Hydrogen Absorption Kinetics of the Transition-Metal-Chloride

Jun 15, 2012 - This corresponds to the simple morphology of Figure 6e. ..... (34) Because it is desirable to reduce the bulk diffusion length/particle...
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Hydrogen Absorption Kinetics of the Transition-Metal-ChlorideEnhanced NaAlH4 System Mark P. Pitt,*,†,§ Per E. Vullum,‡ Magnus H. Sørby,† Hermann Emerich,∥ Mark Paskevicius,§ Craig E. Buckley,§ John C. Walmsley,‡,⊥ Randi Holmestad,‡ and Bjørn C. Hauback† †

Physics Department, Institute for Energy Technology, P.O. Box 40, Kjeller, N-2027 Norway Department of Physics, Norwegian University of Science and Technology, N-7491 Trondheim, Norway § Department of Imaging and Applied Physics, Curtin University of Technology, GPO Box U1987, Perth WA 6845, Australia ∥ Swiss-Norwegian Beamline, European Synchrotron Radiation Facility, BP 220, Grenoble, CEDEX, France ⊥ SINTEF, Materials and Chemistry, N-7465 Trondheim, Norway ‡

ABSTRACT: This study elucidates the role of transition metal (TM) additives in enhancing hydrogen (H) reversibility and hydrogenation kinetics for the NaAlH4 system. The isothermal hydrogen absorption kinetics of the planetary milled (PM) NaAlH4 + xTMCln (TM row 1 = Sc, Ti, V, Cr, Mn, Fe, Co, Ni, Cu; row 2 = Zn, Y, Zr, Nb, Mo, Ru, Rh Pd; row 3 = Pt; 2 < n < 5) and cryo-milled (CM) NaAlH4 + xTMCln (TM row 1 = Ti, V, Cr, Fe, Co, Ni, Cu; 2 < n < 3) systems have been measured at 140 °C and 150 bar system pressure. The variation in hydrogenation kinetics across the TM series for NaAlH4 + xTMCln is strongly dependent on the TM species and additive level, milling technique, and the type, structure, and morphological arrangement of nanoscopic Al1−xTMx phases that are embedded on the NaAlH4 surface. In the most interesting case, the surface-embedded Al1−xTix phases in the TiCl3-enhanced NaAlH4 system perform a dual catalytic function, where the outer Al1−xTix surface performs dissociation/recombination of molecular H2 and the inner Al1−xTix surface allows the distortion of a minor number of Al−H bonds from AlH4− tetrahedra in the vicinity of the subsurface Al1−xTix/NaAlH4 interface. The density of Ti atoms in the subsurface interface (which is Al:Ti composition- and H cycling temperature-dependent) shows the strongest effects on hydrogenation kinetics.

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phere: argon,4,6−8,10,13,15,16,18,19,21 nitrogen,5,9,11,14 or hydrogen;12,17,20,22 starting phase: NaAlH4,4,6,8−11,13−15,20 or NaH + Al;7,12,17,18 and TiCl3 content: 0.9 mol %,4,6 2%,4,6,9−11,13,15,16,19−22 4%,4,6,7,9,12,14,17−19 6%,4,6 9%,4,6 or 10%.8 There is also a wide range of hardness of milling media, including WC,4,6 Si3N4,16 hardened tool steels,19,21,22 high-grade steels,10 and stainless steels,8,9,11−15,17,18,20 which can significantly affect the final milled NaAlH4 particle size. Dependent on these milling conditions, a wide range of TiCl3dependent hydrogenation rates have been reported for TiCl3enhanced NaAlH4, in the range 1.513 to 120 wt % H/h.22 With such variable milling parameters, it is not immediately apparent if the addition of Ti-precursors to NaAlH4 results in ‘typical’ catalytic behavior, where the resultant benefit to kinetic activity scales with Ti content. Furthermore, hydrogenation rates are also dependent on the form of the Ti-precursor, with 2 mol % Ti(OBun)4 enhanced NaAlH4 showing rapid kinetics immediately.23 Hydrogenation rates also show typical dependence on temperature and pressure.23,24 It should be noted that activation energies, Ea, or rate constants related to Na3AlH6

INTRODUCTION The transition metal (TM)-enhanced NaAlH4 system is the prototypical example of catalysis of complex hydrides. For the benchmark Ti-enhanced NaAlH4 system, the earliest report described that the addition of 2 mol % Ti(OBun)4 or TiCl3 to NaAlH4 by wet chemistry-enabled hydrogen reversibility and improved isochronal dehydrogenation kinetics (ca. 85 °C decrease) compared with pristine NaAlH4.1 Absorption studies showed that >80% of hydrogen capacity could be achieved below the NaAlH4 melting temperature (ca. 180 °C) within 1 h at 170 °C and 152 bar.1 H cycling studies showed that over 1 wt % H capacity was lost during the first 35 cycles, predominantly due to a reduction in hydrogenation rate.1 Later studies showed that dry grinding of Ti(OBun)4 with NaAlH4 under an Ar atmosphere improved H release temperatures by ca. 30 °C during isochronal dehydrogenation compared with the addition of Ti(OBun)4 to NaAlH4 by wet chemistry.2 A subsequent study utilized mechanical blending of Ti(OBun)4 with NaAlH4,3 which promoted the most common form of addition of Ti-based additives to NaAlH4, by mechanical ball milling. Reported hydrogenation kinetics from the addition of TiCl3 to NaAlH4 by mechanical milling can be broken down by mill type: vibration/shaker (high energy)4−15 or planetary (lower energy);16−22 milling atmos© 2012 American Chemical Society

