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Synchrotron EXAFS studies of Ti-doped Mg2Ni alloy on the cycling behavior
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y x x x ( 2 0 1 4 ) 1 e8
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Synchrotron EXAFS studies of Ti-doped Mg2Ni alloy on the cycling behavior Y.T. Wang a, C.B. Wan a, R.L. Wang b, X.H. Meng a, X. Ju a,* a
Department of Physics, University of Science and Technology Beijing, Beijing, China Electronic Design Automation Center of Chinese Academy of Sciences (WuXi Branch Office), Jiangsu Research and Development Center for Internet of Things, Wuxi 214135, Jiangsu, China
b
article info
abstract
Article history:
In this study, Mg1.9NiTi0.1 alloy was synthesized by mechanical alloying and its cyclic
Received 9 November 2013
hydrogen storage behaviors were investigated. It was found that titanium substituting
Received in revised form
magnesium in Mg1.9NiTi0.1 alloy notably improves the absorption/desorption kinetics.
9 February 2014
Upon cycling, the kinetic rates of absorption/desorption further increase, whereas the
Accepted 10 March 2014
hydrogen storage capacity decreases. To identify the micro-structural evolution of
Available online xxx
Mg1.9NiTi0.1 alloy during cycling, we used X-ray diffraction, scanning electron microscope, and extended X-ray absorption fine structure. After ball milling, the decrease of MgeNi
Keywords:
atomic interaction lowers the stability of Ti-doped phases and has positive effect on the
EXAFS
absorption/desorption behaviors. After 20 cycles, the decrease in the cycling capacity may
XRD
be attributed to the increasing MgNi2 content. Further studies revealed that some amounts
Hydrogen storage
of Mg2Ni transform into MgNi2, which result in great decrease in effective hydrogen storage
Magnesium based
capacity.
Mechanical alloying and milling
Copyright ª 2014, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
Introduction Hydrogen is the ideal candidate as an energy carrier for mobile and stationary applications [1]. The Mg-based alloys possess good quality functional properties, such as heat-resistance, reversibility, and recyclability [2]. Mg2Ni alloy is considered the most popular alloy of this category. Many efforts have been devoted to this alloy to reduce its desorption temperature, enhance the kinetics and cycle life. Ball milling is used to produce amorphous powders in the range of 60e100 mm to create more active sights for hydrogen penetration [3,4]. However, prolonged milling deteriorates the absorption properties of the alloy because of the oxidation on the surface [5].
Several studies have been done to destabilize Mg2Ni by alloying Mg and Ni with transition metal elements such as Ti, Zr, Cr, Cu, Mn, Co and Pd [6e13]. Ti as the most useful additive element shows a great improvement on the hydrogen storage properties of Mg2Ni alloy. Song et al. [14] increased the hydriding and dehydriding rates of Mg2Ni by mechanical alloying of MgH2 with Ni and Ti. Yang et al. [15] showed in a comparative experimental study that Ti is probably the most promising candidate as a dopant to destabilize Mg2Ni. Lin et al. [13] found that the additions of Ti (5, 7.5 and 10 wt%) improves the hydrogen storage properties of Mg2Ni alloy, especially after 15 h ball milling. Moreover, Huang et al. [8] combined the experiments and first-principles calculations to investigate the Mg2xTixNi (x ¼ 0, 0.5) alloys. They found
* Corresponding author. Tel./fax: þ86 10 62333921. E-mail address: [email protected] (X. Ju). http://dx.doi.org/10.1016/j.ijhydene.2014.03.042 0360-3199/Copyright ª 2014, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
Please cite this article in press as: Wang YT, et al., Synchrotron EXAFS studies of Ti-doped Mg2Ni alloy on the cycling behavior, International Journal of Hydrogen Energy (2014), http://dx.doi.org/10.1016/j.ijhydene.2014.03.042
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Experimental details
a
2.5
2.0
Absorbed H wt%
that the most preferable site of Ti substitution in Mg2Ni lattice is Mg(6i) position and TiNi phase is more stable than Ti-doped phase. Extended X-ray absorption fine structure (EXAFS) analysis is a useful approach to study the structural environment around Ni and Ti atoms in Mg2Ni alloy. Vece et al. [16] investigated the Mg2Ni switchable mirror with EXAFS at the Ni K-edge in the virgin, hydrogenated, and dehydrogenated states. Very recently, Zheng et al. [17] reported that Ti doping destabilizes the Mg2NiH4 system by reducing the hydrogenation enthalpy from 64 kJ/(mol H2) to approximately 40 kJ/ (mol H2). They also used EXAFS to investigate the local structure around Ni and Ti atoms in MgeNieTi thin films. However, to the best of our knowledge, the research on the cyclic stability of Ti-doped Mg2Ni has not been reported. As mentioned in a recent review by Jain et al. [18], cyclic stability is one of the major criteria for applicability of metal/metal hydride systems for reversible hydrogen storage. In this paper, Mg1.9NiTi0.1 alloy was synthesized using a high-energy mixer. Its cyclic behavior was investigated using volumetric absorption and desorption measurements. The micro-structural change of Mg1.9NiTi0.1 alloy during cycling was observed by scanning electron microscope (SEM), X-ray diffraction (XRD), and EXAFS.
1.5
1 cycled 4 cycled 7 cycled 10 cycled 13 cycled 17 cycled 20 cycled
1.0
0.5
0.0 0
10
20
30
40
50
60
Time (min)
b
0.0
10 cycled 13 cycled 17 cycled 20 cycled
1 cycled 4 cycled 7 cycled
-0.2
Desorbed H wt%
2
-0.4
-0.6
Elemental Mg, Ni and Ti powders (200 mesh powders with at least 99.9% purity) were mixed and milled under Ar atmosphere to produce Mg1.9NiTi0.1 alloy (in the glove box). The ratio of Mg:Ni:Ti was 2:1:0.1. The ball to powder weight ratio was 10:1. The powders were milled for 15 h by milling for 30 min and resting for 10 min using a high-energy Spex8000 M Mixer/Mill. Volumetric absorption and desorption measurements were performed on a Sieverts-type apparatus. The sample was hydrogenated at 150 C and dehydrogenated at 300 C with an initial hydrogen pressure of 3 MPa and vacuum, respectively. The powder morphologies were observed by a LEO-1450 scanning electron microscope (SEM). XRD analysis was performed on a Philips diffractometer using a Cu-Ka radiation source (l ¼ 1.5406 A). Data were obtained at 2q values of 10 e90 . Phase compositions and lattice parameters for all cycled alloys were calculated by software Jade 5.0. Relative intensity ratio (RIR) method was used to determine the quantitative composition of the samples [19]. EXAFS spectra were recorded for the Ni K-edge (8333 eV) and Ti K-edge (4966 eV) at Station BL14W1 of the SSRF and at station 1W1B of the BSRF, using a Si (1 1 1) double crystal monochromator. For analyzing EXAFS data to evaluate different structural models, we used IFEFFIT software.
