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The effects of dynamic structural transformations on hydrogenation properties of Mg and MgNi thin films
international journal of hydrogen energy 34 (2009) 5131–5137
Available at www.sciencedirect.com
journal homepage: www.elsevier.com/locate/he
The effects of dynamic structural transformations on hydrogenation properties of Mg and MgNi thin films L. Praneviciusa,*, D. Milciusb, C. Templierc a
Vytautas Magnus University, 8 Vileikos St., LT-44404 Kaunas, Lithuania Lithuanian Energy Institute, 3 Breslaujos St., LT-44403 Kaunas, Lithuania c Universite´ de Poitiers, SP2MI, Te´le´port 2, Bd Marie et Pierre Curie BP 30179, Futuroscope, France b
article info
abstract
Article history:
Mg capped with Al and Ti thin layers and MgxNi films have been sputter-deposited on
Received 19 March 2009
quartz substrates and hydrogenated at 600 kPa for 250 C. A complete fast transformation
Received in revised form
of metallic into hydride phase is registered for the films, demonstrating the dynamic state
17 April 2009
of the internal microstructure under hydrogenation. It leads to local and long-range
Accepted 18 April 2009
restructuring and the fast hydrogenation rate is attributed to the fast hydrogen uptake and
Available online 14 May 2009
transport along columns and grain boundaries of nanocrystallites. A slow H-loading is observed when the dynamic structural relaxation processes are suppressed by internal and
Keywords:
external inhomogeneities such as barrier layers on the surface, new phases in the bulk and
Hydrogenation
impurities. A partial transformation of metallic into hydride phase is registered when the
Thin films
structural formations newly nucleated at the initial stages of hydrogenation suppress
Mg and MgNi
dynamic processes and prevent the H-uptake. ª 2009 International Association for Hydrogen Energy. Published by Elsevier Ltd. All rights reserved.
1.
Introduction
Mg-based materials have attracted extensive attention for application in hydrogen storage technology due to the high storage capacity of the hydride phase. Nearly 7.6 wt% of hydrogen can be stored as MgH2 in Mg and nearly 3.6 wt% hydrogen reacts with Mg2Ni to form Mg2NiH4 in a reversible reaction [1,2]. Unfortunately, the hydride phase of Mg-based materials is very stable and shows poor hydrogen absorption and desorption kinetics even at high temperature. In the last few years, substantial improvement in hydrogen absorption and desorption kinetics has been achieved by several different methods [3–5]. The most effective methods realized so far are based on structural refinements of the material and the addition of proper catalysts. Nanocrystalline and amorphous magnesium-based alloys demonstrate promising
performance as hydrogen storage materials. Nanoscale particles are advantageous over bulk materials as they have a larger solid/gas interface area and shorter hydrogen diffusion paths, yielding potentially faster kinetics for gas absorption and desorption. A proper way to make suitable nanocomposite for hydrogen storage is to produce thin films by sputtering or evaporation methods as it allows controlling the composition and structure of the film layers at a nanometer scale. Thus, thin films have been widely used to study hydrogen behavior properties of the corresponding bulk materials. However, the fundamental difference with respect to bulk samples is the presence of a substrate which might alter kinetics and thermodynamics of the processes. Hydrogen in metals expands the metal host lattice. A thin film cannot expand freely in all dimensions as adhesion forces at the film–substrate interface
* Corresponding author. Tel.: þ370 37 327909; fax: þ370 37 327916. E-mail address: [email protected] (L. Pranevicius). 0360-3199/$ – see front matter ª 2009 International Association for Hydrogen Energy. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.ijhydene.2009.04.034
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international journal of hydrogen energy 34 (2009) 5131–5137
prevent the expansion of the film. It leads to considerable inplane stresses of the order of several GPa [6,7]. Nucleation rates usually involve rapidly varying free energies which are extraordinary sensitive to the conditions under which the experiment is run (pressure, impurity concentrations, as well as temperature). Reproducibility of experimental data is difficult to achieve and the theoretical prediction of nucleation rate is subject to large uncertainties. The effects of stress generation and relaxation on the hydride phase nucleation mechanism and hydrogen storage properties are not fully understood. It has been shown [8,9] that stress relaxation can occur by diffusional deformation wherein mass is transported between the free surface of the film and the grain boundaries. The presence of a passivation layer on the surface of the film can inhibit the diffusional relaxation and causes the inelastic deformation to be dominated by dislocation processes within the grains [10]. Studies of stress relaxation processes have shown [11] that for Al films the presence of a native oxide inhibits diffusional flow between the grain boundaries in the film and the surface and sustains biaxial stress of several hundred MPa at temperatures as high as 400 C. For metals such as Cu, Ag, Pt and Au, which do not form dense oxide barrier layers, the biaxial stresses relax almost completely above the temperatures at which grain boundary diffusion can occur. The role of the surface oxide layer on hydriding– dehydriding properties has been studied in a recent publication [12]. Many attempts have been undertaken to understand the dominant forces driving hydrogen from the surface into the bulk by modifying the surface of magnesium, reducing the particle size, controlling the surface oxidation or adding various elements. However, no general mechanism for the hydrogenation reaction has been proposed. This is because the hydrogenation kinetics depends strongly on many parameters characterizing the surface and bulk material properties, such as particle size, microstructure and impurities content, material fabrication technology and particle morphology. In the present paper we concentrate our studies on the effects of the dynamic structural rearrangements induced under hydrogenation and on the hydrogenation efficiency.
2.
