- Email: [email protected]
Effect of ball milling and cold rolling on hydrogen storage properties of nanocrystalline TiV1.6Mn0.4 alloy
Journal of Alloys and Compounds 484 (2009) 154–158
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Effect of ball milling and cold rolling on hydrogen storage properties of nanocrystalline TiV1.6 Mn0.4 alloy S. Couillaud a , H. Enoki b , S. Amira c , J.L. Bobet a , E. Akiba b , J. Huot c,∗ a
ICMCB, Université Bordeaux 1, 87 avenue du Dr Schweitzer, 33608 Pessac Cedex, France Energy Technology Research Institute, National Institute of Advanced Industrial Science and Technology, AIST Tsukuba Central 5, 1-1-1 Higashi, Tsukuba, Ibaraki, 305-8565 Japan c Institut de Recherche sur l’Hydrogène, Université du Québec à Trois-Rivières, 3351, Boul. des Forges, Trois-Rivières, Québec, Canada G9A 5H7 b
a r t i c l e
i n f o
Article history: Received 2 March 2009 Accepted 10 May 2009 Available online 18 May 2009 Keywords: Cold rolling Ball milling BCC alloys Nanocrystalline Hydrogen absorption
a b s t r a c t The effect of high energy milling and cold rolling on the BCC solid solution TiV1.6 Mn0.4 was investigated. The as-cast alloy presented a nanocrystalline structure with a crystallite size of 17 nm. After 5 h of milling a slight reduction of crystallite size was measured and the BCC solid solution lattice parameter was reduced from 3.082 Å to 3.061 Å. Cold rolling also had for effect a reduction of crystallite size and lattice parameter. The main effect of cold rolling was to produce a highly textured alloy along the (2 0 0) Bragg peak. The as-cast alloy could store a maximum of 3.48 wt.% of hydrogen at 423 K and has a reversible capacity of 1.8 wt.%. The ball -milled and cold rolled samples did not absorb hydrogen even after 10 hydrogenation/dehydrogenation cycles. Possible reasons why ball milled and cold rolled samples did not absorb hydrogen are discussed. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Hydrogen storage is a key for the development of hydrogen economy. For some applications such as portable, stationary, and hybrid tank (high pressure and metal hydride) a storage capacity of a few wt.% of hydrogen is acceptable if the system cost is low and could operate near room temperature. In 1995 Iba and Akiba showed that BCC (body centered cubic) solid solution alloys of the type Ti–V–M (M: transition metal) could have sufficient hydrogen capacity and sorption kinetics for practical applications [1–3]. Since then, many studies have been published on BCC solid solution, particularly on Ti–V–Cr compositions [4–8]. In the case of Ti–V–Mn, it has been investigated by Akiba’s group which found from neutron diffraction that TiV1.1 Mn0.9 forms two hydrides: a mono-hydride with a pseudo-cubic FCC structure and a di-hydride with a CaF2 structure [9–11]. Usually, the mono-hydride is too stable for practical applications thus making the usable capacity smaller than the total capacity. If the mono-hydride could be destabilized without changing the hydrogen capacity and the di-hydride plateau pressure then a solid solution BCC could be very attractive alloys for commercial applications. In a recent study, it has been shown that when TiV0.9 Mn1.1 alloy is energetically milled for 80 h the C14 Laves phase originally present in the as-melted alloy vanishes and is replaced by
a FCC phase while the BCC phase experiences a slight reduction of lattice parameters. In terms of hydrogen storage properties, the milled alloy do not absorbs hydrogen, probably due to the presence of the FCC phase and the high iron contamination (about 1 wt.%). Despite this loss of hydrogen capacity of the milled alloy, it has been shown that TiV1.1 Mn0.9 has a catalytic effect on hydrogen sorption of magnesium [12–14]. The effect of severe plastic deformation (SPD) on BCC Ti–22Al–27Nb alloy has been investigated by Zhang et al [15,16]. They showed that the first hydrogenation (activation) was much faster for the deformed alloy compared to the as-quenched sample. The deformed alloy had also a faster absorption/desorption kinetics. However, the beneficial effect of deformation was lost after a few hydrogenation cycles. In these studies, SPD was obtained by cold rolling or compression. In the case of cold rolling one rolling was performed at 10.5% and 80% thickness reduction. Some of the 80% rolled specimen were further rolled to 10% thickness reduction in a perpendicular direction with respect to the first rolling. In the present work, we investigated the effect of multiple cold rolling on the hydrogen sorption properties of BCC solid solution. The composition TiV1.6 Mn0.4 was selected. A recent study indicated that a composition very close to this one (Ti0.95 V1.70 Mn0.36 ) has a total capacity of 3.80 wt.% and a reversible capacity at 373 K of 2.45 wt.% [17]. 2. Experimental
∗ Corresponding author. Tel.: +1 819 376 5011x3576; fax: +1 819 376 5164. E-mail address: [email protected] (J. Huot). 0925-8388/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2009.05.037
The TiV1.6 Mn0.4 alloy was synthesized by arc melting the constituent pure metals (Ti sponge 99.9%; V chunks 99.7%; Mn chunks 99,98%). The arc melted alloy was
S. Couillaud et al. / Journal of Alloys and Compounds 484 (2009) 154–158
155
of the as-cast and ball milled alloys, a Williamson–Hall [20] plot indicated that the strain is zero within one standard deviation and the crystallite size is identical to the one calculated from the Scherrer formula reported in Table 1. It is clear from Fig. 1 and Table 1 that the effect of high energy milling for 5 h is to decrease the crystallite size (peak broadening) and to reduce the lattice parameters (peaks shifted to higher angles). However, the relative intensities of the peaks do not change. The most drastic effect is seen for the cold rolled sample. The Bragg peak (2 0 0) is very intense compare to the other peaks, indicating a highly textured material. It should be noted that the lattice parameter and crystallite size of the cold rolled sample are identical to the ball milled sample. 3.2. Microstructure
