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Promising gaseous and electrochemical hydrogen storage properties of porous Mg–Pd films under mild conditions Gongbiao Xin,*a Huiping Yuan,a Lijun Jiang,a Shumao Wang,a Xiaopeng Liua and Xingguo Li*b In this paper, the gaseous and electrochemical hydrogen storage properties of 200 nm Mg–Pd thin films with different morphologies have been investigated. The results show that Mg–Pd films become porous
Received 1st April 2015, Accepted 21st April 2015 DOI: 10.1039/c5cp01897h
with the increase of substrate temperature. Porous Mg–Pd films exhibit superior gaseous and electrochemical hydrogen storage behaviors under mild conditions, including rapid hydrogen sorption kinetics, a large hydrogen storage amount, high electrochemical discharge capacity, and a fast hydrogen diffusion rate. The excellent behaviors of porous Mg–Pd films might be ascribed to the significantly shortened hydrogen diffusion paths and the large contact areas between the hydrogen gas and the solid Mg
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phases, which are elucidative for the development and applications of thick Mg–Pd films.
1. Introduction Hydrogen has been considered as a promising energy carrier in future energy systems.1,2 Therefore, various hydrogen storage materials have been investigated in the last few decades.3–10 Among all of these materials, MgH2 has attracted a lot of attention as a solid-state hydrogen storage medium due to its high gravimetric capacity (7.6 wt%), high volumetric density (110 kg m 3), low cost, abundance and encouraging reversibility.11,12 However, the main obstacles limiting the commercial application of MgH2 are its high thermodynamic stability, high dehydrogenation temperature and sluggish hydrogen sorption kinetics.13 Recently, numerous efforts have been devoted to overcome these disadvantages, including the addition of catalysts,14 alloying with other elements,15 and fabricating various nanostructures.16 Compared with these approaches, Mg-based thin films open up a new opportunity to carry out such studies because we can accurately tailor their composition, interface and crystallinity on the nanoscale.17–19 As a result, a lot of work has been conducted to investigate the hydrogen storage properties of Mg-based thin films.20–24 Particularly, the hydrogen storage properties and related applications of binary Mg–Fe, Mg–Ni, Mg–Ti, Mg–La and ternary Mg–Fe–Ti, Mg–Cr–Ti alloy films have been thoroughly a
Department of Energy Materials and Technology, General Research Institute for Nonferrous Metals, Beijing 100088, China. E-mail: [email protected]; Fax: +86-10-82241294; Tel: +86-10-82241241 b Beijing National Laboratory for Molecular Sciences (BNLMS), (The State Key Laboratory of Rare Earth Materials Chemistry and Applications), College of Chemistry and Molecular Engineering, Peking University, Beijing 100871, China. E-mail: [email protected]; Fax: +86-10-62765930; Tel: +86-10-62765930
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investigated by many groups.25–34 However, experiments on Mg-based films reported so far are mostly operated at high temperatures, while it is much more necessary to clarify the hydrogen sorption kinetics under moderate conditions, as the hydrogen storage materials are expected to be applied at ambient temperature.35,36 In our previous work, we have concentrated our interest on improving the hydrogen storage properties of Mg-based thin films under mild conditions, i.e. low temperature and low hydrogen pressure. The results showed that 100 nm Mg–Pd, Pd–Mg–Pd, Mg–Ti–Pd and Mg–Al–Pd films exhibited promising hydrogen sorption properties at room temperature.37–41 However, the hydrogen sorption behaviors will become extremely disappointing when the Mg layer becomes thicker than 100 nm at ambient temperature. Through insertion of 1 nm thin Ti interlayers, the hydrogen storage kinetics of 500 nm thick Mg–Pd films have been successfully promoted.42 Nevertheless, the hydrogen storage capacity of Mg will be inevitably sacrificed due to the addition of heavy weight transition elements. In this study, we have prepared a series of 200 nm Mg–Pd films at different substrate temperatures by magnetron sputtering. The morphology of the Mg–Pd thin films can be controlled by the deposition temperature. The gaseous and electrochemical hydrogen storage properties of 200 nm Mg–Pd films with different structures have been systematically investigated. The results indicated that Mg–Pd films with porous structures exhibited superior gaseous hydrogen sorption kinetics and promising electrochemical hydrogen storage capacities and discharge stability. The results in this study demonstrated that the hydrogen storage properties of thick Mg–Pd films could be remarkably promoted
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by tailoring their structures without sacrificing the hydrogen storage capacity, providing further guidance for the research and application of thick Mg–Pd films.
2. Experimental 2.1.
Film preparation
Pd-capped 200 nm Mg thin films were prepared at four different substrate temperatures using a custom designed direct current (DC) magnetron sputtering system at a background pressure of around 2 10 4 Pa. Prior to deposition, the substrates were heated to 25 1C, 80 1C, 120 1C and 150 1C, respectively. Then, 200 nm Mg layers were first deposited onto Si(110) wafers and glass substrates using an Mg (99.99%) target. Afterwards, a 10 nm Pd layer was coated on top of the Mg layer using a Pd (99.99%) target to protect Mg against oxidation and to catalyze hydrogen dissociation and recombination. The magnetron discharges were generated under an argon pressure of 0.6 Pa with an argon flow rate of 76 sccm. The samples prepared at four different substrate temperatures are designated as 25 1C, 80 1C, 120 1C and 150 1C separately for simplicity. After deposition, the samples were transferred to a vacuum chamber to avoid oxidation. 2.2.
