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Substrate effects on formation and hydrogenation of Mg–Ni films
Applied Surface Science 263 (2012) 202–209
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Substrate effects on formation and hydrogenation of Mg–Ni films Martynas Lelis a,b,∗ , Darius Milcius a,1 , Dag Noréus c a b c
Lithuanian Energy Institute, Centre for Hydrogen Energy Technologies, Breslaujos st. 3, LT-44403 Kaunas, Lithuania Vytautas Magnus university, Faculty of Natural Sciences, Department of Physics, Vileikos st. 8, LT-44404 Kaunas, Lithuania Department of Structural Chemistry, Arrhenius Laboratory, Stockholm University, S-106 91 Stockholm, Sweden
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Article history: Received 8 June 2012 Received in revised form 23 August 2012 Accepted 7 September 2012 Available online 15 September 2012 Keywords: Hydrogen energy Hydrogen storage Mg2 NiH4 Metal hydride Film Substrate modification Plasma
a b s t r a c t Usually metallic films for metal hydrides research applications are deposited on hard and flat substrates such as silicon, magnesium oxide, fused silica or quartz glass. With the film thickness increases all films during hydrogenation deals with typical problems such as film brittleness and cracking. In this paper we demonstrate that metal Mg–Ni films for hydrogen storage can be successfully deposited on to the flexible low surface energy expanded PTFE substrates. The received results for soft substrates (expanded PTFE) are compared to films being deposited on crystalline silicon substrate with and without plasma pretreatment. It is observed that different interface zone between substrate and film has great affect on both film crystallinity and its reaction with hydrogen. It is also demonstrated that modifying surface of the substrate might have affect on film microstructure before and after hydrogenation. © 2012 Elsevier B.V. All rights reserved.
1. Introduction Metal hydrides were discovered in the beginning of the nineteenth-century but for a long time remained interesting only for the fundamental research [1]. However, the development of fundamental understanding and experimental observations of the metal hydride properties revealed their relevance to the industrial applications. For the last decades a lot of the research was focused to use metal hydrides in energy related applications [2–6]. It was shown that advanced solid state hydrogen storage systems potentially can become one of the key elements for the future energy systems, meanwhile rechargeable nickel–metal hydride batteries are already widely spread in the global market and successfully competes with the Li based alternatives [7]. Applications of thin films in metal hydrides research studies have gained separate interest because most of the Physical Vapour Deposition (PVD) technologies allow synthesizing nanostructured materials in much more controllable way in comparison to the chemical methods or semi-mechanical ball milling technique [8]. Furthermore, synthesis of the metallic films and its subsequent hydrogenation opened up possibilities for the in situ and
∗ Corresponding author at: Breslaujos st. 3, LT-44403 Kaunas, Lithuania. Tel.: +370 61 252924/37 401801/37 327910; fax: +370 37 351271. E-mail addresses: [email protected], [email protected], [email protected] (M. Lelis), [email protected] (D. Milcius), [email protected] (D. Noréus). 1 Tel.: +370 37 401801; fax: +370 37 351271. 0169-4332/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.apsusc.2012.09.028
ex situ optical spectroscopy and electrical resistivity measurements which are difficult or even impossible to implement for the ball milled powders [9,10]. Indeed, then Huiberts et al. synthesized and hydrided yttrium and lanthanum films they observed that transition from dihydride to trihydride leads to the significant changes in optical and electrical properties of the films [11]. The later transition from the conductive metallic mirror state to transparent semiconductor or dielectric hydride phase was also observed for other metal hydrides including magnesium–rareearth and magnesium–transition-metal hydrides [12–15]. These observations were important not only from the fundamental point of view but also opened up new areas for the metal hydride film applications such as hydrogen sensors and switchable mirror devices [16–20]. Mg2 NiH4 hydride has theoretical hydrogen content of 3.6 wt.%, good hydrogen absorption–desorbtion cycle stability [21,22], moderate decomposition temperature [23,24] and exhibits optical switching phenomena then produced in thin film form [14,17]. Therefore it was a target of a lot of research both in hydrogen storage and metal-hydride-based devices (hydrogen sensors, switchable mirrors and recently even as possible candidate to replace silicon in solar cells) [25]. Mg2 NiH4 has three crystal phases. Low temperature phases (below 510 K) LT-1 and LT-2 both have monoclinic structure respectively without and with microtwinings; and high temperature (HT) phase has FCC structure [26,27]. Blomqvist and Noréus demonstrated that at room temperature Mg2 NiH4 powders of LT-1 phase has brownish-gray color meanwhile LT-2 phase powders has characteristic orange color [28].