Received: March 29, 2012 Revised: June 15, 2012 Published: June 15, 2012 14205

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cycling31,32 do not appear until after the completion of the first thermal desorption. As such, a more representative comparison of candidate TM-salt additives is one that is conducted on at least the first absorption of hydrogen, after crystalline Al1−xTMx phases have formed/stabilized. This Article is organized as follows: Section A presents the hydrogen absorption kinetics of the NaAlH4 + xTMCln system, comparing all first and second period TMs, along with Pt. Section B studies the composite surface-embedded nanoscopic crystalline (c-) Al 1−x TM x + amorphous (a-) Al 1−x TM x morphology that exists for V and Ti, with particular focus on the hydrogenation kinetics of the NaAlH4 + xVCl3 system. Section C follows the decay in hydrogen absorption kinetics of PM NaAlH4 + 0.02TiCl3 during the first 10 H cycles and utilizes H2/D2 isotope scrambling to discriminate which of the rate-determining steps at the surface is slowing the kinetics of hydrogen sorption in the NaAlH4 system after cycling.

and NaAlH4 formation/decomposition have been calculated based on kinetic models of experimental data7,20,25 across different additives of varying content, 4% TiCl3,7 2% TiCl3,20 and 2% Ti(OBun)4,20,25 utilizing vibration7 and planetary mills.20,25 It is noted that the kinetics of Ti-enhanced NaAlH4 appear to be dictated by the NaAlH4 particle size and the ‘state of dispersion’ of the Ti containing species.20 For additives other than Ti-based precursors, a wide range of hydrogenation rates are reported. A study of the desorption kinetics from milled NaAlH4 + xTMCln5 showed that TiCl3 outperformed every other salt from the first and second period TMs. The effects of VCl3, ZrCl 4, and FeCl3 on the dehydrogenation kinetics of NaAlH4 were also shown to be poor compared with TiCln based additives.26 Further studies identified that ScCl310,14,27,28 and CeCl310,22,27,29,30 were more efficient additives than TiCl3. The most recent study22 shows that TiCl3-enhanced NaAlH4 possesses equally rapid kinetics as CeCl3-enhanced NaAlH4 if the starting material of NaH + Al + xTMCl3 (TM = Ti, Ce) is milled under a reactive atmosphere of high-pressure hydrogen. Presently, no salts other than ScCl3, TiCl3, and CeCl3 have been milled with NaH + Al under a hydrogen atmosphere, and it remains unknown if other TM or rare earth (RE) species will be found that are as effective. A study of RECln-enhanced NaAlH4 (milled under Argon)13 indicates that SmCl3 is more effective than both CeCl3 and TiCl3. Aside from the correct choice of the most active salts, of which there are now several choices, such as CeCl3, SmCl3, ScCl3, and TiCl3, the issue concerning the true nature of catalysis for NaAlH4 remains unresolved. If we assume a simple heterogeneous catalysis model (a catalyst that exists in a different phase from the reactants), then the addition of more catalytic material should result in an increase in activity, until the activity reaches a plateau. This assumption appears valid, as the TiCl3-enhanced NaAlH4 system forms Al1−xTix phases that stabilize in composition after several H cycles.31,32 Aside from the increasing hydrogenation kinetics as a function of TiCl3 content,4 such a ‘typical’ catalytic feature has also been reported for the HfCl4-enhanced NaAlH4 system, where the desorption kinetics continue to increase with the addition of more HfCl4.33 This study utilized a centrifugal type of mill, with agate vial and balls. Of particular focus then is the understanding of why such a large difference in hydrogenation rate occurs dependent on milling technique and how best to obtain TiCl3-enhanced NaAlH4 at the lowest mol % of TiCl3 to ensure the least amount of dead-weight NaCl that occurs due to the mechanochemical reduction of TiCl3 with NaAlH4, subsequently consuming valuable hydrogen capacity. Understanding the hydrogen absorption/desorption kinetics of TMCln-enhanced NaAlH4 is of paramount importance for further engineering the NaAlH4 system for practical application,24 and applying the model of catalytic understanding to other tetravalent complex hydrides, such as the borohydride family, which contains significantly higher hydrogen capacities. The first logical step in this process is an experimental review of candidate TM-salts, and one such study has been previously undertaken.5 Similarly, a review of candidate RE-salts has been performed.13 The study of candidate TM-salts5 performed only the first thermal vacuum desorption and did not present any data on absorption kinetics. With the understanding gained for the TiCl3-enhanced NaAlH4 system, it is clear that initially a nanoscopic surface-embedded Al/a-Al50Ti50 composite is formed after the completion of the milling process,34 and the crystalline Al 1−x Ti x phases typically observed after H