Results and discussion Kinetic properties and cycleability Fig. 1 shows the absorption/desorption behaviors of Mg1.9NiTi0.1 alloy. The hydriding and dehydriding processes mostly completed within 1 h. Upon cycling, the sorption rates
-0.8 0
20
40
60
80
100
120
Time (min) Fig. 1 e Hydrogen cycling absorption and desorption behaviors of Mg1.9NiTi0.1 alloy.
markedly improved. After 15 cycles, the sample could hydrogenate 1.7 wt% in less than 5 min. Therefore, the substitution of Ti certainly improved the absorption/desorption kinetics. Normally every material shows an activation period. The absorbed hydrogen amount increased in the initial cycles and reached to 2 wt% after 7 cycles, indicating that the sample was fully activated after the 7th cycle. However, the amount reduced to 1.73 and 1.65 wt% after 15 and 20 cycles, respectively. This may be related to the remaining hydrogen or a deformation of the crystal structure. The hydrogen desorption curves of Mg1.9NiTi0.1 alloy are illustrated in Fig. 1(b). To make the absorbed hydrogen release thoroughly, the sample was desorbed repeatedly. The hydride alloy could release 22% of the absorbed hydrogen at the initial cycles, but 44% of the absorbed hydrogen after 10 cycles. The increased hydrogen content was attributed to the increase of desorption equilibrium pressure. In fact, the equilibrium pressure altered from 0.2 MPa to almost 0.3 MPa based on original data. The original data also suggested that the rate and the equilibrium pressure desorption increased at the same time along with cycles. Above all, Mg1.9NiTi0.1 alloy showed the highest hydrogen capacity at 7th cycle, whereas its hydrogen storage capacity
Please cite this article in press as: Wang YT, et al., Synchrotron EXAFS studies of Ti-doped Mg2Ni alloy on the cycling behavior, International Journal of Hydrogen Energy (2014), http://dx.doi.org/10.1016/j.ijhydene.2014.03.042
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3
slightly degraded after 20 cycles. In contrast, the absorption/ desorption kinetics both notably improved after 20 cycles.
Structure evolution SEM Fig. 2 exhibits the microstructures for Mg1.9NiTi0.1 alloy before and after prolonged cycles. After ball milling, the particles had some small particles on their surfaces, and their surfaces were quite flat. These fine particles and loose structures had positive effects on the absorption/desorption kinetics. The particles after 20 cycles had large agglomerates consisting of small particles and their surfaces were rough. Because surface agglomerates hindered the hydrogen to diffuse in and out of the microstructures, they degraded the hydrogen storage capacity [20].
X-ray diffraction Fig. 3 demonstrates XRD patterns of as-milled, hydrogenated and dehydrogenated samples. The high background and broad peaks revealed that the ball-milled sample was somewhat amorphous. Only Mg2Ni and Ni phases were present, and no expected TiNi, TiNi3 or Ti2Ni phase was detected in this sample [17]. This result strongly suggested that the Ti dopant substituted for Mg or Ni during ball milling. After the sample hydrogenated at 150 C and 3 MPa H2, no monoclinic Mg2Ni was observed; instead, cubic Mg2NiH4 was found. Hence, almost all Mg2Ni transformed to Mg2NiH4 after hydrogenation. The sharp peaks of Mg2NiH4 phase implied that the hydriding process led to a more ordered crystalline structure. Microstructural changes associated with cycling of Mg1.9NiTi0.1 alloy could also be monitored. Previous researchers have proposed that consecutive cycling increases the unit-cell dimensions of alloys [11,21]. This work presented further results on the effect of phase composition on hydrogen storage capacity. All XRD patterns contained the following phases: Mg2Ni (P 6222), Ni, Mg2NiH0.3 (P 6222), Mg2NiH4 (Fm-3m) and NiH (Fm3m). The lattice parameters and the phase compositions of Mg1.9NiTi0.1 alloy were calculated using software Jade 5.0 (Table 1). The observed unit cell parameters agreed well with A; the reference data: Mg2Ni: a ¼ 5.190 and c ¼ 13.220
Fig. 3 e XRD patterns for Mg1.9NiTi0.1 alloy in ball-milled, hydrogenated and dehydrogenated states. (a) Whole XRD patterns, and (b) enlarged part of Mg2Ni, Mg2NiH0.3, Ni and MgNi2 diffraction.
Fig. 2 e SEM micrographs of the Mg1.9NiTi0.1 after ball milling (a) and after 20 hydridingedehydriding cycles (b). Please cite this article in press as: Wang YT, et al., Synchrotron EXAFS studies of Ti-doped Mg2Ni alloy on the cycling behavior, International Journal of Hydrogen Energy (2014), http://dx.doi.org/10.1016/j.ijhydene.2014.03.042
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Table 1 e Weight fraction of different phases and lattice parameters of Mg1.9NiTi0.1 alloys at different states. State of the sample
Phase content (wt%) Mg2Ni
Ball-milled hydrogenated 1 cycle 5 cycles 10 cycles 20 cycles
Mg2NiH4
Mg2NiH0.3
Ni
NiH
6 12
53 27 29 23 16 11
8 9 12 12 12
47 65 62 12
6 6 7
47 60 58
Unit-cell dimensions
MgNi2
Mg2NiH0.3: a ¼ 5.232 and c ¼ 13.404 A; Mg2NiH4: a ¼ 3.507 A; NiH: a ¼ 3.740 A [22,23]. Interestingly, the Ni content was more than 50 wt% in the milled sample, but was 27 wt% in the hydrogenated sample. For the hydrogenated sample, no Mg2Ni phase was found, suggesting that all monoclinic Mg2Ni transformed to cubic Mg2NiH4. After dehydrogenation, the content of Mg2Ni increased to 62 wt%. These results indicated the formation of a more ordered crystalline structure during absorption/desorption cycling. In addition, a magnified image of the selected XRD patterns was presented in Fig. 3(b), which showed the narrower and higher peaks of Mg2Ni phases after the first cycle. Consequently, the milled Mg1.9NiTi0.1 powder had an amorphous structure, and the realistic content was more than the calculated value. The powder became more ordered during heating, similar as an annealing process. Mg2NiH0.3, which has a similar structure as monoclinic Mg2Ni, appeared at the 5th cycle. The content of Mg2NiH0.3 was nearly 50 wt% and gradually increased during cycling. Its lattice constant is a little bigger than Mg2Ni (Table 1). Thus, the unit cell of the sample increased after a few cycles, in agreement with literature [21]. The peak corresponding to Mg2NiH4 was almost invisible, and thus, the hydrogen remained in the form of Mg2NiH0.3 or NiH. Given that the Mg2NiH0.3 content reached to approximately 50 wt% after 5 cycles, it has little effect on cyclic stability. The NiH phases did not participate in the absorption/desorption reactions because they were found both in hydrogenated and dehydrogenated samples. Thus, the remaining hydrogen in the sample did not deteriorate the reversible hydrogen capacity. Notably, MgNi2 (6 wt%) appeared at the 10th cycle, and its content increased to 12 wt% after 20 cycles. MgNi2 generally does not react with H2. Consequently, the decay of hydrogen capacity after 20 cycles may be attributing to the increasing MgNi2 content. Due to the low crystallinity of the samples and low accuracy of aforementioned XRD results, we need another technique to identify the structural variation of Mg1.9NiTi0.1 alloy. For this purpose, we used EXAFS because this technique does not require longrange order in the samples.
