EFFECT OF MILLING TIME ON THE FIRST ...

CONAMET/SAM-2006

CONAMET/SAM 2006

EFFECT OF MILLING TIME ON THE FIRST HYDROGEN ABSORPTIONDESORPTION CURVES AND STRUCTURAL PROPERTIES OF A MmNi5-Ni ALLOY OBTAINED BY MECHANICAL MILLING M.R. ESQUIVEL 1,2 AND G. MEYER 1,2 1

Comisión Nacional de Energía Atómica-Centro Atómico Bariloche-Avda Bustillo km. 9,5-BarilocheRío Negro-Argentina 2 Consejo Nacional de Investigaciones Científicas y Técnicas [email protected]

ABSTRACT Mechanical milling has become one of the most successful methods for synthesis of alloys. Easy scaling up and low processing cost are mentioned as main advantages over both full equilibrium and chemical synthesis methods. This technique is specially appropriated to obtain alloys used to interact with hydrogen because it produces defects and strain in the alloy microstructure that enhances the reaction. Nevertheless, the effects of mechanical milling on both the microstructure and hydrogen interaction properties on most hydride forming materials are not completely understood. In this work, the effect of mechanical milling on both the structural properties and first hydrogen absorption – desorption curves of a MmNi5-Ni mixture is analyzed. Mixtures of both Mm (Mischmetal, a lanthanides alloy)-Ni and a MmNi5-Ni were treated in a Uni-Ball-II apparatus. The chamber was opened at fixed milling times and samples were withdrawn in a glove box under Ar atmosphere. O2 level was maintained below a 5 ppm level to avoid sample oxidation. Hydrogen-alloy interaction was characterized using a Sievert´s type equipment designed in our laboratory. Absorption – desorption curves were obtained at room temperature, 50 ºC and 90 ºC for samples milled during different periods of time. From these curves, the effect of milling is studied. Scanning electron microscopy (SEM) was used to study the particle size and morphology changes during the evolution of mechanical milling. Phase identity, crystallite size and strain of both MmNi5 and Ni were studied using X-ray diffraction (XRD). These changes on structural parameters were correlated to the effect of milling time and governing mechanisms during mechanical milling. Results from hydriding – dehydriding curves were also correlated to changes on microstructure during sample preparation. Key words: MmNi5 , Mechanical Alloying, Hydrogen

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1. INTRODUCTION The last years have been characterized by an extensive study of AB5 based alloys with potential application in devices for the storage or compression/purification of hydrogen [1,2]. It enhanced the research on both the structural characterization and hydrogen interaction of the alloy family [3,4]. Both chemical [5] and equilibrium methods [6,7] were used to synthesize the alloys studied. A combination of these processes followed by mechanical milling (MM) was also used to obtain the alloy [7]. Analysis of the interaction with hydrogen after synthesis was also presented [8,9]. Nevertheless, there are aspects not fully analyzed in these research works. One of them is the effect of milling time on the structure. It is not only related to the changes of the cell parameters of the alloy but also to the possibility of studying the modifications on the structure due to the effects of milling time. Since milling can produced isotropic or non isotropic effects on the structure [10], it can lead to a different effects on the interaction of the alloy with hydrogen. The other aspect not completely studied is the analysis of the stages present during mechanical milling [11]. It is important because the final size of particles depend on the stage controlling the milling process. Particle size and morphology would depend on either fracture or cold welding dominating the process [11]. This fact would also affect the interaction of the alloy with hydrogen. Then, the effect of milling time on the structure and the interaction with hydrogen should be analyzed simultaneously to conduct to a better understanding of the processing of materials by milling. The potential applications of this study on both research and industrial practice aimed the elaboration of the present work. 2. EXPERIMENTAL Pure Ni (3.80 µm, 99.99%) (Sigma Aldrich) and drilled lumps of Mischmetal (99.7 %) (Alpha Aesar) of nominal composition 52.0 wt% Ce, 25.6 wt% La, 16.9 wt% Pr , 5.5 wt% Nd were mechanically milled under Argon atmosphere. Neutron Activation Analysis (NAA) and Energy Dispersive Spectroscopy (EDS) were used to verify Mischmetal (Mm) chemical composition. Milling was performed in a Uni-Ball-Mill II apparatus (Australian Scientific Instruments). Powder mixture in a proportion of 20% excess of Nickel over the stoichiometric MmNi5

