Edge-functionalized graphene nanoribbon encapsulation to enhance

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Edge-functionalized graphene nanoribbon encapsulation to enhance stability and control kinetics of hydrogen storage materials Liwen F. Wan, Eun Seon Cho, Tomas Marangoni, Patrick Shea, ShinYoung Kang, Cameron Rogers, Edmond Zaia, Ryan R. Cloke, Brandon C. Wood, Felix R. Fischer, Jeffrey J. Urban, and David Prendergast Chem. Mater., Just Accepted Manuscript • Publication Date (Web): 26 Mar 2019 Downloaded from http://pubs.acs.org on March 27, 2019

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Chemistry of Materials

Edge-functionalized graphene nanoribbon encapsulation to enhance stability and control kinetics of hydrogen storage materials Liwen F. Wan,⇤,†,‡,? Eun Seon Cho,⇤,¶,? Tomas Marangoni,§ Patrick Shea,‡ ShinYoung Kang,‡ Cameron Rogers,§ Edmond Zaia,†,k Ryan R. Cloke,§ Brandon C. Wood,‡ Felix R. Fischer,§ Je↵rey J. Urban,⇤,† and David Prendergast⇤,† †The Molecular Foundry, Lawrence Berkeley National Laboratory, Berkeley, CA 94720, USA ‡Materials Science Division, Lawrence Livermore National Laboratory, Livermore, CA 94550, USA ¶Department of Chemical and Biomolecular Engineering, Korea Advanced Institute of Science and Technology (KAIST), Daejeon, 34141, Republic of Korea §Department of Chemistry, University of California, Berkeley, California 94720, USA kDepartment of Chemical and Biomolecular Engineering, University of California, Berkeley, California 94720, USA ?L. F. Wan and E. S. Cho contributed equally to this work E-mail: [email protected]; [email protected]; [email protected]; [email protected] Abstract Hydrogen is a long-term clean energy carrier that enables completely carbon-free energy production. However, practical implementation of hydrogen fuel technologies is restricted due to lack of safe and high-performing storage materials. Here we report Mg

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nanocrystals encapsulated by narrow, bottom-up synthesized graphene nanoribbons (GNRs) as environmentally stable and high-capacity hydrogen storage materials. As an encapsulation medium, GNR o↵er similar functionalities as reduced graphene oxide (rGO) to protect the encapsulated Mg nanocrystals from extensive oxidation, while allowing penetrations of hydrogen. In addition, the GNRs can be edge functionalized to tune the (de-)hydrogenation kinetics, in particular for the processes occurred at the GNR/Mg interfaces. In this work, four di↵erent types of edge-functional groups were introduced into GNRs with the goal of comparing their cycling performances due to edge functionalization. Based on detailed kinetic analysis coupled with firstprinciples calculations, we propose that edge-functional groups can contribute to the reduction of kinetic barriers for surface hydrogen reactions at the interface with the GNR by stabilizing surface defects. Furthermore, the GNR-Mg composite exhibited higher hydrogen storage capacity (7.1 wt.% of hydrogen based on the total composite) compared with current literature while demonstrating long-term air stability. This work suggests that rational design of edge functional groups in graphene derivatives can provide a general and novel paradigm for simultaneous encapsulation and hydrogen storage catalysis in simple metal or complex metal nanocrystals.

As the transition towards renewable energy sources accelerates, the use of reliable and sustainable energy carriers gains critical importance. 1,2 Hydrogen is one such environmentally benign energy carrier — it possesses an intrinsically high specific energy and has the unique potential to realize a carbon-free energy system with no associated greenhouse gas emissions. 3 To practically implement a hydrogen economy, it is imperative to establish a safe and compact hydrogen storage system. Such a system also requires high gravimetric and volumetric capacity, especially for mobile applications such as fuel cell electric vehicles (FCEVs). 4 Magnesium-based simple and complex metal hydrides such as MgH2 and Mg(BH4 )2 have been suggested as candidates for solid-state hydrogen storage. 5,6 Specifically, magnesium hydride (MgH2 ) has attracted attention as a promising hydrogen storage material due to its high gravimetric capacity (7.6 wt.%), fully reversible sorption and des2