Structural characterization and property measurements
The structures, resistance changes, optical transmittance properties, cyclic voltammogram and diffusion coefficients during the hydrogen absorption and desorption process were systematically determined. The structural changes of different Mg–Pd films were studied by power X-ray diffraction (XRD) (Rigaku D/max-200) using monochromated Cu Ka radiation and a y–2y scan. The cross-section and surface morphologies of different Mg–Pd films were examined using scanning electron microscopy (SEM) measurements (Hitachi S4800). The resistance changes during hydrogen absorption were recorded in a gas loading cell equipped with a four-probe resistance measurement, monitored using an Agilent 34401A digital multimeter. The optical transmission measurements at 298 K were performed using a UV-vis recording spectrophotometer (Shimadzu UV-2401PC) with a dual beam measurement system. Hydrogen volumetric loading/unloading measurements were performed using a Sieverts apparatus developed at the National Institute of Standards and Technology (NIST) with a high precision pressure transducer (0.003% FS) over a relatively wide pressure range.29,30 All hydrogen absorption processes were carried out at 0.5 bar H2 pressure at 300 K and 353 K; the hydrogen desorption procedure was studied at a backpressure of 10 Pa at the same temperatures. After each desorption, the samples were evacuated for another 1 h before the next absorption cycle. Electrochemical experiments were carried out at room temperature in 6 M KOH solution using a three-electrode cell. Platinum foil and Hg/HgO were used as the counter and reference electrodes, respectively. Cyclic voltammogram tests were performed between the potential range of 1.2 and 0.45 V with a scan rate of 50 mV s 1. For the constant current
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discharge measurements, all the electrodes were first fully charged at a current value of 1 mA followed by 10 min relaxation, and then discharged at a current value of 0.2 mA until the cut-off voltage of 0 V (versus the Hg/HgO reference electrode) was reached. The corresponding capacity retention rate (Sn) can be obtained as: Sn(%) = Cn/Cmax 100%, where Cn is the discharge capacity of the electrode at the nth cycle, and Cmax is the maximum discharge capacity. The diffusion coefficients of different Mg–Pd films were determined by the electrochemical multipotential steps in 6 M KOH solution. All the samples were first held at a cathodically polarized potential ( 1.10 V vs. Hg/HgO) for 1.5 h and subsequently switched to an anodically polarized potential ( 0.5 V vs. Hg/HgO) for another 1.5 h.
3. Results and discussion 3.1.
Structural characterization
The surface morphologies of 200 nm Mg–Pd films prepared at different substrate temperatures are examined by SEM, as shown in Fig. 1. A uniform and smooth surface was obtained for the Mg–Pd film prepared at room temperature. The compact film was composed of numerous regular Mg nanocrystals, which generally had a hexagonal shape and the particle size was approximately 100 nm. Significant changes in the surface morphologies were observed when increasing the substrate temperatures. Many pores appeared in the bulk Mg–Pd film deposited at 80 1C, and the regular hexagonal Mg nanocrystals became undistinguishable (Fig. 1b). In the case of samples prepared at 120 1C and 150 1C, the dense and compact Mg–Pd thin films transformed to a porous and puffy state, with numerous pores and cracks through the whole Mg matrix. The formation of porous Mg–Pd thin films might result from the partial melting of Mg particles at higher temperatures, leading to their random aggregation during deposition. The cross-section SEM images of different 200 nm Mg–Pd films deposited on silicon substrates are shown in Fig. 2. It can be clearly seen that all the Mg–Pd films presented an identical
Fig. 1 The surface morphologies of 200 nm Mg–Pd films prepared at different substrate temperatures.
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Fig. 2 Cross-section SEM images of 200 nm Mg–Pd films prepared at different substrate temperatures: (a) 25 1C, (b) 80 1C, (c) 120 1C and (d) 150 1C.
thickness of 200 nm. Moreover, the cross-section view of Mg–Pd films prepared at room temperature was quite compact with small Mg grains. After increasing the substrate temperatures to 120 1C and 150 1C, the Mg grains grew much larger, and the cross section profiles became quite coarse, giving further evidence of the morphology transformation from the dense to the porous state. 3.2.
Gaseous hydrogen storage properties
The time dependent resistance changes, R/R0, of different Mg–Pd films during hydrogen absorption in 0.1 MPa H2 at 298 K are shown in Fig. 3a, where R0 is the initial resistance of the metallic film. Before hydrogenation, the resistance values of all the samples were very low because they were in the metallic state. During hydrogenation, the resistance values increased immediately due to the formation of a poor conductive phase of MgH2. As shown in Fig. 3a, no changes were detected after inducing H2 for the 25 1C and 80 1C samples in 15 h, indicating that the hydrogen absorption kinetics of Mg–Pd films prepared at low substrate temperatures were extremely sluggish. In sharp contrast, the 120 1C and 150 1C samples exhibited a significant increase in relative resistance changes immediately after exposure to hydrogen. After 6 h, the resistance changes in the 120 1C and 150 1C samples became saturated, suggesting an almost complete
Fig. 3 (a) The time evolution of relative resistance changes (R/R0) of different Mg–Pd films after hydrogenation in 0.1 MPa H2 at 298 K and (b) optical transmittance properties of different Mg–Pd films after hydrogenation in 0.1 MPa H2 at 298 K for 16 h.