composition and steel balls were set in a stainless steel chamber under Ar atmosphere in a glove box. Particles size and morphology were observed by Scanning Electron Microscopy (SEM). The balls to powder mass relation was 33.5/1. The oxygen level was monitored by a trace anaylizer (Series 3000, Alpha Omega) and kept under 5 ppm to avoid material oxidation during sample manipulation inside the glove box. Representative amount of powder was withdrawn from chamber at different milling times inside a glove box and samples were analyzed by X-ray powder diffraction (XRD). Room temperature X-ray diffraction was achieved on a Philips PW 1710/01 Instrument with Cu Kα radiation (graphite monocromator). Diffraction patterns were analyzed by the Rietveld method [12] using DBWS software [13]. Strain and crystallite size effects were estimated from diffraction peaks by assuming empirically a Gauss distribution and Cauchy (Lorentz) component, respectively [14]. A Sievert´s type equipment was used to measure hydrogen absorption-desorption curves. The sample was placed in the reactor at fixed temperature and a selected initial pressure. Adsorption pressure was fixed at 6000 kPa and desorption pressure at 100 kPa. A PC-based data acquisition system monitored and controlled the experiment variables. Experimental set-up device details can be found elsewhere [15]. 3. RESULTS AND DISCUSSION 3.1 MmNi5-Ni synthesis by combined mechanical milling followed by heating. Mm and Ni were milled during 465 h to synthesize the MmNi5-Ni mixture. Full crystalline sample was obtained after heating at 600 ºC during 5 days. Diffractograms obtained during the synthesis process are shown in Figure 1. As milled and after heating MmNi5-Ni mixtures are shown in Fig. 1.a and 1.b, respectively. For reference, experimental diffraction pattern of pure Ni is also shown in Fig. 1.c

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Figure 1 a) As milled MmNi5-Ni. b) MmNi5-Ni after heating. c) Experimental diffraction pattern of Ni. MmNi5 structure found is in agreement with previous results [16]. Experimental diffraction lines are assigned to eihter MmNi5 or Ni and no secondary phases were found. Sample heated mass percentages were calculated using DBWS program [12] A summary of results including the structure parameters and mass percentages found in the MmNi5-Ni mixture of Fig. 1.b is shown in Table I. Rwp stands for the goodness of the fitting using Rietveld method. Table 1 a (Å) 4.907

MmNi5 c (Å) 3.977

V (Å3) 82.93

Ni a (Å) 3.530

Mass %

Rwp

MmNi5

Ni

94±1

6±1

10

Table 1. Structure parameters and mass percentages of after heating MmNi5-Ni mixture MmNi5-Ni mixture was heated in order to both increase the crystallize sample and release strain. It is clearly observed by comparing Figs 1.a and 1.b. The first one shows wider diffraction peaks due to strain and low crystallite size typical of milling products. After heating, strain is released and crystallite size augmented and thinner peaks are obtained as a result. This sample is used as reference to compare the changes on structure due to milling which are, in turn, correlated to the first hydriding/dehydriding characteristics of samples milled at 60 and 200 h 3.2 Mechanical milling at shorter periods of time. Figure 2 shows the diffractograms of MmNi5-Ni mixture withdrawn at different integrated short milling times.

