume as observed by XRD [START_REF] Gao | A comparison study on Raman scattering properties of α-and β -MnO2[END_REF]. Although all the Raman spectra meas- ured in this work (Fig 3 .15a for the thin dense and Fig. 3.15b for the thick nano-architectured LMO films) present a combination of orthorhombic and rhombohedral modes, the relative intensity of these modes strongly depends on the temperature of the 18 O exchange. The differences are most clearly observable for the nano-architectured thick films (Fig 3 .15b) for which the rhombohedral component is much more pronounced after the exchange at 600 ºC, while the orthorhombic component is the main line observed after the exchange at 500 ºC. Hence the evolution of the symmetric stretching and the JT modes in Raman point towards an increase in oxygen content in the structure after the exchange at 600 °C, leading thus to a decrease in oxygen vacancy concentration in the LMO film.

To confirm this assumption, XRD measurements were carried out on pristine LMO/ LAO dense thin films reproducing the whole thermal history of isotopically exchanged samples, but in air. In Fig. 3.15c, no displacement of the LMO Bragg peak positions is observed when comparing the as-deposited film with the sample annealed at 500 ºC for 9h (to reproduce 1h+8h annealing). In contrast, a shift of the LMO peaks towards higher 2θ values is observed in the case of the sample annealed at 600 °C for 2 h, indicating a decrease in the out-of-plane interplanar distance. The pseudocubic out of plane parameters reported in Table 3.4 indicate no parameter change for the LMO/LAO film annealed at 500 °C and a decrease of 0.04 Å after annealing at 600 °C. Similarly, a cell parameter decrease of 0.06 Å is estimated for the film deposited on STO substrate and annealed at 600 °C in good agreement with an increase of the Mn 4+ , and thus less oxygen vacancies, in the structure.

It can thus be concluded that the surprisingly high oxygen diffusion coefficients measured at 500 °C (Fig. 3.10), mainly for LMO/LAO but also for LMO/STO, result from a higher concentration of oxygen vacancies (and more orthorhombic structure) at this lower temperature. On the other hand, at 600 °C the structure would incorporate more oxygen, become more rhombohedral, leading to a lower concentration of oxygen vacancies, and consequently to a lower diffusion coefficient. These D* are in perfectly agreement with the diffusion coefficients previously measured in the literature for bulk LMO samples at higher temperatures.

Conclusions

In this work we propose a global strategy to optimize the low temperatures (500-600 ºC) oxygen mass transport properties of functional perovskite LMO thin films. We show that exceptionally high oxygen diffusion can be obtained at 500 °C in epitaxial LMO dense thin films, which are oxygen vacancy rich. In addition, growing the LMO dense thin films on LAO single crystals (with a higher mismatch when compared to STO) leads to films with a higher concentration of defects. These defects are indeed favorable for the oxygen diffusion and surface exchange processes and lead to enhanced mass transport coefficients.

Furhtermore,nano-architectured LMO films composed of a continuous bottom dense thin film covered by "finger-like" features were designed to efficiently increase the surface/volume relation, resulting in a larger oxygen incorporation. Moreover, an optimized oxygen migration towards the bottom interface was achieved for the LMO/LAO films.

This unique combination of defect engineering and morphology nano-architecturing shows a great potential for the design of low temperature SOFC electrodes and could be extended in the future to other promising MIEC oxides to obtain optimized functional devices.

Supplementary Information

Model describing oxygen exchange and diffusion A model of an infinite solid slab extending over the region of thickness 2l is assumed. This plane sheet model gives a symmetrical solution for the plane sheet occupying the region -l<z<l. This solution apply also to the sheet 0<z<l when the face z = 0 is non-permeable (see Figure S2). In our case the impermeable surface corresponds to the film-substrate interface, through which no oxygen diffusion occurs. The assumption made is that the rate of exchange is proportional to the difference between the concentration in the gas and the concentration in the surface at any time. The Boundary condition:

-D * ∂ c * (l,t) ∂ z = k * [c * g -c * (l,t)] Equation (1)
And the initial boundary condition as expressed in equation 4:

∂ c * (l,t) ∂ z = 0 Equation (2)
Then the solution to the diffusion equation given by Crank [START_REF] Crank | The Mathematics of Diffusion[END_REF] for a plane sheet is as follows: 18 O isotopic fraction for dense thin films LMO films grown on STO and LAO exchanged at 600 °C for 1 h. The magenta diamond pattern indicates the film/substrate interface region which is not taken into account for the fitting due to the intermixing of ion species from the LMO film and the substrate. The blue grid pattern indicates the second region for the 2 region fit for LMO/LAO. The red (LMO/STO) and blue (LMO/LAO) dashed lines correspond to the best fit to Crank's solution to the diffusion equation for a solid slab. The black line corresponds to the FEM modelling solution with 2 different regions in the LMO film grown on LAO, a defect-rich region at the top of the film and a defect-poor region at the bottom.

c * -c * Bg c * g -c * Bg = 1- ∞ ∑ n = 1 2Lcos(β n z/l)exp(-β 2 n D * t/l 2 ) cosβ n (β 2 n +L 2 +L) Equation (3
Table 3.5 -Average isotopic fraction ( integrated area of the normalized 18 O isotopic fraction divided by the film thickness) of the LMO/STO and LMO/LAO films exchanged at 500 and 600 °C.

Introduction

Resistive switching (RS) denotes the phenomena occurring in capacitor-like heterostructures (metal-insulator/semiconductor-metal, MIM), namely memristors, when a non-volatile change of resistance is produced under the effect of an applied current or electric field [18]. As these resistance changes are reversible, RS is suitable for redox-based resistive switching random access memory (Re-RAM) applications, where different resistance values can be written, read and erased by applying the appropriate voltages. Typically two different states are characterized per device, i.e. a high resistance state (HRS) and a low resistance state (LRS).

As promising candidates for this application, manganites present large programming windows [START_REF] Rubi | Two resistive switching regimes in thin film manganite memory devices on silicon[END_REF][56], i.e. high ratio between HRS and LRS, as well as the possibility of multi-level resistance states [START_REF] Stoliar | Nonvolatile Multilevel Resistive Switching Memory Cell: A Transition Metal Oxide-Based Circuit[END_REF], a clear advantage towards increasing storage density. The manganite of our choice is LaMnO 3+δ (LMO), as it is able to accommodate a wide range of cation and oxygen stoichiometry (La 1-y MnO 3-(3y/2)+δ and LaMn 1-z O 3-(3z/2)+δ ), which leads to changes in its electrical properties [START_REF] Xu | Evidence for a Crucial Role Played by Oxygen Vacancies in LaMnO3 Resistive Switching Memories[END_REF]. Here it is important to notice that for a La/Mn ratio of 1 LaMnO 3+δ corresponds to the simplified formula for cation deficient La 1-E Mn 1-E O 3 , where E=δ/(3+δ). The apparent oxygen excess in LMO films is expected to be compensated by a mixed valence state of the manganese cation (Mn 3+ /Mn 4+ ).

Particularly, RS in LMO has been reported to be larger for oxygen vacancy-rich films [START_REF] Jin | Mechanism study on oxygen vacancy induced resistance switching in Au / LaMnO 3 / SrNb 0[END_REF] [START_REF] Töpfer | LaMnO(3+δ ) Revisited[END_REF]. In this way, depending on the oxygen content (δ), the LMO structure changes from orthorhombic to rhombohedral phase at high δ [START_REF] Rodríguez-Carvajal | Neutron-diffraction study of the Jahn-Teller transition in stoichiometric LaMnO3[END_REF]. This transition which occurs around δ = 0.09 at room temperature for La:Mn ratio 1:1 [START_REF] Xu | Evidence for a Crucial Role Played by Oxygen Vacancies in LaMnO3 Resistive Switching Memories[END_REF][START_REF] Töpfer | Charge transport and magnetic properties in perovskites of the system La-Mn-O[END_REF]. A similar transition from orthorhombic to rhombohedral with increasing oxygen content has also been reported for non-stoichiometric lanthanum manganites (La/Mn̸ =1), for which, following Vegard's law a linear variation of the lattice constants was observed with y and ,z for the orthorhombic samples, while a monotonic decrease of the rhombohedral angle α was observed for the rhombohedral ones [START_REF] Töpfer | Transport and Magnetic Properties of the Perovskites La1-yMnO3 and LaMn1-zO3[END_REF].

In order to grow engineered LMO films, we chose the pulsed injection metal organic chemical vapour deposition (PI-MOCVD) technique, as it allows for a controlled growth of the perovskite phase over large areas (at wafer level) with high film uniformity and conformal coverage using liquid precursors at room temperature and under an inert atmosphere [START_REF] Dubourdieu | Pulsed liquid-injection MOCVD of high-K oxides for advanced semiconductor technologies[END_REF][START_REF] Pla | Engineering of Functional Manganites Grown by MOCVD for Miniaturized Devices[END_REF]. Both conventional MOCVD [START_REF] Bertrand | Low-pressure-MOCVD LaMnO3 very thin films on YSZ (100) optimized for studies of the triple phase boundary[END_REF][START_REF] Bosak | Cation and oxygen nonstoichiometry in RâMnâO (R=La, Nd) bulk samples and thin films[END_REF]109] and PI-MOCVD [START_REF] Vincent | Crystal Structure Transformations of a Magnetoresistive La0.8MnO3-δ Thin Film[END_REF][START_REF] Pignard | Chemical vapor deposition and structural study of La0.8 MnO3âδ thin films[END_REF] have been used for the deposition of epitaxial and polycrystalline LMO and doped LMO thin films enabling the control of the oxygen and cation stoichiometry.

Since the precursor solution is simply prepared by dissolution of metalorganic species, the stoichiometry of the film is easily tuneable by changing their concentration. Besides, the PI-MOCVD technique offers the additional benefit of injecting micro droplets by using an electric valve granting excellent control over the quantity of precursor transferred to the reaction chamber and therefore a good control of the thickness of the films. Hence, by modifying a number of controlled parameters on the process, such as pulse frequency, oxygen partial pressure and temperature, the LMO thin films can be tuned during growth.

In memristors, the electrodes play a crucial role in the RS response. For example, on electrochemical metallization memory chips, ions from the electrode (such as Cu or Ag) [START_REF] Tsuruoka | Forming and switching mechanisms of a cation-migration-based oxide resistive memory[END_REF][START_REF] Guo | Understanding the switchingoff mechanism in Ag+ migration based resistively switching model systems[END_REF] migrate to the other electrode generating a filament across the memristive material. Another example is the case of valence-change memories, in which the nature of the contact varies depending on the difference between the work functions of the electrode and active material, creating an ohmic contact or a Schottcky barrier. Furthermore, some electrodes can be oxidized forming a new interface layer that can also act as oxygen reservoir (e.g. Ti, TiO 2 ) [START_REF] Herpers | Spectroscopic Proof of the Correlation between Redox-State and Charge-Carrier Transport at the Interface of Resistively Switching Ti/PCMO Devices[END_REF][START_REF] Cabout | Role of Ti and Pt electrodes on resistance switching variability of HfO2-based Resistive Random Access Memory[END_REF][START_REF] Schroeder | Resistive switching in a Pt/TiO2/Pt thin film stack -a candidate for a non-volatile ReRAM[END_REF].

Integration of LaMnO 3+δ films on platinized silicon substrates for Resistive Switching applications by PI-MOCVD

The use of Pt as bottom electrode in our LMO-based devices guarantees an inert and ohmic contact, as the work function of LMO is 4.5-5.1 eV [START_REF] Jacobs | Understanding and Controlling the Work Function of Perovskite Oxides Using Density Functional Theory[END_REF] and the one of Pt 5.9-6.2 eV [START_REF] Derry | Work function of Pt(111)[END_REF].

In order to integrate crystalline LMO memristive films in silicon-based devices, we used commercial platinized silicon as the bottom electrode/substrate heterostructure. Nevertheless, these substrates undergo dewetting when exposed to high temperatures for long times. The high temperatures required to grow perovskite thin films by techniques such as PI-MOCVD or PLD, are a drawback because the continuity of the Pt bottom electrode can be lost [122]. In this work, we explore a number of different strategies to integrate LMO films on platinized siliconbased devices by PI-MOCVD, overcoming the challenge of the high temperatures required for their deposition. The growth parameters have been optimized to fabricate homogeneous and dense LMO films with different values of δ. Furthermore, we report the experimental proof of structural changes related to the growth strategy, being able to tune the LMO structure from an orthorhombic to a rhombohedral phase, as well as the changes in Mn n+ valence associated to this structure.

Experimental

LMO deposition conditions

LaMnO 3+δ (LMO) thin films were grown by pulsed injection metal-organic chemical vapour deposition (PI-MOCVD) in a JIPELEC reactor [START_REF] Sénateur | Synthesis and characterisation of YBCO thin films grown by injection-MOCVD[END_REF][START_REF] Szkutnik | Influence of precursor nature on the thermal growth of Tin-Indium oxide layers by MOCVD[END_REF]. The precursor solutions were prepared using tris(2,2,6,6-tetramethyl-3,5-heptanedionato)lanthanum(III) [La(thd) 3 ] and tris(2,2,6,6tetramethyl-3,5-heptanedionato)manganese(III) [Mn(thd) 3 ] commercial metal-organic precursors provided by Strem chemicals, and m-xylene (1,3-dimethylbenzene) solvent from Alfa Aesar. All solutions were prepared with a total metallic precursor concentration of 0.0225 M and a La/Mn precursor ratio of 2, the value of which was previously optimized to grow stoichiometric films with a La/Mn ratio close to 1.

The injection of the liquid precursors was performed using a frequency of 2.5 Hz and an opening time of 2 ms with a solution feeding rate of 0.35 mL/min. The evaporator was thermalized at 250 °C and Ar was used as carrier gas. The total pressure in the chamber was fixed at 5 Torr and oxygen gas was added directly in the reaction chamber to obtain an oxygen partial pressure of 50%. Independent heaters and several thermocouples distributed along the reactor circuit allowed for a well-controlled heat gradient from the injector to the reaction chamber, maximizing the flux of carried precursor. The deposition temperature inside the main chamber (a hot-wall quartz reactor heated by an external furnace) ranged from 500 to 750 °C. The substrates used were 1 cm × 1 cm chips cut from a Pt (150 nm)/TiO 2 (

Structural and electrical characterization

Scanning electron microscopy (SEM) was performed in a Quanta250 environmental SEM FEG from FEI, and SEM FEG ZEISS GeminiSEM 500 to study the surface morphology and determine the LMO thickness using the cross section of the films. The cationic film composition was analyzed by semi-quantitative energy-dispersive X-ray analyses (EDX) using an Oxford Inca Energy detector coupled to the SEM. A combined study in X-ray diffraction (XRD) and Raman spectroscopy was performed to determine the crystal structure of the films and to detect the presence/absence of secondary phases. XRD was measured in grazing incidence configuration (GIXRD) in a 5-circle Rigaku Smartlab diffractometer to enhance the diffraction signal from the polycrystalline films and minimize the signal of the platinized silicon substrate. Raman spectra were collected using a Jobin Yvon/Horiba Labram spectrometer equipped with a liquid nitrogen-cooled CCD detector. Experiments were conducted in the micro-Raman mode in a backscattering geometry using a green laser (λ = 514.5 nm). The silicon spectrum at ambient temperature was always measured and used as reference to calibrate the LMO spectra. The nanostructure growth was further analyzed in cross section by transmission electron microscopy (TEM), a JEOL 2011 equipment operating at 200 kV with a 0.19 nm point-to-point resolution.

X-ray absorption near-edge spectroscopy (XANES) spectra at the Mn K-edge of LMO thin films were collected at the ESRF ID12 beamline (Grenoble, France). Measurements were taken under vacuum at 25 °C in fluorescence mode using a nearly constant 200 mA beam current. Two silicon photodiodes detectors were used to collect the total fluorescence, one in back-scattering geometry and a second diode mounted at 90°with respect to the incident beam. The relative variation in the Mn formal valence was carried out from the experimental recorded inflection point after XANES spectra normalization, and using references found in the literature for other manganite perovskites [START_REF] Cuartero | X-ray absorption and emission spectroscopy study of Mn and Co valence and spin states in TbMn1-xCoxO3[END_REF].

The device fabrication for electrical measurements was performed in clean-room facilities combining laser lithography (Heidelberg instruments µPG 101) to define the top electrode pads (200 µm squared pads) and metal evaporation (Plassys MEB 550 electron gun 10 kW) to grow 100 nm thick Au layer as top metal electrode. The electrical characteristics were measured under ambient conditions within a Faraday cage with microprobe manipulators using a B1500 Agilent semiconductor parameter analyzer.

Integration of LaMnO 3+δ films on platinized silicon substrates for Resistive Switching applications by PI-MOCVD

Results and discussion

Dense and homogeneous LaMnO 3+δ (LMO) thin films with variable oxygen content (δ) have been deposited by pulsed injection metal organic chemical vapour deposition (PI-MOCVD).

The structural transition between orthorhombic and rhombohedral phases has been correlated with δ and the manganese oxidation state of the films. Furthermore, the tuned LMO films integrated in a silicon-based substrate showed resistive switching behavior as a proof-of-concept of the suitability of their use as ReRAM.

Optimisation of the deposition conditions

In this section we present the optimisation of the temperature and number of pulses using a fixed pressure of 5 Torr, a gas mixture composed of Ar 50% and O 2 50% during the deposition step, a pulse injection frequency of 2.5 Hz, and an opening time of the valve of 2 ms. These last two parameters allowed for a good evaporation of the precursors and a constant flux during deposition. Samples grown at 675 °C and ca. 1 h of deposition time (i.e., 10000 pulses) evidenced the thermal instability of the platinum substrate surface. The main issues were Pt dewetting, the formation of pinholes at the LMO film, and/or LMO film cracking due to Pt grain evolution. The surface and the cross section of the heterostructures showing cracking of the sample, holes and Pt percolating to the surface are presented in Figure S1(4.9) and Figure S2(4.10), respectively, in Supporting Information Section4.7. In order to avoid these problems and with the aim of decreasing the time of exposure of the bare Pt surface to high temperatures, three strategies including different heating and deposition steps were proposed and tested, as shown in Figure 4.1.

(I) Single-step strategy: The reactor was heated at 25 °C/min in vacuum (ca. 0.15 Torr) up to the deposition temperature. Once the temperature was reached and stabilized (ca. 10 min), the gas mixture was introduced in the chamber and the injection of precursors started and ran for a limited deposition time (controlled number of pulses). Several deposition temperatures in the range of 650-700 °C were tested.

(II) Single-step strategy with annealing: The heating and deposition procedure was the same as described in strategy I, but adjusting the deposition temperature between 660 and 680 °C. Once the injection was completed, an additional thermal treatment at 500 °C was performed. The samples were let to cool down to the post-deposition annealing temperature at a rate of 10 °C/min. At this temperature the environment was modified (i.e., annealing in oxygen) and the temperature was held for 1 h.

(III) Double-step strategy: This growth procedure was divided in two deposition steps at Integration of LaMnO 3+δ films on platinized silicon substrates for Resistive Switching applications by PI-MOCVD different temperatures. The first step consisted of depositing a stabilizing layer at an intermediate temperature (500 °C) and was followed by a second deposition step at high temperature (750 °C) to achieve the complete crystallization of the film. The heating conditions of the first step were the same as described in strategy I. Next, after a new heating ramp up to 750 °C (25 °C/min), the second deposition step took place once the temperature was reached. Cooling down began right after the injection was completed to minimize the time the sample was held at high temperature.

As common points to all strategies, the mixture of gases (O 2 /Ar) was introduced in the reactor chamber when a temperature 50 °C lower than the deposition temperature was reached. Controlled cool-down conditions varied depending on the desired structure (orthorhombic or rhombohedral). The samples were cooled down to an intermediate temperature (450-400 °C) using a cooling ramp of 10 °C/min in a specific gas environment (either pure Ar or O 2 /Ar mixture in strategy I, and O 2 /Ar mixture in strategies II and III). From this intermediate temperature, the cool-down rate was free and the pressure was kept at 5 Torr using only Ar gas. The influence of the atmosphere during cooling was investigated using strategy I through comparison of the effect of using either O 2 /Ar mixture gas or pure Ar. Figure 4.2 summarizes the LMO film growth as a function of the maximum temperature used and the period of time at which the samples were exposed to temperatures above 500 °C. This representation allows one to define the temperature region in which Pt remains thermomechanically stable, as well as the time limits to obtain dense and homogenous LMO films, i.e., without nanoporosities due to insufficient thickness, or the opening of grains in a flower bouquet-manner for thick films (see Figure 4.11 in Supporting Information Section4.7).

For strategy I, as a first approach samples were deposited at 500, 600 and 700 °C. At 600 °C the LMO layers showed low crystallinity whereas above 700 °C, diffusion of Pt across the film occurred, which may lead to short circuits in the final LMO-based device. The samples corresponding to this strategy are depicted as red squares in Figure 4.2. The points inside the shadowed green region correspond to deposition conditions for which dense and homogeneous LMO films are successfully obtained. The maximum appropriate temperature was 700 °C for times (above 500 °C) shorter than 45 min. This time limit could be increased (up to 50 min) for the deposition temperature range between 640 and 680 °C.

Strategy II was designed to obtain higher δ values in the LMO film by adding a postdeposition annealing step. LMO films were grown following strategy II (represented in Figure 4.2 by blue triangles) respecting the same critical limits established from strategy I. It was proved that despite the addition of a thermal annealing at 500 °C for 1 h, the selected conditions led to dense and homogeneous LMO films.

For strategy III, based on two consecutive deposition steps at 500 and 750 °C, the critical Green circles stand for double step strategy (strategy III). Data in the green shadowed area correspond to successfully grown dense and homogeneous LMO films without instability issues. The red shadowed region corresponds to films with Pt instability problems. Purple area corresponds to columnar growth and opening of the grains, leading to porous films.

Integration of LaMnO 3+δ films on platinized silicon substrates for Resistive Switching applications by PI-MOCVD parameters were: (i) the LMO layer thickness obtained by growth at low temperature (d 1 ) to stabilize the Pt layer, and the LMO layer thickness corresponding to the second growth at 750 °C (d 2 ) that must fully cover the LMO film deposited first; (ii) the time at 750 °C, which must be sufficiently long to allow the bottom LMO layer to be fully crystallized but short enough to avoid Pt dewetting; (iii) the maximum thickness d 1 + d 2 (and the ratio between both thicknesses d 1 /d 2 ) in order to obtain non-porous films (see Figure 4.10 in Supporting Information Section4.7), which is limited by the opening of grains in a flower bouquet-manner through columnar growth.

As previously explained, for strategies I and II, the time above 500 °C was limited to 45/50 min and the temperature should not exceed 700/685 °C during this time. This leads necessarily to a limitation in the LMO film thickness, which was of 60-80 nm for the fixed injection frequency of 2.5 Hz. Nevertheless, this limitation cannot be extrapolated to strategy III, for which the growth is stable for longer times, since the protective layer grown at 500 °C prevents Pt dewetting and Pt grain evolution.

The deposition time at 750 °C in strategy III was varied between 35 and 85 min. It is important to notice that strategy III allowed the growth of highly dense LMO films on Pt at 750 °C for longer deposition times, the Pt film was stable for all samples. In addition, the exposure times at 750 °C were enough to fully crystallise the LMO protective layer. From this point of view, strategy III is more robust than strategies I and II. Nevertheless, large thickness samples (green circle at 85 min above 500 °C in Figure 4.2) exhibit columnar grain opening that makes them inappropriate for device fabrication where flat surfaces are required. Besides, in strategy III, we have to take into consideration the effect of thicknesses d 1 (500 °C) and d 2 (750 °C), bearing in mind that d 1 should be thick enough to operate as protective layer and that d 2 should be continuous enough. Therefore, for this last strategy the minimum deposition time required above 500 °C is ca. 30 min, and the maximum time will be limited by the beginning of columnar opening, around 60 min.

In summary, the outcome of this optimisation process allowed for the determination of the optimal temperature/time range of growth for each strategy, i.e., the time limit before Pt begins dewetting, films begin to crack or nanoporosities begin to form. For films thicknesses smaller than 80 nm, strategy I and II are the most suitable, while strategy III is optimal for film thicknesses above 80 nm. In the latter case, the maximal thickness before columnar grain opening depends on d 1 and d 2 and will be subject of further studies.

Structural and composition characterisation of LMO thin films

In order to properly compare the films obtained by each strategy, structural and chemical studies were carried out by electron microscopy coupled to EDX. Within the resolution limit Furthermore, TEM cross-section observation of the LMO films was performed for strategies I and III, which correspond to the extreme cases (Figure 4.4). The cross-section images corroborated that in both heterostructures the Pt layer is continuous and stable as required; neither dewetting nor Pt diffusion occurred at the interface. The good crystallization of both films was also verified, including the bottom part of the film for the case of strategy III (d1 thickness), which was deposited at a lower temperature. The cross section of strategy III showed more irregular columns, probably due to the difference in growth temperature between the two steps, but no flower bouquet effect is observed for this thickness of ca. 100 nm.

The first step of the phase identification for the LMO films grown by strategies I, II and III was performed by GIXRD. The XRD patterns corresponding to representative LMO films deposited by the three strategies are shown in Figure 4.5.

All diffraction peaks can be attributed to the polycrystalline LMO perovskite phase without any clear preferential orientation and no impurities. Therefore, pure LMO can be obtained through all three deposition strategies. Nevertheless, the orthorhombic and rhombohedral patterns are quite similar and do not allow for an easy discrimination between both structures. Because the structure is strongly linked to the oxygen content, complementary structural characterization was carried out by Raman spectroscopy to identify the LMO phase. As reported by M. N. Iliev et al. [START_REF] Iliev | Raman phonons and Raman Jahn-Teller bands in perovskite-like manganites[END_REF], Raman spectroscopy is very sensitive to small changes in the LMO structure, allowing one to discriminate between the two phases. The orthorhombic phase is Integration of LaMnO 3+δ films on platinized silicon substrates for Resistive Switching applications by PI-MOCVD characterized by sharp Raman lines at high wavenumbers: A g mode (in-phase stretching/outof-phase bending) at 493 cm -1 and B 2g mode (in-phase stretching) at 612 cm -1 , while a weaker A g mode (rotational) is observed around 257 cm -1 . On the other hand, the Raman spectrum of the rhombohedral phase is composed of broad bands centred at ca. 497 and 617 cm -1 , related to the Jahn-Teller distortion while weaker bands are observed at lower wavenumbers, such as the characteristic rotational mode, A 1g , at 217 cm -1 .

Figure 4.6 shows the evolution of the LMO Raman spectrum from an orthorhombic phase (bottom red curve) to a rhombohedral phase (top blue curve). At low wavenumbers, we can follow the evolution from the orthorhombic mode. At higher wavenumbers, a shift can be observed from the A g and B 2g orthorhombic modes centred at 495 and 614 cm -1 , respectively, to the broad Jahn-Teller bands characteristic of mixed-valence manganites with Mn 3+ /Mn 4+ charge and orbital disorder [START_REF] Iliev | Role of Jahn-Teller disorder in Raman scattering of mixed-valence manganites[END_REF], i.e., the first one at ca. 513 cm -1 and the second one split into a main band at ca. 630 cm -1 and a second component at ca. 660 cm -1 [START_REF] Popov | A possible demonstration of cooperative dynamical Jahn-Teller effect in Raman scattering spectrum of rhombohedral LaMnO3[END_REF].

In Figure 4.6 we observe that pure orthorhombic and rhombohedral phases can be obtained for the two extreme deposition conditions. The LMO films obtained by strategy I show a pure orthorhombic phase when cooling down in Ar, which favours a lower oxygen incorporation in the LMO films. When the cooling is performed in O 2 , the films contain a mixture of both phases since δ increases. The annealing step incorporated in strategy II, when performed under O 2 atmosphere, yielded a completely rhombohedral phase. Finally, the films grown using strategy III appear as mainly rhombohedral. Therefore, the Raman results give us an indication of the variation of the oxygen content in the LMO films from the structural phases observed.

The apparent oxygen excess in LMO films is expected to be compensated by a mixed valence state of the manganese cation (Mn 3+ /Mn 4+ ). The variation of the Mn oxidation state for the same series of LMO films measured by Raman was confirmed by XANES (Figure 4.7). The local geometry of Mn was extracted from the pre-edge feature. For all the spectra, the absence of a sharp pre-edge shape reveals an octahedral (O h ) symmetry, i.e., a local symmetry of MnO 6 units, number of coordination = 6, which is in agreement with the perovskite structure [START_REF] Farges | Ab initio and experimental pre-edge investigations of the Mn K -edge XANES in oxide-type materials[END_REF]. The formal valence of Mn was estimated from the Mn K-edge position using reference values reported for the LaMnO 3 compound [START_REF] Cuartero | X-ray absorption and emission spectroscopy study of Mn and Co valence and spin states in TbMn1-xCoxO3[END_REF]. The Mn K-edge position was obtained from the inflection point of the absorption edges, calculated from the second derivative of the curves. The energy values obtained for the four LMO films are detailed in Table 4.1. As expected, the edge position increases with the Mn valence, being the lowest energy value the one of the orthorhombic LMO reference (strategy I, cooling in Ar) and the highest the one of the rhombohedral LMO reference (strategy II, annealing in O 2 ). The relative variation of the Mn formal valence was estimated comparing the shift in the Mn K-edge position with the Integration of LaMnO 3+δ films on platinized silicon substrates for Resistive Switching applications by PI-MOCVD one obtained in the literature [START_REF] Cuartero | X-ray absorption and emission spectroscopy study of Mn and Co valence and spin states in TbMn1-xCoxO3[END_REF]. Cuartero et al. measured an energy shift of 0.4 eV for an increase from Mn 3+ to Mn 4+ . Table 4.1 shows the estimated gradient in Mn formal valence among the four thin films, which is of 0.10 for the two extreme samples (rhombohedral and orthorhombic).

