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Structural, hydrogen storage, and electrochemical properties of Laves phase-related body-centered-cubic solid solution metal hydride alloys
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 e1 1
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Structural, hydrogen storage, and electrochemical properties of Laves phase-related body-centeredcubic solid solution metal hydride alloys K. Young a,*, J. Nei a, D.F. Wong a,b, L. Wang a,b a b
BASF/Battery Materials-Ovonic, 2983 Waterview Drive, Rochester Hills, MI 48309, USA Department of Chemical Engineering and Materials Science, Wayne State University, MI 48202, USA
article info
abstract
Article history:
Structure, gaseous phase hydrogen storage, and electrochemical properties of a series of
Received 13 November 2013
Laves phase-related BCC solid solution metal hydride alloys with BCC/C14 ratios ranging
Received in revised form
from 0.09 to 8.52 were studied. Some properties are correlated to the phase abundance and
2 January 2014
V-content in the alloy with monotonic evolutions, for example, lattice constant, phase
Accepted 22 January 2014
abundance, and hydrogen storage pressure. Other properties such as gaseous phase ca-
Available online xxx
pacities, PCT hysteresis, high-rate dischargeability, and bulk hydrogen diffusion correlate better with the C14 phase crystallite size, which are considered to be more related to the
Keywords:
synergetic effect between main and secondary phases. In contrast with conventional metal
Hydrogen absorbing alloys
hydride alloys used in NiMH batteries, the electrochemical discharge capacities of these
Metal hydride electrode
alloys are not between the maximum and the reversible hydrogen storage measured in the
Laves phase alloys
gaseous phase. The current study’s alloys have electrochemical capacities that are
Body-centered-cubic alloys
insensitive to composition but have much room for improvement, with high-rate dis-
Phase segregation
chargeabilities that are superior compared to other commercially available alloys. With further research, these alloys show potential for high-rate battery applications. Copyright ª 2014, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
Introduction Iba and Akiba described the “Laves phase-related bodycentered-cubic (BCC) solid solution” as a family of metal hydride (MH) alloys with a two-phase microstructure composed of a BCC phase and a Laves phase (mostly C14) and a general formula of ABx, where A is an element or any combination of elements from Group 4A (mostly Ti), B is from Group 5A, 6A, or 7A (mainly V), and x is between 1 and 6 [1,2]. The BCC phase, with its higher V-content, has higher melting temperature and solidifies before the Laves phases (C14, for example) do. This
phenomenon produces the following evolution in microstructure with the increase in C14/BCC ratio (Fig. 5 in Ref. [3]): C14 appears at the grain boundaries / sections of C14 phase connect to form a three-dimensional network / the alloy matrix shifts to C14, and BCC phase forms its own threedimensional network, a cross-sectional view showing embedded, isolated BCC islands / the C14 matrix continues to expand, and BCC phase forms fish-bone-like inclusions within the matrix. The high density of phase boundaries makes it possible to combine the advantages of BCC (high hydrogen storage capability [4,5]) and C14 phases (good absorption kinetics [6], easy formation due to its brittleness
* Corresponding author. Tel.: þ1 248 293 7000; fax: þ1 248 299 4520. E-mail addresses: [email protected], [email protected] (K. Young). 0360-3199/$ e see front matter Copyright ª 2014, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijhydene.2014.01.134
Please cite this article in press as: Young K, et al., Structural, hydrogen storage, and electrochemical properties of Laves phaserelated body-centered-cubic solid solution metal hydride alloys, International Journal of Hydrogen Energy (2014), http:// dx.doi.org/10.1016/j.ijhydene.2014.01.134
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[7e9], and high surface catalytic activity [10]) by increasing the synergetic effect between the two phases. The density of phase boundaries also promotes the formation of coherent and catalytic interfaces between the boundaries, improving hydrogen absorption [11]. Similar two-phase microstructures involving BCC phase are also seen in other systems, for example, C15/BCC in Ni-free alloys [12e14] and TiNi/BCC in TieVeNi ternary alloys [15e17]. Both A-site (Ti) and B-site (V) substitution studies were performed on a number of mixed phase alloys, evaluating alloy properties for gaseous phase hydrogen storage applications. The A-site substitutions Zr [18e21] and Hf [8] were found to increase C14 phase abundance significantly, while Y only increased it slightly [22] and La promoted C15 phase [23]. The B-site substitutions Mn [19], Fe [24e26], Ni [27], Si [28,29], Nb [30,31], Ta [32], and B [33] promoted C14 phase, and Al [34], Cr [27], Ce [35], Mo [32,36], and W [32] had the opposite effect. Soon after the gaseous phase hydrogen storage characteristics of this alloy family were reported in 1995 [7,18], its electrochemical properties and potential for nickel/metal hydride (Ni/MH) battery applications attracted researchers’ attention [21,27,37e42]. In particular, researchers from three institutes in China have made continuous efforts in improving the gaseous and electrochemical hydrogen storage properties of this family of alloys. Through a series of composition refinements [35,36,43e47], Wang, Pan, and their coworkers at Zhejiang University obtained an alloy formula (Ti0.8Zr0.2)(V0.5Mn0.1Cr0.16Ni0.2Fe0.04)5 with superb high-rate dischargeability (HRD) performance. Zhao and his coworkers at Chinese Academy of Science in Changchun worked on AB3 [48e52] and AB2 [53e55] alloys and discovered the formulas and (Ti0.79Zr0.21)(V0.36 (Ti0.68Zr0.32)(V0.47Cr0.13Ni0.33Mn0.07)3 Mn0.15Ni0.42Cr0.07)2 to have good balance among various electrochemical performance properties. Yu and his coworkers from Chinese Academy of Science in Shanghai reported that gaseous phase storage capacity and kinetics can be improved by reducing the amount of C14 phase; however, the alloys become more difficult to activate [5,56,57]. Although many studies focusing on this family of alloys have been performed in the past, the studies offer somewhat mixed conclusions. For example, while Ovonic used BCC as a
secondary phase in a C14-based alloy to protect the surface from alkaline corrosion [40], other groups used C14 as a secondary phase to improve activation and charge-transfer properties. Furthermore, while Chai et al. claimed a BCC:C14 abundance ratio of 1:1 yielded optimized capacity results [51], Gue´guen et al. showed that only a small amount of C14 is needed to improve gaseous phase storage [58]. Reviewing the previous studies, one cannot find an optimized C14 abundance for electrochemical performance. Therefore, a systematic study of alloys with a broad range of BCC/C14 ratio is of interest to researchers in the field and presented in this paper.
Experimental setup Arc melting was performed under continuous argon flow with a non-consumable tungsten electrode and a watercooled copper tray. Before each run, a piece of sacrificial titanium underwent a few meltingecooling cycles to reduce the residual oxygen concentration in the system. Each 12-g ingot was re-melted and turned over several times to ensure uniformity in chemical composition. Chemical composition for the ingot was analyzed using a Varian Liberty 100 inductively coupled plasma optical emission spectrometer (ICP-OES). A Philips X’Pert Pro X-ray diffractometer (XRD) was used to study the microstructure, and a JEOLJSM6320F scanning electron microscope (SEM) with energy dispersive spectroscopy (EDS) capability was used to study the phase distribution and composition. Gaseous phase hydrogen storage characteristics for each sample were measured using a Suzuki-Shokan multi-channel pressureeconcentrationetemperature (PCT) system. In the PCT analysis, each sample (a single piece of ingot with newly cleaved surface and a weight of about 2 g) was first activated by a 2-h thermal cycle between 300 C and room temperature at 2.5 MPa H2 pressure. PCT isotherms at 30, 60, and 90 C were then measured. Details of electrode preparations as well as the measurement methods have been reported previously [59,60].
Table 1 e Design compositions (in bold) and ICP results in at.%. Alloy P1 P2 P3 P4 P5 P6 P7 P8
Design ICP Design ICP Design ICP Design ICP Design ICP Design ICP Design ICP Design ICP
Ti
Zr
V
Cr
Mn
Fe
Co
Ni
Al
B/A ratio
x Value in (1)
14.7 15.1 14.5 14.7 14.3 14.7 14.1 14.1 13.9 14.0 13.8 14.0 13.7 14.2 13.6 13.8
13.8 13.9 11.1 11.0 8.9 8.8 7.0 6.9 5.5 5.3 4.2 4.3 3.1 3.0 2.1 2.0
20.0 20.2 25.6 25.4 30.1 31.0 33.8 34.1 37.0 36.9 39.7 38.6 42.0 41.3 44.0 43.1
6.0 3.8 7.7 7.3 9.0 7.5 10.1 10.0 11.1 11.5 11.9 11.4 12.6 12.1 13.2 13.1
12.6 12.8 11.2 10.9 10.2 9.8 9.3 9.0 8.6 8.1 7.9 8.1 7.4 7.0 6.9 6.5
1.2 1.3 1.6 1.6 1.9 1.9 2.1 2.1 2.3 2.3 2.5 2.7 2.6 2.7 2.7 2.8
2.6 2.7 2.3 2.4 2.1 2.1 1.9 1.9 1.8 1.8 1.6 1.8 1.5 1.6 1.4 1.4
28.5 29.5 25.6 26.2 23.1 23.7 21.1 21.4 19.5 19.7 18.0 18.6 16.8 17.6 15.7 16.7
0.6 0.7 0.5 0.5 0.5 0.5 0.4 0.5 0.4 0.5 0.4 0.5 0.3 0.5 0.3 0.6
2.50 2.45 2.91 2.89 3.32 3.26 3.73 3.76 4.14 4.19 4.55 4.46 4.96 4.81 5.37 5.33
0.7 1.0 1.3 1.6 1.9 2.2 2.5 2.8
Please cite this article in press as: Young K, et al., Structural, hydrogen storage, and electrochemical properties of Laves phaserelated body-centered-cubic solid solution metal hydride alloys, International Journal of Hydrogen Energy (2014), http:// dx.doi.org/10.1016/j.ijhydene.2014.01.134
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8.100
2.980
8.075
2.975
BCC, a
8.050 8.025
2.970 2.965
C14, c
8.000
2.960
7.975
2.955
4.950
2.950
4.925
2.945
C14, a
4.900 1
2
3
BCC Lattice Constant (Å)
C14 Lattice Constants (Å)
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2.940 4
5
6
7
8
Alloy Number
Fig. 2 e Lattice constants a and c from C14 phase and a from BCC phase as functions of alloy number.
Fig. 1 e XRD patterns using Cu-Ka as the radiation source for alloys P1 (a), P2 (b), P3 (c), P4 (d), P5 (e), P6 (f), P7 (g), and P8 (h). Vertical lines are used to indicate shifts in both C14 and BCC peaks.
Results and discussion Alloy composition Eight alloys with a general formula of
Ti0.4þx/6Zr0.6x/6Mn0.44Ni1.0Al0.02Co0.09(VCr0.3Fe0.063)x, x ¼ 0.7 to 2.8
(1)
were prepared by arc melting for this study. Design compositions of these alloys are listed in Table 1, and the design principle is described as follows. The alloy formula with x ¼ 2.5 is close to the formula developed by Pan et al. [46], and the extra Al and Co at the optimized ratio is added to improve
C14 electrochemical properties [60,61]. While the extra Al and Co can improve C14 electrochemical properties, they can also affect BCC properties. Al is known to expand the BCC unit cell; however, it lowers capacity [62]. Co suppresses pulverization, improves the cycle stability and HRD in a BCC-based alloy, but reduces capacity [45]. Therefore, there is a trade-off in the amounts of Al and Co. In this study, we reduce the amount of Co and Al while increasing V, when the main phase shifts from C14 into BCC. The alloy formula with x ¼ 0.7 is close to Ovonic’s AB2 formula developed in the 1980s [39]. With the increase in Vcontent (x), the Ti- and Zr-contents have to be altered accordingly to maintain the metal-to-hydrogen (MeH) bond strength in the C14 phase. Assuming the BCC/C14 ratio is approximately 1:1 and that one-fifth of the V-content in BCC phase tends to appear in the C14 phase [45], the amount of V that goes into C14 phase is about one-sixth of the total amount. According to the results from previous experiments, the target decrease in Zr (weaker MeH bond) at the expense of Ti is roughly the same as the amount of increase in V (stronger MeH bond) [63] in the C14 phase. Therefore, onesixth of x is subtracted from the Zr-content and added to the Ti-content to maintain the same AB2 stoichiometry and MeH bond strength. Pd is also a possible modifier [49,54] but is too expensive to be considered for practical application.