Fig. 1. X-ray powder diffraction patterns of TiV1.6 Mn0.4 after casting, after ball milled for 5 h (BM 5 h), and after 150 cold rolls (CR150). Miller indices of each Bragg peak is indicated. turned over three times and remelted in order to ensure high homogeneity. The alloy button was inserted between two stainless steel (316) plates and rolled in air in a conventional rolling mill with 75 mm diameter rolls. In the first few rolls the thickness was regularly reduced by about 50% until a thickness of about 0.5 mm was achieved. Thereafter the sample was folded in two and rolled again thus giving a 50% thickness reduction on each roll. The prepared samples were stored in air before hydrogen activation and sorption measurements. Ball milling was performed with the same starting alloy in a Spex shaker model 8000. A vial of internal volume of 55 cm3 was used with 3 stainless steel balls. The powder to ball weight ratio was 10. In the case of milling, all handling and storing of the products were performed in an argon filled glove box. The hydrogen activation and sorption measurements were carried out on a Sievert-type apparatus. Before the first pressure–composition–temperature (PCT) each sample was evacuated by turbo molecular pump at 423 K for 12 h. All PCT measurements were performed at 423 K. The powder X-ray diffraction (XRD) patterns were made at room temperature using a Rigaku D-max diffractometer, with Cu-K␣ radiation. Using JADE software [18], the full width at half maximum (FWHM) of each peak was evaluated, taking into account the instrumental line broadening by subtracting the FWHM of a silicon reference line. Crystallite sizes and microstrain were evaluated by a Williamson–Hall plot. In this type of plot, the reciprocal peak widths (FWHM) cos are plotted versus 2sin , where is the diffraction angle. The slope (m) and intercept point (b) respectively give the microstrain (ε) and crystallite size () by the relations ε = m/4 and = 0.9/b where is the X-ray wavelength. For patterns that had only one significantly intense Bragg peak the crystallite sizes were evaluated from the Scherrer equation: = 0.9/(FWHM) cos [19]. The microstructures were analyzed by transmission electron microscopy (TEM) with selected area electron diffraction (SAED) and energy-dispersive X-ray spectroscopy (EDX) on a JEOL JEM 2000FXII. The hydrogen sorption properties were measured with a computer controlled Sievert’s-type apparatus.
3. Results and discussion 3.1. Crystal structure Fig. 1 shows the XRD patterns of TiV1.6 Mn0.4 in the as-cast state, after 5 h of milling (BM 5 h) and after 150 cold rolls (CR150). All patterns indicate a BCC solid solution structure. The lattice parameters and crystallite size of each sample are given in Table 1. As the CR150 pattern has only one measurable peak, the crystallite size was calculated by the Scherrer formula using the (2 0 0) peak. In the case
The microstructure of TiV1.6 Mn0.4 in as-cast, ball milled for 5 h, and after 150 cold rolls was investigated by TEM. Fig. 2 shows the bright field and dark field images of all samples. For as-cast TiV1.6 Mn0.4 , the dark field image shows crystallite size that is of the same order of magnitude as the value reported in Table 1. This validates the use of Scherrer equation for estimation of crystallite size. Moreover, a close inspection of the dark field image indicates a columnar structure where the columns are approximately the diameter of the crystallite size found by X-ray diffraction. This structure is most probably due to the directional solidification of the alloy after arc-melting. In the case of TiV1.6 Mn0.4 ball milled for 5 h, the dark field image clearly shows the crystallite size. Because of the thickness of the particle in its middle, the bright spots are easier to see on the particle’s edge. The biggest bright spot is about 30 nm diameter and the smallest visible ones less than 3 nm with most of the bright spots between 10 nm and 20 nm diameter, in good agreement with the crystallite size value estimated from X-ray powder diffraction reported in Table 1. For TiV1.6 Mn0.4 cold rolled 150 times, the bright field image clearly shows the pile-up of dislocations. The dark field image shows that, contrary to the as-cast and ball milled samples, the crystallites tend to be aligned along dislocations. Most of the crystallites are of the order of 5 nm to 10 nm, a value smaller than the one estimated from X-ray powder diffraction. This small discrepancy may be due to the significant texture of the alloy. The chemical composition of each alloy as measured by TEMEDX is reported in Table 2. The as-cast alloy is slightly titanium rich and manganese poor compared to the nominal composition but this is most probably due to small discrepancies between different batches of as-cast alloys. In the case of cold rolled sample the composition is very close to the nominal stoichiometry. The ball milled sample has a significant contamination by iron caused by attrition of the milling tools. The uncertainty on iron concentration is high because iron is not uniformly distributed in the alloy thus making a high dispersion on the iron concentration measurements. Iron contamination was also seen for other BCC solid solution [21,22]. As demonstrated by Santos et al [21], metallic iron is not seen in the X-ray diffraction pattern because iron atoms enter in the BCC solid solution. It is worthy to note that he relative atomic abundance of ball milled alloy when iron atoms are disregarded is: Ti 35 at.%, V 53 at.%, and Mn 11 at.%.
Table 1 Lattice parameter and crystallite size of TiV1.6 Mn0.4 after casting, after ball milled for 5 h, and after 150 cold rolls.
Table 2 Average alloy composition, in atomic %, as measured by TEM-EDX. One standard deviation on the last significant digit is indicated in parenthesis.