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transformation from Mg to MgH2. Compared to other samples, 150 1C films presented the fastest hydrogen absorption kinetics at ambient temperature, clearly demonstrating that the porous structure could significantly improve the hydrogen sorption behaviors of Mg–Pd films. The corresponding optical transmittance properties of different Mg–Pd films after hydrogenation in 0.1 MPa H2 at 298 K for 16 h are shown in Fig. 3b. For the as-prepared samples, their transmittance values were close to zero due to the shiny metallic state. After hydrogen loading, the 120 1C and 150 1C samples transformed from a reflecting shiny state into a transparent state due to the formation of MgH2, which is theoretically a large band gap insulator that is transparent in the visible spectrum. It suggested that a lot of hydrogen was absorbed at ambient temperature for these two samples. However, 25 1C and 80 1C samples remained opaque after hydrogenation in 16 h, with low transmittance changes of about 2%, indicating that negligible hydrogen absorption took place. It should be noted that the 150 1C sample possessed the largest transmittance change value, approximately 12%, confirming the successful transformation from Mg to MgH2, which was consistent with the electric resistance changes discussed above. To review the structural variations of different Mg–Pd films during the hydrogen absorption process, XRD patterns of the as-prepared and hydrogenated samples are shown in Fig. 4. The patterns of all the as-prepared samples were quite similar, consisting of a well-defined Mg(002) peak at B34.61, a broad peak at around 401 of Pd(111), and the Pd(200) peak at 461. After exposure to 0.1 MPa hydrogen at room temperature for 16 h, the Mg(002) peaks of the 80 1C, 120 1C and 150 1C samples disappeared and the corresponding MgH2(110) peaks (2y = 28.41) became significant, suggesting the transformation from the Mg phase to the b-MgH2 phase. But for the 25 1C sample, the Mg(002) peak still existed after 16 h hydrogenation, demonstrating that the hydrogenation kinetics of this sample was much slower than that of other samples and little hydrogen was absorbed. XRD results also showed that the hydrogen sorption properties of 200 nm Mg–Pd films could be dramatically improved by fabricating porous structures, in good agreement with the kinetics results. The hydrogen absorption behaviors of different 200 nm Mg–Pd films during the first cycle at 300 K and 353 K are compared, respectively, as shown in Fig. 5. During the first absorption cycle at 300 K, the 25 1C and 80 1C samples could absorb about 4.5% and 5.5% mass fraction hydrogen in 16 h, separately. For the samples prepared at the substrate temperatures of 120 1C and 150 1C, the absorption kinetics was significantly promoted. Around 7% mass fraction of hydrogen could be absorbed for the 120 1C sample after exposure to H2 for 16 h. It should be mentioned that the hydrogen absorption rate of the 150 1C sample was extremely fast, and it only took about 1 h to achieve the theoretical value of 7.6% mass fraction hydrogen, indicating a complete transformation from Mg to MgH2. When increasing the temperature to 353 K, the hydrogen sorption properties of all the Mg–Pd films improved notably. Approximately 4%, 5% and 6% mass fraction hydrogen could
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Fig. 4 XRD patterns of different 200 nm Mg–Pd films: (a) as-prepared samples and (b) after hydrogenation in 0.1 MPa H2 at 298 K for 16 h.
be absorbed in 1.5 h for the 25 1C, 80 1C and 120 1C samples, respectively. Obviously, the 150 1C sample also presented the most promising sorption behavior, which could absorb the theoretical value of 7.6% mass fraction hydrogen in only 10 min, much faster than that of other samples. In order to further investigate the hydrogen sorption kinetics of different Mg–Pd samples, the hydrogen absorption and desorption cycles of all the samples at 300 K have been studied. Unfortunately, only the 150 1C sample exhibited the expected hydrogenation reversibility, whereas the 25 1C, 80 1C and 120 1C samples could not release any absorbed hydrogen under the same conditions. Fig. 6 illustrates the hydrogen sorption cycles of the 150 1C sample at 300 K. The hydrogen absorption processes were carried out at 0.5 bar H2 pressure, and hydrogen desorption procedures were studied at a backpressure of 10 Pa. The jaggedness of some absorption and desorption curves is due to instrumental noise. The results showed that the 150 1C sample presented reasonable (de)hydrogenation reversibility. No activation period was necessary before reaching the largest hydrogen absorption amount. In the first cycle, only small partially absorbed hydrogen (B1%) could be released during
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Fig. 5 The hydrogen absorption kinetics of different 200 nm Mg–Pd films during the first cycle: (a) at 300 K and (b) at 353 K.
the dehydrogenation process. However, the residual hydrogen could be totally released during the following 1 h evacuation treatment. As a result, about 6% mass fraction hydrogen could still be absorbed in 2 h during the second absorption cycle. Apparently, the maximum hydrogenation capacity decreased with the increasing sorption cycles. After 6 cycles, only around 2% mass fraction hydrogen could be reversibly absorbed and desorbed, which might be ascribed to the volume expansion of Mg films during consecutive hydrogen sorption cycles. Moreover, it is quite amazing that the hydrogen release kinetics was much faster than that of the absorption process, as the dehydrogenation process of Mg based materials is usually extraordinarily sluggish compared to absorption behavior in most of the cases. The hydrogen sorption results indicated that the porous structures of the 150 1C sample could dramatically accelerate the hydrogen sorption kinetics of Mg thin films and improve the reversible hydrogen storage properties. 3.3.
Electrochemical hydrogen storage properties
Fig. 7 shows the cyclic voltammograms (CV) of different 200 nm Mg–Pd films conducted in 6 M KOH solution between 1.2 and 0.45 V versus Hg/HgO electrode. The weak anodic wave coupled
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potentials more negative than 1.1 V.44 The intense anodic peaks between 0.5 V and 0.8 V were observed for all the samples, which were related to the hydrogen desorption behaviors. As indicated in Fig. 7, the hydrogen absorption peaks at around 1.1 V were observed more distinctly for the 120 1C and 150 1C samples, while they were not detected for the 25 1C and 80 1C samples. In addition, the CVs of the 120 1C and 150 1C samples displayed a notable shoulder peak at about 0.8 V before the main H-desorption peak, which was probably attributed to the presence of two types of hydrogenation sites, one located in the extended Mg–Pd region and the other located in the pure Pd layer.39 It should also be noted that the hydrogen desorption potentials shifted to the negative direction; meanwhile, the hydrogen absorption potentials moved to the positive side with an increase in substrate temperatures, suggesting that both the hydrogen absorption and desorption processes could be significantly facilitated by fabricating the porous film structures. The discharge behaviors of different Mg–Pd films at 0.2 mA were systematically investigated, as shown in Fig. 8. Before the discharge process, all the samples were first charged to their fully hydrogenated state using a current of 0.1 mA. Obviously, an activation process was required for all the Mg–Pd samples before attaining their largest discharge capacities. It required
Fig. 6 The hydrogen absorption and desorption cycles of different 200 nm Mg–Pd films at 300 K.