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When heating across the LT → HT transition at 510 K, both phases transforms to HT phase and change their color to black. Coiling down of HT phase Mg2 NiH4 powders leads to the transition to LT-2 phase and reappearance of orange color independent of what phase powders were before heating up. At the same article authors demonstrate that both conductivity and crystal structure of Mg2 NiH4 powders is affected by the applied mechanical stress [28]. All these features of Mg2 NiH4 hydride and particularly its phase changes resulting into different conductivity and optical properties boosted research for Mg2 NiH4 synthesis in thin film form. It was expected that hydrogenation properties of Mg2 Ni alloy thin films will be similar to powders, whereas experiments showed that situation with thin films is more complicated. Contrary to expectation our earlier study has shown that hydrogenation of Mg2 Ni films formed on quartz and calcium fluorite substrates results into pseudo cubic phase (instead of LT-1 for hydrogenation temperatures below 510 K) which remains unchanged over all temperature range from room temperature up to its decomposition at around 550 K [29,30]. Gremaud et al. investigated optical properties of Mgy Ni1−y Hx gradient thin films which were hydrided at room temperature and also reported yellow-to-red colors suggest that it could have been LT-2 or pseudo cubic phase (experimentaly measured hydride structure was not reported) [31]. But Jain et al. showed that formation of crystalline monoclinic Mg2 NiH4 phase is also possible [32] whereas Westerwaal et al. claimed that in film form Mg2 NiH4 hydride is amorphous [33,34]. Interestingly, later group has demonstrated that hydrogenation of Mg2 Ni alloy starts by formation of ␣ hydride phase (Mg2 NiH0.3 solid solution) at the interface with the substrate followed by the  hydride phase which also forms at the interface and proceeds to the film surface [33]. Up to the best of our knowledge there were no systematic approaches to investigate what factors introduced by magnetron sputtering technique or thin film interaction with substrate are responsible for the dissimilarities between Mg2 NiH4 hydride properties in film and powder forms. However, it is highly believable that it is substrate and its interaction with the thin film that becomes the main factor affecting hydrogenation properties. Experimental observations indicate that when hydrogen enters the film the in-plane expansion is limited and this introduces the in-plane stress which changes the thermodynamic properties compared to the corresponding bulk material [35,36]. Another group demonstrated that hydrogenation of Mg thin films deposited on porous and non-porous substrates can have effect for the micro-strains inside the film which alters Mg film hydrogenation properties [37]. Considering later information in this work we wanted to test if using dissimilar substrates and varying their pretreatment conditions we could have significant effects on hydrogenation properties of Mg2 Ni films deposited by magnetron sputtering.
2. Materials and methods In this work we used PVD-75 magnetron sputtering system manufactured by Kurt J. Lesker company with two independent 3 size Torus DC powered magnetrons. The system was equipped with cryopump which allows to achieve a base vacuum level down to 5 × 10−8 mbar (5.05 × 10−6 Pa). For sputtering we used 99.99% purity magnesium and 99.999% purity nickel targets obtained from Kurt J. Lesker Company. As it was already mentioned during our earlier experiments [29] as substrates we have used amorphous quartz and polycristalline calcium fluorite. Therefore, in order to check if usage of different substrates have effects on Mg2 NiH4 hydride formation properties we have chosen two types of substrate with completely opposing properties from those which were used
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Fig. 1. SEM images of expanded PTFE substrates: (a) expanded PTFE substrate washed with acetone/alcohol and (b) expanded PTFE substrate was washed with acetone/alcohol and treated with pulsed DC plasma.
during earlier experiments. First of them was polished prime grade silicon wafers with (1 1 1) surface orientation received from Siegert Consulting e.K. In contrast to quartz and CaF2 , silicon wafers have flat and periodical surface even at the atomic scale which is guaranteed by (1 1 1) crystalline orientation. Furthermore, it is known that at higher temperatures (around 600 ◦ C) magnesium and silicon can form magnesium silicates [38] which can suppose much higher Mg–Ni film and substrate interaction comparing to film–substrate pairs used in earlier studies. For the second substrate we were looking for substrate with as low surface energy as possible. We have chosen to use expanded PTFE from W.L. Gore & Associates Inc. Expanded PTFE is widely known for very low free surface energy and apart from that it also (i) exhibits very good temperature stability in −268 ◦ C to +315 ◦ C temperature range; (ii) is dimensionally stable with no aging or degradation; (iii) chemically inert-resistant to all media in the 0–14 pH range, except molten alkali metals and element fluorine; (iv) possible operating pressure from vacuum up to 200 bar. The combination of low surface energy, chemical inertness of the expanded PTFE substrate as well as its flexible microstructure is expected to result in much smaller micro-strains which are formed in Mg–Ni–H films. For both types of substrates we have used three different surface pre-treatment procedures. For the first group of samples Si and expanded PTFE substrates were washed in acetone in ultrasonic bath, dried out under dry air flow and then wiped out with cloth rinsed in ethyl alcohol. Substrates for the second group of samples in addition to washing as described above before film deposition were affected by DC power source generated plasma (no contact with air between plasma treatment and thin film
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Fig. 2. SEM images of as deposited Mg–Ni thin films on Si substrates: (a) surface of Mg–Ni thin film on washed silicon, (b) cross section of Mg–Ni thin film on washed silicon, (c) surface of Mg–Ni thin film on washed and DC plasma treated silicon, (d) cross section of Mg–Ni thin film on washed and DC plasma treated silicon, (e) surface of Mg–Ni thin film on washed and pulsed DC plasma treated silicon and (f) cross section of Mg–Ni thin film on washed and pulsed DC plasma treated silicon.
deposition); i.e. constant negative potential of 900 V was applied to the substrate holder for 15 min at 2 × 10−2 mbar pressure prior to the film deposition. DC power source ignited plasma and constantly attracted positive Ar+ ions which induced mainly ballistic interaction with the substrate. Substrates for the third group were washed and treated with pulsed DC power source generated plasma. In pulsed DC regime power source used long 650 V negative potential with shorting out and reversing the target voltage to roughly 100 V for 5 s at 20-kHz rate. This reversal of voltage periodically attracts electrons from plasma and prevents charge buildup on the sample holder as well as substrate surfaces and increases process efficiency. Exposure of substrate for electron bombardment also can induce modification in electronic structure of substrate which is usually short term for metals and semiconductors and long term polymeric materials [39,40]. We have chosen to use plasma based pretreatment of the substrate because it is quite often used in thin film deposition experiments, therefore it is important to know if it has any effect on hydrogenation properties. The films were deposited by 15 min of co-sputtering of Mg and Ni at 6 × 10−3 mbar argon pressure using constant 0.5 A and 0.2 A
magnetron currents respectively. All films were produced without any top cap layer in order to form natural oxide on the top surface of the film and in this way to pasivate film surface and to enhance the hydrogenation process at the interface. All as-deposited thin films were analyzed by using stylus profiler (Ambios XP-200), X-ray diffraction (Bruker D8) and Scanning Electron Microscopy (JOEL JSM-5600 with Bruker Quad 5040 EDS system) techniques. After initial characterization all films were hydrogenated for 72 h at 20 bar hydrogen pressure and constant 180 ◦ C temperature. After hydrogenation thin films were reanalyzed with XRD and SEM techniques. The received data was analyzed in two aspects: (i) how microstructure of the hydrogenated films change compared to the initial samples; and (ii) is there any disparity between crystal structures of hydrogenated Mg–Ni thin films deposited on dissimilar substrates. In particular it was interesting to see if it is possible to receive the low temperature phase (LT-1 and/or LT-2) magnesium nickel hydride which was not achievable during earlier experiments when films were deposited on quartz and calcium fluoride substrates [29].
