In situ observation of hydride nucleation and selective growth in magnesium thin-films with environmental transmission electron microscopy

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In situ observation of hydride nucleation and selective growth in magnesium thin-films with environmental transmission electron microscopy Magnus Hamm a, Marian David Bongers a, Vladimir Roddatis a, Stefan Dietrich b, Karl-Heinz Lang b, Astrid Pundt b,* a

Institute of Materials Physics, University of Goettingen, Friedrich-Hund-Platz 1, 37077, Goettingen, Germany Institute of Applied Materials (IAM-WK), Karlsruhe Institute of Technology KIT, Engelbert-Arnold-Straße 4, 76131, Karlsruhe, Germany

b

highlights
Successful

graphical abstract

in-situ

hydrogen

loading of Mg films in the environmental TEM.
Hydride formation is accompanied by local nanocrystallisation with LAGBs.
FEM simulations give insights in the local stress distribution.
Hydride

phase

formation

is

controlled by the local stress and by the kinetics.
a

model

describes

transformation

the

including

phase the

microstructural change.

article info

abstract

Article history:

The comprehension of solute-induced phase transformations is crucial in a variety of

Received 7 August 2019

research fields such as catalysis, memory switching or energy storage. We study solute-

Received in revised form

induced phase transformations on the model magnesium-hydrogen (MgeH) system

4 October 2019

which provides high lattice expansion during the phase transformations. In situ precipi-

Accepted 10 October 2019

tation and growth of MgH2 is analyzed in an environmental transmission electron mi-

Available online xxx

croscope (ETEM), combining electron energy loss spectroscopy (EELS) and various imaging techniques. It is found that the Mg-hydride (MgH2) formation proceeds through the for-

Keywords:

mation of nanocrystals that are separated by low-angle grain boundaries. This change in

In situ environmental transmission

the microstructure is easily detectable in the ETEM and allows studying the growth of the

electron microscopy

hydride phase. The EELS results confirm the direct match between the nanocrystalline

Magnesium

microstructure and the MgH2 phase. We attribute this microstructural change to large

Hydrogen

strains and stresses between the matrix and the MgH2 created during the transformation

* Corresponding author. E-mail address: [email protected] (A. Pundt). https://doi.org/10.1016/j.ijhydene.2019.10.057 0360-3199/© 2019 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved. Please cite this article as: Hamm M et al., In situ observation of hydride nucleation and selective growth in magnesium thin-films with environmental transmission electron microscopy, International Journal of Hydrogen Energy, https://doi.org/10.1016/ j.ijhydene.2019.10.057

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Phase formation

process. Half-spherical as well as finger-like regions of MgH2 is observed in the film.

Diffusion

Combing the ETEM studies with finite element method simulations on the local stress

Mechanical stress

distribution in the lamella suggests an influence of local stresses on the growth behavior. As the FEM simulations reveal, the local stress distribution depends on the shape of the hydrided region. This includes stresses that occur after the hydride nucleation and during the growth of the hydride. Hydrogen diffusion is suggested to be fast along the Pd/MgH2 interphase as well as along the high angle grain boundaries and less fast in low angle grain boundaries in MgH2, in contrast to the very slow diffusion in the MgH2 grains. This paper highlights the interdependency of the hydride growth and its self-created local stress fields in an in situ ETEM study. We consider these results not to be limited to the MgeH system, but being of more general nature. © 2019 Hydrogen Energy Publications LLC. Published by Elsevier Ltd. All rights reserved.

Introduction Magnesium (Mg) is of research interest in varied fields: It has received renewed interest as a lightweight construction material, e.g. in the automobile industry [1e3], as well as a suitable material for orthopedic applications [4]. At the beginning of the century, high temperature superconductivity was discovered in MgB2 [5e7]. Its critical temperature of Tc ¼ 39 K remains the highest known value of all conventional superconductors. Recently, Mg was found to be promising for applications in active plasmonics by switching into its hydride phase (MgH2) and back in a hydrogen atmosphere [8,9]. Magnesium, its hydride phase and magnesium alloys are further intensely researched for energy storage applications. On the one side the materials are discussed for rechargeable batteries [10e12], while on the other side MgH2 is a promising material for the storage of hydrogen as an energy carrier [13e15]. A research field where, up until now, magnesium was rarely applied is the study of solute-induced phase transformation. These processes are of importance for catalysts [16], in energy storage applications [17,18], memory switching [19,20] and nanoparticle synthesis [21]. A. Baldi et al. and T. Naranayan et al. uses palladium (Pd) clusters in a hydrogen atmosphere to study the phase transformation in situ by scanning transmission electron microscopy - electron energy loss spectroscopy (S)TEM-EEL S [22,23]. The PdeH system is a well-known model system for the hydride formation and therefore was a logical research focus for these first studies [24]. However, important influences of phase transformations on physical properties are strains and stresses, which are formed by the incorporation of solute atoms. These lead to plasticity [25e33], as well as changes in the thermodynamic and kinetic behavior of phase transformations [34e37]. The phase transformation in the PdeH system is accompanied by a relative modest volume expansion of about DV V0 ¼ 19% [38]. The bulk MgeH system has a higher volume expansion of DV V0 ¼ 32% during its phase transformation. It is expected that the larger volume expansion results in higher strains and stresses. Thus, the in situ study of the phase transformation in the MgeH system promises to deliver a clearer picture of the influence of a phase transformation on the appearance of intrinsic strain and stress and also, of the influence of this