Table 2 e Lattice constants, unit cell volumes, phase abundances, and crystallite sizes of C14 and BCC phases of alloys P1eP8 from XRD analysis. FWHM and XS denote full-width at half maximum and crystallite size, respectively. C14, a ( A) C14, c ( A) C14, a/c C14 unit cell C14 wt.% C14 (103) C14 XS ( A) BCC, a ( A) BCC wt.% BCC (200) BCC XS ( A) volume ( A3) FWHM FWHM P1 P2 P3 P4 P5 P6 P7 P8
4.9525 4.9386 4.9292 4.9138 4.9165 4.9078 4.9064 4.9025
8.0690 8.0470 8.0306 8.0040 8.0044 7.9897 7.9835 7.9728
0.6138 0.6137 0.6138 0.6139 0.6142 0.6143 0.6146 0.6149
171.40 169.97 168.98 167.37 167.56 166.66 166.44 165.95
91.6 85.9 80.4 72.1 64.7 53.5 45.8 10.5
0.457 0.376 0.351 0.316 0.319 0.405 0.363 0.427
207 262 285 325 321 289 273 225
2.9568 2.9576 2.9601 2.9621 2.9633 2.9649 2.9655 2.9664
8.4 14.1 19.6 27.9 35.3 46.5 54.2 89.5
e 0.197 0.365 0.391 0.464 0.648 0.604 0.689
e 714 298 274 224 174 167 144
Please cite this article in press as: Young K, et al., Structural, hydrogen storage, and electrochemical properties of Laves phaserelated body-centered-cubic solid solution metal hydride alloys, International Journal of Hydrogen Energy (2014), http:// dx.doi.org/10.1016/j.ijhydene.2014.01.134
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ICP results from the ingot samples are listed in Table 1. Other than some uneven distribution of Cr in the alloys with lower Cr-content (alloys P1 and P3), ICP results are very close to the target compositions. The measured B/A ratio varies from 2.45 to 5.33 and is within 3% of the design value.
XRD structure analysis XRD patterns of the eight alloys are shown in Fig. 1. Two sets of diffraction peaks, C14 and BCC, are observed. As the Vcontent in the alloy increases, the BCC peaks increase in intensity and shift to lower angles, while the C14 peaks behave the opposite. Lattice constants of both phases calculated from the XRD patterns are listed in Table 2 and plotted in Fig. 2. As the amount of V increases, both a and c in C14 phase decrease, while a in BCC phase increases. In C14 phase, the decrease in lattice parameter c is faster than the decrease in a, which
results in an increasing a/c aspect ratio with increasing Vcontent in the alloys. Therefore, a lower pulverization rate during hydride/dehydride cycling is expected from alloys with higher V-content in this study [64e66]. Lattice parameter a in BCC phase increases from 2.9568 to 2.9664 A as the V-content in the alloy increases, which is smaller than the optimized value of 3.042 A corresponding to a maximized hydrogen storage capacity [67], leaving room for potential improvement in storage capacity in future studies. Phase abundance and crystallite sizes of each phase are listed in Table 2. These values were obtained from full pattern fitting of the XRD data using the Rietveld method and Jade 9 software. As the V-content in the alloy increases, the abundance of C14 phase decreases monotonically from 91.6 to 10.6 wt.% while the abundance of BCC phase increases. At the same time, the C14 crystallite size first increases and then decreases, while the BCC crystallite size decreases
Fig. 3 e SEM back-scattering electron images for alloys P1 (a), P2 (b), P3 (c), P4 (d), P5 (e), P6 (f), P7 (g), and P8 (h). Chemical compositions in the numbered areas measured by EDS are listed in Table 3. Please cite this article in press as: Young K, et al., Structural, hydrogen storage, and electrochemical properties of Laves phaserelated body-centered-cubic solid solution metal hydride alloys, International Journal of Hydrogen Energy (2014), http:// dx.doi.org/10.1016/j.ijhydene.2014.01.134
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Table 3 e Summary of EDS results. All compositions are in at.%. Compositions of main AB2 phase are in bold. Fig. 3a-1 2 3 4 5 6 Fig. 3b-1 2 3 4 Fig. 3c-1 2 3 4 Fig. 3d-1 2 3 4 Fig. 3e-1 2 3 4 5 Fig. 3f-1 2 3 4 Fig. 3g-1 2 3 4 Fig. 3h-1 2 3 4 5
Ti
Zr
V
Cr
Mn
Fe
Co
Ni
Al
B/A
Phase
12.8 13.9 17.8 19.5 2.1 5.8 15.5 18.6 5.1 6.9 18.0 20.9 5.6 9.2 18.1 20.4 4.7 0.8 20.2 22.1 32.4 5.0 0.9 22.0 29.7 5.5 2.3 23.1 35.5 5.8 3.4 22.4 36.3 35.1 5.7 1.6
17.1 15.3 10.8 9.1 87.4 1.6 15.1 11.9 0.5 66.3 13.1 10.9 0.3 59.5 13.3 10.5 0.1 94.1 11.0 9.3 4.6 0.0 95.9 9.2 3.3 0.1 90.2 8.3 3.3 0.0 87.4 9.0 2.4 8.2 0.1 93.1
20.5 20.6 20.7 19.1 3.0 56.6 18.6 18.2 58.8 8.5 20.1 19.2 58.0 10.4 19.4 19.2 58.6 2.0 18.7 19.4 10.0 58.9 1.3 20.7 17.6 58.5 3.6 20.0 9.0 59.4 4.1 19.5 8.8 16.1 58.7 2.8
5.0 4.3 2.8 2.2 0.9 9.1 5.6 4.3 15.2 2.4 4.1 3.3 16.7 2.0 5.0 4.2 20.2 0.5 5.0 4.1 1.4 20.0 0.3 4.5 4.2 19.6 1.0 4.3 1.2 18.3 1.1 2.4 1.3 1.4 19.2 1.0
15.0 13.4 11.0 9.7 1.8 14.9 11.4 9.3 12.7 4.4 9.4 7.8 10.8 4.5 9.6 8.1 9.0 0.5 9.0 7.1 3.9 8.4 0.3 7.9 5.0 7.7 0.7 7.3 3.2 7.4 0.9 7.1 3.9 4.5 6.9 0.3
1.8 1.6 1.0 0.8 0.3 1.1 2.1 1.3 1.6 0.6 2.0 1.6 2.0 0.8 2.9 1.9 1.8 0.2 3.3 2.3 1.1 2.3 0.0 3.2 1.8 2.3 0.2 3.6 1.8 2.8 0.2 2.8 3.3 2.0 2.7 0.2
2.9 2.9 2.5 2.1 0.4 1.3 2.9 2.5 1.1 0.9 2.7 2.1 1.0 1.0 2.8 2.4 0.9 0.3 2.9 2.2 2.2 0.8 0.0 2.5 2.3 0.9 0.1 2.8 2.8 1.1 0.1 1.8 3.6 1.4 1.0 0.1
24.2 27.3 32.7 36.9 3.9 9.4 28.2 33.2 4.8 9.9 29.8 33.5 5.2 12.2 28.3 32.5 4.5 1.6 29.0 32.9 43.4 4.5 1.1 29.4 35.3 5.0 1.9 30.0 42.3 5.0 2.7 33.8 39.0 30.8 5.3 0.9
0.6 0.7 0.7 0.7 0.2 0.3 0.6 0.7 0.3 0.1 0.7 0.7 0.3 0.2 0.7 0.7 0.2 0.1 0.8 0.7 1.0 0.2 0.1 0.6 0.8 0.4 0.1 0.7 0.9 0.2 0.1 1.0 1.3 0.5 0.5 0.1
2.34 2.42 2.50 2.50 0.12 12.53 2.27 2.28 16.88 0.37 2.21 2.14 15.93 0.45 2.19 2.23 19.83 0.05 2.20 2.19 1.70 19.02 0.03 2.21 2.03 16.86 0.08 2.19 1.58 16.24 0.10 2.18 1.58 1.31 16.26 0.06
AB2 AB2 AB2 AB2 ZrO2 BCC AB2 AB2 BCC ZrO2 AB2 AB2 BCC Zr AB2 AB2 BCC Zr AB2 AB2 ZrxNiy BCC Zr AB2 AB2 BCC Zr AB2 ZrxNiy BCC Zr AB2 ZrxNiy ZrxNiy BCC Zr
monotonically. The measured BCC/C14 ratio increases from 0.09 to 8.52 at a more dramatic rate compared to the design ratio (about 0.24e0.57).
SEM/EDS microstructure analysis Microstructures for this series of alloys were studied using SEM, and the back-scattering electron images (BEI) are presented in Fig. 3. Chemical compositions of several areas with different contrasts were studied by EDS analysis, and the results are summarized in Table 3. During cooling, BCC phase with high V-content solidifies first to form a threedimensional framework while the rest of the liquid solidifies into Laves phase as the alloy cools further. Darker regions in the SEM micrographs are BCC phase (lower average atomic mass due to lower Zr-content) and represent the intersection of a plane and a three-dimensional framework with three major axes perpendicular to each other (Fig. 4). In Fig. 3a from Alloy P1, the BCC phase starts to appear as small dark inclusions (Fig. 3a spot 6, abbrev. as 3a-6) of about 2e3 mm in size. The matrix is composed of AB2 phases with composition variation mainly in Ti/Zr/Ni-contents (3a-1 to 3a-4). Besides, occasional ZrO2 inclusion can be also found (3a-5). As the V-
content in the alloy increases, BCC phase expands, and its shape evolves from dots to fish-bones and then to islands. Since the designed B/A ratios in these alloys (column 12 in Table 1) are well above 2.5 (the approximate maximum solubility of AB2 phase on the B-rich side [68]) and the BCC phase contains almost no Zr, the extra Zr is solidified individually. With the increase in V-content in the alloy, the extra Zr changes its form of existence from ZrO2 (3a-5 and 3b-4) to metallic Zr (3c-4, 3d-4, 3e-5, 3f-4, 3g-4, and 3h-5). In contrast, there are cases where metallic Zr-inclusions found in Si-free alloys are replaced by occasional dark spots of ZrO2 in Sisubstituted alloys [69]. V is a better oxygen salvager compared to Zr, and the oxide for V is light and floats on the surface to form slag, which separates out from the liquid by the tundish. As such, the extra Zr is less oxidized with higher V-content in the alloy. In alloys P2eP4, the AB2 matrix is more homogeneous in composition when compared to alloy P1. For alloys with higher V-content (alloys P5, P7, and P8), an intergrain region with a higher Ti-content and a darker contrast (3e-3, 3g-2, 3h-2, and 2h-3), comprising a mixture of TiNi, ZrNi, Zr7Ni10, and Zr9Ni11, can be found; the inter-grain region has been studied extensively by transmission electron microscope and was reported in a previous work [70,71]. This mixed phase
Please cite this article in press as: Young K, et al., Structural, hydrogen storage, and electrochemical properties of Laves phaserelated body-centered-cubic solid solution metal hydride alloys, International Journal of Hydrogen Energy (2014), http:// dx.doi.org/10.1016/j.ijhydene.2014.01.134
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of ZrxNiy is a product of solid-state transformation after the Laves phase has solidified. Compositions of C14 phases determined by EDS are listed in bold in Table 3. In the calculation of B/A stoichiometry, V is assumed to occupy the B-site based on a previous study of Vsubstitution in a series of C14 MH alloys [63]. The Zr-, Ti-, V-, Mn-, Ni-contents in C14 phase are plotted in Fig. 5a as functions of alloy number, which increases as the V-content in the alloy increases. With the increase in overall V-content, the Zrand Mn-contents in C14 phase decrease, the Ti- and Nicontents increase, and the V-content remains at the same level. The increase in Ti/Zr ratio as the overall V-content increases corresponds well with the change in design of overall Ti/Zr ratio and causes the reduction in C14 unit cell volume seen in the XRD analysis. The decrease in Mn is due to the Mn-content reduction in the overall composition as the overall V-content increases. In spite of a decrease in Nicontent in the overall composition with the increase in Vcontent, the Ni-content in C14 phase increases due to the increase in BCC phase abundance, which is low in Ni-content and consequently leaves more Ni for the C14 phase. The stoichiometry of AB2 phase vs. the alloy number is also plotted in Fig. 5a. The B/A ratio decreases rapidly at first and then stabilizes to around 2.2 as the overall stoichiometry increases from 2.50 to 5.37. The Ti-, V-, Cr-, Mn-, and Ni-contents in BCC phase are plotted against the alloy number in Fig. 5b. As the overall Vcontent and BCC abundance increase in the alloy, the Ti-, Vand Ni-contents stay at the same level, the Cr-content increases, and the Mn-content decreases. The increase in Cr and the decrease in Mn are from the changes in overall composition and cause the unit cell to expand slightly due to the larger
Fig. 4 e Schematic of a three dimensional-framework representing grain structure of BCC phase in Laves phaserelated BCC solid solution metal hydride alloys. The shaded area presents the projected image on the SEM micrograph.
size of Cr. The relatively low Ti-content (4.7e5.8 at.% compared to 28.4 at.% in Ref. [67]) is the main reason for the lower than optimized lattice constant value. Therefore, increasing the Ti-content in BCC phase is essential for improving the storage capacity.
Gaseous phase study Gaseous phase hydrogen storage properties of the alloys were studied by PCT. All samples are activated with one thermal cycle in the presence of hydrogen. Consequent measurements did not change the PCT characteristic significantly. The resulting absorption and desorption isotherms measured at 30 and 60 C are shown in Fig. 6. Information obtained from the PCT study is summarized in Table 4. Due to the multi-phase and highly disordered nature of these alloys, a flat plateau is not observed. Therefore, the desorption equilibrium pressure at 0.5 wt.% hydrogen storage is used as a reference point in order to compare the thermodynamic properties of these alloys. The alloys, with the exception of alloys P1 and P8,
Fig. 5 e Chemical composition and B/A ratio for C14 phase (a) and composition for BCC phase (b) as functions of alloy number. In Fig. 5b, the amount of V is divided by 2 for the purpose of easier comparison.