Experimental technique
Films made of 1.5-mm thick Mg capped with 100-nm thick Ti and Al layers or MgxNi films were deposited on quartz glass substrates in two vacuum systems. The surface roughness of substrates was investigated using AFM and did not exceed 10 nm. The C-free MgNi films were deposited in a vacuum device where the ultimate pressure obtained using cryo pumps was around 106 Pa. Another vacuum device pumped by rotary and diffusion oil pumps, with a lower ultimate pressure of 2 103 Pa, was used to produce the C-contaminated MgNi films. Both vacuum systems were equipped with DC magnetrons. Metallic Mg (99.98%), Ni (99.995%), Ti (99.995%) and Al (99.95%) disks of 80-mm diameter each produced by Kurt J. Lesker Company were used as target materials for the magnetron sources. Ar (99.999%) delivered by
AGA Gas Gmbh was injected into the vacuum chamber and its flow was controlled to keep a constant pressure of 0.4 Pa while magnetron targets were sputter-cleaned for 15 min, the quartz substrate being protected by a shutter. The argon pressure was then reduced to 8 102 Pa and the shutter removed before proceeding to the Mg or MgNi films deposition. Without vacuum interruption, the deposited Mg films were subsequently capped with Ti and Al layers. The thickness of the capping layer was selected to ensure continuous film structure without holes and cracks. Mg/Ni atom ratio in MgNi films was chosen slightly above 2 expecting the formation of the Mg2Ni alloy, in order to avoid the formation of metallic MgNi2 that does not absorb hydrogen. Despite the substrate not being heated during deposition, the temperature increased up to 80–120 C, depending on deposition parameters. The film deposition rate was determined from the slope of the sample weight change, measured using a microbalance with a weight uncertainty of 2 mg which corresponds to an uncertainty of 10 nm on the thickness. The deposition rates of Mg, Ti and Al were 0.6, 0.45 and 0.3 nm s1, respectively After deposition, the films were exposed to air and moved for hydrogenation into a stainless steel cell which was thoroughly flushed with pure H2 (99.999%). Steady-state temperature 250 C and the hydrogen pressure 600 kPa were reached in 7 min. Heating was turned off after hydrogenation, Ar gas was injected and samples were cooled down over 16 h. The microstructure of the films was analyzed using X-ray diffraction (XRD) with the 2Q angle in the range 20–90 using CuKa radiation in steps of 0.05 . The average crystallite dimension of films, D, was calculated using the formula D ¼ 0:9lðb cosQÞ1 neglecting the microstrain, where l is the X-ray wavelength, Q is the Bragg diffraction angle and b is the full-width of the peak after correcting for the instrument broadening. The identification of phases has been performed using Crystallographica Search-Match program (Oxford Cryosystem Ltd, 1996–2003) based on Powder Diffraction Data [13]. Scanning electron microscopy (SEM, JEOL JSM-6300) in parallel with atomic force microscopy (AFM) was used to investigate the surfaces of both as-deposited and hydrogenated films. After hydrogenation, Elastic Recoil Detection Analysis (ERDA) with 35 MeV Cl7þ ions and Nuclear Reaction Analysis (NRA) with 15N2þ ions were performed on the films at the Dresden–Rossendorf Tandem accelerator. The ERDA measurements have been used to determine the concentrations and depth profiles of all elements, especially the light elements like O [14]. Hydrogen was also registered with ERDA using a solid state detector preceded by an Al foil to stop other ion species but with a poor depth resolution. The analyzed area was about 2 1 mm2. Energy dispersive X-ray spectroscopy (EDS) was also used to have access to the elemental composition of the matrix elements.
3.
Experimental results
3.1.
Hydrogenation of Mg films
Fig. 1 includes XRD patterns of sputter-deposited Mg films after different stages of treatment.
international journal of hydrogen energy 34 (2009) 5131–5137
Fig. 1 – XRD patterns of Mg films after different stages of treatment: a – as-deposited; b – after 3 h of hydrogenation; c – Al-capped Mg film after 6 h of hydrogenation; d – Ticapped Mg film after 0.5 h of hydrogenation and e – with Ti layers on both sides and after 3 h of hydrogenation.
The XRD pattern of as-deposited Mg film (Fig. 1a) demonstrates the nanocrystalline microstructure with broad characteristic peaks for h-Mg (002) at 34.5 and h-Mg (101) at 36.9 . The evaluated size of crystallites is in the range 75–90 nm. It is seen that after hydrogenation for 3 h at 250 C, the Mg is partially converted into MgH2 with characteristic peaks at 35.9 and 40.2 (Fig. 1b). It is observed that the synthesis of the hydride phase takes place during the initial stage of hydrogenation. The amplitude of the main peak at 35.9 is obtained after a few minutes of hydrogenation and does not change over time. The characteristic peak at 40.2 appears after 0.5 h of hydrogenation and its intensity slowly increases over time. The evolution of the Mg diffraction peaks demonstrates that restructuring processes occur upon hydrogenation. The broadening of the diffraction peaks as hydrogenation proceeds is related to the introduction of a high concentration of defects in the dominant material and is due to the decrease of the average crystallite size (in the range 45–60 nm) and the increase of the lattice strain. Ti and Al-capped Mg films exhibit different behavior under hydrogenation. The MgH2 phase is not registered after hydrogenation for 6 h at 250 C in Al-capped Mg film (Fig. 1c). However, the Ti-capped Mg film is completely converted into MgH2 after hydrogenation for 0.5 h at the same temperature (Fig. 1d). Mg film with thin Ti layers on both sides is only partially transformed into MgH2 after hydrogenation over 3 h at 250 C (Fig. 1e). It indicates that hydride synthesis kinetics depends not only on the properties of the external surface contacting with hydrogen. Structural and chemical properties of the film and the substrate interface region may significantly modify hydrogenation kinetics. Complementary information about structural transformations under hydrogenation has been obtained from the changes of surface topography in SEM as these changes can be related to the processes acting in the bulk.
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The surface view of the as-deposited Mg film suggests (Fig. 2a) that it has a columnar structure and exhibits many well-defined boundaries which are open and may serve as channels for the transportation of hydrogen from the surface into the bulk. After hydrogenation, the surface becomes enclosed and does not show open channels (Fig. 2b). The cross-sectional SEM view (Fig. 2b, insert) reveals the presence of a thin dense layer on the surface of hydrogenated Mg film. It may be attributed to the MgH2 grown during the initial stages of hydrogenation which blocks the ensuing hydrogen transport. It explains the fast appearance of small amount of MgH2 phase at the beginning of hydrogenation and the subsequent slow hydrogenation rate. A thin Al layer deposited on the top of the Mg film blocks hydrogen uptake. After exposure to air, this layer is covered by a 2–5-nm thick film of the natural Al2O3 which is nonpermeable to H. Thus, the hydride phase is not registered even after 6 h of hydrogenation at 250 C (Fig. 1c). SEM also reveals the presence of blister-like protrusions (Fig. 2c) which could result from hydrogen diffusing through the surface barrier layer and nucleating at the Al2O3–Al interface. Hydrogen dissolved in Al precipitates and forms bubble. Mg film hydrogenation properties change significantly if the film is capped with a Ti layer. It is completely converted into MgH2 after 0.5 h of hydrogenation at 250 C. In many cases, the hydrogenated film is lifted from the substrate after cooling (Fig. 2d). It suggests that Ti on the surface of Mg catalyzes the dissociative adsorption of molecular hydrogen, and dynamic restructuring processes initiated by surface chemistry prevent formation of a barrier layer on the surface and drive hydrogen atoms from the surface to the film– substrate interface boundary. When thin Ti layers are on both sides of the Mg film, it also modifies the kinetics of restructuring processes. The difference of chemical potentials between the free surface and the interface decreases so that the flow of hydrogen from the surface is less in the Ti/Mg/Ti than in the Mg/Ti structure.