Fig. 7 Cyclic voltammograms of different 200 nm Mg–Pd films in 6 M KOH with a scan rate of 50 mV s 1.
to the cathodic peak located between 0.2 V and 0.3 V were ascribed to a Pd(II) oxide/hydroxide surface formation and reduction process.43 The intense cathodic peaks between 1.0 V and 1.1 V were attributed to the hydrogen absorption process, and the hydrogen evolution reaction (HER) occurred at
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Fig. 8 Discharge capacities (a) and the corresponding capacity retention rates (b) of different 200 nm Mg–Pd films with respect to the cycle number.
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approximately 60 and 70 cycles for the 25 1C and 80 1C samples to reach their maximum discharge capacities, respectively. In comparison, only 35 and 10 activation cycles were needed for the 120 1C and 150 1C samples separately, demonstrating that the activation process of Mg–Pd films could be drastically reduced when their structures became porous. Other than the activation performance, the discharge capacity and the cyclic stability of the electrodes are also very important in the practical application of Ni–MH batteries. As indicated in Fig. 8, the discharge capacities of the 25 1C, 80 1C and 120 1C samples were quite identical, with a maximum discharge capacity of about 500 mA h g 1. The 150 1C sample presented a much larger discharge capacity than that of other samples, the maximum capacity of which was achieved to B720 mA h g 1 after 10 activation cycles. It has to be mentioned that all the Mg–Pd samples exhibited superior cyclic stability, and almost no capacity loss was observed for all the samples. After 100 discharge cycles, more than 70% of their largest discharge capacities could be maintained for all the Mg–Pd films, demonstrating promising applications as anodes of Ni–MH batteries. We all know that Mg will be easily corroded and dissolved in the alkaline solution, so the cyclic properties of Mg-based powder electrodes are quite disappointing. In this study, due to the existence of a Pd cap layer, the anti-corrosive properties of Mg films are significantly improved, leading to promising cyclic stability. The hydrogen diffusion coefficient was also a critical kinetics parameter for hydrogen storage materials. Hagi’s model was applied to analyze the current with respect to the discharge time to calculate the hydrogen diffusion coefficient.45,46 When the discharge time t is long enough (43L2/p2DH), the relationship between the anodic current Id and the hydrogen diffusion coefficient DH can be described as: ln(Id) = p2DH/4d2t + constant, where d is the thickness of the film. The hydrogen diffusion coefficients of different Mg–Pd films during the discharge process can be obtained from the slopes of the curves of ln(Id) as a function of t, as shown in Fig. 9. It can be clearly seen that the hydrogen diffusion coefficient of the 150 1C
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sample was 1.62 10 14 cm2 s 1, which was much larger than that of other films, indicating that the porous network structure of the 150 1C sample was considerably favorable for the hydrogen diffusion through the bulk Mg film. 3.4.
Discussion
According to the results presented above, it could be definitely concluded that the morphologies of Mg–Pd thin films played an important role in affecting their gaseous and electrochemical hydrogen storage properties. Porous thin Mg–Pd films prepared at higher substrate temperatures demonstrated superior gaseous and electrochemical hydrogen storage performances under mild conditions, which might be explained as follows. Hydrogen sorption processes in hydrogen storage materials have characteristics of solid-state transformation, with kinetics of the transformation depending on thermodynamic driving forces and nucleation barriers.29 It is critical for hydrogen storage materials to deliver hydrogen from a gas phase to its interiors to form a hydride phase during the absorption process, and remove hydrogen from a hydride phase to a gas phase during the desorption process. The diffusion of hydrogen in a pure Mg film is really fast, however, its diffusion in Mg hydride is typically slow and is a limiting factor of the transformation. In order to overcome these limitations, measures of increasing the solid/gas surface area and decreasing the hydrogen diffusion paths are very effective. In this paper, porous structures can be formed for the Mg–Pd films prepared at higher deposition temperatures. The porous and network structures significantly increase the contact areas between the solid Mg phase and the hydrogen gas phase, notably facilitating the diffusion process of hydrogen atoms during absorption and desorption processes. In addition, numerous channels and cracks dispersed through the bulk Mg film can be served as ‘‘highways’’ to deliver fastly the hydrogen atoms between the Mg and Mg hydride phases, remarkably improving the hydrogen sorption kinetics and electrochemical hydrogen storage capacities. Furthermore, the bulk Mg films can be successfully divided into many small islands due to the presence of large amounts of channels; thus, the hydrogen diffusion paths are significantly shortened and the small grains of Mg can be quickly charged/discharged before a critical thickness of rate damping MgH2 is reached. Consequently, the porous Mg–Pd films can exhibit the most promising gaseous and electrochemical hydrogen storage behaviors under mild conditions.
4. Conclusions
Fig. 9 The hydrogen diffusion coefficients of different 200 nm Mg–Pd films at 298 K.
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In summary, a series of 200 nm Mg–Pd thin films have been prepared at different substrate temperatures using magnetron sputtering. Their gaseous and electrochemical hydrogen storage properties under mild conditions have been investigated. The results showed that loose and porous film structures could be obtained at higher deposition temperatures. 200 nm Mg–Pd films prepared at 150 1C exhibited the most promising gaseous and electrochemical hydrogen storage properties, including rapid hydrogen sorption kinetics, large hydrogen storage amount, high
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electrochemical discharge capacity, and fast hydrogen diffusion rate. The excellent behaviors of porous Mg–Pd films might be ascribed to the significantly shortened hydrogen diffusion paths and the large contact areas between the hydrogen gas and the solid Mg phases. The results in our study demonstrated that the hydrogen storage properties of thick Mg–Pd films could be improved by fabricating the porous structures without sacrificing the hydrogen storage capacity, which might open up new opportunities for the research and application of thick Mg–Pd films in the future.
Acknowledgements The authors acknowledge the MOST of China (No. 2010CB631301, 2009CB939902 and 2012CBA01207) and the NSFC (No. U1201241 and 51071003). The authors also sincerely acknowledge Prof. L. A. Bendersky and Dr Eric Lass in NIST for their assistance with material measurements.
