MH battery prepared by centrifugal casting and gas atomization

Journal of Alloys and Compounds 496 (2010) 669–677

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Study of AB2 alloy electrodes for Ni/MH battery prepared by centrifugal casting and gas atomization K. Young a,∗ , J. Koch a , T. Ouchi a , A. Banik b , M.A. Fetcenko a a b

Energy Conversion Devices Inc./Ovonic Battery Company, 2983 Waterview Drive, Rochester Hills, MI 48309, USA Special Metal Corporation, 100 Industry Lane, Princeton, KY 42445, USA

a r t i c l e

i n f o

Article history: Received 29 January 2010 Received in revised form 17 February 2010 Accepted 19 February 2010 Available online 3 March 2010 Keywords: Hydrogen absorbing materials Transition metal alloys Metal hydride electrode Electrochemical reactions

a b s t r a c t Centrifugal casting and gas atomization processes were applied on multiple phase AB2 alloys and compared to the conventional melt-and-cast. Four different compositions were chosen for this study. The roles of Zr, Mn, Cr, and Ni in various battery aspects are identified. Cooling speed was found to be crucial for the C14 and C15 phase abundances. As the cooling speed increased from 102 to 104 degrees per second, a higher percentage of C15 was found. The centrifugal casting process provided better cycle life and low temperature performance with the only trade-off being slower activation. The gas atomization process can achieve lower production cost due to the elimination of a grinding procedure and extended cycle life, but suffers from higher bulk oxygen content and thicker surface oxide, and thus inferior in all battery performance characteristics other than cycle life and charge retention. © 2010 Elsevier B.V. All rights reserved.

1. Introduction While most of the commercial nickel metal hydride (Ni/MH) batteries use misch metal based AB5 alloys as the metal hydride electrode, Laves phase based AB2 alloys offer an opportunity in achieving higher energy density [1–25]. However, AB2 alloy made by conventional melt-and-cast (MC) process is difficult to activate and its cycle life is inferior [8,26–45]. Various manufacturing processes have been tried before to improve the battery performance using AB2 electrode [18,46–54]. In this paper, we focus on two alternative processes: centrifugal casting (CC) and gas atomization (GA). Results from melt-spin and mechanical alloys with higher quenching rates will be published elsewhere. All of the earlier literature of CC prepared hydrogen storage alloys or rare-earth based magnets were from US Patent publications [55–61]. The process involves pouring of melting liquid on a rotating surface and the subsequently dispersing liquid onto a cooled surface by using a centrifugal force. The final product can be either a layer [56] or flakes [57]. The typical cooling speed is about 1–3 × 103 degree per second for a flake thickness of 0.1–0.3 mm [57,58]. The main purpose of applying this technique is to improve the phase uniformity within the alloy [56,57,59,61], and to extend cycle life [56]. It can also be designed to fabricate small secondary phase grains finely dispersed in the flakes [58]. Centrifugal impact

∗ Corresponding author. Tel.: +1 248 293 7000/854 2355; fax: +1 248 299 4520/853 4296. E-mail address: [email protected] (K. Young). 0925-8388/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2010.02.161

atomization, a similar process with a liquid stream hitting the surface of rotating blades, was also used to make amorphous powder capable of storing hydrogen [62]. The GA process requires passing a low-viscosity liquid through a small orifice with an applied pressure difference. The liquid is further disintegrated into a spray of fine droplets by the impinging the liquid metal stream with high pressure gas. The droplets solidify into spherical particles during their travel in the inert gas environment [63]. The typical quenching speed is a 3-5 ×104 degree per second [64]. While most of the early work on applying gas atomization to hydrogen storage alloys were focused on misch metal based AB5 [64–92], some work on AB2 alloys were also reported [93–98]. The advantage of a gas atomization process is to reduce production costs by eliminating the pulverization process. In addition, the spherical shape of the gas atomized powder increases packing density [55,80,87], and reduces stress created in the alloy particle during hydride–dehydride process to enhance cycle life [68,74,78,86,92]. The fast quenching rate of this process allows incorporation of some elements, such as Cr [85], Fe and Cu [68] into AB5 , and enables the selection of crystallinity of the main phase [70,94] and certain secondary phases [80,95]. The shortcomings of the gas atomization process are twofold: a thicker surface oxide layer hindering the activation process [64,68,83,88,92,94] and a higher oxygen content in the bulk to reduce hydrogen storage capacity [68,86,92]. There was also a two-stage powder preparation method which combines gas atomization and centrifugal casting techniques for AB5 hydrogen storage alloy [99]. This process enhances the cycle life and allows the incorporation of Cr into the matrix.