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Fig. 3. Film thickness measurement results for Mg–Ni films deposited on Si substrate after different pre-treatment.
3. Results and discussion Before film deposition we used SEM to see how surfaces of silicon and expanded PTFE substrates look after they were washed with acetone/alcohol and treated with pulsed DC plasma. As it might be expected prime grade silicon wafers with and without plasma treatment were too flat for SEM to reveal any morphology change. Meanwhile, SEM analysis of expanded PTFE substrates with and without plasma treatment reveals extensive changes (Fig. 1) which are induced by bombardment with ions and electrons. Later observations allow to assume that by using plasma treatment we will alter the interface region between substrate and thin film even if none of the later two are changed. SEM images of metallic Mg–Ni films deposited on silicon substrates are shown in Fig. 2. From these images we can identify that plasma treatment has observable effect on film morphology. In Fig. 2a we can see that Mg–Ni film is rugged having small (50–100 nm) ridges. Similar ridges just of smaller dimensions (20–30 nm) still can be identified in Fig. 2c and are completely absent in Fig. 2e. These observations indicate that plasma treatment improves film homogeneity, augment its density and makes film flatter which means it repeats silicon substrate surface better than plasma untreated sample. Such changes can be attributed to the removal of organic contamination from the surface and/or strengthening of the film–substrate bond. The qualitative statements of augmented film density are supported by the quantitative film thickness measurement data observed with stylus profiler (Fig. 3). The cross-section images of Mg–Ni films deposited on silicon reveal that in all cases it has characteristic columnar structure (Fig. 2). The diameter of individual columns seems to be smallest for Mg–Ni deposited on untreated Si and progressively increases for DC and pulsed DC treated samples. If we would assume that substrate temperature increases significantly during plasma treatment then such tendency agrees with film growing model proposed Thorton [41]. Nevertheless, in contrast to Westerwaal et al. who also investigated structure of Mg–Ni film on Si substrate [33] at the available SEM resolution we did not noticed that Mg–Ni films would have clear smaller interfacial grains. SEM images for Mg–Ni films deposited on expanded PTFE substrate are provided in Fig. 4. Images reveal that Plasma treatment of expanded PTFE substrates not only brakes structural “strings”
Fig. 4. SEM images of as deposited Mg–Ni thin films on expanded PTFE substrates: (a) Mg–Ni thin film on washed, (b) Mg–Ni thin film on washed and DC plasma treated and (c) Mg–Ni thin film on washed and pulsed DC plasma treated expanded PTFE substrates.
but also significantly improves film adhesion to the expanded PTFE which results in better surface coverage with the film. This can be explained by the extremely low surface energy of untreated expanded PTFE and its increase induced by possible bond braking, crosslinking, carbonisation and other events which are present during plasma treatment. It was not possible to use surface profiler for films deposited on expanded PTFE substrates therefore the film thickness was roughly estimated by comparing SEM images of uncoated and coated expanded PTFE. The measured rough thickness of the film deposited on plasma untreated substrate was 200 nm, meanwhile for the plasma treated samples it was estimated around 350 nm.
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Fig. 5. SEM images of as hydrogenated Mg–Ni thin films on different substrates: (a) Mg–Ni thin film on washed Si, (b) Mg–Ni thin film on washed expanded PTFE, (c) Mg–Ni thin film on washed and DC plasma treated Si, (d) Mg–Ni thin film on washed and DC plasma treated expanded PTFE, (e) Mg–Ni thin film on washed and pulsed DC plasma treated Si and (f) Mg–Ni thin film on washed and pulsed DC plasma treated expanded PTFE.
Elemental composition measurements by EDS revealed that for all samples Mg and Ni atomic concentration respectively were close to 71 at.% and 29 at.%, respectively. X-ray diffraction analysis of asdeposited Mg–Ni thin films on both silicon and expanded PTFE substrates has revealed that in all cases deposited Mg–Ni film is formed in amorphous phase. After hydrogenation surface characteristics of Mg–Ni thin films deposited on silicon substrates have changed drastically (Fig. 5). For the substrate which did not had any plasma pre-treatment (Fig. 5a) we can see that the film is flatter and probably more homogeneous than it was before the hydrogenation, meanwhile SEM images of Mg–Ni films on plasma pre-treated silicon substrates (Fig. 5c and e) presuppose that during hydrogenation they have undergone through much higher mass transport phenomena which induced various structures at the surface and therefore it is reasonable to assume that these two samples have adsorbed more hydrogen than the first one. Also by the surface morphology changes (signs of expansion) we can predict that plasma treatment was favorable for the formation of magnesium nickel hydride.