All presented results confirm that it is possible to grow dense polycrystalline LMO films by PI-MOCVD adjusting the deposition conditions. Furthermore, we can tune the oxygen content and the resulting Mn oxidation state, leading to a structure transition from orthorhombic to rhombohedral phase. Considering the advantages and limitations of the three deposition strategies, despite the narrow temperature and time ranges required to avoid Pt instability, strategy I allows for the growth of LMO thin films up to 80 nm providing a very good control of the crystal structure. On the other hand, when rhombohedral films are suitable for the required application, the double-step strategy has proven to be more robust, allowing for the growth of thicker samples by adjusting the thicknesses d 1 and d 2 . Once the films were obtained we proceeded to the fabrication of MIM structures by evaporating squared Au electrodes of 200 × 200 µm 2 . Figure 4.8 presents the current-voltage characteristics of an orthorhombic LMO device. After a forming step, reproducible and reversible clockwise bipolar RS response was attained. The "set" and "reset" values were around -0.7 V and +0.6 V, respectively, which are considerably lower than those reported for epitaxial LMO (-1.5 V and +3.0 V) [START_REF] Jin | Mechanism study on oxygen vacancy induced resistance switching in Au / LaMnO 3 / SrNb 0[END_REF]. While the ratio of resistance of the "ON" and "OFF" states was of the same order as the one reported in the literature (over 7.3 in our case for polycrystalline LMO and between 1.8 and 18 depending on anneal conditions for epitaxial LMO [START_REF] Jin | Mechanism study on oxygen vacancy induced resistance switching in Au / LaMnO 3 / SrNb 0[END_REF]). Therefore, LMO films have proven their suitability as functional material for ReRAM applications.

Conclusion

After solving the issues related to the thermomechanical instability of Pt in the substrate, LMO films have been integrated on a platinized silicon substrate by using PI-MOCVD to fabricate perovskite-based ReRAM devices. To do so, three deposition strategies have been developed.

Integration of LaMnO 3+δ films on platinized silicon substrates for Resistive Switching applications by PI-MOCVD

All of them offer the possibility of growing dense and homogeneous LMO films suitable for resistive switching applications. Furthermore, the oxygen content can be tailored by adjusting the deposition parameters, such as temperature, number of pulses and atmosphere used during cooling. The transition from the orthorhombic to the rhombohedral phase has been correlated with the oxygen content and the Mn oxidation state of the LMO films. Orthorhombic films, as well as films with a mixture of phases have been obtained by using strategy I. With strategy II, which includes annealing in oxygen atmosphere, it was possible to obtain a purely rhombohedral structure. The third double-step strategy has proven to be a very robust method to grow mainly rhombohedral samples of comparable and higher thicknesses than the other strategies by adjusting the thicknesses d1 and d2. Therefore, by combining growth parameters in PI-MOCVD and wisely choosing the best deposition strategy we can tune the characteristics of the LMO films for functional devices. Finally, using Au and Pt as top and bottom electrodes, respectively, the functional resistive switching properties of the optimized LMO films has been validated.

Acknowledgements

This work has been partially supported by the ANR funded "Alps Memories" project (ANR-15-CE24-0018) and has been performed with the help of the "Plateforme Technologique Amont" de Grenoble, with the financial support of the CNRS Renatech network. In addition, this work benefited from the facilities and expertise of the OPE)N(RA characterization platform of FMNT (FR 2542, fmnt.fr) supported by CNRS, Grenoble INP and UGA and from the facilities, and the scientific and technical assistance of the CMTC platform. This project has also received funding from the European Union's Horizon 2020 research and innovation program under the Marie Skłodowska-Curie Grant Agreement No. 746648 -PerovSiC (for D.P.) and from the "Microswitch" French ANR project (ANR-14-ACHN-0012) (for Mo.Bu.). This chapter is divided in 2 main sections. In the first section the fabrication procedures for the standard devices, and for the devices designed for advanced characterisation, are described. Section 5.1.1 presents the fabrication flows used for the memristive devices with standard top electrodes, while section 5.1.2 presents those used for the fabrication of advanced characterisation-compatible devices in Top-Bottom configuration. Devices were prepared in two configurations, with different metals and using different deposition strategies, as follows:

Supplementary Information

1. Devices in Top-Top configuration were fabricated on LMO/SrTiO 3 , LMO/LaAlO 3 and LMO/Si 3 N 4 /SiO 2 /Si films. Ti and Pt were selected as top electrodes for standard devices. Au was the electrode of choice for C-AFM experiments.

2. Devices in Top-Bottom 'sandwich like' configuration were built on commercial platinised silicon substrates. M, M/LMO/Pt heterostructures were prepared using four different metals, being M: Au, Pd, TiN and Ti.

The advanced characterisation devices (Section 5.1.2) allowed the use of techniques such as hard X-ray photoelectron spectroscopy (HAXPES) in Synchrotron radiation facilities. HAX-PES had previously proven successful to study buried interfaces, such is the case for TiN/HfOx [START_REF] Bertaud | In-operando and non-destructive analysis of the resistive switching in the Ti/HfO 2 /TiN-based system by hard x-ray photoelectron spectroscopy[END_REF]or Au/FeSTO/NbSTO [START_REF] Lenser | Formation and Movement of Cationic Defects During Forming and Resistive Switching in SrTiO 3 Thin Film Devices[END_REF]. In this work it was used to study the changes in the elemental composition and charge distribution of the LMO/(Au or TiN) interface [START_REF] Meunier | Resistive switching in a LaMnO 3+δ /TiN memory cell investigated by operando hard X-ray photoelectron spectroscopy[END_REF] both in ex situ and operando conditions. The experimental results led to the publication of an article, which is presented in Appendix A. More details on the HAXPES technique are given in Section 5.1.2, as they are necessary to understand the specific requirements for the fabrication of the advanced characterisation-compatible devices. In section 5.2, a general description of the electrical measurements for resistive switching characterisation used during the project for the standard devices is presented. This covers the description of current-voltage (I-V) sweeps and chronoamperometry measurements (I(t)) applying a fixed voltage, as well as the procedure for general pulsed measurements and, in particular, pulsed resistance switching hysteresis loops (RSHL). The standard fixed parameters applied in the measurements are explained.

Fabrication of Top-Top and Top-Bottom devices for resistive switching

In order to study the behaviour of LMO as active material for resistive switching, different metals were chosen as top electrodes for the test devices. Gold (Au) was selected as top electrode as noble metal, providing an ohmic contact, and was evaporated on samples grown by strategies SI and SIII. Au has also been broadly used as electrode for other studies on the resistive switching of perovskites [28,[START_REF] Sun | Coexistence of the bipolar and unipolar resistive switching behaviours in Au/SrTiO3 /Pt cells[END_REF]. TiN and Ti were chosen as oxidisable metals as they can potentially act as oxygen reservoir [START_REF] Jeon | Stabilized resistive switching behaviors of a Pt/TaO x /TiN RRAM under different oxygen contents[END_REF][START_REF] Fujimoto | TiO2 anatase nanolayer on TiN thin film exhibiting high-speed bipolar resistive switching[END_REF][START_REF] Balatti | Multiple Memory States in Resistive Switching Devices Through Controlled Size and Orientation of the Conductive Filament[END_REF][START_REF] Xu | Bipolar switching behavior in TiN/ZnO/Pt resistive nonvolatile memory with fast switching and long retention[END_REF]. Since TiN is currently used in the production at the integration scale in microelectronics (down to the 45 nm technology), mainly as barrier metal in semiconductor devices, TiN/LMO(SI)/Pt has been the main heterostructure used in this study.

Standard electrode fabrication

Device fabrication was performed in clean room facilities (Upstream Technology Platform (PTA), Minatec, Grenoble). In order to fabricate devices, different pathways can be followed. The simplest option is to use a hard mask and to deposit the electrodes by metal evaporation. This method had been previously used in the preliminary research in our team, to fabricate Ti/( 18 O)LMO/Pt top-bottom devices. In order to give flexibility to the devices design, and to 5.1 Fabrication of Top-Top and Top-Bottom devices for resistive switching 101 fabricate custom devices adapting them to our needs, laser lithography (Heidelberg instrument µPG 101, available in PTA) was used as the patterning tool in this thesis (and in general in the ALPS memories project). The use of laser lithography was possible since the minimal lateral feature in our patterns was of 10 µm, larger than the resolution limit imposed by the laser (theoretical resolution limit 0.6 µm using the highest precision laser writehead available in the service). As the maximal lateral feature in our patterns is of 500 µm the writemode of the laser was chosen consequently (in Heidelberg instrument µPG 101: Writemode III, 5 µm minimum structure size). The photoresist choice in lithography depends on the resolution required. Positive photoresists are those which structure becomes more soluble in a resist developer when exposed to UV light, while negative resists harden when exposed to UV light and therefore generate a mask which is inverse of the exposed region. In our case, the AZ1512 positive photoresist for wet etching from MicroChemicals was selected for the photolithographic procedures. The AZ1512 photoresist is suitable for a resolution above 2 µm and has a very high concentration of the photoactive compound, which contributes to a fast development. The standard fabrication process, as shown in The metallization, depending on the target metal, was performed either by evaporation in a MEB 550 Plassys evaporator (Au, Pd, Pt, Ti) using different recipes depending on the metal; or by sputtering, in a physical vapour deposition (PVD) CT100 Alliance concept in the case of TiN. Furthermore in the case of the Ti electrodes, the Ti is 10 nm thick and is coated with a 90 nm-thick film of Pt to prevent the oxidation of Ti coming from the air-exposed surface. The lift-off procedure was optimized for Au as follows: metallised samples were introduced in an acetone beaker placed in a water bath for 10-30 min at 40 °C. The beaker containing the samples immersed in acetone was extracted from the water bath and, in some cases, to help release the remnant metal, soft bubbling by the side of the sample was done with a The electrode masks were designed with Layout Editor software. The top-bottom configuration mask is composed by 4 rows of 14 electrodes of four different sizes, (50x50 µm 2 , 100x100 µm 2 , 150x150 µm 2 and 200x200 µm 2 ) for two different metals, as shown in Figure 5.3 a. The design including electrodes of 4 different sizes had as goal to facilitate the test of size-dependent phenomena. The alignment of a two layer virtual mask is required in the lithography step. In order to perform the virtual mask alignment, coordinates are calibrated in the laser lithography. The positions of the electrodes in the virtual mask (GDS File) are linked to the physical position of the evaporated electrodes using the camera integrated in the instrument. The rows of 14 electrodes were accommodated, centred in the homogeneous area The patterns used for the Top -Top masks (Figure 5.3 b) are designed to allow electrical measurements as a function of the pad size, of the pad separation distance, and also statistical measurements of several devices for equidistant pads. All of these measurements were performed using Pt -Pt, Ti -Pt and Ti -Ti combinations. The vertical rectangles in the bottom left corner of the mask were designed for transfer length measurements (Pt-Pt and Ti-Ti), to extract the resistivity of the films, the contact resistance and transfer length for the M/LMO interfaces.

Fabrication of devices compatible with advanced characterisation

In HAXPES measurements, changes in the oxidation state of La and/or Mn atoms can be followed by the displacement of their binding energy. The goal of the experiment was to study changes in the oxidation state occurring on the top M/LMO interface related to resistive switching. Operando experiments to follow changes in the oxidation state of the species at the interface were carried out, in which the sample was analysed in its pristine state, as well as in the different resistance states. The photoelectrons detectable depth is of tens of nm, therefore this experiment imposes to work with very thin electrodes in order to focus on the described region of interest, i.e. the top-electrode/LMO interface. Standard characterisation devices (with 100 nm thick electrodes) were not suitable for hard X-ray photoelectron spectroscopy (HAXPES) since the incident radiation needs to cross through the top electrode into the LMO film and the photoelectrons need to cross this depth back up to the detector, as shown in Figure 5.4. Electrodes with a thickness of 5 nm and a surface area of 100x100 µm 2 were designed The fabrication flow for a batch of four devices implicates three alignments of the lithographic virtual masks. Each sample contains four devices (see Figure 5.6) fabricated with one type of metal: Au, Pd or TiN.

Previous to the device fabrication, the LMO film was coated with a 50 nm thick sputtered Si 3 N 4 layer (by PVD in the CT100 Alliance concept), process during which the Pt contact area (visible in Figure 5.4) was covered with a hard mask.

The fabrication flow is presented in Figure 5.5 and divided in three cycles. As previously described for the standard devices fabrication, the samples need to be cleaned before the fabrication process. Then, Figure 5.5 a, the photoresist of choice (AZ1512) is spin coated (2000 rpm, 60s) and baked at 100 °C for 90 s. The laser lithography and resist development (Figure 5.5 b) are performed in the same way as described for standard devices. Figure 5.5 c, represents the reactive ion etching (RIE) step. This step is preformed in an inductively coupled plasma (ICP) etcher (Plasmalab100 from Oxford), equipped with a plasma radio frequency generator (13.56 MHz and 0 to 1200 W) that controls the plasma properties. The RIE constitutes the critical step for the fabrication of these devices. The optimisation of the RIE recipe After the RIE process optimisation the optimal conditions were used, except for the last batch of devices. For this last batch the etch time at 300 W had to be doubled due to a decay in the ICP energy. The 'Plasma Ar soft' parameter (see in the table 5.1) was also extended to 20 s, since it seemed to reduce the amount of F -(coming from the use of CH 2 F 2 and SF 6 ) when the etched areas were measured by Auger spectroscopy (by Dr. E. Martinez, CEA-Leti).

To ensure the cleanliness of the surface all the remaining resist had to be removed (step d, Figure 5.5 d). This cleaning is the same procedure that would be followed in a Au lift-off. During the fabrication flow optimisation performed in test samples (not to be used as devices) it was noticed that the resist hardened after RIE and, if reused for a second lithography step, the metal lift-off left flakes of resist and metal, independently of the lift-off conditions. This may be due to resist crystallisation or partial crystallisation when submitted to the reactive ion etching conditions.

In the second cycle (steps e to h) fresh resist is spin coated on the sample, and a 2nd lithography step is performed, this time to define an area larger than the active one, as shown in Figure 5.6 a. This area is of 400 x 600 µm 2 and it is off-centre. If the active area of 100 x 100 µm 2 is taken as reference, the electrode border is separated 150 µm on the free side and 300 µm from the contact electrode, as can be observed in Figure 5.6. This design came from the need of separating the bulk part of the electrode (contact area) from the analysis area, to avoid interference in the HAXPES measurements. This was particularly critical in the case of Au, since the Au signal (Au 4f) overlaps in HAXPES measurements with the La signal (La 3d). In the metallisation step (step g) a highly conformal 5 nm thick film of metal is deposited in this area.

The third cycle, Figure 5.5 i to l, corresponds to a third spin-coating step, the alignment and lithography of the contact electrodes area and the evaporation of the metals, overlapping in a 100 x 400 µm 2 area over the 5 nm electrode fabricated in cycle 2. The resulting devices are shown in the 3D model and the picture in Figure 5.6.

Standard electric characterisation of the devices

The devices with electrodes fabricated following the procedure detailed in section 5.1.1 were used for the standard electric characterisation of the devices. An Agilent B1500A semiconductor device parameter analyser was used for all the measurements, using a two electrode configuration, applying voltage to one of the electrodes while grounding the other one. For both Top-top and Top-Bottom devices the grounded electrode is Pt as it was the common 110 Fabrication of LaMnO 3±δ -based memristive devices and electrical measurement protocols metal for all devices. Four types of electrical measurements were performed (Figure 5.7), as detailed:

• a) Standard voltage (V) bipolar sweeps: This procedure consists on a 0 V → +V max → 0 V → -V max → 0 V voltage sweep, as shown in Figure 5.7 a. It is performed using V steps of either 0.005 V or 0.010 V, depending on the |V max | value (step size 0.005 V if V max < 2 V and 0.010 V if V max > 2 V). The number of steps in each 0 V → |V max | → 0 V loop is limited to 300 steps.

• b) Chronoamperometry (I(t)): In this procedure (Figure 5.7 b) a constant V is applied and the evolution of the current vs time is measured.

• c) Pulsed measurements: Voltage pulses with a certain pulse height (voltage) and pulse length (duration) are applied to the device, as indicated in Figure 5.7 c. Write and erase pulses are those exceeding V SET and V RESET, while read pulses allow to measure the resistance (current) value after a write/erase pulse has been applied. The pulse duration can be varied, but it is limited to a minimum of 150 µs (in the Agilent B1500A Semiconductor Device Parameter Analyser used for the measurements).

• d) Pulsed measurements used for acquiring Resistance Hysteresis Switching Loops (RHSL): This measurement is performed by applying a staircase of V pulses followed by consecutive read pulses, as shown in Figure 5.7 d. Increasing V pulses are applied until the resistance change is given followed by decreasing V pulses back to 0 V. Next, increasing V pulses in negative polarity are applied until the device switches back, and finally decreasing V pulses in negative polarity are applied back to 0V. Since the read pulses are applied after each write/erase pulse, the high resistance and low resistance states can be determined. This measures both the current (resistance) at a certain voltage (equivalent to the I-V standard bipolar sweeps) and the remanent resistance (measured for each read pulse). 114

Resistive switching in TiN/LMO/Pt heterostructures This chapter presents the resistive switching (RS) characteristics measured in TiN/LMO/Pt micrometric devices. First, the LMO thin film sample and the TiN/LMO/Pt devices used for the electrical measurements are described in Section 6.1. Next, in section 6.2 the measurement of the initial resistance state (Ri) is used to show a statistical analysis of the pristine state. These measurements served to establish a classification among devices related to their initial state. The electrical characterisation of the devices is presented next. First, the conditions required for the device activation (soft forming or initialisation) prior to the resistive switching operation are explained in Section 6.3. Then, the devices are compared depending on their activation process. Different activation and operation protocols led to different I-V characteristics, which are presented in Section 6.4. Once the most convenient operation protocol is chosen, the dependence of the RS with the pad size is evaluated. Next, in Section 6.6 the coexistence between threshold switching and bipolar switching is presented. Section 6.7 is dedicated to the study of optimized devices, including the analysis of a selection of the best operating devices for each size, the performed cyclability tests, and a study of the conduction mechanisms and a qualitative model for RS in these heterostructures. Finally, the TEM and EELS analysis of some pristine and operated devices is shown and described. As presented in Chapter 4, the growth of the LMO films was performed using three different strategies. To understand the relationship between the growth strategy and the switching response, a comparison of the TiN/LMO/Pt heterostructures containing LMO films grown by two different strategies (single step SI and double step SIII) is presented in Section 6.8. Finally,to identify the role of the nature of the metal electrode in the switching characteristics of the devices, the TiN/LMO/Pt devices are compared with Au/LMO/Pt heterostructures operated following the same protocol. The results are presented and compared in Section 6.9.

The TiN/LMO/Pt heterostructure

Most of the results presented in this chapter (Sections 6.2 to 6.7 and Section 6.9) have been obtained for heterostructures in which the LMO film was deposited using Strategy I (single step strategy, SI) (described in Chapter 4) followed by a cooling in oxygen atmosphere. Strategy I was the strategy selected as the protocol to grow the films is easier, faster and requires lower temperatures (650-675 ºC), limiting the potential dewetting of the bottom Pt. The LMO films fulfil all the selection criteria for device fabrication in terms of low porosity, no cracking, and good homogeneity. They also present a crystalline structure identified as a single perovskite by GIXRD and Raman spectroscopy, as previously described in Chapter 4. The LMO films selected for the preparation and characterisation of TiN/LMO/Pt devices were composed by a mixture of orthorhombic and rhombohedral phases, which can be easily obtained by PI-MOCVD. The LMO films have the following characteristics: a La/Mn ratio of 1.1 ± 0.1, measured by EDX and calibrated using the RBS composition calibration obtained from LMO/Si 3 N 4 /SiO 2 /Si samples grown in the same conditions; a thickness of 75 nm measured by TEM and a roughness of 3.66 ± 0.07 nm, obtained by AFM (measurements shown in Figure 6.1).

The pristine LMO/Pt film was analysed by Time of flight secondary ion mass spectroscopy (ToF-SIMS, here referred to as SIMS). A SIMS depth profile across the LMO sample and reaching the Pt bottom electrode is presented in Figure 6.2 a. After the first few surface nanometres, the LaO -, MnO -and O -ions are constant throughout the LMO film, while the Pt - SIMS counts are negligible until the bottom interface is reached. A sharp interface between the LMO film and the Pt bottom electrode is measured. No large differences in the LaO -/MnO - (A/B) ratio are observed in the 'bulk' of the pristine film, as shown in Figure 6.2 b. However, a possible small deviation in the concentration ratios is observed in the first nm, being both LaO -/MnO -and MnO -/O -lower close to the top surface compared to the bulk of the film. Surface segregation and B-site rich regions close to the surface are very common features which have been observed for perovskites of many different compositions, both in the form of polycrystalline bulk samples [START_REF] Druce | Surface termination and subsurface restructuring of perovskite-based solid oxide electrode materials[END_REF] and as epitaxial thin films [START_REF] Chen | Segregated Chemistry and Structure on (001) and (100) References Surfaces of (La 1-x Sr x ) 2 CoO 4 Override the Crystal Anisotropy in Oxygen Exchange Kinetics[END_REF]. They have been observed for annealed samples, but also in 'as deposited' films. Thus, it can be concluded that the LMO composition is homogeneous in depth for the LMO thin film with small deviations in the 10 nm closest to the surface, as expected. Figure 6.2 c corresponds to a 2D images of the total SIMS ions measured for the 50 x 50 µm 2 area analysed. No inhomogeneities, no edge effect and no pinholes are observed over the entire surface measured. In addition, 3D images of the O-and Pt-ions measured are shown in Figure 6.2 d and Figure 6.2 e. A clear separation of the O -and Pt -containing regions with a very flat, clean and homogeneous interface proves that there is no dewetting nor intermixing of species. The initial resistance (Ri) of the TiN stacks was measured for 53 devices (out of the 56 available devices). In Figure 6.3 c and d the data shows the variability of the Ri with the size of the electrodes. The initial resistance ranges between several ohms to tens of kΩ. In Figure 6.3 c, the initial resistance is represented as the cumulative probability by pad size, and in Figure 6.3 d the statistical distribution of the same data is presented in a frequency diagram. It can be observed that the Ri variability is smaller for the 100 x 100 µm 2 and 150 x150 µm 2 than for the two other device sizes (50 x 50 µm 2 and 200 x 200 µm 2 ). This effect can be explained by the presence of defects or inhomogeneities in the LMO film, or at the interface between the metal and the LMO, which can affect the contact resistance and electrical field distribution. The density of defects can have more or less influence depending on the size of the electrodes. While the larger pad sizes have more probability of encountering a region with a high defect density, for the smaller size pads the change in resistance resulting from the presence of inhomogeneities is statistically higher.

Analysis of pristine TiN/LMO/Pt devices

As shown by Figure 6.3 d, the LMO-based devices have been classified as a function of their initial resistance, as follows:

• High Ri (HRi), devices presenting initial resistance above 1 kΩ. They constitute 21% over the total number of devices.

• Low Ri (LRi), constitute 60% of the total number of devices, comprising devices which initial resistance is either between 8 Ω and 50 Ω (47% of the total devices) or between 100 Ω and 300 Ω (13% of the total devices).

• Very low Ri, devices which Ri is lower than 8 Ω. The operation of these devices requires very high currents and thus they were discarded for operation purposes. It is possible that these devices are 'short-circuited', meaning that the top and bottom electrodes are in contact.

The data extracted from the operated devices will be analysed separately for HRi devices and LRi devices, as this feature strongly influences the RS behaviour. It is important to notice that HRi and LRi do not correspond to the high resistance state (HRS) and low resistance state (LRS) during device operation, but to the initial resistance state. I-V sweeps were performed to characterise the RS behaviour of the TiN/LMO/Pt stacks, as described in section 5.2. The Pt electrode was always grounded for the measurement of the devices. The sweep rates used were either of 5 mV/step for sweeps from 0 to ±2 V or of 10 mV/step for sweeps above 2 V. In this way the number of steps was limited to 300 steps for a loop from 0 V → | V max | → 0 V. The optimal working conditions were defined according to the following four criteria:

1) The endurance of the devices should be maximized.

2) The HRS/LRS ratio should be as high as possible.

3) The HRS and LRS should be stable, with a high retention. This feature was favoured over high HRS/LRS ratios.

4) The current compliance (Icc) should be as low as possible, imposing in all cases an upper limit of 100 mA.

Activation of the devices: soft forming vs initialisation

In order to operate the devices an activation protocol was required. This protocol was found to be different for the two families of devices defined (HRi and LRi). The resistive switching community is still discussing about the correct terminology [START_REF] Williams | Summary of the Faraday Discussion on New memory paradigms: Memristive phenomena and neuromorphic applications[END_REF], but in general lines two types of RS activation processes are possible, namely the 'forming or electroforming ' and the 'initialisation'. A 'forming' process describes a sudden change in resistance at a certain voltage, considerably larger than the operation voltages (V SET , V RESET ), and typically by the formation of a filament between the electrodes. On the other hand, an 'initialisation' process is described as a gradual change between the initial resistance state (IRS) and a HRS (or LRS), which occurs progressively and at voltages comparable to the operation voltages of the device [START_REF] Williams | Summary of the Faraday Discussion on New memory paradigms: Memristive phenomena and neuromorphic applications[END_REF], i.e. the V SET and V RESET (as defined in Chapter 5). During the initialisation process the resistance values of the HRS and LRS states vary in a continuous manner, separating gradually, until stable states are reached. The different activation conditions for the two families of devices are summarised in table 6.1.

Activation of the HRi devices

The HRi devices activation requires what will be named hereafter as a 'soft forming', corresponding to a sudden change of resistance occurring at a voltage equal to, or slightly larger than the operation voltage. The forming voltage in our case is of 2 V, for devices working with operation voltages ranging in absolute value from 1 to 2 V in both polarities. The 'soft forming' protocol for the HRi devices has been defined as follows:

The devices were operated by bipolar voltage sweeps (0 V → +V max → 0 V → -V max → 0 V) as shown in Figure 6.4. Although, it is possible to limit the current compliance to 100 mA in positive and negative polarity, as shown in the example in Figure 6.4, it is convenient to reduce the Icc as much as possible to protect devices from break-down during the 'soft forming' procedure. In the established protocol (chosen to maximise the devices endurance), the current compliance (Icc) is limited to 40 mA for positive voltages to ensure that the device does not overcome this value, and to 60 mA in the negative polarity to allow enough current for the RS to occur. For most devices a small opening occurs during the first ± 1.5 V cycle, as shown in Figure 6.4 t1(a,b), followed by a 'threshold-like switching' that occurs when sweeping below -1.5 V, as shown in Figure 6.4 t2(a,b). It should be noted that depending on the device, the voltage values required for the forming to take place can be somewhat higher than in this example.

In our TiN/LMO/Pt devices the step shown in Figure 6.4 t2(a,b) leads to a change in resistance from a HRi to a lower resistance state, which implicates the coexistence of the observed 'threshold-like' and bipolar resistive switching. If 'pure' threshold-type switching had taken place, the resistance after the cycle would equal the initial resistance, differing from the case depicted here. The change of resistance after the cycle was verified by remeasuring the resistance (current) at 100 mV. After this, if a + 2 V voltage sweep is applied in the positive polarity, a sudden jump to a LRS occurs (Figure 6.4 t3 a and b). This complete three-step process (t1 to t3) constitutes the so-called 'soft forming'. This switching regime is then stable and reproducible with two well-defined states, switching in positive polarity from HRS to LRS and in negative polarity from LRS to HRS. This sense of switching is defined in the literature as counter clock wise (CCW) switching.

In summary, the HRi devices 'soft forming' presents two requirements. First, a 'thresholdlike' switching needs to take place at negative polarity, for which the current compliance is set to 50-60 mA. Second, a larger positive bias (larger than the |V| applied in negative polarity to find the threshold) must be applied right after with the Icc set to a maximum of 40 mA. The limitation in the Icc in the positive polarity is done to prevent the device breakdown and to keep the current smaller than the one reached in the negative regime. Note that during the 'soft forming' more than one cycle may be required to provoke the resistance drop in the positive polarity and that while the initial resistance of these devices is in the order of kΩ, the LRS was always smaller than 500 Ω and the HRS is most often in the order of the hundreds of Ω for stable devices.

Activation of LRi devices

Two different protocols have been defined for the activation of the LRi devices (see summary Table 6.1). The first one consists of an 'initialisation' and the second one is a 'soft forming' that takes place for similar operation conditions to the ones described above for the activation of HRi devices.

For the 'initialisation' process, presented here in Figure 6.5, bipolar cycling is performed from the beginning, i.e. the apparition of the threshold feature is not necessary. The 'initialisation voltage' is the voltage used to obtain a first separation of two different resistance states (Figure 6.5 t1a,b) and is here in the range of ± 1 V. In Figure 6.5 t1a,b to t2a,b an evolution of the LRS I-V characteristics is observed together with the appearance of the HRS. A small increase of the LRS resistance occurs during the first sweeps (from approx. 15 to 20 Ω), and the LRS acquires a small curvature in the R-V plot, showing a less ohmic behaviour. This increase in R was observed mainly for very low initial resistance devices (Ri 8 to 50 Ω) and could be related to the apparition of a barrier as the ohmic contact is lost. A gradual increase in the HRS/LRS ratio is observed when cycling in a consecutive manner keeping the same voltage window (Figure 6.5 t1-t3 a,b) until a stable regime is reached (Figure 6.5 t4 a and b). In the example the Icc was limited to 40 mA in both polarities. Posterior tests proved that an asymmetric current limit is beneficial to maximise the HRS/LRS, hence the general Icc limit to 60 mA in the negative polarity and to 40 mA in positive polarity was imposed as the standard protocol (see summary Table 6

.1).

A visual model of the mechanism of activation has been introduced in subsection 6.7.4 along with the model for the RS operation. Given the differences in the activation process for the devices with different initial resistance, i.e. a sudden vs a progressive separation of the HRS and LRS for the HRi and LRi devices respectively, we propose the hypothesis that the Resistive switching in TiN/LMO/Pt heterostructures the HRi is caused by the existance of an interlayer at the TiN/LMO interface resulting from oxygen incorporation into the TiN during the TiN deposition (device fabrication). In the case of the LRi devices this layer would be discontinuous, or even non existent (see Figures 6.19a and 6.20a). In both cases, the TiN would scavenge O from the LaMnO 3+δ generating an oxygen-depleted LaMnO 3+δ ' region at the top interface where δ ' < δ (see Figures 6.19a and 6.20a). LaMnO 3+δ ' is expected to be less conducting than LaMnO 3+δ . Therefore, for the HRi devices, when the top electrode is polarized negatively, oxygen would drift from the oxidized TiN interlayer back into the LMO film (see Figures 6.19a) and the total resistance decreases. While for the LRi devices there would be a small drop in resistance under positive bias upon the homogenisation of the oxygen content in LMO driven by the electric field (or the formation of a conducting filament, see 6.20a), while under negative bias the depletion of oxygen in the LMO would reappear (or the oxygen rich filmament will dissolve as shown in Figure 6.20b), and extra oxygen could be incorporated to the 'oxidized TiN interlayer' either from the electrode itself or from the environment.