strain and stress on the phase transformation, especially in nanomaterials. Recently Sterl et al. combined near-field scattering microscopy with atomic force microscopy and single-particle far -field spectroscopy to study the phase transformation in single Mg nanodiscs [9]. Their results were experimentally impressive. However, they did not offer interpretation of the local phase changes. Aberration-corrected environmental transmission electron microscopy (ETEM) [39] is a modern technique to study in detail the interaction of materials with a defined gaseous atmosphere. The gas is directly introduced into the ETEM without the need for using a specialized TEM holder. Therefore, ETEM is a convenient method to study the phase transformation in the MgeH system in situ. This technique was not applied before to the MgeH system. Beattie et al. and Paik et al. studied MgH2 in the TEM [40,41] and found that MgH2 is destabilized under their electron beam and decomposes to Mg. This suggests moderate conditions for TEM studies on the formation of MgH2 from Mg in a hydrogen gas environment. Such studies have not been performed in situ in the ETEM, yet. In this study, dark-field (DF) imaging, electron diffraction imaging, and electron energy loss spectroscopy (EELS) is used for in situ investigations of the MgeH system. EELS allow distinguishing between Mg and MgH2 by their plasmon energy, as shown by Paik et al. [41,42]. These studies allow direct visualization of the hydride formation and the related microstructural changes, in Mg thin films.

Experimental A Mg film with a thickness of QUOTE (650 ± 20) nm was prepared by cathode argon-ion-beam sputtering, the used sputtering system has been presented before [43]. Above the Mg film a 35 nm thick Pd film was deposited as hydrogen entry gate and as (oxygen) protection layer. A cross sectional TEM lamella was prepared from the thin film sample by focused ion beam. The final lamella is shown in Fig. 1. Two windows are predominately thinned in the cross-sectional slice. They are marked inFig. 1by ‘window 1’ and ‘window 2’. Window 2 is the thinnest part of the lamella, recognizable by the removed FIB-Pt on top. In the following, we refer to the lateral direction

Please cite this article as: Hamm M et al., In situ observation of hydride nucleation and selective growth in magnesium thin-films with environmental transmission electron microscopy, International Journal of Hydrogen Energy, https://doi.org/10.1016/ j.ijhydene.2019.10.057

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Fig. 1 e REM image of the TEM lamella used in the in situ ETEM experiment. Two separate windows were thinned into the lamella of which two are marked in the image. The 650 nm thick Mg film was grown on top of a Si substrate. Above the Mg film an 35 nm thick Pd film was deposited as a hydrogen dissociation catalyst and an (oxygen) protection layer. The lamella was prepared in a FIB by thinning with Ga-ions, the remaining FIB Pt is on top of the Pd film.

or in-plane direction, which is in parallel to the Mg/Pd interface. It is aligned to the left/right direction in the TEM images. The vertical direction or out-of-plane direction is orthogonal to the Mg/Pd interface and aligned to the top/bottom direction in the TEM images. For a better understanding the vertical and lateral direction are also marked inFig. 1. After the preparation, the lamella was transferred to the ETEM with exposure to air. This leads to a desired oxidation of the free Mg surfaces on the lamella sides. According to the literature, an oxide layer of about 2.5 nm thickness forms on Mg during the first QUOTE 15 min of air exposure [44e46]. This layer thickness remains stable within the used transfer time. In situ investigations were carried out in a FEI Titan 80-300 ETEM operated at 300 kV and equipped with a Cs-image corrector and a post-column Quantum 965ER Gatan Image Filter (GIF). For in situ hydrogen loading experiments, the hydrogen flux was set to the maximum value of the ETEM, resulting in a hydrogen gas pressure of approximately 6:5 hPa (6:5 mbar). This pressure is expected to allow for in-situ hydrogenation of the Mg film. Bright- and dark-field images were taken at certain (given) times after the maximum pressure was reached. In between the acquisitions the electron beam was blanked. This procedure minimizes the electron exposure to the Mg film, which is reported to destabilize the MgH2 phase [40,41]. Dark-field imaging was applied to investigate the crystal structure and its change in the hydrogenation process. Darkfield images were taken with the central and selected diffraction spots, which were defined in the initial state of the sample on a diffraction pattern of the Mg film. EELS were applied to distinguish the areas of Mg and MgH2 in the film by their plasmon peaks [41,42]. For this purpose, rectangular spectrum image maps were acquired with 5 nm