3.2.
Hydrogenation of MgNi films
The as-deposited C-free MgNi film demonstrates X-ray amorphous structure (Fig. 3a). This is in agreement with results previously obtained under similar conditions [15]. Upon hydrogenation for 3 h at 250 C, the dominant phase becomes Mg2Ni with a small amount of Mg2NiH4 (Fig. 3b). The presence of Mg2Ni is mainly indicated through its h-(003) at 19.9 and h-(006) at 40.5 reflections. This is an indication that the Mg2Ni has a textured structure with the c-axis of the hexagonal lattice in the direction of the layer normal with a ¼ 0.5216 nm and c ¼ 1.343 nm. The small amount of Mg2NiH4 with characteristic peaks at 23.9 and 39.5 nucleates at the initial stages of hydrogenation. It is consistent with a preferential growth of Mg2NiH4 which takes place at the film–substrate interface [16]. It is possible that the enthalpy is more negative and more favorable for synthesis of Mg2NiH4 in the region closer to the substrate than in the rest of the film and the interface residual stresses direct hydrogen flow to that region. The ensuing hydrogenation process proceeds by slow growth of the same phase, but at the grain boundaries [17].
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international journal of hydrogen energy 34 (2009) 5131–5137
Fig. 2 – SEM views of Mg films: a – as-deposited; b – after 3 h of hydrogenation; c –Al-capped and after 6 h of hydrogenation and d – Ti-capped and after 0.5 h of hydrogenation.
C-contaminated MgNi films exhibit different behavior upon hydrogenation. Mg2Ni is the dominant phase in the as-deposited film (Fig. 3c) and is entirely transformed into Mg2NiH4 after hydrogenation at 250 C for 0.5 h (Fig. 3d). The changes of film colour have been registered when hydrogenation is performed in a transparent quartz tube. After 2 min of hydrogenation, as sample temperature reaches 353 K, the shiny metallic film colour changes into dark brown. After 5 min of hydrogenation, as the sample temperature reaches 433 K, the film changes colour into orange and becomes transparent indicating that stoichiometric Mg2NiH4 compound is formed. Upon further
Fig. 3 – XRD patterns of MgNi films: C-free as deposited (a) and after hydrogenation during 3 h (b), C-contaminated as deposited (c) and after hydrogenation during 0.5 h (d).
hydrogenation up to 250 C the film is transparent and changes its colour into reddish. Film colour changes into orange after cooling. The SEM surface views of hydrogenated MgNi films are included in Fig. 4a and b. The XRD diffraction patterns of corresponding samples are shown in Fig. 3b and d, correspondingly. The C-free hydrogenated MgNi film has a smooth surface with the dark patches indicating areas lifted at the film–substrate interface (Fig. 4a). The SEM cross-sectional view (Fig. 4a, insert) shows that the microstructure of the interface region differs from the bulk and the film is separated from the substrate. It supports the above-made suggestion that the preferential nucleation of Mg2NiH4 takes place at the film–substrate interface and induced stresses that lift the film from the substrate. The C-contaminated hydrogenated MgNi films are completely transformed into Mg2NiH4 (Fig. 3d) and stressdriven relaxation processes crush the thin film during cooling as shown in Fig. 4b. Fig. 5 includes the distribution profiles of C, O and H atoms in the near-surface region of C-free MgNi film hydrogenated during 6 h at 250 C measured using ERDA. The corresponding NRA profiles across all film thickness have already been published [18]. It is seen that the nearsurface region (w100 nm) of hydrogenated film is highly contaminated by oxygen (about 40 at%) and a small amount (about 2 at%) of C. The main source of oxygen and carbon impurities on the surface may be desorbed water and hydrocarbons from the walls of the experimental device during hydrogenation. Beyond this contaminated near-surface region, the hydrogen appears homogeneously (about 40 at%) distributed in the bulk.
international journal of hydrogen energy 34 (2009) 5131–5137
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Fig. 4 – SEM views of hydrogenated MgNi films: a – C-free after 3 h of hydrogenation and b – C-contaminated after 0.5 h of hydrogenation.
Energy dispersive X-ray spectroscopy (EDS) performed on as-deposited C-contaminated MgNi films gives a carbon concentration about 6–8 at% and an oxygen concentration less than 10 at%. After 0.5 h of hydrogenation at 250 C, this film is converted into stoichiometric Mg2NiH4 compound including 3–5 at% of C and still about 10 at% of O. Experimental results give the indication that C impurities in MgNi films change their hydrogenation properties. It seems that C impurities in MgNi film prevent formation of the barrier oxide layer and support dynamic structural rearrangements. A tentative explanation of the observations may be as-follows: incident oxygen atoms during interaction with C-contaminated surface react with C atoms and form CO2 molecules which are desorbed. For C-free MgNi films, the oxygen arriving at the surface forms oxides and builds a barrier layer hindering ensuing hydrogenation. The surface barrier layer blocks relaxation processes and inhibits dynamic behavior of grains.
4.