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K. Young et al. / Journal of Alloys and Compounds 496 (2010) 669–677 Table 1 Compositions of design, mother ingot, gas atomized powder, and centrifugal casting flakes in atomic percentage determined by ICP.

A1 target A1 mother ingot A1 gas atomized A1 CC A2 target A2 mother ingot A2 gas atomized A2 CC A3 target A3 mother ingot A3 gas atomized A3 CC A4 target A4 mother ingot A4 Gas atomized A4 CC

Fig. 1. SEM secondary electron image of typical morphology of gas atomization powder.

2. Experimental Conventional MC process was performed in a tilt-poured vacuum induction furnace using 25 kg of raw material in an elemental form (except for Sn which is already present in the Zircaloy) to achieve the designed composition. The system includes a MgO crucible, an alumina tundish, and a steel pancake mold. The average thickness of the ingot is about 4 cm. The ingot first went through a hydride–dehydride process to reduce the size and create internal cracks, followed by a centrifugal pulverization to a −200 mesh powder. The centrifugal casting was done by tilt-pouring a 2 kg of molten liquid (melt from raw material) through a tundish with a small nozzle onto the flat surface of a spinning turntable (50 cm diameter and 2 cm thick) made from 25 kg of copper with no water cooling (for a schematic diagram of the set-up, see Fig. 1 in Ref. [55]). The rotational speed of the turntable was about 100 rpm. The obtained flakes were collected from a pan installed at the bottom of the vacuum chamber. The flakes then went through a mechanical grinding process using a motorized mortar-and-pestle to achieve a −200 mesh powder. Gas atomization was performed by induction melting using a pre-alloyed master alloy from conventional melt-and-cast. The molten alloy is poured into a tundish and nozzle assembly. Inert gas impinges on the molten stream dispersing the stream into fine droplets which are collected at the base of the atomization tower. The spherical nature of the powder achieved directly from gas atomization process can be best represented by the scanning electron microscope (SEM) micrograph in Fig. 1. The bulk of the particles are spherical in morphology. Isolated particles can contain satellite particles which adhere to the surface during solidification. The final −100 mesh powder was obtained by sifting. The spherical shape of the GA powder allows us to use a larger particle without concern of punching through the separator from sharp corners normally seen using MC negative particles. The chemical composition of each sample was examined by a Varian Liberty 100 inductively-coupled plasma (ICP) system. A Philips X’Pert Pro X-ray Diffractometer (XRD) was used to study each alloys microstructure, and a JOEL-JSM6320F SEM with X-ray Energy-Dispersive Spectroscopy (EDS) capability was used to study phase distribution and composition. Pressure–Concentration–Temperature (PCT) characteristics were measured using a Suzuki-Shokan multi-channel PCT system. In the PCT analysis, each sample was first activated by a 2-h thermal cycle between 300 ◦ C and room temperature at 25 atm H2 pressure. Details of both electrode and cell preparations as well as measurement methods have been reported before [100].