Please cite this article in press as: Young K, et al., Structural, hydrogen storage, and electrochemical properties of Laves phaserelated body-centered-cubic solid solution metal hydride alloys, International Journal of Hydrogen Energy (2014), http:// dx.doi.org/10.1016/j.ijhydene.2014.01.134
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 e1 1
demonstrate comparable equilibrium pressures and reversibility. Both maximum and reversible capacities increase first and then decrease as the average V-content in the alloy increases. Hysteresis of the PCT isotherm is defined as ln(Pa/Pd), where Pa and Pd are the absorption and desorption equilibrium pressures at 0.5 wt.% hydrogen storage, respectively. In general, the hystereses measured at both 30 and 60 C decrease first and then increase as the average V-content in the alloy increases. PCT hysteresis represents the energy required to elastically distort the lattice at the metal/hydride interface. Smaller hysteresis and higher capacities are observed in the mid-range of x values (V-content) in this series, which corresponds to alloys with larger crystallite sizes; the connection is not clear.
7
Due to the steepness in the PCT isotherms, a true plateau pressure cannot be identified for the alloy samples. For these highly disordered MH alloys, the desorption equilibrium pressures at 0.5 wt.% hydrogen storage at 30, 60, and 90 C were used to estimate the changes in enthalpy (DH) and entropy (DS) by the equation
DG ¼ DH TDS ¼ RT ln P
(2)
where R is the ideal gas constant and T is the absolute temperature. Results of these calculations are listed in Table 4. The values may not be accurate and can only be used for comparison among these alloys. As the V-content in the alloy
Fig. 6 e PCT isotherms of alloys P1, P2, P3, P4 at 30 C (a) and 60 C (c), and P5, P6, P7, and P8 at 30 C (b) and 60 C (d). Open and solid symbols are for absorption and desorption curves, respectively. Please cite this article in press as: Young K, et al., Structural, hydrogen storage, and electrochemical properties of Laves phaserelated body-centered-cubic solid solution metal hydride alloys, International Journal of Hydrogen Energy (2014), http:// dx.doi.org/10.1016/j.ijhydene.2014.01.134
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Table 4 e Summary of gaseous phase and thermodynamic properties.
P1 P2 P3 P4 P5 P6 P7 P8
Des. pressure @ 0.5%, 30 C (MPa)
Des. pressure @ 0.5%, 60 C (MPa)
PCT hysteresis @ 0.5%, 30 C
PCT hysteresis @ 0.4%, 60 C
Max. cap. @ 30 C (wt.%)
Rev. cap. @ 30 C (wt.%)
DH (kJ mol1)
DS (J mol1 K1)
0.170 0.077 0.065 0.060 0.064 0.082 0.060 0.003
0.683 0.277 0.232 0.272 0.218 0.305 0.251 0.074
0.77 0.38 0.28 0.28 0.31 0.27 0.70 2.54
0.46 0.46 0.34 0.27 0.37 0.45 0.64 6.54
0.61 0.95 1.00 1.16 1.32 1.21 1.12 1.08
0.44 0.80 0.83 1.04 1.17 1.05 0.94 0.61
38.9 35.8 35.6 35.7 34.3 36.7 40.0 44.6
133 116 114 113 109 120 127 131
increases, both DH and DS first decrease and then increase. With the decreases in C14 unit cell volume and abundance and with the increases in BCC unit cell volume and abundance, a monotonic decrease in DH is expected but not observed. Therefore, the hydride stability is not determined by composition or phase distribution. DS is an indication of how far the MH system is from a perfect, ordered situation. The theoretical value of DS is the entropy of hydrogen gas, which is close to 135 J mol1 K1. For all alloys, especially ones with x values in the mid-range, the DS values higher than 135 J mol1 K1 indicate a higher degree of disorder in the MH system, i.e., the hydrogen occupation sites are far from being ordered. Gaseous phase properties (capacity, hysteresis, heat of formation) in the current study cannot be correlated to either chemical composition or phase abundance; however, they show close correlations to the crystallite size of C14 phase. Smaller C14 crystallites (either low or high V-contents) yield lower gaseous phase hydrogen storage capacities. The real mechanism of this correlation is not clear and needs further investigation.
P7, and P8 exhibit some degree of capacity degradation in the first 13 cycles. Cycle stability worsens as the V-content in the alloy increases. High-rate capacities (50 mA g1) and full capacities (4 mA g1) measured at the 4th and 2nd cycle, respectively, are listed in Table 5. Both capacities decrease and then increase slightly with increasing overall V-content. If the gaseous phase capacity is converted to a theoretical electrochemical one and plotted versus the alloy number along with the electrochemical storage capacities (Fig. 7), a large discrepancy develops between the capacity evolutions in the gaseous phase and in the electrochemical environment. In this plot, the gaseous phase capacity is converted to electrochemical capacity by
1 wt.% of H2 ¼ 268 mAh g1
(3)
While both gaseous phase capacities (full and high-rate) increase to a maximum and decrease as the V-content in the alloy increases, the changes in the electrochemical capacities are relatively small. This result is very different from the conclusion drawn previously from C14 MH alloy studies, where the electrochemical capacity was observed to always fall in between the gaseous phase maximum and reversible capacities [59,72e74]. For the BCC/C14-mixed system in the current study, the electrochemical capacity is not as sensitive to the changes in composition, phase abundance, or microstructure compared to the gaseous phase capacity. The electrochemical capacities of these alloys (157e172 mAh g1) are stable with cycling but much lower than the highest capacity reported by Yu et al. (814 mAh g1) from a Ti40V30Cr15Mn15 BCC alloy. In their case, the capacity soon dropped below
Electrochemical measurement Discharge capacity of each alloy was measured in a floodedcell configuration against a partially pre-charged Ni(OH)2 positive electrode. No alkaline pretreatment was applied before the half-cell measurement. Each sample electrode was charged at a constant current density of 50 mA g1 for 10 h and then discharged at a current density of 50 mA g1 followed by two pulls at 12 and 4 mA g1. From the cycling data, alloys P1, P2, and P3 demonstrate slower activation, and alloys P6,
Table 5 e Summary of electrochemical and magnetic properties.
P1 P2 P3 P4 P5 P6 P7 P8
4th cycle cap @ 50 mA g1 (mAh g1)
2nd cycle cap @ 4 mA g1 (mAh g1)
4th cycle HRD
Activation cycle to reach max. cap.
Diffusion coefficient @ RT (1010 cm2 s1)
Exchange current @ RT (mA g1)
Saturated magnetic susceptibility (memu g1)
161.1 158.2 157.6 150.8 151.7 159.2 166.0 166.7
162.0 161.0 160.7 157.4 156.7 168.7 170.0 172.0
0.990 0.990 0.988 0.983 0.985 0.987 0.989 0.989
8 6 2 2 2 2 2 2
1.29 1.22 1.11 1.08 1.05 1.01 1.15 1.54
39 42 59 35 38 42 39 33
480 345 426 328 425 188 212 306
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small (either at low or high V-content), the D and HRD values are larger. Therefore, in this series of BCC/C14-mixed alloys, the high-rate performance is limited by the bulk transport of hydrogen and enhanced by smaller crystallite size. Compared to a previously studied high-rate C14-predominated MH alloy containing 10 at.% V [69], the D values from the C14/BCCmixed alloys in the current study are lower but the HRD and Io values are considerably higher. In fact, the Io values in the current study are even higher than those from commercial AB5 and A2B7 MH alloys [3]. Further magnetic susceptibility measurement indicates much higher saturated values than those from the conventional MH alloys used in NiMH battery, which is an indication of the larger amount of metallic nickel clusters imbedded in the surface oxide of these alloys. Therefore, BCC/C14-mixed MH alloys may be suitable for high-rate application if both bulk-diffusion and electrochemical capacity can be improved. Fig. 7 e Hydrogen storage capacities converted from gaseous phase hydrogen storage and as measured electrochemically as functions of alloy number.
100 mAh g1 after only 10 cycles because part of the high discharge capacity may come from the metal oxidation [5]. Half-cell HRD of each alloy, defined as the ratio of discharge capacity measured at 50 mA g1 to that measured at 4 mA g1. Within four cycles, HRDs of all alloys are stabilized. The HRD values measured at the stabilized 4th cycle are listed in Table 5. As the V-content in the alloy increases, HRD first decreases and then increases. In order to study the origin of HRD evolution, both bulk diffusion coefficient (D) and surface exchange current (Io) were measured electrochemically. The details of both parameters’ measurements were previously reported [75], and the values are listed in Table 5. While the D value has a close correlation to HRD, Io does not. The D values are plotted together with the crystallite size of C14 phase as functions of alloy number in Fig. 8. When the crystallites are
Fig. 8 e Plots of diffusion coefficient and C14 phase crystallite size as functions of alloy number. The trends indicate that hydrogen diffuses more easily within an alloy with smaller C14 crystallites.
Conclusions A series of Laves phase-related body-centered-cubic solid solution metal hydride alloys with BCC/C14 ratios ranging from 0.09 to 8.52 were studied. Although the electrochemical capacities of these alloys are low and the main limiting factors require further investigation, the high-rate dischargeabilities of these alloys are very high due to the synergetic effect between C14 and BCC phases and the good surface catalytic nature. As the V-content increases in the alloy, properties related to the changes in V-content and/or phase abundance vary monotonically while others vary in one direction and then the opposite (mainly due to the change in primary/secondary phase contact area, i.e., the synergetic effect).
references
[1] Iba H, Akiba E. Hydrogen absorption and microstructure in BCC alloys with C14-type Laves phase. J Jpn Inst Met 1994;58:1225e32. [2] Akiba E, Iba H. Hydrogen absorption by Laves phase related BCC solid solution. Intermetallics 1998;6:461e70. [3] Young K, Nei J. The current status of hydrogen storage alloy development for electrochemical applications. Materials 2013;6:4574e608. [4] Inoue H, Arai S, Iwakura C. Crystallographic and electrochemical characterization of TiV4xNix alloys for nickelemetal hydride batteries. Electrochim Acta 1996;41:937e9. [5] Yu XB, Wu Z, Xia BJ, Xu NX. A TieV-based bcc phase alloy for use as metal hydride electrode with high discharge capacity. J Chem Phys 2004;121:987e90. [6] Chen N, Li R, Zhu Y, Liu Y, Pan H. Electrochemical hydrogenation and dehydrogenation mechanisms of the TieV base multiphase hydrogen storage electrode alloy. Acta Metall Sin 2004;40:1200e4. [7] Iba H, Akiba E. The relation between microstructure and hydrogen absorbing property in Laves phase-solid solution multiphase alloys. J Alloys Compd 1995;231:508e12. [8] Ro¨nnebro E, Nore´us D, Sakai T, Tsukahara M. Structural studies of a new Laves phase alloy (Hf,Ti)(Ni,V)2 and its very stable hydride. J Alloys Compd 1995;231:90e4.
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10
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 e1 1
[9] Tsukahara M, Takahashi K, Mishima T, Isomura A, Sakai T. V-based solid solution alloys with Laves phase network: hydrogen absorption properties and microstructure. J Alloys Compd 1996;236:151e5. [10] Qiu S, Chu H, Zhang Y, Sun D, Song X, Sun L, et al. Electrochemical kinetics and its temperature dependence behaviors of Ti0.17Zr0.08V0.35Cr0.10Ni0.30 alloy electrode. J Alloys Compd 2009;471:453e6. [11] Iba H, Akiba E. Hydrogen absorption and modulated structure in TieVeMn alloys. J Alloys Compd 1997;253e254:21e4. [12] Basak S, Shashikala K, Kulshreshtha SK. Hydrogen absorption characteristics of Zr substituted Ti0.85VFe0.15 alloy. Int J Hydrogen Energy 2008;33:350e5. [13] Shashikala K, Banerjee S, Kumar A, Pai MR, Pillai CGS. Improvement of hydrogen storage properties of TiCrV alloy by Zr substitution for Ti. Int J Hydrogen Energy 2009;34:6684e9. [14] Yang XW, Li JS, Zhang TB, Hu R, Xue XY, Fu HZ. Role of defect structure on hydrogenation properties of Zr0.9Ti0.1V2 alloy. Int J Hydrogen Energy 2011;36:9318e23. [15] Tsukahara M, Takahashi K, Mishima T, Sakai T, Miyamura H, Kuriyama N, et al. Metal hydride electrodes based on solid solution type alloy TiV3Nix (0 & x & 0.75). J Alloys Compd 1995;226:203e7. [16] Tsukahara M, Takahashi K, Mishima T, Isomura A, Sakai T. Heat-treatment effects of V-based solid solution alloy with TiNi-based network structure on hydrogen storage and electrode properties. J Alloys Compd 1996;243:133e8. [17] Tsukahara M, Takahashi K, Mishima T, Isomura A, Sakai T. Vanadium-based solid solution alloys with threedimensional network structure for high capacity metal hydride electrodes. J Alloys Compd 1997;253e254:583e6. [18] Huot J, Akiba E, Ogura T, Ishido Y. Crystal structure, phase abundance and electrode performance of Laves phase compounds (Zr, A)V0.5Ni1.1Mn0.2Fe0.2 (A h Ti, Nb or Hf). J Alloys Compd 1995;218:101e9. [19] Yoshida M, Akiba E. Hydrogen absorbing-desorbing properties and crystal structure of the ZreTieNieMneV AB2 Laves phase alloys. J Alloys Compd 1995;224:121e6. [20] Huot H, Akiba E, Ishido Y. Crystal structure of multiphase alloys (Zr,Ti)(Mn,V)2. J Alloys Compd 1995;231:85e9. [21] Kuriiwa T, Tamura T, Amemiya T, Fuda T, Kamegawa A, Takamura H, et al. New V-based alloys with high protium absorption and desorption capacity. J Alloys Compd 1999;293e295:433e6. [22] Miao H, Gao M, Liu Y, Lin Y, Wang J, Pan H. Effects of Y substitution for Ti on the microstructure and electrochemical properties of TieVeFe-based hydrogen storage alloys. J Electrochem Soc 2007;154:A1010e4. [23] Qiu S, Chu H, Zhang J, Zhang Y, Sun L, Xu F, et al. Effect of La partial substitution for Zr on the structural and electrochemical properties of Ti0.17Zr0.08exLaxV0.35Cr0.1Ni0.3 (x ¼ 0e0.04) electrode alloys. Int J Hydrogen Energy 2009;34:7246e52. [24] Cho S, Enoki H, Akiba E. Effect of Fe addition on hydrogen storage characteristics of Ti0.16Zr0.05Cr0.22V0.57 alloy. J Alloys Compd 2000;307:304e10. [25] Dou T, Wu Z, Mao J, Xu N. Application of commercial ferrovanadium to reduce cost of TieV-based BCC phase hydrogen storage alloys. Mater Sci Eng A 2008;476:34e8. [26] Santos SF, Huot J. Hydrogen storage in TiCr1.2(FeV)x BCC solid solutions. J Alloys Compd 2009;472:247e51. [27] Yu J, Liu B, Cho K, Lee J. The effects of partial substitution of Mn by Cr on the electrochemical cycle life of TieZreVeMneNi alloy electrodes of a Ni/MH battery. J Alloys Compd 1998;278:283e90. [28] Yan Y, Chen Y, Liang H, Wu C, Tao M. The effect of Si on V30Ti35Cr25Fe10 BCC hydrogen storage alloy. J Alloys Compd 2007;441:297e300.