Discussion
Hydrogen molecules dissociate on metal surfaces, dissolve atomically into the substrate, and interact with the atoms in the bulk. Adsorption–desorption processes intimately control the degree to which surfaces are contaminated. The H-uptake
Fig. 5 – Distribution profiles of H, C and O atoms in the near-surface region of the C-free MgNi film after 3 h of hydrogenation.
mechanism depends not only on the surface atom composition and hydrogen concentration in the bulk but also on the prevalent microstructure; i.e., the mixture of crystal defects and impurities, grain boundaries, grain size, grain orientation, and phase composition present within the metal matrix. Under hydrogenation, the synthesis of new phases activates the internal structure and initiates relaxation processes. In this way, the H-uptake rate may be modified as the film properties change. A fast and complete transformation of metal into the hydride phase is observed for Ti-capped Mg and for Ccontaminated MgNi films when the dynamic state of grains and stress-driven relaxation processes are not suppressed during hydrogenation. If newly nucleated structural formations inhibit the dynamic state, the hydrogenation rate sharply decreases and only partial transformation of metallic into hydride phase is registered. The purposefully formed Al2O3 barrier on the surface, which suppresses the stress relaxation processes [18], sharply diminishes the hydrogenation rate. For C-free MgNi films, a fast H-uptake takes place during the initial stages of hydrogenation. The nucleation of a thin Mg2NiH4 layer near the film–substrate interface region increases the chemical potential of the region and the H-flow from the surface into the bulk decreases. As the dynamic structural rearrangements subside, a dense surface oxide barrier layer is formed. It blocks the H-uptake into the bulk and hinders stress-driven relaxation processes. The system becomes kinetically locked with low hydrogenation rate. The C-contaminated MgNi films are fully hydrogenated after only 10–15 min at 250 C. It seems that oxygen adsorption and carbon desorption processes intimately control the degree to which surfaces are contaminated. Carbon impurities prevent formation of the surface oxide barrier layer as arriving oxygen reacts with carbon, forming volatile CO2. The fast H-uptake may be explained in the following way. It is well-established that an adsorption of reactive atoms renders metallic surfaces dynamic. The surface chemistry is directed to optimize bonding between reacting species and it results in surface diffusion of adatoms over distances of many atomic bonds. Continuous reactive adsorption makes the surface structure continually adjust to changes in the chemical environment, which leads to local and long-range
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international journal of hydrogen energy 34 (2009) 5131–5137
restructuring, depending on the properties of the near-surface region. In this way, the surface layer becomes both chemically and physically distinct from the underlying bulk material. Thus, a difference in chemical potentials is established between chemically activated surface, grain boundaries and film–substrate interface region. The excess of the surface chemical potentials relative to the grain boundaries and the interface region produce a net flux of H atoms into the grain boundaries and to the interface region. It therefore generates a compressive stress in the grains and across film thickness. If the stress exceeds the limit of plasticity, stress relaxation occurs through the motion of dislocations. If dislocation motion is prevented, which occurs at the grain boundaries, formation of subgrains within the original grain structure takes place. The development of a cell structure occurs by entanglement of dislocations on various slip systems. Subgrains develop by (dynamic) recovery. In this way the uptake of hydrogen atoms through the grain boundaries in open contact with the highly activated surface occurs. Thus, the high surface chemical potentials are not only imposed on the external surface but also on the internal grain interfaces. The diffusivity along boundaries is significantly higher (by more than three orders of magnitude) than the lattice diffusivity of reactive species [19]. An estimate of internal stresses gives that accommodated biaxial stresses are on the order of few Gpa [20] which exceeds yield stress. Thus, compressive stress provides a strong driving force for the plasticity in the grains. The plastic deformation of grains supports the dynamic state of grains and their fragmentation and inhibits formation of continuous oxide layer. Consequently, a fast inward diffusion of hydrogen through the grain boundaries takes place.
5.
Conclusions
Hydrogenation of MgNi film is fast and complete if their surface layer is permeable for H atoms. The dynamic state of this surface is connected with structural transformations in the bulk under hydrogenation. Hydrogen is thus driven from the surface region through the grain boundaries of nanocrystallites to the film–substrate interface where the chemical potential of the local microstructure is diminished. Barrier layers on the surface inhibit transport of H atoms from the surface into the bulk, sustain relaxation processes and suppress activated state. A partial transformation of metallic into hydride phase is registered when hydrides are nucleated at the initial stages of hydrogenation, subsequently inhibiting the dynamic processes in the bulk. Small quantities of impurities may change behavior of hydrogen on the surface and in the bulk and can induce large differences in the hydrogen uptake mechanism and the amount of hydride formed.
Acknowledgments This work has been financially supported by FP6 project (No. 038941; HySIC) and by the European Research Training
Network (Contract No. MRTN-CT-2004-512443 HyTRAIN). Partial fundings by the European Commission DG Research (SES6-2006-518271/NESSHY) and the Lithuanian Science Foundation (H2technologijos, 2008–2010) are also gratefully acknowledged.
references
[1] Yamamura S, Kasahara S, Takata M, Sugawara Y, Sakaka MJ. Imaging of the electron density distributions of hydrogen in LiH and LiOH by maximum entropy method. J Phys Chem Solids 1999;60(10):1721–4. [2] Pozzo M, Alfe` D. Hydrogen dissociation and diffusion on transition metal (¼Ti, Zr, V, Fe, Ru, Co, Rh, Ni, Pd, Cu, Ag)doped Mg(0001) surfaces. Int J Hydrogen Energy 2009;34(4): 1922–30. [3] Cermak J, Kral L. Hydrogenation of Mg and two chosen Mg–Ni alloys. Int J Hydrogen Energy 2008;33(24):7464–70. [4] Orimo S, Zu¨ttel A, Ikeda K, Saruki S, Fukunaga T, Fujii H, et al. Hydriding properties of the MgNi-based systems. J Alloys Compd 1999;293–295:437–42. [5] Qu J, Wang Y, Xie L, Zheng J, Liu Y, Li X. Hydrogen absorption–desorption, optical transmission properties and annealing effect of Mg thin films prepared by magnetron sputtering. Int J Hydrogen Energy 2009;34(4): 1910–5. [6] Patah A, Takasaki A, Szmyd JS. Influence of multiple oxide (Cr2O3/Nb2O5) addition on the sorption kinetics of MgH2. Int J Hydrogen Energy 2009;34(7):3032–7. [7] Gremaud R, Borgschulte A, Chacon C, van Mechelen JLM, Schreuders H, Zuttel A. Structural and optical properties of MgxAl1xHy gradient thin films: a combinatorial approach. Appl Phys A 2006;84(1–2):77–85. [8] Vinci RP, Zielinski EM, Bravman JC. Thermal strain and stress in copper thin films. Thin Solid Films 1995;262(1–2): 142–53. [9] Keller R-M, Baker SP, Arzt E. Quantitative analysis of strengthening mechanisms in thin Cu films: effects of film thickness, grain size, and passivation. J Mater Res 1998;13(5): 1307–17. [10] Thouless MD, Rodbell KP, Cabral CI. Effect of a surface layer on the stress relaxation of thin films. J Vac Sci Technol A 1996;14(4):2454–61. [11] Gao H, Zhang L, Nix WD, Thompson CV, Arzt E. Crack-like grain-boundary diffusion wedges in thin metal films. Acta Materialia 1999;47(10):2865–78. [12] Pranevicius L, Wirth E, Milcius D, Lelis M, Pranevicius LL, Bacianskas A. Structure transformations and hydrogen storage properties of co-sputtered MgNi films. Appl Surf Sci 2009;255(11):5971–4. [13] JCPDS data cards. Swarthmore, PA: International Center of Diffraction Data; 1988. [14] Grigull S, Kreissig U, Huber H, Assmann W. Elementdependent ERDA probing depths using different detection systems. Nucl Instr Meth In Physics Research Section B 1997; 132(4):709–17. [15] Borsa DM, Lohstroh W, Gremaud R, Rector JH, Dam B, WijnGaarden RJ, et al. Critical composition dependence of the hydrogenation of Mg2d Ni thin films. J Alloys Compd 2007;428(1–2):34–9. [16] Ouyang LZ, Wang H, Chung CY, Ahn JH, Zhu M. MgNi/Pd multilayer hydrogen storage thin films prepared by dc magnetron sputtering. J Alloys Compd 2006;422(1–2): 58–61.