21.5 20.6 21.0 21.6 21.5 21.4 20.6 21.1 19.5 18.9 19.0 19.5 21.5 20.8 20.9 20.7

Ti

V

Ni

Cr

Mn

Co

Al

Sn

Fe

12.0 12.4 12.0 11.7 12.0 11.8 12.2 12.1 14.0 14.1 14.1 13.6 12.0 12.0 12.2 12.2

10.0 10.1 9.8 10.0 10.0 9.4 10.0 10.0 10.0 10.1 10.0 9.9 10.0 10.1 10.0 10.0

32.2 32.5 32.7 32.1 36.2 36.3 36.0 36.4 36.2 36.7 36.4 36.9 32.2 32.6 32.5 32.8

8.5 8.5 8.6 8.6 4.5 4.9 4.7 4.5 4.5 4.5 4.8 4.5 6.5 6.7 6.7 6.4

13.6 13.6 13.6 13.8 13.6 13.8 14.4 13.9 13.6 13.6 13.6 13.4 15.6 15.6 15.6 15.5

1.5 1.4 1.4 1.5 1.5 1.7 1.3 1.4 1.5 1.5 1.4 1.5 1.5 1.4 1.3 1.4

0.4 0.6 0.6 0.7 0.4 0.4 0.5 0.4 0.4 0.4 0.5 0.4 0.4 0.5 0.5 0.6

0.3 0.2 0.2 0.2 0.3 0.2 0.2 0.2 0.3 0.1 0.1 0.2 0.3 0.2 0.2 0.2

0.0 0.1 0.1 0.1 0.0 0.1 0.1 0.1 0.0 0.1 0.1 0.1 0.0 0.1 0.1 0.1

capacity, plateau pressure, slope, and hysteresis can be measured from the 30 ◦ C PCT isotherms obtained on alloys prepared by conventional melt-and-cast, and plotted in Fig. 2. Alloy A3 with the smallest Zr/Ti ratio has the highest plateau pressure. Alloy A4 with the largest Mn content moves the average electron density closer to the bottom of the parabola curve and corresponds to the largest PCT hysteresis, which agrees well with our previous model [105]. Powders were fabricated from conventional melt-and-cast plus grinding, centrifugal casting plus grinding, and gas atomization. The chemical composition determined by ICP for each batch of powder is listed in Table 1 together with the original design intent. Small deviations in Zr and Ti may come from the non-uniformity of the produced alloys. The elevations in Al and Fe contents were the result of contamination from the alumina tundish and steel mold. Lower Sn levels were due to the particular Zircaloy lot with varying Sn content received from the vendor.

3. Results and discussion Four compositions, A1, A2, A3, and A4 were chosen for this study of preparation methods with different cooling speeds. The design compositions are listed in Table 1. The concentrations of Zr, Ti, and V were optimized through the balance between capacity and high rate dischargeability (HRD) [101,102]. Compositions of Cr and Mn were optimized for consideration between corrosion resistance and ease of formation [103,104]. Using Co, Al, and Sn as additives was the result of an orthogonal array study [100]. The hydrogen storage

Fig. 2. PCT isotherms measured at 30 ◦ C for alloys A1, A2, A3 and A4 prepared by conventional melt-and-cast.

K. Young et al. / Journal of Alloys and Compounds 496 (2010) 669–677

Fig. 3. XRD spectra using Cu-K␣ as the radiation source for alloys A1 by conventional melt-and-cast (a), centrifugal casting (b), gas atomization (c), A3 by conventional melt-and-cast (d), centrifugal casting (e), and gas atomization (f).

3.1. Microstructure The microstructure of alloys A1 and A3 prepared by three different methods were studied using XRD analysis. The resulting spectra are plotted in Fig. 3 and the important parameters obtained from this analysis are listed in Table 2. In the case of the A1 composi-

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tion, the crystal structure shifts from mostly C14-dominating, to a C14/C15 mixture as the cooling speed increases. A small amount of the TiNi phase becomes visible in the GA sample (Fig. 3c). The CC sample gives the largest lattice constants while the MC sample shows the smallest lattice constants. The a/c ratio remains the same independent of the cooling rate. The full-width at half-maximum (FWHM) of the XRD C14 (1 0 3) peak and (2 0 1) peak clearly demonstrates a narrowing trend (corresponding to a larger grain size by the Scherrer equation [106]) as the cooling speed increases. The microstructure evolution of the A3 composition is similar to, but not exactly the same, as the A1 composition. Both powders show varying amounts of C14, C15 and TiNi phases. The abundance of C15 phase increases with the increase in cooling rate while the abundance of TiNi phase remains roughly the same. The C14 unit cell volume in A3 comparing three different preparation methods is on the same order as A1: CC > GA > MC. It is thus speculated that without the consideration of possible oxygen contamination in the surface and/or in the bulk of the alloy powder, the hydrogen storage capacity should follow the same order as the unit cell volume when comparing similar alloys [107]. The change in lattice parameters in the A3 composition is not isotropic, which is different from the A1 composition. As the cooling speed increases, the a/c ratio first increases and then decreases. Therefore, we cannot conclude there is an influence of cooling rate on the a/c ratio. From Table 2, we can also observe the FWHM of the XRD peak of the A3 composition decreases first and then increases as the cooling speed increases. Therefore, we reach the conclusion that while both CC and GA processes can increase the size of C14 grain, it is not clear which one gives rise to the larger grain size. In this range of cooling rate (102 –104 ), it is not the only factor dictating the grain size since it is known that higher cooling rates (106 ) will reduce the grain size [108].