Once this point has been reached and the 'initialisation' process is completed the two stable resistance states are accessible. The initial (t1) and stable (t3 or t4) HRS and LRS states for both the HRi and LRi devices are plotted in Figure 6.6 in Log I vs Log V scale to evaluate the evolution of the resistance states linearity. The slope for each linear range of the graph is included in the figure and the different behaviour of the two initial states (devices HRi and LRi) can be appreciated. The characteristics of the HRi devices show a strong non-linear behaviour in the pristine state (Figure 6.6 a t1) and a change in slope is observed when the device is activated. In positive polarity, for V > 0.2 V, the differences in slope from t1 to t3 are: m = 4 at t1, while m = 3 at t3 in HRS t3 and m=1.4 in LRS t3 . Similar changes are observed in the negative polarity for V < -0.2 V. While the slopes of the transient states (for t1) are comparable for the HRS t1 and the LRS t1 , they are different for the final HRS and LRS at t3 (see values in Figure 6.6 a). The LRi device (Figure 6.6 b ) starts with pristine states which are almost linear (slopes m= [1-1.1]) and which evolve gradually upon cycling to less linear states (2 slopes in HRS). In Figure 6.6 b t1 each transient state can be fitted to a single slope (LRS t1 m=1, HRS t1 m=1.1) while at t4 the HRS shows a change in slope to very high values (m=10) at V 0.6 V in positive polarity and a small change in slope to m=1.9 at V -0.6 V in negative polarity.

The non-linear pristine states in the HRi devices, are in agreement with the potential existence of an oxide barrier in the interface TiN/LMO. After activation, this barrier may break or filaments may form across it (during the soft forming process) which would explain the HRS being lower than the pristine state.

In the second type of activation process, which we consider as 'soft forming', the devices Resistive switching in TiN/LMO/Pt heterostructures were also operated using bipolar voltage sweeps (0 V → +V max → 0 V → -V max → 0 V) as shown in Figure 6.7. In this case an increase in resistance is given consecutively in the first positive polarity sweeps (Figure 6.7 t0, 0-5), after which a 'threshold-like' switching occurs in the negative polarity (Figure 6.7 t0, 6-7). In this example, the current compliance at t0 is Icc-= Icc+ = 100 mA, value which is not reached under the positive bias. In the standard protocol, in order to protect the devices, the Icc in positive polarity was set to a current low enough to prevent the device break-down and in negative polarity was set to the minimal current necessary to obtain an increase in resistance, typically +Icc = 40 mA and -Icc = 60 mA (as in the soft forming of the HRi devices, see summary Table 6.1). So far, the difference between the 'soft forming' procedures in LRi and HRi lies on the need of the resistance to increase at least to hundreds of ohms in the LRi devices at the beginning of the activation process as shown in Figure 6.7 t0, 0-5. After this step the process is identical to a 'soft forming' for a HRi device (see Figure 6.4). It must be noted that it is possible to cycle in the clockwise (CW) regime before the 'soft forming' process is completed, but this regime is not stable and it is only possible to cycle up to a maximum of 5 to 10 cycles until the hysteresis is lost, as shown in Figure 6.7 t1 and t2. Therefore, no further attention will be paid to cycling before activation.

Once the 'soft forming' is completed the current compliance is fixed either as Icc-= Icc+, Icc-> Icc+, or Icc-< Icc+. As previously indicated for the HRi devices, the first situation (Icc-= Icc+) leads to a risk of device breakdown during operation and to low endurance values (≤10 cycles). The second condition (Icc-> Icc+) results in CCW cycling and large endurance (>10 cycles), while the third condition (Icc-< Icc+) leads to CW cycling, for which a cyclability of 10 to 15 cycles was tested, presenting very narrow cycles (small HRS/LRS ratio) as shown in Figure 6.8. In this figure the final step of the device activation through a 'soft forming' is shown. Step 0 shows the last transient LRS t' before the drop in resistance to the operation HRS. The cycles (steps 1 to 4) are performed using bipolar voltage sweeps (0 V → +2V → 0 V → -2V → 0 V) and with the fixed current compliance relation Icc+ = 60 mA and Icc-= 40 mA leading to CW operation.

It must be noted that this protocol for CW cyclability can also work for HRi devices, although it is not shown here. For LRi devices an initialisation in CW regime is possible by imposing the +Icc > -Icc condition during activation. Devices initialised in such a manner are further studied in section 6.4. The activation protocols are pad size independent, the activation required seems to be related to the initial resistance state. The totality of the HRi devices undergo 'soft forming' while in the LRi devices the possibility of an initialisation seems related to an initial resistance state sufficiently high (devices with Ri [50,300] Ω). Table 6.1 shows a summary of the activation protocols for HRi and LRi devices. The table collects the current compliance conditions imposed during the activation for 'soft forming' and 'initialisation' and after activation to obtain CW or CCW cycling.

In summary, the asymmetry in current limitation was found to be crucial for obtaining reliable and reproducible resistive switching behaviour among different devices. Allowing a higher current to pass in negative than in positive polarity leads to CCW operation, while an inversion of this relation (i.e. higher Icc in positive) leads to an inversion of the I-V curve cycling sense, leading the device to work in CW instead. The CCW devices have better potential for RS since the HRS/LRS ratios are larger (maximum HRS/LRS for CCW devices of ≈40, and maximum HRS/LRS for CW devices of 1.7). In the following section (section 6.4) a brief evaluation of devices operated in CW mode is presented. Moving forward our efforts will be focused on understanding RS in CCW regime, as it will be presented in section 6.5 and forward.

CW and CCW operation of the TiN/LMO/Pt devices

As explained in the previous section, it is possible to operate the TiN/LMO/Pt devices in CW or CCW cycling mode depending on the current compliance relation between the positive and negative voltage (see summary Table 6.1). The possibility of cycling CW with very narrow cycles (small window opening) after a 'soft forming' was presented in Figure 6.8 where five consecutive cycles were plotted. The change of switching sense from CCW to CW can be achieved in the devices that have already been initialised, as well as in the pristine devices, independently of the initial resistance range.

Here, only the cycling results concerning 'initialised' LRi devices are presented. In Figure 6.9 three devices with different sizes are represented. LRi devices turning CW have very small HRS/LRS ratios (here ~1.2) and so are the differences between the windows obtained for devices of different sizes. The resistance values are in the order of 100 Ω.

Interestingly, if after switching CW the Icc relation is inverted to Icc+ < Icc-, it is possible to go back to a larger HRS/LRS CCW switching, as shown in Figure 6.9. There are changes in the I-V characteristics provoked by the reversed switching. The process requires two increases in Icc (first in positive Icc to turn CW and then in negative to turn CCW). The devices will not cycle for more than 5 to 10 cycles if lower current compliances are used. This proves that the devices do not work exactly in the same manner as before, hence an irreversible change has been induced to the device when the Icc was increased.

It must be noticed that each time the regime is changed the operating currents of the devices need to be increased, compromising the stability of the devices as the higher the current the higher the Joule effect and the probability to damage the device.

This behaviour in which a device can switch CW and CCW was also observed by [START_REF] Muenstermann | Coexistence of Filamentary and Homogeneous Resistive Switching in Fe-Doped SrTiO3 Thin-Film Memristive Devices[END_REF] et. al. in Fe doped STO. In their devices the sense of cycling was determined by the maximum voltage reached in negative polarity. In those devices they found coexistence of filamentary and homogeneous resistive switching, as one of the switching behaviours exhibited scaling with the pad size.

After analysing the operation in CW mode, and given the promising HRS/LRS ratios observed after initialisation on CCW operated devices (values average [1.4-20]) compared to the very small ratios found in CW operated devices [1.2-1.7], we concluded that CCW devices show better potential for resistive switching applications, and therefore they will be the focus of further investigation.

Devices size dependence evaluation.

In order to try to explain the phenomena behind RS in the TiN/LMO/Pt heterostructures the data will be treated separately depending on the pristine state, i.e. the HRi and LRi devices will be analysed separately. In this section, the pad size dependence will be evaluated. First, the HRS and LRS values of all measured devices are compared. To do so, the states are read at 100 mV (V READ ), at the 5th cycle after activation. Then, the comparison of the devices operation is performed for a series of selected devices. One device of each size for each of the families (HRi and LRi) is selected. Pad size dependent HRS and LRS would hint towards an interfacial or volumetric phenomenon, while the absence of pad size dependence would be in agreement with a filamentary nature of resistive switching. The studies were performed on devices sharing operation history, which details are indicated for each study.

Statistics for all devices (HRi and LRi)

As has been presented for the pristine sample description, there is a dispersion of initial resistance values (Figure 6.3 d), with most of the devices presenting a low initial resistance state (60% of the devices present a LRi between 8 and 300 Ω), but with 21% devices in a much higher initial resistance value (HRi > 1 kΩ).

To judge if the size feature affects the RS behaviour all the data obtained from the first operation of the devices after the activation protocol are studied. Here, the HRi 'soft forming' CCW devices and the LRi 'initialised' CCW devices are compared (see Table 6.1). As presented in Table 6.2 the largest HRS/LRS ratio is observed in 200 x 200 µm 2 HRi devices (HRS/LRS ratio of 164). In the table the mean and median value for the HRS/LRS ratios have been calculated. While mean values are affected by outliers (single data points too high or too low compared to the other obtained values), the median gives the numeric value separating the superior and inferior halves in a set of data, and is not affected by outliers. The HRS/LRS median values for HRi devices are comparable for all devices ranging from 5.5 to 10. Nevertheless for LRi devices a small decrease in the median HRS/LRS value from 4 to 1.9 is observed. To determine if there is a real size dependence the median values were selected and plotted in Figure 6.10 where a log-log R vs pad area representation permits to evaluate the resistance to area relation. Since the represented data for the HRS and LRS are far from fitting a slope of -1, the dependence of the full pad area with the resistance can be discarded.

In summary, in this statistical study the HRS/LRS ratios as well as the HRS and the LRS of all operated devices were evaluated, for the four available pad sizes and for the two device families (HRi and LRi). The relation of the median values of resistance vs area showed that there is no direct dependence of the full pad size with the HRS and LRS resistances neither for the HRi nor for the LRi devices. The switching scenario is most probably filamentary. Nevertheless small variations can be observed hinting towards a potential mixture of mechanisms contributions, such as a multifilamentary scenario.

In order to better study the potential pad size dependence of the devices, the population was reduced and only devices that could cycle more than 10 times were further studied. A device from each size has been chosen and will be analysed in the following subsection (6.5.0.2)

Study of representative devices of each size

TiN/LMO/Pt devices of the four different pad sizes were selected as representative if they accomplish the two following requirements:

1. Devices must have stable cycling for more than 10 cycles 2. Devices must have comparable history before the switching, i.e. they have been submitted to the same activation/operation protocols.

The devices are again separated according to their initial resistance i.e. the HRi devices and the LRi devices are analysed separately. By choosing among the representative devices for each size, the cumulative probability of the HRS and LRS resistance for N cycles (N stands for the number of consecutive cycles in the same voltage range and with the same voltage step) can be plotted for each device. This type of representation gives the probability of finding a device at a certain resistance state or at a lower value. The two sets are presented in Figure 6.11 a and b for HRi devices and Figure 6.11 c and d for LRi devices showing that for each device, for all cycles, the HRS and LRS are always clearly separated, and both states can always be clearly distinguished.

For the HRi devices, the 150 x 150 µm 2 device (dark blue and light blue triangles in Figure 6.11 a and b) presents the largest dispersion in the HRS and LRS values while a lower dispersion is observed for smaller devices, making them more suitable for memristive applications. For the LRi devices, the highest HRS values measured correspond to the 150 x 150 µm 2 device (dark blue triangles in Figure 6.11 c and d). In this case the dispersion of the resistance values is much smaller than for the device of the same size (150 x 150 µm 2 ) but with an initial HRi. The HRS and LRS states can be clearly differentiated for all devices, having a low dispersion in all cases. Nevertheless, the LRS and HRS values for 50 x 50 µm 2 devices are the lowest among all devices (dark green diamonds in Figure 6.11 c and d). The HRS values of 50 x 50 µm 2 devices are comparable to the LRS values for bigger devices (such as the 100x 100 µm 2 devices, light purple circles in Figure 6.11 c and d).

Representative I-V curves of the selected devices from the two families (HRi and LRi devices) have been plotted separately to compare the behaviour of both defined families. Figure 6.12 shows the four pad sizes, first separately, and then overlapped for comparison. In the I-V curves presented in Figure 6.12 the same CCW operation is observed for all the HRi and the LRi devices. It should be noted that V max and V max' are used as the maximum positive and negative voltages reached, which in some case can be different in absolute value for different polarities, as shown in Figure 6.12 for the 50 x 50 µm 2 devices.

In the case of the initialisation (LRi devices) the I-V cycles are always performed as follows: 0 V → +V max → 0 V → -V max' → 0 V, so at the beginning of each cycle the devices 6.3 -HRS/LRS ratio for representative HRi and LRi devices for the 12 th consecutive cycle start in HRS. During operation, shown in Figure 6.12, both types of devices start in HRS. In step 1 a quarter of a cycle (0 V → +V max ) is performed, during which the switching to an LRS occurs. During step 2 (a quarter of a cycle: +V max → 0 V) the devices are in LRS. Step 3 corresponds to the 0 V → -V max , where the switch from LRS to HRS occurs, and in the last quarter of a cycle (-V max' → 0 V) the devices are in the HRS.

From the two datasets shown in Figure 6.12 (HRi and LRi devices) the obtained HRS/LRS ratios are presented in Table 6.3.

For these two datasets in particular, the HRS/LRS ratios in HRi are in general larger than in LRi.

When the I-V curve shape is evaluated (Figure 6.12), it can be observed that the HRi and LRi families behave very similarly, presenting changes in resistance which are quite gradual. While HRi devices present small jumps in resistance at voltages around 0.5-0.7 V (typical from filamentary behaviour), LRi devices show in general curves with a more gradual change in resistance. For the LRi devices the curves are comparable independently of the pad size. Further analysis of the I-V curve shape of the devices will be presented for optimised devices in subsection 6.7.4 with the aim of determining the conducting mechanisms of the TiN/LMO/Pt heterostructure in the HRS and LRS.

Summarising, the analysis was focused here on representative devices cycled more than ten times and sharing activation and operation history. Through the cumulative probability chart it was clarified that the smaller devices from both families have less variability in the HRS values, while the LRS remains invariant with the exception of 150 x150 µm 2 devices. Nevertheless, the biggest HRS/LRS ratios obtained are those of the largest devices (200 x 200 µm 2 and 150 x 150 µm 2 ). From the evaluation of the I-V curves shape and the HRS/LRS values, obtained for the selected devices on their 12 th consecutive cycle, smaller HRS/LRS ratios are found for LRi devices while the curve shapes are quite comparable with the exception of some small abrupt jumps in current observed for the HRi devices.

Threshold RS coexistence with Bipolar Resistive Switching

The TiN/LMO/Pt devices presenting CCW bipolar resistive switching have so far shown promise to operate in a stable manner with relatively high HRS/LRS windows. It may have caught the eye of the attentive reader that threshold-like features are also present in several occasions. These features appear at high voltages in the I-V curves, such as those appearing in HRi devices before the 'soft forming' (Figure 6.4) or in LRi devices after the initialisation Figure 6.12 a in the 200 x 200 µm 2 and 150 x 150 µm 2 devices. It is important to notice that for the HRi devices and LRi devices the threshold voltage in negative polarity is in the same range: V th [-1.1,-1.5]. While in the LRi devices threshold RS frequently appears when a current of ≈30 mA is reached, in HRi devices it normally requires currents larger than 50 mA after the 'soft forming' event. This 'threshold-like' switching is therefore a common feature to both families of devices. Although this feature has been found more frequently in negative polarity, probably due to the asymmetric limitation in Icc, it can also be found in positive polarity, as presented here in Figure 6.13. It is found more frequently in LRi devices and it has never been observed prior to the 'soft forming' in HRi devices. Given the shape asymmetry and the threshold voltage value difference in positive and negative polarities, we have designated this process as an asymmetric threshold switching. In the cycles shown in the Figure 6.13 (1 → 8) bipolar resistive switching is given counter-clockwise (CCW) and 'threshold-like' switching is given clockwise (CW). We propose this CW threshold switching to happen in a filamentary fashion, although the nature of the filament remains unclear.

In this case, these threshold-like characteristics could be related to heating, in which case the dissipation of heat could be easier in the bottom Pt electrode, as it is a continuous electrode, than in the top electrode pads, which have micrometric area sizes. An example of resistance change which could be provoked by heating would be the local change of structure from orthorhombic to rhombohedral, for the LMO oxide. With the temperature increase related to Joule effect the structure could locally change to rhombohedral reducing its resistivity, and when the heat is dissipated the structure would go back to orthorhombic. The example is described in terms of rhombohedral and orthorhombic for simplicity, as in fact the sample is a mixture of phases or distorted orthorhombic phase, as explained in Chapter 4.In the literature, threshold switching in LMO was reported on LMO single crystal by Nath et. al. [START_REF] Nath | Electric field driven destabilization of the insulating state in nominally pure LaMnO 3[END_REF], as mentioned in the introduction (section 1.3.5). In their study they show symmetric threshold switching and proposed that it is given due to the transition in LMO from insulating to a high resistance bad metallic phase. According to the authors this reversible transition is given between two coexistent phases below T JT . These phases could be as suggested in the work of Zhou et. al. [START_REF] Zhou | Paramagnetic phase in single-crystal LaMnO3[END_REF], where resistivity changes are reported when the structure changes from an orbitally disordered orthorhombic to an orbitally ordered orthorhombic phase. It must be mentioned that Nath et. al. discarded the Joule effect as the driving force behind this threshold switching, attributing the phase change origin to the effect of the applied field.

Optimised devices

To evaluate the stability of the TiN/LMO/Pt memristive devices, additional electrical characterisation was performed. For this purpose devices operated only in CCW regime both from the HRi and LRi families were selected among those representative of each family. First, to evaluate the endurance of the switching windows, we repeated cycles for more than 25 times, and up to 300 times in some cases. Along this section we refer to the devices cyclability, and not the endurance, as the limit of times that a device could be cycled was not reached. Then, preliminary tests of pulsed cycles in the form of resistance switching hysteresis loops (RSHL) were performed and, finally, the retention of the devices was evaluated.

I-V characteristics devices with high cyclability.

Since the data reported in literature, as well as the required number of cycles for a memory to enter the market are in the order of 10 6 cycles [169], the number of cycles reached in this study remains as a mere proof of concept of the cyclability, as the objective of this work was not to have the highest endurance possible but just to confirm that the switching observed is stable and reproducible. It should be pointed out that a much larger endurance is expected as the measurements were always stopped before failure of the devices.

The evolution of I-V curves was extracted form devices switched more than 30 times (up to 300 consecutive cycles). In Figure 6.14 representative I (in log scale) vs V plots of the LRi devices (left column) and HRi devices (right column) are presented, showing a minimum of 25 cycles, a number considered sufficient to describe the evolution and stabilisation of the switching characteristics. It can be observed that the curves increasingly overlap when N increases (from green to red curves). After these first cycles the HRS and LRS states are generally stable and so is the programming window (HRS/LRS). A SET voltage (V SET ) of around 0.7 V is observed for all the devices as well as a RESET voltage (V RESET ) equal to 1 V (or up to 1.5 V in some particular cases). In addition, the previously described threshold-like characteristics are often observed in the negative polarity.

The general features of each family of devices presented in Figure 6.14 are analysed next. Starting with the curve shape, while the change from LRS (steps 2 and 3) to HRS (steps 1 and 4) in both families of devices is quite gradual after the stabilisation period, the main difference is that for the HRi devices the first 20 cycles showed abrupt jumps in current in the positive polarity, with the exception of the 100 x 100 µm 2 devices (Figure 6.14 right). These abrupt jumps, together with the non scaling of the resistance with the electrode area, would point towards a filamentary or multifilamentary type switching. On the other hand, the operation in negative polarity for all HRi devices (Figure 6.14 right) shows an abrupt jump combined with a soft I-V curve evolution from the LRS to the HRS from the first to the last cycle, with the exception of the 50 x 50 µm 2 device, that shows no jump after the first cycles.

In Figure 6.14 asymmetric polarization is used to operate most of the best performing devices. In fact, we have found that the induction of a voltage asymmetry larger in the negative polarity (|-V max' | > |V max |) enhances the HRS/LRS ratio. In this case a maximum increase of the HRS/LRS ratio of 32% has been obtained 25 cycles after changing to asymmetric voltage conditions. In Figure 6.14 the HRi 100 x 100 µm 2 device exemplifies this increase (the voltage values were: -V = -1.2 V; V = 1 V). Another test, in which the opposite asymmetry (|-V max' | < |V max |) was evaluated is represented in Figure 6.14 for the LRi 50 x 50 µm 2 device. In this case the device showed no proof of increase of the HRS/LRS ratio. It must be noted that Icc is reached in some cases, mainly when the threshold regime appears in the negative polarity, but occasionally it is also reached for the positive polarity (where it is limited to 40 mA to prevent device breakdown). This cyclability test confirms our prediction that the CCW operated devices present stable HRS and LRS and relatively large HRS/LRS ratios when cycled multiple times. The cyclability of the TiN/LMO/Pt memristive devices has been represented by plotting the HRS and LRS values (extracted from the I-V curves of Figure 6.14 at V = 100 mV) as a function of the number of cycles (where N is the number of consecutive cycles). In general, the stabilisation of the device behaviour is achieved in less than 30 cycles for all devices, as it can be seen in Figure 6.15 for the LRi representative devices (left column) and HRi representative devices (right column).

For the HRi devices the HRS/LRS maximum window is defined by the high resistance and low resistance values achieved in the forming step (Figure 6.4 t3 (a,b)). Most often, after the stabilisation (cycling 5 to 20 times) the operation HRS/LRS ratio becomes smaller, presenting an evolution and stabilisation to a slightly higher LRS and to a lower HRS, as it can be observed in Figure 6.15. The 200 x 200 µm 2 , 150 x 150 µm 2 and 50 x 50 µm 2 HRi devices show a decrease of the HRS from the first cycle, and the decrease in HRS/LRS upon operation until the states stabilize after 15 to 30 cycles. Contrary, the 100 x 100 µm 2 device shows an increase in the HRS during the first cycles performed up to ± 1.5 V after the 'soft forming', a larger HRS/LRS and also a larger HRS/LRS variability (first increase then decrease on the ratio). The following cycles (after cycle 30) were performed using an asymmetric voltage sweep ( -V = -1.2 V; V = 1 V), and the HRS/LRS ratio decreased, but the HRS stabilises. Although one may argue that the stability could come from the continuous cycling, the fact that in the asymmetric voltage regime the current compliance is not reached is in itself an advantage.

For the LRi devices the maximum HRS/LRS ratio is obtained upon cycling, and it is equal or larger than the value obtained just after the initialisation (example Figure 6.15 200 x 200 µm 2 device). This ratio can be modified as explained before: a) by choosing asymmetric V values or b) by increasing the current compliance (although this is risky as it can lead to device breakdown) or c) by increasing the voltage window symmetrically whilst the Icc is not reached.

Resistance hysteresis switching loop (RHSL) measurements

The TiN/LMO/Pt devices present a RS with large cyclability and HRS/LRS ratios between 2 and 10 when using bipolar voltage sweeps. However, the devices capability to undergo successive cycling is insufficient to prove the devices suitability for memory applications. In order to have a proper proof of concept, pulsed measurements were performed, applying write pulses of 0.1 s in incremental positive voltage from 0.250 mV to 1 V, followed by decreasing voltage pulses from 1 V to 0 V, and then again applying the correspondent negative pulses down to -1 V, and back to 0 V. After each write pulse, a read pulse of 0.1 V and 0.1 s duration was applied. The results obtained in a 50 x 50 µm 2 LRi device after cycling are represented in Figure 6.16. The obtained window is comparable to the HRS/LRS found performing bipolar voltage sweeps on the same device and on other similar devices (HRS/LRS = 2.8). The I-V curve reconstruction from the write pulses is shown in Figure 6.16 b. Two stable states of resistance HRS ≈ 110 Ω and LRS ≈ 40 Ω (magenta dots) are observed in Figure 6.16 c corresponding to the read resistance at each write voltage pulse.

Future work using a fast pulse system capable of applying pulses in shorter intervals (shorter than 0.1 s) has been envisioned to determine the write/erase operation speed. Preliminary results obtained in the available equipment (Agilent B1500A) allowed us to determine that there is a compromise between the applied voltage (pulse height) and duration of the pulse (pulse length). We have been able to estimate that the minimum write time is smaller than 0.1 s for a resistance change at 0.8 V, but to reach the optimal HRS/LRS ratio requires a larger pulse length of up to 2 s duration at the same voltage. The same switching range (HRS/LRS = 2.8) can be obtained at 0.9 V for a pulse length of 0.1 s.

Retention times

A general feature of all TiN/LMO/Pt devices tested in this work is that the stable state is the LRS, independently of the initial resistance. Several devices were left in LRS and found in LRS when measured after a couple of days. Retention time measurements were performed for the HRS in some cases, since when a certain device, which had been left in HRS was measured on later days, the resistance value had decreased either to the LRS or to intermediate resistance values. The evolution of the HRS with time was measured for 14 h for a LRi device after initialisation. The measurement was performed by reading the resistance value of a 50 x 50 µm 2 device at 100 mV every 60 s. This measurement showed the preservation of the resistance state which varied of only a 1.7% over the HRS initial value. This may indicate that HRS states can be stable for a certain time, (lager than 14h but smaller than 'some days'), since as mentioned at the beginning of the section the HRS tends to decay.

Conduction mechanisms and qualitative model

In this section the potential conduction mechanisms governing the LRS, the HRS and the switching phenomena are discussed. Several reviews on conducting mechanisms classify these in two groups: electrode limited and bulk limited [START_REF] Chiu | A review on conduction mechanisms in dielectric films[END_REF][START_REF] Lim | Conduction Mechanism of Valence Change Resistive Switching Memory: A Survey[END_REF]. Among electrode limited mechanisms we find Schottky emission, Fowler-Nordheim tunnelling, and direct tunnelling. Among bulk limited mechanisms, we find Poole-Frenkel emission, ohmic conduction, space charge limited conduction (SCLC), ionic conduction, hopping conduction and trap-assisted tunnelling.

Different representations of the I-V curves can be plotted to fit the different possible conduction mechanisms taking into account the equations describing the current density (listed in table 6.4). The most convenient representation to elucidate the conduction mechanisms implicated is obtained through the representation of both I and V absolute values in logarithmic scale. The curves previously presented in Figure 6.14 (subsection 6.7.1) are the ones here represented in logarithmic scale for I and V.

Given the previous separation in two families, the HRi devices and LRi devices were analysed separately. The slopes for each linear part of the curves are presented in Figure 6.17 for HRi devices and in Figure 6.18 for LRi devices. A common feature to both types of devices and for all sizes, is that LRS presents ohmic conduction (slope m = 1) for low voltages. Depending on the device, a second slope was observed and fitted for voltages lager than 0.2 V in positive polarity, or smaller than -0.2 V in negative polarity, which hints towards a different conduction mechanism. The conduction mechanism of the HRS remains under discussion. The HRS shows always a linear part with a slope of 1 or close to 1, but the voltage range Table 6.4 -Equations expressing the field dependence of several possible conduction mechanisms (from [START_REF] Chiu | A review on conduction mechanisms in dielectric films[END_REF][START_REF] Lim | Conduction Mechanism of Valence Change Resistive Switching Memory: A Survey[END_REF]). Reproduced from K.Maas. 'La 2 NiO 4+d , a Mixed Ionic-Electronic Conductor for Interface-Type Valence Change Memories' 2019. [START_REF] Maas | La2NiO4+d, a Mixed Ionic-Electronic Conductor for Interface-Type Valence Change Memories[END_REF] Conduction Mechanism Current density dependence on electric field and temperature

Comments

Complete expression

Simplified expression (A and B are constants)

Interface limited conduction:

(1) Schottky (Sch) emission

J SchE = A * T ² exp   -q φ B -qE 4πε k B T   J SchE ∝ T ² exp A √ E T -B
Important at high T

(2) Fowler-Nordheim (FN) tunnelling

J FN = q 2 8πφ B E 2 exp -8π √ 2qm * 3hE φ 3/2 B J FN ∝ E² exp -A E Measured at low T (3) Direct tunnelling J DT ≈ exp -8π √ 2qm * φ B 3h κt ox,eq J DT ∝ exp -A.κt ox,eq
Measured at low T, t ox,eq < 3 nm

Bulk limited conduction (4) Poole-Frenkel (PF) emission

J PF = qµN C E exp -q(φ T -qE 4πε k B T J PF ∝ E exp A √ E T -B
For high T and high E

(5) Ohmic conduction

J Ohm = σ E = qµN C E exp -(E C -E F ) k B T J Ohm ∝ E exp -A T
Always present at very low E (6) Space-charge limited conduction (SCLC) a. Ohmic region

I ∝ V J ∝ E Ohmic-like region at low E b. Trap-unfilled region I ∝ 9 8 ε i µθ V 2 d 3 J SCLC ∝ E 2
Child's square law region, in the case of a single discrete trap level c. Trap-filled region

I ∝ V n , n > 2 J SCLC ∝ E n , n > 2
Steep increase of current at high field (7) Ionic conduction

J ionic ∝ E T exp -∆ G * k B T J ionic ∝ E T exp -A T (8) Hopping conduction
For low E a. Nearest neighbour hopping (NNH)

J NNH = σ 0 E exp( -T 0 T ) J NNH ∝ E exp -A T b. Variable-range hopping (VRH) J V RH = σ 0 E exp( -T 0 T ) 1/4 J V RH ∝ E exp -A T 1/4
(9) Trap-assisted tunnelling (TAT) up to which this linearity remains varies. This part corresponds either to ohmic conduction at low voltage when m = 1, or could correspond to a Schottky emission or Poole-Frenkel mechanism when the slope is in the range of 1.2-1.9. The common tendency found in these curves shows that at voltages larger than 0.4 V the slope becomes m 2, and then increases even more for voltages above 0.7 V. This behaviour could correspond to a space charge limited conduction (SCLC) mechanism, which would be at the origin of the RS in these TiN/LMO/Pt devices, as occurs for other perovskite-type oxides such as (Ba,Sr)(Zr,Ti)O 3 [START_REF] Xia | Field-induced resistive switching based on space-charge-limited current[END_REF] or STO with Pt electrodes [START_REF] Choi | Reversible resistive switching of SrTiOx thin films for nonvolatile memory applications[END_REF]. The SCLC would be possible due to the existence of electron traps which can absorb injected electrons, traps which could come from oxygen vacancies, La vacancies or interstitial ions. [START_REF] Choi | Reversible resistive switching of SrTiOx thin films for nonvolatile memory applications[END_REF] The HRS could correspond to the situation where the traps in the oxide are filled with electrons, while the LRS would correspond to the situation where all these electrons are released. The region in which these traps are filled or emptied would correspond to the top part of the LMO layer in direct contact with the film. However, to univocally confirm this hypothesis temperature dependent measurements should be carried out, which were not possible during the course of this thesis due to time and equipment limitations.