Looking at SEM images of hydrogenated Mg–Ni films which were deposited on expanded PTFE substrates we can see that the crosscut dimensions of strings in Fig. 5b is roughly 1.5 times bigger than in Fig. 4a and this reveal high level of thin film expansion. The measured crosscut dimension increase of films deposited on plasma treated expanded PTFE samples is slightly smaller – 1.3 times. These results show that Mg–Ni films deposited on expanded PTFE substrate can absorb hydrogen and expand significantly in all directions without braking or falling apart from the substrate. When metallic film is deposited on hard flat substrate (such as Si) during the hydrogenation and/or dehydrogenation processes it is hard to avoid film cracking and falling off of the substrate [37]. Therefore, later results of expanded PTFE substrate usage for hydride film formation can be useful in trying to find proper substrates for the hydrides for functional devices such as hydrogen sensors. X-ray diffraction data for hydrogenated Mg–Ni films on silicon and expanded PTFE substrates are shown in Figs. 6 and 7 respectively. As a general trend we can notice that XRD diffractograms of hydrogenated samples on differently pretreated Si substrates
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Fig. 6. X-ray diffractograms of hydrogenated Mg–Ni thin films deposited on silicon substrates with different pre-treatment.
differ meanwhile, those for the expanded PTFE substrates looks rather similar. The diffraction patterns of hydrogenated Mg–Ni films deposited on untreated and pulsed DC plasma treated Si substrates have all of the strongest peaks related to the ␣-phase Mg2 NiH0.3 hydride (solid solution phase). The diffraction pattern of hydrogenated Mg–Ni film deposited on DC plasma treated Si substrate has several additional peaks those interpretations and indexing is not straightforward. However, there are several factors which allow assuming that later sample (DC plasma treated) has a coexistence of two crystal phases, namely ␣ hydride phase Mg2 NiH0.3 and monoclinic low temperature  hydride phase Mg2 NiH4 . First of all, to advocate such assumption we can notice the presence of the same peaks as in former (untreated and pulsed DC plasma treated) samples which were already attributed to the Mg2 NiH0.3 phase. In powder diffraction case monoclinic Mg2 NiH4 LT phase is easily recognized by its characteristic duplet peaks at roughly 23.3◦ and 24.1◦ which are coming from (−1 1 2) and (1 1 2) crystallographic planes respectfully. In our case for the thin film sample we observe only single peak at 23.3◦ and there is no second peak at 24.1◦ , but this can be explained by the preferential crystal structure orientation of Mg2 NiH4 crystal phase. The preferential orientation of thin films deposited by magnetron sputter technique is a well known phenomena which can limit observable X-ray diffraction peaks down to a single or several reflections from closely oriented planes [42–44]. In the case of monoclinic Mg2 NiH4 crystal phase the orientations of (−1 1 2) and (1 1 2) crystallographic planes are differing significantly and this can explain the absence of several intensive peaks of monoclinic Mg2 NiH4 phase which are present in the corresponding powder diffraction patterns at 24.1◦ (1 1 2), 39.1◦ (0 2 4), 39.5◦ (2 2 0) and 40◦ (2 0 4). The provided assumption of presence of monoclinic Mg2 NiH4 phase is further supported by the experimental observation that both of the strongest additionally present peaks at roughly 23.3◦ and 37.8◦ can be respectively attributed to the reflections from relatively close crystallographic planes (−1 1 2) and (−2 0 4). The same is applied for the appearance of smaller peaks at 19.5◦ (larger shoulder), 31.4◦ , 32.9◦ , 48.7◦ and 55.7◦ . Furthermore, the intensive peak at 37.8◦ only can be attributed to the monoclinic Mg2 NiH4 phase as none of the Mg, Ni, Mg2 Ni, MgNi2 , Mg2 NiH4 , Mg2 NiH0.3 or corresponding oxides
phases have X-ray diffraction peaks at this angle. Summarizing observations from Fig. 6 we can see that according to the crystal phase information the highest hydrogen content (partial beta hydride) is received for the hydrogenated Mg–Ni film which was deposited on DC plasma treated silicon substrate, whereas films deposited on other substrates has less hydrogen containing phase. This is in good agreement with predictions which were made looking at SEM images (Fig. 5). As it can be seen from Fig. 7 all three Mg–Ni films deposited on expanded PTFE substrate during hydrogenation formed pseudo cubic Mg2 NiH4 hydride, i.e. different crystal structure (in comparison to films deposited on Si) even if the hydrogenation conditions were fixed for all six samples. The observed cubic phase of Mg2 NiH4 in general is similar to those reported in our earlier studies with films deposited on quartz and CaF2 substrates [29]. It is important to point up that  magnesium nickel hydride is formed only for the sample which was deposited on DC plasma pre-treated substrate and no clear signs crystalline  hydride formation was observed for other two samples which were deposited on differently pretreated silicon substrates. Another striking result is that after DC plasma pre-treatment we were able to form low temperature monoclinic Mg2 NiH4 hydride which was not observed on expanded PTFE as well as on other substrates which were used in earlier study [29]. The observed results (especially for Si substrates) indicate that probably there is a range of interface region conditions under which we witness best formation of crystalline Mg2 NiH4 hydride as well as preference for monoclinic or cubic hydride phase formation. These conditions can include several factors amongst those we can think about: (i) proper adhesion strength which leads to the certain amount of in-plane stress during hydrogenation, (ii) selected topography and temperature of substrate surface which leads to preferential growth of certain microstructures (for example columns of optimal diameter and/or formation of smaller grain interfacial zones), (iii) presence of particular crystal phases at the interface (for example magnesium silicates) and/or elements which could act as hydrogenation catalyst and/or centers for hydride phase nucleation. Other factors can be equally important though more experimental data and systematic studies are needed in order point out most important contributors as well as to test if these statements are also valid for other hydrogen absorbing materials.
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Fig. 7. X-ray diffractograms of hydrogenated Mg–Ni thin films deposited on expanded PTFE substrates with different pre-treatment.
4. Conclusions During current study it was qualitatively demonstrated that at least for Mg–Ni films using of substrates with opposing properties and various preparation methods can lead to different conditions in the interface region between substrate and the film which could lead to changes in properties of hydride formation. These observations indicate that for the controllable formation of the particular metal hydride film its substrate might not be a completely freely selectable variable. Therefore, our main conclusion is that it is important that metal hydride film and its substrate should be considered as an integral system where both elements perform synergistically.