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[17] Wirth E, Munnik F, Pranevicius LL, Milcius D. Dynamic surface barrier effects on hydrogen storage capacity in Mg–Ni films. J Alloys Compd 2009;475(1–2):917–22. [18] Pranevicius L, Wirth E, Milcius D, Pranevicius LL, Kanapickas A. Effects of surface dynamic behavior on hydrogen storage properties of sputter-deposited MgNi films. Surf Coat Technol 2009;203(8):998–1003.
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[19] Fielitz P, Borchardt G, Schmucker M, Schneider H, Willich P. Measurement of oxygen grain boundary diffusion in mullite ceramics by SIMS depth profiling. Appl Surf Sci 2003;203–204: 639–43. [20] Laudahn U, Pundt A, Bicker M, Hulsen U, Geyer U, Wagner T, et al. Hydrogen-induced stress in Nb single layers. J Alloys Compd 1999;293–295:490–4.
Available at www.sciencedirect.com
journal homepage: www.elsevier.com/locate/he
The effects of dynamic structural transformations on hydrogenation properties of Mg and MgNi thin films L. Praneviciusa,*, D. Milciusb, C. Templierc a
Vytautas Magnus University, 8 Vileikos St., LT-44404 Kaunas, Lithuania Lithuanian Energy Institute, 3 Breslaujos St., LT-44403 Kaunas, Lithuania c Universite´ de Poitiers, SP2MI, Te´le´port 2, Bd Marie et Pierre Curie BP 30179, Futuroscope, France b
article info
abstract
Article history:
Mg capped with Al and Ti thin layers and MgxNi films have been sputter-deposited on
Received 19 March 2009
quartz substrates and hydrogenated at 600 kPa for 250 C. A complete fast transformation
Received in revised form
of metallic into hydride phase is registered for the films, demonstrating the dynamic state
17 April 2009
of the internal microstructure under hydrogenation. It leads to local and long-range
Accepted 18 April 2009
restructuring and the fast hydrogenation rate is attributed to the fast hydrogen uptake and
Available online 14 May 2009
transport along columns and grain boundaries of nanocrystallites. A slow H-loading is observed when the dynamic structural relaxation processes are suppressed by internal and
Keywords:
external inhomogeneities such as barrier layers on the surface, new phases in the bulk and
Hydrogenation
impurities. A partial transformation of metallic into hydride phase is registered when the
Thin films
structural formations newly nucleated at the initial stages of hydrogenation suppress
Mg and MgNi
dynamic processes and prevent the H-uptake. ª 2009 International Association for Hydrogen Energy. Published by Elsevier Ltd. All rights reserved.
1.
Introduction
Mg-based materials have attracted extensive attention for application in hydrogen storage technology due to the high storage capacity of the hydride phase. Nearly 7.6 wt% of hydrogen can be stored as MgH2 in Mg and nearly 3.6 wt% hydrogen reacts with Mg2Ni to form Mg2NiH4 in a reversible reaction [1,2]. Unfortunately, the hydride phase of Mg-based materials is very stable and shows poor hydrogen absorption and desorption kinetics even at high temperature. In the last few years, substantial improvement in hydrogen absorption and desorption kinetics has been achieved by several different methods [3–5]. The most effective methods realized so far are based on structural refinements of the material and the addition of proper catalysts. Nanocrystalline and amorphous magnesium-based alloys demonstrate promising
performance as hydrogen storage materials. Nanoscale particles are advantageous over bulk materials as they have a larger solid/gas interface area and shorter hydrogen diffusion paths, yielding potentially faster kinetics for gas absorption and desorption. A proper way to make suitable nanocomposite for hydrogen storage is to produce thin films by sputtering or evaporation methods as it allows controlling the composition and structure of the film layers at a nanometer scale. Thus, thin films have been widely used to study hydrogen behavior properties of the corresponding bulk materials. However, the fundamental difference with respect to bulk samples is the presence of a substrate which might alter kinetics and thermodynamics of the processes. Hydrogen in metals expands the metal host lattice. A thin film cannot expand freely in all dimensions as adhesion forces at the film–substrate interface
* Corresponding author. Tel.: þ370 37 327909; fax: þ370 37 327916. E-mail address: [email protected] (L. Pranevicius). 0360-3199/$ – see front matter ª 2009 International Association for Hydrogen Energy. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.ijhydene.2009.04.034
5132
international journal of hydrogen energy 34 (2009) 5131–5137
prevent the expansion of the film. It leads to considerable inplane stresses of the order of several GPa [6,7]. Nucleation rates usually involve rapidly varying free energies which are extraordinary sensitive to the conditions under which the experiment is run (pressure, impurity concentrations, as well as temperature). Reproducibility of experimental data is difficult to achieve and the theoretical prediction of nucleation rate is subject to large uncertainties. The effects of stress generation and relaxation on the hydride phase nucleation mechanism and hydrogen storage properties are not fully understood. It has been shown [8,9] that stress relaxation can occur by diffusional deformation wherein mass is transported between the free surface of the film and the grain boundaries. The presence of a passivation layer on the surface of the film can inhibit the diffusional relaxation and causes the inelastic deformation to be dominated by dislocation processes within the grains [10]. Studies of stress relaxation processes have shown [11] that for Al films the presence of a native oxide inhibits diffusional flow between the grain boundaries in the film and the surface and sustains biaxial stress of several hundred MPa at temperatures as high as 400 C. For metals such as Cu, Ag, Pt and Au, which do not form dense oxide barrier layers, the biaxial stresses relax almost completely above the temperatures at which grain boundary diffusion can occur. The role of the surface oxide layer on hydriding– dehydriding properties has been studied in a recent publication [12]. Many attempts have been undertaken to understand the dominant forces driving hydrogen from the surface into the bulk by modifying the surface of magnesium, reducing the particle size, controlling the surface oxidation or adding various elements. However, no general mechanism for the hydrogenation reaction has been proposed. This is because the hydrogenation kinetics depends strongly on many parameters characterizing the surface and bulk material properties, such as particle size, microstructure and impurities content, material fabrication technology and particle morphology. In the present paper we concentrate our studies on the effects of the dynamic structural rearrangements induced under hydrogenation and on the hydrogenation efficiency.