J TAT = A exp -8π √ 2qm * 3hE φ 3/2 T J TAT ∝ exp -A E Measured at low T J is the current density, A = 4πm * k 2 B h 3 is the Richardson constant, q is the electronic charge, m * is the effective electron mass in the oxide, k B is Boltzmann's constant, h is Planck's constant, T is the absolute temperature, E = V d is
A very simple phenomenological model could be suggested in which a region of the LMO film is oxygen depleted, with a higher concentration of oxygen vacancies. This oxygendepleted region could be at the LMO/TiN interface, since TiN attracts oxygen and tends to oxidise and could have a higher resistivity as the Mn-O-Mn-O chains, through which electronic conductivity is given in manganites, would be disrupted [START_REF] Nian | Evidence for an Oxygen Diffusion Model for the Electric Pulse Induced Resistance Change Effect in Transition-Metal Oxides[END_REF].

The model for the HRi devices is shown in Figure 6.19 and the model for LRi devices is shown in Figure 6.20. The models include the activation process which is different for the two types of devices. In HRi devices Figure 6.19a, a TiNyOx barrier could be the origin of the high initial resistance values. When a negative voltage is applied, oxygen drifts into the LMO 3+δ ′ region creating conducting paths and a small decrease in resistance occurs. In Figure 6.19 b the 'soft forming' is completed by applying a positive bias (larger than the precedent negative bias), which will lead to the formation of filaments in the TiNyOx destroying the initial barrier. From this point on the barrier is considered broken (although it may have a contribution that could be the reason for larger HRS/LRS in HRi devices). By the end of this step, the lowest LRS possible has ben reached. In Figure 6.19 c the operation RS mechanism is presented. Under a negative bias the devices switch from LRS to HRS. The LMO 3+δ ′ region has a large density of traps (which could be the oxygen and/or cation vacancies) leading to a SCLC mechanism in the HRS.

In LRi devices there is no barrier in the pristine state, as shown in Figure 6.20 a. However the the TiN electrode contains a small amount of oxygen, as previously seen by EELS. The oxygen-depleted region still exists at the interface, and either it broadens or becomes further depleted upon cycling, until the initialisation is completed (Figure 6.20 b). In the operation regime, Figure 6.20 c, when a positive voltage is applied, oxygen will drift towards the top electrode and enrich this region with O ions (either in a filamentary or an interfacial fashion). This oxygen drift could generate fast conduction paths across the oxygen-depleted region, and therefore the resistance would switch from HRS to LRS as a consequence of the formation of conductive LMO filaments. The oxygen enrichement near the top electrode probably also provokes a TiN oxidation, but without forming a continous and resistive layer (as the preexistent layer in the HRi state). If a voltage is then applied with a negative polarity, these fast paths would be destroyed, and the resistance will again change from LRS to HRS. This hypothesis would be in agreement with bulk limited conduction mechanisms where the 'bulk' region in the HRS would be the oxygen-depleted LMO, and the traps would correspond to oxygen vacancies. In the LRS, if the oxygen vacancies are filled with oxygen ions, no traps (or a smaller amount of traps) would exist, leading to a higher electronic conduction and lower resistance. Although we are aware that the real electrochemical processes occurring at the interface might be much more complicated, this phenomenological model allows for a simple explanation, which is in good agreement with the experimental observations. In Appendix A, TiN/LMO/Pt structures analysed by HAXPES are presented which highlight the role of the TiN as an oxygen scavenger. The measurements were performed during a soft forming procedure, and they showed the incorporation of oxygen into the TiN layer under positive bias through the changes in the binding energy of Ti 1s (and supported by the changes in the O 1s). This incorporation of oxygen into the TiN corresponds to the first steps of the 'soft forming', where the general resistance rises as the oxygen excess in the LMO decreases. Then, the 'soft forming' could be completed by the formation of a Vö filament (or several). From this point on the model works as described above.

In addition, the model presented here, is in agreement with the C-AFM tests performed by B.Meunier and presented in Appendix B, where the C-AFM experiments showed the change of Mn valence from +3.6 to +3.1, and the oxygen ions drift to the surface with the change on resistance. The main difference could be the cycling sense, as the curves observed in this previous work are CW. However the cycling sense cannot be compared in a straightforward manner, since the experiments were performed in a very different configuration, and with very different applied fields, with both electrodes on top of the film: Au (grounded) and a diamond coated Si tip biased. In this work, Meunier et.al. [START_REF] Meunier | Microscopic Mechanisms of Local Interfacial Resistive Switching in LaMnO 3+δ[END_REF] identified the conducting mechanism as SCLC. In order to validate the presented model, TEM was performed in the sample of study, intending to observe oxygen rich/depleted regions in the LMO or a barrier in the TiN/LMO interface. In this section TEM images of the focus ion beam (FIB) lamella prepared and analysed by D. Cooper (CEA-Leti) in a FEI TITAN Themis microscope at an electron acceleration of 200 kV are presented. These lamella were prepared from a pristine device and from two operated LRi devices, one in HRS and one in LRS. Figure 6.21 a shows a large cross section of the pristine device lamella. The observed stack consisting of the TiN top-electrode, the LMO film and the Pt bottom electrode is reasonably flat and homogeneous in thickness. The white rectangle, in the center of the image, indicates the region of interest (ROI), where electron energy loss spectroscopy image (EELS) analysis was performed. Figure 6.21 b and d show the oxygen (red) and titanium (green) composition maps detected by EELS, respectively, superimposed to the ROI TEM image. At the centre, Figure 6.21 c corresponds to the combined composition map for O and Ti, thanks to the superposition it can be seen that Ti presence is limited to the TiN electode, as well as the presence of oxygen in the electrode, and a considerably homogeneous distribution of the oxygen in the LMO film. In Figure 6.21 e a high angle annular dark field (HAADF) image (containing the region where EELS was performed) is presented, while Figure 6.21 f shows the ROI corresponding to the O and Ti composition maps, (on a cut extracted from e). The EELS analysis of the samples reveals a certain amount of oxygen present within the TiN layer, particularly at a distance of about 30 nm on top of the TiN/LMO interface. As this oxygen-rich region is present in the pristine state, it can be concluded that it was formed during the deposition of the TiN film in the cleanroom. Finally, in Figure 6.21 g the same area observed through annular dark field ADF is presented. As can be observed in the images the different TEMmapping methods offer different contrast, i.e. for HAADF the large Z elements are the brightest while for ADF the lightest elements are the brightest, in HAADF images Pt is the lightest film followed by LMO and TiN while this is reversed in ADF.

In Figure 6.22 the STEM (scanning transmission electron microscopy) images obtained in ADF mode for two different lamella are compared. Figure 6.22 a corresponds to a section of the lamella extracted from a pristine (LRi) device, while Figure 6.22 b corresponds to a section of a cycled device left in HRS, and is presented for comparison. As had already been shown in Figure 6.21, an oxygen-rich region is present in the TiN electrode for both devices. As this feature already exists in the pristine state sample, it can be concluded that it was formed during the deposition of the TiN film in the cleanroom, and not during the device operation. It can be observed that the device has undergone operation without traces of Ti electromigration nor any other evidence of physical damage. composition maps for same two lamella. By comparing the HRS and pristine state images in Figure 6.22 no evidence of the formation or presence of a TiNyOx or TiOx layer can be perceived. Furthermore, no evidence of variations of the oxygen concentration have been found across the LMO film, which lead to think that these variations are small, as expected. The fact that oxygen rich/depleted regions have not been observed does not invalidate the hypothesis of their existence, but it proves that the changes are to small to be observed by simple EELS mapping. It should be noted that the present observations have the limitation of corresponding to very small areas compared to the size of the device and to two different devices. For the future, in order to try to observe very small changes on a localized region, in-situ TEM experiments are proposed, in which the same cross-section region would be monitored before and after the application of voltage to the device. Nevertheless operando experiments are extremely challenging and a dedicated sample holder which can be electrically contacted is required.

6.8 Comparison of the memristive characteristics of TiN/ LMO/Pt devices for LMO films grown by different strategies

The previous sections 6.2 to 6.7.4 described the switching characteristics of the TiN/LMO/Pt devices in which the LMO film had been grown by the single step strategy (SI). Although less extensively, TiN/LMO/Pt devices with LMO grown by the double-step strategy (SIII) were characterised as well. These LMO films were grown in two steps: a first step at 500 °C and a second step at 750 °C. In this case the characterisation was mostly focused on the comparison between both strategies, aiming to assess if the memristive devices can perform in the same manner. In the SIII case, two simultaneously grown samples with TiN/LMO/Pt devices were tested, but the number of devices analysed is much smaller. From the first sample (sample 1) the pristine states of 13/56 electrodes were measured being 10/14 100 x 100 µm 2 devices and the two remaining a 150 x 150 µm 2 and a 200 x 200 µm 2 device. From the second sample (sample 2) 9/56 electrodes pristine states were measured ( 7 were 100 x 100 µm 2 devices, and the two remaining were 50 x 50 µm 2 and 150 x 150 µm 2 devices). It was found that the initial resistances of devices from sample 1 were higher (6 -28 kΩ) than those from sample 2 (62 -600 Ω and one device 27 kΩ), and the devices could only be successfully activated through 'soft forming'.

The devices could be cycled previous to the activation (as described in subsection 6.3.1.2). In the case of the LMO (SIII) based devices, this regime was more stable than for strategy I based devices (cyclability > 20 cycles). In the example a HRS/LRS ratio of 2.3 was found for a 50 x 50 µm 2 device, which is comparable to the values found for the HRi devices in the SI sample in the CCW regime. The retention time of the HRS of a 200 x 200 µm 2 device was tested for 14 h previous to the 'soft forming' (data not shown). The measurement conditions were 100 mV applied every 60 s with a Icc 1 mA, that was never reached. A slow decay in time was mesured: the resistance started at 24 kΩ decayed of 16% in the first 3h and of of 19% over the initial HRS over 6h.The device showed a much stronger decay in the last hour of measurement where it dropped by 1.4 kΩ. This may indicate that the states obtained previous to the 'soft forming' have no retention, even if they showed better HRS/LRS than the ones obtained from devices in the SI sample. However, more retention tests would be required to have a good statistic and extract reliable conclusions.

A test imposing a 'soft forming' activation for CCW operation was carried out on a SIII device (200x200 µm 2 ). The same operation conditions as for SI devices were imposed (+Icc = 40 mA and -Icc = 60 mA). In this case we find the threshold curve in the negative polarity that was characteristic of the devices starting in HRi (not shown here). Following the 'soft forming' protocol, previously presented in Section 6.3.1 in Figure 6.4, the devices 6.9 Memristive characteristics of Au/LMO/Pt devices 157 operate effectively in CCW regime. Another device of 50x50 µm 2 was tested (following the same 'soft forming') and cycled 50 times from 0 V → +V max → 0 V → -V max → 0 V being, V max = 1 V and -V max =-1.2 V. The LRS and HRS were 25 Ω and 195 Ω respectively. This was the largest HRS/LRS ratio obtained (7.8) in HRi 50 x50 µm 2 devices among both strategies (Figure 6.23 b1).

It is important to notice that the I-V sweeps for the SI and SIII devices of the same size operating CCW after activation by 'soft forming' are similar, as shown in Figure 6.23. Some particular features are identified in the I-V characteristics which are similar to those found for the SI based devices. For the SIII devices, in positive polarity a resistance change composed by a sharp jump plus a soft slope occurs. The jump is particularly evident in the 200x200 µm 2 curves (Figure 6.23 a and b). For both strategies in negative polarity differences are found for larger and smaller devices. The 50 x 50 µm 2 devices show a gradual increase in resistance that is reached after the V RESET , behaviour previously observed in the SI series of devices (Section 6.7.1, Figure 6.14 ). The 200 x 200 µm 2 show rather abrupt changes of resistance in the negative polarity.

Two resistance switching hysteresis loops were performed for a 50 x 50µm 2 SIII device working in CCW regime, as previously shown for a SI LMO device in Section 6.7.2. In order to compare both heterostructures in the same operation regime the same RHSL conditions were used, i.e. write pulses of 0.1 s in incremental positive voltage from 0.250 mV to 1 V, followed by decreasing voltage pulses from 1 V to 0 V and then again increasing negative values to -1 V and back to 0 V. Between the write pulses, read pulses of 0.1 V and 0.1 s duration were applied. Figure 6.24 shows the possibility to access two different resistance states by changing the reset voltage. The first cycle (steps 1 to 4) was performed starting in a HRS (HRS 1 = 204 Ω) after a -1.1 V pulse (Figure 6.24 a step 0). Two cycles were performed in the +1/ -1 V regime (Steps 1 to 8), the change in voltage range led to a second HRS (HRS 2 = 96 Ω). HRS retention obtained in these conditions was verified for 100 s in between the cycles.

In conclusion, when the devices are activated following the same protocol and operated in the same conditions the same RS behaviour is found for both heterostructures independently of the growth strategy (SI or SIII) used for the deposition of the LMO thin film.

Memristive characteristics of Au/LMO/Pt devices

TiN has been the top electrode of choice so far to study the RS on the LMO/Pt stack. This electrode is expected to play an 'active' role in the RS phenomenon since it has been reported to scavenge oxygen from manganites. [START_REF] Baek | In situ TEM observation on the interface-type resistive switching by electrochemical redox reactions at a TiN/PCMO interface[END_REF] The TiNyOx can act as oxygen reservoir, [START_REF] Zhang | Bipolar Resistive Switching Characteristics of HfO2/TiO2/HfO2 Trilayer-Structure RRAM Devices on Pt and TiN-Coated Substrates Fabricated by Atomic Layer Deposition[END_REF], exchanging oxygen back and forth with the sandwiched manganite. On the other hand, the contact nature of noble metals such as Pt or Au with manganites is ohmic [START_REF] Borgatti | Chemical insight into electroforming of resistive switching manganite heterostructures[END_REF][START_REF] Tsubouchi | High-Throughput Characterization of Metal Electrode Performance for Electric-Field-Induced Resistance Switching in Metal/Pr0.7Ca0.3MnO3/Metal Structures[END_REF] and Au is not expected to favour the oxygen scavenging from the LMO film as it is a noble metal and will not oxidise easily. To understand the role of the electrodes and of the metal/LMO interfaces, it is important to test devices from the same exact sample in the same exact operation conditions, but with Au top electrodes (instead of TiN) and compare the memristive behaviour of the Au/LMO/Pt devices with that of the TiN/LMO/Pt devices.

Au and TiN electrodes were micro fabricated on the same SI LMO layer sample as described in Chapter 5, and that has been used to describe the RS curves in this chapter. The initial resistances of the Au devices tested are all in the low (8 -160 Ω) or very low range (6-8 Ω). This fact can be interpreted taking into account our mechanism model (Figures 6. 19 and 6.20), as in this case no oxygen-depleted LMO interface would form at the Au/LMO interface, as the O from LMO would not be attracted by the oxygen affinity of the Au. The contact is ohmic as proved by the linear relation I-V and the obtained slope of 1 in the log I vs log V curves (not-shown here).

The first operation tests on Au/LMO/Pt devices rise considerably in resistance. Before the 'soft forming' protocol is applied, the I-V curves show asymmetric threshold switching and a certain aperture HRS/LRS (≈ 1.5 measured at 100 mV). as shown in Figure 6.25 a when the devices were tested using Icc + = Icc -. The asymmetry may be related to the difference in top and bottom electrodes.

Secondly, (Figure 6.25 b) to be able to compare these curves with the best studied TiN/LMO/Pt devices, the operation conditions that induce CCW mode were imposed, i.e. Icc + < Icc -. In Figure 6.25 b the operation in CW regime with 2 differentiated states, HRS (1820 Ω) and LRS (1450 Ω) that are stable in the absence of applied electric field (read at 100 mV). However, the HRS/LRS ratio is small (e.g. 1.3) and the device operates CW, since the 'soft forming' had not been performed at this point. It must be noted that both the threshold regime observed in Figure 6.25 a and the I-V curve shape at high V seen in Figure 6.25 b, present a certain similarity with the previously mentioned work of Nath et.al. [START_REF] Nath | Electric field driven destabilization of the insulating state in nominally pure LaMnO 3[END_REF], supporting that this behaviour is derived from the LMO film in the stack. The separation of the states in Figure 6.25 b can be induced by the current asymmetry if we assume that the driving force moving oxygen vacancies in the LMO matrix is larger in the negative polarity (larger applied field). This may lead to the forced accumulation of Vö in the top interface changing the resistance of the LMO in the interface.

As shown in Figure 6.25 c, by completing the 'soft forming' protocol (Section 6.3.1) operating the device CCW is possible. However, the resistances measured in this operation regime were always lower than 70 Ω and the HRS/LRS ratios of 2, smaller than the average for TiN/LMO(SI)/Pt heterostructure. This regime could correspond to the formation of a filament trough the whole LMO film, explaining the drop to very low resistance states.

In conclusion, we can impose the same operation regime (sense of cycling) on Au/LMO/Pt heterostructures at the cost of considerably reducing the HRS/LRS ratios. Cycling has not been tested for more than 10 times. Finally, the percentage of devices that were effectively activated vs tested is much smaller (only 2/14), meaning that for most of the Au/LMO/Pt devices resistive switching was not achieved. This confirms the very important role of the TiN electrode in the switching of LMO based devices, acting as an oxygen reservoir and allowing it to drift between the TiN and LMO layers. The combination of this electrode with a perovskite oxide is the key of the RS described in this chapter.

Conclusions

• The initial resistance state differences extracted from the analysis of a large array of devices allowed to classify the devices in two families, namely HRi and LRi devices, that require different activation protocols. Once the activation has been carried out, if the same operation conditions are imposed, the behaviour of the two families is comparable.

• The TiN/LMO/Pt memristive devices can perform in CCW and in CW mode by controlling the Icc+ and Icc-. However, the CCW mode shows a larger programming window and better stability.

• The evaluation of pad size dependence led to the conclusion that the variations in HRS and LRS do not scale with the whole pad area. The mechanism is not purely interfacial, and could be either filamentary or mixed filamentary/interfacial.

• Threshold switching has been found as a common feature in the TiN/LMO/Pt devices before and after activation (initialisation or soft forming); and it has been found in both voltage polarities and seems to be related to a heating effect.

• Optimised devices show cycling up to at least 300 cycles, relatively large HRS/LRS ratios, stable resistance states and retention times of more than 14 h. In these devices, up to 30 cycles may be necessary to stabilize the HRS and LRS values.

• RSHL measurements were carried out for optimised devices showing that the minimum write time is smaller than 0.1 s for 0.8 V for a minimum HRS/LRS, but the optimal HRS/LRS change requires larger pulses (up to 2 s) at the same voltage. The same switching range can be obtained using 0.9 V pulses of 0.1 s.

• In a first approximation the RS in the TiN/LMO/Pt has been explained by a simple phenomenological model consisting first in the formation of an oxygen-depleted region of LMO (LMO 3+δ ′ ) with a larger resistivity than the LMO matrix ( LMO 3+δ ), leading to the existence of electron traps (related to oxygen vacancies, cation vacancies or insterstitials). In the LRS oxygen ions would fill this region, possibly creating filaments across this region and connecting and disconnect the 'bulk' LMO with the TIN electrode.

• Comparing heterostructures containing films grown by two different strategies (chapter 4), if the devices are activated following the same protocol and operated in the same conditions the same resistive switching behaviour is found independently of the growth technique.

• Au/LMO/Pt heterostructures show linear I-V characteristics in the pristine state, as expected from the ohmic contact Au/LMO.

• The Au/LMO/Pt heterostructures can also be cycles CW or CCW, however they show small HRS/LRS ratios and low cyclability.

• The comparison between the Au/LMO/Pt and the TiN/LMO/Pt memristive devices confirmed the very important role of the TiN electrode in the switching, acting as an oxygen reservoir and allowing it to drift between the TiN and LMO layers.

Concluding summary and prospects

Valence change memories are interesting candidates for new generation memories that could perform the function of a RAM memory with the added feature of non-volatility. In this thesis we have focused on LaMnO 3+δ as a candidate material based on its flexible oxygen stoichiometry accomodated by the valence change in Mn (Mn 3+ /Mn +4 ).

The proper window of deposition conditions for the growth of LMO by PI-MOCVD was found. The growth of LMO using a precursor solution with a La/Mn ratio of 2 leads to the obtention of films in the perovskite phase with a La/Mn composition ratio close to 1 and δ > 0. For temperatures below 700 °C (preferably 650 -675 °C) and by the injection of 2500 to 6000 pulses of precursor solution it was possible to obtain homogeneous dense LMO films with a surface roughnesses below 2 nm.

Two lines of research were carried out:

1. On the one hand, the optimized LMO deposition conditions were used to grow highly textured and epitaxial films on single crystal substrates at 700 °C. These films were dedicated to oxygen ion transport studies.

2. On the other hand, the growth conditions were readapted to integrate LMO films on Pt for memristive devices, which were fabricated using the grown LMO thin films. These devices were used to test resistive switching in metal/LMO/Pt heterostructures and to perform advanced characterisation studies.

In the first line of research presented, the morpho-structural characteristics of the films were studied. By varying the film thickness an evolution of the roughness from 0.5 nm (~10 nm films) to 1.6 nm (~60 nm films) was observed. The TEM cross section of the films revealed the formation of colums on the LMO film emerging from a 10-12 nm thick continuous film. The combined observations from top view (SEM and AFM) and cross section indicate that the columns width increases with the sample thickness from ~10 nm (thinnest samples) up to ~20 nm (thickest samples). The structural analysis revealed the LMO films to be a mixture of orthorhombic and rhombohedral phases. In addition, by cooling down in different atmospheres an evolution of the film structure could be observed, with an increasing rhombohedral contribution when the films are cooled in an oxygen-rich environment. Thus, it is possible to easily control the thickness, roughness and oxygen content of the films by tuning the deposition parameters.

The oxygen transport in highly textured thin dense and nano-architectured LMO films was studied in detail. The high density of defects present in the LMO films grown on LaAlO 3 compared to the high quality epitaxial films grown on SrTiO 3 substrates, contribute to the enhanced functional properties of the films. It has been shown that the oxygen diffusion at 500 °C is extraordinary high due to a high concentration of oxygen vacancies in the material. Nano-architectured LMO with a high catalytic surface area were prepared maximizing the surface/volume features by growing 'finger-like' LMO films presenting nanocolumnar features on top of a dense continuous LMO film. The results of this study have direct implications for the use of LMO as memristive material, but also for other applications, such as cathode for Solid Oxide Fuel Cells.

In the second and main line of research, the LMO films grown by PI-MOCVD were succesfully integrated on platinised silicon substrates,where the Pt serves as bottom electrode for the memristive devices, through three different growth strategies. SI: growth in a single step at a temperature in the 650-675 °C range; SII: growth in a single step in the same temperature range followed by an anneal during the cool down process to tune the oxygen content; and SIII: growth in two steps, the first at 500 °C and the second at 750 °C. These strategies allowed to find different windows of deposition time and temperature to grow LMO on platinised silicon while avoiding dewetting of the Pt. In addition, the different conditions enable the tunnability of the oxygen content in the LMO films grown on platinised substrates.

The fabrication of top-bottom devices was performed combining laser lithography and metal evaporation. Standard devices were fabricated with Pt as bottom electrode and TiN and Au as top electrodes. Advanced characterisation compatible devices were designed to perform HAXPES measurements, and their fabrication permited to succesfully study the changes in Mn oxidation state in the top interface upon device operation. Top-Top devices were prepared with Ti and Pt electrodes on Si 3 N 4 /SiO 2 /Si, STO and LAO substrates. These will be used for future work on LMO memristive devices.

Extensive electrical characterisation was performed on the TiN/LMO/Pt heterostructures in standard devices. The devices were classified according to their initial resistance state and analysed separately. The two families, namely HRi and LRi devices, required different activation protocols before operation. Nevertheless, after activation the resistive switching behaviour of the two families is very similar. The operation of the TiN/LMO/Pt devices was performed both in CCW and CW regime. The CCW showed better stability and larger programming windows (up to a maximum HRS/LRS ratio in the order of 100). The comparison of devices with four different pad sizes permitted to see a small statistical variation of the HRS and LRS with the pad area. However, this variation does not correspond to the dependence expected from interfacial behaviour. This hints either towards a filamentary (possibly multifilamentary) scenario or a mixed interfacial/filamentary mechanism (mainly filamentary) scenario. The operation of the devices with high currents (40-60 mA) led to the observation of threshold-like switching. These threshold-like features seem to be filamentary. Although the exact nature of the filament remains unclear, a possible phase change between orthorhombic and rhomobohedral LMO phases, resulting from a heating effect, has been proposed.

The optimised devices operated in CCW mode have shown cyclability of up to at least 300 times, with low cycle to cycle variability and relatively large HRS/LRS (~13 for the best performing device). The retention times of the HRS were tested up to 14 h, while the LRS has proven to be always stable. They could be operated in pulsed mode in the same range of voltages and reproducing windows of HRS/LRS in the same range than those seen in the I-V sweep mode.

The conduction mechanism of the devices was studied through log-log I-V curves. The analysis hints to a possible SCLC mechanism, although further studies will be needed to identify it unequivocally. A simple phenomenological model was proposed in which first the formation of an oxygen depleted region of LMO (LMO 3+ δ ' ) with a larger resistance than the LMO matrix ( LMO 3+ δ ) would exist at the LMO/TiN interface. By applying voltage there would be a drift of oxygen ions to this region, possibly by the formation of filaments, which would connect and disconnect the 'LMO bulk' and the top TiN electrode. The precence of oxygen vacancies (or cation vacancies or interstitials) in the interface region would explain the presence of electron traps governing the SCLC mechanism in the HRS.

If Au is used instead of TiN as top electrode most of the devices cannot be operated, and those which can show a small memristive window and low cyclability. The comparison between the Au/LMO/Pt and the TiN/LMO/Pt memristive devices confirmed the very important role of the TiN electrode in the switching, acting as an oxygen reservoir and allowing it to drift between the TiN and LMO layers.

In summmary, the combination of the two lines of research led to better understanding on the oxygen transport properties of LMO. The integration of LMO in top-bottom devices on silicon based substrates was fulfilled taking a step forward towards compatibilizing the structure with CMOS fabrication, and the TiN/LMO/Pt devices showed promissing performances for memory applications.

Prospects

Despite the promising results of this work, further experiments would be required to better understand resistive switching in LMO based valence change memories.

Starting from the essential, the evaluation of the proposed model and its modification according to empirical experience could be of great interest.

The validation of the proposed model could be done by combining in-situ TEM with electrical measurements and EELS on a pristine lamella, to assess the drift of the oxygen atoms/vacancies in the layers. This would allow not only to validate the hypothesis of the existance of an oxygen rich/oxygen poor region in the LMO film, but also to reveal further information regarding the nature of the switching, to validate if it is indeed a mixed filamentary/interfacial phenomenon, or a simple filamentary scenario.

Further operando experiments, such as HAXPES or XANES, combined with electrical measurements, could help quantify the valence change of the Mn during the switching. Adjusting the penetration depth by angle resolved HAXPES to control the variation on the analysed depth of the devices, such as the ones fabricated for operando measurments, this depth resolution could be used to localize the valence changes in depth.

Electrical measurements on a larger number of devices, increasing the statistics, could reveal tendencies or particularities of the families of devices not seen in the small amount of data precedently aquired. Moreover, the performance of electrical measurements in smaller size devices, with several more rows of different sizes, could shed light on the dependence of the resistive switching in TiN/LMO/Pt with the pad size. Also, electrical measurements in temperature aiming to clarify the conduction mechanisms, allowing to discard either the temperature dependent mechanism or the temperature independent mechanisms, would be of interest.

Regading the analysis of the potential of these devices, and considering the requirements for mass production, it would be necessary to assess the challenges for the integration of LMO in memristive devices, such as the fabrication temperature, which should be compatible with current microelectronics (200 °C -400 °C in BEOL), or determining how the miniaturisation of the devices to competitive sizes would affect the performance of these devices in terms of read and write voltage, current limts etc. non-destructive operando characterization of functional devices. [16-18] Indeed, compared to conventional XPS with laboratory AlKα (1486.6 eV) or MgKα (1253.6 eV) sources, the use of hard X-ray synchrotron radiation (2.4 -12 keV) gives information from a region up to 5 times deeper inside the sample and even deeper using inelastic background analysis. [19,20] HAXPES has been successfully implemented during resistive switching operation to study buried interfaces of different oxide-based heterostructures such as Ti/HfOx, [21-26] TiN/Hf 0.5 Zr 0.5 O 2 , [27] Pt/Ti/Pr 0.5 Ca 0.5 MnO 3 , [22] Au/FeSTO/NbSTO, [28] and TiN/Pt/NiO. [29] Real operando measurements, consisting in biasing the device while performing HAXPES, are therefore enlightening in order to understand the changes in elemental composition and charge distribution inside the device. Such a method was previously used for studying interface polarization in piezoelectric oxide-based high-density capacitors. [30] In the present contribution, we study the case of LMO memristive devices grown on platinized silicon and covered by TiN, a technologically-relevant top electrode. Operando HAXPES measurements were performed at different resistance states, carefully induced by in situ electrical characterization. The results reveal the chemical modifications and reactions occurring at the LMO/TiN interface.