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Acknowledgements This work has been financially supported by the Nordic Energy Research project “Nordic Center of Excellence for Hydrogen Storage Materials”. References
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Contents lists available at SciVerse ScienceDirect
Applied Surface Science journal homepage: www.elsevier.com/locate/apsusc
Substrate effects on formation and hydrogenation of Mg–Ni films Martynas Lelis a,b,∗ , Darius Milcius a,1 , Dag Noréus c a b c
Lithuanian Energy Institute, Centre for Hydrogen Energy Technologies, Breslaujos st. 3, LT-44403 Kaunas, Lithuania Vytautas Magnus university, Faculty of Natural Sciences, Department of Physics, Vileikos st. 8, LT-44404 Kaunas, Lithuania Department of Structural Chemistry, Arrhenius Laboratory, Stockholm University, S-106 91 Stockholm, Sweden
a r t i c l e
i n f o
Article history: Received 8 June 2012 Received in revised form 23 August 2012 Accepted 7 September 2012 Available online 15 September 2012 Keywords: Hydrogen energy Hydrogen storage Mg2 NiH4 Metal hydride Film Substrate modification Plasma
a b s t r a c t Usually metallic films for metal hydrides research applications are deposited on hard and flat substrates such as silicon, magnesium oxide, fused silica or quartz glass. With the film thickness increases all films during hydrogenation deals with typical problems such as film brittleness and cracking. In this paper we demonstrate that metal Mg–Ni films for hydrogen storage can be successfully deposited on to the flexible low surface energy expanded PTFE substrates. The received results for soft substrates (expanded PTFE) are compared to films being deposited on crystalline silicon substrate with and without plasma pretreatment. It is observed that different interface zone between substrate and film has great affect on both film crystallinity and its reaction with hydrogen. It is also demonstrated that modifying surface of the substrate might have affect on film microstructure before and after hydrogenation. © 2012 Elsevier B.V. All rights reserved.
1. Introduction Metal hydrides were discovered in the beginning of the nineteenth-century but for a long time remained interesting only for the fundamental research [1]. However, the development of fundamental understanding and experimental observations of the metal hydride properties revealed their relevance to the industrial applications. For the last decades a lot of the research was focused to use metal hydrides in energy related applications [2–6]. It was shown that advanced solid state hydrogen storage systems potentially can become one of the key elements for the future energy systems, meanwhile rechargeable nickel–metal hydride batteries are already widely spread in the global market and successfully competes with the Li based alternatives [7]. Applications of thin films in metal hydrides research studies have gained separate interest because most of the Physical Vapour Deposition (PVD) technologies allow synthesizing nanostructured materials in much more controllable way in comparison to the chemical methods or semi-mechanical ball milling technique [8]. Furthermore, synthesis of the metallic films and its subsequent hydrogenation opened up possibilities for the in situ and
∗ Corresponding author at: Breslaujos st. 3, LT-44403 Kaunas, Lithuania. Tel.: +370 61 252924/37 401801/37 327910; fax: +370 37 351271. E-mail addresses: [email protected], [email protected], [email protected] (M. Lelis), [email protected] (D. Milcius), [email protected] (D. Noréus). 1 Tel.: +370 37 401801; fax: +370 37 351271. 0169-4332/$ – see front matter © 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.apsusc.2012.09.028
ex situ optical spectroscopy and electrical resistivity measurements which are difficult or even impossible to implement for the ball milled powders [9,10]. Indeed, then Huiberts et al. synthesized and hydrided yttrium and lanthanum films they observed that transition from dihydride to trihydride leads to the significant changes in optical and electrical properties of the films [11]. The later transition from the conductive metallic mirror state to transparent semiconductor or dielectric hydride phase was also observed for other metal hydrides including magnesium–rareearth and magnesium–transition-metal hydrides [12–15]. These observations were important not only from the fundamental point of view but also opened up new areas for the metal hydride film applications such as hydrogen sensors and switchable mirror devices [16–20]. Mg2 NiH4 hydride has theoretical hydrogen content of 3.6 wt.%, good hydrogen absorption–desorbtion cycle stability [21,22], moderate decomposition temperature [23,24] and exhibits optical switching phenomena then produced in thin film form [14,17]. Therefore it was a target of a lot of research both in hydrogen storage and metal-hydride-based devices (hydrogen sensors, switchable mirrors and recently even as possible candidate to replace silicon in solar cells) [25]. Mg2 NiH4 has three crystal phases. Low temperature phases (below 510 K) LT-1 and LT-2 both have monoclinic structure respectively without and with microtwinings; and high temperature (HT) phase has FCC structure [26,27]. Blomqvist and Noréus demonstrated that at room temperature Mg2 NiH4 powders of LT-1 phase has brownish-gray color meanwhile LT-2 phase powders has characteristic orange color [28].