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and 10 nm step sizes. The size of the spectrum image map was chosen to cover an area that contains the Si substrate, the Mg film and the Pd layer. Fig. 2 shows an example of this data presentation. Within the green rectangle given in Fig. 2 a, the spectrum image map (see Fig. 2 b) contains an EEL spectrum (as exemplarily shown in Figure c) at every position. Each EEL spectrum covers an energy loss range of 100 eV starting at approximately 10 eV. The corresponding low loss EEL spectrum, cf. Fig. 2 c, in the plasmon region contains the Mg peak (for a spectrum determined in the blue marked region) and a second peak originating from Mg plural scattering. A low loss EEL spectrum in the region marked with a red box in Fig. 2 b, shows the MgH2 peak. Both spectra contain the zero loss peak ZLP. At energy losses of approximately 51 eV and higher, the Mg L3,2 edge occurs if a sufficiently high signal-to-noise ratio is obtained (cf. b). The EEL spectra were used to distinguish between the hydrided from the non-hydrided film volume. For this purpose, contributions of the plasmon peaks of Mg at 10:5 eV and of MgH2 at 14:6 eV [41] were calculated with the ‘multiple linear least-squares’ (MLLS) fitting function in DigitalMicrograph (DM). This process yields hydrided, non-hydrided and mixed areas. Further details about the EELS acquisition and evaluation are given in the SI.

Results Fig. 3 provides dark field images of the Mg film before and after hydrogen absorption. Two different magnifications are given for the state after hydrogenation: Fig. 3 b provides the whole vertical film extension, while Fig. 3 c shows an area of about 300 nm  300 nm. Before hydrogenation, the Mg film reveals a columnar grain microstructure with fiber texture (see also a). Fiber texture typically occurs for Mg films on Si substrates [43,47]. After hydrogenation, the microstructure of the lamella has noticeably changed. Fig. 3 b shows tiny light “dots” which reveal nano-sized grains in Fig. 3 c. The size of these nano
grains is below 20 nm. Some of the nano-grains show Moire fringes. These fringes demonstrate different nano-grains laying on top of each other in the direction of the beam and verify the small nano-grain size, also in the lamella depth. As the a-phase solubility of H in Mg is negligibly small, it is very likely to assume that the nanograins reflect the hydride phase of Mg. This matching can be demonstrated by EELS, which allows distinguishing between the hydrided and the non-hydrided volume in the Mg film by investigating the phase-related plasmons peaks [41,42]. An EEL spectrum map was acquired in a selected area of the TEM lamella, at the position shown in Fig. 4 with a green rectangle. In the bright field image of Fig. 4, the bright contrast areas reveal the initial columnar grain structure of the a-Mg phase while the darker contrast area contains nanograins. An evaluation of the EEL spectrum image with MLLS fitting yields the positions of a-Mg related plasmons (shown in green) and MgH2 related plasmons (shown in blue) presence, within the green rectangle. Comparison of the image contrast with the EELS maps verifies the presence of nanograins in the very region of the hydride. The

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Fig. 2 e a) HAADF image of the Mg film on the Si substrate, including the Pd cap layer and FIBePt (PteC). The green rectangle indicates the position where the spectrum image maps were taken. b) A corresponding EEL spectrum image that contains individual EEL spectra for every measured position. c) Example of two representative individual EEL spectra in the highlighted areas shown in b) by a blue and red box. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

a-Mg part possesses columnar grain structure. This suggests a direct link between the hydrogenation process and the nanocrystallization in Mg thin films. Fig. 4 reveals regions (exemplarily highlighted by white dashed frame) in the EEL spectrum map, that contain hydrided and non-hydrided volume by the simultaneous appearance of both types of plasmons. This finding suggests a locally three dimensional growth of the hydride in the Mg film and not the migration of a two dimensional front. This finding can be confirmed by the evaluation of the MgH2 volume over time (see SI). To analyze the reversibility of the observed microstructural change, the sample was exposed to environmental conditions for 5 days, after in situ hydrogen loading. This certainly allows

for H-unloading of the partially hydrided film [48]. Afterward, the sample was transferred back into the ETEM. The related STEM bright-field image is shown in Fig. 5. The microstructural changes introduced by the hydrogenation process are still present. In Fig. 5, the nanograin microstructure is visible in the upper part of the film. In this upper part, the MgH2 appeared during the partial hydrogenation experiment. In the lower part of the film, the columnar microstructure of the Mg film is still observed. This part remained in the a-Mg phase, during the partial hydrogenation experiment (see window 1 in Fig. 1). An EEL spectrum map was acquired, as exemplarily shown with green color code in Fig. 5. No MgH2 plasmon related signal was found in this area, and also not the sample. This confirms the dehydrogenation of the lamella by air