Ex-situ electrical measurements were performed beforehand on a simple device composed of a platinized silicon substrate (Pt/TiO 2 /SiO 2 /Si), a LMO thin film (70 nm thick) grown by pulsed injection metal-organic chemical vapor deposition and a 200×200 m² TiN top electrode (inset of Fig. 1). A detailed description of the growth conditions and thin film characterization can be found elsewhere (strategy I, O 2 cooling). [31] Figure S1-A.5(a) shows the initialization step (soft electroforming) followed by the resistance change obtained when cycling the device in the -2/+2 V range. The first time a positive voltage is applied, the resistance increases i.e "initializing" the device setting it in a high resistance state (HRS). Then, the resistance of the sample increases by applying a negative voltage and comes back to the same resistance value when a positive voltage is applied, confirming the bipolar memristive behavior of the device, and the ability of the LMO structure to reversibly change its resistance state.

In order to match the HAXPES requirements, an optimized device was designed. In this case, the TiN electrode should be thin enough to allow the photoelectrons passing through and the remote electrical contacts should be employed to avoid any obstruction of the active area. The resulting design is presented in Figure A.2 (a). The LMO thin film (100 nm-thick, grown following strategy III of Ref. [31] ) is partially covered by a 50 nm thick Si 3 N 4 insulating layer deposited by sputtering (CT100 Alliance Concept). The structural characterization of the LMO film (XRD and Raman spectroscopy) exhibiting the coexistence of both rhombohedral and orthorhombic phases can be found in Figure S1-A.5 of the Supporting Information. 100x100 µm² squared craters were etched through the Si 3 N 4 layer and subsequently covered The HAXPES experiments were performed at the Galaxies beamline, SOLEIL synchrotron facility (St-Aubin, France). The pressure in the analysis chamber remained around 10 -8 -10 -7 mbar during the whole experiment. The photon energy was set to 7994.51 eV (as determined from the Fermi level of the Au top electrode at zero bias) with a take-off angle of 50°. The incident photon energy was selected with a Si(333) double crystal monochromator, and the photons were focused to a 30×80 µm 2 beam spot using a toroidal mirror. The photon and analyzer bandwidths were both set to 150 meV providing an overall energy resolution of 210 meV. The Au 4f, Ti 1s and O 1s core-levels were recorded after each change in resistance, while still biasing the sample. The La 3d core-level could not be recorded due to the poor signal-to-noise ratio and the large energy range required, while the manganese core-levels

196

Resistive switching in a LaMnO 3+δ /TiN memory cell investigated by operando hard X-ray photoelectron spectroscopy

were not analyzed because of intrinsically low cross-sections and severe overlapping with the Au photoemission lines. A Shirley-type background was subtracted from all the spectra. The inelastic mean free path estimated for the different core levels through the LMO/TiN stack provides an average sampling depth of 25 nm. [32] The initial spectra (at 0 V and initial resistance state -IRS) of Ti 1s and O 1s are shown in Figure A.2(c) and (d). Peak fitting was performed using a pseudo-Voigt function with 70% and 20 % of Lorentzian peak for Ti 1s and O 1s respectively. The Ti 1s photoemission spectrum consists of three contributions. The component at the lowest binding energy (4965.6 eV) corresponds to bulk TiN while the two other contributions are attributed to TiO x N y (4966.6 eV) and TiO x (4968.2 eV) resulting from the oxidation of TiN. [33,34] The larger width of the oxide peaks (FWHM = 2.6 eV) compared to the TiN peak (FWHM = 1.3 eV) is attributed to inhomogeneities in the oxides composition. The composition of the top as-deposited electrode has been calculated to be of 32% TiN and 68% of oxides (26% TiO x N y and 42% TiO x ). The O 1s spectrum in Figure A.2(d) is dominated by the SiO 2 contribution coming from the Si 3 N 4 oxidized surface due to the large footprint of the incident beam on the surface. Two other contributions are identified at 530.6 eV and 530.3 eV assigned to Ti-O bonds (from TiO x and TiO x N y ) and to LMO, respectively. A more complete HAXPES characterization of the sample (presenting La 3d 3/2 , Mn 3s, O 1s and N 1s in different regions) can be found in Figure S2-A.6 of Supporting Information.

We now address the results of the operando measurements of the LMO/TiN memory cell. The Au 4f, O 1s and Ti 1s core-levels were recorded while biasing the device with different voltages following the sequence: 0 V 2.5 V 5V -2.5 V -5 V. The core level spectra for each step are presented in the Supporting Information (Au 4f in Figure S3-A.7, Ti 1s in Figure S4-A.8 and O 1s in Figure S5-A.9). During this sequence, the sample resistance changes from the initial resistance state (R = 6.2) to the HRS (R = 52) at +5 V and back to a LRS (R = 12) at -5 V. deeper the layer, the smaller the corresponding HAXPES shift (i.e. V Au > V TiN > V LMO ). The voltage applied to the device is progressively lost in the series resistances corresponding to each material or interface of the stack. Considering the Au 4f peak from the top electrode, the total resistance (R tot ) of the device and the external resistance can be calculated with the following potential divider relationship: being R ext = 2.9Ω the external resistance and V appl the applied voltage. Therefore, for V appl = +5 V, the Au 4f energy shift yields V Au = 4.73 V and the total resistance drop can be calculated to be R tot = 53.7 Ω , meaning that the device is in HRS. Similarly, for V appl = -5 V, V Au = -1.0 V and R tot = 3.6 Ω , corresponding to the LRS case. These resistance values are comparable with those experimentally measured (R HRS = 52 and R LRS = 12 Ω ), the differences in LRS being due to leakage current inside the material. This highlights the good reliability of the experimental set-up and the relevance of the spectroscopic measurements in a specific resistance state of the device. When switching the device to HRS, i.e. when applying a positive bias (2.5 and 5 V), a potential drop is observed between the TiN and LMO layers and between the TiN and Au electrode. During the RESET event (IRS to HRS) the resistances increase, first the TiN/LMO resistance at 2.5 V and then the TiN/Au resistance. During the SET event (HRS to LRS) at -2.5 V the potential drop between TiN and LMO strongly decreases (i.e. V TiN -V LMO 0 V) indicating a decrease of the TiN(TiO x N y )/LMO interfacial contact resistance. .3(c) shows the evolution of all Ti components after normalization of their area by the total Ti 1s area (which is assumed to be constant). During the RESET event (from 0 V to 5 V), the Ti 1s area from TiN decreases while the oxide contributions increase indicating the oxidation of the positively biased TiN electrode. Similar behavior of the Ti-O is observed in the O 1s spectra in Figure S5-A.9 as well as a decrease of the O(LMO) contribution. Comparing each of the Ti oxide, the TiO x N y (partially oxidized TiN) ratio rises from 26% to 36% while the TiOx (fully oxidized TiN) ratio only rises from 42% to 44%. This is indicative of the partial oxidation of the TiN electrode, remaining moderate during the RESET event. This oxygen drift from the LMO to the top electrode, in the form of O 2-anions, is driven by the electric field. Conversely, during the SET event (applying -2.5 and -5 V), each contribution goes back to a value close to its initial value. A reversed oxygen drift, from TiO x N y back into the LMO is observed driven by the electric field. The consequence of this redox reaction at the interfacial TiO x N y layer is the increase/decrease of the potential drop previously observed at this interface. The exchange of oxygen ions between LMO and the TiN electrode is described in Figure A. 4.

R tot = R ext • (1 = V Au V appl -V Au )
In this work we presented operando HAXPES measurements for investigating the resistive switching mechanisms taking place in a TiN/LMO/Pt memory cell under application of electrical bias. The use of HAXPES has successfully demonstrated its ability to observe buried interfaces and fingerprint relevant modifications occurring during resistive switching. In particular, spectra recorded during the process reveal changes in the composition of both the TiN electrode and the LMO itself. Fitting of the Ti 1s spectra (and confirmed by the O 1s spectra) clearly indicates a partial oxidation of the TiN electrode and reduction of LMO surface during RESET followed by the reverse reactions during SET. The mechanism can be described by a redox reaction triggered by the oxygen drift at the TiN/LMO interface due to the external bias. The TiO x N y interlayer plays the role of an oxygen getter during RESET, expelling oxygen back in LMO during the SET process. Thanks to careful operando spectroscopic analysis, we proposed a redox model to fully describe the physiochemical mechanism occurring in a potential technologically-relevant memory cell consisting in a LaMnO 3+δ active layer and a TiN top electrode. S1 (a) Grazing incidence X-ray diffraction (GIXRD) 2θ scans performed using a Rigaku-Smartlab diffractometer equipped with a rotating anode to characterize the crystal structure of LMO films grown following the double strategy presented in Ref. [31]. Experimental pattern for LMO thin film grown on Pt/TiO 2 /SiO 2 /Si. * diffraction line from Pt substrate. Rhombohedral (R-3c for La 0.95 Mn 0.95 O 3 from ICDD: 04-007-6350). (b) Raman spectra collected using a Jobin Yvon/Horiba Labram spectrometer equipped with a liquid nitrogen-cooled CCD detector. Experiments were conducted in the micro-Raman mode in a backscattering geometry using a green laser (λ = 514.5 nm). The Raman spectra were calibrated using a silicon reference spectrum at room temperature (theoretical position = 520.7 cm -1 ).In green, positions of orthorhombic Raman lines and in red, of rhombohedral Raman lines. More details in Rodriguez-Lamas et al. [31] Even if the theoretical shape of the beam consists of a 30 mm wide and 150 mm long ellipse, the real exposed area is way larger because of the beam tail and X-ray diffusion through the surface material. (b) N 1s spectrum recorded over the region of interest. The main contribution of the signal comes from the insulating Si 3 N 4 layer. The two peaks at lower binding energy are attributed to Ti-N-Ti bond in TiN and in N-TiO 2 matrix (where N as been substituted by O). Even if no consensus has yet been achieved for the N 1s peak around 400 eV, it is commonly admitted that it may come from oxynitride species such as Ti-O-N and Ti-N-O. (c) Comparison of the La 3d 3/2 , O 1s and Mn 3s recorded on the bare surface and over the complete structure. The two O 1s contributions of the bare LMO surface corresponds have been reported to come from the LMO bulk and the very surface. The O 1s spectrum recorded over the whole device has been described in Fig. 2. However, the present measurements have been carried out with a Take-off angle of 50°i nstead of 60°in the main text. Interestingly, more important Ti-O contribution is observed in the more surface sensitive case (at 50°) confirming the fitting model for the Ti-O and O (LMO) contributions. The energy splitting of the Mn 3s core level has been used in our previous work to study the LMO surface and Mn valence state. The gap between the Mn 3s peaks is inversely proportional to the Mn valence. In the present case on the bare surface the Mn 3s splitting indicates a Mn valence close to 3+ on the bare surface and a mixed Mn 2+/3+ under the Si 3 N 4 insulating layer and TiN top electrode. This observation is in agreement with the presence of oxidized TiN and Si 3 N 4 we previously observed at the interface. During the device fabrication, oxido-reduction reactions occur between the LMO surface and the deposited material resulting to a slightly understoichiometric LaMnO 3-δ surface.

Sample characterization

Details on operando measurements

B.1 Abstract

Manganite perovskites exhibit promising resistive switching properties, for which the understanding of the related microscopic physicochemical changes taking place is still rather scarce. In this work the resistance of a LaMnO 3+δ thin film has been locally tuned within a range of two orders of magnitude using conductive-Atomic Force Microscopy. With the use of X-ray photoemission electron microscopy it has been possible to simultaneously unravel composition and work function changes related to changes in the LaMnO 3+δ resistance state. The resistance change is found to be triggered by oxygen ions diffusing to the surface, where they remain adsorbed. Concomitant to this oxygen displacement, the Mn oxidation state is reduced from +3.6 to +3.1, while the work function decreases by 0.28 eV. We discuss the effect of these physicochemical modifications on the conduction mechanism, which is in agreement with a space-charge limited conduction (SCL) mechanism where the current is restrained by the density of traps at the interface. We show that the resistive switching in the material can be described as a change of the transport regime from a trap-free SCL conduction to a trapcontrolled SCL conduction, depending on the oxygen content in the material.

B.2 Introduction

The rising demand for data storage and miniaturized electronic devices pushes the present Flash technology to its limits. 1,2 New oxide-based resistive switching memories have emerged as alternative non-volatile resistive random access memories (RRAMs) to increase the density data storage, as well as enhance the devices endurance and retention. 3 Among the plethora of oxide structures exhibiting resistive switching (RS) response, perovskites, such as Among the number of different mechanisms that can lead to RS in perovskites, the valence change mechanism (VCM), which controls the behavior of ionic-electronic resistive switches, is one of the most common and promising ones. 12 This VCM can occur through the creation of conducting oxygen vacancy filaments between the two metal electrodes (filamentary VCM) or the accumulation of oxygen vacancies at the electrode-oxide interface (interfacetype VCM). 13 Several chemical and physical processes have been reported to be involved in the VCM-based memristive characteristics in perovskites, e.g. redistribution of point defects, oxygen exchange with the environment, charge-carrier transport at interfaces, and electronic trapping or tunneling effects. 14 However, few works are devoted to correlate and disentangle them, due to the difficulty in obtaining accurate compositional, structural and electronic information with nanometric resolution in the regions of interest. 15-20 Direct experiments correlating the local surface chemistry, nanostructure and electronic transport modifications taking place during interfacial RS in perovskites are yet missing. This lack of understanding of the coupled physicochemical processes is particularly acute for the case of polycrystalline perovskites integrated in silicon substrates in the thin film form.

Mixed-valence manganites have attracted much attention last decades due to their unique electric (e.g. Mott-Hubbard interaction, double exchange interaction, metal-to-insulator transitions), optical (e.g. spectral weight transfer in optical conductivity), magnetic (e.g. magnetic phase transitions, magnetic ordering temperature), and transport (e.g. colossal magnetoresistance effect, Jahn-Teller distortion effect, ion conduction) properties derived from their large capability of defect chemistry engineering. 21,22 In particular, properties are strongly correlated to the Mn valence state and the oxygen stoichiometry, e.g. metal-insulator transition in La 0.8 Sr 0.2 MnO 3 induced by oxygen vacancies. 20,23,24 Among the mixed-valence manganites materials, thin films of LaMnO 3±δ have been extensively studied as cathode materials for solid oxide fuel cells due to their high oxygen reduction catalytic activity. 25 As many manganites, LaMnO 3±δ thin films can be obtained over a large composition range, particularly by varying the oxygen content. The oxygen deficiency (i.e. LaMnO 3-δ ) occurs via the creation of oxygen vacancies in the material, whereas the apparent oxygen excess is ascribed to the formation of La or/and Mn cation vacancy sites (i.e. La 1-x x ,Mn 1-y y O 3 , viz. LaMnO 3+δ ), where represents cation vacancies and charge neutrality remains fulfilled thanks to the presence of a mixture of Mn 3+ /Mn 4+ oxidation states. 23,26-28 Due to steric effects, it is commonly admitted that LaMnO 3+δ cannot present interstitial oxygen in its structure. This work delves into the physicochemical changes occurring at the LaMnO 3+δ (LMO) surface during reversible interfacial RS. We combine conductive-atomic force microscopy Microscopic mechanisms of local interfacial resistive switching in LaMnO3+δ (c-AFM), X-ray photoemission electron microscopy (XPEEM) and Raman spectroscopy to elucidate the compositional, structural and electrical modifications involved in the switching response. 25 nm-thick dense LMO thin films grown on a SiN x /SiO 2 /Si heterostructure were switched from a low resistance state (LRS) to a high resistance state (HRS) using c-AFM. Being able to selectively switch micron-scale areas at different resistance states, the local work function, the lattice system, and the chemical state of all LMO constituent elements at the surface level have been solved and directly related to the resistive states. The results reported here are of crucial importance enabling a better understanding of interfacial-type RS in polycrystalline manganite films and advancing toward the development of new nanoengineered perovskite-based memristive devices.

B.3 Experimental Section B.3.1 Material Deposition

LMO films 25 nm-thick were grown by pulsed injection metal-organic chemical vapor deposition (PI-MOCVD) on top of single crystal (100)-oriented silicon chips, covered by stress free bilayers of 100 nm-thick thermal SiO 2 and 300 nm-thick SiN x grown by low-pressure chemical vapor deposition. An homogenous solution of 2,2,6,6-tetra-methylheptane-3,5-dionate metal precursors, i.e. La(thd) 3 and Mn(thd) 3 , was prepared using a [La(thd) 3 ]/[Mn(thd) 3 ] ratio equal to 2. The optimum growth conditions for these dense polycrystalline films were at 973 K with background O 2 pressure fixed to 5 Torr, using a pulse frequency of 2.5 Hz and 4000 pulses.

B.3.2 Instrumentation

Grazing incidence X-ray diffraction (GIXRD) 2θ scans were performed using a Rigaku-Smartlab diffractometer equipped with a rotating anode to characterize the crystal structure of LMO films.

AFM and c-AFM were performed in a Dimension Icon from Bruker located in a nitrogenfilled glove box (H 2 O and O 2 content below 3 ppm). Highly p-doped diamond coated Si tips (CDT-NCHR from Nanosensors) with a radius in the range of 20-50 nm, a nominal spring constant of 80 N/m and an electrical resistance of 6 kΩ were used for the RS measurements. The applied force on the surface of the film was set to the 0.3-0.9 µN range to ensure efficient electrical contact and low mechanical pressure (<1 GPa). Prior to the electrical measurements, the sample was heated at 390 K for 1h to remove the moisture and ensure a clean LMO surface. Electrical measurements were performed inside the glove box using a Keithley 2635B. The

B.3 Experimental Section

211

I-V characteristics were studied by sweeping from zero to positive, then to negative voltages and back to zero (i.e. 0 V 20 V 0 V -20 V 0 V) at around ± 600 mV/s and grounding a semi-infinite Au electrode located at approximately 100 µm from the tip. It is worthy to mention that first several I-V measurements were performed at different distances from the semi-infinite top electrode, and then an optimized distance was selected (i.e. ~100 µm), ensuring a reproducible switching response and avoiding leakage currents to the Au electrode (which occurred for very short distances). The sweep rate was fixed at ± 600mV/s. The scanning c-AFM mode was used to change the resistance state in a micrometric area at the virgin LMO surface. The writing, erasing and reading voltages used were + 4 V, -6 V and ± 0.5 V, respectively, and a pulse duration of 0.5 s.

The study of the local structure changes was also performed by Raman spectroscopy. Particularly, the variation of the Jahn-Teller effect with the Mn 4+ /Mn 3+ ratio was investigated for both resistance states. Raman spectra were collected at room temperature in a backscattering geometry using a Jobin Yvon/Horiba Labram spectrometer equipped with a liquid nitrogen cooled CCD detector. The 488-nm Ar + laser line with an incident power of 1.2 mW was used for excitation. The light was focused to a spot size of ∼1 µm in diameter by a x100 objective. The Raman spectra were calibrated using a silicon reference spectrum at room temperature (theoretical position = 520.7 cm -1 ).

The local surface analyses were performed by XPEEM with a high transmission, aberrationcompensated energy filter (double hemispherical analyzer) and a focused monochromatic Al Kα source (1486.6 eV). This microscopic photoemission technique provides information about both the work function and the chemical state of specific elements from a given area of interest at the surface. For the local work function analyses, spectromicroscopy measurements were performed where full field, energy-filtered images using secondary electrons are recorded. Photoelectron images were acquired within an energy range E-EF = 1.5 -10 eV (0.025 eV energy steps and 4 s acquisition time per image). For the chemical state analysis, the instrument was switched to the microspectroscopic mode in which photoemission spectra are collected by closing a field aperture either on the HRS or the LRS region. Prior to the analysis, the LMO surface was cleaned under UHV during 25 s by gentle ion sputtering to selectively remove the adventitious carbon contamination without damaging the oxide surface. The sample was then transferred in a vacuum transport vessel to the XPEEM chamber. Secondary electron images were used to localize the area where c-AFM was previously performed.

B.3.3 Data Treatment of Photoemission Spectra

The following methodology was used to analyze the La 3d, O 1s, Mn 2p and Mn 3s core levels spectra recorded. A Shirley-type background was subtracted excluding the plasmon and Auger regions from the La 3d surrounding. Quantification was performed based on the La 3d, O 1s and Mn 2p peaks. The manganese stoichiometry was calculated using the Mn 2p core level, due to its better signal to noise ratio compared to the Mn 3s spectrum. The area of each contribution was normalized by the corresponding cross-section of its core level, 29 the inelastic mean free path calculated for bulk LMO, and the transmission function of the analyzer. Peak fitting of La 3d 5/2 , O 1s and Mn 3s was performed using Pseudo-Voigt function. Although this function does not separate the effect of core-hole lifetime and instrumental broadening (since the width is the same for both Lorentzian and Gaussian contributions), they are commonly used and are considered as an accurate enough approximation for the compositional study of the LMO surface. S3 in Supporting Information) reveal a single rhombohedral phase of polycrystalline LMO without preferential orientation and with no secondary phase. Further details on the observed modifications in the Raman spectra and possible identification of phases (orthorhombic and rhombohedral) for LMO thin films with different oxygen can be found in the work by Rodríguez-Lamas et al. 30 The cation stoichiometry of the film was determined by EDX as La/Mn = 1.02 (the corresponding SEM image in an asymmetric bipolar counter-eightwise-type switching without requiring an electroforming step. In all cases noticeable differences were observed between the SET and RESET processes. An abrupt SET from LRS to HRS occurred for positive polarity at 6.5 V. Conversely, a gradual RESET was achieved by progressively decreasing the voltage in the negative polarity.

B.4 Results and Discussion

To analyze the electrical transport mechanisms, the logarithmic representation of the I-V curve is presented in Figure B.2c. For negative voltage, the log I -log V plot presents three linear regions with different slopes (m) characteristic of a space charge limited conduction (SCLC) mechanism. 31,32 At low voltage (region I) the thermal generation of free carriers inside the material dominate. As the voltage increases (region II), the density of injected carriers becomes larger than the thermally generated carriers and the SCLC regime, controlled by the density of trapped charges (m > 2), appears. 33 The injected carrier density increases with the applied voltage until the traps are completely filled leading to the sharp rise of the current (region III). The resulting trap-free SCLC regime is characterized by a quadratic Child's law (m ~2) (region IV). 34-36 For positive voltage, the behavior is similar, except for certain differences in the slope in LRS (m ̸ = 2) and the absence of the ohmic regime. These differences can be explained by the nature of the P(degenerate)/p junction between the tip and the material which is electronblocking and hole-passing type when positively biased. A particular case was the first sweep, in which a first regime with a slope ca. 1.4 is observed when operating at low voltages. We can surmise that this electronic behavior results from the heterogeneous contact formed at the nanoscale between the LMO surface in its virgin state and the highly doped tip, i.e. contact between two materials with different carrier density. 31 Thus, the resistive switching in our material can be described by a SCLC-like mechanism as a change of the transport regime from a trap-free SCL conduction to a trap-controlled SCL conduction, corresponding to the LRS and HRS, respectively. This behavior was corroborated by performing "read" measurements for both resistance states, i.e., I-V curves with small ranges of voltage (from 0 to 0.5 V). Section S.2 in Supporting Information (Figures S4 andS5) presents more details of the I-V curves analysis for the multiple cycles conducted. The ratio between the HRS and LRS was of approximately 10 2 . Both resistance states are stable and no change in resistance was observed after several days under different atmospheres, i.e. inside the glove box, under UHV and in air environment. Pulse measurements varying the voltage magnitude were carried out to optimize the SET and RESET voltages for LMO. After each pulse, the resistance value was read applying a V read = 0.5 V pulse. For writing voltages lower than 3 V, there was no significant change in the value of resistance read after (not shown). However, as shown in Figure B.3, the resistance increases by up to 4 orders of magnitude after applying a 4 V pulse for 0.5 s. On the other hand, we observe that a larger negative bias of -6 V was required to come back to the LRS state. The difference of behavior between negative and positive voltage is in agreement with the asymmetry observed in the I-V responses We have thus been able to define two characteristic voltages V SET = 4 V and V RESET = -6 V, allowing reproducible switching between the two resistance states.

To selectively obtain switched micron-scale regions, a 3x3 µm² area of virgin LMO surface was scanned by c-AFM using these two characteristic voltages, defining two concentric squares at different resistance states. The resulting resistance map is represented in of 3x3 µm² at HRS where VSET was applied, and (iii) an inner square of 1x1 µm² at LRS where V SET and V RESET were consecutively applied.

B.4.3 Physicochemical Modification of the Surface

Being able to selectively obtain micron-scale regions at different resistance states opens the way to various relevant characterization studies to investigate the corresponding surface chemistry and electronic structure. To this end the chemical changes of the surface were probed by XPEEM by performing local spectroscopic measurements in 30x30 m² HRS and virgin regions. The core level spectra of each LMO component, i. The O 1s spectrum consists of two contributions. Although many studies on perovskite oxides assign these two contributions to "bulk" and "surface" oxygen 37 , given the high sur- face sensitivity of the XPEEM technique ( & 4 nm), we consider more accurate to label the low binding energy component as "sub-surface" LMO (instead of bulk) and the high binding energy component as "surface" oxygen. The first contribution corresponds to the oxygen of the first atomic layers of the material which may not have the same environment as the bulk ones, whereas the second contribution can come from oxygen compounds adsorbed on the surface (such as water, hydrocarbon contamination, etc.), or due to a drift of the perovskite oxygen from the LMO sub-surface to the surface. In this sense, oxygen movement from the sub-surface to the surface was observed in other oxides using 18 O tracking, 38 and single-negatively-charged oxygen (O -) has been shown to be thermodynamically stable as an adsorbed specie on the surface of transition-metal oxides. 39 In the virgin state (LRS, top line of Figure B.5), the main contribution to the O 1s core level corresponds to the doubly-negatively-charged O 2-ions from the LMO sub-surface. After the resistance change, a strong increase of the surface contribution is observed in the HRS region. Since resistive switching was induced in a very low oxygen partial pressure atmosphere, the surface oxygen necessarily comes, mostly, from the LMO material itself. This is in agreement with the relative elemental concentrations in the LMO matrix as determined by the compositional analysis (see Table B.1).

Only the sub-surface contribution is considered for the oxygen quantification since the surface contribution arises from oxygen atoms adsorbed on the interface, i.e. outside the lattice. A closer analysis of the O 1s spectrum shows that the composition of the virgin surface (in LRS) presents an excess of oxygen, as expected for the LaMnO 3+δ films, whereas in the HRS region the composition is approximately the one of stoichiometric LaMnO 3.0 (δ=0), which means that the material has been reduced (by a valence change of the Mn cation, since La possesses a very stable 3+ oxidation state). Even more interestingly, due to the increase of the surface contribution, the total oxygen amount is maintained constant in both cases, which indicates that the oxygen removed from the sub-surface, remains adsorbed on the surface instead of being expelled. This oxygen movement is driven by the positive bias applied on the surface by the AFM tip following the electric field lines. The sharp SET is thus related to a chemical reduction process leading to the formation of a depletion region, which would be governed by the SCLC-like mechanism, as observed in the LMO switching characteristics.

The La 3d photoelectron peak of lanthanum compounds usually shows very prominent satellites features due to core hole screening processes. 40-43 Indeed, after 3d-ionization, the initially unoccupied 4f-levels are pulled down close to the valence orbitals of the neighborhood, allowing thus charge transfer. This results in two possible configurations of the ionic system depending on whether the screening charge is transferred to a 4f-orbital or not. 44 The resulting X-ray photoelectron spectrum is therefore a combination of those final states composed by a main peak without charge transfer (pink line in Figure B.5), and two satellite peaks with charge transfer and bonding/anti-bonding states (grey lines in Figure B.5). 45-47 In particular, for the La 3d 5/2 core level component, the ratio between the main peak and the satellites is significantly different in the two regions. In the HRS spectrum, the relative intensity of the main peak increases compared to the LRS region. This is in good agreement with the removal of oxygen from LMO explained above, which leads to a diminution of the O 2p electron density. The reduction of the La surrounding environment leads to less favorable conditions for charge transfer and thus to a diminution of the satellite intensity.

The photoemission spectra of 3d transition-metal complexes such as Mn presents various multiplet structures. Mn 3s core levels are split into two contributions. This 3s doublet occurs for the unfilled 3d transition metal complexes and is due to the presence of two possible spin states, i.e. spin-up and spin-down, resulting on 3s-photoemission with two different associated 3s3d electrostatic exchange interactions. The energy separation between the high-spin and the low-spin in the 3s core-level signal may be expressed by the Van Vleck's theorem as: 48 ∆E 3s = 0.2(n + 1)G 2 (3s3d) (1) Where n is the number of unpaired 3d-electrons and G 2 the Slater exchange integral. The energy separation increases linearly against n as long as the nature of the ligands is the same. 49 In this way, a linear relationship between ΔE 3s and the oxidation state of the transition-metal has been demonstrated for simple oxides, 50 and extended to mixed valence oxides (for instance Mn 2+ /Mn 3+ in Mn 3 O 4 51 ). Data reported in the literature give an average multiplet splitting of 5.4 eV for pure Mn 3+ and 4.4 eV for pure Mn 4+ . 50,51 In the present work the initial value of the energy splitting in LRS is 4.81 eV, which is in between the values for Mn3+ and Mn4+. This indicates the presence of a mixture of both valences and confirms the nonstoichiometric nature of the virgin LMO thin films. By using Equation 1 we calculate the Mn valence state in LaMnO 3+δ surface to be of 3.6. This Mn 3+ /Mn 4+ mixed valence is responsible for the electrical conduction properties at the LRS, where traps are filled and free charge carriers are transferred between Mn and O atoms. After the resistance change, the energy splitting rises up to 5.29 eV. This latter value close to pure Mn 3+ confirms the stoichiometry of LMO in HRS shown in Table 1. Hence, a decrease of the Mn 4+ content leads to a drastic decrease of the electrical conductivity. Beyond the simple confirmation of the reduction occurring in LMO, we clearly demonstrate that the electrical properties of the films and charge neutrality during the switching response are provided by a change in the manganese oxidation state to accommodate the local non-stoichiometry caused by the oxygen ions movement. The corresponding redox reaction can be depicted by the following equation: being the cation vacancies Taking advantage of the strong Jahn-Teller (J-T) character of rhombohedral LMO, we focus here on elucidating the microstructural changes involved in the switching response. Micro-Raman spectroscopy was used to investigate the evolution of the J-T distortion in the same two areas (LRS and HRS). Figure 6 presents the most relevant region of the LMO Raman spectra (550-750 cm -1 ), as a large part of the LMO signal is screened by the intense Raman signal arising from the Si substrate.