EXAFS results Figs. 4 and 5 present the EXAFS spectra of Mg1.9NiTi0.1 alloy. Fig. 4(a) shows the Ti K-absorption edge spectra of Mg1.9NiTi0.1 alloy in the as-milled state, and at the 10th and 20th cycle. The pre-edge structure of Ti disappeared after 15 h of ball milling [13], demonstrating that Ti was in an intermediate valence state in all samples. Since no individual TieNi phase was detected in the XRD results, the Ti dopant substituted for Mg or Ni, and its electronic structure changed during milling.
aMg2 Ni
cMg2 Ni
aMg2 NiH4
aMg2 NiH0:3
cMg2 NiH0:3
aNiH
5.208
13.108 13.271 13.309 13.198 13.173 13.145
6.496 6.500 6.529 6.509 6.512
5.281 5.259 5.259 5.256 5.244
13.27 13.383 13.419 13.379 13.423
3.792 3.789 3.788 3.798 3.772
5.189 5.21 5.217 5.221
Fourier transforms (FTs) is the standard tool used for frequency separation. This operation transforms each sinusoidal A) space. component in an FT modulus peak from k ( A1) to R ( The outcome function is a radial distribution. The interatomic distances can be estimated within an error of 0.01e0.02 A for the first shell of the atoms [24]. Based on Fig. 3, the hexagonal structure of Mg2Ni has space group P 6222 (No. 180) [25]. The input parameters corresponding to the first shell of the considered chemical species (Ni) are listed in Table 2. Figs. 4b and 5a present the FTs of the k3-weighted EXAFS spectra of all Mg1.9NiTi0.1 samples. As seen in Fig. 4(b), the second radial distribution function (RDF) peak was located at approximately 2.5 A (without phase shift corrections). This RDF peak corresponded to the TieNi, TieMg and TieTi path contributions. The first RDF peak, located at about 2 A, corresponded to NieNi, NieMg(6i), NieMg(6f) and NieTi bond distances in the first coordination shell of Ni (indicated by an arrow in Fig. 5(a)). Based on recent reports, the most preferable site of substitution of Ti in Mg2Ni lattice is Mg [26e28]. The best-fit values were also obtained by replacing Mg(6i) atoms with Ti atoms. Figs. 4c and 5b show the experimental and refined k3c (k) EXAFS data. The best fitting results are presented in Table 2. The fit parameters showed that Ti was surrounded by 1.89 Ni atoms at 2.62 A, 0.7 Ti atoms at 2.82 A, and 0.8 Mg atoms at 2.99 A. Comparing with the theoretical values, the local structure of Ti was most similar to that of Mg. On the other hand, the lattice expansion can be attributed to the increase of NieNi and NieMg distances, and the shortest distance was the NieTi bond (2.62 A). In order that the phase expanded, the Ti atoms must replace only the inside Mg atoms of the hcp lattice. Mg(6i) atoms (yellow) occupied the inside of an hcp lattice in the structure (see Fig. 6(a)), thus the Ti atoms most likely substituted the Mg atoms at the position 6i, in agreement with the report of Huang [8]. In addition, the NieTi bond was the strongest bond compared with the NieNi and NieMg bonds. The strong bonding between Ni and Ti weakened the NieMg atomic interactions. As the MgeNi atomic interaction is the most dominant to support the structural framework of unit cell, the decrease of MgeNi atomic interaction lowered the stability of Ti-doped phases. The variation in the interatomic distances of Mg1.9NiTi0.1 alloy during cycling is shown in Fig. 5(c). The NieMg(6i) and NieMg(6f) distances slightly increased after 10 cycles and further increased after 20 cycles. There might be two reasons for the lattice expansion. One is the structural relaxation induced by the hydriding and dehydriding reactions, and another is the formation of Mg2NiH0.3. The difference between NieTi (R ¼ 2.621 A) and TieNi (R ¼ 2.685 A) distances was
Please cite this article in press as: Wang YT, et al., Synchrotron EXAFS studies of Ti-doped Mg2Ni alloy on the cycling behavior, International Journal of Hydrogen Energy (2014), http://dx.doi.org/10.1016/j.ijhydene.2014.03.042
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a
5
1.0
relative intensity (Arb. Unit)
0.8
0.6
ball-milled 10 cycled 20 cycled
0.4
0.2
0.0 4.90
4.95
5.00
5.05
5.10
5.15
5.20
Photon Energy (kev)
b
0.05
Ti-Ni Ti-Ti Ti-Mg
FT(k3·χ(κ))
0.04
ball-milled 10 cycled 20 cycled
0.03
0.02
0.01
0.00 0
2
4
6
8
R(Å)
c
1.5
Exp Cal
1.0
k3 c(k)
0.5 0.0
-0.5 -1.0
-1.5 2
4
6
8
10
-1
k(Å ) Fig. 4 e (a) Ti K-edge spectra for Mg1.9NiTi0.1 alloy at different cycles, (b) Radial distribution functions (RDF), and (c) FFT EXAFS signal and the best fit curve.
Fig. 5 e (a) RDF of NiK-edge as a function of cycle number for Mg1.9NiTi0.1 alloy. (b) FFT EXAFS signal and the best fit curve. (c) The near-neighbor atomic distances.