2.
Experimental technique
Films made of 1.5-mm thick Mg capped with 100-nm thick Ti and Al layers or MgxNi films were deposited on quartz glass substrates in two vacuum systems. The surface roughness of substrates was investigated using AFM and did not exceed 10 nm. The C-free MgNi films were deposited in a vacuum device where the ultimate pressure obtained using cryo pumps was around 106 Pa. Another vacuum device pumped by rotary and diffusion oil pumps, with a lower ultimate pressure of 2 103 Pa, was used to produce the C-contaminated MgNi films. Both vacuum systems were equipped with DC magnetrons. Metallic Mg (99.98%), Ni (99.995%), Ti (99.995%) and Al (99.95%) disks of 80-mm diameter each produced by Kurt J. Lesker Company were used as target materials for the magnetron sources. Ar (99.999%) delivered by
AGA Gas Gmbh was injected into the vacuum chamber and its flow was controlled to keep a constant pressure of 0.4 Pa while magnetron targets were sputter-cleaned for 15 min, the quartz substrate being protected by a shutter. The argon pressure was then reduced to 8 102 Pa and the shutter removed before proceeding to the Mg or MgNi films deposition. Without vacuum interruption, the deposited Mg films were subsequently capped with Ti and Al layers. The thickness of the capping layer was selected to ensure continuous film structure without holes and cracks. Mg/Ni atom ratio in MgNi films was chosen slightly above 2 expecting the formation of the Mg2Ni alloy, in order to avoid the formation of metallic MgNi2 that does not absorb hydrogen. Despite the substrate not being heated during deposition, the temperature increased up to 80–120 C, depending on deposition parameters. The film deposition rate was determined from the slope of the sample weight change, measured using a microbalance with a weight uncertainty of 2 mg which corresponds to an uncertainty of 10 nm on the thickness. The deposition rates of Mg, Ti and Al were 0.6, 0.45 and 0.3 nm s1, respectively After deposition, the films were exposed to air and moved for hydrogenation into a stainless steel cell which was thoroughly flushed with pure H2 (99.999%). Steady-state temperature 250 C and the hydrogen pressure 600 kPa were reached in 7 min. Heating was turned off after hydrogenation, Ar gas was injected and samples were cooled down over 16 h. The microstructure of the films was analyzed using X-ray diffraction (XRD) with the 2Q angle in the range 20–90 using CuKa radiation in steps of 0.05 . The average crystallite dimension of films, D, was calculated using the formula D ¼ 0:9lðb cosQÞ1 neglecting the microstrain, where l is the X-ray wavelength, Q is the Bragg diffraction angle and b is the full-width of the peak after correcting for the instrument broadening. The identification of phases has been performed using Crystallographica Search-Match program (Oxford Cryosystem Ltd, 1996–2003) based on Powder Diffraction Data [13]. Scanning electron microscopy (SEM, JEOL JSM-6300) in parallel with atomic force microscopy (AFM) was used to investigate the surfaces of both as-deposited and hydrogenated films. After hydrogenation, Elastic Recoil Detection Analysis (ERDA) with 35 MeV Cl7þ ions and Nuclear Reaction Analysis (NRA) with 15N2þ ions were performed on the films at the Dresden–Rossendorf Tandem accelerator. The ERDA measurements have been used to determine the concentrations and depth profiles of all elements, especially the light elements like O [14]. Hydrogen was also registered with ERDA using a solid state detector preceded by an Al foil to stop other ion species but with a poor depth resolution. The analyzed area was about 2 1 mm2. Energy dispersive X-ray spectroscopy (EDS) was also used to have access to the elemental composition of the matrix elements.
3.
Experimental results
3.1.
Hydrogenation of Mg films
Fig. 1 includes XRD patterns of sputter-deposited Mg films after different stages of treatment.
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Fig. 1 – XRD patterns of Mg films after different stages of treatment: a – as-deposited; b – after 3 h of hydrogenation; c – Al-capped Mg film after 6 h of hydrogenation; d – Ticapped Mg film after 0.5 h of hydrogenation and e – with Ti layers on both sides and after 3 h of hydrogenation.