Fig. 3 e a) Dark field images of the Mg-film before hydrogen absorption. Columnar grains are observed. b) Dark field images of the Mg-film after hydrogen absorption and hydride formation as confirmed by EELS (see Fig. 4). Tiny light “dots” are observed. Figure c) focusses on an enlarged region showing individual nanocrystals of less than 20 nm. This demonstrates nanocrystallization by hydrogen absorption in these films. Please cite this article as: Hamm M et al., In situ observation of hydride nucleation and selective growth in magnesium thin-films with environmental transmission electron microscopy, International Journal of Hydrogen Energy, https://doi.org/10.1016/ j.ijhydene.2019.10.057

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Fig. 4 e HAADF image and EEL spectrum image of the partially hydrided Mg lamella. The contrast was enhanced to see more subtle variations in the Mg film area. The green rectangular marks the area where the EEL spectrum map was taken. The resulting EEL spectrum image is set on top of STEM image to facilitate the comparison. The EEL spectrum allows separating the volumes of Mg (green) and of MgH2 (blue). The comparison between the spectrum image and the STEM image reveals a matching between the nanograin microstructure and the hydrided volume. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

exposure. But, the nanograins are still present. This demonstrates the stability of the hydrogen induced microstructural change, even after exposure to air. The time development of the hydride phase formation process is shown for window 2 of the lamella in a DF image series in b) e f). The related diffraction pattern before and after hydrogenation are provided in Fig. 6 a and g, respectively. The hydrogen pressure was approximately 6:5 hPa. At time t ¼ 0 s (Fig. 6 b), hydrogen was already introduced into the ETEM. Between the acquisitions of the DF images the beam was blanked. Therefore, the exact moment of the hydride nucleation could not be detected. After t ¼ 1000 s (Fig. 6 c) a larger region of bright contrast accompanied with nanograin microstructure is observed. As presented before, this relates to the MgH2 phase. Fig. 6 g shows the related diffractogram, showing peak broadening and stronger ring appearance. In c), the MgH2 phase is visible as a quasi-rectangle ranging from top-to-bottom of the film. No other hydride precipitate is observed in window 2. Upon further hydrogenation this precipitate grows. This further growth was found to be halfcircular shaped, as shown in Fig. 6 d,e. After almost t ¼ 7000 s window 2 is converted into the nanograin MgH2 phase, cf. in Fig. 6 f. After the nanograin MgH2 phase reaches the border of the FIB-thinned window 2, a different growth behavior is found. The hydride grows along the Mg/Pd interface in the upper part of the Mg layer. This “finger”-growth time development is shown in Fig. 7aed. Figs. 6 and 7 relate to the same time scale. The MgH2 finger grows through the thicker part between windows 2 and 1 into window 1 and arrives here at tz 4000 s. At this time in window 2, the hydride still grows (see Fig. 6eef). The MgH2 finger growth in the Mg film adjacent to the Pdinterface (in lateral film direction, Fig. 7 bed is fast,

compared to the growth of the hydride precipitate in the vertical film direction (shown in Fig. 6). A diffusion coefficient D can be estimated for the two different directions (along the Mg/Pd interface and orthogonal

Fig. 5 e HAADF image and EEL spectrum image of the Pd/ Mg/Si lamella after hydrogen loading and subsequent exposure to air. The EEL spectrum map (presented in green color signal) demonstrates the strong Mg plasmons peak contribution. Only Mg contributions are detected in the EEL spectrum image, MgH2 is not present any more after air exposure. This confirms MgH2 unloading by air exposure. The residual columnar structure (bottom part of the film) as well as the nano-grain microstructure (top part of the film) is visible. Thus, the microstructural change remains even after H-unloading. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

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Fig. 6 e DF TEM images of the growth of the nanograin MgH2 phase in window 2 of the TEM lamella. b) At time t ¼ 0 s the film predominately consists of a-Mg phase. At this time, mainly the crystal structure of the Mg film is visible, as shown in a diffraction image of the Mg phase a). Afterward, a first hydride can be seen by the contrast difference between the nanograin hydride and the fiber structure of Mg (see image b). At first, the hydride grew through the complete Mg film c) and then in a half sphere in lateral direction d) e f). After tz6800 s the whole window 2 was covered by the nanocrystalline MgH2, as seen in image f). A diffraction image of the formed nanograin MgH2 phase is shown in g). The arrowed rings marked with italic 1,2 and 3 correspond to (101), (200) and (111) planes of MgH2, respectively.