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EXPERIMENTAL DETAILS NaAlH4 was purchased from Albemarle (lot no. 22470404-01, >93% purity). All transition-metal-chloride precursors were purchased from Sigma-Aldrich Chemicals (>99.99% purity). At all times, all powders have been handled under an inert Ar atmosphere in a dry glovebox, with 90% of full capacity is achieved in ca. 10 min. The kink in the isotherm during the first minute is due to exothermic self-heating of the sample. Unless otherwise explicitly discussed, all rates in the following discussion are from the ‘fast’ component of the isotherm, which do not include the exothermic self-heating component. Figure 2a−c shows the isothermal absorption kinetics at 140 °C and 150 bar system pressure during the first absorption of hydrogen for period 1 14207

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Figure 3. Isothermal absorption kinetics at 140 °C and 150 bar system pressure for: (a) the first absorption of hydrogen for period 2 transition metals for PM NaAlH4 + 0.02TMCln, (b) period 1, second absorption of hydrogen for PM NaAlH4 + 0.1TMCln, and (c) period 1, second absorption of hydrogen for CM NaAlH4 + 0.1TMCln.

Figure 4. (a) ‘Fast’ hydrogenation rate measured from the isothermal absorption kinetics for PM NaAlH4 + 0.02TMCln (Periods 1 and 2; first absorption), PM NaAlH4 + 0.1TMCln (Periods 1 and 2; first absorption. Period 1; second absorption. Pt; first absorption), and CM NaAlH4 + 0.1TMCln (Period 1; second absorption). (b) Hydrogen storage capacity (wt % H) achieved after 10 h of absorption for PM NaAlH4 + 0.02TMCln (Periods 1 and 2; first absorption), PM NaAlH4 + 0.1TMCln (Period 1; first absorption), and CM NaAlH4 + 0.1TMCln (Period 1; second absorption).

the PM NaAlH4 + 0.02ScCl3 system reaches only 1.94 wt % H, suggesting that: (i) incomplete reduction of the TMCln occurs for Sc and Ti, (ii) Al-rich Al1−xTMx species form, which consume Al that would otherwise be available to NaAlH4 formation, or (iii) a fraction of the Al1−xTMx species forms slightly below the outer powder grain surface or TM atoms have diffused below the outer surface of the Al1−xTMx phase, rendering that part of the surface inactive to molecular H2 dissociation/recombination. Our previous studies31,32 demonstrate that the c-Al85Ti15 solid solution is the dominant Al-rich Ti-containing phase in the H-cycled PM NaAlH4 + xTiCl3 system, and that all Cl atoms are accounted for and complete reduction has occurred. Similarly for Sc, all originally added Cl atoms are accounted for as NaCl,38 with a c-Al87.6Sc12.4 solid solution. The significant proportion of Al-rich solid solutions that have formed demonstrate that point ii is the most plausible explanation for the decreased H capacity for Sc and Ti, consistent with quantitative phase analysis (QPA), showing that the mol % of NaAlH4 created is lower than expected, 57.04 and 79.33%, respectively, for Sc and Ti. Significantly higher quantities of unreacted Na3AlH6 and Al are present for the ScCl3-enhanced NaAlH4 sample, ca. twice as much unreacted Na3AlH6 compared with TiCl3-enhanced NaAlH4 and ca. four times as much unreacted Al. Incomplete formation of NaAlH4

during H cycling has been previously discussed in detail.39 The strong decay in hydrogen capacity across both period 1 and 2 TM species is consistent with the formation of multiple Al containing phases observed.38 Inspection of the variation of absorption rate as a function of TMCln for period 1 TM species in Figure 4a reveals that for PM NaAlH4 + xTMCln, ‘atypical’ catalytic behavior is observed for Sc and Ti, with 2 mol % TMCl3 displaying faster kinetics than 10 mol %. This variation is reversed for V, where 10% VCl3 displays over twice the absorption rate of 2% VCl3enhanced NaAlH4. The remainder of the series from Cr−Zn all display ‘typical’ catalytic behavior, where for all cases 10 mol % TMCln displays faster absorption kinetics compared with 2% TMCln enhanced NaAlH4. Comparing the first and second absorption of hydrogen for PM NaAlH4 + 0.1TMCln also reveals that the absorption rate only changes with cycling for the early TM species Sc−V. For Cr−Zn, the absorption rate changes little during the first two cycles. This suggests that the composite surface-embedded morphology of c-Al1−xTMx + aAl1−xTMx phases observed for Ti and V31 plays a significant 14208

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The FeCl3-enhanced NaAlH4 system appears to be a unique case among all of the TM species studied, where the absorption rate appears almost invariant in the 3 to 4 wt % H/h range, with little change in rate among PM and CM, H cycling, or FeCl3 content. This is consistent with our previous studies showing that all Fe is consumed as an amorphous Al1−xFex phase (average composition of Al88.5Fe11.5 by QPA for 10 mol % FeCl3), embedded on the NaAlH4 surface as 5 mol %) samples dramatically degrades kinetic performance.