Alloy
Lattice parameter a (Å) (±0.001)
Crystallite size (nm) (±1)
Alloy
Ti
V
Mn
Fe
As-cast Ball milled 5 h Cold rolled 150 times
3.082 3.061 3.061
17 11 13
Nominal composition (calculated) As-cast Ball milled 5 hours Cold rolled 150 times
33.3 36.5(3) 32(2) 34(1)
53.3 54(1) 48(3) 53(1)
13.3 10(1) 10(2) 13(1)
– – 10(6) –
156
S. Couillaud et al. / Journal of Alloys and Compounds 484 (2009) 154–158
Fig. 2. TEM micrographs of TiV1.6 Mn0.4 after casting (top), after ball milled for 5 h (middle), and after 150 cold rolls (bottom). Micrographs on the left are bright field images and micrographs on the right are dark field images.
3.3. Hydrogen storage As for BCC the hydrogen sorption kinetics is usually quite fast, only the pressure–composition isotherms were measured. No activation process was performed before the first measurement, except for pumping with a turbo molecular pump at 423 K for 12 h. Fig. 3 shows the first cycle of the as-cast TiV1.6 Mn0.4 alloy. The maximum capacity reached is 1.74 H/M which correspond to 3.48 wt.% of hydrogen. The fact that the equilibrium pressure goes up to 0.09 MPa at H/M = 0.5 and then falls to 0.02 MPa at H/M = 0.7 may be due to the high hysteresis of the mono-hydride formation or to the slow kinetic due to the presence of a thin oxide layer. In the case of a slow hydrogenation kinetic, the reaction is so slow that in fact, the system is not in true equilibrium before the pressure is increased for the next data point. When the applied pressure gets much higher than the true equilibrium pressure then the driving force of the reaction is strong enough and the reaction could proceed. This may be the cause for the drop in pressure in the first part of the hydrogenation PCT curve.
As shown by Nakamura et al [9–11], the first plateau correspond to the two phase region of the solid-solution phase and the monohydride phase and goes up to H/M ≈ 0.7 which agrees with the present result. The second plateau corresponds to the formation of the di-hydride phase and terminates at H/M ≈ 1.7 which again agrees with previously reported values [9–11]. The same hydrogenation procedure was applied for TiV1.6 Mn0.4 ball milled 5 h and TiV1.6 Mn0.4 cold rolled 150 times. For both alloys no hydrogen absorption was registered. However, this may be due to a slow activation process. For the cold rolled sample it may be argued that the loss of capacity is due to the formation of a thick oxide layer. However, as ball milling was performed in argon and all handling were done in a glove box, formation of oxide layer should have been minimal and this sample should have maintain the capacity of as-cast alloy which was not the case. Moreover, TEMEDX did not show important amount of oxygen in both cold rolled and ball milled samples. Usually, multiple cycling could break the surface oxide and bring a higher capacity. To test this, all samples were cycles 9 or 10 times at 423 K between 10 MPa and 0.01 MPa of hydrogen pressure. In Fig. 4
S. Couillaud et al. / Journal of Alloys and Compounds 484 (2009) 154–158
Fig. 3. Pressure–composition–temperature (PCT) curve at 423K of as-cast TiV1.6 Mn0.4 . First cycle.
we present the PCT curves after cycling for as-cast TiV1.6 Mn0.4 and ball milled TiV1.6 Mn0.4 . For the as-cast sample, as the mono-hydride is too stable to be desorbed in these measuring conditions, only the di-hydride plateau is recorded. The absorption and desorption plateaus correspond with the ones recorded at the first cycle (Fig. 3) indicating no change in the di-hydride stability after 9 cycles. Form this curve, the reversible capacity is estimated to be about 0.9 H/M or 1.8 wt.%. The plateaus pressures and capacity are in good agreement with the corresponding values reported by Challet et al for the composition Ti0.95 V1.70 Mn0.36 [17]. In the case of ball milled TiV1.6 Mn0.4 no hydrogenation takes place even after 10 cycles as could be seen in Fig. 4. A similar situation was seen in the case of TiV0.9 Mn1.1 ball milled for 80 h [22]. A definitive explanation was not proposed but formation of a face centered cubic (FCC) alloy upon milling and iron contamination was envisioned. In the present case, ball milling was relatively short and the BCC structure was kept during milling. On the other hand, iron contamination is much higher in the present case than for TiV0.9 Mn1.1 ball milled for 80 h. The lattice parameter of the ball milled sample is much smaller than the as-cast alloy but a geometrical explanation could not be sustained because Ti–V–Mn BCC
157
alloy with lattice parameter down to 2.988 Å have been shown to absorb hydrogen [11]. The TiV1.6 Mn0.4 cold rolled 150 times was also submitted to 10 hydrogenation/dehydrogenation cycles. The PCT curve after 10 cycles (not shown) is identical to the one reported in Fig. 4 for the ball milled alloy. Therefore, for cold rolled as well as for ball milled samples cycling does not improve the hydrogen sorption properties and both alloys are totally inert to hydrogen. The fact that iron contamination is not present in the cold roll sample means that this is not the reason for loss of hydrogen capacity. The reason why cold rolled and ball milled samples do not absorb hydrogen while the as-cast alloy has a good hydrogen storage capacity is still not clear. As discussed before, the fact that the lattice parameter of cold rolled and ball milled alloys is reduced compare to the as-cast alloy may be a reason but from a purely geometrical point of view, the interstitial sites remain big enough to accommodate hydrogen atoms. However, it is probably not a coincidence that rolled and ball milled alloys have exactly the same lattice parameters. Both processes induce deformation in the alloy but they are not totally equivalent. For example, cold roll produce low angle grain boundaries while ball milling generate high angle grain boundaries. Also, as seen in the diffraction patterns the cold rolled sample is highly textured while the ball milling do not induce any texture. More research is needed in order to understand the mechanism responsible for this loss of hydrogen storage capacity. We are planning similar study for other TiV1−x Mnx stoichiometry to check if the same phenomenon is present in these alloys. 4. Conclusions From this study, we found that as-cast TiV1.6 Mn0.4 has a nanocrystalline structure and could absorb hydrogen up to 3.48 wt.% and with a reversible capacity of 1.8 wt.%. The effect of extended cold rolling as well as energetic ball milling is a reduction of crystalline size and lattice parameter but no change in the crystal structure. Both processes gave the same lattice parameter and crystallite size despite the fact that the ball milled alloy had a significant contamination by iron which entered in the BCC solid solution crystal structure. Therefore, it seems that the alloy has an intrinsic limit for lattice parameter and crystallite size. Unfortunately, neither ball milled sample nor cold rolled sample absorb hydrogen even after 10 cycles of hydrogen pressurization (10 MPa) and vacuum at 423 K. The reason for this significant loss of hydrogen capacity is still unknown but iron contamination and reduction of lattice parameters are probably not the cause. Further investigation is needed to solve this problem. Acknowledgements The authors gratefully acknowledge the financial support from the Natural Resources Canada and New Energy and Industrial Technology Development Organization (NEDO) under “Development for Safe Utilization and Infrastructure of Hydrogen”. J.H. would like to thanks the Japanese Society for Promotion of Science (JSPS) for a visiting scientist fellowship. References
Fig. 4. Pressure–composition–temperature (PCT) curves at 423K of as-cast TiV1.6 Mn0.4 after 9 cycles and of ball milled TiV1.6 Mn0.4 after 10 cycles.