0.064 A for the ball-milled sample, whereas it amounted to 0.006 A for the 20 cycled sample. This result may be due to a more ordered local environment of Ti atoms after 20 cycles. In addition, the s2 value of TieNi bond decreased with the increase of cycle number, indicating a more ordered structure of
the alloy. After 20 cycles, the nearest neighbor bond-length of Ni slightly decreased (see Fig. 5(a)). Nevertheless, an increase of that of Ti was shown in Fig. 4(b). All the NieMg, NieTi, TieTi and TieMg distances showed a slight increase (Fig. 5(c)). The increasing distances were ascribed to the absorption/
Please cite this article in press as: Wang YT, et al., Synchrotron EXAFS studies of Ti-doped Mg2Ni alloy on the cycling behavior, International Journal of Hydrogen Energy (2014), http://dx.doi.org/10.1016/j.ijhydene.2014.03.042
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Table 2 e Results of Mg1.9NiTi0.1 alloys Ni and Ti K-edges EXAFS fitting. Sample Mg2Ni (theory)
MgNi2 (theory) Mg1.9NiTi0.1
Mg1.9NiTi0.1-10cycled
Mg1.9NiTi0.1-20cycled
Atomic bond
Coordination number
Bond length ( A)
A2) s2 (103
NieNi NieMg(6i) NieMg(6f) NieNi NieMg NieNi NieMg(6i) NieTi NieMg(6f) TieNi TieTi TieMg(6f) NieNi NieMg(6i) NieTi NieMg(6f) TieNi TieTi TieMg(6f) NieNi NieMg(6i) NieTi NieMg(6f) TieNi TieTi TieMg(6f)
2 4 4 6 6 1.74 3.63 0.43 3.93 1.89 0.73 0.87 2.15 3.64 0.41 4.36 1.98 0.80 0.88 2.65 4.42 0.37 6.25 2.85 1.56 0.34
2.60 2.65 2.69 2.42 2.83 2.66 2.70 2.62 2.74 2.69 2.82 2.99 2.59 2.71 2.70 2.73 2.68 2.85 3.06 2.54 2.78 2.80 2.82 2.81 2.90 3.09
e e e e e 1.23 1.55 4.34 1.17 14.75 2.16 4.43 3.43 2.87 0.78 0.98 9.22 4.22 1.43 6.35 2.54 1.75 9.14 4.37 1.92 1.06
desorption process. These results revealed that the reduction in the first RDF peak of Ni was most likely caused by the decrease of the NieNi distance. At the same time, the coordination number of NieNi and NieMg increased. Therefore, the unit cell swelled monotonously in the direction of the c axis, while it shrank monotonously in the band a and b directions (see Fig. 6(b)). The NieNi and NieMg bond distances were close to those of MgNi2 (P63/mmc). Thus, the transformation of Mg2Ni into MgNi2 was most likely taken place after 20 cycles. Since MgNi2 generally does not interact with hydrogen, it has negative effect on the hydrogen capacity. On basis of the crystallographic data of Mg2Ni, Ni should be surrounded by 2 Ni neighbors at 2.60 A, 4 Mg(6i) neighbors at 2.65 A and 4 Mg(6f) neighbors at 2.69 A (see Fig. 6(a) in Ref. [29]). A In the MgNi2 alloy, a Ni is surrounded by 6 Ni atoms at 2.42 and 6 Mg atoms 2.83 A (see Fig. 6(c) in Ref. [30]). For the 10
cycled sample, our fitting found that each Ni atom was surrounded by 2.2 Ni atoms at 2.59 A, 3.64 Mg(6i) atoms at 2.71 A, and 4.36 Mg(6f) atoms at 2.73 A. These values were close to those of crystalline Mg2Ni. Thus, the hydrogenation/dehydrogenation process led to a more ordered lattice, which agreed well with the aforementioned XRD results. After 20 cycles, the NieNi and NieMg interatomic distances altered to 2.54 A and 2.8 A, and the coordination numbers of NieNi and NieMg significantly increased. Consequently, some Mg2Ni phases transformed from P6222 to P63/mmc (Fig. 6). Moreover, the strong bonding between Ni and Ti decreased, and the NieNi atomic interaction increased. Therefore, MgNi2 phases were formed in the 20 cycled sample. The phase variation during cycling is proposed in Fig. 6. Through the absorption and desorption processes, monoclinic Mg2Ni in the Mg1.9NiTi0.1 alloy transformed to cubic Mg2NiH4,
Fig. 6 e Evolution of the crystal structures for Mg1.9NiTi0.1 alloy during cycling. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.) Please cite this article in press as: Wang YT, et al., Synchrotron EXAFS studies of Ti-doped Mg2Ni alloy on the cycling behavior, International Journal of Hydrogen Energy (2014), http://dx.doi.org/10.1016/j.ijhydene.2014.03.042
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and then changed back to monoclinic Mg2Ni. After a few cycles, the volume of crystal lattice increased because of the formation of Mg2NiH0.3. After continuous cycles, Mg2Ni reacted with the extra Ni in the alloy and transformed to MgNi2. Consequently, some Mg2Ni transformed to MgNi2, and its deformation products were no longer functional in the following de-/re-hydrogenation experiments.
Conclusion All samples were characterized with respect to the microstructural transformations during hydridingedehydriding cycles. The following conclusions can be drawn: 1. The hydrogen absorption/desorption properties of Mg1.9NiTi0.1 alloy were noticeably improved by continuous cycles. The hydrogen capacity reached maximum at the 7th cycle, but decreased after 20 cycles. 2. After ball milling, the submission of Mg(6i) by Ti increased the NieNi and NieMg distances. The metastable Ti-doped Mg2Ni phases were formed and substantially enhanced the absorption/desorption kinetics. However, the unit cell swelled monotonously in the vertical direction, while it shrank monotonously in the horizontal directions after consecutive cycles. Such lattice changes resulted in the formation of MgNi2 phases and great decrease in effective hydrogen storage capacity.
Acknowledgments This project was granted financial support from National Natural Science Foundation of China (No. 11087011). The authors want to thank BL14W1 station in SSRF and 1W1B station in BSRF for assistance with extended X-ray absorption fine structure.
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[8] Huang LW, Elkedim O, Nowak M, Chassagnon R, Jurczyk M. Mg2xTixNi (x ¼ 0, 0.5) alloys prepared by mechanical alloying for electrochemical hydrogen storage: experiments and firstprinciples calculations. Int J Hydrogen Energy 2012;37:14248e56. [9] Jain P, Jain A, Vyas D, Verma R, Khan SA, Jain IP. The effects of Ni and Mg2Ni interlayer on hydrogenation properties of Pd sandwiched Mg films. J Alloys Compd 2011;509:2105e10. [10] Zhang YH, Zhao DL, Li BW, Qi Y, Guo SH, Wang XL. Hydrogen storage behaviours of nanocrystalline and amorphous Mg20Ni10xCox(x ¼ 0e4) alloys prepared by melt spinning. Trans Nonferrous Met Soc China 2010;20:405e11. [11] Zhang YH, Li BW, Ma ZH, Guo SH, Qi Y, Wang XL. Improved hydrogen storage behaviours of nanocrystalline and amorphous Mg2Ni-type alloy by Mn substitution for Ni. Int J Hydrogen Energy 2010;35:11966e74. ¨ zdemir G, Baksan B. Electrochemical [12] Anik M, Akay I, O hydrogen storage performance of MgeTieZreNi alloys. Int J Hydrogen Energy 2009;34:9765e72. [13] Lin CK, Wang CK, Lee PY, Lin HC, Lin KM. The effect of Ti additions on the hydrogen absorption properties of mechanically alloyed Mg2Ni powders. Mater Sci Eng A 2007;449:1102e6. [14] Song MY, Kwak YJ, Lee SH, Song J, Mumm DR. Enhancement of hydrogen-storage performance of MgH2 by Mg2Ni formation and hydride-forming Ti addition. Int J Hydrogen Energy 2012;37:18133e9. [15] Yang HB, Yuan HT, Ji JT, Sun H, Zhou ZX, Zhang YS. Characteristics of Mg2Ni0.75M0.25 (M ¼ Ti, Cr, Mn, Fe, Co, Ni, Cu and Zn) alloys after surface treatment. J Alloys Compd 2002;330:640e4. [16] Di Vece M, Van Der Eerden AMJ, Grandjean D, Westerwaal RJ, Lohstroh W, Nikitenko SG, et al. Structure of the Mg2Ni switchable mirror: an EXAFS investigation. Mater Chem Phys 2005;91:1e9. [17] Zheng Q, Pivak Y, Mooij LPA, van der Eerden AMJ, Schreuders H, De Jongh PE, et al. EXAFS investigation of the destabilization of the MgeNieTi (H) system. Int J Hydrogen Energy 2012;37:4161e9. [18] Jain IP, Lal C, Jain A. Hydrogen storage in Mg: a most promising material. Int J Hydrogen Energy 2010;35:5133e44. [19] Hubbard CR, Snyder RL. RIR-measurement and use in quantitative XRD. Powder Diffr 1988;3:74e7. [20] Lin C-K, Huang S-M, Jhang Y-H. Effects of cyclic hydridingedehydriding reactions of Mg2Ni alloy on the expansion deformation of a metal hydride storage vessel. J Alloys Compd 2011;509:7162e7. [21] Song MY. Phase separation of Mg2Ni by hydridingedehydriding cycling. J Alloys Compd 1999;282:297e301. [22] Yvon K, Schefer J, Stucki F. Structural studies of the hydrogen storage material Mg2NiH4. 