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Promising gaseous and electrochemical hydrogen storage properties of porous Mg–Pd films under mild conditions Gongbiao Xin,*a Huiping Yuan,a Lijun Jiang,a Shumao Wang,a Xiaopeng Liua and Xingguo Li*b In this paper, the gaseous and electrochemical hydrogen storage properties of 200 nm Mg–Pd thin films with different morphologies have been investigated. The results show that Mg–Pd films become porous
Received 1st April 2015, Accepted 21st April 2015 DOI: 10.1039/c5cp01897h
with the increase of substrate temperature. Porous Mg–Pd films exhibit superior gaseous and electrochemical hydrogen storage behaviors under mild conditions, including rapid hydrogen sorption kinetics, a large hydrogen storage amount, high electrochemical discharge capacity, and a fast hydrogen diffusion rate. The excellent behaviors of porous Mg–Pd films might be ascribed to the significantly shortened hydrogen diffusion paths and the large contact areas between the hydrogen gas and the solid Mg
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phases, which are elucidative for the development and applications of thick Mg–Pd films.
1. Introduction Hydrogen has been considered as a promising energy carrier in future energy systems.1,2 Therefore, various hydrogen storage materials have been investigated in the last few decades.3–10 Among all of these materials, MgH2 has attracted a lot of attention as a solid-state hydrogen storage medium due to its high gravimetric capacity (7.6 wt%), high volumetric density (110 kg m 3), low cost, abundance and encouraging reversibility.11,12 However, the main obstacles limiting the commercial application of MgH2 are its high thermodynamic stability, high dehydrogenation temperature and sluggish hydrogen sorption kinetics.13 Recently, numerous efforts have been devoted to overcome these disadvantages, including the addition of catalysts,14 alloying with other elements,15 and fabricating various nanostructures.16 Compared with these approaches, Mg-based thin films open up a new opportunity to carry out such studies because we can accurately tailor their composition, interface and crystallinity on the nanoscale.17–19 As a result, a lot of work has been conducted to investigate the hydrogen storage properties of Mg-based thin films.20–24 Particularly, the hydrogen storage properties and related applications of binary Mg–Fe, Mg–Ni, Mg–Ti, Mg–La and ternary Mg–Fe–Ti, Mg–Cr–Ti alloy films have been thoroughly a
Department of Energy Materials and Technology, General Research Institute for Nonferrous Metals, Beijing 100088, China. E-mail: [email protected]; Fax: +86-10-82241294; Tel: +86-10-82241241 b Beijing National Laboratory for Molecular Sciences (BNLMS), (The State Key Laboratory of Rare Earth Materials Chemistry and Applications), College of Chemistry and Molecular Engineering, Peking University, Beijing 100871, China. E-mail: [email protected]; Fax: +86-10-62765930; Tel: +86-10-62765930
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investigated by many groups.25–34 However, experiments on Mg-based films reported so far are mostly operated at high temperatures, while it is much more necessary to clarify the hydrogen sorption kinetics under moderate conditions, as the hydrogen storage materials are expected to be applied at ambient temperature.35,36 In our previous work, we have concentrated our interest on improving the hydrogen storage properties of Mg-based thin films under mild conditions, i.e. low temperature and low hydrogen pressure. The results showed that 100 nm Mg–Pd, Pd–Mg–Pd, Mg–Ti–Pd and Mg–Al–Pd films exhibited promising hydrogen sorption properties at room temperature.37–41 However, the hydrogen sorption behaviors will become extremely disappointing when the Mg layer becomes thicker than 100 nm at ambient temperature. Through insertion of 1 nm thin Ti interlayers, the hydrogen storage kinetics of 500 nm thick Mg–Pd films have been successfully promoted.42 Nevertheless, the hydrogen storage capacity of Mg will be inevitably sacrificed due to the addition of heavy weight transition elements. In this study, we have prepared a series of 200 nm Mg–Pd films at different substrate temperatures by magnetron sputtering. The morphology of the Mg–Pd thin films can be controlled by the deposition temperature. The gaseous and electrochemical hydrogen storage properties of 200 nm Mg–Pd films with different structures have been systematically investigated. The results indicated that Mg–Pd films with porous structures exhibited superior gaseous hydrogen sorption kinetics and promising electrochemical hydrogen storage capacities and discharge stability. The results in this study demonstrated that the hydrogen storage properties of thick Mg–Pd films could be remarkably promoted
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by tailoring their structures without sacrificing the hydrogen storage capacity, providing further guidance for the research and application of thick Mg–Pd films.
2. Experimental 2.1.
Film preparation
Pd-capped 200 nm Mg thin films were prepared at four different substrate temperatures using a custom designed direct current (DC) magnetron sputtering system at a background pressure of around 2 10 4 Pa. Prior to deposition, the substrates were heated to 25 1C, 80 1C, 120 1C and 150 1C, respectively. Then, 200 nm Mg layers were first deposited onto Si(110) wafers and glass substrates using an Mg (99.99%) target. Afterwards, a 10 nm Pd layer was coated on top of the Mg layer using a Pd (99.99%) target to protect Mg against oxidation and to catalyze hydrogen dissociation and recombination. The magnetron discharges were generated under an argon pressure of 0.6 Pa with an argon flow rate of 76 sccm. The samples prepared at four different substrate temperatures are designated as 25 1C, 80 1C, 120 1C and 150 1C separately for simplicity. After deposition, the samples were transferred to a vacuum chamber to avoid oxidation. 2.2.