Both Raman spectra are composed by two lines at 615 cm -1 and 663 cm -1 ascribed to the J-T distortion. 52,53 By comparing them, we observe an increase of the first contribution in the HRS region compared to the virgin one. As reported by Dubey and co-workers, this feature increases with decreasing oxygen content, thus increasing the concentration of Mn 3+ . 54 This observation confirms the previous XPEEM results, which pointed out to a change in the Mn 3+ /Mn 4+ ratio between the two resistance states (lower Mn oxidation state for the HRS). Thus, it has been experimentally confirmed that the J-T distortion is enhanced by the presence of more Mn 3+ in HRS. Considering that micro Raman is a bulk sensitive characterization technique, these observations confirm the reduction of Mn 4+ and clearly show that the space charge layer created by removing oxygen causes also important micro structural perturbations of the LMO lattice. Complementary, photoelectron images series across the photoemission threshold were acquired. 58,59 The work function represented refers to the energy difference between vacuum and the highest filled electronic state of the valence band (which is not the same as the Fermi level) and may be different for semiconductor materials such as LMO. The states screened correspond to filled states. The diminution of the oxygen content and the reduction of the manganese oxidation state from Mn 3+ / 4+ to Mn 3+ in HRS yields a slightly decrease of the WF, indicating a shrinkage of the band hybridization (i.e. a smaller overlapping between Mn 3d and O 2p bands) which results in an upward shift of the Fermi level. 58 The distinct WF values measured at both resistance states are in accordance with the charge carrier transport properties previously described.

In summary, the combination of different advanced spectroscopic techniques allowed us to disentangle the physicochemical fingerprints related to the interfacial memristive behavior of LMO. The pristine LMO is characterized by the presence of a Mn 3+ /Mn 4+ mixture which insures the electroneutrality of the material. Upon the application of positive bias, part of the oxygen is removed from the sub-surface to the surface, increasing the density of traps in the LMO and switching to HRS that is governed by a SCLC transport mechanism. Both the stoichiometry calculation and the Mn 3s splitting reveals that the surface composition is almost purely Mn 3+ . Hence, we have proved that the increase of the resistance due to the movement of oxygen comes together with a Mn redox reaction to compensate the global charge, enhancing the Jahn-Teller effect by increasing the Mn 4+ concentration. We have demonstrated that by applying a negative electric field to the surface oxygen, it is gradually re-incorporated into the LMO both from the adsorbed surface and from the surrounding LMO oxide. The amount of adsorbed oxygen is very similar to the oxygen removed during the SET stage, which explains the resistance value after the RESET is the same as for the virgin state.

B.5 Conclusion

In this work we have demonstrated that the resistance of a LaMnO 3+δ thin film can be reversibly changed within a range of two orders of magnitude using c-AFM. The possibility of changing the resistance state in a spatially defined region allowed us to combine several complementary characterization techniques to investigate the mechanism involved in the RS response. The XPEEM measurements highlight the diminution of the oxygen content is correlated with the resistance increase. Interestingly, the oxygen is not expelled from the material but remains absorbed on the surface allowing the reverse process to occur. This loss of oxygen has direct influence both on the Mn valence, which goes from 3.6 to 3.1, and on the work function that decreases by 0.28 eV. The initial presence of cation vacancies is then compensated by the apparition of oxygen vacancies resulting in a rise of the Fermi level. We also explained the I-V characteristics measured by c-AFM in agreement with a space-charges limited current model relating them with the physicochemical changes observed. In HRS, the current is strongly controlled by the presence of traps at the surface resulting from the extraction of oxygen under positive bias. These traps are filled when oxygen is re-incorporated into the sub-surface leading to a trap-free SCL mechanism. Globally these results give a complete and comprehensive picture of the mechanisms involved in interfacial resistive switching occurring at the surface of LaMnO 3+δ thin films. fréquence de 2,5 Hz, en utilisant un temps d'ouverture de l'électrovanne de 2 ms, et un débit d'alimentation en solution de 0,35 ml/min. La température de dépôt a été variée entre 500 °C et 750 °C. Alors que l'Ar a été choisi comme gaz vecteur, un mélange à 50 %/50 % d'O 2 et Ar avec une pression totale de 5 Torr a été imposé dans la chambre de réaction pendant la croissance de l'échantillon. Cette pression totale a été maintenue pendant le refroidissement, avec la possibilité d'effectuer un recuit in situ en variant les proportions d'O 2 et d'Ar à convenance.

Ces études sur la croissance du LMO ont été réalisées sur des substrats de Si 3 N 4 /SiO 2 /Si fournis par l'Institut de Microélectronique de Barcelone (IMB-CNM) et des substrats monocristallins de SrTiO 3 (001) (STO) et de LaAlO 3 (001) (LAO) fournis par Crystech.

Différentes séries d'échantillons ont été préparées pour établir les paramètres de croissance qui permettent l'élaboration des couches minces de LMO adaptées pour la fabrication des dispositifs. Les paramètres étudiés ont été i) le ratio La/Mn dans la solution de précurseurs, ii) la température de dépôt dans la chambre de réaction, iii) l'effet de l'épaisseur de la couche mince, et iv) l'effet des différents environnements de refroidissement. Tout d'abord, l'étude de la composition des couches minces en fonction du rapport La/Mn dans la solution de précurseurs a conduit à la détermination d'une gamme optimale pour laquelle la croissance de la pérovskite LMO s'effectue sans la présence d'impuretés. Le rapport La/Mn dans la solution de précurseurs doit ainsi être limité à 1,5 si la croissance est réalisée à 750 °C, tandis qu'un rapport La/Mn égal à 2 permet la croissance de LMO stoechiométrique en dessous de 700 °C. Les caractérisations chimiques combinant la spectroscopie de rétrodiffusion de Rutherford et la spectroscopie de rayons X à dispersion d'énergie ont permis de comprendre comment l'élément La est l'espèce limitant la croissance, et comment un certain excès de La est toléré par la structure avant l'apparition des phases La 2 O 3 (comme le montre les expériences de diffraction des rayons X à incidence rasante). Ensuite, l'étude de la température de croissance a permis de trouver un compromis permettant d'abaisser la température de dépôt tout en préservant la stabilité du substrat, pour un rapport La/Mn dans les couches minces proche de 1. Une dépendance claire du rapport La/Mn avec la température de dépôt a été identifiée. Des températures de croissance inférieures à 700 °C ont conduit à des couches de LMO qui répondent mieux aux exigences morpho-structurelles. Un excès d'oxygène nominal a été identifié dans tous les échantillons (LaMnO 3+δ avec δ > 0). L'excès d'oxygène implique la présence de lacunes La et Mn dans les couches. Enfin, l'étude de l'incorporation d'oxygène dans la couche à travers la modification de l'atmosphère de refroidissement a prouvé qu'il était possible d'ajuster la teneur en oxygène, comme en témoigne le changement de structure mesuré par spectroscopie Raman.

Une optimisation supplémentaire des conditions de dépôt a été nécessaire pour les croissances sur des substrats en silicium platiné (substrats Vin Karola en Pt/TiO 2 /SiO 2 /Si) compte tenu de l'instabilité thermique à haute température que présentent ces substrats. Pour ce faire, trois stratégies de dépôt ont été développées, offrant chacune la possibilité de produire des couches de LMO denses et homogènes adaptées aux applications de commutation résistive. En outre, la teneur en oxygène peut être ajustée par les paramètres de dépôt, tels que la température, le nombre d'impulsions dans l'injection du précurseur, et l'atmosphère utilisée pendant le refroidissement. La transition de la phase orthorhombique à la phase rhomboédrique a été corrélée avec la teneur en oxygène et l'état d'oxydation du Mn des couches minces de LMO. Des couches orthorhombiques, ainsi que des couches avec un mélange de phases ont été obtenues en utilisant la stratégie I, consistant en une croissance à une température limitée en-dessous de 675 °C pendant un temps inférieur à 30 min. Avec la stratégie II, qui comprend un recuit sous atmosphère d'oxygène, il était possible d'obtenir une structure purement rhomboédrique. La stratégie III est divisée en deux étapes de croissance, ce qui permet d'arriver à des épaisseurs plus importantes. Néanmoins, les couches produites ont toujours une structure rhomboédrique. Par conséquent, en combinant les paramètres de croissance par PI-MOCVD et en choisissant judicieusement la meilleure stratégie de dépôt, nous pouvons ajuster les caractéristiques des couches minces de LMO pour les dispositifs fonctionnels. Après avoir résolu les problèmes liés à l'instabilité thermomécanique des substrats platinés, des couches minces de LMO ont été intégrées sur ce substrat pour fabriquer des dispositifs verticaux ReRAM à base de pérovskite.

D.1.3 Transport d'oxygène dans les couches épitaxiées

Les études de transport d'oxygène dans des couches minces de LMO ont été réalisées dans une sélection de couches épitaxiées ou hautement texturées, correspondant à des couches qui ont suivi une croissance sur des substrats monocristallins. Ces études ont été réalisées en combinant des techniques de caractérisation morpho-structurelles (microscopie à force atomique, microscopie électronique à balayage, microscopie électronique à transmission, spectroscopie Raman, et diffraction de rayons X) avec des techniques de caractérisation chimiques comme la spectrométrie de masse d'ions secondaires (Tof-SIMS). Ces travaux ont permis de proposer une stratégie globale pour optimiser les propriétés de transport de masse d'oxygène à basse température (500-600 ºC) des couches minces fonctionnelles de pérovskite LMO. Une diffusion d'oxygène exceptionnellement élevée, avec un coefficient de diffusion D* > 1.2 x 10 -16 cm 2 /s, peut être obtenue à 500 °C dans des couches minces épitaxiées denses en LMO, qui sont riches en lacunes d'oxygène. De plus, la croissance des couches minces de LMO denses sur des monocristaux de LAO (avec un désaccord de maille plus élevé par rapport au STO) conduit à des films avec une concentration plus élevée de défauts. Ces défauts sont en effet favorables aux processus de diffusion d'oxygène et d'échange en surface et conduisent à des coefficients de transport de masse améliorés.

De plus, des structures de LMO nano-architecturées composées d'une couche mince dense à partir de laquelle des nanostructures « digitées » ont été formées pour augmenter efficacement le rapport surface/volume, ce qui se traduit par une plus grande incorporation d'oxygène. De plus, une migration optimisée de l'oxygène vers l'interface inférieure a été obtenue pour les films LMO/LAO.

Cette combinaison unique d'ingénierie des défauts et de nano-architectures morphologiques montre un grand potentiel pour la conception d'électrodes pour piles à combustible à oxyde solide 'SOFC' basse température et pourrait être étendue à l'avenir à d'autres oxydes conducteurs ioniques -électroniques mixtes 'MIEC' prometteurs pour obtenir des dispositifs fonctionnels optimisés.

D.1.4 Commutation résistive dans les hétérostructures à base de LMO/Pt

Après l'obtention des couches minces de LMO intégrées sur silicium platiné, des dispositifs de type 'top-bottom' ont été fabriqués en salle blanche afin de réaliser une analyse de performance électrique. Ces dispositifs ont une disposition verticale dans laquelle des électrodes supérieures (TiN, Au) ont été déposées par évaporation sur la surface du LMO, et où la couche de Pt, sur laquelle la synthèse du LMO a eu lieu, constitue l'électrode inférieure.

La fabrication des dispositifs pour les mesures électriques nécessite de l'utilisation d'équipements fonctionnant dans un environnement de salle blanche. Le procédé de fabrication combine la lithographie laser, l'évaporation des métaux et les procédés de décollage. La lithographie laser utilise le système 'Heidelberg instruments μPG 101', et a permis d'adapter le masque à nos besoins. Dans ce travail, nous avons conçu un réseau de 4 colonnes d'électrodes de différentes tailles, à savoir 200 x 200 µm 2 , 150 x 150 µm 2 , 50 x 50 µm 2 et 100 x 100 µm 2 , réparties sur 14 rangées, avec un total de 56 électrodes TiN. Des électrodes métalliques en TiN de 100 nm d'épaisseur ont été déposées dans un équipement Alliance Concept CT100 par dépôt physique en phase vapeur (PVD) dans la salle blanche de la PTA. Le réseau d'électrodes a été répété deux fois sur la surface de l'échantillon, afin d'obtenir un deuxième motif identique de 56 électrodes en or. Puisque l'or est un métal noble, il devrait permettre d'identifier le rôle du TiN dans l'empilement TiN/LMO pendant la commutation résistive. Les électrodes d'or d'une épaisseur de 100 nm ont été évaporées sous vide dans un évaporateur MEB 550 Plassys dans la salle blanche de la PTA. Les caractérisations électriques ont été réalisées sur une station de mesure Karl Suss avec des manipulateurs à microsonde utilisant un analyseur de paramètres semi-conducteur B1500 Agilent.

Dans une première étude, nous nous concentrons sur l'hétérostructure TiN/LMO/Pt. Ensuite nous avons établi une comparaison avec la hétérostructure Au/LMO/Pt en utilisant les

1 1 1. 2 2 1. 3 4 1. 3 . 1 8 1. 3 . 2 9 1. 3 . 3 11 1. 3 . 4 12 1. 3 . 5 16 1. 4

 1122343183293311341235164 Towards the smallest library ever built . . . . . . . . . . . . . . . . . . Memristive devices as candidate for new generation memories . . . . . Resistive switching: concept, classification and terminology . . . . . . Valence change memories (VCM) . . . . . . . . . . . . . . . . . . Oxygen ion transport in VCMs . . . . . . . . . . . . . . . . . . . Resistive switching in manganites . . . . . . . . . . . . . . . . . . LaMnO 3±δ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Resistive switching in LaMnO 3 . . . . . . . . . . . . . . . . . . . Scope of the thesis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

1. 2 3 Figure 1 . 1 -

 2311 Figure 1.1 -Memory pyramid, shows from top to bottom the fastest to slowest memories, and more expensive to cost reduced memories. The base length representing the storage capacity. Non-Volatile memories and volatile memories are separated by a gap. Filling this gap is one of the potential goals of RRAMs

Figure 1 . 2 -

 12 Figure 1.2 -Description of different operation modes according to switching conditions. a) Unipolar resistive switching. b) Bipolar counter clockwise resistive switching c) Threshold resistive switching and d) Complementary resistive switching. In all the figures the low resistance state is marked in red and the high resistance state in blue. The current compliance can be limited on one or both polarities of the device.

Figure 1 . 3 -

 13 Figure 1.3 -Perovskite structure a) Ideal cubic perovskite (Pm3m) b) Deformed orthorhombic perovskite (Pnma) c) Deformed rhombohedral perovskite (R-3c)

Figure 1 . 4

 14 Figure 1.4 -a) LMO perovskite structure diagram as a function of the oxygen content in terms of delta (δ ) for La/Mn ratio of 1 LaMnO 3±δ. Reprinted from 'LaMnO 3+δ Revisited'[100] b) Phase diagram of perovskite-type LaMnO 3 (from La 2 O 3 -Mn 2 O 3 ) re-drawn from Roosmalen et. al.[97]

Figure 1 . 5 -

 15 Figure 1.5 -Polarized Raman spectra of a) orthorhombic LaMnO3, where the main Raman modes are indicated and b) of rhombohedral LaMnO 3 where the main Raman lines are indicated. Extracted from Iliev et. al.[101]

Figure 1 . 6

 16 Figure 1.6 -a) Cell parameter dependence with temperature for LaMnO 3 , the orthorhombic structure with pseudo cubic cell parameters a o = c o (1), b o (2), and angle β o , the rhombohedral structure with pseudocubic cell parameters a r (3) and angle α r (6) and last the cubic structure with cell parameter a c (4) . Reproduced from Rudskaya et. al.[105] b) Resistivity evolution in temperature of a LaMnO 3 single crystal. Reproduced and reformated from Zhou et.al.[START_REF] Zhou | Paramagnetic phase in single-crystal LaMnO3[END_REF]

( a )

 a Growth process summary. (b) Temperature control in the PI-MOCVD JIPELEC reactor.

Figure 2 . 1 -

 21 Figure 2.1 -JIPELEC reactor schematics and growth process summary. 2.1aThe reactor is divided roughly in three main areas: 1 Evaporator, 2 transport line and 3 main chamber. 2.1b The table describes the physical location per section of said termocouples and the set temperature fixed in the temperature controllers is indicated when corresponding to a fixed parameter.

Figure 2 . 2 -

 22 Figure 2.2 -Images of the LMO thin film from precedent growth conditions, which characteristics were the starting point for deposition optimization . a) Top SEM view and b) Cross section view of the same LMO thin film sample.

Figure 2 .

 2 Figure 2.3 a presents the relation established at 750 °C of the La/Mn ratio incorporated to the film against the La/Mn ratio in the precursor solution. This relationship allowed to verify

Figure 2 . 4 -

 24 Figure 2.4 -Composition optimisation of the LMO/Si 3 N 4 /SiO 2 /Si thin film samples: SEM top view images of the LMO films grown by the injection of 10000 pulses using different La/Mn ratios in the precursor solution.

Figure 2 . 5 -

 25 Figure 2.5 -GIXRD diffractograms of the LMO/Si 3 N 4 /SiO 2 /Si films with different stoichiometry, La/Mn in the solution 2.5, 2 and 1.5 from top to bottom. Samples details: 10000 pulses and growth temperature 750°C. PDF patterns Pbnm: 00-035-1353and R-3c: 01-078-5243 for structure reference. Substrate Bragg peak is indexed by S.

Figure 2 . 6 -

 26 Figure 2.6 -Variation of thickness and La/Mn atomic ratio in La 1-x MnO 3-δ thin films with deposition temperature. Reproduced from Pignard et. al.[START_REF] Pignard | Chemical vapor deposition and structural study of La0.8 MnO3âδ thin films[END_REF]

Figure 2 . 7 -

 27 Figure 2.7 -La/Mn ratio in the thin film dependence with deposition temperature in the 650 °C -750 °C range. The green area stands for the composition range to grow LaMnO 3+δ perovskite phase at 850°C , as established by VanRoosmalen et. al. [97].

Figure 2 . 8 -

 28 Figure 2.8 -Morphology changes observed in SEM top view of the LMO/Si 3 N 4 /SiO 2 /Si films grown at different temperatures. Insets show AFM topography images linked to height scale. Grain size and roughness (rms) are indicated below each image.

Figure 2 . 9 -

 29 Figure 2.9 -Raman spectra of the LMO/Si 3 N 4 /SiO 2 /Si films grown at different temperatures a) 6000 pulses films b) 10000 pulses films

Figure 2 .

 2 Figure 2.10 -Morpho-structural changes observed in SEM top view of a) LMO/Si 3 N 4 /SiO 2 /Si films grown with different thicknesses and with different O 2 gas atmosphere during the cooling.

Figure 2 .

 2 Figure 2.11 -AFM topography of 10 nm thick LMO/Si 3 N 4 /SiO 2 /Si thin films grown at 700 °C and cooled in different atmospheres.

Figure 2 . 12 -

 212 Figure 2.12 -Raman of LMO/Si 3 N 4 /SiO 2 /Si films growth at 700 °C by increasing the number of pulses and in two different cool down environments.

Figure 3 . 1 -Figure 3 . 2 -

 3132 Figure 3.1 -Secondary electron top view image of films grown by the injection of 2500 pulses, 6000 pulses and 10000 pulses. In the AFM insets topographic images associated to heigth scale. a) LMO films on LAO b) LMO films on STO

  Figure 3.4. The Raman modes corresponding to othorhombic (O) and rhombohedral (R) LMO, have been reported in literature and the reference spectra has been presented in Figure 1.5 in Chapter 1. The sharp mode at 487 cm -1 coresponds to LAO, as do the ones 123 cm -1 and

Figure 3 . 3 -

 33 Figure 3.3 -XRD diffractograms (θ -2θ ) for films of different thickness: a) LMO films grown on LAO and b) LMO films grown on STO. LMO and LAO are indexed in pseudocubic system and STO in cubic system. The insights on the right are presented for a clearer view of peak position evolution.

Figure 3 . 4 -

 34 Figure 3.4 -Raman spectra obtained for the LMO films grown on LAO with different thicknesses. Orthorhombic (green) and rhombohedral (navy blue) reference lines are indexed.

Figure 3 . 5 -

 35 Figure 3.5 -XRD θ -2θ diffractograms of the LMO films grown on LAO of different thicknesses paired two by two to compare the oxygen content variations given by the difference in cool down environement

Figure 3 . 6 -

 36 Figure 3.6 -Raman Spectroscopy spectra show the evolution of the structure of the LMO/LAO films with thickness and in two cool down conditions. O 2 in legend stands for 50% Ar/O2 environment while Ar refers to 100% Ar environment. The theoretical positions of the Raman lines for orthorhombic LMO (olive green) and rhombohedral LMO royal blue) are indicated in dashed lines. The curves are labeled with the number of pulses used to grow each sample.

  SI 3.16, section Model describing oxygen exchange and diffusion for further details of the physical model used).

Figure 3 . 7 -

 37 Figure 3.7 -Structural characterization of the dense thin LMO films. XRD θ-2θ patterns obtained for as-deposited films: a) LMO/STO, b) LMO/LAO. LMO and LAO peaks are indexed in the pseudocubic system, , and STO in the cubic system. SEM (left) and AFM (right) top views including the AFM height scale (right), providing a topography image and a roughness reference of the films after the exchange at 600 °C: c) LMO/STO, d) LMO/LAO, respectively. TEM bright field cross section image of the films exchanged at 500 °C: e) LMO/STO, 12 nm thick, along the [001] zone axis for STO and f) LMO/LAO, 11 nm thick, along the [001] zone axis for LAO.

Figure 3 . 8 -

 38 Figure 3.8 -Structural characterization of the nano-architectured LMO films. XRD θ-2θ patterns obtained for the as-deposited films: a) LMO/STO, b) LMO/LAO. LMO and LAO peaks are indexed in the pseudocubic system and STO in the cubic system. SEM (left) and AFM (right) top views including the AFM height scale (right), providing a topography image of the films after the exchange at 600 °C: c) LMO/STO and d) LMO/LAO. TEM bright field cross section image of the films exchanged at 600 °C: e) LMO/STO, ~54 nm thick, along the [001] zone axis for STO and f) LMO/LAO, ~66 nm thick, along the [001] zone axis for LAO. Both films present a bottom dense and continuous epitaxial film (approx. 13-14 nm thick) covered with "finger-like" columnar grains separated from each other by 1-2 nm distance.

Figure 3 . 9 -

 39 Figure 3.9 -Scheme of the isotopic exchanges in a) LMO dense films, d) LMO "finger-like nanoarchitectured" films. Comparative normalized18 O isotopic fraction for LMO films grown on STO and LAO: dense thin films exchanged b) at 500 °C for 8 h and c) at 600 °C for 1 h; "finger-like nano-architectured" films exchanged e) at 500 °C for 8 h and d) at 600 °C for 1 h. The magenta diamond pattern indicates the film/substrate interface region which is not taken into account for the fitting due to the intermixing of ion species from the LMO film and the substrate. The blue pattern stands for the dense part of the film (either totality of thin dense film, or continuous part of nanoarchitectured films).

Figure 3 .

 3 Figure 3.10 -Arrhenius plots of the oxygen tracer: a) surface exchange and b) diffusion coefficients for LMO/STO and LMO/LAO epitaxial films grown by PI-MOCVD (this work). Comparison with data obtained for LaMnO3 and La0.9MnO3 pellets from Berenov et. al.[135] .

  SI

  3.17), could come from local stoichiometry variations close to the top surface. Since the oxygen self-diffusivity (D*) is proportional to the vacancy diffusivity (Dv) and to the oxygen vacancy concentration,

Figure 3 .

 3 Figure 3.11 -HRTEM cross section images of the LMO films exchanged at 500 °C. The epitaxial growth of the LMO film can be observed on top of the STO (LAO) substrate (cube on cube). a) 12 nm epitaxial LMO/STO film along the STO [001] zone axis. SAED inset in the STO [001] zone axis (LMO [010] zone axis), b) Fourier transform showing the reciprocal space related to the same section. The inverse Fourier transform was performed applying a mask to show only the planes perpendicular to the surface, c) 10 nm epitaxial LMO/LAO film along the LAO [001] zone axis, d) Fourier transform showing the reciprocal space related to same section. The inverse Fourier transform was performed applying a mask to show only the planes perpendicular to the surface. No dislocations nor defects are observed for the LMO/STO film, while a large number of defects are observed for the LMO/LAO film.

Figure 3 .

 3 Figure 3.12 -GPA analysis of LMO/STO (a) HRTEM image of the 12 nm epitaxial LMO/STO film along the STO [001] zone axis. (b) Corresponding GPA strain image of the LMO layer with respect to the STO substrate along the [10 1]LMO // [100]STO. (c) Corresponding GPA strain image of the LMO layer with respect to the STO substrate along the [101]LMO // [010]STO

Figure 3 . 13 -

 313 Figure 3.13 -HRTEM cross-section images of a ~54 nm LMO/STO film presenting a "fingerlike" microstructure. a) LMO/STO heterostructure where the perovskite structure can be observed for the highly textured LMO film and cubic substrate (STO). b) SAED along the LMO [010] zone axis (considering the Pnma structure). c) High resolution bright field image of the LMO/STO interface. d) Fourier transform inset showing the reciprocal space corresponding to the interface; inverse Fourier transform applying a mask to show only the planes perpendicular to the surface.

Figure 3 . 14 -

 314 Figure 3.14 -HRTEM cross-section images of a ~66 nm LMO/LAO film presenting a "fingerlike" microstructure. a) the LMO/LAO stack where the perovskite structure can be observed for the highly textured LMO film and LAO pseudocubic substrate. b) SAED along the LMO [010] zone axis (considering Pnma structure). c) High resolution bright field image of the LMO/LAO interface. d) Fourier transform inset showing the reciprocal space corresponding to the interface; inverse Fourier transform applying a mask to show only the planes perpendicular to the surface.

Figure 3 . 15 -

 315 Figure 3.15 -Raman spectra of the pristine and exchanged a) dense thin and b) nano-architectured LMO/LAO films. The position of both lines in the pristine spectrum is indicated by the grey dashed lines.The narrow line at 485 cm-1 is related to the LAO substrate. c) XRD patterns obtained for the pristine and annealed dense epitaxial LMO/LAO thin film with LMO and LAO peaks indexed in the in the orthorhombic and pseudocubic systems, respectively. For all three figures, the pattern in brown corresponds to the as-deposited (pristine) LMO film, the black one was obtained after annealing (exchanged) at 500 °C and the blue one after annealing (exchanged) at 600 °C.

Figure 3 .

 3 Figure 3.16 -Figure SI 1. Schematics of the plane sheet model used for the solution to the diffusion equationThe boundary conditions equation (3) are related to the rate of transfer of the diffusing species across the surface.The assumption made is that the rate of exchange is proportional to the difference between the concentration in the gas and the concentration in the surface at any time. The Boundary condition: -D * ∂ c * (l,t)

  ) where β n (n=1, 2, . . . ) are the nonnegative roots of β tan β = L, L = lk * /D * SIMS ions depth profiles Fig. SI 2a corresponds to one of the depth profile measurements carried out on the LMO/STO film and Fig. SI 2b to a measurement carried out on the LMO/LAO film, both of them after the 18 O exchange at 500 ºC. All the ionic species corresponding to LaO 2-ions (La 16 O 2-+La 16 O 18 O +La 18 O 2 ), MnO -ions (Mn 16 O -+Mn 18 O -), and TiO -ions (Ti 16 O -+ Ti 18 O -), as well as the 18 O -and 16 O -ions are visible across the LMO film. An abrupt interface between LMO and STO (LAO) is observed. It is characterized by the abrupt decay of LaO 2-and MnO -ions and the corresponding sharp increase of TiO -(AlO 2 2-) ion signal.The dash-dot line at distance zero stands for the interface calculated as the midpoint of the decaying slope of the MnO -ions. The distance between the LMO surface and the zero line corresponds to 12 and 11 nm for the LMO/STO and LMO/LAO thin films, respectively.

Figure 3 . 17 -

 317 Figure 3.17 -SIMS ion depth profiles as a function of the distance for the samples exchanged at 500 °C for 8 h: a) LMO/STO and b) LMO/LAO films. The secondary ions plotted correspond to 18 O, 16 O, total LaO 2 -ions (La 16 O 2 -+La 16 O 18 O +La 18 O 2 ), total MnO -ions (Mn 16 O -+Mn 18 O -) and a) total TiO -ions (Ti 16 O -+ Ti 18 O -)or b) total AlO 2 -ions (Al 16 O 2 -+Al 16 O 18 O +Al 18 O 2 -). The zero position corresponds to the LMO/STO and LMO/LAO interface. The magenta diamond pattern indicates the film/substrate interface region which is not taken into account for the fitting due to the intermixing of ion species from the LMO film and the substrate. The blue pattern indicates the epitaxial LMO dense film.

Figure 3 .

 3 Figure 3.18 -Figure SI 4 Comparative normalized18 O isotopic fraction for dense thin films LMO films grown on STO and LAO exchanged at 600 °C for 1 h. The magenta diamond pattern indicates the film/substrate interface region which is not taken into account for the fitting due to the intermixing of ion species from the LMO film and the substrate. The blue grid pattern indicates the second region for the 2 region fit for LMO/LAO. The red (LMO/STO) and blue (LMO/LAO) dashed lines correspond to the best fit to Crank's solution to the diffusion equation for a solid slab. The black line corresponds to the FEM modelling solution with 2 different regions in the LMO film grown on LAO, a defect-rich region at the top of the film and a defect-poor region at the bottom.

Figure 4 . 1 -

 41 Figure 4.1 -Growth diagrams of the three deposition strategies. The deposition steps correspond to temperature 'plateaux' at 5 Torr using a gas mixture of 50% O 2 + 50% Ar. I) single-step strategy; II) single-step strategy with an annealing treatment (i.e., anneal in oxygen) at 500 °C and 5 Torr; III) two-step strategy scheme.

Figure 4 . 2 -

 42 Figure 4.2 -Representation of the range of appropriate deposition conditions. Maximum temperature used versus time of exposure to temperatures above 500 °C. Red squares stand for single step strategy (strategy I), blue triangles for single step strategy with in situ annealing (strategy II).Green circles stand for double step strategy (strategy III). Data in the green shadowed area correspond to successfully grown dense and homogeneous LMO films without instability issues. The red shadowed region corresponds to films with Pt instability problems. Purple area corresponds to columnar growth and opening of the grains, leading to porous films.