1. Cubic hightemperature structure. Inorg Chem 1981;20:2776e8. [23] Schubert K, Anderko K. Kristallstruktur von NiMg2 und AuMg2. Naturwissenschaften 1951;38:259. [24] Wan CB, Ju X, Wang YT. EXAFS characterization of TiVCrMn hydrogen storage alloy upon hydrogen absorptionedesorption cycles. Int J Hydrogen Energy 2012;37:990e4. [25] Zolliker P, Yvon K, Jorgensen JD, Rotella FJ. Structural studies of the hydrogen storage material magnesium nickel hydride (Mg2NiH4). 2. Monoclinic low-temperature structure. Inorg Chem 1986;25:3590e3. [26] Zhang J, Huang YN, Mao C, Peng P, Shao YM, Zhou DW. Ab initio calculations on energetics and electronic structures of
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cubic Mg3MNi2 (M ¼ Al, Ti, Mn) hydrogen storage alloys. Int J Hydrogen Energy 2011;36:14477e83. [27] Mustafa A. Improvement of the electrochemical hydrogen storage performance of Mg2Ni by the partial replacements of Mg by Al, Ti and Zr. J Alloys Compd 2009;486:109e14. [28] Lu¨ GL, Chen LH, Wang LB, Yuan HT. Study on the phase composition of Mg2xMxNi (M ¼ Al, Ti) alloys. J Alloys Compd 2001;321:L1e4.
[29] Schefer J, Fischer P, Ha¨lg W, Stucki F, Schlapbach L, Didisheim JJ, et al. New structure results for hydrides and deuterides of the hydrogen storage material Mg2Ni. J LessCommon Met 1980;74:65e73. [30] Komura Y, Tokunaga K. Structural studies of stacking variants in Mg-base Friauf-Laves phases. Acta Crystallogr Sect B Struct Sci 1980;36:1548e54.
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Synchrotron EXAFS studies of Ti-doped Mg2Ni alloy on the cycling behavior Y.T. Wang a, C.B. Wan a, R.L. Wang b, X.H. Meng a, X. Ju a,* a
Department of Physics, University of Science and Technology Beijing, Beijing, China Electronic Design Automation Center of Chinese Academy of Sciences (WuXi Branch Office), Jiangsu Research and Development Center for Internet of Things, Wuxi 214135, Jiangsu, China
b
article info
abstract
Article history:
In this study, Mg1.9NiTi0.1 alloy was synthesized by mechanical alloying and its cyclic
Received 9 November 2013
hydrogen storage behaviors were investigated. It was found that titanium substituting
Received in revised form
magnesium in Mg1.9NiTi0.1 alloy notably improves the absorption/desorption kinetics.
9 February 2014
Upon cycling, the kinetic rates of absorption/desorption further increase, whereas the
Accepted 10 March 2014
hydrogen storage capacity decreases. To identify the micro-structural evolution of
Available online xxx
Mg1.9NiTi0.1 alloy during cycling, we used X-ray diffraction, scanning electron microscope, and extended X-ray absorption fine structure. After ball milling, the decrease of MgeNi
Keywords:
atomic interaction lowers the stability of Ti-doped phases and has positive effect on the
EXAFS
absorption/desorption behaviors. After 20 cycles, the decrease in the cycling capacity may
XRD
be attributed to the increasing MgNi2 content. Further studies revealed that some amounts
Hydrogen storage
of Mg2Ni transform into MgNi2, which result in great decrease in effective hydrogen storage
Magnesium based
capacity.
Mechanical alloying and milling
Copyright ª 2014, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
Introduction Hydrogen is the ideal candidate as an energy carrier for mobile and stationary applications [1]. The Mg-based alloys possess good quality functional properties, such as heat-resistance, reversibility, and recyclability [2]. Mg2Ni alloy is considered the most popular alloy of this category. Many efforts have been devoted to this alloy to reduce its desorption temperature, enhance the kinetics and cycle life. Ball milling is used to produce amorphous powders in the range of 60e100 mm to create more active sights for hydrogen penetration [3,4]. However, prolonged milling deteriorates the absorption properties of the alloy because of the oxidation on the surface [5].
Several studies have been done to destabilize Mg2Ni by alloying Mg and Ni with transition metal elements such as Ti, Zr, Cr, Cu, Mn, Co and Pd [6e13]. Ti as the most useful additive element shows a great improvement on the hydrogen storage properties of Mg2Ni alloy. Song et al. [14] increased the hydriding and dehydriding rates of Mg2Ni by mechanical alloying of MgH2 with Ni and Ti. Yang et al. [15] showed in a comparative experimental study that Ti is probably the most promising candidate as a dopant to destabilize Mg2Ni. Lin et al. [13] found that the additions of Ti (5, 7.5 and 10 wt%) improves the hydrogen storage properties of Mg2Ni alloy, especially after 15 h ball milling. Moreover, Huang et al. [8] combined the experiments and first-principles calculations to investigate the Mg2xTixNi (x ¼ 0, 0.5) alloys. They found
* Corresponding author. Tel./fax: þ86 10 62333921. E-mail address: [email protected] (X. Ju). http://dx.doi.org/10.1016/j.ijhydene.2014.03.042 0360-3199/Copyright ª 2014, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
Please cite this article in press as: Wang YT, et al., Synchrotron EXAFS studies of Ti-doped Mg2Ni alloy on the cycling behavior, International Journal of Hydrogen Energy (2014), http://dx.doi.org/10.1016/j.ijhydene.2014.03.042
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Experimental details
a
2.5
2.0
Absorbed H wt%
that the most preferable site of Ti substitution in Mg2Ni lattice is Mg(6i) position and TiNi phase is more stable than Ti-doped phase. Extended X-ray absorption fine structure (EXAFS) analysis is a useful approach to study the structural environment around Ni and Ti atoms in Mg2Ni alloy. Vece et al. [16] investigated the Mg2Ni switchable mirror with EXAFS at the Ni K-edge in the virgin, hydrogenated, and dehydrogenated states. Very recently, Zheng et al. [17] reported that Ti doping destabilizes the Mg2NiH4 system by reducing the hydrogenation enthalpy from 64 kJ/(mol H2) to approximately 40 kJ/ (mol H2). They also used EXAFS to investigate the local structure around Ni and Ti atoms in MgeNieTi thin films. However, to the best of our knowledge, the research on the cyclic stability of Ti-doped Mg2Ni has not been reported. As mentioned in a recent review by Jain et al. [18], cyclic stability is one of the major criteria for applicability of metal/metal hydride systems for reversible hydrogen storage. In this paper, Mg1.9NiTi0.1 alloy was synthesized using a high-energy mixer. Its cyclic behavior was investigated using volumetric absorption and desorption measurements. The micro-structural change of Mg1.9NiTi0.1 alloy during cycling was observed by scanning electron microscope (SEM), X-ray diffraction (XRD), and EXAFS.