[29] Huang Z, Cuevas F, Liu X, Jiang L, Wang S, Latroche M, et al. Effects of Si addition on the microstructure and the hydrogen storage properties of Ti26.5V45Fe8.5Cr20Ce0.5 BCC solid solution alloys. Int J Hydrogen Energy 2009;34:9385e92. [30] Martı´nez A, dos Santos DS. Influence of the substitution of V by Nb in the structure and properties of hydrogen absorption/desorption of TiCr1.1V0.9 alloy. J Alloys Compd 2012;536(Suppl. 1):S231e5. [31] Martı´nez A, dos Santos DS. Hydrogen absorption/desorption properties in the TiCrV based alloys. Mater Res 2012;15:809e12. [32] Kamegawa A, Shirasaki K, Tamura T, Kuriiwa T, Takamura H, Okada M. Crystal structure and protium absorption properties of TieCreX alloys. Mater Trans 2002;43:470e3. [33] Uno M, Takahashi K, Maruyama T, Muta H, Yamanaka S. Hydrogen solubility of BCC titanium alloys. J Alloys Compd 2004;366:213e6. [34] Shirasaki K, Kuriiwa T, Tamura T, Kamegawa A, Takamura H, Okada M. Effects of Al addition on structures and protium absorptionedesorption properties of TieCr alloys. Mater Trans 2002;43:1173e7. [35] Mi J, Lu¨ F, Liu X, Jiang L, Li Z, Wang S. Enhancement of cerium and hydrogen storage property of a low-cost TieV based BCC alloy prepared by commercial ferrovanadium. J Rare Earths 2010;28:781e4. [36] Liu Y, Zhang S, Li R, Gao M, Zhong K, Miao H, et al. Electrochemical performances of the Pd-added TieV-based hydrogen storage alloys. Int J Hydrogen Energy 2008;33:728e34. [37] Sapru K, Hong K, Fetcenko MA, Venkatesan S, Inventors; Energy Conversion Devices, Inc., Assignee. Hydrogen storage materials and methods of sizing and preparing the same for electrochemical applications. United States patent US 4551400; 1985 Nov 5. [38] Reichman B, Venkatesan S, Fetcenko MA, Jeffries K, Stahl S, Bennett C, Inventors; Energy Conversion Devices, Inc., Assignee. Activated rechargeable hydrogen storage electrode and method. United States patent US 4716088; 1987 Dec 29. [39] Venkatesan S, Reichman B, Fetcenko MA, Inventors; Energy Conversion Devices, Inc., Assignee. Enhanced charge retention electrochemical hydrogen storage alloys and an enhanced charge retention electrochemical cell. United States patent US 4728586; 1988 Mar 1. [40] Ovshinsky SR, Fetcenko MA, Ross J. A nickel metal hydride battery for electric vehicles. Science 1993;260:176e81. [41] Lee H, Lee K, Lee J, Lee J. The Ti-based metal hydride electrode for NieMH rechargeable batteries. J Alloys Compd 1996;239:63e70. [42] Kim JS, Paik CH, Cho WI, Cho BW, Yun KS, Kim SJ. Corrosion behaviour of Zr1exTixV0.6Ni1.2M0.2 (M ¼ Ni, Cr, Mn) AB2-type metal hydride alloys in alkaline solution. J Power Sources 1998;75:1e8. [43] Zhang QA, Lei YQ, Wang CS, Wang FS, Wang QD. Structure of the secondary phase and its effects on hydrogen-storage properties in a Ti0.7Zr0.2V0.1Ni alloy. J Power Sources 1998;75:288e91. [44] Zhu Y, Pan H, Gao M, Liu Y, Wang Q. A study on improving the cycling stability of (Ti0.8Zr0.2)(V0.533Mn0.107Cr0.16Ni0.2)4 hydrogen storage electrode alloy by means of annealing treatment: II. Effects on the electrochemical properties. J Alloys Compd 2003;348:301e8. [45] Chen L, Guo R, Lei Y, Li L, Wang Q. Structural and electrochemical characteristics of V2.1TiNi0.5Hf0.05M0.152 (M ¼ Co, Cr) hydrogen storage electrode alloys. Mater Chem Phys 2005;92:554e8. [46] Pan H, Li R, Liu Y, Gao M, Miao H, Lei Y, et al. Structure and electrochemical properties of the Fe substituted TieV-based hydrogen storage alloys. J Alloys Compd 2008;463:189e95.
Please cite this article in press as: Young K, et al., Structural, hydrogen storage, and electrochemical properties of Laves phaserelated body-centered-cubic solid solution metal hydride alloys, International Journal of Hydrogen Energy (2014), http:// dx.doi.org/10.1016/j.ijhydene.2014.01.134
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 e1 1
[47] Gao M, Zhang S, Miao H, Liu Y, Pan H. Pulverization mechanism of the multiphase TieV-based hydrogen storage electrode alloy during charge/discharge cycling. J Alloys Compd 2010;489:552e7. [48] Chai Y, Zhao M. Structure and electrochemical properties of Ti0.25V0.35Cr0.40exNix (x ¼ 0.05e0.40) solid solution alloys. Int J Hydrogen Energy 2005;30:279e83. [49] Qiao Y, Zhao M, Li M, Zhu X, Cao G. Microstructure and electrochemical performance of Ti0.17Zr0.08V0.34Pd0.01Cr0.1Ni0.3 electrode alloy. Scr Mater 2006;55:279e82. [50] Qiao Y, Zhao M, Zhu X, Cao G. Effect of cerium on microstructure and electrochemical performance of TieVeCreNi electrode alloy. J Rare Earths 2007;25:341e7. [51] Chai YJ, Zhao MS, Wang N. Crystal structural and electrochemical properties of Ti0.17Zr0.08V0.35Cr0.10Ni0.30xMnx (x ¼ 0e0.12) alloys. Mater Sci Eng B 2008;147:47e51. [52] Wang N, Chai Y. Electrochemical properties and reaction mechanism of TieV multiphase alloy at high temperature. Rare Met Mater Eng 2011;40:1519e21. [53] Yin W, Zhao M. Structure and electrochemical characteristics of TiV1.1Mn0.9Nix (x ¼ 0.1e0.7) alloys. Electrochim Acta 2007;52:2723e8. [54] Li S, Zhao M, Wang L, Liu Y, Wang Y. Structures and electrochemical characteristics of Ti0.26Zr0.07V0.24Mn0.1Ni0.33Mox (x ¼ 0e0.1) hydrogen storage alloys. Mater Sci Eng B 2008;150:168e74. [55] Li S, Zhao M, Zhai J, Wang F. Effect of substitution of chromium for nickel on structure and electrochemical characteristic of Ti0.26Zr0.07V0.24Mn0.1Ni0.33 multi-phase hydrogen storage alloy. Mater Chem Phys 2009;113:96e102. [56] Yu X, Wu Z, Li F, Xia B, Xu N. Body-centered-cubic phase hydrogen storage alloy with improved capacity and fast activation. Appl Phys Lett 2004;84:3199e201. [57] Feng S, Wu Z, Yu X, Chen C, Chen K, Chen T, et al. Effect of Zr substitution for Ti on Ti19.5V40Mn16.2Cr9.8Ni14.5 alloy I: microstructure and hydrogen storage property. Rare Met Mater Eng 2007;36:794e8. [58] Gue´guen A, Joubert JM, Latroche M. Influence of the C14 Ti35.4V32.3Fe32.3 Laves phase on the hydrogenation properties of the body-centered cubic compound Ti24.5V59.3Fe16.2. J Alloys Compd 2011;509:3013e8. [59] Young K, Fetcenko MA, Li F, Ouchi T. Structural, thermodynamic, and electrochemical properties of TixZr1x(VNiCrMnCoAl)2 C14 Laves phase alloys. J Alloys Compd 2008;464:238e47. [60] Young K, Fetcenko MA, Koch J, Morii K, Shimizu T. Studies of Sn, Co, Al, and Fe additives in C14/C15 Laves alloys for NiMH battery application by orthogonal arrays. J Alloys Compd 2009;486:559e69. [61] Young K, Ouchi T, Reichman B, Mays W, Regmi R, Lawes G, et al. Optimization of Co-content in C14 Laves phase multicomponent alloys for NiMH battery application. J Alloys Compd 2010;489:202e10.
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[62] Yoo J, Shim G, Yoon J, Cho S. Effects of substituting Al for Cr in the Ti0.32Cr0.43V0.25 alloy on its microstructure and hydrogen storage properties. Int J Hydrogen Energy 2009;34:1463e7. [63] Young K, Fetcenko MA, Li F, Ouchi T, Koch J. Effect of vanadium substitution in C14 Laves phase alloys for NiMH battery application. J Alloys Compd 2009;468:482e92. [64] Young K, Ouchi T, Fetcenko MA. Pressureecompositionetemperature hysteresis in C14 Laves phase alloys: part 1. Simple ternary alloys. J Alloys Compd 2009;480:428e33. [65] Young K, Ouchi T, Mays W, Reichman B, Fetcenko MA. Pressureecompositionetemperature hysteresis in C14 Laves phase alloys: part 2. Applications in NiMH batteries. J Alloys Compd 2009;480:434e9. [66] Young K, Ouchi T, Fetcenko MA. Pressureecompositionetemperature hysteresis in C14 Laves phase alloys: part 3. Empirical formula. J Alloys Compd 2009;480:440e8. [67] Yan Y, Chen Y, Liang H, Zhou X, Wu C, Tao M, et al. Hydrogen storage properties of VeTieCreFe alloys. J Alloys Compd 2008;454:427e31. [68] Chen KC, Peterson EJ, Thoma DJ. HfCo2 Laves phase intermetallicsdpart I: solubility limits and defect mechanisms. Intermetallics 2001;9:771e83. [69] Young K, Ouchi T, Huang B, Reichman B, Blankenship R. Improvement in 40 C electrochemical properties of AB2 metal hydride alloy by silicon incorporation. J Alloys Compd 2013;575:65e72. [70] Bendersky LA, Wang K, Boettinger WJ, Newbury DE, Young K, Chao B. Examination of multiphase (Zr,Ti)(V,Cr,Mn,Ni)2 NieMH electrode alloys: part II. Solid-state transformation of the interdendritic B2 phase. Metall Mater Trans A 2010;41:1891e906. [71] Boettinger WJ, Newbury DE, Wang K, Bendersky LA, Chiu C, Kattner UR, et al. Examination of multiphase (Zr,Ti)(V,Cr,Mn,Ni)2 NieMH electrode alloys: part I. Dendritic solidification structure. Metall Mater Trans A 2010;41:2033e47. [72] Sun D, Jiang J, Lei Y, Liu W, Wu J, Wang Q, et al. Effects of measurement factor on electrochemical capacity of some hydrogen storage alloys. Mater Sci Eng B 1995;30:19e22. [73] Young K, Nei J, Huang B, Ouchi T, Fetcenko MA. Studies of Ti1.5Zr5.5V0.5(MxNi1x)9.5 (M ¼ Cr, Mn, Fe, Co, Cu, Al): part 2. Hydrogen storage and electrochemical properties. J Alloys Compd 2010;501:245e54. [74] Young K, Chao B, Huang B, Nei J. Studies on the hydrogen storage characteristic of La1xCex(NiCoMnAlCuSiZr)5.7 with a B2 secondary phase. J Alloys Compd 2014;585:760e70. [75] Li F, Young K, Ouchi T, Fetcenko MA. Annealing effects on structural and electrochemical properties of (LaPrNdZr)0.83Mg0.17(NiCoAlMn)3.3 alloy. J Alloys Compd 2009;471:371e7.