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orption mechanism, and low cost. 4,7,8 However, the use of MgH2 for hydrogen storage has been limited by sluggish kinetics, and most critically by the oxidative instability of nanoscale metallic Mg particles. Although support materials can address the instability issue to some degree, they often add significant dead mass that does not participate in hydrogen storage, and long-term stability remains an issue. 9–11 Therefore, we sought a new support material that could o↵er superior capacity and stability over commercially relevant timescales while contributing a minimum of additional mass. Recently, we reported a multilaminate composite of reduced graphene oxide (rGO) and Mg nanocrystals, which achieved a remarkable hydrogen capacity along with excellent airstability and kinetic enhancement for hydrogen sorption. 12 This result stems from selective hydrogen transport in graphene oxide (GO) via intrinsic defects, which enables GO to protect Mg nanocrystals from extensive oxidation while simultaneously mediating hydrogen permeation. 13–15 In this multilaminate composite, the rGO sheets that encapsulate the Mg crystals exhibited two discrete e↵ects, serving not only as a gas-selective membrane, but also as an agent for improvement of the hydrogen sorption kinetics. 12 Further kinetic enhancements were provided by the shortened hydrogen di↵usion lengths in the nanoscale Mg crystals. Based on this paradigm, any graphene derivative that retains a suitable gas permeability through defect sites or intrinsic pores would be a suitable encapsulating matrix for Mg crystals. Moreover, this approach is quite general and may be applied to any other metal or complex hydride. As an extension of our initial study, we herein demonstrate the synthesis and the hydrogen absorption/desorption kinetics of nanodispersed Mg crystals embedded in an alternative encapsulant based on a bottom-up synthesized graphene nanoribbon (GNR) matrix. Bottom-up synthesized GNRs are high-aspect-ratio, quasi-one-dimensional graphitic materials that can be prepared by a controlled polymerization of molecularly defined and tunable small molecule precursors. 16,17 These flexible nanomaterials feature precisely controlled widths of approximately 1 nm and a high edge-to-surface ratio. As an important advan-

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tage over more disordered graphene derivatives, the electronic structure and local chemistry of GNRs can be precisely tuned by introducing heteroatoms or functional groups at predetermined positions in the molecular precursor that come to line the edges of the resulting GNRs. 17,18 In this work, a series of pristine, nitrogen- and fluorenone-doped chevron GNRs are used to prepare hybrid GNR-Mg crystal nanocomposites. These are tested for overall performance and stability, with kinetic analysis performed to understand the key decomposition mechanisms. Furthermore, first-principles simulations are used to elucidate how functionalizing the GNR edges impacts the (de-)hydrogenation process of the GNR-Mg nanocomposite. Similar to our previous work, 12 the nanoscale pores formed in a random GNR network provide a selective permeation path for dihydrogen molecules, while blocking other gas molecules such as oxygen. We propose that atomically thin GNRs are promising candidates to encapsulate the active Mg nanocrystals while minimizing dead mass.

Results and Discussion Synthesis of GNR-Mg nanocomposite Four related types of GNRs with di↵erent edge-functional groups were prepared to investigate their suitability as nanocrystal support materials and the e↵ect of heteroatom incorporation on hydrogen uptake performance of the composite. The synthesized materials included an unfunctionalized chevron GNR (cGNR), nitrogen edge-doped N2 -cGNR and N4 -cGNR as well as fluorenone GNR (f-cGNR) as depicted in Figure 1(a). These GNRs share the same chevron-type backbone structure, allowing for investigation of the role of the heteroatoms. Mg nanocomposites with each type of GNR were synthesized via a Rieke-like reaction as previously reported; 12,19,20 the nanocomposites will henceforth be referred to as cGNRMg, N2 -cGNR-Mg, N4 -cGNR-Mg and f-cGNR-Mg, identified according to the encapsulating GNR used (Figure 1(b)). As determined by synthetic mass balance, the composites contained less than 2 wt.% of GNRs and up to 98% of Mg metal such that dead mass could 4


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Figure 1: (a) Chemical structures of cGNR, N2 -cGNR, N4 -cGNR, and f-cGNR, (b) a schematic cartoon of GNR/Mg composite, (c) XRD pattern of GNR/Mg with as-synthesized and after 6 months air-exposed (color coding: predicted XRD pattern of Mg (red), Mg(OH)2 (green) and MgO (blue)), (d) TEM images of GNR/Mg (Inset of (d): di↵raction pattern where the hexagonal dots are matched to Mg (100)), (e) TEM qualitative elemental mapping of GNR/Mg representing dark field image, carbon and Mg mapping.

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be minimized for hydrogen storage. Regardless of the edge-functional groups present in the GNRs, all composites exhibited excellent air stability, even after 6 months exposure to ambient conditions. X-ray di↵raction (XRD) results confirm that all GNR-Mg composites studied here display an explicit Mg crystalline structure that was unchanged after 6 months in air (Figure 1(c)). Remarkably, these composites boast significantly enhanced oxidative stability relative to the Mg crystals multi-laminated with rGO layers reported in our previous study, which exhibited air-stability over only a 3-month period. 12 The Mg metal state was verified via selected area electron di↵raction (SAED) using transmission electron microscopy (TEM) (Figure 1(d)). In contrast with the previously reported ⇠3 nm Mg nanocrystals encapsulated by rGO sheets, the size of individual Mg nanocrystals in the GNR composite was not readily distinguishable via TEM, despite a distinct Mg crystal structure from SAED. Based on TEM analysis, the Mg crystals appear to form a continuous structure rather than individual particles, while Scherrer analysis of XRD peak width estimates the crystalline domain size at approximately 14-15 nm on average, with little variation between GNR types (Table S1). Presumably, this e↵ect is related to the quantity and location of functional groups on GNR layers, which represent a key chemical departure from composites formed with GO sheets. A limited number of nitrogen or ketone moieties exist only along the edge of these GNR molecules, whereas abundant oxygen functional groups are located on GO sheets throughout the plane. In our synthetic approach, Mg2+ precursors were stabilized by those functional groups, followed by co-reduction. Sites where Mg2+ and the functional composites interacted strongly then acted as nucleation and growth points in a Rieke-like reaction. Therefore, it is possible that because the functional groups were fewer in number and localized only on the edge of GNR layers, growth of larger sheets of polycrystalline Mg domains was promoted, in direct contrast to the previous rGO-Mg composite. Further support for this hypothesis is given by scanning transmission electron microscopy-energy dispersive X-ray spectroscopy (STEM-EDX) mapping, which shows that Mg crystals are uniformly distributed over GNR layers as shown in Figure 1(e). Once the