Structural characterization and property measurements
The structures, resistance changes, optical transmittance properties, cyclic voltammogram and diffusion coefficients during the hydrogen absorption and desorption process were systematically determined. The structural changes of different Mg–Pd films were studied by power X-ray diffraction (XRD) (Rigaku D/max-200) using monochromated Cu Ka radiation and a y–2y scan. The cross-section and surface morphologies of different Mg–Pd films were examined using scanning electron microscopy (SEM) measurements (Hitachi S4800). The resistance changes during hydrogen absorption were recorded in a gas loading cell equipped with a four-probe resistance measurement, monitored using an Agilent 34401A digital multimeter. The optical transmission measurements at 298 K were performed using a UV-vis recording spectrophotometer (Shimadzu UV-2401PC) with a dual beam measurement system. Hydrogen volumetric loading/unloading measurements were performed using a Sieverts apparatus developed at the National Institute of Standards and Technology (NIST) with a high precision pressure transducer (0.003% FS) over a relatively wide pressure range.29,30 All hydrogen absorption processes were carried out at 0.5 bar H2 pressure at 300 K and 353 K; the hydrogen desorption procedure was studied at a backpressure of 10 Pa at the same temperatures. After each desorption, the samples were evacuated for another 1 h before the next absorption cycle. Electrochemical experiments were carried out at room temperature in 6 M KOH solution using a three-electrode cell. Platinum foil and Hg/HgO were used as the counter and reference electrodes, respectively. Cyclic voltammogram tests were performed between the potential range of 1.2 and 0.45 V with a scan rate of 50 mV s 1. For the constant current
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discharge measurements, all the electrodes were first fully charged at a current value of 1 mA followed by 10 min relaxation, and then discharged at a current value of 0.2 mA until the cut-off voltage of 0 V (versus the Hg/HgO reference electrode) was reached. The corresponding capacity retention rate (Sn) can be obtained as: Sn(%) = Cn/Cmax 100%, where Cn is the discharge capacity of the electrode at the nth cycle, and Cmax is the maximum discharge capacity. The diffusion coefficients of different Mg–Pd films were determined by the electrochemical multipotential steps in 6 M KOH solution. All the samples were first held at a cathodically polarized potential ( 1.10 V vs. Hg/HgO) for 1.5 h and subsequently switched to an anodically polarized potential ( 0.5 V vs. Hg/HgO) for another 1.5 h.
3. Results and discussion 3.1.
Structural characterization
The surface morphologies of 200 nm Mg–Pd films prepared at different substrate temperatures are examined by SEM, as shown in Fig. 1. A uniform and smooth surface was obtained for the Mg–Pd film prepared at room temperature. The compact film was composed of numerous regular Mg nanocrystals, which generally had a hexagonal shape and the particle size was approximately 100 nm. Significant changes in the surface morphologies were observed when increasing the substrate temperatures. Many pores appeared in the bulk Mg–Pd film deposited at 80 1C, and the regular hexagonal Mg nanocrystals became undistinguishable (Fig. 1b). In the case of samples prepared at 120 1C and 150 1C, the dense and compact Mg–Pd thin films transformed to a porous and puffy state, with numerous pores and cracks through the whole Mg matrix. The formation of porous Mg–Pd thin films might result from the partial melting of Mg particles at higher temperatures, leading to their random aggregation during deposition. The cross-section SEM images of different 200 nm Mg–Pd films deposited on silicon substrates are shown in Fig. 2. It can be clearly seen that all the Mg–Pd films presented an identical
Fig. 1 The surface morphologies of 200 nm Mg–Pd films prepared at different substrate temperatures.
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Fig. 2 Cross-section SEM images of 200 nm Mg–Pd films prepared at different substrate temperatures: (a) 25 1C, (b) 80 1C, (c) 120 1C and (d) 150 1C.
thickness of 200 nm. Moreover, the cross-section view of Mg–Pd films prepared at room temperature was quite compact with small Mg grains. After increasing the substrate temperatures to 120 1C and 150 1C, the Mg grains grew much larger, and the cross section profiles became quite coarse, giving further evidence of the morphology transformation from the dense to the porous state. 3.2.
Gaseous hydrogen storage properties
The time dependent resistance changes, R/R0, of different Mg–Pd films during hydrogen absorption in 0.1 MPa H2 at 298 K are shown in Fig. 3a, where R0 is the initial resistance of the metallic film. Before hydrogenation, the resistance values of all the samples were very low because they were in the metallic state. During hydrogenation, the resistance values increased immediately due to the formation of a poor conductive phase of MgH2. As shown in Fig. 3a, no changes were detected after inducing H2 for the 25 1C and 80 1C samples in 15 h, indicating that the hydrogen absorption kinetics of Mg–Pd films prepared at low substrate temperatures were extremely sluggish. In sharp contrast, the 120 1C and 150 1C samples exhibited a significant increase in relative resistance changes immediately after exposure to hydrogen. After 6 h, the resistance changes in the 120 1C and 150 1C samples became saturated, suggesting an almost complete
Fig. 3 (a) The time evolution of relative resistance changes (R/R0) of different Mg–Pd films after hydrogenation in 0.1 MPa H2 at 298 K and (b) optical transmittance properties of different Mg–Pd films after hydrogenation in 0.1 MPa H2 at 298 K for 16 h.