Table 2 Summary of XRD results for alloys A1 and A3 by three different preparation methods. Composition #

Preparation

Lattice cont. a in C14 (Å)

Lattice cont. c in C14 (Å)

a/c ratio

C14 unit cell volume (Å3 )

FWHM of C14 (103)

FWHM of C14 (201)

Phase summary

A1 A1 A1 A3 A3 A3

Conventional MC Centrifugal cast Gas atomization Conventional MC Centrifugal cast Gas atomization

4.9785 4.9815 4.9799 4.9630 4.9868 4.9637

8.1222 8.1261 8.1247 8.0853 8.0774 8.1023

0.6129 0.6130 0.6129 0.6138 0.6174 0.6126

174.34 174.63 174.49 172.47 173.95 172.88

0.308 0.229 0.205 0.270 0.213 0.240

0.292 0.222 0.212 0.248 0.215 0.222

Mainly C14 C14 + C15 C14 + C15 + TiNi C14 + C15 + TiNi C14 + C15 + TiNi C14 + C15 + TiNi

Table 3 Chemical composition determined by EDS in the circled area in Figs. 4–7. Area

Zr

Ti

V

Cr

Mn

Co

Ni

Sn

Al

O

A1-1 A1-2 A1-3 A1-4 A1-5 A1-6 A2-1 A2-2 A2-3 A2-4 A2-5 A3-1 A3-2 A3-3 A3-4 A3-5 A4-1 A4-2 A4-3 A4-4 A4-5 A4-6

96.7 62.4 20.2 20.2 18.8 31.8 97.4 20.4 21.8 22.1 32.6 25.7 20.4 20.9 14.4 32.8 24.2 21.3 18.7 21.4 20.7 33.5

0 12.9 13.4 13.9 28.9 0 0 24.6 12.7 10.8 0 18 13.6 11.1 28.4 0 13.3 11.6 24.6 11.8 12.5 0

0 2.3 9.2 8.6 0 0 0 0.1 8.0 10.3 0 0.5 6.6 12.7 1.5 0 5.7 11 0.9 10.4 9.9 0

0 1.4 6.0 5.7 0 0 0 0.3 3.7 4.8 0 0 2.1 6.3 0.4 0 2.9 6.5 0.3 6.0 5.1 0

0 3.6 12.9 12.4 0 0 0 3.7 13.4 16 0 3.5 10.5 16.9 4.8 0 9.9 16.4 4.2 16.1 15.8 0

0 0.4 1.5 1.6 0 0 0 1.5 1.7 2.0 0 0.7 1.1 2.0 1.3 0 0.8 1.7 1.1 1.5 1.7 0

0 17.1 35.4 36.3 52.3 0 0 49.4 37.3 32.8 0 44.8 45.0 29.2 48.5 0 38.8 30.3 48.4 31.2 32.7 0