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GNR-Mg nanocomposites are successfully synthesized, the existence of GNR layers (at minimum, only a single atomic layer of GNR is required) e↵ectively blocks Mg from interacting: (1) with itself, avoiding self-agglomeration and (2) with the environment, such as oxygen and water, preventing extensive oxidation. Hypothetically, the strength of the GNR-Mg interaction will contribute to the e↵ectiveness of the protection provided by GNR and the overall stability of the GNR-Mg nanocomposite. In this regard, functionalization of the GNR edges o↵ers additional tunability to enhance the interaction between GNR and Mg nanoparticles, thereby creating mechanically more stable composites.

Hydrogen sorption performance of Mg crystals encapsulated by GNRs Hydrogen absorption and desorption properties for each GNR-Mg were tested using a Sieverts PCT-Pro instrument at 15 bar H2 and 0 bar, respectively, and the results are shown in Figure 2(a), (b) (additional absorption/desorption results at di↵erent temperatures are shown in Figure S1.). The formation of MgH2 upon absorption and restoration to Mg upon desorption were confirmed by XRD (Figure 2(c)). Regardless of hetero-atom functional moieties of GNR, all the GNR-Mg nanocomposites readily absorbed hydrogen and exhibited high gravimetric capacity, reaching 7.10-7.29 wt.% — to our knowledge, the best value ever reported for Mg-based hydride materials when calculated based on total material amount (Table 1). The absorption and desorption patterns varied comparatively little for GNRs with di↵erent functional groups. However, the f-cGNR-Mg and N4 -cGNR-Mg composites showed slightly better sorption kinetics. In the following discussions, we focus on N2 -cGNR-Mg and N4 -cGNR-Mg composites as representative cases, since these two have the same functional group (pyridine) at the cGNR edges yet show distinct hydrogen sorption performance. To closely study the kinetics of these cGNR-Mg composites, the activation energies for hydrogenation/dehydrogenation reactions were computed by fitting the rates of hydrogen absorption/desorption to the Arrhenius law, and the acquired results are presented in Figure 3 7



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Figure 2: (a) Hydrogen absorption of each composite at 200 C and 15 bar H2 , (b) hydrogen desorption of each composite at 300 C and 0 bar, (c) XRD spectra of GNR-Mg after hydrogenated and dehydrogenated (the bottom bars represent an XRD pattern of Mg (red), MgH2 (pink), Mg(OH)2 (green) and MgO (blue))

Table 1: Maximum gravimetric hydrogen storage capacities of all composites. Absorption

Desorption

N4 -cGNR-Mg

7.28

-6.80

N2 -cGNR-Mg

7.15

-6.91

cGNR-Mg

7.10

-6.73

f-cGNR-Mg

7.29

-7.01

rGO-Mg1

6.50

-6.12

1

From Ref. 12

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(N2 -cGNR-Mg and N4 -cGNR-Mg) and Figure S2 (cGNR-Mg and f-cGNR-Mg). Conventionally, simple models for a nucleation and growth process, such as the Johnson-MehlAvrami equation, have been employed to explain the hydrogenation/dehydrogenation mechanism. 21,22 However, in previous studies it has been found that one single model proved unsuccessful in covering the entire reaction spectrum, implying a complicated energy barrier with a combination of di↵erent mechanisms. 12,22 Therefore, the temperature-dependent hydrogen absorption/desorption rates — identified as the rate of change in absorbed/desorbed hydrogen amount — were directly fitted to the Arrhenius law without any assumptions about the hydrogenation/dehydrogenation mechanism, and the results are shown in Figure 3 and S2 (R2 correlation values for the fits are also shown as an assessment of fit accuracy). Upon absorption, all cGNR-Mg systems show similar activation energies (Ea ⇠ 90 kJ·mol 1 ) over the entire range of 1-5 wt.% H2 , which implies that the overall H2 uptake kinetics is governed either by a single rate-limiting step or by competing rate-limiting steps that have the same activation energies. The observed activation energies are similar to barriers in previous studies of H2 dissociation on a Mg surface (Ea = 86.8 – 115.8 kJ·mol 1 ) 23–28 and of H solid-state di↵usion through MgH2 (Ea = 95.5 kJ·mol 1 ). 29 Therefore, we expect that either surface or bulk processes may actively contribute to the overall hydrogenation kinetics of the GNR-Mg composites. However, the decay of the absorption rates with wt.% H2 (Figure 3, panels a and e), despite the relative consistency of the uptake barrier across this same range, appears to be most consistent with expectations from non-steady-state di↵usion. This favors bulk di↵usion through MgH2 as the most probable explanation for the dominant rate limitation during absorption. In addition, functionalizing the cGNR edges seems to have minor impact on the overall hydrogenation kinetics, likely due to the unchanged bulk behavior. In contrast to absorption, more pronounced di↵erences are observed for H2 desorption among di↵erent GNR-Mg composites. First of all, the desorption rates (see examples in Figure 3, panels b and f) indicate the existence of at least two kinetic regimes. In particular, the initially high rate (prior to the sharp decline) corresponds to easily accessible hydrogen