4. 4 89 Figure 4 . 3 -

 48943 Figure 4.3 -SEM surface images of the LMO films deposited by, A) strategy I, B) strategy II, and C) strategy III. The grain size increases from strategy I to III (i.e., 15, 18, 22 nm, respectively). All presented LMO thin films are of the same thickness of ca. 100 nm.

  Figure 4.3 shows the comparison of the surface films morphology with the same thickness (80-100 nm) grown by strategy I (Figure 4.3A), strategy II (Figure 4.3B) and strategy III (Figure 4.3C). All films are polycrystalline and highly compact with average grain sizes of 15 ± 4 nm, 18 ± 4 nm, and 22 ± 4 nm, respectively. Moreover, the LMO films are homogeneous and do not exhibit porosity, cracking or dewetting.

Figure 4 . 4 -

 44 Figure 4.4 -TEM cross-section images from LMO films grown by strategy I (A) and strategy III (B). A) A continuous and homogeneous film is observed. The Si 3 N 4 coating marked with * comes from device fabrication. B) A homogeneous crystallization of LMO from bottom to top was achieved in the second deposition step.

Figure 4 . 5 -

 45 Figure 4.5 -GIXRD patterns obtained for LMO thin films grown by strategy I (cooled in Ar and cooled in O 2 ), strategy II (1 h post-anneal in O 2 ) and strategy III. The orthorhombic (Pnma for La 0.974 Mn 0.974 O 3 ) and rhombohedral (R-3c for La 0.95 Mn 0.95 O 3 ) patterns have been added for comparison (ICDD: 01-087-2015 and 04-007-6350, respectively).

Figure 4 . 6 -

 46 Figure 4.6 -Raman spectra of LMO films obtained by the three deposition strategies under different conditions of cooling or annealing. An evolution of the structure from orthorhombic to rhombohedral phases is observed (dashed lines are guidelines for the eyes).

Figure 4 . 7 -

 47 Figure 4.7 -XANES absorption in Mn K-edge data of LMO films obtained by the three deposition strategies under different conditions of cooling or annealing (same samples as in Figure 4.6). The inset corresponds to an enlargement of the region close to the edge. An evolution of the structure from orthorhombic to rhombohedral phases is observed in good agreement with Raman spectra.

Table 4 . 1 -

 41 Mn K-edge position for the four representative LaMnO 3+δ samples obtained by strategies I, II and III and estimated variation in the Mn formal valence taking as reference the orthorhombic sample.

Figure 4 . 8

 48 Figure 4.8 -A) Cross section of the LMO-based MIM structure. B) Resistive switching cycles, I-V characteristics detailing the "set" and "reset" values.

Figure 4 . 9 -

 49 Figure 4.9 -Figure S1: Surface images from secondary electrons (SE) and backscattered electrons (BSE) of a LMO film deposited by strategy I at 690 °C of maximum temperature, showing pores (black in BSE) and cracks (white in BSE) due to the dewetting of the Pt layer.

Figure 4 . 10 -

 410 Figure 4.10 -Figure S2: SEM cross section of the LMO thin film structure, deposited by strategy III (25000 pulses at 500 °C and 5000 pulses at 750 °C), showing the contrast in composition (BSE mode). The discontinuity in the Pt films is clearly visible forming holes and with Pt interdiffusion as indicated in the image. These holes provoke the appearance on pinholes at the LMO surface.

Figure 4 . 1 100 5. 1 . 1 100 5. 1 . 2 103 5. 2

 4110011100121032 Figure 4.11 -Figure S3: Columnar opening in a flower bouquet manner, in a sample deposited by strategy III (25000 pulses at 500 °C and 5000 pulses at 750 °C) A) SEM top view where the separation of the columns is visible. B) Columnar growth and opening cross section by BSE.

Figure 5 .

 5 1 consists of several steps: a) Cleaning of the samples, rinsing consecutively in acetone, isopropanol and H 2 O, followed by N 2 drying (Figure 5.1 a) b) Spin coating (2000rpm, 60s) of the AZ1512 positive photoresist followed by a 90 s bake at 100 °C on a hot plate (Figure 5.1 b) c) Laser lithography, in which the resist is exposed drawing the pattern of interest, followed by resist development: (35 s soft stirring in AZ developer diluted in distilled water 1:1) and 40 s to 1 min rinse in H 2 O (Figure 5.1 c) d) Metallization (deposition of 100 nm-thick electrodes) (Figure 5.1 d) and finally e) Lift-off process (Figure 5.1 e).

Figure 5 . 1 -

 51 Figure 5.1 -Standard fabrication flow presenting a) cleaning of the LMO film/substrate sample, b) spin coating step, c) result of lithography and development of resist, d) result of metal evaporation, e) result of lift off and f) 3D diagram of the Top-Bottom configuration sample.

Figure 5 . 2 -

 52 Figure 5.2 -Pictures of samples after devices fabrication. a) LMO/Pt sample with TiN and Au top electrodes (Top-Bottom). Sample with Ti and Pt top electrodes on b) LMO/Si 3 N 4 /SiO 2 /Si and on c) LMO/STO

Figure 5 . 3

 53 Figure 5.3 -a) Diagram of electrodes virtual masks for top-bottom configuration. b) Diagram of electrodes virtual mask for top-top configuration. The masks are designed in Layout Editor for the deposition of two different metals .

Figure 5 . 4 -

 54 Figure 5.4 -HAXPES experiment scheme: 3D cross-section of a device adapted for HAXPES experiments showing the incident hard X rays (energy 4.95 to 45 keV) and the emitted photoelectrons collected by the detector. Other emitted radiation as X-ray diffraction is depicted as well.

Figure 5 . 5 -

 55 Figure 5.5 -Fabrication flow of advanced characterisation compatible devices. HAXPES compatible devices required of a 3 layered mask and fabrication procedure of several steps, divided here in 3 cycles. Cycle 1: a) first spin coating step, b) result of 1st lithography and development of resist, c) result of reactive ion etch (RIE), d) result of resist lift-off. Cycle 2: e) resist spin coating, f) laser lithography performed with alignment steps to properly overlap the mask layers, g) Evaporation of 5 nm of metal, and h) performance of metal lift off. Cycle 3: i) spin coating of positive resist, j) aligned laser lithography for the 3rd mask layer and development, k) second metal evaporation (100 nm) and l) lift off.

Table 5 . 1 -

 51 RIE recipe optimisation. Four tested recipes to obtain the optimised conditions for the elimination of the passivating layer. Boxes highlighted in yellow indicate the changes for each recipe relative to the original one (previously developed by Dr. D.Pla).

Figure 5 . 6 -

 56 Figure 5.6 -HAXPES devices and mask. a) Three-layers mask for advanced characterisation compatible devices. b) 3D schematics of the heterostructure including the substrate, LMO film, Si 3 N 4 50 nm coating, active area and contact area covered by metal. c) Real picture of the fabricated devices with mask-device overlap on the top device. d) 3D cross section showing the coating of 5 nm on the active area and the top electrode for contact separated 300 µm from contact area of 100 nm thickness.

Figure 5 . 7 -

 57 Figure 5.7 -Diagrams representing the standard electrical characterisation of the LMO-based devices. a) I-V sweep: voltage vs time representation of the standard voltage bipolar sweep. The current is continuously measured. b) Chronoamperometry (I vs t) measurements: a fixed voltage is applied for a certain time duration. The current is continuously measured.c) Trains of pulses: in the diagram two different possibilities are shown: alternated positive and negative pulses or consecutive pulses of the same polarity. The pulse height (voltage) and width (duration), as well as the time between pulses can be adapted according to the needs of the experiment. d) Particular case of pulsed measurement in which the pulses sequence allows the study of the resistance switching hysteresis loops (RSHL) through the measurement of the remanent resistance values after each pulse.

Figure 6 . 1 -

 61 Figure 6.1 -AFM and SEM images of the LMO sample (Strategy I) used to prepare most of the characterised devices. AFM and SEM composed images a) large magnification where grain size and shape can be observed and b) lower magnification image of same region shows surface homogeneity.

Figure 6 . 2

 62 Figure 6.2 -a) LaO -, MnO -, O -and Pt -ions depth profile as a function of the distance from the top surface, measured for a 75 nm LMO film grown on Pt/TiO2/SiO2/Si. b) LaO -/MnO -and MnO -/O -ratios vs distance from the top surface. c) 2D image of the total SIMS ions measured for the 50 µm x 50 µm area in the centre of the SIMS crater. d) 3D representations of oxygen ions (O -) and e) Pt -ions through the LMO film and TiO 2 /SiO 2 layers.

  Devices in Top-Bottom configuration were prepared by depositing TiN and Au electrodes on the top of LMO thin film samples (60-80 nm) obtained by strategy I. The size and distribution of square pads were obtained following the protocol presented in Chapter 5. Figure 6.3 a and b show a picture and the schematic distribution of the metallic electrodes. Two series of electrodes are visible: on the left, TiN electrodes arranged in 4 columns corresponding to 4 6.2 Analysis of pristine TiN/LMO/Pt devices different pad sizes, and on the right, Au contacts with the same pad size distribution. The pad sizes are 100 x 100 µm 2 , 50 x 50 µm 2 , 150 x 150 µm 2 and 200 x 200 µm 2 for both metals. Each column is composed of 14 rows.

Figure 6 . 3 119 6. 3

 631193 Figure 6.3 -a) Optical picture of a 1x1 cm 2 LMO thin film sample deposited using the strategy I (SI) with TiN and Au electrodes deposited on top. b) Sketch of the LMO sample with intercalated Au and TiN electrodes in the centre reproducing the real sample configuration. For clarity, dimensions of the side of the squared pads are indicated on top, and electrodes materials at the bottom. c) Cumulative probability representation of the initial resistance (Ri) values for the TiN devices. d) Frequency diagram of the initial resistance as a function of the pad size for the TiN devices.

Figure 6 . 4 -

 64 Figure 6.4 -Step by step 'soft forming' protocol in TiN/LMO/Pt HRi devices. The evolution with voltage of the a) current , and b) resistance ( both in logarithmic scale) are represented for the three steps (t1 to t3) of the 'soft forming', followed by the first operation cycle (t4). Bipolar sweeps performed as 0 V → +V max → 0 V → -V max → 0 V. Step t1: first opening in negative polarity. Step t2: opening of a 'threshold-like' voltage switch in negative polarity. Step t3: opening in positive polarity. Double headed arrow indicates the polarity in which the sweep started in t1. Darker single sense arrows indicate the sweep direction.

Figure 6 . 5 -

 65 Figure 6.5 -Cycle by cycle 'initialisation' protocol in TiN/LMO/Pt LRi devices. The evolution with voltage of a) current (in logarithmic scale), and b) resistance are represented for the different steps of the 'initialisation'. Progressive opening and an increase in the HRS/LRS ratio developed during the first 4 cycles. t1: first opening and t2-t4: gradual opening. The green arrow indicates the beginning of the CCW cycle.

Figure 6 . 6

 66 Figure 6.6 -I-V curves (in log-log scale ) corresponding to the first (t1) and last (t3 or t4) steps of the activation process for the a) HRi and b) LRi devices. HRS t1,t3 and LRS t1,t3 stand for transient high and low resistance states at t1. HRS and LRS stand for the stable HRS and LRS reached after the activation process

Figure 6 . 7 -Figure 6 . 8 -

 6768 Figure 6.7 -LRi device cycling during the 'soft forming' protocol, represented in resistance versus voltage on the left, and current (in logarithmic scale) versus voltage on the right. From top to bottom: t0) first resistance increase given in the positive polarity (0 to 3) and 1st opening of full cycle (4 to 7); t1) CW characteristic curve represented for differentiation of LRS and HRS; and t2) cycling instability in standard Icc conditions (Icc->Icc+)

Figure 6 . 9

 69 Figure 6.9 -a) Clock wise (CW) cycling represented as I-V curves (I in logarithmic scale) obtained for LRi devices of different sizes: 150x150 µm 2 , 100x100 µm 2 and 50x50 µm 2 . b) Counter clock wise cycling (CCW) obtained after clock wise cycling (CW), obtaining symmetric I-V curves.

Table 6 . 2 -

 62 Statistics extracted from the HRi and LRi functional devices. Number of functional measured devices indicated as a fraction and as percentage. Maximum and minimum values for the HRS/LRS ratio are presented as well as the mean and median HRS/LRS ratio values.

Figure 6 . 10 -

 610 Figure 6.10 -Resistance vs area for a) HRi and b) LRi devices. Resistance and area represented in logarithmic scale, slope =-1 marked as guide for the eyes in grey dotted lines.

Figure 6 . 11 -

 611 Figure 6.11 -Cumulative probability of the HRS and LRS states for the optimised devices: for devices starting in HRi a) for all size representative devices and b) excluding the data of the 150x150 µm 2 devices (for clarity); for devices starting in LRi c) for all size representative devices and d) excluding the data of the 50x50 µm 2 devices (for clarity).

Figure 6 . 12 -

 612 Figure 6.12 -Twelfth sweep curve measured for the a) LRi devices and b) HRi devices of all sizes. The voltage sweep was carried out always from 0 V → +V max → 0 V → -V max' → 0 V. The green arrow indicates the beginning of the cycling and its sense.

Table

Figure 6 . 13 -

 613 Figure 6.13 -Asymmetric threshold resistive switching a) 150x150 µm 2 device b) 100x100 µm 2 device c) 50x50 µm 2 device.

Figure 6 .

 6 Figure 6.14 -I-V characteristics of devices with high cyclability. I (log scale) vs V for LRi devices (left column): 25 cycles for the 200 µm 2 devices, 15 cycles for the 150 x 150 µm 2 devices, 25 cycles for the 100 x 100 µm 2 devices and, 10 cycles for the 50 x 50 µm 2 devices; and HRi devices (right column): 100 cycles for the 200 x 200 µm 2 devices, 30 cycles for the150 x 150 µm 2 devices, 25 cycles for the100 x 100 µm 2 devices and 25 cycles for the50 x 50 µm 2 devices

Figure 6 . 15 -Figure 6 . 16 -

 615616 Figure 6.15 -Cyclability plot corresponding to more than 20 cycles for devices presenting different pristine state LRi devices (left): 200 x 200 µm 2 device 54 cycles, 150 x 150 µm 2 device 30 cycles, 100 x 100 µm 2 device 50 cycles, 50 x 50 µm 2 device 30 cycles; HRi devices (right): 200 x 200 µm 2 device 205 cycles, 150 x 150 µm 2 device 65 cycles, 100 x 100 µm 2 device 52 (30+22) cycles and 50 x 50 µm 2 device 50 cycles. (These are the same devices as the ones represented in Figure 6.14)

  the electric field across the oxide, with V the applied voltage and d the oxide thickness, φ B = qφ B is the Schottky barrier height energy, ε is the dielectric permittivity of the oxide, κ = ε ε 0 is the relative dielectric permittivity of the oxide layer, t ox,eq is the equivalent oxide thickness, µ is the carrier mobility, N C is the density of states in the conduction band, Φ T = qφ T is the energy of the electron traps with respect to the conduction band edge, θ is the ratio of free and shallow trapped charge, ∆G * is the free activation enthalpy, σ 0 is the electrical conductivity (at T = T 0 ), and finally E C , E F and E V are the conduction band, Fermi and valence band energies, respectively.

Figure 6 . 17 -

 617 Figure 6.17 -HRi device RS characteristics representation of I vs V in log scale. Slopes per sections are fitted in the graph.

Figure 6 . 18 -

 618 Figure 6.18 -LRi device RS characteristics representation of I vs V in log scale . Slopes per sections are fitted in the graph.

Figure 6 . 19 -Figure 6 . 20 -

 619620 Figure 6.19 -Phenomenological model scheme of activation and RS in the HRi devices. Color code: Light brown corresponds to low resistance TiN(O); dark brown corresponds to oxidazed and resistant TiNyOx. Pink LMO 'with (δ'< δ) corresponding to a more resistant LMO.Violet, LMO corresponding to a less resistant LMO

Figure 6 .

 6 22 c and d shows the O and Ti EELS

Figure 6 .

 6 Figure 6.21 -TEM images of a pristine lamella obtained by focus ion beam (FIB). a) Large cross section image of the device. ROI in the centremarked in white square. b) O EELS map superimposed on ROI, c) O and Ti EELS combined composition map d) Ti map superimposed on ROI. e) HAADF section of the ROI, f) O superimposed and Ti superimposed on ROI , and g) ADF image of the same section of the ROI.

Figure 6 . 22 -

 622 Figure 6.22 -Comparison between ADF STEM pictures for a lamella prepared from a HRS and from a pristine device. Cross section lamellas prepared from a a) HRS device and from a b) pristine device. EELS measured O and Ti composite image of the lamella from the c) HRS device and from the d) pristine device.

Figure 6 . 23 -Figure 6 . 24 -

 623624 Figure 6.23 -Comparison of I-V characteristics (in logarithmic scale) for LMO-based devices prepared from SI and SIII a) SIII LMO-based HRi 200 x 200 µm 2 device, 27 consecutive cycles performed between 2 V and -2 V with a 10 mV step. b) SI LMO-based HRi 200x200 µm 2 device, 20 consecutive cycles performed between 2 V and -2 V with a 5 mV step the Icc limited for first cycle to 75 mA and after to 100 mA in positive and negative polarity. c) SIII LMO-based LRi 50 x 50 µm 2 1 + 50 consecutive cycles performed between 1.2 V and -1.2 V with a 5 mV step (the first cycle corresponded to the last step of the 'soft forming') d) SI LMO-based LRi 50x50 µm 2 device, 12 consecutive cycles between 1 V and -1 V with a 5 mV step. (Icc 60 mA in negative and 40 mA in positive when not specified)

Figure 6 .

 6 Figure 6.25 -I-V characteristics on Au/LMO/Pt devices a) 10 cycles of threshold switching in symmetric operation Icc 100 mA for a 100 x 100 µm 2 device. b) 10 cycles of asymmetric operation to a soft switching regime for the same 100x100 µm 2 device (Icc-= 100 mA and Icc+ =40 mA). c) 6 cycles of CCW regime after resistance abrupt drop (TiN 'sof forming' protocol) for the same 100x100 µm 2 device.

Figure A. 1

 1 Figure A.1 -I/V vs V performed ex-situ on a Pt/LMO/TiN stack with the soft initializing step (blue) and the resistive switching response (black)

Figure A. 2 -

 2 Figure A.2 -(a) sample description (b) picture of a mounted sample (c)(d)HAXPES spectra of Ti 1s and O 1s for the stack TiO x /TiN/TiO x N y /LaMnO 3+δ . in the initial resistance state at 0 V.

Figure A. 3

 3 FigureA.3(a) shows the peak shifts, relative to 0 bias, of Au 4f, O 1s (LMO) and Ti 1s (TiN) for different applied voltages and resistance states. The energy shift is proportional to the potential drop in the material for a given voltage. Thus, as depicted in Figure A.3(b) the deeper the layer, the smaller the corresponding HAXPES shift (i.e. V Au > V TiN > V LMO ). The voltage applied to the device is progressively lost in the series resistances corresponding to each material or interface of the stack. Considering the Au 4f peak from the top electrode, the total resistance (R tot ) of the device and the external resistance can be calculated with the following potential divider relationship:

Figure A. 3 -

 3 Figure A.3 -(a) HAXPES shift of Au 4f, O 1s and Ti 1s depending on the applied voltage (b) equivalent circuit of the structure with R ext = 2.9 Ω (c) relative areas of all Ti contributions depending on the applied bias and the resistance state

Figure A. 4 -

 4 Figure A.4 -Sketch of the proposed oxygen exchange process occurring during RESET ( VSET > 0) and SET (VRESET < 0). Inserts: the nominal composition of the TiN electrode based on the quantification of the Ti 1s area ratios shown in Figure3(c).

Figure A. 5

 5 Figure A.5 -FigureS1(a) Grazing incidence X-ray diffraction (GIXRD) 2θ scans performed using a Rigaku-Smartlab diffractometer equipped with a rotating anode to characterize the crystal structure of LMO films grown following the double strategy presented in Ref.[31]. Experimental pattern for LMO thin film grown on Pt/TiO 2 /SiO 2 /Si. * diffraction line from Pt substrate. Rhombohedral (R-3c for La 0.95 Mn 0.95 O 3 from ICDD: 04-007-6350). (b) Raman spectra collected using a Jobin Yvon/Horiba Labram spectrometer equipped with a liquid nitrogen-cooled CCD detector. Experiments were conducted in the micro-Raman mode in a backscattering geometry using a green laser (λ = 514.5 nm). The Raman spectra were calibrated using a silicon reference spectrum at room temperature (theoretical position = 520.7 cm -1 ).In green, positions of orthorhombic Raman lines and in red, of rhombohedral Raman lines. More details in Rodriguez-Lamas et al.[31]

Figure A. 6 -

 6 Figure A.6 -(a) Sketch describing the two characterized surface regions. One corresponds to the as-deposited bare surface while the other corresponds to the whole device with TiN top electrode.Even if the theoretical shape of the beam consists of a 30 mm wide and 150 mm long ellipse, the real exposed area is way larger because of the beam tail and X-ray diffusion through the surface material. (b) N 1s spectrum recorded over the region of interest. The main contribution of the signal comes from the insulating Si 3 N 4 layer. The two peaks at lower binding energy are attributed to Ti-N-Ti bond in TiN and in N-TiO 2 matrix (where N as been substituted by O). Even if no consensus has yet been achieved for the N 1s peak around 400 eV, it is commonly admitted that it may come from oxynitride species such as Ti-O-N and Ti-N-O. (c) Comparison of the La 3d 3/2 , O 1s and Mn 3s recorded on the bare surface and over the complete structure. The two O 1s contributions of the bare LMO surface corresponds have been reported to come from the LMO bulk and the very surface. The O 1s spectrum recorded over the whole device has been described in Fig.2. However, the present measurements have been carried out with a Take-off angle of 50°i nstead of 60°in the main text. Interestingly, more important Ti-O contribution is observed in the more surface sensitive case (at 50°) confirming the fitting model for the Ti-O and O (LMO) contributions. The energy splitting of the Mn 3s core level has been used in our previous work to study the LMO surface and Mn valence state. The gap between the Mn 3s peaks is inversely proportional to the Mn valence. In the present case on the bare surface the Mn 3s splitting indicates a Mn valence close to 3+ on the bare surface and a mixed Mn 2+/3+ under the Si 3 N 4 insulating layer and TiN top electrode. This observation is in agreement with the presence of oxidized TiN and Si 3 N 4 we previously observed at the interface. During the device fabrication, oxido-reduction reactions occur between the LMO surface and the deposited material resulting to a slightly understoichiometric LaMnO 3-δ surface.

Figure A. 7 -

 7 Figure A.7 -Energy shift of the Au 4f doublet depending the applied voltage and the resistance state.The shift should be proportional with the voltage, thus for 2.5 V and 5 V we measure, respectively --eV and --eV difference. However, the energy shift for -2.5 V is larger than for -5 V due to the change of the resistance change occurring at those voltages.

Figure A. 8 -

 8 Figure A.8 -Fit of Ti 1s for different applied bias at different steps of operando measurement. Note that fit of -5 V and final 0 V are the same only shifted by 0.8 eV.

Figure A. 9 -

 9 Figure A.9 -(a) Fit of O 1s for different applied bias at different steps of operando measurement. (b) and (c) show the evolution the position and the area ratio of this contributions. It is in very good agreement with the observation made in Fig.3 based on Ti 1s.

B. 2

 2 Introduction 209 (Pr,Ca)MnO 3 , 4 SmNiO 3 , 5 SrZrO 3 , 6-8 LaSrMnO 3 , 9,10 as well as the thoroughly-studied SrTiO 3 11 have already shown memristive behavior.

B. 4 . 1

 41 Characterization of LMO Thin FilmThe surface morphology of fully dense LMO thin films determined by AFM in the tappingmode is presented in Figure B.1a. It shows grains of several tens of nanometers in diameter with a smooth surface morphology (roughness rms value 1.07 nm). The GIXRD pattern shown in Figure B.1b -according to the 04-007-6350 file of the International Center for Diffraction Data (ICDD)-, Raman spectra and high-resolution transmission electron microscopy (respectively Figure S2 and Figure

  Figure B.8 in Supporting Information).

B. 4 . 2

 42 Figure B.2a presents a sketch of the fabricated Au/LMO/Si-based memristive heterostructure, while Figure B.2b shows the hysteretic current-voltage (I-V) characteristics measured for LMO by cyclic voltammetry (the curve corresponds to the first sweep). LMO exhibits

Figure B. 1 -

 1 Figure B.1 -Thin film characterization of the LaMnO 3 /Si3N x /SiO 2 /Si heterostructure by (a) 5x5 µm² AFM topographic image obtained in the tapping-mode (insert: 250 x 250 nm²), and (b) experimental GIXRD pattern (8: Si substrate) compared with the ICDD-04-007-6350 file of LMO (La 0.95 Mn 0.95 O 3 space group R-3c, a = 5.52419 Å, c = 13.33110 Å).

Figure B. 2 -

 2 Figure B.2 -(a) Sketch of the top-top memristive device configuration, where the Au top electrode was grounded and the AFM tip on LMO thin film was biased. (b) I-V characteristics obtained for LMO in the described configuration, and (c) Plot of log(I) vs. log(V) for the positive applied voltage (right) and negative applied voltage (left) showing the different regime of SCLC-like mechanism.

  Figure B.4a, where three different regions are observed: (i) the virgin surface, (ii) an outer square

Figure B. 4 -

 4 Figure B.4 -(a) Resistance map of the LMO surface after defining two concentric squares at different resistance states by c-AFM, i.e. an outer square at HRS and an inner square at LRS and (b) Resistance profile taken along the green dashed line, detailing the different resistance states.

  Figure B.4b shows a 0.5 µm large averaged resistance profile extracted from the dotted green line crossing the scanned surface (see Figure B.4a). The first microns correspond to the virgin region, which was at initial LRS. After 2 µm, the resistance changes along the concentric squares are observed. In the central region, a LRS with a resistance value close to the initial one was attained, confirming the HRS/LRS ratio of 10 2 and the possibility of obtaining reversible switching also in the scanning mode. As shown, in LMO the changes in resistance occur mainly below the tip and laterally (horizontally) for a radius of about 100 nm around the tip. No modifications are observed in the region between the tip and Au electrode for larger distances.

  e. O 1s, La 3d 5/2 and Mn 3s, of the two regions (virgin and HRS) are shown in Figure B.5. It is noteworthy to mention that in this work only the La 3d 5/2 core level is presented because the other spin-orbit La 3d 3/2 component overlaps with the Auger La M 4,5 N 4,5 N 4,5 peak.

B. 4

 4 Figure B.5 -XPEEM analysis of the La 3d 5/2 , O 1s and Mn 3s core levels of the LMO surface. The top row corresponds to a virgin region and the bottom row after -10 V has been applied by c-AFM (HRS).

La 1 -

 1 x x (Mn 3.6+ ) y 1-y O 3 (LRS) La 1-x x (Mn 3.1+ ) y 1-y O 3-0.25 O 2-(HRS)

Figure B. 6 -

 6 Figure B.6 -Raman spectra of the LMO film collected at room temperature in virgin and HRS regions ( = 488 nm).

  Figure B.7c illustrates the schematic density of states diagrams of both HRS and LRS of our LMO films based on the model proposed by Jacobs et al. and Mellan et al..58,59 The work function represented refers to the energy difference between vacuum and the highest filled electronic state of the valence band (which is not the same as the Fermi level) and may be different for semiconductor materials such as LMO. The states screened correspond to filled states. The diminution of the oxygen content and the reduction of the manganese oxidation state from Mn 3+ / 4+ to Mn 3+ in HRS yields a slightly decrease of the WF, indicating a shrinkage of the band hybridization (i.e. a smaller overlapping between Mn 3d and O 2p bands) which results in an upward shift of the Fermi level.58 The distinct WF values measured at both resistance states are in accordance with the charge carrier transport properties previously described.

Figure B. 7 -

 7 Figure B.7 -(a) Work function mapping of the LMO surface after resistance has been locally changed by c-AFM, (b) histogram showing the difference in work function, and (c) schematic DOS of partially filled 3d shell and the effect of removing oxygen from the LMO.