Figure 2. Samples obtained at different integrated milling times. a) 2h. b) 4h. c) 8 h. d) 10 h. As observed, the diffraction lines width increases with time. The changes in structure parameters are summarized in Table 2. Rwp stands for the goodness of the fitting using Rietveld Method. Both a and c parameters of MmNi5 increase as milling time increases from 0 h to 4h reaching relative change values of 0.20% and 0.22%. Ni parameter a also increases a 0.25% in the same milling period. At approximately the same integrated milling time of 5h, the behavior of MmNi5 a-parameter is different for that reported in LaNi5 [17-18] and multisubstituted LaNi5 [19]. In the first case, a parameter decreases with milling an c parameter remains approximately the same [17]. In the second and third case, a parameter decreases slightly while c parameter increases with milling [18-19]. Table 2 Milling time (h) 2 4 8 10 20 60

MmNi5 a c V (Å) (Å) (Å3) 4.915 3.984 83.36 4.917 3.986 83.46 4.911 3.984 83.23 4.917 3.989 83.53 4.907 3..993 83.28 4.888 4.001 82.83

Table 2. Structure percentages.

Ni Mass % a MmNi5 Ni (Å) 3.538 95.0 5.0 3.539 94.0 6.0 3.536 95.0 5.0 3.540 94.0 6.0 3.537 93.0 7.0 3.524 95.0 5.0

parameters

and

Rwp 18 18 18 17 16 16

mass

This indicates that dumbbells formation does not occur at this stage in this work. Its presence would lead to the decrease of a parameter and an increase of c parameter [19]. The formation of vacancies in crystalline structure can not be responsible for these changes at this state since it would lead to the a decrease in both parameters [19]. Although atomic site interchange between A (La, Ce, Nd, Pr) and B (Ni) atoms could be responsible for the increment in cell parameters, it could not seem to be achieved in this work, either because of the amount of energy supplied or the intrinsic characteristics of this low energy mill.

CONAMET/SAM-2006 Then, a reduction in the long range order can be the explanation for the observed increment in a and c parameters. It is related to the fact that damage affects more to small particles resulting in a larger strain density [17]. As a result, an increment on the lattice parameters is observed. Figure 3 displays a zoom of the 40 -46 º 2θ range of the diffractograms shown in Figure 2.

Figure 4. Samples obtained at different integrated milling times. a) 20 h. b) 60 h. c) 200 h.

Figure 3. Detail of the 40 – 46 º 2θ range of diffractograms shown in Fig. 2. A displacement towards lower angles is observed for (020) and (002) MmNi5 diffraction lines. It is also a proof that a and c parameters of this alloy increases with milling time. The dashed lines in figure are shown as a “guide to the eye”. There is a slight asymmetry in the MmNi5 peaks indicating that milling produces plastic deformations and stacking faults. The mechanical milling also seems to proceed isotropically, since there is no preferential direction affected. 3.3. Mechanical milling at longer periods of time Figure 4 shows the diffractograms of MmNi5-Ni mixture withdrawn at milling times longer than 20 h. Although the determination of the cell parameters is difficult because of the increase of the width of the peaks, the values obtained in Table 2 indicate a clear tendency towards the decrease of a parameter and the increase of c parameter. It implies a change in microstructure due to milling. This behavior coincides with these reported of similar alloys [1819]. Although atomic site-interchange between A and B species was observed in high energy milling devices [19], it could be the explanation for the behavior at long time milling in a low energy milling device such as the one used here.

As milling time progresses, the formation of a bump in the 20-40º 2θ range indicates the progressive amorphization of the alloy. MmNi5 shows a marked asymmetry of the diffraction peaks indicating that stacking faults and directional strains are taking place at this stage. Ni diffraction line is not longer visible at times higher than 10 h and no further analysis regarding to it can be done at this stage. 3.4 Governing processes during milling. Figures 5 and 6 shows mosaics of SEM micrographs of the MmNi5-Ni mixture milled at different integrated milling times and at the same image magnification. As observed in Figure 5, average particle size decrease as milling progresses from un milled samples to those milled 10 h. Faceted faces indicated by white circles show that brittle fracture governs the process in this milling stage. Predominance of this process is observed up to 10 h. Unlike images shown in Fig. 5, cold welding process predominates after milling times longer than 30 h. It is shown in Figure 6 the clear increment in particles size at milling times between 30 h and 200 h. Cold welded particles are identified with white circles. Agglomeration is also clearly shown in Figures 6 c) and d).