[1] E. Akiba, H. Iba, Hydrogen absorption by Laves phase related BCC solid solution, Intermetallics 6 (1998) 461–470. [2] H. Iba, E. Akiba, The relation between microstructure and hydrogen absorbing property in Laves phase-solid solution multiphase alloys, Journal of Alloys and Compounds 231 (1995) 508–512. [3] H. Iba, E. Akiba, Hydrogen absorption and modulated structure in Ti–V–Mn alloys, Journal of Alloys and Compounds 253–254 (1997) 21–24.
158
S. Couillaud et al. / Journal of Alloys and Compounds 484 (2009) 154–158
[4] S.-W. Cho, et al., The hydrogen storage characteristics of Ti–Cr–V alloys, Journal of Alloys and Compounds 288 (1999) 294–298. [5] M. Okada, et al., Ti–V–Cr b.c.c. alloys with high protium content, Journal of Alloys and Compounds 330–332 (2002) 511–516. [6] X.B. Yu, et al., Effect of Cr content on hydrogen storage properties for Ti–Vbased BCC-phase alloys, International Journal of Hydrogen Energy 29 (2004) 1377–1381. [7] H. Pan, et al., Effect of Cr on the structural and electrochemical properties of TiV-based two phase hydrogen storage alloys, Journal of Alloys and Compounds 404–406 (2005) 669–674. [8] S.-W. Cho, et al., Hydrogen absorption–desorption properties of Ti0.32Cr0.43V0.25 alloy, Journal of Alloys and Compounds 430 (2007) 136–141. [9] Y. Nakamura, E. Akiba, New hydride phase with a deformed FCC structure in the Ti–V–Mn solid solution-hydrogen system, Journal of Alloys and Compounds 311 (2000) 317–321. [10] Y. Nakamura, et al., Crystal structure of two hydrides formed from a Ti–V–Mn BCC solid solution alloy studied by time-of-flight neutron powder diffraction - a NaCl structure and a CaF2 structure, Journal of Alloys and Compounds 316 (2001) 284–289. [11] Y. Nakamura, E. Akiba, Hydriding properties and crystal structure of NaCl-type mono-hydrides formed from Ti–V–Mn BCC solid solutions, Journal of Alloys and Compounds 345 (2002) 175–182. [12] X. Liu, et al., Thermal stabilization and hydrogen storage properties of Mg-40 wt.%Ti0.28Cr0.50V0.22 composite prepared by mechanical milling, International Journal of Hydrogen Energy 32 (2007) 965–968.
[13] H. Gu, Y. Zhu, L. Li, Characterization of hydrogen storage properties of Mg30 wt.% Ti1.0 V1.1 Mn0.9 composite, Journal of Alloys and Compounds 424 (2006) 382–387. [14] A.L. Yonkeu, et al., Kinetic investigation of the catalytic effect of a body centered cubic-alloy TiV1.1Mn0.9 (BCC) on hydriding/dehydriding properties of magnesium, Journal of Alloys and Compounds 460 (2008) 559–564. [15] L.T. Zhang, et al., Hydrogen absorption and desorption in a B2 single-phase Ti-22Al-27Nb alloy before and after deformation, Acta Materialia 49 (2001) 751–758. [16] L.T. Zhang, et al., Effects of cold-rolling on the hydrogen absorption/desorption behavior of Ti-22Al-27Nb alloys, Materials Science and Engineering A 329–331 (2002) 362–366. [17] S. Challet, M. Latroche, F. Heurtaux, Hydrogenation properties and crystal structure of the single BCC (Ti0.355 V0.645 )100−x Mx alloys with M = Mn, Fe, Co, Ni (x = 7, 14 and 21), Journal of Alloys and Compounds 439 (2007) 294–301. [18] JADE, MDI Materials Data, 2008. [19] H.P. Klug, L.E. Alexander, X-ray Difffraction Procedures, John Wiley & Sons, New York, 1974. [20] G.K. Williamson, W.H. Hall, X-ray line broadening from filed aluminium and wolfram, Acta Metallurgica 1 (1953) 22–31. [21] S.F. Santos, et al., Mechanical and reactive miling of a TiCrV BCC solid solution, Journal of Metastable and Nanocrystalline Materials 20–21 (2004) 291–296. [22] J. Huot, H. Enoki, E. Akiba, Synthesis, phase transformation, and hydrogen storage properties of ball-milled TiV0.9Mn1.1, Journal of Alloys and Compounds 453 (2008) 203–209.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Effect of ball milling and cold rolling on hydrogen storage properties of nanocrystalline TiV1.6 Mn0.4 alloy S. Couillaud a , H. Enoki b , S. Amira c , J.L. Bobet a , E. Akiba b , J. Huot c,∗ a
ICMCB, Université Bordeaux 1, 87 avenue du Dr Schweitzer, 33608 Pessac Cedex, France Energy Technology Research Institute, National Institute of Advanced Industrial Science and Technology, AIST Tsukuba Central 5, 1-1-1 Higashi, Tsukuba, Ibaraki, 305-8565 Japan c Institut de Recherche sur l’Hydrogène, Université du Québec à Trois-Rivières, 3351, Boul. des Forges, Trois-Rivières, Québec, Canada G9A 5H7 b
a r t i c l e
i n f o
Article history: Received 2 March 2009 Accepted 10 May 2009 Available online 18 May 2009 Keywords: Cold rolling Ball milling BCC alloys Nanocrystalline Hydrogen absorption