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When heating across the LT → HT transition at 510 K, both phases transforms to HT phase and change their color to black. Coiling down of HT phase Mg2 NiH4 powders leads to the transition to LT-2 phase and reappearance of orange color independent of what phase powders were before heating up. At the same article authors demonstrate that both conductivity and crystal structure of Mg2 NiH4 powders is affected by the applied mechanical stress [28]. All these features of Mg2 NiH4 hydride and particularly its phase changes resulting into different conductivity and optical properties boosted research for Mg2 NiH4 synthesis in thin film form. It was expected that hydrogenation properties of Mg2 Ni alloy thin films will be similar to powders, whereas experiments showed that situation with thin films is more complicated. Contrary to expectation our earlier study has shown that hydrogenation of Mg2 Ni films formed on quartz and calcium fluorite substrates results into pseudo cubic phase (instead of LT-1 for hydrogenation temperatures below 510 K) which remains unchanged over all temperature range from room temperature up to its decomposition at around 550 K [29,30]. Gremaud et al. investigated optical properties of Mgy Ni1−y Hx gradient thin films which were hydrided at room temperature and also reported yellow-to-red colors suggest that it could have been LT-2 or pseudo cubic phase (experimentaly measured hydride structure was not reported) [31]. But Jain et al. showed that formation of crystalline monoclinic Mg2 NiH4 phase is also possible [32] whereas Westerwaal et al. claimed that in film form Mg2 NiH4 hydride is amorphous [33,34]. Interestingly, later group has demonstrated that hydrogenation of Mg2 Ni alloy starts by formation of ␣ hydride phase (Mg2 NiH0.3 solid solution) at the interface with the substrate followed by the  hydride phase which also forms at the interface and proceeds to the film surface [33]. Up to the best of our knowledge there were no systematic approaches to investigate what factors introduced by magnetron sputtering technique or thin film interaction with substrate are responsible for the dissimilarities between Mg2 NiH4 hydride properties in film and powder forms. However, it is highly believable that it is substrate and its interaction with the thin film that becomes the main factor affecting hydrogenation properties. Experimental observations indicate that when hydrogen enters the film the in-plane expansion is limited and this introduces the in-plane stress which changes the thermodynamic properties compared to the corresponding bulk material [35,36]. Another group demonstrated that hydrogenation of Mg thin films deposited on porous and non-porous substrates can have effect for the micro-strains inside the film which alters Mg film hydrogenation properties [37]. Considering later information in this work we wanted to test if using dissimilar substrates and varying their pretreatment conditions we could have significant effects on hydrogenation properties of Mg2 Ni films deposited by magnetron sputtering.
2. Materials and methods In this work we used PVD-75 magnetron sputtering system manufactured by Kurt J. Lesker company with two independent 3 size Torus DC powered magnetrons. The system was equipped with cryopump which allows to achieve a base vacuum level down to 5 × 10−8 mbar (5.05 × 10−6 Pa). For sputtering we used 99.99% purity magnesium and 99.999% purity nickel targets obtained from Kurt J. Lesker Company. As it was already mentioned during our earlier experiments [29] as substrates we have used amorphous quartz and polycristalline calcium fluorite. Therefore, in order to check if usage of different substrates have effects on Mg2 NiH4 hydride formation properties we have chosen two types of substrate with completely opposing properties from those which were used
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Fig. 1. SEM images of expanded PTFE substrates: (a) expanded PTFE substrate washed with acetone/alcohol and (b) expanded PTFE substrate was washed with acetone/alcohol and treated with pulsed DC plasma.
during earlier experiments. First of them was polished prime grade silicon wafers with (1 1 1) surface orientation received from Siegert Consulting e.K. In contrast to quartz and CaF2 , silicon wafers have flat and periodical surface even at the atomic scale which is guaranteed by (1 1 1) crystalline orientation. Furthermore, it is known that at higher temperatures (around 600 ◦ C) magnesium and silicon can form magnesium silicates [38] which can suppose much higher Mg–Ni film and substrate interaction comparing to film–substrate pairs used in earlier studies. For the second substrate we were looking for substrate with as low surface energy as possible. We have chosen to use expanded PTFE from W.L. Gore & Associates Inc. Expanded PTFE is widely known for very low free surface energy and apart from that it also (i) exhibits very good temperature stability in −268 ◦ C to +315 ◦ C temperature range; (ii) is dimensionally stable with no aging or degradation; (iii) chemically inert-resistant to all media in the 0–14 pH range, except molten alkali metals and element fluorine; (iv) possible operating pressure from vacuum up to 200 bar. The combination of low surface energy, chemical inertness of the expanded PTFE substrate as well as its flexible microstructure is expected to result in much smaller micro-strains which are formed in Mg–Ni–H films. For both types of substrates we have used three different surface pre-treatment procedures. For the first group of samples Si and expanded PTFE substrates were washed in acetone in ultrasonic bath, dried out under dry air flow and then wiped out with cloth rinsed in ethyl alcohol. Substrates for the second group of samples in addition to washing as described above before film deposition were affected by DC power source generated plasma (no contact with air between plasma treatment and thin film
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Fig. 2. SEM images of as deposited Mg–Ni thin films on Si substrates: (a) surface of Mg–Ni thin film on washed silicon, (b) cross section of Mg–Ni thin film on washed silicon, (c) surface of Mg–Ni thin film on washed and DC plasma treated silicon, (d) cross section of Mg–Ni thin film on washed and DC plasma treated silicon, (e) surface of Mg–Ni thin film on washed and pulsed DC plasma treated silicon and (f) cross section of Mg–Ni thin film on washed and pulsed DC plasma treated silicon.
deposition); i.e. constant negative potential of 900 V was applied to the substrate holder for 15 min at 2 × 10−2 mbar pressure prior to the film deposition. DC power source ignited plasma and constantly attracted positive Ar+ ions which induced mainly ballistic interaction with the substrate. Substrates for the third group were washed and treated with pulsed DC power source generated plasma. In pulsed DC regime power source used long 650 V negative potential with shorting out and reversing the target voltage to roughly 100 V for 5 s at 20-kHz rate. This reversal of voltage periodically attracts electrons from plasma and prevents charge buildup on the sample holder as well as substrate surfaces and increases process efficiency. Exposure of substrate for electron bombardment also can induce modification in electronic structure of substrate which is usually short term for metals and semiconductors and long term polymeric materials [39,40]. We have chosen to use plasma based pretreatment of the substrate because it is quite often used in thin film deposition experiments, therefore it is important to know if it has any effect on hydrogenation properties. The films were deposited by 15 min of co-sputtering of Mg and Ni at 6 × 10−3 mbar argon pressure using constant 0.5 A and 0.2 A
magnetron currents respectively. All films were produced without any top cap layer in order to form natural oxide on the top surface of the film and in this way to pasivate film surface and to enhance the hydrogenation process at the interface. All as-deposited thin films were analyzed by using stylus profiler (Ambios XP-200), X-ray diffraction (Bruker D8) and Scanning Electron Microscopy (JOEL JSM-5600 with Bruker Quad 5040 EDS system) techniques. After initial characterization all films were hydrogenated for 72 h at 20 bar hydrogen pressure and constant 180 ◦ C temperature. After hydrogenation thin films were reanalyzed with XRD and SEM techniques. The received data was analyzed in two aspects: (i) how microstructure of the hydrogenated films change compared to the initial samples; and (ii) is there any disparity between crystal structures of hydrogenated Mg–Ni thin films deposited on dissimilar substrates. In particular it was interesting to see if it is possible to receive the low temperature phase (LT-1 and/or LT-2) magnesium nickel hydride which was not achievable during earlier experiments when films were deposited on quartz and calcium fluoride substrates [29].