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Figure 3: Kinetic analysis of hydrogen absorption and desorption for (a-d) N2 -cGNR-Mg and (e-h) N4 -cGNR-Mg; (a), (b), (e), (f) present the change of sorption rate (defined in terms of wt.% H2 per hour) and (c), (d), (g), (h) show the activation energy, depending on the amount of hydrogenation absorbed or desorbed within each sample.

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that depletes more rapidly than the remainder of the sample. Given that the melting temperature of bulk MgH2 is 327 C, we assume that this easy-to-access hydrogen lies within a melted or disordered near-surface region, which expands into the material at higher temperatures. The transition to a second kinetic regime upon further desorption is associated with decomposition of the remaining material. In addition, faster desorption kinetics are observed for f-cGNR-Mg and N4 -cGNR-Mg compared to the other samples in Figure 2(b), indicating that the specific functionalization can meaningfully a↵ect desorption rates. Focusing on the second kinetic regime (deeper dehydrogenation), the activation energies derived from the Arrhenius equation show distinct profiles in the range of 1-5 wt.% of H2 between cGNR-Mg, N2 -cGNR-Mg and N4 -cGNRMg, f-cGNR-Mg. Notably, these di↵erences can provide clues into the dominant kinetic mechanism for hydrogen release. As shown in Figure 3 (panels d and h), the lower activation energy obtained for N2 -cGNR-Mg (⇠ 140 kJ·mol 1 ) compared to N4-cGNR-Mg (⇠ 160 kJ·mol 1 ) suggests that the former should exhibit a faster desorption rate if one assumes that the rate is determined exclusively by the barrier. However, this clearly contradicts the desorption measurement in Figure 2. Instead, we conclude that the pre-exponential factor in the Arrhenius equation is meaningfully di↵erent for these two composites. Since the pre-factor is largely determined by the collision frequencies, we can expect it is the surface processes that ultimately contribute to the discrepancies observed between the activation energy and the desorption rate. Intuitively, the surface processes — in this case, either H2 formation during dehydrogenation or H di↵usion on the surface — will be a↵ected by any surface modifications induced by GNR-Mg interactions. Therefore, it is critical to examine how the hetero-atoms at the GNR edges a↵ect the surface structure of Mg or MgH2 and consequently regulate the hydrogen sorption kinetics. The next sections explore these e↵ects in detail.

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Interfacial structure of GNR-Mg nanocomposite during (de-)hydrogenation In an e↵ort to gain further insight into the chemical interactions at the interface between the GNR and the surface of the Mg nanocrystal, we carry out density functional theory (DFT) calculations, focusing primarily on how nitrogen functionalization of the GNR edges impacts the (de-)hydrogenation process in the GNR-Mg nanocomposite. Similar behavior is predicted for oxygen edge-functionalization (appended in the supplemental information), consistent with the hydrogen sorption data in Figure 2 (a) and (b). Edge structures of free-standing cGNR We first examine the edge structure of free-standing, pristine and nitrogen-doped cGNRs. Only hydrogen termination is considered at the GNR edges because it is the most experimentally relevant condition during (de-)hydrogenation. We estimate the thermodynamic stability of cGNR using the following equation, 30,31

GH (µH ) =

1 (EGN R 2a

NC Egraphene 2

NH E H2 2

N H µH )

(1)

where a is the length of cGNR, EGN R is the total energy of cGNR, Egraphene is the total energy of graphene within its primitive cell setting (2 carbon atoms per unit cell), EH2 is the total energy of a H2 molecule. NC and NH are the total number of carbon and hydrogen atoms used to model cGNR. µH is the chemical potential of a hydrogen atom that is referenced to a hydrogen molecule in the gas phase at the appropriate absolute temperature and partial pressure, which we approximate in the ideal limit as:

µH2 = H (T )

H (0)

T S (T ) + kB T ln(

P ) P

(2)

Here H and S are the standard enthalpy and entropy of H2 at P =1 bar, respectively. 32 In Figure 4, we plot the formation energy of cGNR as a function of hydrogen chemical

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potential, µH . A window for µH values relevant to (de-)hydrogenation is marked by the vertical dashed lines in Figure 4. It is known that DFT relaxations predict metastable structures corresponding to local energy minima. Here we consider di↵erent C-H bond geometries at the GNR edges and compare their formation energies. The a11 structure refers to single H termination on every carbon atom at the edges. Three a22 structures are tested, in which each carbon is terminated by two hydrogen atoms. In the high symmetry case (a22 -symm), the in-plane mirror symmetry is preserved. In the buckled geometries, this mirror symmetry is broken and the carbon atoms at the edges can move out of plane. The formation energy di↵erence is negligible with respect to buckling, as shown in Figure 4.