The PM and H-cycled NaAlH4 + 0.02TiCl3 sample contains all Ti atoms in kinetically functional crystalline Al1−xTix phases, producing rapid hydrogenation kinetics of ca. 23 wt % H/h, corresponding to the morphology in Figure 6a. Even though the proportion of c-Al86Ti14 in the CM and H-cycled NaAlH4 + 0.1TiCl3 sample holds more Ti atoms in the required kinetically active c-Al1−xTix phase compared with the PM and H-cycled NaAlH4 + 0.02TiCl3 sample, the surface-embedded ‘closed’ morphology between c-Al86Ti14 and a-Al86.5Ti13.5 in the CM sample (Figure 6b) renders the sample effectively inactive. This ‘closed’ morphology is highly undesirable and appears to be only controllable through the choice of milling apparatus that is used. The use of a vibration/shaker (SPEX8000) mill produces significantly faster kinetics for 9 mol % TiCl3-enhanced NaAlH4

Figure 8. (a) Isothermal absorption kinetic profiles for the PM NaAlH4 + 0.02TiCl3 system recorded during the first 10 H cycles. (b) Decay in the ‘fast’ rate for PM NaAlH4 + 0.02TiCl3 absorption kinetics during the first 10 H cycles. 14211

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Figure 9. Isotopic H2/D2 scrambling for a 1 bar mixture at ambient temperature on PM NaAlH4 + 0.02TiCl3 in the desorbed state after (a) the first absorption of hydrogen and (b) the tenth absorption of hydrogen. Panel c compares the HD formation rates over the first 7 days, showing that there is a moderate decrease in the HD formation rate with H cycling.

embedded ‘closed’ morphology at TiCl3 content >5 mol % for PM and H-cycled NaAlH4 + xTiCl3 samples is highly undesirable and likely produces an inactive a-Al1−xTix (x < 0.15)/NaAlH4 subsurface interface. It has been noted in previous studies that for 2 mol % Ti(OBun)4-enhanced NaAlH4, significant reductions in hydrogen capacity and hydrogenation kinetics are observed as a function of H cycling during the first 40 H cycles.1,41 It is presently unclear what the origins of the decreases in capacity and rate are or if they could be related to the presence of Ti-poor a-Al1−xTix (x < 0.15) phases at low mol % Ti(OBun)4 content. From our previous study of the PM and H-cycled NaAlH4 + 0.02TiCl3 system,32 Ti-poor a-Al1−xTix (x < 0.15) phases are absent at the 2 mol % TiCl3 additive level, and we observe excellent hydrogenation kinetics of ca. 23 wt % H/h. (See Figures 2a and 4.) In this section, we study the decay in the capacity and hydrogenation kinetics of PM NaAlH4 + 0.02TiCl3 during the first 10 H cycles. Because Ti-poor a-Al1−xTix (x < 0.15) phases are not present to explain this decay, we follow the evolution of the c-Al1−xTix phases with H cycling for PM NaAlH4 + 0.02TiCl3 and correlate their variation in composition to the hydrogenation kinetics. We also attempt to recover the lost H capacity by annealing to form Al poorer c-Al1−xTix phases. Isotopic H2/D2 scrambling is utilized after the first and tenth H cycles, to discriminate if the decay in capacity and hydrogenation kinetics is due to (i) a loss of Ti active molecular H2 dissociation/ recombination sites in c-Al1−xTix on the outer powder grain surface or (ii) if the subsurface c-Al1−xTix/NaAlH4 interface is becoming Ti-poor with H cycling by consumption of the Tirichest c-Al2Ti phase observed during the early H cycles. i. Decay of Hydrogenation Rate during the First 10 H cycles for PM NaAlH4 + 0.02TiCl3. The absorption kinetic isotherms for PM NaAlH4 + 0.02TiCl3 are shown for the first 10 hydrogen cycles in Figure 8a. The decay in hydrogenation rate as a function of H cycles is shown in Figure 8b. It can be observed that the hydrogenation rate has decayed rapidly by the second absorption to ca. 9.7 wt % H/h and continues to slowly decay to ca. 6.6 wt % H/h by the tenth H cycle. Ca. 0.2 wt % H