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Fig. 3. Film thickness measurement results for Mg–Ni films deposited on Si substrate after different pre-treatment.
3. Results and discussion Before film deposition we used SEM to see how surfaces of silicon and expanded PTFE substrates look after they were washed with acetone/alcohol and treated with pulsed DC plasma. As it might be expected prime grade silicon wafers with and without plasma treatment were too flat for SEM to reveal any morphology change. Meanwhile, SEM analysis of expanded PTFE substrates with and without plasma treatment reveals extensive changes (Fig. 1) which are induced by bombardment with ions and electrons. Later observations allow to assume that by using plasma treatment we will alter the interface region between substrate and thin film even if none of the later two are changed. SEM images of metallic Mg–Ni films deposited on silicon substrates are shown in Fig. 2. From these images we can identify that plasma treatment has observable effect on film morphology. In Fig. 2a we can see that Mg–Ni film is rugged having small (50–100 nm) ridges. Similar ridges just of smaller dimensions (20–30 nm) still can be identified in Fig. 2c and are completely absent in Fig. 2e. These observations indicate that plasma treatment improves film homogeneity, augment its density and makes film flatter which means it repeats silicon substrate surface better than plasma untreated sample. Such changes can be attributed to the removal of organic contamination from the surface and/or strengthening of the film–substrate bond. The qualitative statements of augmented film density are supported by the quantitative film thickness measurement data observed with stylus profiler (Fig. 3). The cross-section images of Mg–Ni films deposited on silicon reveal that in all cases it has characteristic columnar structure (Fig. 2). The diameter of individual columns seems to be smallest for Mg–Ni deposited on untreated Si and progressively increases for DC and pulsed DC treated samples. If we would assume that substrate temperature increases significantly during plasma treatment then such tendency agrees with film growing model proposed Thorton [41]. Nevertheless, in contrast to Westerwaal et al. who also investigated structure of Mg–Ni film on Si substrate [33] at the available SEM resolution we did not noticed that Mg–Ni films would have clear smaller interfacial grains. SEM images for Mg–Ni films deposited on expanded PTFE substrate are provided in Fig. 4. Images reveal that Plasma treatment of expanded PTFE substrates not only brakes structural “strings”
Fig. 4. SEM images of as deposited Mg–Ni thin films on expanded PTFE substrates: (a) Mg–Ni thin film on washed, (b) Mg–Ni thin film on washed and DC plasma treated and (c) Mg–Ni thin film on washed and pulsed DC plasma treated expanded PTFE substrates.
but also significantly improves film adhesion to the expanded PTFE which results in better surface coverage with the film. This can be explained by the extremely low surface energy of untreated expanded PTFE and its increase induced by possible bond braking, crosslinking, carbonisation and other events which are present during plasma treatment. It was not possible to use surface profiler for films deposited on expanded PTFE substrates therefore the film thickness was roughly estimated by comparing SEM images of uncoated and coated expanded PTFE. The measured rough thickness of the film deposited on plasma untreated substrate was 200 nm, meanwhile for the plasma treated samples it was estimated around 350 nm.
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Fig. 5. SEM images of as hydrogenated Mg–Ni thin films on different substrates: (a) Mg–Ni thin film on washed Si, (b) Mg–Ni thin film on washed expanded PTFE, (c) Mg–Ni thin film on washed and DC plasma treated Si, (d) Mg–Ni thin film on washed and DC plasma treated expanded PTFE, (e) Mg–Ni thin film on washed and pulsed DC plasma treated Si and (f) Mg–Ni thin film on washed and pulsed DC plasma treated expanded PTFE.
Elemental composition measurements by EDS revealed that for all samples Mg and Ni atomic concentration respectively were close to 71 at.% and 29 at.%, respectively. X-ray diffraction analysis of asdeposited Mg–Ni thin films on both silicon and expanded PTFE substrates has revealed that in all cases deposited Mg–Ni film is formed in amorphous phase. After hydrogenation surface characteristics of Mg–Ni thin films deposited on silicon substrates have changed drastically (Fig. 5). For the substrate which did not had any plasma pre-treatment (Fig. 5a) we can see that the film is flatter and probably more homogeneous than it was before the hydrogenation, meanwhile SEM images of Mg–Ni films on plasma pre-treated silicon substrates (Fig. 5c and e) presuppose that during hydrogenation they have undergone through much higher mass transport phenomena which induced various structures at the surface and therefore it is reasonable to assume that these two samples have adsorbed more hydrogen than the first one. Also by the surface morphology changes (signs of expansion) we can predict that plasma treatment was favorable for the formation of magnesium nickel hydride.