1.5
1 cycled 4 cycled 7 cycled 10 cycled 13 cycled 17 cycled 20 cycled
1.0
0.5
0.0 0
10
20
30
40
50
60
Time (min)
b
0.0
10 cycled 13 cycled 17 cycled 20 cycled
1 cycled 4 cycled 7 cycled
-0.2
Desorbed H wt%
2
-0.4
-0.6
Elemental Mg, Ni and Ti powders (200 mesh powders with at least 99.9% purity) were mixed and milled under Ar atmosphere to produce Mg1.9NiTi0.1 alloy (in the glove box). The ratio of Mg:Ni:Ti was 2:1:0.1. The ball to powder weight ratio was 10:1. The powders were milled for 15 h by milling for 30 min and resting for 10 min using a high-energy Spex8000 M Mixer/Mill. Volumetric absorption and desorption measurements were performed on a Sieverts-type apparatus. The sample was hydrogenated at 150 C and dehydrogenated at 300 C with an initial hydrogen pressure of 3 MPa and vacuum, respectively. The powder morphologies were observed by a LEO-1450 scanning electron microscope (SEM). XRD analysis was performed on a Philips diffractometer using a Cu-Ka radiation source (l ¼ 1.5406 A). Data were obtained at 2q values of 10 e90 . Phase compositions and lattice parameters for all cycled alloys were calculated by software Jade 5.0. Relative intensity ratio (RIR) method was used to determine the quantitative composition of the samples [19]. EXAFS spectra were recorded for the Ni K-edge (8333 eV) and Ti K-edge (4966 eV) at Station BL14W1 of the SSRF and at station 1W1B of the BSRF, using a Si (1 1 1) double crystal monochromator. For analyzing EXAFS data to evaluate different structural models, we used IFEFFIT software.
Results and discussion Kinetic properties and cycleability Fig. 1 shows the absorption/desorption behaviors of Mg1.9NiTi0.1 alloy. The hydriding and dehydriding processes mostly completed within 1 h. Upon cycling, the sorption rates
-0.8 0
20
40
60
80
100
120
Time (min) Fig. 1 e Hydrogen cycling absorption and desorption behaviors of Mg1.9NiTi0.1 alloy.
markedly improved. After 15 cycles, the sample could hydrogenate 1.7 wt% in less than 5 min. Therefore, the substitution of Ti certainly improved the absorption/desorption kinetics. Normally every material shows an activation period. The absorbed hydrogen amount increased in the initial cycles and reached to 2 wt% after 7 cycles, indicating that the sample was fully activated after the 7th cycle. However, the amount reduced to 1.73 and 1.65 wt% after 15 and 20 cycles, respectively. This may be related to the remaining hydrogen or a deformation of the crystal structure. The hydrogen desorption curves of Mg1.9NiTi0.1 alloy are illustrated in Fig. 1(b). To make the absorbed hydrogen release thoroughly, the sample was desorbed repeatedly. The hydride alloy could release 22% of the absorbed hydrogen at the initial cycles, but 44% of the absorbed hydrogen after 10 cycles. The increased hydrogen content was attributed to the increase of desorption equilibrium pressure. In fact, the equilibrium pressure altered from 0.2 MPa to almost 0.3 MPa based on original data. The original data also suggested that the rate and the equilibrium pressure desorption increased at the same time along with cycles. Above all, Mg1.9NiTi0.1 alloy showed the highest hydrogen capacity at 7th cycle, whereas its hydrogen storage capacity
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3
slightly degraded after 20 cycles. In contrast, the absorption/ desorption kinetics both notably improved after 20 cycles.
Structure evolution SEM Fig. 2 exhibits the microstructures for Mg1.9NiTi0.1 alloy before and after prolonged cycles. After ball milling, the particles had some small particles on their surfaces, and their surfaces were quite flat. These fine particles and loose structures had positive effects on the absorption/desorption kinetics. The particles after 20 cycles had large agglomerates consisting of small particles and their surfaces were rough. Because surface agglomerates hindered the hydrogen to diffuse in and out of the microstructures, they degraded the hydrogen storage capacity [20].
X-ray diffraction Fig. 3 demonstrates XRD patterns of as-milled, hydrogenated and dehydrogenated samples. The high background and broad peaks revealed that the ball-milled sample was somewhat amorphous. Only Mg2Ni and Ni phases were present, and no expected TiNi, TiNi3 or Ti2Ni phase was detected in this sample [17]. This result strongly suggested that the Ti dopant substituted for Mg or Ni during ball milling. After the sample hydrogenated at 150 C and 3 MPa H2, no monoclinic Mg2Ni was observed; instead, cubic Mg2NiH4 was found. Hence, almost all Mg2Ni transformed to Mg2NiH4 after hydrogenation. The sharp peaks of Mg2NiH4 phase implied that the hydriding process led to a more ordered crystalline structure. Microstructural changes associated with cycling of Mg1.9NiTi0.1 alloy could also be monitored. Previous researchers have proposed that consecutive cycling increases the unit-cell dimensions of alloys [11,21]. This work presented further results on the effect of phase composition on hydrogen storage capacity. All XRD patterns contained the following phases: Mg2Ni (P 6222), Ni, Mg2NiH0.3 (P 6222), Mg2NiH4 (Fm-3m) and NiH (Fm3m). The lattice parameters and the phase compositions of Mg1.9NiTi0.1 alloy were calculated using software Jade 5.0 (Table 1). The observed unit cell parameters agreed well with A; the reference data: Mg2Ni: a ¼ 5.190 and c ¼ 13.220
Fig. 3 e XRD patterns for Mg1.9NiTi0.1 alloy in ball-milled, hydrogenated and dehydrogenated states. (a) Whole XRD patterns, and (b) enlarged part of Mg2Ni, Mg2NiH0.3, Ni and MgNi2 diffraction.