capacity is lost by the tenth H cycle. Several possibilities exist that could explain the significant decrease in hydrogenation rate: (i) powder grain sintering has occurred, which has increased NaAlH 4 particle size/diffusion length, (ii) a proportion of the c-Al1−xTix phase(s) has moved slightly subsurface or been covered by another phase, (iii) Ti atoms on the outer surface of the c-Al1−xTix phase(s) have diffused in from the surface, or (iv) Ti atoms at the subsurface c-Al1−xTix/ NaAlH4 interface have diffused away/decreased their density, decreasing the number of Al−H bonds that are affected in AlH4− units in the immediate vicinity of the interface. Point (i) requires particle size measurements to be made by TEM; points (ii) and (iii) will produce a similar effect, a decrease in the molecular H2 dissociation/recombination rate; and point (iv) can be identified by the variation of c-Al1−xTix composition with H cycling. Figure 8b shows the c-Al1−xTix phases that we have identified by TEM and synchrotron X-ray diffraction after H cycles 1, 2 and 10.32 It is immediately apparent that point (iv) is the most plausible, with the Ti-richest phase, c-Al2Ti, present after the first hydrogenation but subsumed by the second H cycle to a significantly Ti-poorer majority c-Al89Ti11 composition. We have measured PM NaAlH4 + 0.1TiCl3 TEM particle size distributions after milling and two H cycles and find no discernible change in the size distributions, indicating that powder grain sintering is highly unlikely during H cycling at ca. 140 °C (nor do we observe powder necking in our TEM images). Figure 9a−c shows isotopic H2/D2 scrambling measurements on PM NaAlH4 + 0.02TiCl3 after the first and tenth H cycles. Samples are measured in the desorbed state to avoid H release from NaAlH4 disproportionation.40 Under a 1.0 bar H2/D2 50:50 molar ratio mixture at ambient temperature, it can be observed that full equilibrium mixing (H2 + D2 ↔ 2HD) has occurred after ca. 7 days. The HD scrambling rate can also be observed to moderately decrease by the tenth H cycle (both HD formation rates are compared in Figure 9c), indicating that the number of active Ti sites for molecular H2 dissociation/ recombination has decreased with H cycling. The HD 14212

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order of magnitude better, with >1.0 Å Al−H bond lengthening possible.41 On this basis, more electronegative elements such as Fe and Mo are proposed to lengthen B−H bonds.46 Whereas Fe and Mo are known to possess efficient surfaces for B2H6 dissociation47 (and thus breaking of B−H bonds), only minor/ negative effects on B−H lengthening are observed for Fe (minor B−H lengthening of 0.05 Å48) and Mo (B−H contraction49) in B-containing molecular structures. Other phases such as metal-oxides (including Fe2O3, Cr2O3, SiO2, Al 2 O 3 , and MgO) are known to be efficient at B 2 H 6 dissociation;47 however, metal-oxides are known to react chemically with borohydrides,45 destroying the original catalytic phase. As such, whereas it initially appears to be unlikely that TM atoms will produce the necessary lengthening of B−H bonds in BH4− units, both DFT modeling and experimental research across the TM rows are required to better assess the viability of nanoparticulate B1−xTMx phases as catalysts for tetravalent complex borohydrides. ii. Annealing the PM NaAlH4 + 0.02TiCl3 System to Improve Hydrogen Capacity. Studies of the NaAlH4 + xTiCl3 system at temperatures >Tm (ca. 180 °C for unadulterated NaAlH4) are sparse in the literature. Two historical reasons to study the system at such high temperatures were (i) the annealing of milled samples to aid in the identification of Al1−xTix phases formed, because until our study,34 the state and location of Ti after milling was unknown (no obvious crystalline Ti containing phases could be observed in diffraction data), and (ii) potentially recovering Al from the majority cAl85Ti15 phase to increase the wt % H storage capacity. With point (i) already extensively investigated,34 in this section, we focus on point (ii). It must be stated clearly here that no molten state/transition is observed at Tm (ca. 180 °C) for TiCl3enhanced NaAlH4 (NaAlH4 is gone well before 180 °C), and only a solid phase transition can be observed up until >300 °C, where NaH reflections start to lose intensity from the diffraction pattern (completely gone by 400 °C31) and vaporous Na is produced, commensurate with physical segregation into individual Al powder particles. Similar physical segregation can be observed for pure NaAlH4 if it is melted at Tm and held for only minutes, producing individual millimeter scaled globules of Al that are physically segregated from NaH particles. On the basis of our in situ annealing data,31 the H-cycled PM NaAlH4 + 0.02TiCl3 system has been annealed after three H cycles at 290 °C under ultrahigh vacuum. Once the sample reached 290 °C, it was then cooled back to 140 °C; it did not remain at 290 °C for more than a few minutes. Figure 10 reveals that the subsequent absorption after annealing has increased capacity to ca. 3.88 wt % H after 12 h of absorption. The increase in capacity is commensurate with the formation of a low d-spacing shoulder on Al (111),31 indicating the formation of a higher proportion of the ordered L12 Al3Ti phase after annealing. Ordering reflections from the D022 and D023 Al3Ti superstructures are evident at >400 °C. This indicates that Al is removed from the majority of c-Al1−xTix composition (c-Al89Ti11 to c-Al85Ti15 dependent on number of H cycles), leaving Al3Ti and subsequently more Al to react with Na3AlH6 to form NaAlH4 and give the extra ca. 0.61 wt % H storage capacity. There has been no significant gain or loss in the absorption rate after annealing compared with the typical rate after three H cycles. Figure 10 also shows that after 6 H cycles the capacity has returned to 3.26 wt % H, similar to the capacity observed at the first absorption, indicating that a larger