3.3 0 0 0 0 0 2.6 0 0 0 0 6.8 0 0 0 0 3.3 0 0.6 0 0 0

0 0 1.5 1.3 0 0 0 0 1.4 1.3 0 0 0.8 0.9 0.8 0 1.1 1.2 1.3 1.6 1.7 0

0 0 0 0 0 68.2 0 0 0 0 67.4 0 0 0 0 67.2 0 0 0 0 0 66.5

B/A ratio

Average e/a

Phase

0.33 1.98 1.93 1.10

5.21 6.79 6.82 7.14

1.22 1.90 2.04

7.16 6.87 6.74

Zr Zr + AB2 AB2 AB2 TiNi ZrO2 Zr TiNi AB2 AB2

1.29 1.94 2.13 1.34

6.83 7.17 6.60 7.14

TiNiSnx AB2 AB2 TiNi

1.67 2.04 1.31 2.01 2.01

6.77 6.62 7.09 6.64 6.71

TiNiSnx AB2 TiNi AB2 AB2 ZrO2

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Fig. 4. SEM backscattering electron image cross-section micrographs for alloy A1 prepared by conventional melt-and-cast (a), centrifugal casting (b), and gas atomization (c). The bar on the right lower corner indicates a length of 20 ␮m. Numbers in each micrograph indicate that EDS analyses were done in these areas and chemical compositions are summarized in Table 3.

The phase compositions in these products made by different preparation methods were further analyzed by using SEM backscattering electron imaging and the results are shown in Fig. 4 (A1), Fig. 5 (A2), Fig. 6 (A3), and Fig. 7 (A4). In the case of MC samples, areas with distinctive contrast corresponding to different average atomic weight can be easily identified. The chemical composition in the area within the numbered circle has been studied by EDS analysis and the results are listed in Table 3. The main phase is AB2 -stoichiometric. The phase with a higher average electron density (e/a) value is believed to be C15 phase. Similar local phase alteration due to local composition variation in AB2 alloy has been reported before [95]. The phase with a B/A ratio substantially below 2.0 is either a TiNi phase (no Sn) [109,110] or ZrTiSn phase (high Sn content) [100] secondary phase. The solubility of V in TiNi phase is almost zero, which is consistent with our previous report [100]. Occasionally bright spots from metallic Zr and dark spots of ZrO2 are also observed. In the case of micrograph from CC samples, the phase segregation is smaller than the EDS detectable area. Other than some bright spots (Zr) and dark spots (ZrO2 ), no other large contrast can be seen. All GA samples show uniform contrast except for a small amount of dark spots (ZrO2 ) due to the resolution limits of the SEM analysis. High-resolution transmission electron microscope micrograph from similar GA-

prepared C14/C15 mixed alloys shows a clear grain size of about 50–100 Å [98]. 3.2. Half-cell measurement In the flooded half-cell measurement, all electrodes were activated in a 110 ◦ C 30% KOH solution for 4.5 h before being connected to a counter electrode. They were first discharged with a small current (C/70 rate). Such obtained initial capacity is understood to result from the absorption of hydrogen generated from the metal oxidation during the hot alkaline activation process [111]. The amount of initial discharge has been correlated to the ease with which the surface metal can be oxidized and is listed in Table 4. Comparing three different preparation methods vs. four different compositions, the MC samples all show similar good activation behavior. With the CC process, the A1 and A2 compositions show much less initial discharge and imply a greater resistance to activation, which can be related to the relatively lower Ti and Mn contents in these two alloys. In the comparison of the GA samples, only the A1 composition shows a small initial discharge due to its highest Cr content protecting the alloy surface from oxidation [104]. In A1 and A2 composition, CC samples are more difficult to activate than the GA samples due to the even thicker native oxide

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Fig. 5. SEM backscattering electron image cross-section micrographs for alloy A2 prepared by conventional melt-and-cast (a), centrifugal casting (b), and gas atomization (c). The bar on the right lower corner indicates a length of 20 ␮m. Numbers in each micrograph indicate that EDS analyses were done in these areas and chemical compositions are summarized in Table 3. Table 4 Composition, preparation methods, half-cell, and full-cell measurement results. The first three specific capacity data are in mAh/g and peak powers are in W/kg. The high rate dischargeability (HRD) is defined by the ratio between the capacities measured at C/7 and C/70 rates. Composition #

Preparation

First Disch.