Fig. 7 e DF TEM images of the growth of MgH2 finger along the Mg/Pd interface in window 1 of the lamella, cf. Fig. 1. The finger grew outwards from the initial hydride which had formed in window 2, cf. Fig. 6. The growth of MgH2 along the interface (lateral direction) is faster than into the depth (vertical direction) of the Mg film a) e c). After the Mg/Pd interface was covered by MgH2, the hydride grew further into the Mg c) and d). Please cite this article as: Hamm M et al., In situ observation of hydride nucleation and selective growth in magnesium thin-films with environmental transmission electron microscopy, International Journal of Hydrogen Energy, https://doi.org/10.1016/ j.ijhydene.2019.10.057

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to it), by evaluating the respective growth length l over the 2 time t in one dimension Dz12 lt : The diffusion coefficient along the Mg/Pd interface, in lateral film direction, is estimated to be Dlateral z2,1015 m2/s. The orthogonal diffusion coefficient into the Mg layer, in vertical film direction, was found to be Dvertical z1017 m2/s, about two orders of magnitude smaller. This result hints in the same direction as measurements of N. Teichmann et al., who found a faster lateral diffusion in Mg thin films compared to the vertical diffusion coefficients [49]. The in situ hydrogenation TEM studies on columnar grain Mg-films reveal different aspects: a) a half-circular MgH2 phase growth from the Pd-cap layer into the Mg film, b) a finger-like growth of the MgH2 phase along the interface between the Pd-cap layer and the Mg film, c) a destabilization of the MgH2 phase in air and d) a nanograin microstructure related to the MgH2 phase. The observations a) - c) will be addressed in the following with regard to finite element method (FEM) simulations on films containing hydrides with the respective global morphologies. Thus, in the simulations we neglect the nanograin

presence. These FEM simulations allow calculating the local mechanical stresses and strains, for idealized systems. Observation d) will be discussed afterward. a) The growth behavior of the Mg hydride phase, cf. Figs. 6 and 7, is sketched in Fig. 8 to show the global hydride morphologies. The hydride phase most likely nucleates at the Mg/ Pd interface of the sample, right where the hydrogen enters.1 However, this nucleation state was not imaged in the ETEM experiments due to the blank beam condition. The hydride grows into the Mg film, first right to the bottom (Fig. 8 a). This results in a rectangular shaped hydride volume in the TEM lamella, as sketched in Fig. 8 a. The further growth is found to be half-circular in the cross sectional TEM lamella (Fig. 8 b). This observed half-circular hydride growth fits to the growth model developed by Uchida et al. [43]. Uchida et al. suggested the formation of nuclei on the surface of the Mg thin film, which grow half-spherical until a closed hydride layer is formed. In a TEM-lamella, such a half-spherical growth state is observed as the presence of a half-circular bright contrast region. The growth model of Uchida et al. does not predict the appearance of MgH2 fingers, see Figs. 7 and 8 c. The observed fingers show a fast lateral growth at the Mg/Pd interface while the vertical MgH2 growth is about two orders of magnitude slower. The directional dependency of the growth found in the ETEM experiment can be explained by considering the influence of hydrogen-induced local stresses in the Mg matrix, as well as kinetic arguments. This will be shown in the following. Hydrogen-induced stresses sii ðcH Þ affect the phase transformation in metal thin films. A thermodynamic description of the stress-influence was recently given by Wagner et al. [37]. They added a global stress contribution to the chemical potential:   cH þ E0  EHH cH  n0 aH sii ðcH Þ mH ¼ RTln r  cH

Fig. 8 e Schematic drawing of the MgH2 growth behavior found in the ETEM experiments. a) the initial nucleation was not imaged, but the hydride quickly forms a cylinder through the Mg layer. b) the hydride grows roughly halfcircular through window 2. c) the hydride forms fingers at the Mg/Pd interface after it hits the thicker edges around window 2. d) The fingers grow rapidly along the Mg/Pd interface, while growth into the film is slower.

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(1)

The first three terms contain the gas constant R, the absolute temperature T, the hydrogen concentration cH , and the maximum number of possible hydrogen sites r. These terms contain the well-known terms of the entropy contribution, the site energy E0 in the related lattice as well as the HeH interaction energy EHH [52,53]. The right-hand term contains the partial molar volume of sites occupied by interstitial species n0 , the lattice expansion factor aH and the sum of normal stresses sii (given in Einstein notation). This term describes how hydrogen-induced stress in a thin film may influence mH . For example, a purely compressive (negative) stress would shift mH to higher values, thereby destabilizing the hydride phase. Wagner et al. applied Eq. (1) for an average lateral stress in the thin films. To describe the hydride growth behavior in the ETEM lamella this average stress is of minor relevance. We suggest here that the local stress distribution is of importance: If the local stress state changes, the local chemical potential Note that the sides (in e beam direction) of the sample oxidize during the transfer to the ETEM. These oxide layers, on magnesium, are reported to be impenetrable for hydrogen, cf [50,51]. 1