Figure 4: Calculated edge formation energy for cGNR. The vertical lines indicate various temperature-pressure conditions that are relevant to the (de-)hydrogenation process. The a11 configuration refers to single hydrogen termination at the cGNR edges and the optimized geometry remains planar. Three a22 configurations are tested, as discussed in the main text, and their optimized structures are shown in the red, green, and blue boxes, respectively. Note here that the energy di↵erence for the blue and green lines are so small that the green line plotted in the figure is e↵ectively covered by the blue line. The insets show the side views of each configuration. The carbon and hydrogen atoms are shown as brown and pink, respectively. At room temperature and 1 bar H2 pressure (indicated by the orange dashed line in Figure 4), a11 and a22 -symm configurations show similar formation energies. The most stable, A lower room-temperature configuration is given by the a22 -buckled structure, with a 60 meV/˚ 13



formation energy compared to the a11 structure. At elevated temperatures, such as the dehydrogenation temperature near 600 K, the hydrogen chemical potential (Eqn. 2) decreases, and as a result, the formation energy of the cGNR increases (Eqn. 1). Specifically, at 600 K and 1 bar, the chemical potential of hydrogen drops -0.156 to -0.38 eV (black dashed line in Figure 4) and the energetically most preferred edge structure is shown to have the a11 configuration. Holding the temperature constant, as pressure increases, the hydrogen chemical potential also increases. The a11 configuration remains most stable up to 15 bar, which is the typical hydrogenation pressure used for the GNR-Mg nanocomposite. Edge structures of freestanding N4 -cGNR When nitrogen substitution is considered at the cGNR edges, it a↵ects the hydrogen termination of the edge carbons, particularly the carbon atoms next to the nitrogen. Following a process similar to Eqn. 1, we examine the free energy of N4 -cGNR as a function of hydrogen chemical potential under various temperature and pressure conditions. By comparing the relative formation energies for di↵erent edge terminations (details in the supporting information), it is found that the a22 configuration is always less stable compared with other configurations over the entire range of hydrogen chemical potentials. This is very di↵erent than the case for cGNR, where the a22 configuration (buckled) shows lower formation energy at room temperature, ambient pressure. The reason is basically due to the unfavorable two-hydrogen termination for the carbon atoms next to the nitrogen at the edges. In the mixed configurations, where the carbon atom next to nitrogen is terminated with single H and the rest of carbon atoms at the edges are -CH2 terminated, the GNR becomes thermodynamically more stable, however, these formation energies are still higher than the a11 configuration in the range of hydrogen chemical potentials of interest. Overall, we expect that N4 -cGNR is -CH terminated upon synthesis and at (de-)hydrogenation conditions.

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Interfacial structure of N4 -cGNR-Mg To study the role of nitrogen functionalization on hydrogenation in the cGNR/Mg nanocomposite, we place the N4 -cGNR next to a three-atomic-layer, hcp [0001] Mg substrate. Interestingly, through interaction with the Mg metal surface, the thermodynamic stability of N4 -cGNR is shifted to favor the two-hydrogen edge termination, as shown in Figure 5(a). Here the formation energy of the graphene sheet is calculated as,

GH (µH ) =

1 (EGN R 2a

EM g(M gH2 )

NC Egraphene 2

NH E H2 2

NH µH

NN µN )

(3)

where EM g(M gH2 ) is the total energy of the Mg (or MgH2 ) substrate in vacuo. At room temperature, ambient pressure, the mixed configuration (single H on the carbon atoms next to nitrogen and two H on the rest of edge carbon atoms, shown as inset in Figure 5) shows 80 meV/˚ A lower energy compared to the a11 configuration. The existence of a Mg substrate can, first of all, create direct chemical interactions with the edge-bound hydrogen atoms that are pointing towards the Mg surface. The charge transfer from Mg (4.43 e per slab from Bader charge analysis 33 ) to the carbon sheet also provides a chance to electronically stabilize the two-hydrogen termination at the edges. As a result of these interfacial interactions, the two-hydrogen termination becomes energetically more favorable up to 600 K. The possible transformation between -CH and -CH2 configurations at the GNR edges is interesting because it implies potential (initial) hydrogen cycling via the GNR edges. We have examined the kinetic barrier towards -CH and -CH2 switching at zero temperature using the climbing image nudged-elastic-band (NEB) method, 34,35 as shown in the supporting information (Figure S11). Starting from the -CH2 configuration, the activation energy to transform into the -CH configuration is approximately 0.5 eV. As the hydrogen leaves the carbon plane, it di↵uses into the Mg layer and stays at the interstitial site on the Mg surface. The activation energy here is approximated at 0 K. At elevated temperatures, we expect the transformation between -CH and -CH2 configurations at the N4 -cGNR edges can occur at