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Structural, hydrogen storage, and electrochemical properties of Laves phase-related body-centeredcubic solid solution metal hydride alloys K. Young a,*, J. Nei a, D.F. Wong a,b, L. Wang a,b a b
BASF/Battery Materials-Ovonic, 2983 Waterview Drive, Rochester Hills, MI 48309, USA Department of Chemical Engineering and Materials Science, Wayne State University, MI 48202, USA
article info
abstract
Article history:
Structure, gaseous phase hydrogen storage, and electrochemical properties of a series of
Received 13 November 2013
Laves phase-related BCC solid solution metal hydride alloys with BCC/C14 ratios ranging
Received in revised form
from 0.09 to 8.52 were studied. Some properties are correlated to the phase abundance and
2 January 2014
V-content in the alloy with monotonic evolutions, for example, lattice constant, phase
Accepted 22 January 2014
abundance, and hydrogen storage pressure. Other properties such as gaseous phase ca-
Available online xxx
pacities, PCT hysteresis, high-rate dischargeability, and bulk hydrogen diffusion correlate better with the C14 phase crystallite size, which are considered to be more related to the
Keywords:
synergetic effect between main and secondary phases. In contrast with conventional metal
Hydrogen absorbing alloys
hydride alloys used in NiMH batteries, the electrochemical discharge capacities of these
Metal hydride electrode
alloys are not between the maximum and the reversible hydrogen storage measured in the
Laves phase alloys
gaseous phase. The current study’s alloys have electrochemical capacities that are
Body-centered-cubic alloys
insensitive to composition but have much room for improvement, with high-rate dis-
Phase segregation
chargeabilities that are superior compared to other commercially available alloys. With further research, these alloys show potential for high-rate battery applications. Copyright ª 2014, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
Introduction Iba and Akiba described the “Laves phase-related bodycentered-cubic (BCC) solid solution” as a family of metal hydride (MH) alloys with a two-phase microstructure composed of a BCC phase and a Laves phase (mostly C14) and a general formula of ABx, where A is an element or any combination of elements from Group 4A (mostly Ti), B is from Group 5A, 6A, or 7A (mainly V), and x is between 1 and 6 [1,2]. The BCC phase, with its higher V-content, has higher melting temperature and solidifies before the Laves phases (C14, for example) do. This
phenomenon produces the following evolution in microstructure with the increase in C14/BCC ratio (Fig. 5 in Ref. [3]): C14 appears at the grain boundaries / sections of C14 phase connect to form a three-dimensional network / the alloy matrix shifts to C14, and BCC phase forms its own threedimensional network, a cross-sectional view showing embedded, isolated BCC islands / the C14 matrix continues to expand, and BCC phase forms fish-bone-like inclusions within the matrix. The high density of phase boundaries makes it possible to combine the advantages of BCC (high hydrogen storage capability [4,5]) and C14 phases (good absorption kinetics [6], easy formation due to its brittleness
* Corresponding author. Tel.: þ1 248 293 7000; fax: þ1 248 299 4520. E-mail addresses: [email protected], [email protected] (K. Young). 0360-3199/$ e see front matter Copyright ª 2014, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijhydene.2014.01.134
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[7e9], and high surface catalytic activity [10]) by increasing the synergetic effect between the two phases. The density of phase boundaries also promotes the formation of coherent and catalytic interfaces between the boundaries, improving hydrogen absorption [11]. Similar two-phase microstructures involving BCC phase are also seen in other systems, for example, C15/BCC in Ni-free alloys [12e14] and TiNi/BCC in TieVeNi ternary alloys [15e17]. Both A-site (Ti) and B-site (V) substitution studies were performed on a number of mixed phase alloys, evaluating alloy properties for gaseous phase hydrogen storage applications. The A-site substitutions Zr [18e21] and Hf [8] were found to increase C14 phase abundance significantly, while Y only increased it slightly [22] and La promoted C15 phase [23]. The B-site substitutions Mn [19], Fe [24e26], Ni [27], Si [28,29], Nb [30,31], Ta [32], and B [33] promoted C14 phase, and Al [34], Cr [27], Ce [35], Mo [32,36], and W [32] had the opposite effect. Soon after the gaseous phase hydrogen storage characteristics of this alloy family were reported in 1995 [7,18], its electrochemical properties and potential for nickel/metal hydride (Ni/MH) battery applications attracted researchers’ attention [21,27,37e42]. In particular, researchers from three institutes in China have made continuous efforts in improving the gaseous and electrochemical hydrogen storage properties of this family of alloys. Through a series of composition refinements [35,36,43e47], Wang, Pan, and their coworkers at Zhejiang University obtained an alloy formula (Ti0.8Zr0.2)(V0.5Mn0.1Cr0.16Ni0.2Fe0.04)5 with superb high-rate dischargeability (HRD) performance. Zhao and his coworkers at Chinese Academy of Science in Changchun worked on AB3 [48e52] and AB2 [53e55] alloys and discovered the formulas and (Ti0.79Zr0.21)(V0.36 (Ti0.68Zr0.32)(V0.47Cr0.13Ni0.33Mn0.07)3 Mn0.15Ni0.42Cr0.07)2 to have good balance among various electrochemical performance properties. Yu and his coworkers from Chinese Academy of Science in Shanghai reported that gaseous phase storage capacity and kinetics can be improved by reducing the amount of C14 phase; however, the alloys become more difficult to activate [5,56,57]. Although many studies focusing on this family of alloys have been performed in the past, the studies offer somewhat mixed conclusions. For example, while Ovonic used BCC as a
secondary phase in a C14-based alloy to protect the surface from alkaline corrosion [40], other groups used C14 as a secondary phase to improve activation and charge-transfer properties. Furthermore, while Chai et al. claimed a BCC:C14 abundance ratio of 1:1 yielded optimized capacity results [51], Gue´guen et al. showed that only a small amount of C14 is needed to improve gaseous phase storage [58]. Reviewing the previous studies, one cannot find an optimized C14 abundance for electrochemical performance. Therefore, a systematic study of alloys with a broad range of BCC/C14 ratio is of interest to researchers in the field and presented in this paper.
Experimental setup Arc melting was performed under continuous argon flow with a non-consumable tungsten electrode and a watercooled copper tray. Before each run, a piece of sacrificial titanium underwent a few meltingecooling cycles to reduce the residual oxygen concentration in the system. Each 12-g ingot was re-melted and turned over several times to ensure uniformity in chemical composition. Chemical composition for the ingot was analyzed using a Varian Liberty 100 inductively coupled plasma optical emission spectrometer (ICP-OES). A Philips X’Pert Pro X-ray diffractometer (XRD) was used to study the microstructure, and a JEOLJSM6320F scanning electron microscope (SEM) with energy dispersive spectroscopy (EDS) capability was used to study the phase distribution and composition. Gaseous phase hydrogen storage characteristics for each sample were measured using a Suzuki-Shokan multi-channel pressureeconcentrationetemperature (PCT) system. In the PCT analysis, each sample (a single piece of ingot with newly cleaved surface and a weight of about 2 g) was first activated by a 2-h thermal cycle between 300 C and room temperature at 2.5 MPa H2 pressure. PCT isotherms at 30, 60, and 90 C were then measured. Details of electrode preparations as well as the measurement methods have been reported previously [59,60].
Table 1 e Design compositions (in bold) and ICP results in at.%. Alloy P1 P2 P3 P4 P5 P6 P7 P8
Design ICP Design ICP Design ICP Design ICP Design ICP Design ICP Design ICP Design ICP
Ti
Zr
V
Cr
Mn
Fe
Co
Ni
Al
B/A ratio
x Value in (1)
14.7 15.1 14.5 14.7 14.3 14.7 14.1 14.1 13.9 14.0 13.8 14.0 13.7 14.2 13.6 13.8
13.8 13.9 11.1 11.0 8.9 8.8 7.0 6.9 5.5 5.3 4.2 4.3 3.1 3.0 2.1 2.0
20.0 20.2 25.6 25.4 30.1 31.0 33.8 34.1 37.0 36.9 39.7 38.6 42.0 41.3 44.0 43.1
6.0 3.8 7.7 7.3 9.0 7.5 10.1 10.0 11.1 11.5 11.9 11.4 12.6 12.1 13.2 13.1
12.6 12.8 11.2 10.9 10.2 9.8 9.3 9.0 8.6 8.1 7.9 8.1 7.4 7.0 6.9 6.5
1.2 1.3 1.6 1.6 1.9 1.9 2.1 2.1 2.3 2.3 2.5 2.7 2.6 2.7 2.7 2.8
2.6 2.7 2.3 2.4 2.1 2.1 1.9 1.9 1.8 1.8 1.6 1.8 1.5 1.6 1.4 1.4
28.5 29.5 25.6 26.2 23.1 23.7 21.1 21.4 19.5 19.7 18.0 18.6 16.8 17.6 15.7 16.7
0.6 0.7 0.5 0.5 0.5 0.5 0.4 0.5 0.4 0.5 0.4 0.5 0.3 0.5 0.3 0.6
2.50 2.45 2.91 2.89 3.32 3.26 3.73 3.76 4.14 4.19 4.55 4.46 4.96 4.81 5.37 5.33
0.7 1.0 1.3 1.6 1.9 2.2 2.5 2.8
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8.100
2.980
8.075
2.975
BCC, a
8.050 8.025
2.970 2.965
C14, c
8.000
2.960
7.975
2.955
4.950
2.950
4.925
2.945
C14, a
4.900 1
2
3
BCC Lattice Constant (Å)
C14 Lattice Constants (Å)
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2.940 4
5
6
7
8
Alloy Number
Fig. 2 e Lattice constants a and c from C14 phase and a from BCC phase as functions of alloy number.
Fig. 1 e XRD patterns using Cu-Ka as the radiation source for alloys P1 (a), P2 (b), P3 (c), P4 (d), P5 (e), P6 (f), P7 (g), and P8 (h). Vertical lines are used to indicate shifts in both C14 and BCC peaks.
Results and discussion Alloy composition Eight alloys with a general formula of
Ti0.4þx/6Zr0.6x/6Mn0.44Ni1.0Al0.02Co0.09(VCr0.3Fe0.063)x, x ¼ 0.7 to 2.8
(1)
were prepared by arc melting for this study. Design compositions of these alloys are listed in Table 1, and the design principle is described as follows. The alloy formula with x ¼ 2.5 is close to the formula developed by Pan et al. [46], and the extra Al and Co at the optimized ratio is added to improve
C14 electrochemical properties [60,61]. While the extra Al and Co can improve C14 electrochemical properties, they can also affect BCC properties. Al is known to expand the BCC unit cell; however, it lowers capacity [62]. Co suppresses pulverization, improves the cycle stability and HRD in a BCC-based alloy, but reduces capacity [45]. Therefore, there is a trade-off in the amounts of Al and Co. In this study, we reduce the amount of Co and Al while increasing V, when the main phase shifts from C14 into BCC. The alloy formula with x ¼ 0.7 is close to Ovonic’s AB2 formula developed in the 1980s [39]. With the increase in Vcontent (x), the Ti- and Zr-contents have to be altered accordingly to maintain the metal-to-hydrogen (MeH) bond strength in the C14 phase. Assuming the BCC/C14 ratio is approximately 1:1 and that one-fifth of the V-content in BCC phase tends to appear in the C14 phase [45], the amount of V that goes into C14 phase is about one-sixth of the total amount. According to the results from previous experiments, the target decrease in Zr (weaker MeH bond) at the expense of Ti is roughly the same as the amount of increase in V (stronger MeH bond) [63] in the C14 phase. Therefore, onesixth of x is subtracted from the Zr-content and added to the Ti-content to maintain the same AB2 stoichiometry and MeH bond strength. Pd is also a possible modifier [49,54] but is too expensive to be considered for practical application.