a b s t r a c t The effect of high energy milling and cold rolling on the BCC solid solution TiV1.6 Mn0.4 was investigated. The as-cast alloy presented a nanocrystalline structure with a crystallite size of 17 nm. After 5 h of milling a slight reduction of crystallite size was measured and the BCC solid solution lattice parameter was reduced from 3.082 Å to 3.061 Å. Cold rolling also had for effect a reduction of crystallite size and lattice parameter. The main effect of cold rolling was to produce a highly textured alloy along the (2 0 0) Bragg peak. The as-cast alloy could store a maximum of 3.48 wt.% of hydrogen at 423 K and has a reversible capacity of 1.8 wt.%. The ball -milled and cold rolled samples did not absorb hydrogen even after 10 hydrogenation/dehydrogenation cycles. Possible reasons why ball milled and cold rolled samples did not absorb hydrogen are discussed. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Hydrogen storage is a key for the development of hydrogen economy. For some applications such as portable, stationary, and hybrid tank (high pressure and metal hydride) a storage capacity of a few wt.% of hydrogen is acceptable if the system cost is low and could operate near room temperature. In 1995 Iba and Akiba showed that BCC (body centered cubic) solid solution alloys of the type Ti–V–M (M: transition metal) could have sufficient hydrogen capacity and sorption kinetics for practical applications [1–3]. Since then, many studies have been published on BCC solid solution, particularly on Ti–V–Cr compositions [4–8]. In the case of Ti–V–Mn, it has been investigated by Akiba’s group which found from neutron diffraction that TiV1.1 Mn0.9 forms two hydrides: a mono-hydride with a pseudo-cubic FCC structure and a di-hydride with a CaF2 structure [9–11]. Usually, the mono-hydride is too stable for practical applications thus making the usable capacity smaller than the total capacity. If the mono-hydride could be destabilized without changing the hydrogen capacity and the di-hydride plateau pressure then a solid solution BCC could be very attractive alloys for commercial applications. In a recent study, it has been shown that when TiV0.9 Mn1.1 alloy is energetically milled for 80 h the C14 Laves phase originally present in the as-melted alloy vanishes and is replaced by
a FCC phase while the BCC phase experiences a slight reduction of lattice parameters. In terms of hydrogen storage properties, the milled alloy do not absorbs hydrogen, probably due to the presence of the FCC phase and the high iron contamination (about 1 wt.%). Despite this loss of hydrogen capacity of the milled alloy, it has been shown that TiV1.1 Mn0.9 has a catalytic effect on hydrogen sorption of magnesium [12–14]. The effect of severe plastic deformation (SPD) on BCC Ti–22Al–27Nb alloy has been investigated by Zhang et al [15,16]. They showed that the first hydrogenation (activation) was much faster for the deformed alloy compared to the as-quenched sample. The deformed alloy had also a faster absorption/desorption kinetics. However, the beneficial effect of deformation was lost after a few hydrogenation cycles. In these studies, SPD was obtained by cold rolling or compression. In the case of cold rolling one rolling was performed at 10.5% and 80% thickness reduction. Some of the 80% rolled specimen were further rolled to 10% thickness reduction in a perpendicular direction with respect to the first rolling. In the present work, we investigated the effect of multiple cold rolling on the hydrogen sorption properties of BCC solid solution. The composition TiV1.6 Mn0.4 was selected. A recent study indicated that a composition very close to this one (Ti0.95 V1.70 Mn0.36 ) has a total capacity of 3.80 wt.% and a reversible capacity at 373 K of 2.45 wt.% [17]. 2. Experimental
∗ Corresponding author. Tel.: +1 819 376 5011x3576; fax: +1 819 376 5164. E-mail address: [email protected] (J. Huot). 0925-8388/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2009.05.037
The TiV1.6 Mn0.4 alloy was synthesized by arc melting the constituent pure metals (Ti sponge 99.9%; V chunks 99.7%; Mn chunks 99,98%). The arc melted alloy was
S. Couillaud et al. / Journal of Alloys and Compounds 484 (2009) 154–158
155
of the as-cast and ball milled alloys, a Williamson–Hall [20] plot indicated that the strain is zero within one standard deviation and the crystallite size is identical to the one calculated from the Scherrer formula reported in Table 1. It is clear from Fig. 1 and Table 1 that the effect of high energy milling for 5 h is to decrease the crystallite size (peak broadening) and to reduce the lattice parameters (peaks shifted to higher angles). However, the relative intensities of the peaks do not change. The most drastic effect is seen for the cold rolled sample. The Bragg peak (2 0 0) is very intense compare to the other peaks, indicating a highly textured material. It should be noted that the lattice parameter and crystallite size of the cold rolled sample are identical to the ball milled sample. 3.2. Microstructure