Figure 5. Fracture predominance at low milling times. Faceted particles indicating brittle fracture

CONAMET/SAM-2006 are shown in circles. a) Un milled. b) Milled 2 h. c) Milled 10 h.

Figure 6. Cold welding process predominance at milling times higher than 30 h. a) 30 h. b) 100 h. c) 200 h. Cold welded particles are identified with white circles. 3.5. Effect of milling time on MmNi5-Ni properties The amount of the energy supplied related to the integrated milling time evidences changes the particle size and morphology. Identified governing milling mechanisms are successively brittle fracture at times shorter than 20 h, equilibrium between this process and cold welding between 20 h and 30 h and finally an increment in particle size due to cold welding at times longer than 30 h. Milling produces different changes in the MmNi5 structure at times shorter than 10 h and longer than this value. Crystalline parameters increase at short milling times in a behavior no observed previously in similar alloys. It is attributed to a diminution in the long range order. Structure parameters change this tendency as milling time increases. It is attributed to atomic interchange of A and B atoms. Strain introduced in samples treated at different milling times is summarized in table 3. As observed, strain increases with milling time. Strain is not distributed isotropically. If relative values are considered, milling produces a higher effect on (011) diffraction line than in (110). The higher strain as milling time increases is related to the better adsorption and desorption properties evidenced by samples milled at higher times. Table 3 Milling time (h) 2 10 60 100

MmNi5 Diffraction line (011)

Strain (%) 0.049 0.098 0.21 0.23

Diffraction line (110)

Strain (%) 0.038 0.068 0.12 0.13

Table 3. Strain introduced in samples treated at different integrated milling times

3.6 Hydrogen absorption-desorption curves The effects of integrated milling time on the hydrogen absorption desorption characteristics were studied at 25 ºC, 50 and 90 ºC. The effect of integrated milling time on first hydriding-dehydriding curves is selected instead of studying samples activated by successive hydriding-dehydriding cycles. This procedure was preferred in order to discriminate the changes on microstructure produced by mechanical milled from those produced by sample cycling. The first hydriding curves for samples milled 60 h and 200 h at 90 ºC are shown in Figure 7. As observed, the hydrogen adsorption not only is quicker in the sample milled 200 h than in the sample milled 60 h but also the amount of hydrogen absorbed is higher. Similar behavior is obtained at 25 ºC and 50 ºC. It is an indication that milling introduces new fresh surfaces and defects in the material which enhances the interaction with hydrogen. Although it is not studied here, an easier activation would be possible in samples milled for longer times.

Figure 7. First absorption curves for samples treated at different integrated milling times. Figure 8 shows the desorption kinetics of samples treated at different integrated milling times. Desorption rate is quicker in sample milled for 200 h. It is due to the fact that this sample presents a higher amount of strain that than of 60 h as shown later. The behavior of samples desorbed at 25 and 90 ºC is similar to that of Figure 8.