formation rate has decreased by ca. 33%. However, this moderate decrease in HD formation rate by itself cannot explain the ca. 70% decrease in hydrogenation rate by the 10th H cycle. As such, the most plausible explanation of the decrease in hydrogenation kinetics for PM NaAlH4 + 0.02TiCl3 with H cycling is a decrease in both (a) active molecular H 2 dissociation/recombination Ti sites on the outer Al1−xTix surface (minor effect) and (b) the Ti atom density in the subsurface c-Al1−xTix/NaAlH4 interface (major effect), as cAl2Ti is subsumed to form c-Al89Ti11, which eventually transforms to a majority c-Al85Ti15 composition by the tenth H cycle. This decrease in interfacial Ti density subsequently decreases the number of Al−H bonds that are affected locally in the immediate vicinity of the subsurface c-Al1−xTix/NaAlH4 interface. The only other TM we have cycled to 10 H cycles is the CrCl3 enhanced NaAlH4 system. We observe a similar decay in hydrogenation rate, implying a similar leaching of Cr atoms from the outer Al1−xCrx surface and the subsurface cAl1−xCrx/NaAlH4 interface with H cycling. The fact that both effects ((a) and (b) above) appear to be necessary to explain the decrease in hydrogenation rate is also sensible from the point of view of Al2Ti conversion to predominantly Al89Ti11; both the inner and outer surfaces and the interior of the Al2Ti phase are leached of Ti as it becomes Al-richer. The bifunctional catalytic nature (H2 dissociation/recombination and Al−H bond distortion) of crystalline Al1−xTix phases we observe in this study shows that distortion/breaking of Al−H bonds at the subsurface Al1−xTix/NaAlH4 interface is the major influence upon hydrogenation kinetics. The destabilization/lengthening of Al−H bonds as a function of transition-metal has been recently investigated utilizing density functional (DFT) models.41 Whereas the models in this study41 are not physically realistic (simple mono transition-metal substitution into the NaAlH4 unit cell does not occur for Ti on a modeling42,43 or experimental basis31,32,34,38,40,44), the maxima in Al−H distortion occurs around Sc and Ti and can be directly correlated with the maxima in absorption kinetics for H-cycled PM NaAlH4 + 0.02TiCl3 we report here. (See Figure 4a.) Such a correlation implies that the DFT model of Al−H lengthening41 is representative of how TM atoms affect local AlH4− units in the vicinity of the subsurface Al1−xTMx/NaAlH4 interface and is consistent with our HD scrambling study, showing that the major effect on degradation of hydrogenation kinetics for H-cycled PM NaAlH4 + 0.02TiCl3 occurs due to Ti loss at the subsurface Al1−xTix/NaAlH4 interface. Noting the bifunctional catalytic nature (H2 dissociation/ recombination and Al−H bond distortion) of crystalline Al1−xTix phases for Ti-enhanced NaAlH4, the analogous catalysts for tetravalent borohydride phases (such as LiBH4) are isolated surface-embedded nanoparticles of transition-metal boride (B1−xTMx) phases. These surface-embedded B1−xTMx nanoparticles must dissociate/recombine molecular H2 and also distort/break B−H bonds at the subsurface B1−xTMx/LiBH4 interface, and both functions must allow H release below the borohydride melting temperature. Catalytic phases must be added as nanoparticles due to the high reactivity of metalchlorides (Cl− substitution into the parent borohydride structure) and metal-oxides with borohydrides.45 An analogous DFT study of the effects of TM atoms on B−H bond lengthening is therefore required. DFT investigations on the effects of mono substitution of Ti into the LiBH4 structure show minimal 0.02 to 0.04 Å B−H bond lengthening.46 In comparison, the effects of Ti on Al−H bonds in NaAlH4 are an 14213

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Figure 11. Comparison of isothermal absorption kinetic rates for vibration milled (VM) and planetary milled (PM) NaAlH4 + xTiCl3 samples. VM conducted at 125 °C/88−79 bar;4 PM conducted at 140 °C/150 bar.