Cap @ C/7

Cap @ C/70

A1 A1 A1 A2 A2 A2 A3 A3 A3 A4 A4 A4

MC CC GA MC CC GA MC CC GA MC CC GA

29 3 11 29 6 32 39 49 35 36 37 36

386 376 319 372 384 346 369 375 331 398 397 354

426 415 367 389 401 366 378 385 349 409 409 377

HRD 91% 91% 87% 96% 96% 95% 98% 97% 95% 97% 97% 94%

formed during solidification with less surface area created by CC process. The discharge capacities at a higher rate (C/7) and a lower rate (C/70) together with their ratio, which is a measurement of high rate dischargeability, are listed in Table 4. The order of the full capacity of the MC alloys is: A1 (426) > A4 (409) > A2 (389) > A3 (378 mAh/g), which totally agrees with the order of the maximum

2 C (30 ◦ C) 67% 65% 47% 67% 79% 52% 68% 71% 66% 26% 47% 51%

0.5 C (−10 ◦ C) 88% 87% 67% 65% 86% 68% 84% 83% 67% 51% 68% 70%

PP @ 80%RT 136 133 97 157 147 115 152 142 125 115 130 100

PP @ 80%0 ◦ C 90 85 55 90 94 79 103 90 85 73 80 20

CR @ 7 days 76% 80% 94% 77% 70% 65% 61% 59% 65% 73% 70% 73%

Cycle life 425 550 550 380 550 470 250 350 470 300 350 490

hydrogen storage capacities found in the PCT isotherms (Fig. 2). The order of the CC samples follows the same trend. The order of the GA samples, A4 > A1 ∼ A2 > A3, is slightly different due to the more catalytic nature of the active surface of the GA-A4 sample judging from the large first discharge current. When comparing the full capacity among three preparation methods with the same composition, the

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Fig. 6. SEM backscattering electron image cross-section micrographs for alloy A3 prepared by conventional melt-and-cast (a), centrifugal casting (b), and gas atomization (c). The bar on the right lower corner indicates a length of 20 ␮m. Numbers in each micrograph indicate that EDS analyses were done in these areas and chemical compositions are summarized in Table 3.

values from the CC process can be higher (A2 and A3), lower (A1) than, or about the same (A4) as the MC samples. GS samples have the lowest capacity due to the higher oxygen content in the bulk of GA sample [68,86,92]. Comparing the HRD of the MC samples with different composition, the order (A3 > A4 > A2 > A1) is almost the same as the plateau pressure in the PCT isotherm (A3 > A2 > A1 > A4) except for the rank of A4. This deviation may come from the different surface reactivity in the wet chemistry. When comparing the HRD among the three different preparation methods, the MC and the CC process produce almost identical results, which are always higher than those from the GS samples due to the difference in surface oxide [64,68,83,88,92,94]. 3.3. Ni/MH battery test All 12 batches of powder were made into negative electrode by roll-mill compaction on expanded Ni substrate without any binder. C-size cylindrical cells were made with paste Ni(OH)2 positive electrode, PP/PE grafted separator, and 30% KOH as electrolyte. The construction of the cell used a stapled top current collector and is not for high power application. These cells were formed by a 5-day 60 ◦ C heat formation followed by a 6-cycle electrical forma-

tion. The results from various tests are listed in Table 4. The room temperature high rate performance was studied by comparing the capacity from a 2 C rate discharge normalized to the capacity from a 0.2 C rate discharge. While the MC samples all show similar 2 C performance, the CC samples show mixed results and the GA samples are in general worse than either one of the other preparation methods, presumably due to the high oxygen content of the GA powder surface. The low temperature performance was evaluated by normalizing the capacity obtained from a −10 ◦ C and 0.5 C to the capacity obtained from a 30 ◦ C and 0.2 C rate. When comparing Ni/MH batteries from the same composition, the CC preparation again shows mixed results, as some increase (A2), some decrease (A4), while the others remain about the same as compared to the MC samples. Most of the GA samples are inferior to the other two methods in the low temperature performance. Peak power, or specific power, as a function of state of charge (SOC) was measured using a pulse discharge method [100]. The room temperature peak powers at the 20th cycle vs. SOC % for 4 Ni/MH batteries with negative electrodes made from the MC process with the A1, A2, A3 and A4 compositions are plotted in Fig. 8. The peak power decreases when more energy from the battery is depleted. The order of peak power at room temperature is A2 > A3 > A1 > A4 is similar to the order of PCT plateau

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Fig. 7. SEM backscattering electron image cross-section micrographs for alloy A4 prepared by conventional melt-and-cast (a), centrifugal casting (b), and gas atomization (c). The bar on the right lower corner indicates a length of 20 ␮m. Numbers in each micrograph indicate that EDS analyses were done in these areas and chemical compositions are summarized in Table 3.