1

  Structuration of the thesis. . . . . . . . . . . . . . . . . . . . . . . . . . . . Morphologic changes at different temperatures observed by SEM and AFM . 2.9 Raman of LMO/Si3N4/SiO2/Si growth at different temperatures . . . . . . . 2.10 Morpho-structural changes observed in SEM top view . . . . . . . . . . . . . 2.11 AFM topography of 10 nm films (LMO/Si3N4/SiO2/Si) . . . . . . . . . . . 2.12 Raman of LMO/Si3N4/SiO2/Si . . . . . . . . . . . . . . . . . . . . . . . . . 3.1 Morphoestructural analysis of the thickness effects . . . . . . . . . . . . . . Raman spectra obtained from LMO films grown on LAO with different thicknesses. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.5 Study of the influence o the cool down environment by XRD . . . . . . . . . 3.6 Raman Spectroscopy evolution study of the LMO structure with cool down environment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.7 Structural characterization of the dense thin LMO films. . . . . . . . . . . . . 3.8 Structural characterization of the nano-architectured LMO films . . . . . . . 3.9 Scheme of the isotopic exchanges . . . . . . . . . . . . . . . . . . . . . . . 3.10 Arrhenius plots of the oxygen tracer . . . . . . . . . . . . . . . . . . . . . . 3.11 HRTEM cross section images of the LMO films . . . . . . . . . . . . . . . . 3.12 GPA analysis of LMO/STO . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.13 HRTEM cross-section images of LMO/STO film presenting a "finger-like" microstructure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.14 HRTEM cross-section images of LMO/LAO film presenting a "finger-like" microstructure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.15 Raman spectra of the pristine and exchanged LMO/LAO films . . . . . . . . 3.16 Figure SI 1. Schematics of the plane sheet model used for the solution to the diffusion equation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.17Fig. SI 2 SIMS ion depth profiles as a function of the distance . . . . . . . . 3.18 Figure SI 3 Comparative normalized 18 O isotopic fraction for dense thin films LMO films grown on STO and LAO . . . . . . . . . . . . . . . . . . . . . .

	xiv List of Tables xxii 3.4 xx GND NEXAFS UV	Ground Near edge x ray absorption fine structure Ultra violet	List of Figures List of Tables xxi List of Tables
	GPA NVM V	Geometric phase analysis Non-volatile memory Voltage
	HAADF OIC VCM(s)	High angle anular dark field Oxygen ion conductor Valence change memory(ies)
	HAXPES PCMO V max	Hard X ray photoelectron spectroscopy Pr 1-x Ca x MnO 3 Maximum voltage
	HRi PCRAM V read	High initial resistance Phase change random access memory Read voltage
	HRS PECVD V RESET	High resistance state Plasma enhanced chemical vapor deposition Reset voltage
	HRSt PI-MOCVD V SET	Transient high resistance state Pulsed injection metal organic chemical vapour deposition Set voltage
	HRTEM PTA V th	High resolution transmission electron microscopy Plateforme Technologique Amont Threshold voltage
	I PVD V write	Current Physical vapor deposition Write voltage
	IC RAM XANES	Integrated circuit Random access memory X ray absorption near edge spectroscopy
	Icc RBS XAS	Current compliance Rutherford backscattering spectrometry X ray absorption spectroscopy
	ICP ReRAM XPS	Inductive coupled plasma Redox based resistive Random Access Memory X ray photoelectron spectroscopy
	IEDP Ri XRD	Isotope Exchange Depth Profile Initial resistance X ray diffraction
	IMB -CNM RIE XRR	Institute of Microelectronics of Barcelona -Centro Nacional de Reactive ion etch X ray reflectometry
	rms YSZ	Microelectronica Root mean square roughness Yttria stabilized Zirconia
	IMEP-LaHC ROI	Institute of Microelectronics, Electromagnetism and Photonics -Region of interest
	ROM	Laboratory of hyperfrequency and characterization Read only memory
	IRDS RRAM	International Roadmap for Devices and Systems Resistive random access memory
	JT RS	Jahn-Teller Resistive Switching
	LAO RSHL	LaAlO 3 Resistance switching hysteresis loops
	LCMO SCLC	La 1-x Ca x MnO 3 Space charge limited conduction
	LCO SEM	LaCoO 3 Scanning electron microscopy
	LMGP SIMS	Laboratoire des matériaux et du génie physique Secondary Ion Mass Spectrometry/ Time-of-Flight Secondary Ion Mass
	LMO	LaMnO 3 Spectrometry (see also: ToF-SIMS)
	LRi SOFC(s)	Low initial resistance Solid oxide fuel cell(s)
	LRS SRAM	Low resistance state Static random access memory
	LRSt STEM	Transient low resistance state Scanning transmission electron microscopy
	LSCO STO	La 1-x Sr x CoO 3-δ SrTiO 3
	LSMO TCM	La 1-x Sr x MnO 3 Thermochemical memory
	MIEC TEM	Mixed ionc-electronic conductor Transmission electron microscopy
	MIM Th S	Metal -Insulator -Metal Threshold Switching
	MOCVD T JT	Metal organic chemical vapour deposition Jahn-Teller temperature
	MOSFET ToF-SIMS	Metal-oxide-semiconductor field effect transistor Time-of-Flight Secondary Ion Mass Spectrometry
	MtM UCCS	More than Moore Unité de catalyse et chimie du solide
	NAND URS	Flash (NAND gate flash memory) Unipolar resistive switching
	Nb:STO USB	Nb:SrTiO3 (Nb doped SrTiO3) Universal serial bus

xxxiii 1.1 Memory pyramid . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.2 Description of different resistive switching according to switching conditions. 1.3 Perovskite structure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.4 LMO structure diagram as a function of the oxygen content and perovskite type phase diagram . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.5 Polarized Raman spectra of LaMnO 3 . . . . . . . . . . . . . . . . . . . . . .

1.6 Cell parameter dependence with temperature for LaMnO and resistivity evolution in temperature of a LaMnO single crystal. . . . . . . . . . . . . . . . . 2.1 JIPELEC reactor schematics. . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2 Deparature film characteristics . . . . . . . . . . . . . . . . . . . . . . . . . 2.3 Incorporation of La and Mn into the LMO film and calibration of the La/Mn ratio . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.4 Secondary electrons SEM images of first optimisation LMO/Si-based substrate 2.5 GIXRD diffractograms of LMO/Si 3 N 4 /SiO 2 /Si films. . . . . . . . . . . . . . 2.6 Variations of thickness and La/Mn atomic ratio in La 1-x MnO 3-δ with temperature 2.7 Atomic ratio La/Mn dependence with temperature. . . . . . . . . . . . . . . 2.8 3.2 XRR of LMO films grown on single crystal substrates . . . . . . . . . . . . . 3.3 Study of influence of thickness by XRD . . . . . . . . . . . . . . . . . . . .

  1.3 Resistive switching: concept, classification and terminology to move faster/easier, however the drift of cations cannot be discarded and can also occur in some cases. The drift of the vacancies, driven by the applied electric field, leads to a local accumulation of oxygen vacancies. The local accumulation of oxygen vacancies affects the valence state of the surrounding cations, so that electroneutrality is maintained. This valence change may lead to a local change in the properties of the electrode/oxide heterostructure or the oxide. The effect of oxygen anions motion and consequent valence change in the neighbouring environment was identified by Waser et. al in 2006. They showed changes are located around the structural defects in the crystal lattice in the nanoscale, in the metal oxide and close to the electrode[34][19]. This mechanism is particularly favourable in the case of non stoichiometric ternary transition metal oxides, where the cation lattice presents mixed valence states.

Table 1 .

 1 1 -Selection of articles representative of the 'State of the art' on manganite based VCMs(2004)(2005)(2006)(2007)(2008)(2009)(2010)(2011)(2012)(2013)(2014)(2015)(2016)(2017)(2018)(2019)(2020). Electrodes biased and ground (GND) can be in top-top or top-bottom configuration. Probe stands for measurments performed contacting the sample surface with a metallic tip directly, while Tip stands for C-AFM tip. RS stands for the resistive switching type: BRS stands for bipolar resistive switching, URS stands for unipolar resistive switching and Th S stands for threshold switching.

	Composition	Electrode		RS	Reference	Year
		Configuration Bias	GND
	Pr 0.7 Ca 0.3 MnO 3	Top-Bottom	Au/	YBCO BRS	Xin et.al.[76]
			YSZ
	Pr 0.7 Ca 0.3 MnO 3	Top-Bottom	Au	Pt	BRS	Kim et.al.[77]
	Pr 0.7 Ca 0.3 MnO 3	Top-Bottom	Pt	W	BRS	Liu et. al.[78]
	Pr 0.7 Ca 0.3 MnO 3	Top-Bottom	Pt	W	BRS	Liu et. al.[79]
	Pr 0.7 Ca 0.3 MnO 3	Top-Bottom	Au	Cu	BRS Norpoth et.al.[80]
	PCMO-CMO-PCMO	Top-Bottom	Au/Al	Pt	BRS	Lee et.al.[81]
	Pr 0.7 Ca 0.3 MnO 3	Top-Top	Hg	Al	BRS	Lee et al[29]
	Pr 0.48 Ca 0.52 MnO 3	Top-Bottom	Rh/	SRO	BRS Arndt et.al.[82]
			YSZ
	La 0.5 Ca 0.5 MnO 3	Top-Bottom	probe	probe BRS Duhalde et.al.[68]
	La 0.7 Ca 0.3 MnO 3	Top-Bottom	Ag	Pt/Ag BRS Shang et. al.[69]
	La 0.5 Ca 0.5 MnO 3	Top-Top	Au	Au	BRS Villafuerte et.al.[70]
	La 0.8 Ca 0.2 MnO 3	Top-Top	Ag	Ag	BRS Sheng et.al.[71]
	La 0.67 Ca 0.33 MnO 3	Top-Top	Ag	Ag	BRS Belogolovskii[72]
	La 0.7 Ca 0.3 MnO 3	Top-Bottom	Ti/	Pt	BRS	Liu et.al.[73]
			TiOx
	La 2/3 Ca 1/3 MnO 3	Top-Bottom	Ti/Cu	n-Si	BRS Rubi et.al.[74]
	La 1/3 Ca 2/3 MnO 3	Top-Bottom	Ti	n-Si	BRS Acevedo et.al.[75]
	Nd 0.7 Ca 0.3 MnO 3	Top-Top	Ag	Ag	BRS	Hsu et.al.[86]
	La 0.67 Sr 0.33 MnO 3	Top-Top	Ag	Ag	BRS Huang et.al.[62]
	La 0.7 Sr 0.3 MnO 3	Top-Top	Au	Au	BRS	Lau et.al.[60]
	La 0.7 Sr 0.3 MnO 3	Top-Top	In,Pt	In,Pt	BRS	Xu et.al.[63]
	La 0.7 Sr 0.3 MnO 3	Top-Bottom	W	Pt	BRS	Lee et.al.[64]
			(probe)
	(La,Sr)MnO 3	Top-Top	Ag	Ag	BRS Manca et.al.[65]

Table 2 .

 2 1 -Sample identifiers and sample growth conditions corresponding to the first optimisation of LMO/Si 3 N 4 /SiO 2 /Si.

Table 2 .

 2 2 -Composition and thickness data obtained from the composition calibration series of LMO/Si 3 N 4 /SiO 2 /Si

	Figure 2.3 -a)

Table 2 .

 2 3 -Sample identifiers and sample growth conditions corresponding to the temperature/thickness optimisation of the LMO/Si 3 N 4 /SiO 2 /Si films.

Table 2 .

 2 4 -Study of composition and thickness variations with deposition temperature. La/Mn ratio in the LMO/Si 3 N 4 /SiO 2 /Si films grown at temperatures ranging from 650 °C to 750 °C obtained by EDX and RBS. Oxygen atomic percentage (oxygen/total) and film thickness estimated by RBS.

Table 3 .

 3 1 -Sample identifiers including number of pulses and cool down environment and substrate used.

Table 3 .

 3 3 -Cell parameters of LMO films on LAO

Table 3 .

 3 4 -Out of plane interatomic distance for LMO/LAO thin dense films as deposited, and after annealing at 500°C and 600°C

Table 6 .

 6 1 -Summary of device activation protocols for the HRi and LRi devices including type of activation, current compliance conditions and cycling sense.

Table B .

 B 1 -Results of XPS quantification in the two resistance state regions

Figure B.3 -Pulse I-V response on LMO surface, top: the resistance read after each step with V read = 0.5 V, bottom: pulse sequence, pulse duration = 0.5 s.

Acknowledgments ACKNOWLEDGMENTS

This work has been supported by the ANR funded project "Alps Memories" (ANR-15-CE24-0018). We acknowledge SOLEIL for provision of synchrotron radiation facilities (Proposal A.2 Main 199 No. 20170169). The authors want to thank Odette Chaix-Pluchery and Hervé Roussel for Raman and X-ray diffraction measurements. Sample fabrication has been performed with the help of the "Plateforme Technologique Amont" of Grenoble with the financial support of the CNRS Renatech network. In addition, this work benefited from the facilities and expertise of the OPE)N(RA characterization platform of FMNT (FR 2542, fmnt.fr) supported by CNRS, Grenoble INP and UGA.

Acknowledgment

The authors would like to thank Laetitia Rapenne and Florence Robaut for their collaboration. This work has been supported by the ANR funded project "Alps Memories" (ANR-15-CE24-0018). Facilities of the NanoCharacterisation PlatForm (PFNC) used in this work are supported by the "Recherches Technologiques de Base" Program of the French Ministry of Research. This project has received funding from the European Union's Horizon 2020

Appendix A

Resistive switching in a LaMnO 3+δ /TiN memory cell investigated by operando hard X-ray photoelectron spectroscopy This chapter reproduces the integral content of the published article: 'Resistive switching in a LaMnO 3+δ /TiN memory cell investigated by operando hard X-ray photoelectron spectroscopy' Benjamin Meunier 1 , Eugénie Martinez 2 , Raquel Rodriguez-Lamas 1 , Dolors Pla 1 , Monica Burriel 1 , Michel Boudard 1 , Carmen Jimenez 1 , Jean-Pascal Rueff 3,4 , Olivier Renault 2 1 Univ. Grenoble Alpes, CNRS, Grenoble INP*, LMGP, F-38000 Grenoble, France * Institute of Engineering Univ. Grenoble Alpes 2 Univ. Grenoble Alpes, CEA, LETI, 38000 Grenoble, France 3 Synchrotron-SOLEIL, BP 48, Saint-Aubin, F-91192 Gif-sur-Yvette, France 4 Sorbonne Université, CNRS, Laboratoire de Chimie Physique -Matière et Rayonnement (LCPMR), F-75005 Paris, France

A.1 Abstract

Transition metal oxides are promising candidates in the development of valence change memories thanks to their ability to present valence change mechanism. The resistive switching mech-Resistive switching in a LaMnO 3+δ /TiN memory cell investigated by operando hard X-ray photoelectron spectroscopy anism of TiN/LaMnO 3+δ /Pt devices was investigated by operando hard X-ray photoelectron spectroscopy after careful in situ electrical characterization. The results presented here highlight the oxygen exchange process at the TiN/LMO interface. The active TiN top electrode acts as an oxygen getter, pumping O 2-anions which are attracted by the positive bias and repelling them under negative bias. This drift of charged defects is correlated with variations of the interfacial resistance. Our results confirm the critical role of the TiN/LMO interface as well as of oxygen drift in the resistive switching behaviour of such devices.

A.2 Main

Motivated by the growing interest in novel non-volatile resistive random-access memories (RRAMs), many new oxide-based heterostructures have emerged as potential candidates to build alternative memristive devices. Among the plethora of oxides exhibiting valence change mechanism, a number of perovskites, such as (Pr,Ca)MnO 3 , [1] SmNiO 3 , [2] SrZrO 3 , [3-5] LaSrMnO 3 , [6] as well as the thoroughly-studied SrTiO 3 [7] have shown memristive behavior. Despite their promising functional properties (high speed, large ON/OFF ratio, large programming window, high endurance), fundamental knowledge and a thorough understanding of the RS mechanisms involved are in most cases still rather scarce. This is due to the difficulty to get element-specific information from the resistive material in an operating device.

The non-stoichiometric transition metal oxide LaMnO 3+δ (LMO) shows a resistive switching behavior which was shown to be governed by the oxygen vacancy content of the material. [8,9] The apparent oxygen excess in LMO is ascribed to the formation of La and/or Mn cation vacancy sites (i.e. La 1-x x ,Mn 1-y y O 3 , viz. LaMnO 3+δ ), where represents cation vacancies and charge neutrality remains fulfilled thanks to the coexistence of Mn 3+ /Mn 4+ oxidation states. [10-13] The presence of Mn in a mixed-valence state leads to a self-doping mechanism of the material. [14] We recently investigated the microscopic mechanisms of local interfacial RS in LaMnO 3+δ by conductive atomic force microscopy (c-AFM) [15] and highlighted that the change in resistance is associated with a drift of oxygen ions to the surface and the concomitant variation in the Mn valence state, which results in modifications in the transport mechanism. However, the configuration used presents several specificities (e.g. low current density, bare surface, large active area . . . ) that do not match the requirements of a standard functional device. Thus, to go a step further in the understanding of the memristive behavior of LMO-based cells, we have now tested more realistic operating devices, in which the active oxide layer is sandwiched between a top-and a bottom-electrode. Yet, the broad features at ~620 cm -1 and ~660 cm -1 are related to the LMO rhombohedral phase, corresponding to the Jahn-Teller distortion and to an E g anti-stretching mode, respectively. For the orthorhombic phase a single narrow line corresponding to the B 2g mode would have been expected at 611 cm -1 . The details and instruments of the main characterisation techniques and the equipments which were used during the thesis are described in this appendix.

C.1 Thin film characterization C.1.1 X-Ray Diffraction

The identification of the phase, structure and orientation of the films was performed for LMO films grown on single crystal substrates (i.e. STO, LAO) trough X-ray diffraction (XRD) in Bragg Brentano θ -2θ configuration. The used equipment was a Bruker D8 Advance serie II diffractometer with monochromatic CuKα radiation (λ = 1.5406 Å)and 1D detector (Bruker Techniques and equipments Lynxeye) with a resolution in 2θ of 0.06°. In general the diffractograms were collected from 10 to 80 or from 20 to 110 °(2θ ) depending on the region of interest (ROI) of the diffractogram. The collection was performed with a step of 0.11°and 2 s by step counting time. The measurement was performed on the spinning sample.

The offset and displacement of the diffractograms have been corrected using the substrate signals as references, in the DIFFRAC.SUITE EVA program from Bruker. The diffractograms were represented with Origin2018b and are plotted in waterfall mode introducing an offset in intensity for clarity.

In the case of pollycristalline samples, a 5-circle Rigaku Smartlab diffractometer was used to obtain grazing incidence XRD (GIXRD) patterns. This permited to enhance the diffraction signal from the polycrystalline films and minimize the signal of the silicon based substrates.

Thickness, roughness and density of films grown on single crystal were also determined by X-Ray Reflectivity (XRR) with a RIGAKU Smartlab diffractometer. The data treatment and fitting simulations were performed using DIFFRACT.SUITE LEPTOS program from Bruker.

C.1.2 Transmission Electron Microscopy

For samples grown on single crystals: Lamellas in cross-sections of the LMO/substrate stacks were prepared by Laetitia Rapenne using the semi-automated polishing tripod technique with the MultiPrepTM system (Allied High Tech Products, Inc.). The PIPS II from GATAN system was used for the final polishing. Transmission electron microscopy (TEM) and high resolution TEM (HRTEM) images were recorded with a JEOL JEM 2010 LaB6 microscope operating at 200 kV with a 0.19 nm pointto-point resolution.

For devices in Pt-Silicon substrates:

Lamellas were prepared by FIB in the CEA LETI facilities by D. Cooper. The crosssection lamellas were analysed in TITAN Ultimate TEM (50 pm resolution) where images were obtained in anular dark field (ADF) and high angle ADF (HAADF). Scanning TEM (STEM) in ADF mode was used to obtain high resolution and contrast images of the devices. High dispersion electron energy loss spectroscopy (EELS) was used to do a compositional analysis in cross section of the elements distribution on the device for O (K) and Ti (L) in all 3 lamellas. For the pristine and high resistance state lamellas energy signals of the La(M) Mn(L) and N(K) were collected additionally to obtain a full element cartography of the cross section.

C.1 Thin film characterization

C.1.3 Secondary Electron Microscopy

Scanning electron microscopy (SEM) was performed in a Quanta250 environmental SEM FEG from FEI, and in a SEM FEG ZEISS GeminiSEM 500 to study the surface morphology and determine the LMO thickness using the cross section of the films. The cationic film composition was analyzed by semiquantitative energy-dispersive X-ray analyses (EDX) using an Oxford Inca Energy detector coupled to the Quanta250 SEM.

In SEM Quanta250 images were aquired at 5-10 kV in high vacuum mode at a working distance of 10 mm. While EDX was performed in semi-quantitative mode using a Cu reference for calibration. Microanalysis was performed at a working distance of 11 mm and unless specified otherwise, with a beam energy of 15 kV.

In SEM Gemini SEM 500 images were acquired with an energy of 3 kV in high vacuum mode and at a working distance of 2.6 mm.

C.1.4 Rutherford backscattering spectrometry

To perform a thorough analysis of the composition of the LMO films, a series of samples were analysed by Rutherford backscattering spectrometry (RBS). This experiment was performed at the 'Centro de micro-analisis de materiales' (CEMAM) in Universidad Autonoma de Madrid.

RBS is a very powerful analytical technique based on the measure of a backscattered beam of high energy ions. RBS is based in the elastic collision of high energy ions with the atoms constituting the material without nuclear reaction. This technique is highly suitable to quantify the composition in the case of thin films and stacks. The energy of the backscattered ions is proportional to the atomic number of the element they collided with. RBS presents a nanometric depth resolution since the ions lose energy during the flight path through the sample.

RBS is very sensible to high weight elements that will be detected at different scattered energies. The number of backscattered ions increase with the atomic number. The simulation of the spectra gives the composition and thickness of each layer in the stacks.

RBS is performed with alpha particles (He +4 ) accelerated to 2 MeV and a detector placed at a backscattered angle of 175°referred to the incident beam. Light elements like O or N are visible in the RBS spectra, but the intensity and sensibility is much lower than heavy atoms. To enhance the sensitivity to these elements, RBS is performed at energies where the nuclear reaction are possible. In this way, the signal from these elements is stronger. The simulation are performed by using the cross section of the reaction of the elements at the incident energy. In our case, an incident beam of 3.045 MeV was used. The spectra can be treated thanks to SIMNRA software (https://home.mpcdf.mpg.de/~mam/), developed by

Techniques and equipments

Max-Planck-Institut für Plasmaphysik. The simulation obtained by the standard RBS allowed to quantify the cationic concentration and thickness. The simulated profile were then used to quantify the oxygen content in the 3 MeV RBS spectra. The simulated stacks should fit both spectra.

The LMO/Si 3 N 4 /SiO 2 /Si heterostructures are well suited to be analysed by this technique since the elements from the substrate do not overlapp in energy with the elements from the film.

C.1.5 Atomic Force Microscopy

Atomic force microscopy (AFM) topography maps in tapping mode were obtained in an AFM (Veeco) Digital Instruments Dimension 3100 through the Nanoscope interface, and the Gwyddion software was used for data treatment (extraction of the root mean square (rms) roughness data). The resonance frequency was in the range of 270 ±10 kHz.

The results presented in Annex B were measured by our project partners in a different setup.

C.1.6 X-ray absorption spectroscopy

X-ray absorption near-edge spectroscopy (XANES) allows the study of the absroption spectra correspondent to atomic core level transitions given when an incident photoelectron excites said atomic core level. When the energy of an X-Ray is transferred to a core-level electron of an atom (K,L or M shell), and then ejected from the atom, the atom remains in an 'excited state'. The emission of a photo-electron leads to the relaxation of the excited state. The photo electron receives the excess energy. The absorption edges correspond to the core-level energies of the atom, and are specific to the exitation and the element.

XANES spectra at the Mn K-edge of LMO thin films were collected at the ESRF ID12 beamline (Grenoble, France). Measurements were taken under vacuum at 25 °C in fluorescence mode using a nearly constant 200 mA beam current. Two silicon photodiodes detectors were used to collect the total fluorescence, one in back-scattering geometry and a second diode mounted at 90°with respect to the incident beam. The relative variation in the Mn formal valence was carried out from the experimental recorded inflection point after XANES spectra C.1.7 X ray photoelectron spectroscopy X ray photoelectron spectroscopy (XPS) and hard x ray photoelectron spectroscopy (HAX-PES) results presented in Annexes Aand B were obtained by our partners in different setups.

Techniques and equipments

HAXPES results were obtained in two syncrotron campaings in GALAXIES line in Soleil and BL15XU line in SPring 8

C.1.8 Electrical Characterization

The resistivity of the samples was obtained using a four probe measurement setup.

Current-Voltage (I-V) characteristics and Chronoamperometry (I-t) were obtained using a source-measurement unit (Agilent B1500) using two tungsten-carbide probes operated in a voltage-control mode in a Karl-Suss probe station (Figure C.2).

C.1.9 Raman Spectroscopy

Raman spectra were recorded using a Jobin Yvon/Horiba Labram spectrometer equipped with a liquid nitrogen cooled CCD detector. The 488 nm excitation line of an Ar + laser was focused C.1 Thin film characterization 237 to a 1 µm 2 spot on the sample using an objective x100. The laser power on the sample surface was close to 0.9-1 mW. Spectra were calibrated at room temperature using a silicon reference sample with a theoretical position of 520.7 cm -1 .

Appendix D Summary (French) /Résumé de la thèse en français A summary of the thesis manuscript in french is contained in this appendix.

D.1 Contenu de la Thèse

Cette thèse porte sur l'étude de la commutation résistive des manganites de lanthane (LaMnO 3 ), phénomène associé à la présence d'ions oxygène mobiles dans sa structure. Elle a pour but l'évaluation du LaMnO 3 (LMO) comme candidat pour le matériau actif dans les dispositifs microélectroniques connus comme Mémoire Résistive (Memresist). Pour la première fois des dispositifs du type mémoire verticaux constitués d'une couche polycristalline de LMO déposée entre deux électrodes métalliques ont été fabriqués et caractérisés. Ces dispositifs ont apporté la preuve de la commutation résistive dans cette hétérostructure à base de LMO, ce qui nous a permis d'étudier les propriétés physicochimiques du matériau dans les différents Summary (French) /Résumé de la thèse en français états de résistance pour mieux comprendre les mécanismes impliqués dans le processus de commutation.

Ce résumé comporte d'abord une introduction (section D.1.1) au domaine des mémoires non-volatiles, et en particulier des mémoires à changement de valence (MCV), qui permet de présenter le sujet et la pérovskite choisie pour cette étude. Ensuite (section D.1.2), la croissance des couches minces de LMO sur plusieurs substrats, sur du nitrure de silicium (Si 3 N 4 ), sur des substrats monocristallins et sur silicium platiné, a était décrite. À partir de ces croissances, plusieurs axes de recherche se dégagent et sont brièvement décrits dans ce résumé. L'axe principal (section D.1.4) a pour objectif la fabrication et l'étude des dispositifs pour la commutation résistive utilisant le LMO intégré sur du silicium platiné. L'axe secondaire (section D.1.3) présente l'étude du transport d'oxygène dans la manganite de notre choix à partir d'un système simplifié, c'est-à-dire dans des couches de LMO épitaxiées ou hautement texturées. Cette deuxième étude est fondamentale puisqu'elle vise la compréhension des mécanismes d'échange et de diffusion des ions oxygène conduisant aux changements de valence dans les MCV. Les dispositifs électroniques de nouvelle génération doivent s'adapter à des nouvelles contraintes et spécifications comme une rapidité d'opération et une capacité de stockage plus élevées, tout en réduisant la consommation d'énergie requise pour chaque opération. Dans ce contexte, les mémoires basées sur la commutation résistive se présentent comme des candidats prometteurs pour la nouvelle génération de mémoires non-volatiles (MNV). Le terme « commutation résistive » fait référence aux phénomènes réversibles liés aux changements de résistance sous l'effet d'un champ électrique appliqué. Ces changements sont non-volatiles, quand les différents états de résistance accessibles (deux ou plus) sont préservés en l'absence de champ électrique. [19] Les mémoires de type 'random access memory' (RAM) actuellement utilisées sont des mémoires volatiles qui nécessitent un apport d'énergie pour préserver les informations qu'elles contiennent, mais les opérations d'écriture et de lecture sont très rapides. C'est dans ce contexte que les mémoires RAM résistives (RRAM) présentent leur plus grand potentiel, car elles permettent de rajouter le caractère non-volatile aux mémoires RAM, tout en préservant des rapidités d'opération comparables à celles des SRAM et DRAM actuelles (entre 5 et 20 ns) [12][13][14].

Nous nous intéressons à des RRAM à base de MCV dont le mode opératoire est lié au D.1 Contenu de la Thèse 241 changement de résistance d'un matériau ou d'une interface. Ces dispositifs offrent la possibilité de construire des mémoires plus denses, facilement industrialisables, tout en réduisant la consommation d'énergie des procédés d'écriture et de lecture.

En particulier, les mémoires RRAM basées sur un procédé d'oxydoréduction ou ReRAM (redox based resistive random access memories) présentent un mécanisme semblable aux MCV, mais pour lesquelles il n'est pas nécessaire de soumettre le dispositif au processus de formation à haute tension pour être activé. Les ReRAM permettraient aussi de présenter une commutation résistive du type filamentaire ou interfaciale.

Le mécanisme de commutation dans les MCV est lié au mouvement d'oxygène ou de lacunes d'oxygène (O/Vö). Dans ces mémoires, le transport d'oxygène est considéré comme le phénomène physique qui limite et détermine la cinétique de commutation. Il peut être simulé par un modèle de 'drift-diffusion'.

Parmi les oxydes ternaires, les manganites ont reçu un intérêt croissant en tant que matériau pour la commutation résistive car ils présentent une réponse de commutation remarquable. Le composé parent de la famille des manganites est le LaMnO 3±δ (LMO). Dans ce matériau, le lanthane présente un état d'oxydation +3 alors que le manganèse présente un état d'oxydation mixte Mn 3+ /Mn +4 qui dépend de la stoechiométrie en oxygène (3±δ) et qui assure la neutralité de charge. Cette capacité à adapter valence et à la teneur en oxygène suggère que le LaMnO 3±δ a le potentiel d'être intégré dans les MCVs.

Le LMO présente une structure de pérovskite déformée. Il peut être orthorhombique ou rhomboédrique en fonction de la teneur en oxygène. D'après Van Roosmalen et. al. [START_REF] Van Roosmalen | Phases in the Perovskite-Type LaMnO3+δ Solid Solution and the La2O3-Mn2O3 Phase Diagram[END_REF], tel que publié dans le diagramme de phase La 2 O 3 -Mn 2 O 3 la structure pérovskite LaMnO 3±δ peut s'adapter à des rapports La/Mn variables entre 0,908 et 1,202 (à 850 °C dans l'air). De plus, la conductivité du LMO présente des variations liées aux modifications de sa structure. Tandis que la phase orthorhombique a été signalée comme phase isolante, le LMO rhomboédrique montre une meilleure conductivité.

D.1.2 Croissance des couches minces

Les couches minces utilisées dans cette thèse ont été obtenues dans un réacteur JIPELEC de croissance par réaction chimique en phase vapeur à partir des organométalliques évaporés par injection pulsée, (PI-MOCVD) [START_REF] Sénateur | Synthesis and characterisation of YBCO thin films grown by injection-MOCVD[END_REF][START_REF] Szkutnik | Influence of precursor nature on the thermal growth of Tin-Indium oxide layers by MOCVD[END_REF]. Comme décrit en détail dans nos travaux publiés [START_REF] Rodriguez-Lamas | Integration of LaMnO3+δ films on platinized silicon substrates for resistive switching applications by PI-MOCVD[END_REF], nous avons travaillé avec une solution de précurseurs commerciaux tris(2,2,6,6tetramethyl-3,5-heptanedionato)lanthanum (III) [La(thd)3] et tris(2,2,6,6-tetramethyl-3,5-heptanedionato) manganèse (III) [Mn(thd)3], fournis par 'Strem chemicals' dans du m-xylène (1,3-dimethylbenzene) fourni par 'Alfa Aesar'. La concentration totale de la solution est de 0,0225 M, avec un rapport de La/Mn égal à 2. L'injection du précurseur dans l'évaporateur a été réalisée à une