Table 2 e Lattice constants, unit cell volumes, phase abundances, and crystallite sizes of C14 and BCC phases of alloys P1eP8 from XRD analysis. FWHM and XS denote full-width at half maximum and crystallite size, respectively. C14, a ( A) C14, c ( A) C14, a/c C14 unit cell C14 wt.% C14 (103) C14 XS ( A) BCC, a ( A) BCC wt.% BCC (200) BCC XS ( A) volume ( A3) FWHM FWHM P1 P2 P3 P4 P5 P6 P7 P8
4.9525 4.9386 4.9292 4.9138 4.9165 4.9078 4.9064 4.9025
8.0690 8.0470 8.0306 8.0040 8.0044 7.9897 7.9835 7.9728
0.6138 0.6137 0.6138 0.6139 0.6142 0.6143 0.6146 0.6149
171.40 169.97 168.98 167.37 167.56 166.66 166.44 165.95
91.6 85.9 80.4 72.1 64.7 53.5 45.8 10.5
0.457 0.376 0.351 0.316 0.319 0.405 0.363 0.427
207 262 285 325 321 289 273 225
2.9568 2.9576 2.9601 2.9621 2.9633 2.9649 2.9655 2.9664
8.4 14.1 19.6 27.9 35.3 46.5 54.2 89.5
e 0.197 0.365 0.391 0.464 0.648 0.604 0.689
e 714 298 274 224 174 167 144
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ICP results from the ingot samples are listed in Table 1. Other than some uneven distribution of Cr in the alloys with lower Cr-content (alloys P1 and P3), ICP results are very close to the target compositions. The measured B/A ratio varies from 2.45 to 5.33 and is within 3% of the design value.
XRD structure analysis XRD patterns of the eight alloys are shown in Fig. 1. Two sets of diffraction peaks, C14 and BCC, are observed. As the Vcontent in the alloy increases, the BCC peaks increase in intensity and shift to lower angles, while the C14 peaks behave the opposite. Lattice constants of both phases calculated from the XRD patterns are listed in Table 2 and plotted in Fig. 2. As the amount of V increases, both a and c in C14 phase decrease, while a in BCC phase increases. In C14 phase, the decrease in lattice parameter c is faster than the decrease in a, which
results in an increasing a/c aspect ratio with increasing Vcontent in the alloys. Therefore, a lower pulverization rate during hydride/dehydride cycling is expected from alloys with higher V-content in this study [64e66]. Lattice parameter a in BCC phase increases from 2.9568 to 2.9664 A as the V-content in the alloy increases, which is smaller than the optimized value of 3.042 A corresponding to a maximized hydrogen storage capacity [67], leaving room for potential improvement in storage capacity in future studies. Phase abundance and crystallite sizes of each phase are listed in Table 2. These values were obtained from full pattern fitting of the XRD data using the Rietveld method and Jade 9 software. As the V-content in the alloy increases, the abundance of C14 phase decreases monotonically from 91.6 to 10.6 wt.% while the abundance of BCC phase increases. At the same time, the C14 crystallite size first increases and then decreases, while the BCC crystallite size decreases
Fig. 3 e SEM back-scattering electron images for alloys P1 (a), P2 (b), P3 (c), P4 (d), P5 (e), P6 (f), P7 (g), and P8 (h). Chemical compositions in the numbered areas measured by EDS are listed in Table 3. Please cite this article in press as: Young K, et al., Structural, hydrogen storage, and electrochemical properties of Laves phaserelated body-centered-cubic solid solution metal hydride alloys, International Journal of Hydrogen Energy (2014), http:// dx.doi.org/10.1016/j.ijhydene.2014.01.134
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Table 3 e Summary of EDS results. All compositions are in at.%. Compositions of main AB2 phase are in bold. Fig. 3a-1 2 3 4 5 6 Fig. 3b-1 2 3 4 Fig. 3c-1 2 3 4 Fig. 3d-1 2 3 4 Fig. 3e-1 2 3 4 5 Fig. 3f-1 2 3 4 Fig. 3g-1 2 3 4 Fig. 3h-1 2 3 4 5
Ti
Zr
V
Cr
Mn
Fe
Co
Ni
Al
B/A
Phase
12.8 13.9 17.8 19.5 2.1 5.8 15.5 18.6 5.1 6.9 18.0 20.9 5.6 9.2 18.1 20.4 4.7 0.8 20.2 22.1 32.4 5.0 0.9 22.0 29.7 5.5 2.3 23.1 35.5 5.8 3.4 22.4 36.3 35.1 5.7 1.6
17.1 15.3 10.8 9.1 87.4 1.6 15.1 11.9 0.5 66.3 13.1 10.9 0.3 59.5 13.3 10.5 0.1 94.1 11.0 9.3 4.6 0.0 95.9 9.2 3.3 0.1 90.2 8.3 3.3 0.0 87.4 9.0 2.4 8.2 0.1 93.1
20.5 20.6 20.7 19.1 3.0 56.6 18.6 18.2 58.8 8.5 20.1 19.2 58.0 10.4 19.4 19.2 58.6 2.0 18.7 19.4 10.0 58.9 1.3 20.7 17.6 58.5 3.6 20.0 9.0 59.4 4.1 19.5 8.8 16.1 58.7 2.8
5.0 4.3 2.8 2.2 0.9 9.1 5.6 4.3 15.2 2.4 4.1 3.3 16.7 2.0 5.0 4.2 20.2 0.5 5.0 4.1 1.4 20.0 0.3 4.5 4.2 19.6 1.0 4.3 1.2 18.3 1.1 2.4 1.3 1.4 19.2 1.0
15.0 13.4 11.0 9.7 1.8 14.9 11.4 9.3 12.7 4.4 9.4 7.8 10.8 4.5 9.6 8.1 9.0 0.5 9.0 7.1 3.9 8.4 0.3 7.9 5.0 7.7 0.7 7.3 3.2 7.4 0.9 7.1 3.9 4.5 6.9 0.3
1.8 1.6 1.0 0.8 0.3 1.1 2.1 1.3 1.6 0.6 2.0 1.6 2.0 0.8 2.9 1.9 1.8 0.2 3.3 2.3 1.1 2.3 0.0 3.2 1.8 2.3 0.2 3.6 1.8 2.8 0.2 2.8 3.3 2.0 2.7 0.2
2.9 2.9 2.5 2.1 0.4 1.3 2.9 2.5 1.1 0.9 2.7 2.1 1.0 1.0 2.8 2.4 0.9 0.3 2.9 2.2 2.2 0.8 0.0 2.5 2.3 0.9 0.1 2.8 2.8 1.1 0.1 1.8 3.6 1.4 1.0 0.1
24.2 27.3 32.7 36.9 3.9 9.4 28.2 33.2 4.8 9.9 29.8 33.5 5.2 12.2 28.3 32.5 4.5 1.6 29.0 32.9 43.4 4.5 1.1 29.4 35.3 5.0 1.9 30.0 42.3 5.0 2.7 33.8 39.0 30.8 5.3 0.9
0.6 0.7 0.7 0.7 0.2 0.3 0.6 0.7 0.3 0.1 0.7 0.7 0.3 0.2 0.7 0.7 0.2 0.1 0.8 0.7 1.0 0.2 0.1 0.6 0.8 0.4 0.1 0.7 0.9 0.2 0.1 1.0 1.3 0.5 0.5 0.1
2.34 2.42 2.50 2.50 0.12 12.53 2.27 2.28 16.88 0.37 2.21 2.14 15.93 0.45 2.19 2.23 19.83 0.05 2.20 2.19 1.70 19.02 0.03 2.21 2.03 16.86 0.08 2.19 1.58 16.24 0.10 2.18 1.58 1.31 16.26 0.06
AB2 AB2 AB2 AB2 ZrO2 BCC AB2 AB2 BCC ZrO2 AB2 AB2 BCC Zr AB2 AB2 BCC Zr AB2 AB2 ZrxNiy BCC Zr AB2 AB2 BCC Zr AB2 ZrxNiy BCC Zr AB2 ZrxNiy ZrxNiy BCC Zr
monotonically. The measured BCC/C14 ratio increases from 0.09 to 8.52 at a more dramatic rate compared to the design ratio (about 0.24e0.57).
SEM/EDS microstructure analysis Microstructures for this series of alloys were studied using SEM, and the back-scattering electron images (BEI) are presented in Fig. 3. Chemical compositions of several areas with different contrasts were studied by EDS analysis, and the results are summarized in Table 3. During cooling, BCC phase with high V-content solidifies first to form a threedimensional framework while the rest of the liquid solidifies into Laves phase as the alloy cools further. Darker regions in the SEM micrographs are BCC phase (lower average atomic mass due to lower Zr-content) and represent the intersection of a plane and a three-dimensional framework with three major axes perpendicular to each other (Fig. 4). In Fig. 3a from Alloy P1, the BCC phase starts to appear as small dark inclusions (Fig. 3a spot 6, abbrev. as 3a-6) of about 2e3 mm in size. The matrix is composed of AB2 phases with composition variation mainly in Ti/Zr/Ni-contents (3a-1 to 3a-4). Besides, occasional ZrO2 inclusion can be also found (3a-5). As the V-
content in the alloy increases, BCC phase expands, and its shape evolves from dots to fish-bones and then to islands. Since the designed B/A ratios in these alloys (column 12 in Table 1) are well above 2.5 (the approximate maximum solubility of AB2 phase on the B-rich side [68]) and the BCC phase contains almost no Zr, the extra Zr is solidified individually. With the increase in V-content in the alloy, the extra Zr changes its form of existence from ZrO2 (3a-5 and 3b-4) to metallic Zr (3c-4, 3d-4, 3e-5, 3f-4, 3g-4, and 3h-5). In contrast, there are cases where metallic Zr-inclusions found in Si-free alloys are replaced by occasional dark spots of ZrO2 in Sisubstituted alloys [69]. V is a better oxygen salvager compared to Zr, and the oxide for V is light and floats on the surface to form slag, which separates out from the liquid by the tundish. As such, the extra Zr is less oxidized with higher V-content in the alloy. In alloys P2eP4, the AB2 matrix is more homogeneous in composition when compared to alloy P1. For alloys with higher V-content (alloys P5, P7, and P8), an intergrain region with a higher Ti-content and a darker contrast (3e-3, 3g-2, 3h-2, and 2h-3), comprising a mixture of TiNi, ZrNi, Zr7Ni10, and Zr9Ni11, can be found; the inter-grain region has been studied extensively by transmission electron microscope and was reported in a previous work [70,71]. This mixed phase
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of ZrxNiy is a product of solid-state transformation after the Laves phase has solidified. Compositions of C14 phases determined by EDS are listed in bold in Table 3. In the calculation of B/A stoichiometry, V is assumed to occupy the B-site based on a previous study of Vsubstitution in a series of C14 MH alloys [63]. The Zr-, Ti-, V-, Mn-, Ni-contents in C14 phase are plotted in Fig. 5a as functions of alloy number, which increases as the V-content in the alloy increases. With the increase in overall V-content, the Zrand Mn-contents in C14 phase decrease, the Ti- and Nicontents increase, and the V-content remains at the same level. The increase in Ti/Zr ratio as the overall V-content increases corresponds well with the change in design of overall Ti/Zr ratio and causes the reduction in C14 unit cell volume seen in the XRD analysis. The decrease in Mn is due to the Mn-content reduction in the overall composition as the overall V-content increases. In spite of a decrease in Nicontent in the overall composition with the increase in Vcontent, the Ni-content in C14 phase increases due to the increase in BCC phase abundance, which is low in Ni-content and consequently leaves more Ni for the C14 phase. The stoichiometry of AB2 phase vs. the alloy number is also plotted in Fig. 5a. The B/A ratio decreases rapidly at first and then stabilizes to around 2.2 as the overall stoichiometry increases from 2.50 to 5.37. The Ti-, V-, Cr-, Mn-, and Ni-contents in BCC phase are plotted against the alloy number in Fig. 5b. As the overall Vcontent and BCC abundance increase in the alloy, the Ti-, Vand Ni-contents stay at the same level, the Cr-content increases, and the Mn-content decreases. The increase in Cr and the decrease in Mn are from the changes in overall composition and cause the unit cell to expand slightly due to the larger
Fig. 4 e Schematic of a three dimensional-framework representing grain structure of BCC phase in Laves phaserelated BCC solid solution metal hydride alloys. The shaded area presents the projected image on the SEM micrograph.
size of Cr. The relatively low Ti-content (4.7e5.8 at.% compared to 28.4 at.% in Ref. [67]) is the main reason for the lower than optimized lattice constant value. Therefore, increasing the Ti-content in BCC phase is essential for improving the storage capacity.
Gaseous phase study Gaseous phase hydrogen storage properties of the alloys were studied by PCT. All samples are activated with one thermal cycle in the presence of hydrogen. Consequent measurements did not change the PCT characteristic significantly. The resulting absorption and desorption isotherms measured at 30 and 60 C are shown in Fig. 6. Information obtained from the PCT study is summarized in Table 4. Due to the multi-phase and highly disordered nature of these alloys, a flat plateau is not observed. Therefore, the desorption equilibrium pressure at 0.5 wt.% hydrogen storage is used as a reference point in order to compare the thermodynamic properties of these alloys. The alloys, with the exception of alloys P1 and P8,
Fig. 5 e Chemical composition and B/A ratio for C14 phase (a) and composition for BCC phase (b) as functions of alloy number. In Fig. 5b, the amount of V is divided by 2 for the purpose of easier comparison.
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demonstrate comparable equilibrium pressures and reversibility. Both maximum and reversible capacities increase first and then decrease as the average V-content in the alloy increases. Hysteresis of the PCT isotherm is defined as ln(Pa/Pd), where Pa and Pd are the absorption and desorption equilibrium pressures at 0.5 wt.% hydrogen storage, respectively. In general, the hystereses measured at both 30 and 60 C decrease first and then increase as the average V-content in the alloy increases. PCT hysteresis represents the energy required to elastically distort the lattice at the metal/hydride interface. Smaller hysteresis and higher capacities are observed in the mid-range of x values (V-content) in this series, which corresponds to alloys with larger crystallite sizes; the connection is not clear.