pressure (A3 > A2 > A1 > A4), but differs from the half-cell HRD (A3 > A4 > A2 > A1). Although both the peak power and half-cell measurements are performed in an electrochemical environment, the former is in a semi-starve condition with a much higher dis-

Fig. 8. Peak power measured at 20th cycle for Ni/MH batteries made from conventional melt-and-cast A1, A2, A3 and A4 alloys.

charge rate while the later is in a completely flooded situation with a lower discharge rate. It is difficult to directly predict the peak power from the half-cell result. Both peak powers measured at room temperature and 0 C at the 5th cycle are listed in Table 4. The peak power also increases with increasing cycle number in the early period of the service life of the Ni/MH battery [112]. So the peak powers (5th cycle) listed in Table 4 are smaller than the data shown in Fig. 8 (20th cycle). When comparing peak power both at room temperature and 0 ◦ C with the same composition and different preparation methods, values from the CC samples again can be higher, lower than, or similar to the MC samples, and the GA samples give an inferior performance comparing to the other two for the same reason as the high power and low temperature. Charge retention was quantified by dividing the remaining capacity measure at C/5 rate vs. the original capacity before the 7-day storage at room temperature and is listed in Table 4. When comparing the MC samples from different compositions, the A3 cell has the smallest remaining capacity due to the combination of low Zr and Cr content in the formula. There is no clear trend to the effect of cooling rate on charge retention. The chemical composition is a higher sensitivity factor in determining charge retention than the structure, crystalline structure, and size of the constituent phases. Room temperature cycle life was studied using a C/2 charge and C/2 discharge scheme. The end of cycle life was determined when

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Fig. 9. Dependence of cycle life on the thickness (䊉) and length (♦) of the negative electrode of Ni/MH batteries made from gas atomization powder with A1 composition.

60% of the original capacity was reached and is listed in Table 4. Comparing batteries made with electrodes using the MC process, the cycle life is in this order (A1 > A2 > A4 > A3), which follows the general understanding of the role of each element playing in the cycle life (Zr, Cr, Ni to improve while Mn to deteriorate). When comparing the three preparation methods, both the CC and the GA have improved cycle life performance due to their lower or no level of stress accumulated in the final powder because less or no pulverization is required to obtain the final particle size. Theoretically, the GA samples with a spherical shape and no pulverization required in the process should have a better cycle life than the CC sample. We further studied the cycle life of the GA powders and obtained results very different from previous studies on MC and CC samples. In the cases of the MC and the CC, there is no clear correlation between the thickness of the electrode and the cycle life. However, when we prepared a few sets of electrode outside our normal range of electrode thickness (0.35–0.36 mm) and plotted the cycle life as functions of both negative electrode thickness and length as shown in Fig. 9, we discovered a strong dependence of cycle life on electrode thickness. Because we use no binder in electrode preparation and do not apply sintering, when the final electrode thickness (350 ␮m) approaches the diameter of the largest particle (150 ␮m), the physical integrity of the electrode itself faces a large challenge. As the thickness of electrode increases, more smaller particles are available to hold the electrode together and tends to increase the cycle life. Therefore, we reach the conclusion that the GA samples will give the best cycle life performance if a thicker electrode is used. 4. Conclusions We have compared three different fabrication methods (conventional melt-and-cast, centrifugal grinding, and gas atomization) to produce AB2 alloy powder used as negative electrode in the Ni/MH battery. As the cooling speed increases, the amount of C15 phase abundance increases. No obvious trends can be found in the unit cell volume or grain size. The battery performance can be summarized by the radar diagram illustrated in Fig. 10. All data are from the average of four different compositions A1, A2, A3 and A4 and normalized to the highest value in the same category. The formation data is from the reciprocal of number of cycle needed to achieve a stable capacity. Judging from this graph, we can summarize that the centrifugal casting method increases cycle life, improves low temperature performance, but hinders the formation

Fig. 10. Radar graph comparing Ni/MH battery performance with conventional melt-and-cast, centrifugal casting, and gas atomization. All data are from the average of four different alloy compositions.

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