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transformation from Mg to MgH2. Compared to other samples, 150 1C films presented the fastest hydrogen absorption kinetics at ambient temperature, clearly demonstrating that the porous structure could significantly improve the hydrogen sorption behaviors of Mg–Pd films. The corresponding optical transmittance properties of different Mg–Pd films after hydrogenation in 0.1 MPa H2 at 298 K for 16 h are shown in Fig. 3b. For the as-prepared samples, their transmittance values were close to zero due to the shiny metallic state. After hydrogen loading, the 120 1C and 150 1C samples transformed from a reflecting shiny state into a transparent state due to the formation of MgH2, which is theoretically a large band gap insulator that is transparent in the visible spectrum. It suggested that a lot of hydrogen was absorbed at ambient temperature for these two samples. However, 25 1C and 80 1C samples remained opaque after hydrogenation in 16 h, with low transmittance changes of about 2%, indicating that negligible hydrogen absorption took place. It should be noted that the 150 1C sample possessed the largest transmittance change value, approximately 12%, confirming the successful transformation from Mg to MgH2, which was consistent with the electric resistance changes discussed above. To review the structural variations of different Mg–Pd films during the hydrogen absorption process, XRD patterns of the as-prepared and hydrogenated samples are shown in Fig. 4. The patterns of all the as-prepared samples were quite similar, consisting of a well-defined Mg(002) peak at B34.61, a broad peak at around 401 of Pd(111), and the Pd(200) peak at 461. After exposure to 0.1 MPa hydrogen at room temperature for 16 h, the Mg(002) peaks of the 80 1C, 120 1C and 150 1C samples disappeared and the corresponding MgH2(110) peaks (2y = 28.41) became significant, suggesting the transformation from the Mg phase to the b-MgH2 phase. But for the 25 1C sample, the Mg(002) peak still existed after 16 h hydrogenation, demonstrating that the hydrogenation kinetics of this sample was much slower than that of other samples and little hydrogen was absorbed. XRD results also showed that the hydrogen sorption properties of 200 nm Mg–Pd films could be dramatically improved by fabricating porous structures, in good agreement with the kinetics results. The hydrogen absorption behaviors of different 200 nm Mg–Pd films during the first cycle at 300 K and 353 K are compared, respectively, as shown in Fig. 5. During the first absorption cycle at 300 K, the 25 1C and 80 1C samples could absorb about 4.5% and 5.5% mass fraction hydrogen in 16 h, separately. For the samples prepared at the substrate temperatures of 120 1C and 150 1C, the absorption kinetics was significantly promoted. Around 7% mass fraction of hydrogen could be absorbed for the 120 1C sample after exposure to H2 for 16 h. It should be mentioned that the hydrogen absorption rate of the 150 1C sample was extremely fast, and it only took about 1 h to achieve the theoretical value of 7.6% mass fraction hydrogen, indicating a complete transformation from Mg to MgH2. When increasing the temperature to 353 K, the hydrogen sorption properties of all the Mg–Pd films improved notably. Approximately 4%, 5% and 6% mass fraction hydrogen could
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Fig. 4 XRD patterns of different 200 nm Mg–Pd films: (a) as-prepared samples and (b) after hydrogenation in 0.1 MPa H2 at 298 K for 16 h.
be absorbed in 1.5 h for the 25 1C, 80 1C and 120 1C samples, respectively. Obviously, the 150 1C sample also presented the most promising sorption behavior, which could absorb the theoretical value of 7.6% mass fraction hydrogen in only 10 min, much faster than that of other samples. In order to further investigate the hydrogen sorption kinetics of different Mg–Pd samples, the hydrogen absorption and desorption cycles of all the samples at 300 K have been studied. Unfortunately, only the 150 1C sample exhibited the expected hydrogenation reversibility, whereas the 25 1C, 80 1C and 120 1C samples could not release any absorbed hydrogen under the same conditions. Fig. 6 illustrates the hydrogen sorption cycles of the 150 1C sample at 300 K. The hydrogen absorption processes were carried out at 0.5 bar H2 pressure, and hydrogen desorption procedures were studied at a backpressure of 10 Pa. The jaggedness of some absorption and desorption curves is due to instrumental noise. The results showed that the 150 1C sample presented reasonable (de)hydrogenation reversibility. No activation period was necessary before reaching the largest hydrogen absorption amount. In the first cycle, only small partially absorbed hydrogen (B1%) could be released during
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Fig. 5 The hydrogen absorption kinetics of different 200 nm Mg–Pd films during the first cycle: (a) at 300 K and (b) at 353 K.
the dehydrogenation process. However, the residual hydrogen could be totally released during the following 1 h evacuation treatment. As a result, about 6% mass fraction hydrogen could still be absorbed in 2 h during the second absorption cycle. Apparently, the maximum hydrogenation capacity decreased with the increasing sorption cycles. After 6 cycles, only around 2% mass fraction hydrogen could be reversibly absorbed and desorbed, which might be ascribed to the volume expansion of Mg films during consecutive hydrogen sorption cycles. Moreover, it is quite amazing that the hydrogen release kinetics was much faster than that of the absorption process, as the dehydrogenation process of Mg based materials is usually extraordinarily sluggish compared to absorption behavior in most of the cases. The hydrogen sorption results indicated that the porous structures of the 150 1C sample could dramatically accelerate the hydrogen sorption kinetics of Mg thin films and improve the reversible hydrogen storage properties. 3.3.
Electrochemical hydrogen storage properties
Fig. 7 shows the cyclic voltammograms (CV) of different 200 nm Mg–Pd films conducted in 6 M KOH solution between 1.2 and 0.45 V versus Hg/HgO electrode. The weak anodic wave coupled
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potentials more negative than 1.1 V.44 The intense anodic peaks between 0.5 V and 0.8 V were observed for all the samples, which were related to the hydrogen desorption behaviors. As indicated in Fig. 7, the hydrogen absorption peaks at around 1.1 V were observed more distinctly for the 120 1C and 150 1C samples, while they were not detected for the 25 1C and 80 1C samples. In addition, the CVs of the 120 1C and 150 1C samples displayed a notable shoulder peak at about 0.8 V before the main H-desorption peak, which was probably attributed to the presence of two types of hydrogenation sites, one located in the extended Mg–Pd region and the other located in the pure Pd layer.39 It should also be noted that the hydrogen desorption potentials shifted to the negative direction; meanwhile, the hydrogen absorption potentials moved to the positive side with an increase in substrate temperatures, suggesting that both the hydrogen absorption and desorption processes could be significantly facilitated by fabricating the porous film structures. The discharge behaviors of different Mg–Pd films at 0.2 mA were systematically investigated, as shown in Fig. 8. Before the discharge process, all the samples were first charged to their fully hydrogenated state using a current of 0.1 mA. Obviously, an activation process was required for all the Mg–Pd samples before attaining their largest discharge capacities. It required
Fig. 6 The hydrogen absorption and desorption cycles of different 200 nm Mg–Pd films at 300 K.