Looking at SEM images of hydrogenated Mg–Ni films which were deposited on expanded PTFE substrates we can see that the crosscut dimensions of strings in Fig. 5b is roughly 1.5 times bigger than in Fig. 4a and this reveal high level of thin film expansion. The measured crosscut dimension increase of films deposited on plasma treated expanded PTFE samples is slightly smaller – 1.3 times. These results show that Mg–Ni films deposited on expanded PTFE substrate can absorb hydrogen and expand significantly in all directions without braking or falling apart from the substrate. When metallic film is deposited on hard flat substrate (such as Si) during the hydrogenation and/or dehydrogenation processes it is hard to avoid film cracking and falling off of the substrate [37]. Therefore, later results of expanded PTFE substrate usage for hydride film formation can be useful in trying to find proper substrates for the hydrides for functional devices such as hydrogen sensors. X-ray diffraction data for hydrogenated Mg–Ni films on silicon and expanded PTFE substrates are shown in Figs. 6 and 7 respectively. As a general trend we can notice that XRD diffractograms of hydrogenated samples on differently pretreated Si substrates
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Fig. 6. X-ray diffractograms of hydrogenated Mg–Ni thin films deposited on silicon substrates with different pre-treatment.
differ meanwhile, those for the expanded PTFE substrates looks rather similar. The diffraction patterns of hydrogenated Mg–Ni films deposited on untreated and pulsed DC plasma treated Si substrates have all of the strongest peaks related to the ␣-phase Mg2 NiH0.3 hydride (solid solution phase). The diffraction pattern of hydrogenated Mg–Ni film deposited on DC plasma treated Si substrate has several additional peaks those interpretations and indexing is not straightforward. However, there are several factors which allow assuming that later sample (DC plasma treated) has a coexistence of two crystal phases, namely ␣ hydride phase Mg2 NiH0.3 and monoclinic low temperature  hydride phase Mg2 NiH4 . First of all, to advocate such assumption we can notice the presence of the same peaks as in former (untreated and pulsed DC plasma treated) samples which were already attributed to the Mg2 NiH0.3 phase. In powder diffraction case monoclinic Mg2 NiH4 LT phase is easily recognized by its characteristic duplet peaks at roughly 23.3◦ and 24.1◦ which are coming from (−1 1 2) and (1 1 2) crystallographic planes respectfully. In our case for the thin film sample we observe only single peak at 23.3◦ and there is no second peak at 24.1◦ , but this can be explained by the preferential crystal structure orientation of Mg2 NiH4 crystal phase. The preferential orientation of thin films deposited by magnetron sputter technique is a well known phenomena which can limit observable X-ray diffraction peaks down to a single or several reflections from closely oriented planes [42–44]. In the case of monoclinic Mg2 NiH4 crystal phase the orientations of (−1 1 2) and (1 1 2) crystallographic planes are differing significantly and this can explain the absence of several intensive peaks of monoclinic Mg2 NiH4 phase which are present in the corresponding powder diffraction patterns at 24.1◦ (1 1 2), 39.1◦ (0 2 4), 39.5◦ (2 2 0) and 40◦ (2 0 4). The provided assumption of presence of monoclinic Mg2 NiH4 phase is further supported by the experimental observation that both of the strongest additionally present peaks at roughly 23.3◦ and 37.8◦ can be respectively attributed to the reflections from relatively close crystallographic planes (−1 1 2) and (−2 0 4). The same is applied for the appearance of smaller peaks at 19.5◦ (larger shoulder), 31.4◦ , 32.9◦ , 48.7◦ and 55.7◦ . Furthermore, the intensive peak at 37.8◦ only can be attributed to the monoclinic Mg2 NiH4 phase as none of the Mg, Ni, Mg2 Ni, MgNi2 , Mg2 NiH4 , Mg2 NiH0.3 or corresponding oxides
phases have X-ray diffraction peaks at this angle. Summarizing observations from Fig. 6 we can see that according to the crystal phase information the highest hydrogen content (partial beta hydride) is received for the hydrogenated Mg–Ni film which was deposited on DC plasma treated silicon substrate, whereas films deposited on other substrates has less hydrogen containing phase. This is in good agreement with predictions which were made looking at SEM images (Fig. 5). As it can be seen from Fig. 7 all three Mg–Ni films deposited on expanded PTFE substrate during hydrogenation formed pseudo cubic Mg2 NiH4 hydride, i.e. different crystal structure (in comparison to films deposited on Si) even if the hydrogenation conditions were fixed for all six samples. The observed cubic phase of Mg2 NiH4 in general is similar to those reported in our earlier studies with films deposited on quartz and CaF2 substrates [29]. It is important to point up that  magnesium nickel hydride is formed only for the sample which was deposited on DC plasma pre-treated substrate and no clear signs crystalline  hydride formation was observed for other two samples which were deposited on differently pretreated silicon substrates. Another striking result is that after DC plasma pre-treatment we were able to form low temperature monoclinic Mg2 NiH4 hydride which was not observed on expanded PTFE as well as on other substrates which were used in earlier study [29]. The observed results (especially for Si substrates) indicate that probably there is a range of interface region conditions under which we witness best formation of crystalline Mg2 NiH4 hydride as well as preference for monoclinic or cubic hydride phase formation. These conditions can include several factors amongst those we can think about: (i) proper adhesion strength which leads to the certain amount of in-plane stress during hydrogenation, (ii) selected topography and temperature of substrate surface which leads to preferential growth of certain microstructures (for example columns of optimal diameter and/or formation of smaller grain interfacial zones), (iii) presence of particular crystal phases at the interface (for example magnesium silicates) and/or elements which could act as hydrogenation catalyst and/or centers for hydride phase nucleation. Other factors can be equally important though more experimental data and systematic studies are needed in order point out most important contributors as well as to test if these statements are also valid for other hydrogen absorbing materials.
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Fig. 7. X-ray diffractograms of hydrogenated Mg–Ni thin films deposited on expanded PTFE substrates with different pre-treatment.
4. Conclusions During current study it was qualitatively demonstrated that at least for Mg–Ni films using of substrates with opposing properties and various preparation methods can lead to different conditions in the interface region between substrate and the film which could lead to changes in properties of hydride formation. These observations indicate that for the controllable formation of the particular metal hydride film its substrate might not be a completely freely selectable variable. Therefore, our main conclusion is that it is important that metal hydride film and its substrate should be considered as an integral system where both elements perform synergistically.
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