Fig. 1. X-ray powder diffraction patterns of TiV1.6 Mn0.4 after casting, after ball milled for 5 h (BM 5 h), and after 150 cold rolls (CR150). Miller indices of each Bragg peak is indicated. turned over three times and remelted in order to ensure high homogeneity. The alloy button was inserted between two stainless steel (316) plates and rolled in air in a conventional rolling mill with 75 mm diameter rolls. In the first few rolls the thickness was regularly reduced by about 50% until a thickness of about 0.5 mm was achieved. Thereafter the sample was folded in two and rolled again thus giving a 50% thickness reduction on each roll. The prepared samples were stored in air before hydrogen activation and sorption measurements. Ball milling was performed with the same starting alloy in a Spex shaker model 8000. A vial of internal volume of 55 cm3 was used with 3 stainless steel balls. The powder to ball weight ratio was 10. In the case of milling, all handling and storing of the products were performed in an argon filled glove box. The hydrogen activation and sorption measurements were carried out on a Sievert-type apparatus. Before the first pressure–composition–temperature (PCT) each sample was evacuated by turbo molecular pump at 423 K for 12 h. All PCT measurements were performed at 423 K. The powder X-ray diffraction (XRD) patterns were made at room temperature using a Rigaku D-max diffractometer, with Cu-K␣ radiation. Using JADE software [18], the full width at half maximum (FWHM) of each peak was evaluated, taking into account the instrumental line broadening by subtracting the FWHM of a silicon reference line. Crystallite sizes and microstrain were evaluated by a Williamson–Hall plot. In this type of plot, the reciprocal peak widths (FWHM) cos are plotted versus 2sin , where is the diffraction angle. The slope (m) and intercept point (b) respectively give the microstrain (ε) and crystallite size () by the relations ε = m/4 and = 0.9/b where is the X-ray wavelength. For patterns that had only one significantly intense Bragg peak the crystallite sizes were evaluated from the Scherrer equation: = 0.9/(FWHM) cos [19]. The microstructures were analyzed by transmission electron microscopy (TEM) with selected area electron diffraction (SAED) and energy-dispersive X-ray spectroscopy (EDX) on a JEOL JEM 2000FXII. The hydrogen sorption properties were measured with a computer controlled Sievert’s-type apparatus.
3. Results and discussion 3.1. Crystal structure Fig. 1 shows the XRD patterns of TiV1.6 Mn0.4 in the as-cast state, after 5 h of milling (BM 5 h) and after 150 cold rolls (CR150). All patterns indicate a BCC solid solution structure. The lattice parameters and crystallite size of each sample are given in Table 1. As the CR150 pattern has only one measurable peak, the crystallite size was calculated by the Scherrer formula using the (2 0 0) peak. In the case
The microstructure of TiV1.6 Mn0.4 in as-cast, ball milled for 5 h, and after 150 cold rolls was investigated by TEM. Fig. 2 shows the bright field and dark field images of all samples. For as-cast TiV1.6 Mn0.4 , the dark field image shows crystallite size that is of the same order of magnitude as the value reported in Table 1. This validates the use of Scherrer equation for estimation of crystallite size. Moreover, a close inspection of the dark field image indicates a columnar structure where the columns are approximately the diameter of the crystallite size found by X-ray diffraction. This structure is most probably due to the directional solidification of the alloy after arc-melting. In the case of TiV1.6 Mn0.4 ball milled for 5 h, the dark field image clearly shows the crystallite size. Because of the thickness of the particle in its middle, the bright spots are easier to see on the particle’s edge. The biggest bright spot is about 30 nm diameter and the smallest visible ones less than 3 nm with most of the bright spots between 10 nm and 20 nm diameter, in good agreement with the crystallite size value estimated from X-ray powder diffraction reported in Table 1. For TiV1.6 Mn0.4 cold rolled 150 times, the bright field image clearly shows the pile-up of dislocations. The dark field image shows that, contrary to the as-cast and ball milled samples, the crystallites tend to be aligned along dislocations. Most of the crystallites are of the order of 5 nm to 10 nm, a value smaller than the one estimated from X-ray powder diffraction. This small discrepancy may be due to the significant texture of the alloy. The chemical composition of each alloy as measured by TEMEDX is reported in Table 2. The as-cast alloy is slightly titanium rich and manganese poor compared to the nominal composition but this is most probably due to small discrepancies between different batches of as-cast alloys. In the case of cold rolled sample the composition is very close to the nominal stoichiometry. The ball milled sample has a significant contamination by iron caused by attrition of the milling tools. The uncertainty on iron concentration is high because iron is not uniformly distributed in the alloy thus making a high dispersion on the iron concentration measurements. Iron contamination was also seen for other BCC solid solution [21,22]. As demonstrated by Santos et al [21], metallic iron is not seen in the X-ray diffraction pattern because iron atoms enter in the BCC solid solution. It is worthy to note that he relative atomic abundance of ball milled alloy when iron atoms are disregarded is: Ti 35 at.%, V 53 at.%, and Mn 11 at.%.
Table 1 Lattice parameter and crystallite size of TiV1.6 Mn0.4 after casting, after ball milled for 5 h, and after 150 cold rolls.
Table 2 Average alloy composition, in atomic %, as measured by TEM-EDX. One standard deviation on the last significant digit is indicated in parenthesis.