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changes as well. This local change may lead to a driving force guiding the hydride growth. This local stress is modeled using FEM simulations (Comsol Multyphysics® 5.2): A Mg lamella was built with infinite dimensions in x-direction, a height h of 650 nm in z-direction and a thickness d of 170 nm in y-direction. This mimics the experimental situation. The MgH2 hydride is simulated for different, significant situations in correspondence to the sketch of Fig. 8. Further information about the general setup of these simulations, the used mesh, etc. is given in the SI. Fig. 9 shows the sum of simulated stresses sii of a halfspherical hydride nucleus located in the top of a lamella plotted on two cut planes through the lamella. Black contour lines present the initial geometry in the lamella. The volume ratio of the initial nucleus and the matrix is 5,3,105. The calculated stress distribution around the hydride nucleus reveals high tensile components (yellow/white regions) in ydirection close to the surface, while in x-direction the stresses are compressive (red/black regions). Below the nucleus, the stress is of tensile nature. This local stress distribution explains why the hydride grows at first down to the bottom of the film: On the sides in ydirection and below the hydride, the volume expansion of the

Fig. 9 e Sum of axial stresses sii around an initial MgH2 nucleus with a radius of 75 nm. The volume ratio of the initial nucleus and the matrix is 5,3,10¡5. The nucleus itself forms very high stresses (mostly compressive). Below the nucleus and on the sides in y-direction of the Mg lamella, tensile stresses are detected. The stress in the yellow marked regions exceeds the tensile strength of Mg of 170 MPa [54,55]. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

MgH2 phase creates tensile stresses. Tensile stresses reduce the local mH leading to a preferred hydride growth in these regions. This means that after the initial nucleus has formed, mH favors a growth down to the lamella bottom. Hence, the local tensile stresses lead to the formation of a cylinder. The tensile strength of Mg is about ð170 ±30Þ MPa [54,55]. The yellow areas in Fig. 9 indicate regions where the tensile stress is higher than the tensile strength of Magnesium. This suggests that plastic events occur in a rather large volume around the hydride nucleus. We note that the simulations shown in Fig. 9 overestimate the deformed volume, because the plastic deformation of the material in and around the nucleus reduces the local stresses. Therefore, the stress field in real samples is expected to be smaller. Fig. 10 shows the sum of normal stresses sii for a cylinder of MgH2 (c)). The volume ratio of the initial cylinder and the matrix is 0.14. The figure is separated in the compressive component (sii < 0) in the left part and the tensile component (sii > 0 in the right part. In each part, the corresponding opposite component is shown in white color. This missing component can be extrapolated as the simulation is linesymmetrical in the middle of the cylinder. Again, the shown stresses are maximum stress values, as the FEM simulations do not include plasticity. In the simulated cylinder, large compressive stresses appear at the bottom, visible as yellow region in the left part in Fig. 10. The tensile stress (right part in Fig. 10) shows a maximum on the outer side of the cylinder close to the top of the lamella. This large tensile component stabilizes a hydride, cf. equation (1). Therefore, it is a probable reason for the halfspherical growth (Fig. 8 b). The FEM simulation of a halfsphere (see Fig. 9) reveals a large tensile component at the sides, increasing to the bottom. Therefore, a trend towards a cylindrical shape is supported by the local stresses, as observed in the ETEM experiment (cf. Fig. 6). Overall, the FEM simulations show how the local stresses influence the hydride growth, which was found in window 2 of the ETEM experiment (Fig. 6). A half-spherical hydride nucleus forms on top of the lamella and turns into a cylinder (Fig. 6 c). Subsequently, the local stresses lead to a preferred hydride growth in the upper part of the lamella (Fig. 10), forming a half-sphere. However, a developing half-sphere creates stresses which favor a growth back to a cylinder (SI Fig. 7). Hence, the local stresses, which the hydride creates during its growth, lead to a hydride morphology between that of a cylinder and half-sphere. b) As the hydride reaches the edges of window 2 the hydride fingers form. Fig. 11 shows a FEM simulation of the finger morphology. The volume ratio of the initial morphology and the matrix is 0.26. Interestingly, compressive stresses form at the finger tips, as shown by the yellow regions in the left side of Fig. 11. This stress component should hinder their lateral spreading. Further, below the fingers tensile stresses are found, as shown by the yellow regions on the right side in Fig. 11, appearing in the edge between the finger and the cylinder. This tensile stress component would support a more half-spherical growth of the hydride. Overall, the local stress distribution would make a further growth of the fingers unfavorable. Thus, the related local mH cannot explain the growth of the finger.

Please cite this article as: Hamm M et al., In situ observation of hydride nucleation and selective growth in magnesium thin-films with environmental transmission electron microscopy, International Journal of Hydrogen Energy, https://doi.org/10.1016/ j.ijhydene.2019.10.057

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Fig. 10 e Sum of elastic normal stresses sii created by an MgH2 cylinder in the middle of a Mg lamella, as determined by FEM simulations. The volume ratio of the initial cylinder and the matrix is 0.14. The left side shows the compressive component. The highest values are found in the hydride cylinder, especially at the bottom. The right side shows the tensile component. The highest values are found outside of the cylinder, at its top. In real samples these values are expected to be reduced by plasticity. White parts reflect those volumes of the corresponding opposite stress component.