7
Due to the steepness in the PCT isotherms, a true plateau pressure cannot be identified for the alloy samples. For these highly disordered MH alloys, the desorption equilibrium pressures at 0.5 wt.% hydrogen storage at 30, 60, and 90 C were used to estimate the changes in enthalpy (DH) and entropy (DS) by the equation
DG ¼ DH TDS ¼ RT ln P
(2)
where R is the ideal gas constant and T is the absolute temperature. Results of these calculations are listed in Table 4. The values may not be accurate and can only be used for comparison among these alloys. As the V-content in the alloy
Fig. 6 e PCT isotherms of alloys P1, P2, P3, P4 at 30 C (a) and 60 C (c), and P5, P6, P7, and P8 at 30 C (b) and 60 C (d). Open and solid symbols are for absorption and desorption curves, respectively. Please cite this article in press as: Young K, et al., Structural, hydrogen storage, and electrochemical properties of Laves phaserelated body-centered-cubic solid solution metal hydride alloys, International Journal of Hydrogen Energy (2014), http:// dx.doi.org/10.1016/j.ijhydene.2014.01.134
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Table 4 e Summary of gaseous phase and thermodynamic properties.
P1 P2 P3 P4 P5 P6 P7 P8
Des. pressure @ 0.5%, 30 C (MPa)
Des. pressure @ 0.5%, 60 C (MPa)
PCT hysteresis @ 0.5%, 30 C
PCT hysteresis @ 0.4%, 60 C
Max. cap. @ 30 C (wt.%)
Rev. cap. @ 30 C (wt.%)
DH (kJ mol1)
DS (J mol1 K1)
0.170 0.077 0.065 0.060 0.064 0.082 0.060 0.003
0.683 0.277 0.232 0.272 0.218 0.305 0.251 0.074
0.77 0.38 0.28 0.28 0.31 0.27 0.70 2.54
0.46 0.46 0.34 0.27 0.37 0.45 0.64 6.54
0.61 0.95 1.00 1.16 1.32 1.21 1.12 1.08
0.44 0.80 0.83 1.04 1.17 1.05 0.94 0.61
38.9 35.8 35.6 35.7 34.3 36.7 40.0 44.6
133 116 114 113 109 120 127 131
increases, both DH and DS first decrease and then increase. With the decreases in C14 unit cell volume and abundance and with the increases in BCC unit cell volume and abundance, a monotonic decrease in DH is expected but not observed. Therefore, the hydride stability is not determined by composition or phase distribution. DS is an indication of how far the MH system is from a perfect, ordered situation. The theoretical value of DS is the entropy of hydrogen gas, which is close to 135 J mol1 K1. For all alloys, especially ones with x values in the mid-range, the DS values higher than 135 J mol1 K1 indicate a higher degree of disorder in the MH system, i.e., the hydrogen occupation sites are far from being ordered. Gaseous phase properties (capacity, hysteresis, heat of formation) in the current study cannot be correlated to either chemical composition or phase abundance; however, they show close correlations to the crystallite size of C14 phase. Smaller C14 crystallites (either low or high V-contents) yield lower gaseous phase hydrogen storage capacities. The real mechanism of this correlation is not clear and needs further investigation.
P7, and P8 exhibit some degree of capacity degradation in the first 13 cycles. Cycle stability worsens as the V-content in the alloy increases. High-rate capacities (50 mA g1) and full capacities (4 mA g1) measured at the 4th and 2nd cycle, respectively, are listed in Table 5. Both capacities decrease and then increase slightly with increasing overall V-content. If the gaseous phase capacity is converted to a theoretical electrochemical one and plotted versus the alloy number along with the electrochemical storage capacities (Fig. 7), a large discrepancy develops between the capacity evolutions in the gaseous phase and in the electrochemical environment. In this plot, the gaseous phase capacity is converted to electrochemical capacity by
1 wt.% of H2 ¼ 268 mAh g1
(3)
While both gaseous phase capacities (full and high-rate) increase to a maximum and decrease as the V-content in the alloy increases, the changes in the electrochemical capacities are relatively small. This result is very different from the conclusion drawn previously from C14 MH alloy studies, where the electrochemical capacity was observed to always fall in between the gaseous phase maximum and reversible capacities [59,72e74]. For the BCC/C14-mixed system in the current study, the electrochemical capacity is not as sensitive to the changes in composition, phase abundance, or microstructure compared to the gaseous phase capacity. The electrochemical capacities of these alloys (157e172 mAh g1) are stable with cycling but much lower than the highest capacity reported by Yu et al. (814 mAh g1) from a Ti40V30Cr15Mn15 BCC alloy. In their case, the capacity soon dropped below
Electrochemical measurement Discharge capacity of each alloy was measured in a floodedcell configuration against a partially pre-charged Ni(OH)2 positive electrode. No alkaline pretreatment was applied before the half-cell measurement. Each sample electrode was charged at a constant current density of 50 mA g1 for 10 h and then discharged at a current density of 50 mA g1 followed by two pulls at 12 and 4 mA g1. From the cycling data, alloys P1, P2, and P3 demonstrate slower activation, and alloys P6,
Table 5 e Summary of electrochemical and magnetic properties.
P1 P2 P3 P4 P5 P6 P7 P8
4th cycle cap @ 50 mA g1 (mAh g1)
2nd cycle cap @ 4 mA g1 (mAh g1)
4th cycle HRD
Activation cycle to reach max. cap.
Diffusion coefficient @ RT (1010 cm2 s1)
Exchange current @ RT (mA g1)
Saturated magnetic susceptibility (memu g1)
161.1 158.2 157.6 150.8 151.7 159.2 166.0 166.7
162.0 161.0 160.7 157.4 156.7 168.7 170.0 172.0
0.990 0.990 0.988 0.983 0.985 0.987 0.989 0.989
8 6 2 2 2 2 2 2
1.29 1.22 1.11 1.08 1.05 1.01 1.15 1.54
39 42 59 35 38 42 39 33
480 345 426 328 425 188 212 306
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small (either at low or high V-content), the D and HRD values are larger. Therefore, in this series of BCC/C14-mixed alloys, the high-rate performance is limited by the bulk transport of hydrogen and enhanced by smaller crystallite size. Compared to a previously studied high-rate C14-predominated MH alloy containing 10 at.% V [69], the D values from the C14/BCCmixed alloys in the current study are lower but the HRD and Io values are considerably higher. In fact, the Io values in the current study are even higher than those from commercial AB5 and A2B7 MH alloys [3]. Further magnetic susceptibility measurement indicates much higher saturated values than those from the conventional MH alloys used in NiMH battery, which is an indication of the larger amount of metallic nickel clusters imbedded in the surface oxide of these alloys. Therefore, BCC/C14-mixed MH alloys may be suitable for high-rate application if both bulk-diffusion and electrochemical capacity can be improved. Fig. 7 e Hydrogen storage capacities converted from gaseous phase hydrogen storage and as measured electrochemically as functions of alloy number.
100 mAh g1 after only 10 cycles because part of the high discharge capacity may come from the metal oxidation [5]. Half-cell HRD of each alloy, defined as the ratio of discharge capacity measured at 50 mA g1 to that measured at 4 mA g1. Within four cycles, HRDs of all alloys are stabilized. The HRD values measured at the stabilized 4th cycle are listed in Table 5. As the V-content in the alloy increases, HRD first decreases and then increases. In order to study the origin of HRD evolution, both bulk diffusion coefficient (D) and surface exchange current (Io) were measured electrochemically. The details of both parameters’ measurements were previously reported [75], and the values are listed in Table 5. While the D value has a close correlation to HRD, Io does not. The D values are plotted together with the crystallite size of C14 phase as functions of alloy number in Fig. 8. When the crystallites are
Fig. 8 e Plots of diffusion coefficient and C14 phase crystallite size as functions of alloy number. The trends indicate that hydrogen diffuses more easily within an alloy with smaller C14 crystallites.
Conclusions A series of Laves phase-related body-centered-cubic solid solution metal hydride alloys with BCC/C14 ratios ranging from 0.09 to 8.52 were studied. Although the electrochemical capacities of these alloys are low and the main limiting factors require further investigation, the high-rate dischargeabilities of these alloys are very high due to the synergetic effect between C14 and BCC phases and the good surface catalytic nature. As the V-content increases in the alloy, properties related to the changes in V-content and/or phase abundance vary monotonically while others vary in one direction and then the opposite (mainly due to the change in primary/secondary phase contact area, i.e., the synergetic effect).
references
[1] Iba H, Akiba E. Hydrogen absorption and microstructure in BCC alloys with C14-type Laves phase. J Jpn Inst Met 1994;58:1225e32. [2] Akiba E, Iba H. Hydrogen absorption by Laves phase related BCC solid solution. Intermetallics 1998;6:461e70. [3] Young K, Nei J. The current status of hydrogen storage alloy development for electrochemical applications. Materials 2013;6:4574e608. [4] Inoue H, Arai S, Iwakura C. Crystallographic and electrochemical characterization of TiV4xNix alloys for nickelemetal hydride batteries. Electrochim Acta 1996;41:937e9. [5] Yu XB, Wu Z, Xia BJ, Xu NX. A TieV-based bcc phase alloy for use as metal hydride electrode with high discharge capacity. J Chem Phys 2004;121:987e90. [6] Chen N, Li R, Zhu Y, Liu Y, Pan H. Electrochemical hydrogenation and dehydrogenation mechanisms of the TieV base multiphase hydrogen storage electrode alloy. Acta Metall Sin 2004;40:1200e4. [7] Iba H, Akiba E. The relation between microstructure and hydrogen absorbing property in Laves phase-solid solution multiphase alloys. J Alloys Compd 1995;231:508e12. [8] Ro¨nnebro E, Nore´us D, Sakai T, Tsukahara M. Structural studies of a new Laves phase alloy (Hf,Ti)(Ni,V)2 and its very stable hydride. J Alloys Compd 1995;231:90e4.
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10
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 e1 1
[9] Tsukahara M, Takahashi K, Mishima T, Isomura A, Sakai T. V-based solid solution alloys with Laves phase network: hydrogen absorption properties and microstructure. J Alloys Compd 1996;236:151e5. [10] Qiu S, Chu H, Zhang Y, Sun D, Song X, Sun L, et al. Electrochemical kinetics and its temperature dependence behaviors of Ti0.17Zr0.08V0.35Cr0.10Ni0.30 alloy electrode. J Alloys Compd 2009;471:453e6. [11] Iba H, Akiba E. Hydrogen absorption and modulated structure in TieVeMn alloys. J Alloys Compd 1997;253e254:21e4. [12] Basak S, Shashikala K, Kulshreshtha SK. Hydrogen absorption characteristics of Zr substituted Ti0.85VFe0.15 alloy. Int J Hydrogen Energy 2008;33:350e5. [13] Shashikala K, Banerjee S, Kumar A, Pai MR, Pillai CGS. Improvement of hydrogen storage properties of TiCrV alloy by Zr substitution for Ti. Int J Hydrogen Energy 2009;34:6684e9. [14] Yang XW, Li JS, Zhang TB, Hu R, Xue XY, Fu HZ. Role of defect structure on hydrogenation properties of Zr0.9Ti0.1V2 alloy. Int J Hydrogen Energy 2011;36:9318e23. [15] Tsukahara M, Takahashi K, Mishima T, Sakai T, Miyamura H, Kuriyama N, et al. Metal hydride electrodes based on solid solution type alloy TiV3Nix (0 & x & 0.75). J Alloys Compd 1995;226:203e7. [16] Tsukahara M, Takahashi K, Mishima T, Isomura A, Sakai T. Heat-treatment effects of V-based solid solution alloy with TiNi-based network structure on hydrogen storage and electrode properties. J Alloys Compd 1996;243:133e8. [17] Tsukahara M, Takahashi K, Mishima T, Isomura A, Sakai T. Vanadium-based solid solution alloys with threedimensional network structure for high capacity metal hydride electrodes. J Alloys Compd 1997;253e254:583e6. [18] Huot J, Akiba E, Ogura T, Ishido Y. Crystal structure, phase abundance and electrode performance of Laves phase compounds (Zr, A)V0.5Ni1.1Mn0.2Fe0.2 (A h Ti, Nb or Hf). J Alloys Compd 1995;218:101e9. [19] Yoshida M, Akiba E. Hydrogen absorbing-desorbing properties and crystal structure of the ZreTieNieMneV AB2 Laves phase alloys. J Alloys Compd 1995;224:121e6. [20] Huot H, Akiba E, Ishido Y. Crystal structure of multiphase alloys (Zr,Ti)(Mn,V)2. J Alloys Compd 1995;231:85e9. [21] Kuriiwa T, Tamura T, Amemiya T, Fuda T, Kamegawa A, Takamura H, et al. New V-based alloys with high protium absorption and desorption capacity. J Alloys Compd 1999;293e295:433e6. [22] Miao H, Gao M, Liu Y, Lin Y, Wang J, Pan H. Effects of Y substitution for Ti on the microstructure and electrochemical properties of TieVeFe-based hydrogen storage alloys. J Electrochem Soc 2007;154:A1010e4. [23] Qiu S, Chu H, Zhang J, Zhang Y, Sun L, Xu F, et al. Effect of La partial substitution for Zr on the structural and electrochemical properties of Ti0.17Zr0.08exLaxV0.35Cr0.1Ni0.3 (x ¼ 0e0.04) electrode alloys. Int J Hydrogen Energy 2009;34:7246e52. [24] Cho S, Enoki H, Akiba E. Effect of Fe addition on hydrogen storage characteristics of Ti0.16Zr0.05Cr0.22V0.57 alloy. J Alloys Compd 2000;307:304e10. [25] Dou T, Wu Z, Mao J, Xu N. Application of commercial ferrovanadium to reduce cost of TieV-based BCC phase hydrogen storage alloys. Mater Sci Eng A 2008;476:34e8. [26] Santos SF, Huot J. Hydrogen storage in TiCr1.2(FeV)x BCC solid solutions. J Alloys Compd 2009;472:247e51. [27] Yu J, Liu B, Cho K, Lee J. The effects of partial substitution of Mn by Cr on the electrochemical cycle life of TieZreVeMneNi alloy electrodes of a Ni/MH battery. J Alloys Compd 1998;278:283e90. [28] Yan Y, Chen Y, Liang H, Wu C, Tao M. The effect of Si on V30Ti35Cr25Fe10 BCC hydrogen storage alloy. J Alloys Compd 2007;441:297e300.