Fig. 7 Cyclic voltammograms of different 200 nm Mg–Pd films in 6 M KOH with a scan rate of 50 mV s 1.
to the cathodic peak located between 0.2 V and 0.3 V were ascribed to a Pd(II) oxide/hydroxide surface formation and reduction process.43 The intense cathodic peaks between 1.0 V and 1.1 V were attributed to the hydrogen absorption process, and the hydrogen evolution reaction (HER) occurred at
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Fig. 8 Discharge capacities (a) and the corresponding capacity retention rates (b) of different 200 nm Mg–Pd films with respect to the cycle number.
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approximately 60 and 70 cycles for the 25 1C and 80 1C samples to reach their maximum discharge capacities, respectively. In comparison, only 35 and 10 activation cycles were needed for the 120 1C and 150 1C samples separately, demonstrating that the activation process of Mg–Pd films could be drastically reduced when their structures became porous. Other than the activation performance, the discharge capacity and the cyclic stability of the electrodes are also very important in the practical application of Ni–MH batteries. As indicated in Fig. 8, the discharge capacities of the 25 1C, 80 1C and 120 1C samples were quite identical, with a maximum discharge capacity of about 500 mA h g 1. The 150 1C sample presented a much larger discharge capacity than that of other samples, the maximum capacity of which was achieved to B720 mA h g 1 after 10 activation cycles. It has to be mentioned that all the Mg–Pd samples exhibited superior cyclic stability, and almost no capacity loss was observed for all the samples. After 100 discharge cycles, more than 70% of their largest discharge capacities could be maintained for all the Mg–Pd films, demonstrating promising applications as anodes of Ni–MH batteries. We all know that Mg will be easily corroded and dissolved in the alkaline solution, so the cyclic properties of Mg-based powder electrodes are quite disappointing. In this study, due to the existence of a Pd cap layer, the anti-corrosive properties of Mg films are significantly improved, leading to promising cyclic stability. The hydrogen diffusion coefficient was also a critical kinetics parameter for hydrogen storage materials. Hagi’s model was applied to analyze the current with respect to the discharge time to calculate the hydrogen diffusion coefficient.45,46 When the discharge time t is long enough (43L2/p2DH), the relationship between the anodic current Id and the hydrogen diffusion coefficient DH can be described as: ln(Id) = p2DH/4d2t + constant, where d is the thickness of the film. The hydrogen diffusion coefficients of different Mg–Pd films during the discharge process can be obtained from the slopes of the curves of ln(Id) as a function of t, as shown in Fig. 9. It can be clearly seen that the hydrogen diffusion coefficient of the 150 1C
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sample was 1.62 10 14 cm2 s 1, which was much larger than that of other films, indicating that the porous network structure of the 150 1C sample was considerably favorable for the hydrogen diffusion through the bulk Mg film. 3.4.
Discussion
According to the results presented above, it could be definitely concluded that the morphologies of Mg–Pd thin films played an important role in affecting their gaseous and electrochemical hydrogen storage properties. Porous thin Mg–Pd films prepared at higher substrate temperatures demonstrated superior gaseous and electrochemical hydrogen storage performances under mild conditions, which might be explained as follows. Hydrogen sorption processes in hydrogen storage materials have characteristics of solid-state transformation, with kinetics of the transformation depending on thermodynamic driving forces and nucleation barriers.29 It is critical for hydrogen storage materials to deliver hydrogen from a gas phase to its interiors to form a hydride phase during the absorption process, and remove hydrogen from a hydride phase to a gas phase during the desorption process. The diffusion of hydrogen in a pure Mg film is really fast, however, its diffusion in Mg hydride is typically slow and is a limiting factor of the transformation. In order to overcome these limitations, measures of increasing the solid/gas surface area and decreasing the hydrogen diffusion paths are very effective. In this paper, porous structures can be formed for the Mg–Pd films prepared at higher deposition temperatures. The porous and network structures significantly increase the contact areas between the solid Mg phase and the hydrogen gas phase, notably facilitating the diffusion process of hydrogen atoms during absorption and desorption processes. In addition, numerous channels and cracks dispersed through the bulk Mg film can be served as ‘‘highways’’ to deliver fastly the hydrogen atoms between the Mg and Mg hydride phases, remarkably improving the hydrogen sorption kinetics and electrochemical hydrogen storage capacities. Furthermore, the bulk Mg films can be successfully divided into many small islands due to the presence of large amounts of channels; thus, the hydrogen diffusion paths are significantly shortened and the small grains of Mg can be quickly charged/discharged before a critical thickness of rate damping MgH2 is reached. Consequently, the porous Mg–Pd films can exhibit the most promising gaseous and electrochemical hydrogen storage behaviors under mild conditions.
4. Conclusions
Fig. 9 The hydrogen diffusion coefficients of different 200 nm Mg–Pd films at 298 K.
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In summary, a series of 200 nm Mg–Pd thin films have been prepared at different substrate temperatures using magnetron sputtering. Their gaseous and electrochemical hydrogen storage properties under mild conditions have been investigated. The results showed that loose and porous film structures could be obtained at higher deposition temperatures. 200 nm Mg–Pd films prepared at 150 1C exhibited the most promising gaseous and electrochemical hydrogen storage properties, including rapid hydrogen sorption kinetics, large hydrogen storage amount, high
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electrochemical discharge capacity, and fast hydrogen diffusion rate. The excellent behaviors of porous Mg–Pd films might be ascribed to the significantly shortened hydrogen diffusion paths and the large contact areas between the hydrogen gas and the solid Mg phases. The results in our study demonstrated that the hydrogen storage properties of thick Mg–Pd films could be improved by fabricating the porous structures without sacrificing the hydrogen storage capacity, which might open up new opportunities for the research and application of thick Mg–Pd films in the future.
Acknowledgements The authors acknowledge the MOST of China (No. 2010CB631301, 2009CB939902 and 2012CBA01207) and the NSFC (No. U1201241 and 51071003). The authors also sincerely acknowledge Prof. L. A. Bendersky and Dr Eric Lass in NIST for their assistance with material measurements.
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