However, the growth of the fingers can be explained by considering kinetic arguments. The thicker lamella parts around window 2 (cf. Fig. 1) need more hydrogen atoms to be transported from the Pd top layer to form the hydride. As the Mg lamella surfaces are covered by oxides, the hydrogen atoms enter the Mg film only from the top side, through the Pd/Mg interface. This hydrogen gate supports the formation of fingers. We suggest that the finger formation is initiated by the tensile stress component in the upper outer part of the hydride cylinder (see Fig. 10). For the thicker areas of the lamella, the fingers may grow longer simply because hydrogen is only accessible in the immediate vicinity of the Pd/Mg interface.

Further, the diffusion of hydrogen has to be considered. Diffusion of hydrogen through Mg is known to be fast, at room temperature. Values between 1010 m2/s and 1012 m2/s are known from literature [43,56]. These values are similar to hydrogen diffusion constants in Pd. Diffusion through MgH2 is much slower. The reported values for MgH2 lay between 1028  1030 m2/s, at room temperature [57,58]. For most samples the main diffusion path is along grain boundaries, as was recently shown [59]. Most probably, these initial grain boundaries are high angle grain boundaries (HAGBs). Typical values for the grain boundary diffusion in MgH2 thin films are around 1017 m2/s, at room temperature [43,47]. This value

Fig. 11 e Sum of normal stresses sii created by an MgH2 cylinder with fingers in the middle of the Mg lamella. The volume ratio of the initial morphology and the matrix is 0.26. The left side shows the compressive component which destabilizes hydride formation. The highest values are found in the hydride cylinder and are in real samples reduced by the formation of plasticity. In addition, high compressive stresses are found at the outside of the fingers. The right side shows the tensile components which improve hydride growth. The highest values are found below the fingers, especially at the meeting point of a finger and the cylinder. Please cite this article as: Hamm M et al., In situ observation of hydride nucleation and selective growth in magnesium thin-films with environmental transmission electron microscopy, International Journal of Hydrogen Energy, https://doi.org/10.1016/ j.ijhydene.2019.10.057

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agrees well with the diffusion coefficient evaluated for the hydride growth below the finger, in vertical direction. To form hydride below the finger, hydrogen seems to diffuse through the grain boundaries in the MgH2, possibly through the initial HAGBs. This leads to the measured diffusion coefficient. It can be seen in Fig. 7 that below the finger MgH2 grows not in a uniform front but as MgH2 “lines”. These lines might reflect the positions of the initial HAGBs of the columnar Mg-phase. Their presence suggests enhanced heterogeneous hydride formation at initial grain boundaries. Hydride growth below the finger is thermodynamically favored by tensile stress components (see Fig. 11, shown by the tensile components below the initial nucleus in Fig. 9), but hydrogen needs to reach this Mg/MgH2 interface. The transport of hydrogen through the finger takes a long time, because of the comparably small grain boundary diffusion constant in the nanocrystalline MgH2 phase. Hydrogen can diffuse around the finger, with the much higher hydrogen diffusion coefficient in a-Mg or in the plastic zone at the Mg/MgH2 interface. This is what may happen during the initial growth in window 2, where we expect the growth direction just to be governed by the thermodynamic influence of local stresses. However, it seems that as soon as the finger reaches a given length the kinetic component becomes important, as for example found in window 1. We suggest that the finger growth becomes selfpromoting, because of kinetic reasons: Hydrogen propagates too slowly to reach the thermodynamically favored positions below the finger. Hydrogen absorption predominately leads to finger propagation which concurrently increases the hydrogen diffusion length. c) As demonstrated in Fig. 5 the Mg hydride in the sample unloads at ambient conditions. After exposing the sample for five days in air, no hydride is detected in the related EELSpattern. This is surprising as bulk MgH2 is rather stable and the theoretical temperature is Tdes z577 K for hydrogen release from bulk MgH2 at normal pressure [14]. We suggest the high compressive stresses in the hydride’s interior, as determined by FEM simulations, might destabilize the hydride. For a fully hydrided Mg thin film, high compressive stresses are expected (see SI in Ref. [36]). Even if these stresses are reduced by plasticity during the formation of the nanograin phase, some increased stress state should remain, in comparison to a bulk sample. We suggest that this causes the dehydrogenation even at room temperature. It is noteworthy that the hydrideinduced nanograin microstructure remains after unloading. d) The hydride-formation induced irreversible nanocrystallization can be interpreted by combining several material science aspects with the ETEM observations: The lattice strain during the first dihydride formation is extremely large, in the MgeH system. For the fiber textured thin Mg film, vertical expansion of about 25% is expected [48] if relaxation processes by plastic deformation are excluded. This calculated mechanical stress is above the tensile strength of the material (see Fig. 9) and can lead to plasticity. Possibly, this is further enhanced by transformation-induced plasticity (TRIP) [60]. This process is known for phase transformations in steels like e.g. the martensitic transformation and has recently also been discovered for carbides growing in AISI 4140 steels [61]. TRIP conventionally occurs in the material with the lower elastic modulus which, here, is the metallic