The XRD pattern of as-deposited Mg film (Fig. 1a) demonstrates the nanocrystalline microstructure with broad characteristic peaks for h-Mg (002) at 34.5 and h-Mg (101) at 36.9 . The evaluated size of crystallites is in the range 75–90 nm. It is seen that after hydrogenation for 3 h at 250 C, the Mg is partially converted into MgH2 with characteristic peaks at 35.9 and 40.2 (Fig. 1b). It is observed that the synthesis of the hydride phase takes place during the initial stage of hydrogenation. The amplitude of the main peak at 35.9 is obtained after a few minutes of hydrogenation and does not change over time. The characteristic peak at 40.2 appears after 0.5 h of hydrogenation and its intensity slowly increases over time. The evolution of the Mg diffraction peaks demonstrates that restructuring processes occur upon hydrogenation. The broadening of the diffraction peaks as hydrogenation proceeds is related to the introduction of a high concentration of defects in the dominant material and is due to the decrease of the average crystallite size (in the range 45–60 nm) and the increase of the lattice strain. Ti and Al-capped Mg films exhibit different behavior under hydrogenation. The MgH2 phase is not registered after hydrogenation for 6 h at 250 C in Al-capped Mg film (Fig. 1c). However, the Ti-capped Mg film is completely converted into MgH2 after hydrogenation for 0.5 h at the same temperature (Fig. 1d). Mg film with thin Ti layers on both sides is only partially transformed into MgH2 after hydrogenation over 3 h at 250 C (Fig. 1e). It indicates that hydride synthesis kinetics depends not only on the properties of the external surface contacting with hydrogen. Structural and chemical properties of the film and the substrate interface region may significantly modify hydrogenation kinetics. Complementary information about structural transformations under hydrogenation has been obtained from the changes of surface topography in SEM as these changes can be related to the processes acting in the bulk.
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The surface view of the as-deposited Mg film suggests (Fig. 2a) that it has a columnar structure and exhibits many well-defined boundaries which are open and may serve as channels for the transportation of hydrogen from the surface into the bulk. After hydrogenation, the surface becomes enclosed and does not show open channels (Fig. 2b). The cross-sectional SEM view (Fig. 2b, insert) reveals the presence of a thin dense layer on the surface of hydrogenated Mg film. It may be attributed to the MgH2 grown during the initial stages of hydrogenation which blocks the ensuing hydrogen transport. It explains the fast appearance of small amount of MgH2 phase at the beginning of hydrogenation and the subsequent slow hydrogenation rate. A thin Al layer deposited on the top of the Mg film blocks hydrogen uptake. After exposure to air, this layer is covered by a 2–5-nm thick film of the natural Al2O3 which is nonpermeable to H. Thus, the hydride phase is not registered even after 6 h of hydrogenation at 250 C (Fig. 1c). SEM also reveals the presence of blister-like protrusions (Fig. 2c) which could result from hydrogen diffusing through the surface barrier layer and nucleating at the Al2O3–Al interface. Hydrogen dissolved in Al precipitates and forms bubble. Mg film hydrogenation properties change significantly if the film is capped with a Ti layer. It is completely converted into MgH2 after 0.5 h of hydrogenation at 250 C. In many cases, the hydrogenated film is lifted from the substrate after cooling (Fig. 2d). It suggests that Ti on the surface of Mg catalyzes the dissociative adsorption of molecular hydrogen, and dynamic restructuring processes initiated by surface chemistry prevent formation of a barrier layer on the surface and drive hydrogen atoms from the surface to the film– substrate interface boundary. When thin Ti layers are on both sides of the Mg film, it also modifies the kinetics of restructuring processes. The difference of chemical potentials between the free surface and the interface decreases so that the flow of hydrogen from the surface is less in the Ti/Mg/Ti than in the Mg/Ti structure.
3.2.
Hydrogenation of MgNi films
The as-deposited C-free MgNi film demonstrates X-ray amorphous structure (Fig. 3a). This is in agreement with results previously obtained under similar conditions [15]. Upon hydrogenation for 3 h at 250 C, the dominant phase becomes Mg2Ni with a small amount of Mg2NiH4 (Fig. 3b). The presence of Mg2Ni is mainly indicated through its h-(003) at 19.9 and h-(006) at 40.5 reflections. This is an indication that the Mg2Ni has a textured structure with the c-axis of the hexagonal lattice in the direction of the layer normal with a ¼ 0.5216 nm and c ¼ 1.343 nm. The small amount of Mg2NiH4 with characteristic peaks at 23.9 and 39.5 nucleates at the initial stages of hydrogenation. It is consistent with a preferential growth of Mg2NiH4 which takes place at the film–substrate interface [16]. It is possible that the enthalpy is more negative and more favorable for synthesis of Mg2NiH4 in the region closer to the substrate than in the rest of the film and the interface residual stresses direct hydrogen flow to that region. The ensuing hydrogenation process proceeds by slow growth of the same phase, but at the grain boundaries [17].
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Fig. 2 – SEM views of Mg films: a – as-deposited; b – after 3 h of hydrogenation; c –Al-capped and after 6 h of hydrogenation and d – Ti-capped and after 0.5 h of hydrogenation.
C-contaminated MgNi films exhibit different behavior upon hydrogenation. Mg2Ni is the dominant phase in the as-deposited film (Fig. 3c) and is entirely transformed into Mg2NiH4 after hydrogenation at 250 C for 0.5 h (Fig. 3d). The changes of film colour have been registered when hydrogenation is performed in a transparent quartz tube. After 2 min of hydrogenation, as sample temperature reaches 353 K, the shiny metallic film colour changes into dark brown. After 5 min of hydrogenation, as the sample temperature reaches 433 K, the film changes colour into orange and becomes transparent indicating that stoichiometric Mg2NiH4 compound is formed. Upon further
Fig. 3 – XRD patterns of MgNi films: C-free as deposited (a) and after hydrogenation during 3 h (b), C-contaminated as deposited (c) and after hydrogenation during 0.5 h (d).
hydrogenation up to 250 C the film is transparent and changes its colour into reddish. Film colour changes into orange after cooling. The SEM surface views of hydrogenated MgNi films are included in Fig. 4a and b. The XRD diffraction patterns of corresponding samples are shown in Fig. 3b and d, correspondingly. The C-free hydrogenated MgNi film has a smooth surface with the dark patches indicating areas lifted at the film–substrate interface (Fig. 4a). The SEM cross-sectional view (Fig. 4a, insert) shows that the microstructure of the interface region differs from the bulk and the film is separated from the substrate. It supports the above-made suggestion that the preferential nucleation of Mg2NiH4 takes place at the film–substrate interface and induced stresses that lift the film from the substrate. The C-contaminated hydrogenated MgNi films are completely transformed into Mg2NiH4 (Fig. 3d) and stressdriven relaxation processes crush the thin film during cooling as shown in Fig. 4b. Fig. 5 includes the distribution profiles of C, O and H atoms in the near-surface region of C-free MgNi film hydrogenated during 6 h at 250 C measured using ERDA. The corresponding NRA profiles across all film thickness have already been published [18]. It is seen that the nearsurface region (w100 nm) of hydrogenated film is highly contaminated by oxygen (about 40 at%) and a small amount (about 2 at%) of C. The main source of oxygen and carbon impurities on the surface may be desorbed water and hydrocarbons from the walls of the experimental device during hydrogenation. Beyond this contaminated near-surface region, the hydrogen appears homogeneously (about 40 at%) distributed in the bulk.