Figure 10. Isothermal absorption kinetics for PM NaAlH4 + 0.02TiCl3 after annealing the sample under vacuum to 290 °C after the third H cycle. No significant gain in the ‘fast’ rate is observed. The improved capacity is attributable to the formation of a higher proportion of cAl3Ti, which in the subsequent two H cycles is destabilized.

proven to be the kinetically functional phase during the very first thermal decomposition of PM NaAlH4 + 0.1TiCl3,40 whereas Ti-poor a-Al1−xTix (x < 0.15) is shown to be highly detrimental to hydrogenation kinetics in H-cycled PM NaAlH4 + 0.1TiCl3. The degraded kinetics in H-cycled PM NaAlH4 + 0.1TiCl3 occur by virtue of an encapsulating ‘closed’ morphology, where the Ti-poor a-Al1−xTix (x < 0.15) phase lies morphologically behind the active crystalline Al1−xTix phase, producing an inactive subsurface a-Al1−xTix (x < 0.15)/NaAlH4 interface, as described above in Section B. We note that at the 2 mol % TiCl3 additive level for H-cycled PM NaAlH4 + xTiCl3 samples, amorphous Al1−xTix phases do not exist, and we observe only crystalline Al1−xTix phases and observe excellent hydrogenation kinetics of ca. 23 wt % H/h. As such, the question of exactly what amorphous Al1−xTix composition has formed in the VM NaAlH4 + 0.1TiCl3 sample8 is highly relevant. The post-annealed absorption isotherms8 begin to show strong degradation after 175 °C, consistent with the formation of crystalline Al1−xTix phases observed at ca. 177 °C in our isochronal PM NaAlH4 + 0.1TiCl3 annealing data.40 This suggests that initially in the milled state VM and PM samples are structurally quite similar, and initially, it is likely that a-Al50Ti50 has formed in VM NaAlH4 + 0.1TiCl3, which transforms to crystalline Al1−xTix phases in a similar manner to PM NaAlH4 + 0.1TiCl3. However, with the commencement of H cycling, dramatically different catalytic scaling occurs between VM and PM samples, as shown in Figure 11. We note that at rich TiCl3 content >5 mol %, H-cycled PM NaAlH4 + xTiCl3 samples are strongly multiphase, with a mixture of a-Al1−xTix + c-Al1−xTix phases. By QPA, we have determined that the a-Al1−xTix component is Tipoor in composition, with x < 0.15.31 As discussed above in Section B, in the H-cycled state for PM NaAlH4 + 0.1TiCl3 samples, our lack of direct observation of Ti-poor a-Al1−xTix (x < 0.15) by high-resolution TEM (it is always observed by integrated electron diffraction) and the degraded hydrogenation kinetics implies that the Ti-poor a-Al1−xTix (x < 0.15) lies morphologically behind the active c-Al1−xTix phases.

proportion of c-Al3Ti is unstable at ca. 140 °C H cycling temperature. This also indicates that the increasing proportion of c-Al3Ti observed after ten H cycles for PM NaAlH4 + 0.02TiCl332 will eventually stabilize, and an equilibrium between c-Al85Ti15 and c-Al3Ti will occur, with c-Al85Ti15 remaining the majority phase. We note that the TiCl3-rich annealing study of vibrationmilled (VM) NaAlH4 + 0.1TiCl38 stands in contrast with Figure 10. This study8 reports that the hydrogenation kinetics after annealing VM NaAlH4 + 0.1TiCl3 up to 225 °C are strongly degraded. The kinetic degradation has been attributed to the formation of crystalline Al1−xTix phases in the postannealed samples. However, for our H-cycled PM NaAlH4 + 0.02TiCl3 sample, we observe no such kinetic degradation in the postannealed state, and we observe only crystalline Al1−xTix phases at the 2 mol % TiCl3 additive level.32 As such, the conclusion that crystalline Al1−xTix phases are the origin of kinetic degradation warrants close re-examination. This study8 has utilized a high-energy SPEX8000 mill, and as discussed above, vibration/shaker milling (VM) produces ‘typical’ catalysis that scales with TiCl3 content, yielding a linear improvement in kinetic rate as a function of TiCl3 content, as shown in Figure 11. X-ray diffraction data of VM NaAlH4 + 0.1TiCl38 shows a ‘broad reflection’, centered around 40° 2θ (in-between the Al reflections (111) and (200)). The center of this ‘broad reflection’ corresponds to ca. d = 2.25 Å, which lies in the 2.21 to 2.30 Å range, where we have previously observed both Ti-rich amorphous (a-) Al50Ti50 (at d = 2.30 Å in the milled state34 for PM NaAlH4 + 0.1TiCl3) and Ti-poor a-Al1−xTix (x < 0.15) (at d = 2.21 Å in the H-cycled state31 for PM NaAlH4 + 0.1TiCl3). Note that the exact center of the ‘broad reflection’8 cannot be precisely determined without knowledge of the zero offset in the diffraction data. Our previous studies31,40 emphasize that the kinetic performance is highly sensitive to the composition of the amorphous phase. Ti-rich a-Al50Ti50 is 14214