hexagonal a-Mg-matrix. Thus, the TRIP zone is expected to appear in a seam enclosing the MgH2-nucleus. As large mechanical stresses appear, several prismatic glide systems can be activated simultaneously in the Mgmatrix. This may lead to dislocation cross linkage and local hardening. We speculate that this hardening process in the TRIP zone may hinder the further growth of the initial MgH2 nanograins. The dislocation-rich zone can shrink to form lowangle grain boundaries (LAGBs). As the experiments reveal, the hydride phase grows with an orientation relation to the original columnar grains. The diffraction pattern of the MgH2 phase provided in Fig. 6 g shows spots with an ark length of about 14 . Thus, different nanograins are tilted against each other by about ±7 , or by 14 in maximum. According to this, the MgH2 nanograins are suggested to be separated by LAGBs. In a subsequent process, the dislocation-rich TRIP zone can help the nucleation of nanograin hydrides in a heterogeneous process. This supports our interpretation. In a subsequent process, the dislocation-rich TRIP zone can help the nucleation of nanograin hydrides in a heterogeneous process. The transport of hydrogen through the TRIP zone between the hydride grain and the a-matrix is expected to be fast. Considering the different nature of HAGBs and LAGBs, HAGB probably play a more important role for the transport through the MgH2 than LAGB. LAGBs can be described as dislocation arrangements and do provide H-diffusion paths mainly along the dislocation line, but not in the grain boundary plane, like HAGBs do. Generally, diffusion in HAGBs is found to be faster than diffusion in LAGBs [62]. Therefore, we speculate that the hydrogen transport in MgH2 along the newly created LAGBs is slower that through the HAGS. However, in MgH2 both are expected to be significantly faster than diffusion in the grains. We suggest that this TRIP-zone controlled nucleation e and growth process can lead to the observed final appearance of the nanograins in the MgH2 phase. This may also affect the behavior in other Mg-based storage alloys [63].

Conclusions In conclusion, due to the large volume expansion during the Mg to MgH2 phase transformation MgeH thin films can be regarded as a model system for studies on stress-affected phase transformations. The volume expansion creates large strains, stresses and transformation-induced plasticity. ETEM revealed several important results as the phase transformation de facto can be studied in situ. For MgeH, the phases can be easily assigned because of the microstructural change. In combination with FEM-simulations it was found that the phase transformation is strongly influenced by mechanical stresses. While this was already shown by other groups on a global scale, this work emphasizes the importance of the local scale stress arising during the growth process. On the local scale, stress fields lead to local plasticity and nanocrystallization. Additionally, the local stress field is, to our interpretation, one major factor in directing the phase transformation, diffusion kinetics being the other. We suggest that the increased stress state found in Mg thin films with respect to bulk samples, leads to a destabilization of the hydride at

Please cite this article as: Hamm M et al., In situ observation of hydride nucleation and selective growth in magnesium thin-films with environmental transmission electron microscopy, International Journal of Hydrogen Energy, https://doi.org/10.1016/ j.ijhydene.2019.10.057

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ambient conditions. This may be explained by a stress-related change of the hydrogen chemical potential. We consider all of these results not to be limited to the MgeH system, but being of more general nature. ETEM now offers the possibility to study solute-induced phase transformation in details.

Supporting information Additional information is provided on the experimental details, on the evaluation of the diffusion processes, the evaluation of the EELS data and on the FEM simulations. Also, some experimental results are supplied regarding the MgePd interface vicinity.

Author contributions M.H., M.B. and A.P. have developed and supervised the project idea. M.H. and M.B wrote the initial of the manuscript. M.H., M.B. and V.R. performed experimental measurements, and analyzed the data. M.H. performed finite element simulations. S.D. and K.-H.L. contributed to the interpretation of the data. All authors discussed the results and contributed to the final manuscript preparation.

Acknowledgments We gratefully acknowledge financial support by the Deutsche Forschungsgemeinschaft (DFG) via the Heisenberg-grant (PU131/9-2) and the project PU131/10-1,2. We thank Matthias Hahn and Torben Erichsen for support at the ETEM.

Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.ijhydene.2019.10.057.

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Please cite this article as: Hamm M et al., In situ observation of hydride nucleation and selective growth in magnesium thin-films with environmental transmission electron microscopy, International Journal of Hydrogen Energy, https://doi.org/10.1016/ j.ijhydene.2019.10.057