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Figure 5: (a) Calculated edge formation energies for N4 -cGNR next to the Mg (0001) surface. (b) Calculated edge formation energies for N4 -cGNR next to the MgH2 (101) surface. Inset shows the freestanding N4 -cGNR structure with mixed edge configuration. reasonable rates. Once the -CH2 termination changes to -CH, the edge carbon of GNR has the ability to react with incoming hydrogen molecules to revert to the -CH2 state. However, this requires an activation energy of 1.6 eV per H2 molecule. Note here that this activation energy is substantially lower than for splitting a hydrogen molecule on a pristine graphene sheet (> 3 eV), yet it is still much higher than dissociating hydrogen directly on Mg surfaces. Therefore, we do not expect detectable spillover e↵ects on the GNRs. However, the switching between -CH and -CH2 configurations at the GNR edges may continue to occur during cycles, especially when there is a change of temperature, hydrogen partial pressure or structure of the substrate as Mg is hydrogenated. Interfacial structure of N4 -cGNR-MgH2 Once Mg is fully hydrogenated, N4 -cGNR is no longer interfacing with Mg, but instead with MgH2 . This shifts the thermodynamic equilibrium for the edge structure of N4 -cGNR. As shown in Figure 5(b), when N4 -GNR is placed next to the MgH2 (101) surface, it always prefers single-H termination across the entire temperature and pressure range of interest. Combined with the observation at the N4 -cGNR-Mg interface, we expect that during sample 16


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preparations of the N4 -cGNR-Mg nanocomposite, the carbon edge of N4 -cGNR is -CH2 terminated, except for the carbon atoms next to the nitrogen groups. Upon hydrogenation, the edge of N4 -cGNR can become -CH terminated by overcoming certain kinetic barriers associated with the C-H dissociation. The dissociated hydrogen atom di↵uses into the Mg layer to hydrogenate Mg. Towards the end of hydrogenation, the edge of N4 -cGNR will be fully -CH terminated. The -CH/-CH2 switching is then reversed during desorption of N4 cGNR-MgH2 . At the beginning of desorption, N4 -cGNR edges remain -CH terminated and towards the end of desorption, when metallic Mg starts to dominate the surface, the edge of N4 -cGNR will become -CH2 terminated to reach the thermodynamically more stable state.

Role of heteroatom towards hydrogenation The interfacial interaction between N4 -cGNR and Mg not only a↵ects the edge configuration of N4 -cGNR, but also changes the local atomic structure of the Mg surface. As shown in Figure 6(b), the formation of a chemical bond between Mg and the pyridinic N from N4 -cGNR e↵ectively drags the Mg atom out of the plane of the Mg surface. Intuitively, these structural defects created by N4 -cGNR/Mg interaction would have an impact on the hydrogenation process. Here we study H2 dissociation at the interface of N4 -cGNR and the Mg (0001) surface. There are two unique types of Mg atom created by interfacing N4 -cGNR with Mg, which are marked as Mg1 and Mg2 in Figure 6(a,b). Mg1 is coordinated with a single N atom from N4 -cGNR and is moved slightly out of plane. Mg2 is coordinated with two N atoms (bidentate chelation) and is sitting on the Mg surface like an adatom. Due to the di↵erence in coordination number and geometry, these Mg atoms exhibit di↵erent partial charges (+0.4 ⇠ +0.55 from Bader charge analysis) and hence their responses to hydrogenation are also expected to vary. As shown in Figure 6(c), the activation energy is calculated for dissociating a dihydrogen molecule on Mg1 (red curve) and Mg2 (blue curve), respectively. Both of these activation energies are lower than the reference activation energy for splitting dihydrogen on the pristine Mg metal (0001) surface (1.15 eV). We note here 17



that even without GNR, the creation of a Mg vacancy on the Mg (0001) surface will promote H2 dissociation. In Figure 7(a), we compare the activation energy for H2 dissociation on the pristine Mg (0001) surface and on the same surface with a single Mg vacancy. The results show an activation energy di↵erence of ⇠ 0.3 eV, which should lead to a significant di↵erence in reaction rate. We believe that surface imperfections, such as vacancies or adatoms, created or stabilized by interfacing Mg with N4 -cGNR (e↵ectively through the pyridinic N) could be responsible for the increased hydrogen uptake activities of the nanocomposite.

Figure 6: Calculated dihydrogen dissociation energy on two di↵erent Mg atoms, marked as Mg1 and Mg2, at the interface of Mg and N4 -cGNR. (a) is the top view of the optimized N4 -GNR-Mg interfacial structure and (b) is the side view. (c) is the NEB energy profile plotted along the H2 dissociation pathway.