dx.doi.org/10.1021/jp3042654 | J. Phys. Chem. C 2012, 116, 14205−14217

The Journal of Physical Chemistry C

Article

For H-cycled VM NaAlH4 + xTiCl3 samples, this ‘closed’ morphology appears to (i) not exist or (ii) is reversed (‘open’), with excellent hydrogenation kinetics of ca. 35 wt % H/h observed for H-cycled VM NaAlH4 + 0.1TiCl3.4 As we have shown in this study, at the 2 mol % TiCl3 additive level, Hcycled PM NaAlH4 + 0.02TiCl3 samples show only crystalline Al1−xTix phases,32 possess excellent hydrogenation kinetics of ca. 23 wt % H/h, and most importantly, survive hightemperature annealing up to 290 °C, with no degradation of hydrogenation kinetics, as shown in Figure 10. Case (i) above for VM NaAlH4 + 0.1TiCl3 implies that all of the original aAl1−xTix (likely a-Al50Ti50) is converted to crystalline Al1−xTix phases after the thermal decomposition, yielding the morphology in Figure 6a. If this were the case, then the observed kinetic degradation is due to crystalline Al1−xTix phases, as concluded.8 As discussed above in Section C.i, we observe a strong reduction in the hydrogenation kinetics of the PM NaAlH4 + 0.02TiCl3 system with H cycling, attributable to the subsumation and transformation of crystalline Al2Ti during the early H cycles. By the tenth H cycle, the hydrogenation kinetics are reduced by ca. 72%, with only crystalline Al85Ti15 (major proportion) and Al3Ti (minor) evident. A similar reduction is evident in the postannealed hydrogenation rates for VM NaAlH4 + 0.1TiCl3 from ca. 32.6 wt % H/h at 125 °C to ca. 7.2 wt % H/h at 200 °C.8 Case (ii) above for VM NaAlH4 + 0.1TiCl3 implies that not all of the original a-Al1−xTix (likely a-Al50Ti50) is converted to crystalline Al1−xTix phases after the thermal annealing treatment, and that during the thermal treatment a similar transformation occurs to that which is observed for PM NaAlH4 + 0.1TiCl3, with a minor proportion of residual Ti-poor a-Al1−xTix (x < 0.15) remaining.31 In the PM case, the Ti-poor a-Al1−xTix (x < 0.15) encapsulates the active crystalline Al1−xTix phase in a ‘closed’ morphology, as described above in Section B. (See Figure 6b.) To understand why VM NaAlH4 + xTiCl3 shows ‘typical’ catalytic scaling4 (see Figure 11) and how initially rapid hydrogenation kinetics at 10 mol % TiCl3 additive level can be degraded with annealing,8 it is necessary to note that the X-ray diffraction data from annealed VM NaAlH4 + 0.1TiCl3 samples shows multiple-phase c-Al1−xTix + a-Al1−xTix mixtures. As the intensity of the ‘broad reflection’8 in VM NaAlH4 + 0.1TiCl3 (likely a-Al50Ti50) decreases with temperature, a high-angle shoulder on Al (111) develops, consistent with the formation of Ti-poor crystalline Al1−xTix solid solution (x < 0.25), which eventually forms an ordered D022/D023 Al3Ti mixture by ca. 400 °C.31 In this respect, VM NaAlH4 + 0.1TiCl3 behaves in a similar fashion to PM NaAlH4 + 0.1TiCl3 during annealing and will form a composite c-Al1−xTix + a-Al1−xTix morphology. With no obvious differences in the amorphous and crystalline Al1−xTix phases that form, the explanation for the difference in catalytic scaling between VM and PM NaAlH4 + xTiCl3 (see Figure 11) can be found in the H-cycling temperature and the onset of crystalline Al1−xTix formation. At 125 °C, the VM NaAlH4 + 0.1TiCl3 samples8 are ca. 50 °C below the temperature of 177 °C, at which we observe crystalline Al1−xTix formation during isochronal annealing (2 °C/min) of PM NaAlH4 + 0.1TiCl3. We further note that at 140−150 °C isothermal H cycling of PM NaAlH4 + xTiCl3, we observe the development of crystalline Al1−xTix phases after the very first thermal desorption from the milled state. As such, the same behavior for VM NaAlH4 + xTiCl3 can be expected at 125 °C with further H cycling (that was not performed in either of the vibration milling studies4,8), and the ‘typical’ catalytic scaling

curve (Figure 11) will show a decrease in gradient, and the curve will flatten with H cycling. This is also strongly supported with the knowledge that the Ti-rich catalytically active aAl50Ti50 phase which is responsible for the rapid hydrogenation kinetics of VM NaAlH4 + xTiCl3 samples, is thermodynamically unstable compared with the crystalline Al1−xTix phases,50 as discussed in our previous studies.31 In this respect, the suppression of crystalline Al1−xTix formation with H cycling is strongly temperature-dependent and highly dependent on mol % TiCl3 additive level and choice of milling technique. During the early H cycles for VM NaAlH4 + xTiCl3 at temperatures