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Fig. 4 – SEM views of hydrogenated MgNi films: a – C-free after 3 h of hydrogenation and b – C-contaminated after 0.5 h of hydrogenation.
Energy dispersive X-ray spectroscopy (EDS) performed on as-deposited C-contaminated MgNi films gives a carbon concentration about 6–8 at% and an oxygen concentration less than 10 at%. After 0.5 h of hydrogenation at 250 C, this film is converted into stoichiometric Mg2NiH4 compound including 3–5 at% of C and still about 10 at% of O. Experimental results give the indication that C impurities in MgNi films change their hydrogenation properties. It seems that C impurities in MgNi film prevent formation of the barrier oxide layer and support dynamic structural rearrangements. A tentative explanation of the observations may be as-follows: incident oxygen atoms during interaction with C-contaminated surface react with C atoms and form CO2 molecules which are desorbed. For C-free MgNi films, the oxygen arriving at the surface forms oxides and builds a barrier layer hindering ensuing hydrogenation. The surface barrier layer blocks relaxation processes and inhibits dynamic behavior of grains.
4.
Discussion
Hydrogen molecules dissociate on metal surfaces, dissolve atomically into the substrate, and interact with the atoms in the bulk. Adsorption–desorption processes intimately control the degree to which surfaces are contaminated. The H-uptake
Fig. 5 – Distribution profiles of H, C and O atoms in the near-surface region of the C-free MgNi film after 3 h of hydrogenation.
mechanism depends not only on the surface atom composition and hydrogen concentration in the bulk but also on the prevalent microstructure; i.e., the mixture of crystal defects and impurities, grain boundaries, grain size, grain orientation, and phase composition present within the metal matrix. Under hydrogenation, the synthesis of new phases activates the internal structure and initiates relaxation processes. In this way, the H-uptake rate may be modified as the film properties change. A fast and complete transformation of metal into the hydride phase is observed for Ti-capped Mg and for Ccontaminated MgNi films when the dynamic state of grains and stress-driven relaxation processes are not suppressed during hydrogenation. If newly nucleated structural formations inhibit the dynamic state, the hydrogenation rate sharply decreases and only partial transformation of metallic into hydride phase is registered. The purposefully formed Al2O3 barrier on the surface, which suppresses the stress relaxation processes [18], sharply diminishes the hydrogenation rate. For C-free MgNi films, a fast H-uptake takes place during the initial stages of hydrogenation. The nucleation of a thin Mg2NiH4 layer near the film–substrate interface region increases the chemical potential of the region and the H-flow from the surface into the bulk decreases. As the dynamic structural rearrangements subside, a dense surface oxide barrier layer is formed. It blocks the H-uptake into the bulk and hinders stress-driven relaxation processes. The system becomes kinetically locked with low hydrogenation rate. The C-contaminated MgNi films are fully hydrogenated after only 10–15 min at 250 C. It seems that oxygen adsorption and carbon desorption processes intimately control the degree to which surfaces are contaminated. Carbon impurities prevent formation of the surface oxide barrier layer as arriving oxygen reacts with carbon, forming volatile CO2. The fast H-uptake may be explained in the following way. It is well-established that an adsorption of reactive atoms renders metallic surfaces dynamic. The surface chemistry is directed to optimize bonding between reacting species and it results in surface diffusion of adatoms over distances of many atomic bonds. Continuous reactive adsorption makes the surface structure continually adjust to changes in the chemical environment, which leads to local and long-range
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restructuring, depending on the properties of the near-surface region. In this way, the surface layer becomes both chemically and physically distinct from the underlying bulk material. Thus, a difference in chemical potentials is established between chemically activated surface, grain boundaries and film–substrate interface region. The excess of the surface chemical potentials relative to the grain boundaries and the interface region produce a net flux of H atoms into the grain boundaries and to the interface region. It therefore generates a compressive stress in the grains and across film thickness. If the stress exceeds the limit of plasticity, stress relaxation occurs through the motion of dislocations. If dislocation motion is prevented, which occurs at the grain boundaries, formation of subgrains within the original grain structure takes place. The development of a cell structure occurs by entanglement of dislocations on various slip systems. Subgrains develop by (dynamic) recovery. In this way the uptake of hydrogen atoms through the grain boundaries in open contact with the highly activated surface occurs. Thus, the high surface chemical potentials are not only imposed on the external surface but also on the internal grain interfaces. The diffusivity along boundaries is significantly higher (by more than three orders of magnitude) than the lattice diffusivity of reactive species [19]. An estimate of internal stresses gives that accommodated biaxial stresses are on the order of few Gpa [20] which exceeds yield stress. Thus, compressive stress provides a strong driving force for the plasticity in the grains. The plastic deformation of grains supports the dynamic state of grains and their fragmentation and inhibits formation of continuous oxide layer. Consequently, a fast inward diffusion of hydrogen through the grain boundaries takes place.
5.
Conclusions
Hydrogenation of MgNi film is fast and complete if their surface layer is permeable for H atoms. The dynamic state of this surface is connected with structural transformations in the bulk under hydrogenation. Hydrogen is thus driven from the surface region through the grain boundaries of nanocrystallites to the film–substrate interface where the chemical potential of the local microstructure is diminished. Barrier layers on the surface inhibit transport of H atoms from the surface into the bulk, sustain relaxation processes and suppress activated state. A partial transformation of metallic into hydride phase is registered when hydrides are nucleated at the initial stages of hydrogenation, subsequently inhibiting the dynamic processes in the bulk. Small quantities of impurities may change behavior of hydrogen on the surface and in the bulk and can induce large differences in the hydrogen uptake mechanism and the amount of hydride formed.
Acknowledgments This work has been financially supported by FP6 project (No. 038941; HySIC) and by the European Research Training
Network (Contract No. MRTN-CT-2004-512443 HyTRAIN). Partial fundings by the European Commission DG Research (SES6-2006-518271/NESSHY) and the Lithuanian Science Foundation (H2technologijos, 2008–2010) are also gratefully acknowledged.
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