Fig. 2 e SEM micrographs of the Mg1.9NiTi0.1 after ball milling (a) and after 20 hydridingedehydriding cycles (b). Please cite this article in press as: Wang YT, et al., Synchrotron EXAFS studies of Ti-doped Mg2Ni alloy on the cycling behavior, International Journal of Hydrogen Energy (2014), http://dx.doi.org/10.1016/j.ijhydene.2014.03.042
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Table 1 e Weight fraction of different phases and lattice parameters of Mg1.9NiTi0.1 alloys at different states. State of the sample
Phase content (wt%) Mg2Ni
Ball-milled hydrogenated 1 cycle 5 cycles 10 cycles 20 cycles
Mg2NiH4
Mg2NiH0.3
Ni
NiH
6 12
53 27 29 23 16 11
8 9 12 12 12
47 65 62 12
6 6 7
47 60 58
Unit-cell dimensions
MgNi2
Mg2NiH0.3: a ¼ 5.232 and c ¼ 13.404 A; Mg2NiH4: a ¼ 3.507 A; NiH: a ¼ 3.740 A [22,23]. Interestingly, the Ni content was more than 50 wt% in the milled sample, but was 27 wt% in the hydrogenated sample. For the hydrogenated sample, no Mg2Ni phase was found, suggesting that all monoclinic Mg2Ni transformed to cubic Mg2NiH4. After dehydrogenation, the content of Mg2Ni increased to 62 wt%. These results indicated the formation of a more ordered crystalline structure during absorption/desorption cycling. In addition, a magnified image of the selected XRD patterns was presented in Fig. 3(b), which showed the narrower and higher peaks of Mg2Ni phases after the first cycle. Consequently, the milled Mg1.9NiTi0.1 powder had an amorphous structure, and the realistic content was more than the calculated value. The powder became more ordered during heating, similar as an annealing process. Mg2NiH0.3, which has a similar structure as monoclinic Mg2Ni, appeared at the 5th cycle. The content of Mg2NiH0.3 was nearly 50 wt% and gradually increased during cycling. Its lattice constant is a little bigger than Mg2Ni (Table 1). Thus, the unit cell of the sample increased after a few cycles, in agreement with literature [21]. The peak corresponding to Mg2NiH4 was almost invisible, and thus, the hydrogen remained in the form of Mg2NiH0.3 or NiH. Given that the Mg2NiH0.3 content reached to approximately 50 wt% after 5 cycles, it has little effect on cyclic stability. The NiH phases did not participate in the absorption/desorption reactions because they were found both in hydrogenated and dehydrogenated samples. Thus, the remaining hydrogen in the sample did not deteriorate the reversible hydrogen capacity. Notably, MgNi2 (6 wt%) appeared at the 10th cycle, and its content increased to 12 wt% after 20 cycles. MgNi2 generally does not react with H2. Consequently, the decay of hydrogen capacity after 20 cycles may be attributing to the increasing MgNi2 content. Due to the low crystallinity of the samples and low accuracy of aforementioned XRD results, we need another technique to identify the structural variation of Mg1.9NiTi0.1 alloy. For this purpose, we used EXAFS because this technique does not require longrange order in the samples.
EXAFS results Figs. 4 and 5 present the EXAFS spectra of Mg1.9NiTi0.1 alloy. Fig. 4(a) shows the Ti K-absorption edge spectra of Mg1.9NiTi0.1 alloy in the as-milled state, and at the 10th and 20th cycle. The pre-edge structure of Ti disappeared after 15 h of ball milling [13], demonstrating that Ti was in an intermediate valence state in all samples. Since no individual TieNi phase was detected in the XRD results, the Ti dopant substituted for Mg or Ni, and its electronic structure changed during milling.
aMg2 Ni
cMg2 Ni
aMg2 NiH4
aMg2 NiH0:3
cMg2 NiH0:3
aNiH
5.208
13.108 13.271 13.309 13.198 13.173 13.145
6.496 6.500 6.529 6.509 6.512
5.281 5.259 5.259 5.256 5.244
13.27 13.383 13.419 13.379 13.423
3.792 3.789 3.788 3.798 3.772
5.189 5.21 5.217 5.221
Fourier transforms (FTs) is the standard tool used for frequency separation. This operation transforms each sinusoidal A) space. component in an FT modulus peak from k ( A1) to R ( The outcome function is a radial distribution. The interatomic distances can be estimated within an error of 0.01e0.02 A for the first shell of the atoms [24]. Based on Fig. 3, the hexagonal structure of Mg2Ni has space group P 6222 (No. 180) [25]. The input parameters corresponding to the first shell of the considered chemical species (Ni) are listed in Table 2. Figs. 4b and 5a present the FTs of the k3-weighted EXAFS spectra of all Mg1.9NiTi0.1 samples. As seen in Fig. 4(b), the second radial distribution function (RDF) peak was located at approximately 2.5 A (without phase shift corrections). This RDF peak corresponded to the TieNi, TieMg and TieTi path contributions. The first RDF peak, located at about 2 A, corresponded to NieNi, NieMg(6i), NieMg(6f) and NieTi bond distances in the first coordination shell of Ni (indicated by an arrow in Fig. 5(a)). Based on recent reports, the most preferable site of substitution of Ti in Mg2Ni lattice is Mg [26e28]. The best-fit values were also obtained by replacing Mg(6i) atoms with Ti atoms. Figs. 4c and 5b show the experimental and refined k3c (k) EXAFS data. The best fitting results are presented in Table 2. The fit parameters showed that Ti was surrounded by 1.89 Ni atoms at 2.62 A, 0.7 Ti atoms at 2.82 A, and 0.8 Mg atoms at 2.99 A. Comparing with the theoretical values, the local structure of Ti was most similar to that of Mg. On the other hand, the lattice expansion can be attributed to the increase of NieNi and NieMg distances, and the shortest distance was the NieTi bond (2.62 A). In order that the phase expanded, the Ti atoms must replace only the inside Mg atoms of the hcp lattice. Mg(6i) atoms (yellow) occupied the inside of an hcp lattice in the structure (see Fig. 6(a)), thus the Ti atoms most likely substituted the Mg atoms at the position 6i, in agreement with the report of Huang [8]. In addition, the NieTi bond was the strongest bond compared with the NieNi and NieMg bonds. The strong bonding between Ni and Ti weakened the NieMg atomic interactions. As the MgeNi atomic interaction is the most dominant to support the structural framework of unit cell, the decrease of MgeNi atomic interaction lowered the stability of Ti-doped phases. The variation in the interatomic distances of Mg1.9NiTi0.1 alloy during cycling is shown in Fig. 5(c). The NieMg(6i) and NieMg(6f) distances slightly increased after 10 cycles and further increased after 20 cycles. There might be two reasons for the lattice expansion. One is the structural relaxation induced by the hydriding and dehydriding reactions, and another is the formation of Mg2NiH0.3. The difference between NieTi (R ¼ 2.621 A) and TieNi (R ¼ 2.685 A) distances was
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a
5
1.0
relative intensity (Arb. Unit)
0.8
0.6
ball-milled 10 cycled 20 cycled
0.4
0.2
0.0 4.90
4.95
5.00
5.05
5.10
5.15
5.20
Photon Energy (kev)
b
0.05
Ti-Ni Ti-Ti Ti-Mg
FT(k3·χ(κ))
0.04
ball-milled 10 cycled 20 cycled
0.03
0.02
0.01
0.00 0
2
4
6
8
R(Å)
c
1.5
Exp Cal
1.0
k3 c(k)
0.5 0.0
-0.5 -1.0
-1.5 2
4
6
8
10
-1
k(Å ) Fig. 4 e (a) Ti K-edge spectra for Mg1.9NiTi0.1 alloy at different cycles, (b) Radial distribution functions (RDF), and (c) FFT EXAFS signal and the best fit curve.