Alloy
Lattice parameter a (Å) (±0.001)
Crystallite size (nm) (±1)
Alloy
Ti
V
Mn
Fe
As-cast Ball milled 5 h Cold rolled 150 times
3.082 3.061 3.061
17 11 13
Nominal composition (calculated) As-cast Ball milled 5 hours Cold rolled 150 times
33.3 36.5(3) 32(2) 34(1)
53.3 54(1) 48(3) 53(1)
13.3 10(1) 10(2) 13(1)
– – 10(6) –
156
S. Couillaud et al. / Journal of Alloys and Compounds 484 (2009) 154–158
Fig. 2. TEM micrographs of TiV1.6 Mn0.4 after casting (top), after ball milled for 5 h (middle), and after 150 cold rolls (bottom). Micrographs on the left are bright field images and micrographs on the right are dark field images.
3.3. Hydrogen storage As for BCC the hydrogen sorption kinetics is usually quite fast, only the pressure–composition isotherms were measured. No activation process was performed before the first measurement, except for pumping with a turbo molecular pump at 423 K for 12 h. Fig. 3 shows the first cycle of the as-cast TiV1.6 Mn0.4 alloy. The maximum capacity reached is 1.74 H/M which correspond to 3.48 wt.% of hydrogen. The fact that the equilibrium pressure goes up to 0.09 MPa at H/M = 0.5 and then falls to 0.02 MPa at H/M = 0.7 may be due to the high hysteresis of the mono-hydride formation or to the slow kinetic due to the presence of a thin oxide layer. In the case of a slow hydrogenation kinetic, the reaction is so slow that in fact, the system is not in true equilibrium before the pressure is increased for the next data point. When the applied pressure gets much higher than the true equilibrium pressure then the driving force of the reaction is strong enough and the reaction could proceed. This may be the cause for the drop in pressure in the first part of the hydrogenation PCT curve.
As shown by Nakamura et al [9–11], the first plateau correspond to the two phase region of the solid-solution phase and the monohydride phase and goes up to H/M ≈ 0.7 which agrees with the present result. The second plateau corresponds to the formation of the di-hydride phase and terminates at H/M ≈ 1.7 which again agrees with previously reported values [9–11]. The same hydrogenation procedure was applied for TiV1.6 Mn0.4 ball milled 5 h and TiV1.6 Mn0.4 cold rolled 150 times. For both alloys no hydrogen absorption was registered. However, this may be due to a slow activation process. For the cold rolled sample it may be argued that the loss of capacity is due to the formation of a thick oxide layer. However, as ball milling was performed in argon and all handling were done in a glove box, formation of oxide layer should have been minimal and this sample should have maintain the capacity of as-cast alloy which was not the case. Moreover, TEMEDX did not show important amount of oxygen in both cold rolled and ball milled samples. Usually, multiple cycling could break the surface oxide and bring a higher capacity. To test this, all samples were cycles 9 or 10 times at 423 K between 10 MPa and 0.01 MPa of hydrogen pressure. In Fig. 4
S. Couillaud et al. / Journal of Alloys and Compounds 484 (2009) 154–158
Fig. 3. Pressure–composition–temperature (PCT) curve at 423K of as-cast TiV1.6 Mn0.4 . First cycle.
we present the PCT curves after cycling for as-cast TiV1.6 Mn0.4 and ball milled TiV1.6 Mn0.4 . For the as-cast sample, as the mono-hydride is too stable to be desorbed in these measuring conditions, only the di-hydride plateau is recorded. The absorption and desorption plateaus correspond with the ones recorded at the first cycle (Fig. 3) indicating no change in the di-hydride stability after 9 cycles. Form this curve, the reversible capacity is estimated to be about 0.9 H/M or 1.8 wt.%. The plateaus pressures and capacity are in good agreement with the corresponding values reported by Challet et al for the composition Ti0.95 V1.70 Mn0.36 [17]. In the case of ball milled TiV1.6 Mn0.4 no hydrogenation takes place even after 10 cycles as could be seen in Fig. 4. A similar situation was seen in the case of TiV0.9 Mn1.1 ball milled for 80 h [22]. A definitive explanation was not proposed but formation of a face centered cubic (FCC) alloy upon milling and iron contamination was envisioned. In the present case, ball milling was relatively short and the BCC structure was kept during milling. On the other hand, iron contamination is much higher in the present case than for TiV0.9 Mn1.1 ball milled for 80 h. The lattice parameter of the ball milled sample is much smaller than the as-cast alloy but a geometrical explanation could not be sustained because Ti–V–Mn BCC
157
alloy with lattice parameter down to 2.988 Å have been shown to absorb hydrogen [11]. The TiV1.6 Mn0.4 cold rolled 150 times was also submitted to 10 hydrogenation/dehydrogenation cycles. The PCT curve after 10 cycles (not shown) is identical to the one reported in Fig. 4 for the ball milled alloy. Therefore, for cold rolled as well as for ball milled samples cycling does not improve the hydrogen sorption properties and both alloys are totally inert to hydrogen. The fact that iron contamination is not present in the cold roll sample means that this is not the reason for loss of hydrogen capacity. The reason why cold rolled and ball milled samples do not absorb hydrogen while the as-cast alloy has a good hydrogen storage capacity is still not clear. As discussed before, the fact that the lattice parameter of cold rolled and ball milled alloys is reduced compare to the as-cast alloy may be a reason but from a purely geometrical point of view, the interstitial sites remain big enough to accommodate hydrogen atoms. However, it is probably not a coincidence that rolled and ball milled alloys have exactly the same lattice parameters. Both processes induce deformation in the alloy but they are not totally equivalent. For example, cold roll produce low angle grain boundaries while ball milling generate high angle grain boundaries. Also, as seen in the diffraction patterns the cold rolled sample is highly textured while the ball milling do not induce any texture. More research is needed in order to understand the mechanism responsible for this loss of hydrogen storage capacity. We are planning similar study for other TiV1−x Mnx stoichiometry to check if the same phenomenon is present in these alloys. 4. Conclusions From this study, we found that as-cast TiV1.6 Mn0.4 has a nanocrystalline structure and could absorb hydrogen up to 3.48 wt.% and with a reversible capacity of 1.8 wt.%. The effect of extended cold rolling as well as energetic ball milling is a reduction of crystalline size and lattice parameter but no change in the crystal structure. Both processes gave the same lattice parameter and crystallite size despite the fact that the ball milled alloy had a significant contamination by iron which entered in the BCC solid solution crystal structure. Therefore, it seems that the alloy has an intrinsic limit for lattice parameter and crystallite size. Unfortunately, neither ball milled sample nor cold rolled sample absorb hydrogen even after 10 cycles of hydrogen pressurization (10 MPa) and vacuum at 423 K. The reason for this significant loss of hydrogen capacity is still unknown but iron contamination and reduction of lattice parameters are probably not the cause. Further investigation is needed to solve this problem. Acknowledgements The authors gratefully acknowledge the financial support from the Natural Resources Canada and New Energy and Industrial Technology Development Organization (NEDO) under “Development for Safe Utilization and Infrastructure of Hydrogen”. J.H. would like to thanks the Japanese Society for Promotion of Science (JSPS) for a visiting scientist fellowship. References
Fig. 4. Pressure–composition–temperature (PCT) curves at 423K of as-cast TiV1.6 Mn0.4 after 9 cycles and of ball milled TiV1.6 Mn0.4 after 10 cycles.