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Figure 7: (a) Calculated activation energy for dihydrogen dissociation on the pristine Mg (0001) surface (red) and on the (0001) surface with single Mg vacancy (black). (b) Calculated potential energy profiles for H2 recombination on the pristine MgH2 (101) surface (red) and the (101) surface with single Mg vacancy (blue). The energy zero at the starting point is arbitrarily chosen for each pathway. In principle, pyridinic N introduced in bulk graphene could also have strong interaction with the Mg substrate. In Figure S12, we construct a series of interfacial models to examine the interactions between defected graphene and Mg, when graphitic N and pyridinic N functionalization are considered. The Mg-N interaction induced by graphitic N functionalization (N substitution) and single pyridinic N functionalization is not strong enough to create Mg vacancies on the Mg (0001) surface. However, the bi-pyridinic N and tri-pyridinic N can e↵ectively extract a Mg atom out of the surface and leave behind a vacancy site. In addition, oxygen functionalization at the GNR edges shows similar e↵ects to create Mg surface defects, as shown in Figure S13, that can benefit the hydrogenation kinetics.

Role of heteroatom towards dehydrogenation Similar to the case of H2 dissociation on defected Mg surfaces during initial hydrogenation, upon dehydrogenation, potential vacancy sites stabilized by the N4 -cGNR–MgH2 interaction can also facilitate H2 recombination. In Figure 7(b), we compare the calculated potential energy profile for H2 recombination on the pristine and defected MgH2 (101) surfaces. The activation barrier for H2 recombination on the pristine MgH2 (101) surface is 1.93 eV. When 19



a single Mg vacancy is created on the surface, the activation energy to extract a hydrogen molecule reduces by 20% to 1.56 eV. Interestingly, after H2 desorption, the undercoordinated Mg on the surface shows the propensity for dissolution/degradation away from the surface (Figure S14(c)). This implies that another advantage of using the wrapping layer is to protect the active material (in this case Mg) from structural degradation during cycles. In summary, based on first-principles simulations, we argue that the surface defects created or stabilized by the interactions between functionalized cGNRs and Mg (MgH2 ) can promote H2 dissociation (during hydrogenation) and H2 recombination (during dehydrogenation). From the kinetics measurement of H2 adsorption, it is possible that both surface processes (H2 dissociation) and bulk processes (H di↵usion) contribute to the overall kinetics, with the latter probably playing the dominant role. Thus, solely modifying the Mg surface does not necessarily lead to enhanced kinetics. However, during desorption, we expect surface processes to dominate the kinetics, meaning subtle changes in the surface structure and cGNR functionalization can lead to non-negligible e↵ects on the overall performance of desorption.

Conclusion We have demonstrated nanocomposites of Mg crystals encapsulated by di↵erent kinds of GNRs for environmentally-stable and high-capacity hydrogen storage materials. Di↵erent types of GNR layers with varied edge-functional groups were exploited to examine the effect of functional groups and defect density on hydrogen absorption/desorption. All types of GNR-Mg composites exhibited high gravimetric hydrogen capacity — at least 7.1 wt.% of H2 in the total composite mass — and displayed excellent reversibility of Mg to MgH2 and vice versa. Additionally, the GNR-Mg composites exhibited an enhanced oxidative stability relative to the previously studied rGO-Mg, coupled with slower hydrogen absorption/desorption kinetics, which could be attributed to the larger crystallite sizes in the GNR-Mg composites

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(⇠ 15 nm) compared to rGO-Mg system (⇠ 3.5 nm). A number of studies have shown that both the hydrogen sorption capacity and kinetics can strongly depend on the particle size of Mg. 22,36,37 The activation energies (86.5-96 kJ/mol for absorption and 140-180 kJ/mol for desorption) for various GNR-Mg composites obtained here are roughly within the range of previous reported values for bulk Mg system (95-130 kJ/mol for absorption and 120-160 kJ/mol for desorption). 22 Note here that due to experimental uncertainties and fitting errors, large variations in the activation energies can be expected from di↵erent reports. Nevertheless, from this work, we found that the hydrogen absorption kinetics weakly depend on surface and interface chemistry, whereas the desorption kinetics can be tuned by the functional groups at the GNR edges. First-principles calculations reveal possible origins of the improved kinetics. In particular, we propose that the bipyridinic N moieties at the GNR edges strongly interact with Mg with the potential to create or stabilize Mg vacancies on the surface of Mg or MgH2 . The existence of these surface vacancies can promote H2 dissociation on the GNR-wrapped Mg surfaces during hydrogen adsorption and H2 recombination on the MgH2 surfaces during desorption. Other hetero-atom moieties can potentially be introduced into graphene derivatives, leveraging e↵ective defect densities and functional groups to optimize the hydrogen storage characteristics. Based on this paradigm, we suggest that it is possible to design optimized composite materials consisting of simple or complex metal nanocrystals and graphene derivatives, which can enable deployment of this technology for the application of energy storage materials and catalysts.