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Figure 8. First desorption curves for samples treated at different integrated milling times. 4. CONCLUSIONS In this work, a MmNi5-Ni mixture was obtained by combined mechanical milling- thermal treatment temperature. After synthesis mixture shows wide diffraction peaks. Therefore, sample was heated at 600 ºC during 5 days in order to release strain and to growth sample crystallite. This sample was used as a reference in order to evaluate successive structural changes due to milling. Structural changes were studied according to milling time. At times shorter than 10 h, an increase in both a and c parameters is observed. This behavior is different from that observed in similar alloys [17-19]. A change in long range order is suggested as explanation. At times longer than 10 h, a decrease in a parameter and an increase in c parameter are observed. A site interchange between A and B atoms is suggested to be the main reason of these structural changes [19]. Although this mechanism was observed at short times in high energy devices [19], it is reasonable to assume that it occurs at higher times in a low energy device such as the one used here. SEM observations were used to study the governing mechanism during milling. At times shorter than 20 h, brittle fracture is found to dominate the main process. It is accompanied by a reduction in particles size. At times higher than 30 h, cold welding predominates and an increase in particle size is observed. First hydriding and dehydriding curves were used to study the effect of milling time on alloyhydrogen interaction. It was found that both processes are quicker as milling time increases. Non activated samples were used in order to isolate the effect of milling on the structure from the cracking produced by cycling the sample during the activation of the sample. 5.REFERENCES [1] G. Sandrock, J. Alloys Compd. , 293-295, (1999), 877-888.

[2] P. Dantzer, Mat. Sc. Eng. A, 329-331, (2002), 313-320. [3] Z. Dehouche, N. Grimard, F. Laurencelle, J. Goyette, T. K. Bose, J. Alloys Compd., 399, (2005), 224-236. [4] Q. Lin, S. Zhao, D. Zhu, N. Chen, Sol. State Ionics, 136-137, (2000), 663-666. [5] W. Genshi, W. Xianglong and S. Panwen, Int. J. Hydrogen Energy, 14, (1989), 45-47. [6] D. Chartoni, N. Kuriyama, A. Otto, V. Güther, C. Nützenadel, A. Züttel and L. Schlapbach, J. Alloys Compd., 285, (1999), 292-297. [7] J.R. Ares, F. Cuevas, A. Percheron-Guégan, Acta Mater, 53, (2005), 2157-2167. [8] G. Liang, J. Huot, R. Schultz, J. Alloys Compd., 320, (2001), 133-139. [9] M. Jurczyk, K. Smardz, W. Rajewski, L. Smardz, Mat. Sc. Eng. A, 303, (2001), 70-76. [10] S. Enzo, E. Bonetti, I. Soletta, G. Cocco, J. Phys. D : Appl. Phys., 24, (1991), 209-216. [11] C. Suryanarayana, Progress in Materials Science, 46, (2001), 1-184. [12] R. A. Young (Ed.), The Rietveld Method, International Union of Crystallography, Oxford Universtiy Press, (1993). [13] “DBWS, 9411 an upgrade of the DBWS programs for Rietveld Refinement with PC and Mainframe Computers”, J. Appl. Cryst, 28, (1995), 366-367. [14] S. Enzo, E. Bonetti, I. Soletta and G. Cocco: J. Phys. D: Appl. Phys, Vol 24, (1991), p. 210. [15] G. Meyer, D. Rodriguez, F. Castro, G. Fernández, J. Bolcich, Proceedings of the 11 th World Hydrogen Energy Conference, Stuttgart, 1996, 1293-1298 [16] M. Esquivel, F.C. Gennari, J.J. Andrade Gamboa, G. Meyer, Jornadas SAM/CONAMET 2005, Mar del Plata, Argentina, 2005. [17] S. Corré, M. Bououdina, N. Kuriyama, D. Fruchart, G. Adachi, J. Alloys Compd., 292, (1999), pp 166-173. [18] G. Liang, J. Huot, R. Schultz, J. Alloys Compd., 320, (2001), pp 133-139. [19] J.R. Ares, F. Cuevas, A. Percheron-Guégan, Acta Mater, 53, (2005), pp 2157-2167. 6. ACKOWLEDGEMENTS The authors wish to thank to the Comisión Nacional de Energía Atómica of Argentina (Project CNEA P5-PID-95-2) , Secretaría de Ciecia, Técnica y Posgrado Universidad Nacional de Cuyo of Argentina, Consejo Nacional de Investigaciones Científicas y Técnicas of Argentina (Project PIP-6448) and Agencia Nacional de Promoción Científica y Tecnológica of Argentina (Project PICT 12-15065) for partial financial support.