[1] E. Akiba, H. Iba, Hydrogen absorption by Laves phase related BCC solid solution, Intermetallics 6 (1998) 461–470. [2] H. Iba, E. Akiba, The relation between microstructure and hydrogen absorbing property in Laves phase-solid solution multiphase alloys, Journal of Alloys and Compounds 231 (1995) 508–512. [3] H. Iba, E. Akiba, Hydrogen absorption and modulated structure in Ti–V–Mn alloys, Journal of Alloys and Compounds 253–254 (1997) 21–24.
158
S. Couillaud et al. / Journal of Alloys and Compounds 484 (2009) 154–158
[4] S.-W. Cho, et al., The hydrogen storage characteristics of Ti–Cr–V alloys, Journal of Alloys and Compounds 288 (1999) 294–298. [5] M. Okada, et al., Ti–V–Cr b.c.c. alloys with high protium content, Journal of Alloys and Compounds 330–332 (2002) 511–516. [6] X.B. Yu, et al., Effect of Cr content on hydrogen storage properties for Ti–Vbased BCC-phase alloys, International Journal of Hydrogen Energy 29 (2004) 1377–1381. [7] H. Pan, et al., Effect of Cr on the structural and electrochemical properties of TiV-based two phase hydrogen storage alloys, Journal of Alloys and Compounds 404–406 (2005) 669–674. [8] S.-W. Cho, et al., Hydrogen absorption–desorption properties of Ti0.32Cr0.43V0.25 alloy, Journal of Alloys and Compounds 430 (2007) 136–141. [9] Y. Nakamura, E. Akiba, New hydride phase with a deformed FCC structure in the Ti–V–Mn solid solution-hydrogen system, Journal of Alloys and Compounds 311 (2000) 317–321. [10] Y. Nakamura, et al., Crystal structure of two hydrides formed from a Ti–V–Mn BCC solid solution alloy studied by time-of-flight neutron powder diffraction - a NaCl structure and a CaF2 structure, Journal of Alloys and Compounds 316 (2001) 284–289. [11] Y. Nakamura, E. Akiba, Hydriding properties and crystal structure of NaCl-type mono-hydrides formed from Ti–V–Mn BCC solid solutions, Journal of Alloys and Compounds 345 (2002) 175–182. [12] X. Liu, et al., Thermal stabilization and hydrogen storage properties of Mg-40 wt.%Ti0.28Cr0.50V0.22 composite prepared by mechanical milling, International Journal of Hydrogen Energy 32 (2007) 965–968.
[13] H. Gu, Y. Zhu, L. Li, Characterization of hydrogen storage properties of Mg30 wt.% Ti1.0 V1.1 Mn0.9 composite, Journal of Alloys and Compounds 424 (2006) 382–387. [14] A.L. Yonkeu, et al., Kinetic investigation of the catalytic effect of a body centered cubic-alloy TiV1.1Mn0.9 (BCC) on hydriding/dehydriding properties of magnesium, Journal of Alloys and Compounds 460 (2008) 559–564. [15] L.T. Zhang, et al., Hydrogen absorption and desorption in a B2 single-phase Ti-22Al-27Nb alloy before and after deformation, Acta Materialia 49 (2001) 751–758. [16] L.T. Zhang, et al., Effects of cold-rolling on the hydrogen absorption/desorption behavior of Ti-22Al-27Nb alloys, Materials Science and Engineering A 329–331 (2002) 362–366. [17] S. Challet, M. Latroche, F. Heurtaux, Hydrogenation properties and crystal structure of the single BCC (Ti0.355 V0.645 )100−x Mx alloys with M = Mn, Fe, Co, Ni (x = 7, 14 and 21), Journal of Alloys and Compounds 439 (2007) 294–301. [18] JADE, MDI Materials Data, 2008. [19] H.P. Klug, L.E. Alexander, X-ray Difffraction Procedures, John Wiley & Sons, New York, 1974. [20] G.K. Williamson, W.H. Hall, X-ray line broadening from filed aluminium and wolfram, Acta Metallurgica 1 (1953) 22–31. [21] S.F. Santos, et al., Mechanical and reactive miling of a TiCrV BCC solid solution, Journal of Metastable and Nanocrystalline Materials 20–21 (2004) 291–296. [22] J. Huot, H. Enoki, E. Akiba, Synthesis, phase transformation, and hydrogen storage properties of ball-milled TiV0.9Mn1.1, Journal of Alloys and Compounds 453 (2008) 203–209.