Methods Materials: Bis(cyclopentadienyl)magnesium 99.99+% (Cp2 Mg) was purchased from Strem Chemicals. Lithium foil and naphthalene 99% were purchased from Alfa Aesar and Sigma Aldrich, respectively. Tetrahydrofuran (THF) was distilled before use. Bis(cyclopentadienyl)magnesium 99.99+% (Cp2 Mg) was purchased from Strem Chemicals. Lithium foil and naphthalene 99%

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were purchased from Alfa Aesar and Sigma Aldrich, respectively. Tetrahydrofuran (THF) was distilled before use. Synthesis of GNR: cGNR as well as N2 -cGNR and N4 -cGNR were prepared following a reported procedure.[17] The solution-based bottom-up synthesis of f-cGNR is depicted in Scheme S1. Yamamoto step-growth polymerization of 1 followed by Scholl oxidation of the intermediate polymer poly-1 yields f-cGNR in 35% yield over two steps. IR spectroscopy of the f-cGNR indicated conservation of the signal at 1718 cm

1

attributed to ketone C=O

stretching (Figure S4). More details are described in SI. Synthesis of GNR-Mg nanocomposite: The GNR-Mg composites were prepared in an Ar glove box. GNRs (5 mg) were dispersed in THF (10 mL) and sonicated for 1.5 h under an atmosphere of Ar. Lithium naphthalenide solution was prepared from naphthalene (1.92 g, 0.015 mol) and Li (0.144 g, 0.021 mol) in THF (96 mL). A solution of Cp2 Mg (1.848 g, 0.012 mol) in THF (18 mL) was added to the GNR dispersion and stirred for 20 min. The reaction mixture was added to the lithium naphthalenide solution and stirred for 2h at 24 C. The precipitate was collected by centrifugation (10000 rpm, 20 min), washed with THF, and dried under vacuum. Characterization and Instrumentation: All samples for characterization were prepared in air, except the hydrogen sorption test. Size and morphology of the nanocomposites were investigated using Philips CM300 and JEOL 2100F Transmission Electron Microscope (TEM) at 200 kV. X-ray Di↵raction (XRD) patterns were acquired using a Bruker AXS D8 Discover GADDS X-ray Di↵ractometer with Cu K↵ radiation ( = 0.154 nm). Hydrogen absorption/desorption measurement was conducted using HyEnergy PCT Pro-2000 at 15/0 bar of H2 at 200-350 C. DFT calculations: The structures of GNR and the interfaces of GNR-Mg nanocomposite are predicted using density functional theory. 38,39 Planewave representation of the wavefunction is used, as implemented in the VASP software package, 40,41 to simulate the electronic structure of these extended systems. Projector-augmented-wave pseudopoten-

22


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tials 42,43 are used in all simulations and a kinetic energy cuto↵ of 400 eV is used to truncate the planewave basis set. For free-standing GNRs, an orthorhombic simulation cell of ⇠ 25⇥17⇥20 ˚ A3 is used, in which the GNRs are separated along the a-direction and extended along the b-direction. A vacuum of at least 15 ˚ A is added in the c-direction. A k-point grid of 1⇥2⇥1 is used in these free-standing GNRs simulations and upon relaxation, forces are converged to be less than 5 meV/˚ A. For the simulations of GNR-Mg (MgH2 ) interfaces, a ˚3 is used, where a Mg (MgH2 ) substrate is inserted in the simulation box of ⇠ 25⇥17⇥30 A c-direction. The optPBE-vdW type exchange-correlation functional 44,45 is used to model the interfacial structure of GNR and Mg (MgH2 ), with the aim to capture the vdW interactions between the GNRs and Mg (MgH2 ). The climbing image nudged-elastic-band method 34,35 is used to simulate the activation energies towards H2 dissociation and recombination.

Acknowledgement We are grateful to Prof. T. A. Pascal for performing classical molecular dynamics simulations of the GNR-Mg systems and Dr. T. W. Heo for fruitful discussions. This work was supported by the Hydrogen Materials Advanced Research Consortium (HyMARC), established as part of the Energy Materials Network by the U.S. Department of Energy, Office of Energy Efficiency and Renewable Energy, Fuel Cell Technologies Office, under Contract Number DE-AC02-05CH11231 and DE-AC52-07NA27344. First-principles calculations by L.F.W. and D.P. were supported by a User Project at The Molecular Foundry and use of its compute cluster, managed by the High Performance Computing Services Group, at Lawrence Berkeley National Laboratory, which is supported by the Office of Science of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231. A portion of the computational work and analysis by L.F.W., P.T.S., S.K and B.C.W. was performed under the auspices of the U.S. Department of Energy by Lawrence Livermore National Laboratory under Contract Number DE-AC52-07NA27344. E.S.C was supported by International Energy Joint R&D

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Program of the Korea Institute of Energy Technology Evaluation and Planning (KETEP), granted financial resource from the Ministry of Trade, Industry & Energy, Republic of Korea (No. 20188520000570). The views and opinions of the authors expressed herein do not necessarily state or reflect those of the United States Government or any agency thereof. Neither the United States Government nor any agency thereof, nor any of their employees, makes any warranty, expressed or implied, or assumes any legal liability or responsibility for the accuracy, completeness, or usefulness of any information, apparatus, product, or process disclosed, or represents that its use would not infringe privately owned rights.

Supporting Information Available Synthesis details, additional characterizations of the GNR-Mg nanocomposites and simulation models are included in the supporting information. This material is available free of charge via the Internet at http://pubs.acs.org.

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Edge-functionalized graphene nanoribbon encapsulation to enhance

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