Fig. 5 e (a) RDF of NiK-edge as a function of cycle number for Mg1.9NiTi0.1 alloy. (b) FFT EXAFS signal and the best fit curve. (c) The near-neighbor atomic distances.
0.064 A for the ball-milled sample, whereas it amounted to 0.006 A for the 20 cycled sample. This result may be due to a more ordered local environment of Ti atoms after 20 cycles. In addition, the s2 value of TieNi bond decreased with the increase of cycle number, indicating a more ordered structure of
the alloy. After 20 cycles, the nearest neighbor bond-length of Ni slightly decreased (see Fig. 5(a)). Nevertheless, an increase of that of Ti was shown in Fig. 4(b). All the NieMg, NieTi, TieTi and TieMg distances showed a slight increase (Fig. 5(c)). The increasing distances were ascribed to the absorption/
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Table 2 e Results of Mg1.9NiTi0.1 alloys Ni and Ti K-edges EXAFS fitting. Sample Mg2Ni (theory)
MgNi2 (theory) Mg1.9NiTi0.1
Mg1.9NiTi0.1-10cycled
Mg1.9NiTi0.1-20cycled
Atomic bond
Coordination number
Bond length ( A)
A2) s2 (103
NieNi NieMg(6i) NieMg(6f) NieNi NieMg NieNi NieMg(6i) NieTi NieMg(6f) TieNi TieTi TieMg(6f) NieNi NieMg(6i) NieTi NieMg(6f) TieNi TieTi TieMg(6f) NieNi NieMg(6i) NieTi NieMg(6f) TieNi TieTi TieMg(6f)
2 4 4 6 6 1.74 3.63 0.43 3.93 1.89 0.73 0.87 2.15 3.64 0.41 4.36 1.98 0.80 0.88 2.65 4.42 0.37 6.25 2.85 1.56 0.34
2.60 2.65 2.69 2.42 2.83 2.66 2.70 2.62 2.74 2.69 2.82 2.99 2.59 2.71 2.70 2.73 2.68 2.85 3.06 2.54 2.78 2.80 2.82 2.81 2.90 3.09
e e e e e 1.23 1.55 4.34 1.17 14.75 2.16 4.43 3.43 2.87 0.78 0.98 9.22 4.22 1.43 6.35 2.54 1.75 9.14 4.37 1.92 1.06
desorption process. These results revealed that the reduction in the first RDF peak of Ni was most likely caused by the decrease of the NieNi distance. At the same time, the coordination number of NieNi and NieMg increased. Therefore, the unit cell swelled monotonously in the direction of the c axis, while it shrank monotonously in the band a and b directions (see Fig. 6(b)). The NieNi and NieMg bond distances were close to those of MgNi2 (P63/mmc). Thus, the transformation of Mg2Ni into MgNi2 was most likely taken place after 20 cycles. Since MgNi2 generally does not interact with hydrogen, it has negative effect on the hydrogen capacity. On basis of the crystallographic data of Mg2Ni, Ni should be surrounded by 2 Ni neighbors at 2.60 A, 4 Mg(6i) neighbors at 2.65 A and 4 Mg(6f) neighbors at 2.69 A (see Fig. 6(a) in Ref. [29]). A In the MgNi2 alloy, a Ni is surrounded by 6 Ni atoms at 2.42 and 6 Mg atoms 2.83 A (see Fig. 6(c) in Ref. [30]). For the 10
cycled sample, our fitting found that each Ni atom was surrounded by 2.2 Ni atoms at 2.59 A, 3.64 Mg(6i) atoms at 2.71 A, and 4.36 Mg(6f) atoms at 2.73 A. These values were close to those of crystalline Mg2Ni. Thus, the hydrogenation/dehydrogenation process led to a more ordered lattice, which agreed well with the aforementioned XRD results. After 20 cycles, the NieNi and NieMg interatomic distances altered to 2.54 A and 2.8 A, and the coordination numbers of NieNi and NieMg significantly increased. Consequently, some Mg2Ni phases transformed from P6222 to P63/mmc (Fig. 6). Moreover, the strong bonding between Ni and Ti decreased, and the NieNi atomic interaction increased. Therefore, MgNi2 phases were formed in the 20 cycled sample. The phase variation during cycling is proposed in Fig. 6. Through the absorption and desorption processes, monoclinic Mg2Ni in the Mg1.9NiTi0.1 alloy transformed to cubic Mg2NiH4,
Fig. 6 e Evolution of the crystal structures for Mg1.9NiTi0.1 alloy during cycling. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.) Please cite this article in press as: Wang YT, et al., Synchrotron EXAFS studies of Ti-doped Mg2Ni alloy on the cycling behavior, International Journal of Hydrogen Energy (2014), http://dx.doi.org/10.1016/j.ijhydene.2014.03.042
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and then changed back to monoclinic Mg2Ni. After a few cycles, the volume of crystal lattice increased because of the formation of Mg2NiH0.3. After continuous cycles, Mg2Ni reacted with the extra Ni in the alloy and transformed to MgNi2. Consequently, some Mg2Ni transformed to MgNi2, and its deformation products were no longer functional in the following de-/re-hydrogenation experiments.
Conclusion All samples were characterized with respect to the microstructural transformations during hydridingedehydriding cycles. The following conclusions can be drawn: 1. The hydrogen absorption/desorption properties of Mg1.9NiTi0.1 alloy were noticeably improved by continuous cycles. The hydrogen capacity reached maximum at the 7th cycle, but decreased after 20 cycles. 2. After ball milling, the submission of Mg(6i) by Ti increased the NieNi and NieMg distances. The metastable Ti-doped Mg2Ni phases were formed and substantially enhanced the absorption/desorption kinetics. However, the unit cell swelled monotonously in the vertical direction, while it shrank monotonously in the horizontal directions after consecutive cycles. Such lattice changes resulted in the formation of MgNi2 phases and great decrease in effective hydrogen storage capacity.
Acknowledgments This project was granted financial support from National Natural Science Foundation of China (No. 11087011). The authors want to thank BL14W1 station in SSRF and 1W1B station in BSRF for assistance with extended X-ray absorption fine structure.
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Please cite this article in press as: Wang YT, et al., Synchrotron EXAFS studies of Ti-doped Mg2Ni alloy on the cycling behavior, International Journal of Hydrogen Energy (2014), http://dx.doi.org/10.1016/j.ijhydene.2014.03.042