[29] Huang Z, Cuevas F, Liu X, Jiang L, Wang S, Latroche M, et al. Effects of Si addition on the microstructure and the hydrogen storage properties of Ti26.5V45Fe8.5Cr20Ce0.5 BCC solid solution alloys. Int J Hydrogen Energy 2009;34:9385e92. [30] Martı´nez A, dos Santos DS. Influence of the substitution of V by Nb in the structure and properties of hydrogen absorption/desorption of TiCr1.1V0.9 alloy. J Alloys Compd 2012;536(Suppl. 1):S231e5. [31] Martı´nez A, dos Santos DS. Hydrogen absorption/desorption properties in the TiCrV based alloys. Mater Res 2012;15:809e12. [32] Kamegawa A, Shirasaki K, Tamura T, Kuriiwa T, Takamura H, Okada M. Crystal structure and protium absorption properties of TieCreX alloys. Mater Trans 2002;43:470e3. [33] Uno M, Takahashi K, Maruyama T, Muta H, Yamanaka S. Hydrogen solubility of BCC titanium alloys. J Alloys Compd 2004;366:213e6. [34] Shirasaki K, Kuriiwa T, Tamura T, Kamegawa A, Takamura H, Okada M. Effects of Al addition on structures and protium absorptionedesorption properties of TieCr alloys. Mater Trans 2002;43:1173e7. [35] Mi J, Lu¨ F, Liu X, Jiang L, Li Z, Wang S. Enhancement of cerium and hydrogen storage property of a low-cost TieV based BCC alloy prepared by commercial ferrovanadium. J Rare Earths 2010;28:781e4. [36] Liu Y, Zhang S, Li R, Gao M, Zhong K, Miao H, et al. Electrochemical performances of the Pd-added TieV-based hydrogen storage alloys. Int J Hydrogen Energy 2008;33:728e34. [37] Sapru K, Hong K, Fetcenko MA, Venkatesan S, Inventors; Energy Conversion Devices, Inc., Assignee. Hydrogen storage materials and methods of sizing and preparing the same for electrochemical applications. United States patent US 4551400; 1985 Nov 5. [38] Reichman B, Venkatesan S, Fetcenko MA, Jeffries K, Stahl S, Bennett C, Inventors; Energy Conversion Devices, Inc., Assignee. Activated rechargeable hydrogen storage electrode and method. United States patent US 4716088; 1987 Dec 29. [39] Venkatesan S, Reichman B, Fetcenko MA, Inventors; Energy Conversion Devices, Inc., Assignee. Enhanced charge retention electrochemical hydrogen storage alloys and an enhanced charge retention electrochemical cell. United States patent US 4728586; 1988 Mar 1. [40] Ovshinsky SR, Fetcenko MA, Ross J. A nickel metal hydride battery for electric vehicles. Science 1993;260:176e81. [41] Lee H, Lee K, Lee J, Lee J. The Ti-based metal hydride electrode for NieMH rechargeable batteries. J Alloys Compd 1996;239:63e70. [42] Kim JS, Paik CH, Cho WI, Cho BW, Yun KS, Kim SJ. Corrosion behaviour of Zr1exTixV0.6Ni1.2M0.2 (M ¼ Ni, Cr, Mn) AB2-type metal hydride alloys in alkaline solution. J Power Sources 1998;75:1e8. [43] Zhang QA, Lei YQ, Wang CS, Wang FS, Wang QD. Structure of the secondary phase and its effects on hydrogen-storage properties in a Ti0.7Zr0.2V0.1Ni alloy. J Power Sources 1998;75:288e91. [44] Zhu Y, Pan H, Gao M, Liu Y, Wang Q. A study on improving the cycling stability of (Ti0.8Zr0.2)(V0.533Mn0.107Cr0.16Ni0.2)4 hydrogen storage electrode alloy by means of annealing treatment: II. Effects on the electrochemical properties. J Alloys Compd 2003;348:301e8. [45] Chen L, Guo R, Lei Y, Li L, Wang Q. Structural and electrochemical characteristics of V2.1TiNi0.5Hf0.05M0.152 (M ¼ Co, Cr) hydrogen storage electrode alloys. Mater Chem Phys 2005;92:554e8. [46] Pan H, Li R, Liu Y, Gao M, Miao H, Lei Y, et al. Structure and electrochemical properties of the Fe substituted TieV-based hydrogen storage alloys. J Alloys Compd 2008;463:189e95.
Please cite this article in press as: Young K, et al., Structural, hydrogen storage, and electrochemical properties of Laves phaserelated body-centered-cubic solid solution metal hydride alloys, International Journal of Hydrogen Energy (2014), http:// dx.doi.org/10.1016/j.ijhydene.2014.01.134
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 e1 1
[47] Gao M, Zhang S, Miao H, Liu Y, Pan H. Pulverization mechanism of the multiphase TieV-based hydrogen storage electrode alloy during charge/discharge cycling. J Alloys Compd 2010;489:552e7. [48] Chai Y, Zhao M. Structure and electrochemical properties of Ti0.25V0.35Cr0.40exNix (x ¼ 0.05e0.40) solid solution alloys. Int J Hydrogen Energy 2005;30:279e83. [49] Qiao Y, Zhao M, Li M, Zhu X, Cao G. Microstructure and electrochemical performance of Ti0.17Zr0.08V0.34Pd0.01Cr0.1Ni0.3 electrode alloy. Scr Mater 2006;55:279e82. [50] Qiao Y, Zhao M, Zhu X, Cao G. Effect of cerium on microstructure and electrochemical performance of TieVeCreNi electrode alloy. J Rare Earths 2007;25:341e7. [51] Chai YJ, Zhao MS, Wang N. Crystal structural and electrochemical properties of Ti0.17Zr0.08V0.35Cr0.10Ni0.30xMnx (x ¼ 0e0.12) alloys. Mater Sci Eng B 2008;147:47e51. [52] Wang N, Chai Y. Electrochemical properties and reaction mechanism of TieV multiphase alloy at high temperature. Rare Met Mater Eng 2011;40:1519e21. [53] Yin W, Zhao M. Structure and electrochemical characteristics of TiV1.1Mn0.9Nix (x ¼ 0.1e0.7) alloys. Electrochim Acta 2007;52:2723e8. [54] Li S, Zhao M, Wang L, Liu Y, Wang Y. Structures and electrochemical characteristics of Ti0.26Zr0.07V0.24Mn0.1Ni0.33Mox (x ¼ 0e0.1) hydrogen storage alloys. Mater Sci Eng B 2008;150:168e74. [55] Li S, Zhao M, Zhai J, Wang F. Effect of substitution of chromium for nickel on structure and electrochemical characteristic of Ti0.26Zr0.07V0.24Mn0.1Ni0.33 multi-phase hydrogen storage alloy. Mater Chem Phys 2009;113:96e102. [56] Yu X, Wu Z, Li F, Xia B, Xu N. Body-centered-cubic phase hydrogen storage alloy with improved capacity and fast activation. Appl Phys Lett 2004;84:3199e201. [57] Feng S, Wu Z, Yu X, Chen C, Chen K, Chen T, et al. Effect of Zr substitution for Ti on Ti19.5V40Mn16.2Cr9.8Ni14.5 alloy I: microstructure and hydrogen storage property. Rare Met Mater Eng 2007;36:794e8. [58] Gue´guen A, Joubert JM, Latroche M. Influence of the C14 Ti35.4V32.3Fe32.3 Laves phase on the hydrogenation properties of the body-centered cubic compound Ti24.5V59.3Fe16.2. J Alloys Compd 2011;509:3013e8. [59] Young K, Fetcenko MA, Li F, Ouchi T. Structural, thermodynamic, and electrochemical properties of TixZr1x(VNiCrMnCoAl)2 C14 Laves phase alloys. J Alloys Compd 2008;464:238e47. [60] Young K, Fetcenko MA, Koch J, Morii K, Shimizu T. Studies of Sn, Co, Al, and Fe additives in C14/C15 Laves alloys for NiMH battery application by orthogonal arrays. J Alloys Compd 2009;486:559e69. [61] Young K, Ouchi T, Reichman B, Mays W, Regmi R, Lawes G, et al. Optimization of Co-content in C14 Laves phase multicomponent alloys for NiMH battery application. J Alloys Compd 2010;489:202e10.
11
[62] Yoo J, Shim G, Yoon J, Cho S. Effects of substituting Al for Cr in the Ti0.32Cr0.43V0.25 alloy on its microstructure and hydrogen storage properties. Int J Hydrogen Energy 2009;34:1463e7. [63] Young K, Fetcenko MA, Li F, Ouchi T, Koch J. Effect of vanadium substitution in C14 Laves phase alloys for NiMH battery application. J Alloys Compd 2009;468:482e92. [64] Young K, Ouchi T, Fetcenko MA. Pressureecompositionetemperature hysteresis in C14 Laves phase alloys: part 1. Simple ternary alloys. J Alloys Compd 2009;480:428e33. [65] Young K, Ouchi T, Mays W, Reichman B, Fetcenko MA. Pressureecompositionetemperature hysteresis in C14 Laves phase alloys: part 2. Applications in NiMH batteries. J Alloys Compd 2009;480:434e9. [66] Young K, Ouchi T, Fetcenko MA. Pressureecompositionetemperature hysteresis in C14 Laves phase alloys: part 3. Empirical formula. J Alloys Compd 2009;480:440e8. [67] Yan Y, Chen Y, Liang H, Zhou X, Wu C, Tao M, et al. Hydrogen storage properties of VeTieCreFe alloys. J Alloys Compd 2008;454:427e31. [68] Chen KC, Peterson EJ, Thoma DJ. HfCo2 Laves phase intermetallicsdpart I: solubility limits and defect mechanisms. Intermetallics 2001;9:771e83. [69] Young K, Ouchi T, Huang B, Reichman B, Blankenship R. Improvement in 40 C electrochemical properties of AB2 metal hydride alloy by silicon incorporation. J Alloys Compd 2013;575:65e72. [70] Bendersky LA, Wang K, Boettinger WJ, Newbury DE, Young K, Chao B. Examination of multiphase (Zr,Ti)(V,Cr,Mn,Ni)2 NieMH electrode alloys: part II. Solid-state transformation of the interdendritic B2 phase. Metall Mater Trans A 2010;41:1891e906. [71] Boettinger WJ, Newbury DE, Wang K, Bendersky LA, Chiu C, Kattner UR, et al. Examination of multiphase (Zr,Ti)(V,Cr,Mn,Ni)2 NieMH electrode alloys: part I. Dendritic solidification structure. Metall Mater Trans A 2010;41:2033e47. [72] Sun D, Jiang J, Lei Y, Liu W, Wu J, Wang Q, et al. Effects of measurement factor on electrochemical capacity of some hydrogen storage alloys. Mater Sci Eng B 1995;30:19e22. [73] Young K, Nei J, Huang B, Ouchi T, Fetcenko MA. Studies of Ti1.5Zr5.5V0.5(MxNi1x)9.5 (M ¼ Cr, Mn, Fe, Co, Cu, Al): part 2. Hydrogen storage and electrochemical properties. J Alloys Compd 2010;501:245e54. [74] Young K, Chao B, Huang B, Nei J. Studies on the hydrogen storage characteristic of La1xCex(NiCoMnAlCuSiZr)5.7 with a B2 secondary phase. J Alloys Compd 2014;585:760e70. [75] Li F, Young K, Ouchi T, Fetcenko MA. Annealing effects on structural and electrochemical properties of (LaPrNdZr)0.83Mg0.17(NiCoAlMn)3.3 alloy. J Alloys Compd 2009;471:371e7.
Please cite this article in press as: Young K, et al., Structural, hydrogen storage, and electrochemical properties of Laves phaserelated body-centered-cubic solid solution metal hydride alloys, International Journal of Hydrogen Energy (2014), http:// dx.doi.org/10.1016/j.ijhydene.2014.01.134